Carbon Fiber - Cours

1.01Polyacrylonitrile (PAN)-basedCarbon FibersA. SHINDO

Fiber Science Laboratory, Hyogo, Japan

1.01.1 INTRODUCTION 2

1.01.2 PROCESSING 2

1.01.2.1 Precursors for CFs 21.01.2.2 Stabilization 31.01.2.3 Carbonization and High-temperature Heat Treatment 51.01.2.4 Surface Treatment 7

1.01.2.4.1 Anodic oxidation 71.01.2.4.2 Plasma treatment 8

1.01.2.5 Ceramic Coatings 8

1.01.3 STRUCTURE 9

1.01.3.1 High-strength CFs 101.01.3.2 High-modulus CFs 10

1.01.3.2.1 Longitudinal sections 101.01.3.3 Structural Parameters and Density 10

1.01.3.3.1 Morphology of fracture surface 111.01.3.4 Radial Heterogeneity 121.01.3.5 Schematic Structure 131.01.3.6 Chemical Composition 14

1.01.4 MECHANICAL PROPERTIES 14

1.01.4.1 Longitudinal 141.01.4.1.1 Elasticity 141.01.4.1.2 Tensile modulus 151.01.4.1.3 Tensile strength 161.01.4.1.4 Compressive strength 18

1.01.4.2 Transverse 201.01.4.2.1 Transverse modulus 201.01.4.2.2 Torsional modulus 21

1.01.4.3 High-temperature Properties 22

1.01.5 ELECTRIC AND MAGNETIC PROPERTIES 23

1.01.5.1 Electrical Resistance and Thermoelectric Power 231.01.5.2 Electromechanical Properties 241.01.5.3 Magnetoresistance 24

1.01.6 THERMAL PROPERTIES 24

1.01.6.1 Thermal Expansion 241.01.6.2 Thermal Conductivity 25

1.01.7 SURFACE PROPERTIES 26

1.01.7.1 Morphology and Surface Areas 261.01.7.2 Functional Groups 261.01.7.3 Surface Free Energy 28

1

1.01.7.4 Wetting Property 281.01.7.5 Reactivity 30

1.01.8 REFERENCES 31

1.01.1 INTRODUCTION

PAN-based carbon fiber (CF) occupies apremier position among high-performance fi-bers for composites. This was released into themarket in the 1960s (Shindo, 1961). The firstdecade of its use may be regarded as a period ofincubation. During the second and third dec-ades, the CF has experienced remarkablegrowth in producing technology, particularlywith respect to tensile strength. It is consideredthat the factors forming the basis of the tech-nology in the early stage are: (i) the inherentnature of PAN fiber, particularly that whichexhibits preferred orientation of molecularfragments or basic structural units at everystage of heat treatment for carbonization(Shindo, 1961, 1964); (ii) the air-oxidizing sta-bilization treatment (Shindo, 1961, 1964); (iii)suppression of length shrinkage of fibers orstretching for increasing or keeping the pre-ferred orientation of molecular fragments orgraphitic layer planes; (iv) elimination offlaws (Johnson, 1969; Sharp and Burney,1971; Moreton and Watt, 1974; Reynolds andMoreton, 1980); and (v) anodically oxidizingsurface treatment for use as a reinforcement(Courtaulds Ltd., 1964). Watt and Johnson(1969), who started their study in autumn1963 (Anon., 1970), pointed out that stretchingfibers during the stabilization step is effectivefor increasing the Young's modulus and tensilestrength of CFs. The growth of the CF industrywas primarily driven by the development of itsapplication in aerospace and sports industries.Especially, high-performance requirement forCFs in the aerospace industry led CF manu-facturers to the development of CFs withmarkedly higher fracture strain during the lat-ter half of the 1980s. In response to the devel-opment of higher performance CFs and theexpectation for further development of theirapplication, there has been also an advance inscientific studies on that structure and mech-anical properties. In this chapter, such proces-sing, structure, and various properties will bereviewed.

As regards references on CFs, there arebooks written by Fitzer (1985), Watt andPerov (1985), Dresselhaus et al. (1988), Donnetand Bansal (1990), and Peebles (1994).

1.01.2 PROCESSING

PAN-based CFs are manufactured by mak-ing tensioned endless yarns or tows of acrylicfibers passed continuously through a series ofzones for stabilization, carbonization, and sur-face treatment. For making HM CFs, the fibersare subjected to high-temperature heat treat-ment just after the carbonization treatment. Afinal treatment temperature is selected accord-ing to desired properties such as tensile strengthand Young's modulus. Thus, high-strength, in-termediate-modulus, or high-modulus type CFsare produced. The surface-treated CFs areusually sized with epoxy resin.

1.01.2.1 Precursors for CFs

PAN fibers used as precursors for CFs aremade of copolymers containing acrylonitrile inexcess of 90%. Pure PAN has a melting point of326 8C (Dunn and Ennis 1970) or a slightlylower temperature (Layden, 1970), while copo-lymer PAN usually has melting points lowerthan that. Such a high melting point makes thestabilizing treatment of PAN fibers possible.Acrylic fibers of copolymer with comonomer-bearing carboxylic groups such as itaconic acid(Grassie and McGuchan, 1972)

or hydroxylic groups such as hydroxymethyl-acrylic compound (Morita et al., 1973, 1972;Takahagi et al., 1986)

wherein X stands for a radical selected from acertain group, and R is a member of anothercertain group, are recommended for use asprecursor fibers. It is assumed that thesecomonomers are incorporated to initiate nitrile

CH2 C

CH2 COOH

COOH

H2C=C X

R CHOH

Polyacrylonitrile (PAN)-based Carbon Fibers2

polymerization (Grassie and McGuchan,1972), very helpful to stabilization.

As regards the thickness of fiber, the thinnerthe precursor fiber, the higher the preferredorientation of molecules along the fiber axisgenerally becomes. The high molecular orienta-tion in precursor fiber generally brings abouthigh tensile modulus and tensile strength ofCFs. Precursor fibers, therefore, are extensivelystretched during the spinning operation

1.01.2.2 Stabilization

The (oxidative) stabilization is indispensablefor converting PAN precursor fibers into fibersthat can withstand the rigors of the carboniza-tion processing and a crucial step in the manu-facture of PAN-based CFs. One of the moststriking effects of the stabilization treatment isan increase in carbon yield, as shown in Figure 1(Shindo, 1961). The stabilization treatment is

usually performed at temperatures between 200(Shindo, 1961) and 350 8C (Mittal et al., 1997a),in air flows for about 30 min or more. Even inair, if acrylic fibers are rapidly heated, drasticexotherm occurs, which results in violent degra-dation of the fiber. To avoid these phenomena,the heating in air must be started at a tempera-ture below that at which the exotherm begins,i.e., around 200 8C. Incorporation of the como-nomers mentioned above into the polymer isalso very effective for checking violentexotherm reactions (Grassie and McGuchan,1972). In addition, it is desirable that the fibersare heated at increasing temperatures at con-trolled rates.

During the stabilization step, polymerizationof nitrile groups and dehydrogenation from themain chains and addition of oxygen to themoccur to form a stabilized structure in the fiber.As regards nitrile polymerization or cyclization,a number of reactions have been proposed,including both intra- and interchain reactions(Bashir, 1991; Peebles, 1994). Intrachain cycli-zation into condensed rings comprising a fewrings, however, seems to have become the ac-cepted mechanism:

Furthermore, it remains an unsolved problemwhether dehydrogenation from the main chainsof PAN precedes or follows the nitrile poly-merization, or occurs at the same time. Nitrilepolymerization is not only for stabilization, butdehydrogenation or addition of oxygen is es-sential for substantial stabilization. Some struc-tures for the stabilized fiber have been proposed(Usami et al., 1990). One of them is the follow-ing (Takahagi et al., 1986):

C

CN

H2

N

H2 H2

H2 H2

CN CN CN

C C C

C C C C

C C C C

N N

CCC

H2C

N

C C

NH

O

N N N N

O

CN

OH

40% 30% 20% 10%

Figure 1 Gravimetric analysis curves in nitrogengas showing an effect of preoxidation with air at200 8C by increasing the yield of CF. *: preox-idized, *: not preoxidized (after Shindo, 1961).

Processing 3

On heating in air, untensioned copolymerfibers which have been stretched during spin-ning shrink to a considerable extent. The fiber isheld, therefore, under tension as high as thefiber strength allows to check the shrinkage inorder to preserve the orientation. In this case,shrinkage stress develops in the fiber, exhibitinga peak at a temperature between 100 and200 8C, and another peak between 200 and300 8C (Warner et al., 1979). According toGupta and Harrison (1997), intramolecular cy-clization initiates in the amorphous parts of thefiber at around 200 8C, and develops at highertemperatures, randomization of the highly or-iented amorphous phase taking place with pro-gression of the stabilization. Figure 2 shows thestructure of PAN chains and a morphologicalmodel of the PAN fiber (Warner et al., 1979).

Intramolecular cyclization and randomizationin the amorphous phase spread to the crystal-line regions at higher temperatures. The secondrise of stress is due primarily to an entropicrecovering force generated by randomizationof the highly oriented amorphous and crystal-line phases in which stabilization modificationsof polymer molecules progress. In particular,randomization of the amorphous phase primar-ily contributes to the stress development,although randomization of crystalline phasemay also contributes.

In the stabilization process, some sort ofmolecular chain scission is unavoidable. Onthe other hand, at least at the stage at whichstabilization has proceeded to some extent,some sort of intermolecular cross-linkingtakes place between cyclized segments, in parti-cular in a fiber of a copolymer where the co-monomer exerts a strong cyclization-initiatingeffect as mentioned previously. This is consid-ered to have the lateral cohesive force of thefiber (Shindo, 1971).

The stabilization reactions are accompaniedby an exotherm. Figure 3 shows a differentialscanning calorimetry curve of a fiber of anacrylic copolymer, methylacrylate, and itaconicacid on heating in air (Gupta and Harrison,1996). At temperatures above 300 8C, it hasbeen considered that the condensed rings, par-tially dehydrated or oxidized, cross-link to eachother to form aromatic condensed rings. Thus,stabilization reactions proceed to completion ataround 380 8C (Mittal et al., 1994; Gupta andHarrison, 1996). The stabilization reactions canalso proceed to an almost full extent at tem-peratures between around 230 and 290 8C.Shrinkage stress when the fiber is held withoutlength change (Hiramatsu et al., 1972; Warneret al., 1979; Takaku et al., 1981) can be used as

Figure 2 (a) Irregular helical structure of PANchains. (b) Morphological model of PAN fibershowing ordered and disordered regions (repro-duced by permission of Elsevier Science Ltd. from

Carbon, 1996, 34, 1427±1445).

Figure 3 Percentage shrinkage at various loadsalong with DSC behavior on heating in air (18Cmin71) (reproduced by permission of ElsevierScience Ltd. from Carbon, 1996, 34, 1427±1445).


an in situ measure of the extent of stabilization,although the oxygen content in stabilized fibershas often been used as a measure.

When a PAN fiber is cooled down from atemperature on the way to stabilization, theshrinkage stress of the fiber decreases, reachinga constant value at 180 8C. From this observa-tion, Ogawa and Saito (1995) proposed a para-meter (SI value) for evaluating the degree ofoxidative stabilization of a PAN fiber fromshrinkage stresses. Figure 4 shows the contentof oxygen bound chemically in the stabilizedfibers as a function of that parameter. A goodlinear relation is seen between them. In Figure 4,an SI value of 100% corresponds to a chemi-cally bound oxygen of 12 wt.%, which has beenproposed as an optimum value of oxygen to becontained in stabilized fibers for high-perfor-mance carbon fibers. Thus, shrinkage stressesare applicable for evaluating the degree of sta-bilization of an operation for the continuousmanufacture of CFs.

Stabilization of PAN fibers requires the long-est processing times among the heating steps ina line for carbon fiber production. Several ap-proaches to stabilization, therefore, have beenresearched. Some of them are methods basedon making thermostable PAN fibers by treatingwith atmospheres containing nitrogen oxide(Morita et al., 1972), sulfur dioxide (Moritaet al., 1972), active sulfur derived by photolysisof carbonyl sulfide, oxygen, and hydrochloricacid gas (Shindo, 1971), and ammonia and air(Bhat, 1990). Other approaches are pretreat-ment with potassium permanganate (Ko et al.,1988; Mathor, 1994), acetic acid or succinicacid solution (Mittal et al., 1997b), solutions

of amines, basic polyol solutions and cuprouschloride solutions (Peebles, 1994), prior to airoxidation.

1.01.2.3 Carbonization and High-temperatureHeat Treatment

Stabilized fibers are heated in order to becarbonized at temperatures in the range 1000±1600 8C in a nitrogen atmosphere under slighttension. During the carbonization process, car-bon atoms or chains in the fibers condense intocarbon material, increasing the carbon contentin the fiber progressively as seen in Figure 5(Shindo, 1971), while the other elements such ashydrogen, nitrogen, and oxygen are eliminatedthrough the evolution of gases (Shionoya et al.,1972; Watt, 1972).

From gas evolution rate±temperature pro-files, the carbonization process can be dividedinto two regions, below and above about700 8C, although this border temperaturemore or less varies depending on the degree ofstabilization of the fiber and on the heatingcondition. In the region below 700 8C, decom-position of stabilized PANmolecular fragmentsoccurs, and gases such as HCN, H2O, CO2, COand NH3 evolve through primarily intra-molecular reactions between functional groups.Intermolecular cross-linking between the frag-ments also proceeds to form condensed benzenearomatic rings through dehydration (Watt,1970) and elimination reactions of hydrogencyanide (Shionoya et al., 1972). This cor-

Figure 4 Tensile strength of CFs heat-treated to1000 8C in nitrogen for 20 min under a tension of1MPa or 10MPa as a function of SI values of thefibers oxidized in air under a constant length. Theoxidation temperatures were 242, 253, or 264 8C(reproduced by permission of Elsevier Science Ltd.

from Carbon, 1995, 33, 783±788).

Figure 5 Effects of preoxidation and carbonizingatmosphere of HCl gas on numbers of atoms, permonomer unit of PAN, of the composing elementsremaining in a PAN fiber. - - - - - : preoxidized, andcarbonized in HCl gas, Ð: preoxidized, and carbo-nized in nitrogen. &: unpreoxidized, and carbo-

nized in nitrogen (after Shindo, 1971).

Processing 5

responds to the phenomena that Young'smodulus and tensile strength start to increaseat around 300 8C (Shindo, 1961; Watt, 1970) or400 8C (Mittal et al., 1997a).

In the region above 700 8C, the followingintermolecular reactions are supposed tooccur, resulting in the formation of polynucleararomatic molecular fragments through theevolution of N2 (Watt, 1970) and HCN gases(Shionoya et al., 1972):

As the heating temperature rises, those frag-ments increase in extent, bonding between thefragments becomes increasingly firm throughsuch reactions, and the carbon atom networkin the fiber becomes more and more dense,while the volume of open pores formed duringthe course of gas evolution decreases fromaround 1000 8C up to around 1200 8C of HTT(Spencer et al., 1970). At temperatures above1200 8C, closed pores remain in the fibers. Thetensile strength and Young's modulus of thefibers keep increasing up to an HTT of around1500 8C or 1600 8C as the carbon atom conden-sation progresses, as shown in Figure 6 (Mittalet al., 1997a).

Although the carbonization treatment isusually carried out in an inert atmosphere, ithas been observed that HCl gas mixed in theatmosphere below 600 8C causes a considerableincrease in carbon yield and tensile strength ofthe CFs obtained (Shindo, 1971).

For CFs to be high modulus, we have tosubject the fibers to a high temperature, suchas 2500 8C or 2800 8C. During treatment athigher temperatures, aromatic layer planes inCFs increase in extent primarily through thecoalescence of neighboring layers. The increasein the extent of layer planes brings about an

Figure 6 Tensile strength and Young's modulusvs. HTT curves (reproduced by permission of Else-vier Science Ltd. from Carbon, 1997, 35, 1196±

1197).

N NN NN N N N

+ nN2

N N N

N NN N

N N N N

N

+ nHCN


increase of preferred orientation along the fiberaxis. Those increases cause an increase in theYoung's modulus of CFs. Thus, the Young'smodulus keeps increasing with the rise in HTT,whereas the tensile strength decreases (Mittalet al., 1997a).

1.01.2.4 Surface Treatment

1.01.2.4.1 Anodic oxidation

Almost all PAN-based CFs produced areused as reinforcements for resin matrix com-posite materials. These have therefore beentreated by anodic oxidation to improve theiradhesion with the resin matrix, followingcarbonization or high-temperature treatment

because anodic oxidation has advantages overother oxidation techniques in rate, uniformity,and controllability of the degree of oxidation.Although a number of electrolytes could beused for the anodic oxidation, alkaline electro-lytes, for example, sodium hydroxide and am-monium bicarbonate, are recommended to beused because the degradation products formedon the fiber surface dissolve in an alkalineaqueous solution leaving no residues, whichsimplifies the fiber washing, unlike in acidicsolution (Shindo, 1983). In acidic solution, thedegradation products remain as residues on thefiber surface, without dispersing or dissolvinginto the solution.

Figure 7 shows the surface of a fiber imme-diately after an anodic oxidation at 3450C g71

in a sulfuric acid solution and washing withwater, and the surface of the fiber washedwith an alkaline aqueous solution and water.It can be seen in Figure 7(b) that the degrada-tion products have been washed off from thefiber surface. The degradation products areconsidered to be polynuclear aromatic mole-cules with carboxylic and hydroxyl groups ontheir peripheries. In this case, the amount ofdegradation products was about 15%, about70% of the total weight loss of the fiber. Ehr-burger and Donnet (1985) have described that,when an HT-type CF is anodically oxidized inan alkaline electrolyte, carbon dioxide isformed together with degradation productswhich are dissolved in the electrolyte. Darken-ing of alkaline electrolyte solution during theanodic oxidation of CFs has been also observedby Kozlowski and Sherwood (1985). Further-more, Wu et al. (1995) have removed thedegradation residue on the surface of nitricacid-oxidized fibers by washing the fibers withaqueous sodium hydroxide.

Furthermore, epoxy matrix composites ofthe HT and HM CFs oxidized in a continuoustreatment line showed that the ILSS for thefibers oxidized in a sodium hydroxide solutionwas higher than for those in the acidic electro-lyte. The lowering of the interlaminar shearstrength (ILSS) of fiber reinforced plasticsseems to be due to the degradation productresiding on the fiber surface in the acidic solu-tion (Shindo, 1983). Rises in ILSS with cellcurrent and surface atomic concentrations ofcarboxylic acid and hydroxyl groups for CFssubjected to anodic oxidation in alkaline solu-tion have been also reported by Harvey et al.(1987) and Alexander and Jones (1995), respec-tively.

With commercial CFs, anodic oxidation in-creases the surface area slightly and the activesurface area markedly, in addition to formingsurface active groups such as carboxylic acid

Figure 7 Surface of Besfight HT anodically oxi-dized in sulfuric acid solution. (a) Washed withwater. (b) Washed with alkaline solution and then

with water (after Shindo, 1983).

Processing 7

and hydroxide. As seen in Table 13, commer-cially used anodic oxidation increases the spe-cific surface area by 3±4% for HS fibers and byabout 20% for a HM fiber, and active sites by70±90% for HS fibers and by 240% for a HMfiber, whereas it scarcely decreases the tensilestrength (Shindo, 1983).

Harvey et al. (1987) have suggested thatcarboxyl±ester groups are formed at edge sitesin a CF surface with anodic oxidation, whereasketo±enol groups are formed on the basalplanes. Nakahara and Shimizu (1992) oxidizedanodically pyrolytic graphite specimens in analkaline electrolyte solution, tetraethylammo-nium hydroxide, and found that on the edgesurface there was no increase in the -COOH/Cratio, whereas the -OH/C ratio increased gra-dually with increasing electric charge up to5000C cm72, and oxidation was limited onlyto the edge surface.

The anodic oxidation of CFs in an ammo-nium bicarbonate electrolyte, furthermore,causes formation of C±N groups in CF surfaces(Proctor and Sherwood, 1983; Bradley et al.,1994). In this case, increasing the concentrationof ammonium ions by saturating the solutionwith ammonia introduces larger amounts ofsurface nitrogen (Kozlowski and Sherwood,1986). Alexander and Jones found that theformation of amide groups (CONH2), in addi-tion to carboxyl, hydroxyl, and aromatic imidegroups, on the surface of AU fiber, Hercules,during an anodic oxidation in ammonium bi-carbonate electrolyte, and have also supposedthat the amide groups result from reactionof the ammonium ion in the electrolyte withthe surface acidic functionalities (Alexanderand Jones, 1996). With low-modulus CFs,Courtaulds LM and AKZO HTA, amide and/or aromatic amines (PhNH2) seem to be formedduring anodic oxidation. The oxidation ofCourtaulds high-modulus CFs also appears tolead to the formation of aliphatic amine andprotonated amine (Bradley et al., 1994), in ad-dition to amide groups. It is further possible tointroduce amino groups onto CF surface byreaction of caboxylic acid groups in the surfacewith tetraethylenepentamine (Pittman et al.,1997).

1.01.2.4.2 Plasma treatment

In addition to anodic oxidation, there aremany methods available to alter the surfacechemistry of CFs in an attempt to promotechemical bonding between the fiber and resinin composite materials. Among them, plasmatreatment is noticeable from the point of viewof treatment time as short as 15±30 s, which is

very close to that needed for anodic oxidationtreatment, and for being usable to introduceamine groups onto CF surfaces. A low-powerair plasma can introduce alcohol and carboxylgroups onto the surface of a HT CF, AmocoT300, and carboxyl groups onto the surface of aHM CF, Hercules HMU, whereas nitrogen andammonia plasmas can form aromatic amines(ÐNH2), andÐC=NH groups, which havethe potential to bond with epoxy resins (Jonesand Sammann, 1990a). It has been pointed outthat those chemical changes mainly occur at theedge site or defects and not on the basal planes(Jones and Sammann, 1990b).

Furthermore, plasma treatment can coat CFswith polymer. For example, treatment withallylcyanide or xylene/air/argon plasma hasbrought about an increase in the tensilestrength of HT fibers, Hercules and Grafil.This increase is probably due to the filling ofcracks and flaws by the plasma-deposited poly-mer on the fiber surface (Dilsiz et al., 1995).

Amino groups react readily with carboxylgroups on the CF surface. With oxidized CFs,amino groups can also be introduced by reac-tion with carboxyl groups on the CFs (Pittmanet al., 1997). Such a treatment is effective forenhancing adhesion of CFs to polyurethanesand epoxy resin matrices.

1.01.2.5 Ceramic Coatings

CFs can also be used as reinforcements forinorganic matrices (for a review see Peebles,1994). The following two cases of compositesparticularly become an issue. First are the com-posites with brittle matrices having a lowerfracture strain than those of the CFs, and sec-ond are those with matrices which form a brittlecarbide layer through reactions with CFs onfabrication. In the former, cracks formed in thematrix under tension will propagate into thefibers if the interfacial adhesion between fiberand matrix is strong enough to withstand theshear stress at the interface. For a CF with abrittle layer of a thickness above a certain value,a similar mechanism acts in crack propagationfrom the layer into the fiber. CFs, when heatedin a flow of a gaseous mixture of TiCl4, hydro-gen, and argon at a temperature from 800 to1000 8C, are coated with titanium carbide(Honjo and Shindo, 1986a). Adhesion of theTiC layer formed in this way with the coreCF is naturally strong. The relation of thestrength of the TiC-coated fiber and the layerthickness accords to the equation, sm! d71/2.This equation comes from the equation,sm= k(2Eg/pd)1/2, where k is a factor related


to the shape of the crack (Ochiai and Mura-kami, 1979), E and g are the modulus of elas-ticity and the surface energy of the CF,respectively, and d is the thickness of thelayer. The strength of the fiber coated withfilms with thicknesses above 10 nm decreaseswith increasing thickness of the TiC layer ac-cording to the above equation. The strength ofthe fiber, however, regains its original strengthwhen the TiC layer is removed from the fiberwith a solution mixture of HNO3 and HF.

CFs can also be coated with SiC. A SiClayer with a uniform thickness, about 0.5 mm,was formed by heating CFs in a flow of agas mixture of monomethyltrichlorosilane,methane, hydrogen, and argon, at a certainhydrogen concentration, at a temperature ofabout 1200 8C (Honjo and Shindo, 1986b;Shindo and Honjo, 1986). With Torayca T300fiber, the SiC coating resulted in a tensilestrength loss of 40%. In contrast to this, thestrength of the fiber coated with carbon layerwas higher by about 20% than that of the barefiber. These results led Honjo and Shindo todevelop the technique of double and triplecoatings.

Thus, a Torayca T300 fiber yarn was con-tinuously coated with a carbon layer and thenwith a ceramic layer such as TiC(1), TiC(2),SiC, or TiN. Figure 8 shows the variation oftensile strength of the CF coated with twolayers with increasing interfacial debondingstrength between the coating and fiber (Honjoand Shindo, 1986b). Further, the double layercoated fiber exhibited a lower debondingstrength than that of the fiber coated with thecorresponding ceramic layer alone.

It is known from these facts that the adhesivestrength between the carbon layer and ceramiclayer is lower than that between the carbonlayer and CF. In fact, the end of the doublycoated fiber pulled out from resin exhibited anelectron diffraction pattern characteristic of thecarbon coating (Honjo and Shindo, 1986b).Murakami et al. (1986) have also observed,by using an X-ray energy dispersive analyzer,that when the aluminum wire reinforced withcarbon +SiC or carbon+SiC+Ti±B layer-coated fiber fractures under tension, a peelingoff occurs at an interface either between thecarbon and SiC layers, or between the carbonlayer and SiC+Ti±B layer. These results sug-gest that the surface of carbon coating is lessactive than that of the bare fiber. In fact, theASA of the fibers coated with a carbon layeralone was only about one-fifth to one-thirdthat of the bare fiber (Shindo and Honjo,1986).

1.01.3 STRUCTURE

CFs heat-treated in the 1000±2500 8C rangeare composed of turbostratic stacks of aromatic(graphitic) layers (basic structural units), shownin Figure 9, and disordered regions includingcross-linking. The sizes of the layers or stacks,and the orientation along the fiber axis,increase as HTT rises. The aromatic layers con-nected longitudinally with cross-linking and/ordisordered regions are wrinkled at HTTs in the

Figure 8 Variation of tensile strength of doublelayer coated CFs by interfacial debonding stress(reproduced by permission of Elsevier Science Pub-lishing Co., Inc. from `Proceedings of the 1stInternational Conference on Composite Interfaces',

1986, pp. 101±107).

Figure 9 (a) Crystal structure of graphite crystal.(b) Structure of turbostratic carbon.

Structure 9

1500±2000 8C range, and are planar at HTTsabove 2000 8C (Oberlin and Guigon, 1988.).

Most PAN-based CFs are unable to graphi-tize. Nevertheless, the regions of large layerplanes, where the cross-linking with neighborsis weak, can be slightly graphitized. In fact, thearcs (112), (103), and (114), which correspondto the planes of the three-dimensional lattice inthe graphite crystal, have appeared in the elec-tron diffraction photographs of CFs heat-trea-ted at 2800 8C and 3000 8C, though they arefaint, in addition to the arc (101) (Shindo,1961a, 1961b). Further, there are microfibrilsin CFs (Johnson andWatt, 1967; Kwizera et al.,1982). The dimension of microfibrils reachesabout 100 nm or more according to types ofCFs.

1.01.3.1 High-strength CFs

TEM studies have showed that HT CFs aremade of very small BSUs (about 10AÊ ) (Oberlinand Guigon, 1988; Dobb et al., 1995), as shownin Figure 10. The studies performed by allpossible modes of TEM show that high-strength type CFs have very small (about100AÊ ) local molecular orientations of BSU(Oberlin and Guigon, 1988). According tothem BSU, in transverse sections, is less than10AÊ in diameter and thickness. The direction ofthe fringes changes continuously. An intricateand entangled structure is thus formed. On thebasis of these observations, in the model ofmicrostructure proposed by Guigon et al.(1984) for a HT fiber, the BSU are associatedin a zigzag form with tilt and twist boundaries.Thus, they form larger wrinkled sheets. Whentwo randomly crumpled sheets of BSU comeclose enough together, either transversally orlongitudinally, bonding occurs in the faultyareas at the places where they touch eachother. Cross-linking atoms, tetrahedral bonds,etc. that are frozen in at the BSU boundariesare responsible for the fiber cohesion. Amonghigh-strength CFs, Lc (20±25AÊ ) or Z (SAD)(35±40AÊ ) do not vary significantly (Oberlinand Guigon, 1988).

1.01.3.2 High-modulus CFs

1.01.3.2.1 Longitudinal sections

HM CFs are made of isometric distortedlamellae (Oberlin and Guigon, 1988). The la-mellae measured in 11 dark-field images havealways been found to be isometric (La\=Lak).002 lattice fringes for HM CFs show stacks ofquasi perfect fringes (Oberlin and Guigon1988). There are also reviews by Jain and Ab-hiraman (1987) and by Johnson (1987) on thestructure of CFs.

1.01.3.3 Structural Parameters and Density

In general, the structural parameters, suchas degree of preferred orientation along thefiber axis and crystallite sizes, increase, and thed-spacing (002) decreases with increasing HTT.The fiber density increases as the HTT rises.Such structural parameters for types of CFs areshown in Table 1 (Kumar et al., 1993; Dobbet al. 1995). The densities of commercial CFsare shown in Table 2. According to Oberlin andGuigon (1988), whatever the fiber examined,Lc002 reaches 30AÊ , and LaTEM (La by TEM)150±200 AÊ at an HTT of 1700 8C. The latter

Figure 10 (a) 002 Dark-field of a longitudinalsection of Toray 300, 40B 2600 221. (b) Schematicrepresentation of BUS preferred orientation. (c) 002Lattice fringe image of Serofim AXT 2666, to becompared to (b) (reproduced by permission ofElsevier Applied Science Publishers Ltd. from Fibre

Science and Technology, 1984, 20, 177±198).


increases from 200±250AÊ at 2000 8C to 250±300AÊ at 2200 8C, and then reaches about 400AÊ

at 2800 8C, whereas Lc002 remains near to 35AÊ

(Oberlin and Guigon, 1988).

1.01.3.3.1 Morphology of fracture surface

According to Vezie and Adams (1990), thereis a distinct difference in the morphology of afracture surface between lower modulus andhigh modulus PAN-based fibers when observedby a high-resolution SEM. Lower modulus fi-bers, AS-4. T-300 and T-40, G40-700, G45-700,and IM-8, show a fracture surface exhibitingrough, rather poorly defined granular textures,with no indication of sheet-like or fibrillarstructures, whereas a HM fiber, GY-70, ismade up of sheet-like structures. A ªtransitionºstructure of very fine sheets is seen in an inter-mediate modulus T-50 fiber. A low magnifica-tion image of this fiber demonstrates thepresence of core and sheath regions, exceptthat in this case, the rough, granular textureof the core and sheath, typical of HT PAN-based fibers, is separated by fine, radiallyaligned sheets. T-50 fiber with such a ªtransi-

tionº structure between the sheet-like structureand the granular texture has intermediate me-chanical properties.

Kogure et al. (1994) also observed a finegranular microstructure on fracture cross-sectional surfaces of Hercules HM3000 andpostcreep (plastically deformed by 16.5% at2310 8C). Some sheet-like characteristics werealso observed to develop in T±50 fiber (Kumaret al., 1993).

High-resolution SEM micrographs of thetransverse sections of T-40 and M-60J fibersalso do not show the development of sheet-likemorphology, but exhibit a particulate morphol-ogy. The morphology of T-40 fiber appears tobe quite inhomogeneous. Careful examinationof a SEM photograph of this fiber reveals theexistence of some very small particles of about20 nm in size. The particle size in M60J isapproximately 40±50 nm, and this fiber alsoexhibits the absence of sheet-like character,though the modulus is higher than that ofGY-70. Both fibers, however, show three-di-mensional order. Good agreement is observedbetween the sheet thickness and crystallite size,Lc, in GY-70 fiber, both having values of about15 nm (Kumar et al., 1993).

Table 1a Structural parameters of various CFs.

FiberZ(8)

d(002)

Lc

(nm)La

(0)La

(90)Three-dimensional

orderSEM

morphology

T-300 35.1 0.342 1.5 2.2 4.1 No ParticlesT-40 30.2 0.343 1.8 3.0 4.6 No ParticlesT-50 16.4 0.3423 5.3 6.5 17.5 Maybe Some sheet-likeGY-70 9.6 0.3396 14.1 12.0 40.0 Yes Sheet-likeAS-4 36.8 0.342 1.8 3.0 4.0 No ParticlesM40J 21.4 0.3427 3.6 5.5 12.5 NoM60J 9.9 0.3411 7.8 7.0 28.0 Yes ParticlesIM8 0.3431 1.9 3.1 5.1 No

Source: Kumar et al. (1993).Z: Full-width at half-maximum of the (002) azimuthal scan, d(002): d-spacing of crystallites, Lc: Crystallitesize perpendicular to graphitic basal plane, La(0) and La(90): Crystallite sizes along graphitic planeperpendicular and parallel to the fiber axis.

Table 1b Structural parameters of various CFs.

Fiber Z(8) Dd Dc La|| La\ Lc

C1 20.1 55.8 33.3 4.6 7.4 5.9C2 36.1 66.0 86.7 3.0 5.3 1.7X 21.2 50.2 40.0 4.8 7.2 4.7Y 22.0 49.9 40.0 4.8 7.0 4.3Z 30.8 57.5 80.0 3.4 5.9 2.6T1000 31.5 62.2 86.7 2.9 5.2 1.7

Source: Dobb et al. (1995).Dd: Intercrystallite disorder, Dc: Intracrystallite disorder, La||: Crystallite length along the a-axis, La\:Crystallite width perpendicular to the a-axis.

Structure 11

Norita et al. (1988) have found that a newseries of Torayca fibers exhibit higher compres-sive strengths than a conventional series ofTorayca fibers with the same tensile modulus,though the new CFs have thinner diametersthan those of the conventional fibers. Theyhave described that the new CFs have a densermicrostructure consisting of smaller and lessoriented graphitic crystallites than the conven-tional CFs with the same modulus.

During the creep test for which a HerculesHM3000 fiber yarn is strained to 140% at2310 8C, the interlayer spacing reduces from3.427 to 3.406AÊ , the crystallite size, Lc, in-creases from 45 to 57AÊ , and the degree ofpreferred orientation increases. The postcreepfiber, furthermore, exhibits faint (112) reflec-tions, though the as-received fiber does notexhibit them (Kogure et al., 1994).

1.01.3.4 Radial Heterogeneity

Within a carbon fiber filament there is aproperty or structure gradient across the radiusof the filament. Some CFs exhibit a profoundskin-core difference, with more highly orientedand larger crystallites in the skin region, andless oriented and smaller crystallites in the core.

Diefendorf and Tokarsky (1975), Chen andDiefendorf (1872) (on Hercules HMS andHMU), Morita et al. (1977) (on Torayca T300and two fibers with higher strength), and Sa-wada and Shindo (1981) (on Torayca T300 andHercules HMS) found that there was a modulusgradient exhibiting a higher modulus in theouter skin region. Sawada and Shindo (1981)also found the radial gradient of density beinglower in the inner region. Guigon and Oberlin(1986) showed that the crystallites in the surface

Table 2 Mechanical and other properties of types of CFs extracted from manufacturers' data sheets.

Mfr. Fibertype

Filamentcount

Filamentdiameter(mm)

Surfacearea

(m2 g71)

Tensilestrength(MPa)

Tensilemodulus(GPa)

Tensilestrain(%)

Density(g cm73)

Amoco T-300 1k, 12k 7.0 0.45 3650 231 1.4 1.76[Thornel] T-40 12k 5.1 0.5 5650 290 1.8 1.81

T650/42 6k, 12k 5.1 0.5 4620 290 1.6 1.78T-50 3k, 6k 6.5 0.45 2900 390 0.7 1.81

Hexcel AS4 3k, 12k 3930 221 1.7 1.79IM4 12k 4138 276 1.5 1.73IM7 6k, 12k 5379 276 1.8 1.77UHM 3k, 12k 3447 441 0.8 1.87

Mitsubishi TR30 3k 3530 235 1.5 1.79rayon TR50 12k 4900 235 2.1 1.80[Pyrofil] MR50k 12k 5490 294 1.8 1.80

SR50 12k 4220 490 0.9 1.88Sigrafil C30 6.8 3000 230 1.4 1.78

C35 7.0 3200 210 1.4 1.8Tenax HTA 1k, 24k 7.0 3950 238 1.5 1.77

UTS 12k 7.0 4800 240 2.0 1.8IMS 6k, 24k 5.0 5500 290 1.9 1.8UMS 12k 4.7 4500 435 1.1 1.81

Toho HTA 3k, 12k 7.0 3920 235 1.7 1.77rayon ST4 12K 7.0 4810 240 2.0 1.78[Besfight] IM600 12k, 24k 5.0 5790 285 2.0 1.80

HM35 12k 6.7 3240 345 0.9 1.79TM40 12k 6.2 3430 390 0.9 1.85UM68 12k 4.1 3330 650 0.5 1.97

Toray T300* 1 ± 12k 7.0 3530 230 1.5 1.76[Torayca] T300J* 3k, 12k 7.0 4210 230 1.8 1.78* available T700S 12k 7.0 4900 230 2.1 1.80from T800H* 6k 5.0 5490 294 1.9 1.81Soficar T1000G 12k 5.0 6370 294 2.2 1.80

M40J* 6k 5.0 4410 377 1.2 1.77M50J 6k 5.0 4120 475 0.8 1.88M60J 3k, 6k 5.0 3820 588 0.7 1.94X665 6k 5.0 3430 637 0.5 1.98M40* 6k, 12k 7.0 2740 392 0.7 1.81

Zoltek Panex33 48k, 320k 7.4 3600 228 1.78[Panex] Panex30 1552 221 1.75


region of PAN-based fibers were much largerthan those in the center of the fibers. Theabove-mentioned results were obtained by test-ing the CFs etched chemically, except by Die-fendorf and Tokarsky (1975).

Raman spectroscopy can also be a usefultechnique for studying the variation of micro-structure across a CF filament (Morita et al.,1986; Huang and Young, 1995). Two first-order Raman bands have been found in varioustypes of graphite and CFs (Tuinstra and Koe-nig, 1970a, 1970b; Morita et al., 1986), namely1580 cm71 and 1360 cm71 bands. The Ramanband at 1580 cm71 has been attributed to theC±C in-plane stretching mode of the graphiteplanes of an infinite crystal. The 1360 cm71

band is thought to be due to the crystal bound-aries of graphite. The ratio of intensities of twobands, I1360/I1580, can be related to the crystalsize, La (Tuinstra and Koenig, 1970a). Further-more, the bandwidth reflects the orientation ofthe graphite layer plane with respect to the fiberaxis (Katagiri et al., 1988).

Huang and Young (1995) thus found fromthe radial variation of the ratio of intensitiesI1360/I1580 obtained from the Raman spectra oflongitudinally sectioned fibers that the crystal-lite size decreases going from the core to theskin in HMS4 fiber. The half-width at half-maximum intensity of the 1580 cm71 band forthe HMS4 fiber also decreased with increasingdistance from the core of the fiber to the skin, asshown in Figure 11. This implies that there is ahighly oriented skin in HMS4 fiber.

Katagiri et al. (1988) measured the I1360/I1580values in the cross-section of a 2500 8C PAN-based fiber, and set forth that the degree ofgraphitization of the fiber were higher both in

the skin and center regions. The profiles of thefull-width at half-maximum of the 1580 cm71

band for the fiber also supported the resultobtained from the band intensity ratios.

Honjo and Shindo (1986c) anodically oxi-dized Torayca fibers T300 and M40 in an aqu-eous solution of sulfuric acid, and found thatthe filament skin was converted to a transpar-ent thin-walled tube of graphite oxide, and thefilament core was left unoxidized. Furthermore,other filaments of the T300 and M40 fiberswere thinned by anodic oxidation in an alkalinesolution, and then oxidized anodically in anaqueous solution of sulfuric acid. The thinnedT300 filaments were tapered without forminggraphite oxide, while the thinnedM40 filamentsconverted to graphite oxide. Such phenomenaseem to correspond to the structural variationalong the fiber radius indicated by Katagiriet al. (1988).

Huang and Young (1995) took TEM micro-graphs and corresponding SAD patterns fromskin and core regions for CFs T50 and HMS4,and estimated the orientation parameters, i.e.,half-width at half-maximum intensity of the(002) arc of the SAD patterns. From thosevalues, they found that the orientation para-meter decreases from 21.78 in the core to 118 inthe skin for T50 and from 17.28 in the core to14.78 in the skin for HMS4.

In longitudinal sections of some high-modu-lus CFs, a continuum decrease of the transverseradius of curvature of lamellae from the exter-nal surface of the fiber toward the center wasobserved by Oberlin and Guigon (1988). Intransverse sections they also observed the pre-sence of a skin-core texture difference in (002)lattice-fringe images for some HM CFs (havingthe largest transverse radius of curvature oflamellae). That is, the majority of graphiticlayers are parallel to the surface, whereas inthe center the majority of layers are stronglymisoriented.

1.01.3.5 Schematic Structure

Several models of the structures of PAN-based low- or high-modulus CFs have beenproposed during the course of structural studiesin the past (Shindo, 1961; Crawford and John-son, 1971; Fourdeux et al., 1971; Diefendorfand Tokarsky, 1975; Bennet and Johnson,1978; Guigon et al., 1984). They suggest a long-itudinal preferred orientaion, sheath±core ar-rangement in longitudinal and transversesections, undulation, and entanglement ofstacks or lamellae of graphitic layers, and thepresence of voids. Figure 12 shows a 002 lattice

CARBON FIBER SKIN/CORE STRUCTURE

T5025

20

15

10

5

00.0 0.2 0.4 0.6 0.8 1.0

P75

HMS4

Distance from Fibre Core, x/R

Ori

enta

tio

n P

aram

eter

(

)oθ h

Figure 11 Variation of a crystallite orientationparameter, yh, across the fiber determined fromthe SAD patterns (reproduced by permission ofElsevier Science Ltd. from Carbon, 1995, 33, 97±

107).

Structure 13

fringe micrograph of CF M50J, cross-section(Deurbergue and Oberlin, 1992).

Figures 13 and 14 show three-dimensionalstructural models proposed by Bennet andJohnson (1978) and Guigon et al. (1984),respectively. In the Bennet and Johnsonmodel, lamellae are entangled with each otherlongitudinally, a part of a lamella belonging toother lamellae and folded parallel to the fiberaxis being caught with their neighbors transver-sely. In the skin region there are stacks of layersoriented circumferentially. Such a texture withpreferred orientation has, in fact, been observedon a lattice-fringe image (Bennet and Johnson,1979). Guigon and co-worker's model for high-strength CFs is made of a set of isometriclamellae, connected edge-to-edge, which foldparallel to the fiber axis. The structural unit inthis model is a stack of aromatic layers with anextension of about 10AÊ on each edge. Theseunits connect to each other to form columnsalong the fiber axis, though the units are invarious orientations of twist both within andoutside the plane of the unit. In the model forHM CFs, the structural units are larger in sizeand higher in degree of orientation.

1.01.3.6 Chemical Composition

Nitrogen contents of CFs Torayca T300 andT900 are 6.3 and 5.4%, respectively. Oxygencontents of those fibers are 0.9 and 0.8%, re-spectively. Their hydrogen contents are below1% (Trinquecoste et al., 1996).

1.01.4 MECHANICAL PROPERTIES

Mechanical properties for various types ofcommercial CFs extracted from manufacturer'sdata sheets are shown in Table 2.

1.01.4.1 Longitudinal

1.01.4.1.1 Elasticity

PAN-based CFs are elastic, but their stress±strain behaviour is nonlinear (Curtis et al.,1968; Jones and Johnson 1971; Hughes, 1986).

Nukushina et al. (1986, 1989) have observedfor several types of PAN-based CFs that thestress±strain curve of a CF filament obtainedafter the loading of 80% of a fracture stress andunloading had been repeated several times isthe same as that obtained before the loading±unloading cycle. Hayakawa et al. (1990) foundthat tangent and sonic moduli for CFs mono-tonously increased with increasing appliedstress.

Okada et al. (1995) have found that the pre-ferred orientation of crystallites of several kindsof CFs with different HTTs increases linearly as

Figure 12 002 Lattice fringe micrograph of CF M50J, cross-section (reproduced by permission of ElsevierScience Ltd. from Carbon, 1992, 30, 981±987).

Figure 13 Schematic three-dimensional representa-tion of structure in high-modulus PAN-based CF.The layer planes are highly interlinked in bothlongitudinal and transverse direction (reproducedby permission of Kluwer Academic Publishers from

J. Mater. Sci., 1983, 18, 3337±3347).


the strain rises up to about 1.3% on loading,but goes back to the original on unloading. Thevariation in scattering peak width of SAXS,giving information on the width of voids, par-allel to the fiber axis with tensile stress also wasalmost reversible.

Curtis et al. (1968) have found that the dy-namic tensile modulus of HT and HM CFsincreases markedly with increasing tensilestress, and that the increase in dynamic mod-ulus accompanies an appreciable increase incrystallite orientation. Shioya and co-workers(Shioya and Takaku, 1994; Shioya et al., 1996)have measured the X-ray diffraction of CFsunder tensile stress and have found that thepreferred orientation increases with increasingstress. Furthermore, they found that the fibercompliance and the rotational compliance as itscomponent decrease with increasing initial or-ientation, though the extensional compliance asthe other component at smaller tensile stress isalmost constant in the orientation range cover-ing most of the CFs currently available.

1.01.4.1.2 Tensile modulus

The Young's modulus of CFs is consideredto be determined by the volume (size) andmodulis and preferred orientation, relative tothe fiber axis, of aromatic layers (graphitic layerplanes) or crystallites, and the volume andmodulus of disordered regions, includingcross-linking. The volume of the disorderedregions is considered to be very small comparedwith that of the crystallites, and also those

moduli would be lower than those of the crys-tallites. Thus, Young's modulus of the CFs canbe described as a function of size of the crystal-lites (in particular, La) (Oberlin and Guigon,1988). Indeed, a good correlation has beenfound between Young's modulus and layer dia-meter for HM-type CFs by Guigon and Oberlin1984).

In CFs, the degree of preferred orientation ofgraphitic layer planes along the fiber axis alsocontrols Young's modulus, and in addition thedegree of orientation as well as crystallite sizesincrease together as the HTT in the preparationrises. Young's modulus, therefore, can also bealmost well expressed as a function of the de-gree of preferred orientation (Ruland, 1969;Fourdeux et al., 1971; Northolt et al., 1991).

Ruland (1969) interpreted the Young's mod-ulus as a function of orientation of graphiticcrystallites composed of wavy graphitic layerplanes by means of an elastic unwrinklingmodel using the intrinsic elastic constants ofcrystallites. Brydges et al. (1969) showed thatthe tensile moduli of CFs were rather close tovalues calculated on the basis of a polycrystal-line model comprising perfect graphite crystal-lites with a measured amount of crystalliteorientation by assuming uniform stress on crys-tallites. On the basis of the uniform stressmodel, Northolt et al. (1991) have also pro-posed an equation interpreting the modulus ofCFs as a function of the shear modulus andorientation of graphitic crystallites. Okada et al.(1995) have pointed out that the observed va-lues of tensile modulus as a function of tensilestrain are markedly different from the valuescalculated based on the uniform stress model.

Figure 14 (a) Model of high-modulus PAN-based CF. (b) Model of high tensile strength PAN-based CF(reproduced by permission of Elsevier Science Ltd. from Carbon, 1992, 30, 981±987).

Mechanical Properties 15

In general, the Raman bands shift to lowerfrequencies under tensile loading and to higherfrequencies in compression. In most cases, anapproximately linear relationship exists be-tween the Raman band shift and the strain onfiber. The Raman band shift per strain in-creases proportionally with fiber tensile modu-lus (Huang and Young, 1995).

1.01.4.1.3 Tensile strength

At the present time, the ratio of tensilestrength to Young's modulus for the commer-cial CFs in Table 2 is in the range 0.005±0.022.Although the ratio for a graphite whisker(Bacon, 1960) is considered to be 0.02, thevalues for many brittle materials are in therange 0.1±0.2 (Johnson, 1987). ToraycaT1000G (Yamane et al., 1987) having the high-est strength, 6.4GPa, among commercial CFs,exhibits a strength-to-modulus ratio of 0.022.Such low values of strength-to-modulus ratioare attributed to the defective nature of CFs.

With the CFs produced earlier, flaws havebeen observed on fracture surfaces (Johnson,1969; Sharp and Burnay, 1971; Moreton andWatt, 1974). Reynolds and Moreton (1980)have reported that strength-limiting flaws de-velop in CFs from impurity particles, particu-larly after heat treatment to 2500 8C. Even inthe CFs produced later, the inhomogeneity inmicrostructure has been observed under TEM,which would act as a strength-limiting element(Dobb et al., 1995).

The strength of CFs, thus, is gauge-lengthdependent owing to a random distribution offlaws or defects (Moreton, 1969). The strength

at a short gauge length as the intrinsic strengthof CFs, then, was evaluated to obtain a realistictarget to aim at in the pursuit of practicalstrength improvement (Diefendorf and To-karsky, 1975). The tensile strength values ofCFs at very short gauge lengths, i.e., criticallengths, were also necessary for determinationof the fiber±matrix interfacial properties inCFRPs by the fragmentation method (Aslounet al., 1989; Hitchon et al., 1979). Suchstrengths were estimated by means of Weibullanalysis and extrapolation by assuming a linearlogarithmic dependence on gauge length of ten-sile strength.

Table 3 shows the tensile strength and Wei-bull parameters for Torayca fibers estimated byMiwa et al. (1991) and R'Mili et al. (1996). Thestatistical variation in mechanical properties ofT300 fiber has been studied by using monofila-ments and loose bundles of 6000 filaments, andthe strength distribution is analyzed by the two-or three-parameter Weibull model (R'Mili et al.1996). The mean tensile strength of CFs in abundle can be determined by treating the fila-ment-strength distribution from the stress-strain curve of the bundle with Weibull statis-tics (Noguchi et al., 1976; R'Mili et al., 1996).Chi et al. (1984) also determined the filamentstrength distribution from fiber bundle testing.

Fiber-resin transfer lengths of 0.5mm(Hitchen and Phillips, 1979) and 0.6mm (No-guchi et al., 1976) have been obtained, andabout 5, 8.5, and 10GPa strength at thatgauge length have been given to ToraycaT300, T800H, and T1000, respectively.

A method for determining the bendingstrength at a very short gauge length is theelastica test (Jones and Duncan, 1971). Table 4gives the results of the loop test for some CFs

Table 3a Statistical values of tensile strength for CFs.

FiberWeibull

modulus, mScale

parameter s0

(GPa)

Locationparameter sp

(GPa)

Mean strengthat 30mm length

(GPa)Length of links

T-300 5.28 5.02 0.502 3.01 0.714M-40 5.35 3.99 0.800 2.36 0.286

Source: Miwa et al. (1991).

Table 3b Statistical values of tensile strength for CFs.

Fiber Sample length(mm)

Weibull modulus,m

Weibull scaleparameter sL

(GPa)

Mean strengthat 30mm length

(GPa)

T-300 30 5.5 3.2 2.95

Source: R'Mili et al. (1996).


(Trinquecoste et al., 1996). In the loop test, theshape of the stressed loop showed a very goodsuperposition with the theoretical curve (calledelastica). The bending strength is obtained ac-cording to the formula st=Ee based on aperfect elastic material and assuming that thetensile and compressive moduli are identical.Compared to a regular single-fiber tensile test,the working volume of the loop-tested sample ismuch smaller. The loop test strength values inthe table, hence, are twice the commercial va-lues.

It is well known that the strength of manykinds of fibers increases with decreasing fiberdiameter, and such a behavior has been alsoobserved for CFs (Shindo, 1961; Jones, 1971;Jones and Duncan, 1971). Figure 15 shows thetensile strengths at a gauge length of 0.1mm forBesfight and Torayca high-strength CFs as afunction of cross-sectional area (Shindo andSawada, unpublished). These fibers wereselected as CFs having Young's moduli, 240±250GPa, being close to each other. Thosestrength values were obtained by Weibull ana-lysis of strengths of about 40 filaments of gaugelengths 5, 10, 20, and 40mm and extrapolation.It can be seen in Figure 15 that the smaller thefiber diameter, the higher the tensile strengthbecomes. As mentioned in Section 1.01.3.4,many types of CFs have radial heterogeneity.Accordingly, it is considered that the relativevolume of the skin region increases and thestructure of CFs becomes more homogeneousas the cross-sectional area decreases. This isconsidered to be one of the causes of the higherstrength of the thinner CF. Extrapolation of theline of least squares in the figure to a fiberdiameter of 1 mm gives a strength of 16.8GPa.Although the slightly lower value seems muchmore likely for the strength in this case, such astrength may suggest the tensile strength offracture at a disordered region or cross-linkingbetween crystallites along the fiber axis in theskin region of the CFs.

The tensile strength of PAN-based CFs in-creases with HTT, reaches a maximum ataround 1500±1600 8C (depending to operatingconditions), and finally decreases more slowly

at an HTT higher than 2500 8C (Figure 6). Thedecrease in tensile strength is attributed to adecrease in the amount of disordered regionlinking the ordered or crystalline regions (John-son and Tyson, 1970; Mittal et al., 1997a). Asmentioned in Section 1.01.3.3, high-strengthtype and intermediate-modulus type CFs ex-hibited rough, rather poorly defined granulartexture in high-resolution SEM images of ten-sile fracture surfaces. Thus, it is considered thatthe initiation of a fracture in such fibers prob-ably occurs in longitudinal bonds of a disor-dered region between two crystallites, itspropagation being transversal. On the otherhand, some HM CFs showed fracture surfacesexhibiting sheet-like textures. In such CFs, ten-sile fracture may be initiated at a flaw on la-mella or a pinched-in region of lamella. Theinitiation at a flaw on lamella has been ex-plained with the mechanism proposed by Ben-

Table 4 Mechanical properties of CFs from loop tests at RT.

T300 T700S T900 M40

Loop test straina % 2.78 5.04 4.18 1.38Tensile modulusb GPa 230 230 294 392Tensile strengthb GPa 3.50 4.80 5.40 2.70Loop test strengthc GPa 6.4 11.6 12.3 5.4

aMeasured. bFrom literature. cCalculated. Reproduced by permission of Elsevier Science Ltd. fromTrinquecoste et al. (1996).

Figure 15 Tensile strength at a gauge length of0.1mm for Besfight and Torayca high-strength CFsas a function of cross sectional area. B: Besfight. T:Torayca. Subscripts of u and s stand for the fiberswithout and with surface treatment, respectively.The line in the figure represents an equation of leastsquares, y=16.614±0.2766x, where y and x aretensile strength and cross-sectional area, respec-tively (after Shindo and Sawada, unpublished).


nett and Johnson (1983) using the Reynoldsand Sharp (1974) criterion. Their proposal isbased on the fact that the internal and surfaceflaws that initiated failure showed evidence oflarge misoriented crystallites in the walls ofholes. According to Reynolds and Sharp, frac-ture of a CF is caused by basal plane rupture incrystallites at large misorientation angles. Theirmechanism is explained in Figure 16.

1.01.4.1.4 Compressive strength

Several test methods such as tensile recoil(Allen, 1987), unidirectional composite, bro-

ken fiber fragment length, loop, and fiberencapsulated into block tests have been usedfor determining the compressive strength ofCFs (Kozey et al., 1995). The single filamentrecoil technique is applicable to fibers with alower compressive strength than tensilestrength (Allen, 1987). In the broken fiberfragment length method, axial tensile or com-pressive stress is applied to a filament bybending a rectangular resin beam in whichthe filament is embedded near the beam sur-face. The mean tensile strength of the fiber at alength, and mean fragment length in tensionand compression, are used in the calculation ofcompressive strength (Ohsawa et al., 1990;Miwa et al., 1991).

Compressive strengths for various types ofCFs are shown in Table 5. These compressivestrength values were determined from the com-pressive strength of composites by normalizingto 100% fiber (Kumar et al., 1993). As seenthere, the compressive strengths even for thesame fiber are considerably different accordingto the test method or author. The same ten-dency can also be seen in a table by Peebles(1994).

The results of compression tests dependstrongly on the loading geometry and test con-ditions. Methods such as the bending beam andelastica loop tests have, by design, stress fieldsthat are not pure axial compression. Tests withembedded fibers (or a single fiber) in a matrixare liable to be affected by the residual stress onthe tested fibers caused by the matrix shrinkage,and the misalignment of the fibers during thetest (Jiang et al., 1993). Osawa et al. (1990) andMiwa et al. (1991) have estimated the compres-sive strength when the thermal stress workingfrom resin perpendicularly to the fiber±resininterface could be regarded as zero.

The recoil strengths of PAN-based CFs,shown in Table 6 (Dobb et al., 1995), however,

Figure 16 Reynolds and Sharp mechanism oftensile failure. (a) Misoriented crystallite linkingtwo crystallites parallel to the fiber axis. (b) Tensilestress exerted parallel to fiber axis causes layerplane rupture in direction La\, crack developsalong La\ and Lc. (c) Further exertion of stresscauses complete failure of crystallite. Catastrophicfailure occurs if the crack exceeds the critical size inLc or La\ directions (reproduced by permission ofKluwer Academic Publishers from J. Mater. Sci.,

1983, 18, 3337±3347).

Table 5 Compressive properties of CFs.

Commercialname

Tensilemodulus(GPa)

Tensilestrength

(st) (GPa)

Compressivestrength (sc)

(GPa)

sc�st71

Density(g cm73)

T-300a 235 3.2 2.88 0.9 1.76AS-4c 235 3.6 2.69 0.75 1.80T-40a 290 5.7 2.76 0.48 1.81T-1000d 295 7.1 2.76 0.39 1.82IM8c 310 5.17 3.22 0.62 1.80T-50a 390 2.4 1.61 0.67 1.81M40Jd 390 4.4 2.33 0.53 1.77GY-7b 520 1.8 1.06 0.59 1.96M-60Jd 585 3.8 1.67 0.44 1.94

Source: Kumar et al. (1993).aAmoco Thornel; bBASF Celion; cHercules Magnamite; dToray Torayca.


are typically lower than those obtained fromcomposite tests. Jiang et al. (1993) proposed auniversal logistic model which provides a goodfit to the recoil strength distributions obtainedat different gauge lengths. This analytical ex-pression can be used to obtain a physicallymeaningful extrapolated average ªzero gaugelengthº recoil strength.

As shown in Table 7, the failure strains ofCFs obtained in the compression test utilizingthe piezoresistivity by which the axial compres-sion strain of the filaments was accompaniedwere 73% or greater (DeTeresa, 1991). Anycompressive strength calculated assuming lin-ear elastic behavior will most likely be an over-estimate. These compressive strength values,nevertheless, are markedly higher than thoseobtained by the composite and recoil methods.The compressive strengths calculated from thestrains obtained for CFs embedded in resin byPrandy and Hahn (1991) are also fairly high.

As regards the fracture mechanism, it hasbeen reported that PAN-based CFs, which arethin and strong compared with pitch-basedCFs, exhibit buckling and/or kinking fracture(Dobb et al., 1990; Prandy and Hahn, 1991;Kumar et al., 1993); such failure modes are in

contrast to that of more crystalline CFs, such aspitch-based HM CFs, having sheet-like micro-structures, which tend to fail in transverse shearmode. However, there is a report which con-cludes that the single filaments of PAN-basedCFs, AS4, IM6, and IM7, embedded in anepoxy polymer, fail in transverse shear andnot in a microbuckling mode (Boll et al., 1990).

The compressive strength of PAN-based CFsincreases with increasing tensile strength (Dobbet al., 1990) and (002) d-spacing, and withdecreasing tensile modulus, orientation para-meter, crystallite sizes, and void content, re-spectively (Masson and Bourgain, 1992;Kumar et al., 1993; Dobb et al., 1995; Noritaet al., 1988; Sumida et al., 1989).

With boron-ion implantation, the compres-sive strength and torsional modulus of CFs(Table 8) increase by up to 25% and 50%,respectively, while the crystallite size, Lc, de-creases (Matsuhisa et al., 1991). A correlationbetween torsional modulus and compressivestrength was reported for PAN-based CFsand others (Norita et al., 1988).

Dobb and co-workers, furthermore, haveobserved that the compressive strength ofPAN-based CFs linearly increases with increas-

Table 6 Mechanical properties of various CFs (single fiber test).

Commercialname

Fibertype Tensile (st) Compressive (sc) sc�st

71

Modulus(GPa)

Strength(GPa)

Strain(%)

Strength(GPa)

Strain(%)

HM-S Grafila C1 320 2.1 0.36 0.8 0.26 0.38S-g Besfightb C2 220 3.0 1.35 1.5 0.69 0.50

X 370 3.4 0.93 1.4 0.38 0.41Y 330 3.1 0.96 1.3 0.40 0.42Z 280 3.5 1.25 1.9 0.67 0.54

Toraycac T1000 255 5.7 1.87 2.2 0.87 0.39

Source: Dobb et al. (1995).aCoutaulds; bToho rayon; cToray.

Table 7 Compressive properties of CFs.

Commercialname

Tensilestrength, st

(GPa)

Compressivestrength, sc

(GPa)

Compressivestrain(%)

sc�st71

T-1000a 7.03 9.9 73.3 1.41IM7b 5.38 8.1 72.9 1.51T650/35c 4.55 7.7 73.2 1.69G30±500d 3.79 8.3 73.6 2.19Grafil 33±500 4.48 9.0 73.9 2.01AS4b 4.00 8.5 73.7 2.13

Source: DeTeresa (1991).aToray Torayca; bHercules Magnamite; cAmoco Thornel; dBASF Celion.


ing intercrystalline disorder or intracrystallitedisorder, along different lines for the groupconsisting of newly developed fibers, whichhave much improved compressive strength,and the conventional fibers group (Dobb et al.,1995). Norita et al. (1988) and Sumida et al.(1989) have reported that Torayca H and Jseries fibers exhibit fairly high compressivestrengths as well as tensile strength comparedwith the conventional high-modulus fibers.

It would appear from Figure 17 that theability of CFs to withstand compressive stressis directly proportional to the degree of inter-crystallite disorder. The same tendency canalso be seen for the degree of intracrystallitedisorder. According to Dobb et al., although allPAN-based CFs consist of two major phases(i.e., crystallites and disordered material), thereare additional differences, depending on thefiber types, in (i) the distribution of crystallitesand disordered material within the cross-sec-tion, and (ii) the distribution in the sizes and theorientation of crystallites across the cross-sec-tion. The uniformity of the size distributionsuggests that compressive fracture is not gov-erned by a random flaw distribution as in thecase of tensile fracture (Dobb et al., 1995).More likely, failure is determined by somemicrocrystalline structure that is uniformly dis-tributed along the fiber (Boll et al., 1990).Although the total amount of disorderedregions plays an important role in determiningthe compressive strength and failure mechan-isms, it is equally important that the disorderedregion is homogeneously distributed through-out the fiber, as shown in Figure 18. Accordingto Dobb et al. (1995), it is also important thatcrystallites have dimensions below about 5 nmin all directions, and that their preferred orien-tation is maintained.

Norita et al. (1988) proposed the ratio oftensile modulus to shear modulus as a measureof anisotropy in various fibers, and found thatthe compressive strength of unidirectional CF

composites decreases with increasing ratio andwith increasing CF tensile modulus. Strongtransverse cohesive force, like covalent cross-links, are considered to increase the macro-scopic shear modulus and to improve thecompressive properties (Northolt et al., 1991).

1.01.4.2 Transverse

1.01.4.2.1 Transverse modulus

Hayakawa et al. (1994) found that PAN-based CFs can be deformed reversibly by

Figure 17 Compressive failure stress of CFs vs.intercrystallite disorder Dd (reproduced by permis-sion of Elsevier Science Ltd. from Carbon, 1995, 33,

1553±1559).

Table 8 Effect of B+ dose on compressive strength, torsional modulus, andtensile strength of single filament.

Property Fiber B+ dose (ions � cm72)

0 1015 1016 1017

Compressive strength (GPa) T1000G 7.84M40 3.63 5.00 7.16 7.74

Torsional modulus (GPa) T1000G 21.60 ± 31.40 ±M40 14.70 20.60 27.40 25.50

Tensile strength (GPa) T1000G 6.18 ± 7.45 ±M40 3.33 4.02 4.21 4.31

Source: Matsuhisa et al. (1998).


transverse compression up to a displacementcorresponding to about 10% of the diameter,the compressive deformation being more diffi-cult with increasing compressive deformation,and that the transverse modulus decreases withincreasing crystallite orientation or longitudi-nal modulus. As shown in Table 9, the fiberswith longitudinal moduli from 226 to 581GPaexhibited transverse moduli from 13.9 to6.9GPa.

1.01.4.2.2 Torsional modulus

The torsional modulus of PAN-based CFsdecreases with increasing Young's modulus.The torsional strength of the fibers increasesas the torsional modulus increases. The tor-sional strength also increases with increasingtensile strength (Table 10) (Sawada and Shindo,1992; Norita et al., 1988; Sumida et al., 1989).It is considered that the torsional modulusdepends on the size of the graphitic layers orcrystallites, their amount, their degree inradial and axial orientation, and the amountof disordered region, including cross-linking,between the layers or crystallites. It is alsoconsidered here that the transverse cohesiveforce, depending on the amount of disorderedregion, plays an important role in increasing thetorsional modulus and strength. Figure 19shows a torsional fracture surface of CF Bes-fight STA. The shape suggests that the trans-verse cohesive force is high compared with that

Figure 19 A torsional fracture surface of CFBesfight STA, torsional modulus: 24.0GPa,Young's modulus: 263GPa, torsional strength:1.28GPa, tensile strength: 4.04GPa. Views from(a) top and (b) a side (reproduced by permission ofElsevier Science Ltd. from Carbon, 1992, 30, 619±

629).

Figure 18 Lattice-fringe TEM image of CF T1000.Marker bar: 6 nm (reproduced by permission ofElsevier Science Ltd. from Carbon, 1995, 33,

1553±1559).


of high-modulus pitch-based CFs. The ditchesrunning along the fiber axis should be effectivefor diminishing the torsional strength, as thestress tends to concentrate at the bottom of theditches (Sawada and Shindo, 1992). Brydgeset al. (1969) obtained a torsional shear modulusof 24.1GPa for a HM fiber.

Villeneuve et al. (1993) have shown thevariations of the radial strain as a functionof the axial strain, both measured directly onthe TEM images recorded under axial tensileloading. A Torayca T300 filament with 6.2 mmdiameter exhibited a radial contraction as theaxial load was progressively increased, thougha T300 filament with 6.9 mm diameter exhib-ited expansion. Villeneuve et al. (1993a) havetentatively calculated a Poisson ratio at roomtemperature from TEM measurements for the

T300 fiber with a 6.2 mm diameter. The valueis 0.22, whereas the values in the literature arein the range 0.24±0.35. Krucinska and Stypka(1991) have reported a Poissons' ratio of 0.28for Safil.

1.01.4.3 High-temperature Properties

CFs exhibit creep behavior under tension athigh temperatures. Sines et al. (1989) carriedout creep tests on HM3000 yarn specimens(Hercules, 3000 filaments, with a nominal dia-meter of 7 mm), applying stresses ranging from455 to 648GPa at temperatures from 2120 to2430 8C for several hours. Elongation up to140% was measured on these fibers. For

Table 10 Torsional properties of CFs.

FiberYoung'smodulus(GPa)

Torsionalmodulus(GPa)

Tensilestrength(GPa)

Torsionalstrength(GPa)

Torayca an early CFa 201.4 23.72 2.316 1.000Exp. type IIIa 269.1 23.61 4.154 1.052Besfight HTa 235.9 25.03 2.839 1.379

STAa 262.9 24.03 4.044 1.279GIRIO Exp. CFa 273.9 23.82 3.028 0.965S H30a 235.1 25.08 2.474 1.106Magnamite ICMa 441.7 21.85 2.327 0.756

2±4/3Aa 431.4 21.47 2.167 0.740Torayca M40a 558.4 22.55 2.607 0.796Besfight HMa 466.3 23.03 2.833 1.113Torayca T300b 230 16.7

M30b 294 16.7M40b 392 15.7M46b 451 14.7T800b 294 17.6M40Jb 392 18.7

Source: aSawada and Shindo (1992); bNorita et al. (1988); Toray (private communication, 1998).GIRIO: Government Industrial Research Institute at Osaka. S: a company in Japan.

Table 9 Transverse compressive properties of CFs.

Fiber Diameter(mm)

Density(g cm73)

Orientationparameter(17 b�p71)

Longitudinalmodulus(GPa)

Transversemodulus(GPa)

Compres.strength(GPa)

T03 7.4 1.77 0.808 226 13.9 (1.9) ±T10 5.8 1.84 0.830 267 12.7 (0.4) ±M40 5.4 1.83 0.887 379 11.6 (1.1) ±M50 5.3 1.90 0.915 469 9.1 (2.6) ±M60 5.0 2.03 0.945 581 6.9 (1.1) ±BesfightSTIII 6.73 0.779 250 16.8 4.23HM40 6.7 0.901 398 6.8 1.79HMS6X 4.2 0.946 588 5.5 1.57

Source: Hayakawa et al. (1994); Besfight: Fujita et al. (1992).b, Width at half-maximum intensity of (002) peak. Standard deviation.


steady-state creep the apparent activation en-ergy and activation volume were 1058 kJmol71

and 2088AÊ 3, respectively. These values of ap-parent activation energy and apparent activa-tion volume suggest that the creep occurs on acrystallite scale.

An apparent non-Hookean behavior hasbeen observed on some Torayca CFs at tem-peratures above 1000 8C during thermal expan-sion experiments, the heated fibers keeping thememory of their last shape. A similar observa-tion was also made in high-temperature looptests. Over 850 8C, a plastic deformation beginsto occur (Trinquecoste et al., 1996). In thistemperature range, a fairly large amount ofnitrogen is eliminated from the fiber. Accord-ingly, such an apparent inelastic behavior seemsto be induced by a structural rearrangement inthe fiber when heteroatoms, mainly nitrogen,are driven out. For example, CFs T300 andT900 exhibit nitrogen contents of 6.3% and5.4%, respectively, but after heat treatment to2000 8C, both the fibers contain no nitrogen.

1.01.5 ELECTRIC AND MAGNETICPROPERTIES

1.01.5.1 Electrical Resistance andThermoelectric Power

The electrical resistance of CFs, Torayca,Besfight, and Amoco decrease with increasingYoung's modulus. These values are shown inTable 11. Matsubara et al (1995) have mea-sured electrical resistivity and thermoelectric

power at temperatures from 0 to 300K forTorayca intermediate (TH series) (T800H,T-1000G) and HM (MJ series) type (M40J,M46J, M50J, M60J) fibers. According to theirresults, the M-type (J-series) fibers indicate thesemiconductor-like temperature dependence allover the temperature ranges examined, andthe resistivity increases with decreasing elasticmodulus. On the contrary, the T-type fibersexhibit a peak in resistivity around 35K,as shown in Figure 20. The metallic-like

Table 11 Physical properties of selected types of CFs in manufacturers' data sheet.

Mfr.Fibertype

Carboncontent(%)

Tensilemodulus(GPa)

Electricalresistivity

(1073O cm)

Specificheat

(cal. g71 8C71)CTE

(1076 8C71)

Thermalconductivity

(cal. cm71 s71 8C71)

Amoco T-300 92 231 1.8 70.6[Thornel] T-40 94 290 1.45 70.75

T650/42 94 290 1.42 70.75T-50 99 390 0.95 71.13

Tenax HTA 238 1.6 0.17 70.1 4.16 1072

IMS 290 1.45Toho rayon HTA 235 1.5[Besfight] IM400 295 1.4

HM35 345 1.0HM40 380 0.9

Toray T300 93 230 1.7 0.19 70.41 2.56 1072

[Torayca] T300J 94 230 1.5 0.18 70.43 2.236 1072

T400H 94 250 1.6 0.18 70.45 2.526 1072

T700S 93 230 1.6 0.18 70.38 2.246 1072

T800H 96 294 1.4 0.18 70.56 8.396 1072

T1000G 95 294 1.4 0.18 70.55 7.656 1072

Figure 20 Resistivity vs. temperature curves forTorayca T800H. Solid line represents the calcula-tion (reproduced by permission of the authors from

Matsubara et al., (1995).

Electric and Magnetic Properties 23

temperature dependence observed in spite oftheir lower crystallite perfection comparedwith M-type fibers can be explained by consid-ering the Rayleigh wave phonon whose velocityis so small that a number of phonons are excitedeven at liquid helium temperature. Theyexplained the behavior assuming a mixturemodel of the band conduction and two-dimen-sional variable-range hopping conduction.

Figure 21 shows the temperature depen-dence of the thermoelectric power of ToraycaMJ-type fibers (Matsubara et al., 1995). TheM-type fibers exhibited the positive tempera-ture dependence of thermoelectric power (S),and the absolute values decrease with decreas-ing modulus. The T-type fibers showed thetemperature dependence of thermoelectricpower similar to that of the resistivity. Such adependence implies that the wave phononsaffect the thermoelectric power.

1.01.5.2 Electromechanical Properties

Single filament electromechanical behaviorfor a Torayca T300 fiber has been observedby measuring the fractional increase in electri-cal resistance (DR/R0), and stress and strainsimultaneously obtained during static andcyclic tension to failure. Wang and Chung(1997a) have found the following phenomena:(i) DR/R0 increases monotonically with strain±stress, with a slight negative deviation from

linearity, and reaches a value of 0.0019 or0.0021 at a strain of 1.1% under static tension.The strain sensitivity values obtained are 1.8and 1.9; (ii) The strain and DR/R0 obtainedduring cyclic tension are totally reversible atlow values of the stress amplitude (up to58.1% of the fracture stress), but their irrever-sible components increase with stress amplitudeat high values of the stress amplitude; (iii) Theobserved irreversible resistance change is attrib-uted to damage, as supported by the accompa-nying decrease in the elastic modulus. Damagecan occur even in the regime of elastic deforma-tion; (iv) Irreversible strain occurs at stressamplitudes 573.1% of the fracture stress andincreases with increasing stress amplitude;(v) The reversible resistance change is mainlydue to the dimensional change associated withelastic deformation; (vi) The strain sensitivity(reversible DR/R0 per unit reversible strain) is1.9±2.3 and is quite independent of stress orcycle number (from 1 to 2). These values areclose to those obtained by Owston (1970).

The CF embedded in epoxy resin exhibits theopposite behavior from that mentioned above.The embedded fiber showed the decrease of thefractional change in resistance upon tensileloading and its recovery upon subsequent un-loading. The decrease in the fractional resis-tance change is attributed to the reduction ofthe residual compression stress in the fiber hav-ing been induced on curing and cooling of theresin (Wang and Chung, 1997b).

1.01.5.3 Magnetoresistance

Values of average crystallite layer planetransverse magnetoresistance for CFs areshown in Table 12 (Hishiyama et al., 1984).The CFs from PAN fiber without stretchingduring stabilization treatment exhibit negativevalues. The CFs prepared from PAN fiberstretched during stabilization treatment, how-ever, exhibit positive values. Furthermore, itcan be seen that the stretching at 2700 8C iseffective for improving the degree of graphiti-zation (Hishiyama et al., 1991).

1.01.6 THERMAL PROPERTIES

1.01.6.1 Thermal Expansion

The coefficient of thermal expansion alongthe basal plane for single crystal graphite, aa, is71.66 1076K71 around RT, and then in-creases up to +1.26 1076K71 over 1000 8C.The coefficient of thermal expansion in the

Figure 21 Temperature dependence of thermoelec-tric power for high modulus CFs (reproduced bypermission of the authors from Matsubara et al.,

(1995).


direction perpendicular to the basal plane, ac, is& 276 1076+36 1079K71 in the tempera-ture range from RT to 800 8C (Trinquecosteet al., 1996).

The coefficient of longitudinal thermal ex-pansion for CFs decreases with increasingYoung's modulus. This can be apparentlyshown when the CTE values for Torayca andThornel fibers in Table 11 are plotted as afunction of Young's modulus. The correlationbetween Young's modulus and CTE has beenalready observed by Wolff (1987); an abruptchange is present in the slope of the curve atabout 680GPa, where the plotted values are forsingle-crystal graphite and pitch-based CFswith higher Young's moduli than those ofPAN-based CFs.

With longitudinal thermal expansion, ak, ofCFs, T300, T700S, and T900, as shown inFigure 22 (Trinquecoste et al., 1996), all of thefibers behave similarly on increasing tempera-ture: slight contraction from RT to around400 8C; expansion up to 2000 8C; burning athigher temperatures. The contraction in thefirst temperature region is comparable to thenegative value of the aa coefficient of the gra-phite single crystal at low temperatures. Com-parison of the ak coefficients shows that T300and T700S fibers present a very similar beha-vior, while T900 has a much lower ak coeffi-cient.

The ak coefficients up to 1000 8C have alsobeen measured on fibers T300 and M40 byYasuda et al. (1987), obtaining a similar beha-

vior as by Trinquecoste et al. (1996). Villeneuveand Naislain (1993) measured longitudinalthermal expansion for T300 by TEM, and ob-tained its mean coefficients calculated for var-ious temperature ranges, RT to 800 8C.

According to Trinquecoste et al. (1996), theradial thermal expansion coefficient is506 1076K71 for Torayca T900, and thatfor M60J measured with a composite in therange 40±90 8C by Nanjyo et al. (1994) is156 1076 8C71. Villeneuve et al. (1993) ob-tained mean coefficients of radial expansionfor T300 fiber, for various temperature rangesfrom 25 8C to 800 8C based on measurements byTEM, and found that it decreases from17.66 1076 8C71 for a range of 25±200 8C to1.26 1076 8C71 for a range of 600±800 8C.

1.01.6.2 Thermal Conductivity

The thermal conductivity of CFs increaseswith increasing Young's modulus. This can beseen if the values for Torayca fibers in Table 11are plotted as a function of Young's modulus.

Low thermal conductivity PAN-based CFsare used as reinforcements of phenolic resinmatrix composites used as ablative insulatorsin solid rocket motor nozzles and exit cones.The CFs with HTTs of 900 8C, 1120 8C, and1350 8C show thermal conductivities as low as2.0, 4.0, and 13.6Wm71 K71, respectively(Katzman et al., 1994). After Heremans et al.(1985), the thermal conductivity of PAN-basedCFs increases from the order of 1071 to theorder of 102 Wm71 K71 with increasing tem-perature from several tens to several hundredsof degrees Kelvin.

Figure 22 Coefficients of longitudinal thermalexpansion of CF filaments (reproduced by permis-sion of Elsevier Science Ltd. from Carbon, 1996, 34,

923±929).

Table 12 Average crystallite layer plane transversemagnetoresistance (Dr�r71)cr at liquid nitrogen

temperature for CFs.

HTT Treatment (Dr�r71)cr

3000 70.5472900 70.9542800 70.6183000 Stretcheda 26.182830 6.373000 SG-4b 26.202830 9.192700 2.103000 SG-8b 21.182830 9.232700 4.743000 SG-12b 22.992830 4.732700 1.54

Source: Hishiyama et al. (1984).aStretched with a load of 40mg per filament both in stabilization andcarbonization.bPrepared from the CF with treatment of stretcheda,by stretching at 2700 8C with loads of 40mg per filament (SG-4),80mg per filament (SG-8), and 120mg per filament (SG-12).

Thermal Properties 25

1.01.7 SURFACE PROPERTIES

1.01.7.1 Morphology and Surface Areas

The open pore volume (porosity) of CFsincreases with increasing HTT up to around1000 8C, and then decreases to be negligible atHTTs above 1200 8C (Spencer et al., 1970; Is-mail, 1991). The surface of CFs has manywrinkles or ridges running nearly along thefiber axis. Surface roughness results from suchwrinkles or ridges as well as depressions on thefiber surface. The ratios of BET surface areas tothe geometrical areas estimated from laser dia-meters can be regarded as a measure of surfaceroughness. Table 13 shows those areas andsurface roughness for some Torayca and Bes-fight CFs (Shindo, 1983). The Besfight fibershave a circular cross-section, but the Toraycafibers used have a cross-section of broad beanshape. The circumference of the cross-sectionfor the Torayca fibers was then estimated withan approximation that the cross-section has ashape of an ellipse with apparent major andminor axes. Table 13 also shows another mea-sure of geometrical surface areas obtained fromthe densities by a density gradient column test.Both of these values almost coincide with eachother.

Surface area of CFs can be determined bymeasuring the amount of a monomolecularlayer of nitrogen or krypton on the CF surfaceusing the BET method. ASA, which is the areaoccupied by chemically active sites on the fibersurface, can be determined by measuring theamount of oxygen chemisorbed at 300 8C on thesurface of fiber degassed at 1000 8C in a highvacuum, under the assumption that the chemi-sorbed oxygen forms C±O groups at active siteslocated at the basal plane edges of the crystal-lites. Table 13 also shows the active surfaceareas for the Torayca and Besfight CFs withand without surface treatment. Ismail (1987)obtained an ASA of 0.075 for an Amoco T-300 fiber by extrapolating a plot of ASA vs.CO2 produced at 300 8C to zero CO2. A BETarea for the fiber obtained by Ismail was0.56m2 g71. The ASA value is 13% of theBET SA value. The ASAs for the oxidizedfibers in the table are in the range 9±13%.Those for the fibers without surface treatmentare in the range 5±8% for HS type fibers, and3% for HM type fibers. Such low values ofASA correspond to the fiber cuticle composedof graphitic basal planes parallel to the fibersurface. The surface treatment to which M-40has been subjected seems to be severe comparedwith the others.

Takahagi and Ishitani (1988) have proposedthe asymmetric parameter, ac, of the XPS C1s

spectrum as a measure of structural deviationfrom the completeness of the graphite structure.The parameter ac is defined by the ratio a/b,where a is the long tailing component toward ahigher binding energy, and b is the short com-ponent of FWHM (full-width at half-maxi-mum), namely a+b, of the C1s spectrum.They reported that it increases in ToraycaM40 by a surface oxidation treatment. Thisdisordering was not restored after the fiberwas subjected to a thermal treatment at1000 8C in a vacuum, where surface functionalgroups are removed.

1.01.7.2 Functional Groups

XPS is a very powerful technique for char-acterizing functional groups on CF surfaces.Oxidized Torayca CF T300 has markedly largeramounts of carbonyl and hydroxyl groups anda moderately larger amount of carboxyl groupsthan the unoxidized T300 fiber. The concentra-tion of these groups was determined by XPSafter the functional groups had been modifiedby chemical reactions with fluorine com-pounds. The relative amounts of oxygen esti-mated from the amounts of three functionalgroups determined by the surface modificationreaction, however, were about 20% of the ob-served values from a direct measurement of O1sspectra for both the unoxidized and oxidizedCFs. In their previous paper Takahagi andIshitani (1984) concluded that hydroxyl, carbo-nyl, and carboxyl groups are major functionalgroups formed by the surface oxidation treat-ment, whereas the amounts of ester and ethergroups are negligible. The chemical modifica-tion reagents may react only with the functionalgroups in the outermost layer (5±10AÊ ) in theCF surface because of the poor permeability ofreagents into CF. If there are more functionalgroups in a greater depth within the detectiondepth of XPS (30±50AÊ ), the direct O1s mea-surement will give a higher relative intensity(Takahagi and Ishitani 1988).

Nakayama et al. (1990) determined the con-centration of functional groups on the surfaceof weakly and strongly oxidized Torayca T800fibers by curve-fitting the C1s spectrum usingan asymmetric peak shape or a symmetric peakshape. Amounts of groups such as hydroxyl,carboxyl, and amine were estimated by gaschemical modification using fluorine com-pounds (Table 14). Consequently, the real sur-face of the oxidized CF was considered to becomposed of both graphite-like and aliphaticstructures with functional groups.

Estimation of functional groups on CF sur-faces using chemical modification were also


Table 13 Surface properties of CFs.

Fiber

Young'smodulus(GPa)*

SA(m2 g71)

Incrementof SA bySTr (%)

ASA(m2 g71)

Rate of ASA toSA (%)

Incrementof ASA bySTr (%)

Surface area of one filament(1076� m2� m71) evaluated from

Laserdiameter (SAL)

Density andweight (SAD) SA (SAB)

Roughness(SAB±SAL)6 (SAL)

71

(%)

HS X550-Ua 244 0.603 0.046 7.7 19.5 19.5 32.1 65X550-Sa 240 0.630 4.47 0.085 13.4 74 19.4 19.4 33.4 73ST III-Ub 250 0.473 0.025 5.2 21.1 21.1 29.8 41ST III-Sb 247 0.488 3.17 0.049 10.0 92 21.1 21.3 31.4 49

HM M40-Ua 392 0.526 0.014 2.7 20.4 19.6 31.1 53M40-Sa 1.199 128 0.138 11.5 325 20.7 20.2 70.2 238HM40-Ub 0.496 0.013 2.6 20.8 20.2 31.0 49HM40-Sb 0.592 19.3 0.052 8.8 238 20.8 20.6 37.1 78

Source: Shindo (1983).*Manufacturer's data, a Torayca, b Besfight, SA: BET surface area, ASA: Active surface area, STr: Surface treatment.

applied to Hercules AS-4 and Celenese Celion6000. Fluorine-containing compounds and amercuric fluorine-containing compound havebeen used as modification reagents. The resultsobtained are shown in Table 15 (DeVilbiss andWightman 1986). XPS analysis, utilizing deri-vatization reactions with fluorine-containingcompounds, of functional groups on CFs wasalso performed by Chan et al. (1991). Theresults obtained are shown in Table 15. TheCFs used were Hercules IM7, Amoco T40,Amoco T650±42, and BASF Celion G40±800.These CFs are classified as intermediate-mod-ulus grade having a tensile modulus of about280GPa. All the fibers were surface treated andcontained no sizing.

In this case, it is to be noted that though theÐCOOH group is clearly acidic, this is notalways true for the hydroxyl group, RÐOH.The ÐOH group can be acidic if the R group isa phenyl or basic if R is aliphatic. Similarly, theÐC=O is generally slightly basic, but if thereis an alpha hydrogen, such as in CHÐC=O,this hydrogen atom shows acidic characteris-tics. Therefore, the assessment of carbon-sur-face acidity or basicity can only be madetentatively from XPS data, particularly sinceXPS cannot detect hydrogen atoms (Chanet al., 1991). In such a case, surface acid-basefree energy data become useful for deducingwhether either one or both of them contribute

to the basic nature or to the acidic nature, aswill be shown in Table 16. It has also beenpointed out that CFs from various manufac-turers have significant differences in their sur-face acidity (Chun et al., 1992).

1.01.7.3 Surface Free Energy

Chan et al. (1991) made contact-angle mea-surement to estimate surface free energy on theCFs at the same time as XPS examination offunctional groups. In the measurement, methy-lene iodide was used as the probe liquid for theLifshitz±van der Waal's interactions; ethyleneglycol and formamide were used to probe forthe acid±base interactions. The results obtainedare shown in Table 16 (Chan et al., 1991). Suchsurface acid±base free energy data for CFs areapplicable for searching a CF more compatiblewith resin having either acidic, basic, or am-photeric characteristics.

1.01.7.4 Wetting Property

As pointed out by Zielke et al. (1996a), com-mercial surface-treated carbon fibers are moreor less surface-contaminated with adsorbed

Table 14 Surface chemical composition of surface-oxidized Torayca T800H fiber.

Sample O/C6 103 N/C6 103

:COH 7COOH :C-O-C:7COOC:=C=O, etc.

:Si-O- Total 7NH2 =NH=N7

Total

Control ± 2 30 14 46 1 24 25Weakly oxidized 1 5 75 ± 81 3 28 31Strongly oxidized 3 31 244 ± 278 5 29 34

Reproduced by permission of Elsevier Science Ltd. from Nakayama et al., 1990.

Table 15 Concentration of functional groups per hundred surface atoms, based on functional groupderivatization and fluorine or mercury XPS analysis.

Fiber C O N OH 7C=O COOH OH+C=O+COOH

IM7c 81 14 5.2 0.7 1.1 1.8 3.6a

T650±42d 85 11 2.0 0.5 1.4 1.6 3.5a

G40±800e 82 15 2.4 0.7 1.3 1.2 3.2a

T40d 95 2.7 1.9 0.2 1.8 0.6 2.6a

AS-4c 85 11 4.1 0.1 0 0.4 0.5b

Celion 6000e 84 14 1.6 0.2 0.2 0.3 0.7b

Source: aChan et al. (1991); bDeVilbiss and Wightman (1986). c Hercules, d Amoco. e BASE. C, O, and N: total carbon, oxygen, and nitrogendetermined by XPS prior to derivatization.


oxidation products. The contaminations, how-ever, can be removed efficiently by extractionwith boiling water. Such extraction bringsabout an increase of the surface concentrationof carboxyl groups as a consequence of hydro-lysis of anhydrides, which could have beenformed by a drying treatment. In scanningtunneling microphotographs on the micro-and nanoscale, the crevices and pits are moreshapely and defined on the washed fiber sur-faces, and defects in the form of trenches run-ning longitudinally along the fiber axis are alsomore pronounced on the extracted fibers whichhave sharper images (Figure 23). Figure 24shows the functional groups determinedthrough curve fitting of the O1s peak obtainedby the XPS examination of these CFs beforeand after water washing. For the curve fitting,the following functions were considered:C=O groups (group 1), carbonyl oxygenatoms in esters, amides, anhydrides, and oxy-gen atoms in hydroxyls or ethers (group 2), theether oxygen atoms in esters and anhydrides(group 3), and the oxygen atoms in carboxylgroups (group 4). The content of oxygen atomsof group 2 of all fibers varies between 30.4%(fiber (C)) and 52.8% (fiber (E)), and generallysurpasses the content of group 1, for whichvalues between 9.1% (fiber (E)) and 32.1%(fiber (F)) were determined. Contents between5.3% (fiber (F)) and 12.8% (fiber (D)) wereobtained for COOH (group 4). Additionally,water was found in most cases, even for driedfibers. The contribution of water was in therange 0±5.5% in relation to total oxygen.

Although XPS is a valuable method for iden-tifying functional groups in a CF surface, itoffers an integral picture of a more or lessdefined surface volume, but not of the nettwo-dimensional surface. The adhesion of apolymer to the fiber surface is controlled bythe surface structure and chemistry of the CFsused. Accordingly, the information of the sur-face which is most relevant for adhesion can beobtained by contact angle measurement. Fromsuch a point of view, Zielke and co-workershave investigated the surface chemistry of CFsby measuring contact angles of the fibers with

aqueous solutions of different pH values(Zielke et al., 1996a, 1996b, 1996c).

WSL values were estimated from the contactangles obtained, and WSL/pH value diagramswere obtained for extracted surface-oxidizedfibers (unsized Tenax HTA, AS4, T800, andM40) investigated. A schematic presentationof the work of adhesion, WSL, depending onthe pH value, is given in Figure 25. WSL iscomposed of the following fractions (Equations(1) and (2)):

WLWSL +WAB

SL (1)

WABSL =WAB/B

SL +WAB/HSL (2)

where WSL=work of adhesion, WLWSL =work

of adhesion due to Lifshitz±van der Waal'sinteractions, WAB

SL =work of adhesion due toacid±base interactions, and WAB/B

SL =work ofadhesion due to Bronsted acid±base complexes.WLW

SL was estimated from measurements withdiiodomethane. Although fibers of differentheat-treatment temperatures and crystallineperfection were investigated, those values werein a rather narrow range from 57 to 64 mJm72.The two kinds of acid±base interactions makecontributions to WSL. With all fibers, the maincontribution results from hydrogen bonds. Atvery low pH values they are based on carboxyl,hydroquinone, hydroxyl, phenol, and carbonylsurface groups. With increasing pH value, thecarboxyl groups can form Bronsted acid±basecomplexes. The free energy of complex forma-tion causes an increase of WSL. This can beunderstood by the following equation:

WAB/BSL =7ni �DG0

i (3)

where ni=number of acid±base complexes iand DG0

i =Gibb's free energy of formation ofthe acid±base complex i.

The height of the step is not large because thecontribution of hydrogen bonds formed by car-boxyl groups at the lower pH values is simulta-neously reduced by the formation ofcarboxylate anions, and the hydrogen bondswith carboxylate anions are much weaker

Table 16 Acid±base component of surface free energy for CFs (gLWS ) ethyleneglycol/CF (gAB7

SL ) and formamide/CF (gABgSL ) (all units are in mNm71).

Fiber gLWS gAB7SL gAB

SL Rank

IM7 31.4 30.1 32.5 1 (amphoteric)T650±42 30.8 31.7 27.2 2 (basic)G40±800 31.5 24.1 34.8 3 (acidic)T40 33.2 23.8 23.2 4 (amphoteric)

Reproduced by permission of Elsevier Science Ltd. from Chan et al., 1991.

Surface Properties 29

than those with carboxyl groups. The sameinterpretation can be given for the second stepat high pH values at which Bronsted acid±basecomplexes are formed between phenolic OHgroups and the basic solutions (Zielke et al.,1996c).

1.01.7.5 Reactivity

Generally, the oxidative stability of CFs be-comes higher as the HTT or Young's modulusrises. Although CFs are attacked with manykinds of oxidizing reagents such as nitric acid,sulfuric acid, dichromatic acid, kalium-perman-ganate + acid, etc., even at RT, the tempera-tures at which the weight loss of CFs starts inoxygen, air, carbon dioxide, and water vaporare in the range 300±800 8C.

Barr (1995) measured rapid weight loss in airflows at temperatures of 430±470 8C, and long-term weight loss in air at 316 8C and 371 8C forCFs having tensile moduli of 207±348 GPa. The60-hour test of 207±276GPa fibers at 430 8Cresulted in a weight loss of at least 5%, which isan aid in distinguishing between CFs. The mostoxidation-resistant fiber was T-650/50X, whichhad a modulus of 345GPa and gave a weightloss of only 1% at 470 8C. Long-term weightloss after 500 h in air at 371 8C for 207±276Gpafibers ranged from 31% (T-650/42) to 78%(T-300 3K). In contrast, the 345GPa modulus

fibers were so oxidation resistant at 371 8C thatmonitoring was continued for 21 000 h (2.4years). Furthermore, Barr found that theweight losses after 60 h at 430 8C or 450 8C inthe rapid test correlated well with the results oflong-term aging for 1000 or 2000 h at 316 8C.

In the oxidation of carbon materials withgaseous reagents, metallic impurities such asalkali metals and alkaline earth metals acceler-ate the reaction (McKee, 1981). For example,the dependence of oxidation rates on the so-dium content was observed in CFs including

Figure 23 Top views of the surface of CF AS4, unoxidized and unsized, on the micro- and nanoscale. (a)The nonextracted fiber showing less contrast; (b) the extracted fiber with well-contrasted elongatedcrystallites; (c) a region of the nonextracted fiber with a high degree of surface roughness; and (d) theextracted fiber with well-ordered graphitic regions (reproduced by permission of Elsevier Science Ltd. from

Carbon, 1996a 34, 983±998).

Figure 24 Functional groups determined by XPSshowing the relevant oxygen atoms at their corre-sponding binding energies (reproduced by permis-sion of Elsevier Science Ltd. from Carbon, 1996b,

34, 999±1005).


T-300 (Ismail, 1991). Contamination of CFswith sodium compounds is liable to occur inthe process of surface treatment by anodic oxi-dation using an alkaline electrolite containing asodium compound, which is one of the com-mercial treatments.

Stark et al. (1994) have investigated the de-terioration of an HM CF, Torayca M40, re-heated at 3000 8C by means of exposure toatomic oxygen, and found that the Young'smodulus and tensile strength decrease with in-creasing exposure to atomic oxygen, and thatthe CF surface is severely degraded by atomicoxygen.

CFs can be intercalated with various kinds ofelements and compounds. The capability offorming intercalation compounds of PAN-based CFs, however, is low compared withthose of pitch-based CFs and vapor-grownCFs. This is due to the less graphitic structureand circumferential skin structure of the CFs.

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