Carbon 42 (2004) 795–804

www.elsevier.com/locate/carbon

A comparison of the interfacial, thermal, and ablativeproperties between spun and filament yarn type

carbon fabric/phenolic composites

Jong Kyoo Park a,*, Donghwan Cho b, Tae Jin Kang c

a Agency for Defense Development, Composite Laboratory, Yuseong, P.O. Box 35-5, Daejon 305-600, South Koreab Department of Polymer Science and Engineering, Kumoh National Institute of Technology, Kumi, Kyungbuk 730-701, South Korea

c School of Materials Science and Engineering, Seoul National University, San 56-1, Shimlim-Dong, Kwanak-Ku, Seoul 151-742, South Korea

Received 8 September 2003; accepted 14 January 2004

Abstract

In the present paper, the interfacial, thermal, and ablative properties of phenolic composites reinforced with spun yarn type

carbon fabrics (spun C/P composite) and filament yarn type carbon fabrics (filament C/P composite) heat-treated at 1100 �C have

been extensively compared. The interlaminar shear strength, crack growth rate, and fracture surface were studied to evaluate the

interfacial characteristics of the composites using short-beam shear test, double cantilever beam test, and scanning electron

microscopy, respectively. The thermal conductivity and the coefficient of thermal expansion were also measured in the longitudinal

and transverse directions, respectively. To explore the ablative characteristics of the composites in terms of insulation index, erosion

rate, and microscopic pattern of ablation, an arc plasma torch was used. The interfacial properties of the spun C/P composite are

significantly greater than those of the filament C/P composite, with qualitative support of fracture surface observations. It has been

investigated that the presence of protruded fibers in the phenolic matrix of the spun C/P composite may play an important role in

enhancing the properties due to a fiber bridging effect. The longitudinal thermal conductivity of the spun C/P composite is about 7%

lower than that of the filament C/P counterpart. It has been found from the ablation test using arc plasma torch flame that the

erosion rate is 14% higher than that of the filament C/P counterpart. Consequently, all the experimental results suggest that use of

spun yarn type carbon fabrics heat-treated at low carbonization temperature as reinforcement in a phenolic composite may sig-

nificantly contribute to improving the interfacial, thermal, and ablative properties of C/P composites.

� 2004 Elsevier Ltd. All rights reserved.

Keywords: A. Carbon composites; B. Carbonization; C. Thermal analysis; D. Interfacial properties, Thermal conductivity

1. Introduction

During the last decades, polyacrylonitrile (PAN)-

based carbon fiber reinforced phenolic composites have

increasingly replaced rayon-based carbon fiber rein-

forced phenolic composites for a thermal protection

system of reentry vehicles or rocket engine componentsdue to their excellent ablation resistance and mechanical

properties [1]. However, because of high thermal con-

ductivity and low interfacial property with phenolic

resin, use of PAN-based carbon fibers has been limited,

especially in thermal insulation applications. For solid

*Corresponding author. Tel.: +82-42-821-4611; fax: +82-42-821-

2393.

E-mail address: pjkyoo@hanafos.com (J.K. Park).

0008-6223/$ - see front matter � 2004 Elsevier Ltd. All rights reserved.

doi:10.1016/j.carbon.2004.01.046

rocket motor applications, the key requirements of

carbon fiber composite are not only of low thermal

conductivity to minimize the thickness of pyrolyzed

carbon layer but also of high interfacial strength to re-

duce possible catastrophic erosion by abnormal ablation

behavior. Therefore, development of PAN-based carbon

fiber composites having comparable thermal conduc-tivity and interfacial strength with rayon-based carbon

fiber composites has been in great demand. It has been

generally known that the structural alteration of rein-

forcing textiles may change the thermal conductivity

and/or the interfacial properties of composites. It has

been reported that the phenolic composites reinforced

with spun yarn type PAN-based carbon fabrics reduce

the thermal conductivity along the fiber direction about30% in comparison with filament yarn type counterparts

Table 1

The physical properties of stabilized PAN spun yarn

Properties Value

Linear density (Tex) 102

Twist (TPM) 258

Breaking strength (gf ) 972

Elongation at break (%) 11

796 J.K. Park et al. / Carbon 42 (2004) 795–804

[2]. Also, the thermal conductivity and the tensile

strength of PAN-based carbon fiber composites strongly

depend on heat-treatment temperature (HTT) of carbon

fibers [3–5].In many aerospace and military applications using

thermal insulating materials with anti-ablation perfor-

mances, it may be considered that lowering the thermal

conductivity of a composite of interest is primarily more

important than maintaining its mechanical strength.

Lowering the thermal conductivity of a carbon/phenolic

composite may be successfully achieved by controlling

the thermal conductivity of carbon fibers rather thanphenolic matrix because the conductivity of carbon fi-

bers is much greater than that of polymer matrix and it

is also effectively managed through a variety of heat-

treatment processes. Only a few literatures on achieving

carbon fibers with less thermal conductivity have been

reported [3–9]. Recently, Park and Kang [10] reported

that use of low temperature carbon fibers heat-treated at

1100 �C as reinforcement of the phenolic compositeimproves the insulation property of the composite.

The aim of the present study is to understand the

interfacial, thermal, and ablative properties between

phenolic composites reinforced with spun yarn type

carbon fabrics heat-treated at 1100 �C (spun C/P com-

posites) and those of phenolic composites reinforced

with filament yarn type carbon fabrics heat-treated at

1100 �C (filament C/P composites). The interlaminarshear strength and the interlaminar fracture energy of

the composites are studied to explore the effect of the

spun yarn on the interfacial properties. Scanning elec-

tron microphotographic results also provide to qualita-

tively support the fracture behavior after double

cantilever beam tests and the ablation behavior after

ablation tests, respectively.

2. Experimental

2.1. Materials and composites

Stabilized staple PAN fibers (Pyron�, Zoltek Co.,

USA) were used as precursor fibers to prepare spun

carbon yarns for fabricating spun C/P composites. The

length of the staple fiber is 102 mm in average. The sta-

bilized staple PAN fibers were converted into the sta-

bilized PAN spun yarns using a semi-worsted spinning

method. The physical properties of the stabilized PAN

spun yarn are summarized in Table 1. The stabilizedPAN spun yarns were woven into a fabric form with an

eight harness satin texture. The woven spun yarn type

fabrics were proprietarily heat-treated up to 1100 �C at a

heating rate of 1 �C/min with a purging N2 gas in a

batch-type carbonization furnace and then successfully

converted into spun yarn type carbon fabrics. Resol-

type phenolic resin (KC-98�, Kangnam Chemical Co.,

Korea) was used as a matrix for the composites used in

this work. The resin content of spun C/P prepregs was

about 36% (w/w). The spun C/P composites were fab-

ricated at 150 �C for 2 h with an identical curing cycle

using a hydroclave. The debulking process at 105 �C for

1 h was done to remove possible entrapped air and

consequently voids in the resulting composites. A pres-

sure of about 1000 p.s.i. was applied. The resin contentof spun C/P composites after fabrication was about 32%

(w/w).

2.2. Interfacial measurements

In order to explore the interlaminar shear strength of

spun C/P and filament C/P composites, short-beam

shear tests were conducted using a universal testingmachine (Instron 4505) according to ASTM D-2344.

The crosshead speed was 1.3 mm/min and the span-to-

depth ratio was 4. The number of the specimen for each

measurement was 5.

The interlaminar fracture energy was determined by a

compliance method using a double cantilever beam

(DCB) test according to ASTM D-5528. The specimen

dimensions were 100 mm · 20 mm · 5 mm with a mid-ply initial crack of 25 mm in length at one end. The

initial crack was made by a folded sheet of 36 lm thick

aluminum foil interleaved between the two inter-plies

located in the middle of the composite prior to com-

posite processing. The steel hinge tab was firmly adhered

to both upper and lower sides of the pre-cracked spec-

imen. To monitor the location of the interlaminar crack

tip, one lateral side of the specimen was coated withtypewriter correction fluid. Such the thin coating pro-

vided good contrast between the dark crack and the

white intact area of specimen. Individual marks of 1 mm

apart were made on the white background.

The DCB test was performed at a constant crosshead

speed of 0.5 mm/min until the crack propagated to be 20

mm in length. As soon as the crack length became 20

mm, the test machine was immediately stopped and thenthe specimen was unloaded. The load–displacement

curve was recorded. The propagated crack length was

measured using an optical microscope at the magnifi-

cation of 8. The crack growth energy GIC was calculated

by Eq. (1) based on the linear elastic beam theory [11]:

GIC ¼ 3P 2c

2BoCoa

; ð1Þ

Table 2

A summary of various characteristics of staple and filament type car-

bon fibers prepared for the present study

Property Staple type

carbon fiber

Filament type

carbon fiber

Densitya (g/cm3) 1.75 1.75

Diameterb (lm) 7.37 7.45

Tensile strengthc (GPa) 1.93 2.65

Tensile modulusc (GPa) 181 224

Elongationc (%) 1.9 1.7

Electrical resistivityd (lX cm) 5900 7300

Carbon contente (%) 84 86

aDensity measured by a densimeter based on ASTM D-4018.b Fiber diameter measured by SEM assuming a circular cross-sec-

tion.c Single filament tensile test based on ASTM D3379-75.d Electrical resistivity measured using a two probe method.e Carbon content measured by an elemental analyzer.

J.K. Park et al. / Carbon 42 (2004) 795–804 797

where Pc is the fracture load required to extend crack

length a in the specimen width B and C is the compliance

which is the inverse slope of the load (Pc)–displacement

(d) curve (i.e. C ¼ d=Pc).

2.3. Thermal measurements

The thermal conductivities of spun C/P and filament

C/P composite specimens by employing a comparative

steady-state method against a reference sample were

measured using a tailor-made apparatus according toASTM E1225-87. At a steady-state condition, the ther-

mal conductivity was derived by comparing the tem-

perature gradient between a reference sample and a

target sample. The thermal conductivity apparatus was

described in more detail elsewhere [10]. The thermal

conductivity of a target sample, kt, was calculated using

the following Eq. (2) derived from the Fourier’s Law of

heat conduction [12]:

kt ¼ krDTrDTt

� dtdr

� d2r

d2t

K; ð2Þ

where kr is the thermal conductivity of a reference

sample with the diameter dr, DT and d is the temperature

difference and the distance between two junctions of

thermocouples used in sample, respectively. The sub-

script letters t and r designate the target sample and thereference sample, respectively. K is the correction factor

compensating for the effect of the negligible conditions

(usually, the value of K is within the limits of 1.0–1.02).

The standard deviation of k-measurements using the

thermal conductivity apparatus is less ±6%. The mea-

surements were carried out along the directions parallel

and perpendicular to the laminar plane of a composite,

respectively. The temperature difference between thereference and the target sample was 50 �C in average.

The temperature gradient at each end of samples was

generated by the electrical heater.

The thermal expansion behavior of spun C/P and

filament C/P composites was studied using a thermo-

mechanical analyzer (TMA 2100, Dupont). The speci-

men dimensions were 8 mm · 8 mm · 20 mm in

rectangular shape. All thermal expansion measurementswere carried out up to 800 �C with a purging N2 gas

along the directions parallel and perpendicular to the

laminar plane of a composite.

2.4. Ablation measurements

The ablation test was continuously conducted until

the specimen was completely burnt through showing a

pin-hole by an arc plasma torch flame according to

ASTM E285-80. The test specimens were of a rectan-

gular plate form and the dimensions were 100 mm in

length, 100 mm in width, and 10 mm in thickness. Two

thermocouples were firmly attached with phenolic resin

at the center of both back and front faces of a specimen

in order to record the temperature variations as a

function of time during the ablation test. The distance

and the angle between the front surface of the specimenand the nozzle tip of a plasma gun were 30 mm and 90�,respectively. The flame temperature in the present

ablation testing system was estimated to be approxi-

mately 3000 �C and the heat flux was about 28.4 MW/

m2. The erosion rate was calculated by dividing the

specimen thickness before and after the test by a burn-

through time for each specimen. The insulation index

was obtained from the time reaching the temperaturechange of the back-face of the specimen divided by the

specimen thickness.

Microscopic observations on the fiber and the matrix

of the eroded composites after each ablation test were

performed using a scanning electron microscope (SEM,

Model Philips XL30) to qualitatively support the abla-

tion properties studied.

3. Results and discussion

3.1. Physical characteristics

Table 2 shows various properties of staple and fila-

ment type carbon fibers heat-treated up to 1100 �C. Thedifference in the physical and mechanical properties

between staple and filament type carbon fiber comes

from different carbonization processes. Staple and fila-

ments type carbon fibers were prepared by batch-type

and continuous-type carbonization processes identicallyconducted at 1100 �C, respectively. The filament type

carbon fibers were heat-treated under tension during

low-temperature carbonization process, while the staple

type carbon fibers were heat-treated at free state without

applied tension. This led to a low Young’s modulus

of staple type carbon fiber. The tensile strength and

798 J.K. Park et al. / Carbon 42 (2004) 795–804

modulus of filament type carbon fiber are higher than

those of staple type carbon fiber. The continuous car-

bonization process under tension may lead to rapid

development of the incipient graphite structure due tohigher degree of graphitic order and preferred orienta-

tion in the carbon basal planes and also reduction of

microstructural flaws like voids, inclusions, and surface

imperfection in the resulting fiber [13].

In general, the electrical resistivity of a fiber is in-

versely proportional to the thermal conductivity of the

corresponding fiber. The measurement of electrical

resistivity of a fiber is used as a simple screening test [5].The electrical resistivity of the staple type carbon fiber of

5900 lX cm is about 20% lower than that of the fila-

ment type carbon fiber. The staple type carbon fibers

may be expected to have thermal conductivity slightly

higher than the filament type carbon fibers.

Fig. 1 shows SEM observations on the surface

microstructure of staple and filament type carbon fibers.

It is likely that the surface appearance of the staple typecarbon fiber is very similar to that of the filament type

carbon fiber. The striations along the fiber axis on the

fiber surface can be clearly seen due to strong chemical

reaction resulting in severe shrinkage in both carbon

fibers.

3.2. Interfacial properties

The interlaminar shear strengths of spun C/P and

filament C/P composites measured are approximately 30

and 20 MPa, respectively. The interlaminar shear

Fig. 1. Scanning electron microphotographs of fibers used for textile

composites: (a) staple type carbon fiber and (b) filament type carbon

fiber.

strength of the spun C/P composite is 50% greater than

that of the filament C/P composite. This is due to the

protruded fibers surrounded by the phenolic matrix in

the spun C/P composite. They may significantly con-tribute to improving the fiber–matrix adhesion at the

interfaces between the bulk fiber and the bulk polymer

matrix. In other words, the protruded fibers may be

anchored by the phenolic resin in the interfacial region,

leading to a higher interlaminar adhesion along the

thickness direction in the woven spun C/P composite

than that in the woven filament C/P counterpart.

The interlaminar fracture energy of the spun C/Pcomposite obtained using the DCB test is compared

with that of the filament C/P composite, as seen in Fig.

2. The fracture energy release rate or crack growth en-

ergy, so-called GIC, has been calculated from the load–

displacement curve of each test specimen using the

compliance method [14–19]. The result shows that the

crack growth energy of the spun C/P composite is

obviously higher than that of the filament C/P com-posite over the measuring range of crack length. This

may be due to the presence of the protruded fibers at the

surface of spun yarns, which can resist possible delam-

ination occurring in a composite laminate. In the fila-

ment C/P composite, the crack growth energy gradually

increases with increasing crack length at the earlier stage

of crack propagation and then remains constant above

the crack length of 5 mm. On the other hand, the crackgrowth energy decreases with increasing crack length

and then remains constant at the later stage of crack

propagation. As can be seen, the crack growth energy

required for initiating and propagating the crack in the

specimens is relatively higher in the spun C/P composite

than in the filament C/P counterpart. This is because

strong interfacial bonding and fiber bridging effects in

the spun C/P composite including the protruded fibersmay contribute to restricting the propagation of crack.

The fiber bridging effect is originated from the fiber

migration in the adjacent plies during curing with heat

0 5 10 15 200.0

0.5

1.0

1.5

2.0

2.5

3.0

3.5

Cra

ck G

row

th E

nerg

y, G

IC(K

J/m

2 )

Crack Length (mm)

Spun C/P Filament C/P

Fig. 2. Variations of crack growth energy (GIC) as a function of crack

length for spun C/P and filament C/P composites.

Fig. 3. Scanning electron microphotographs showing the fractured

surface for (a) spun C/P and (b) filament C/P composites after the DCB

test.

J.K. Park et al. / Carbon 42 (2004) 795–804 799

and pressure, contributing to preventing the composite

from possible delamination between the interlaminae.

Also, the crack growth of the spun C/P composite is

more effectively restricted by the increased resistanceagainst the bridged fibers upon fracture.

Fig. 3 compares scanning electron microphotographs

observed at the fracture surface between the spun C/P

and filament C/P composites observed after the DCB

test. The spun C/P composite specimen shows the

irregular distribution of fibers and fiber pullout traces at

the fracture surface. It indicates that the protruded fi-

bers in the spun C/P composite apparently contribute toresisting delamination not only by a reinforcing effect

but also by a fiber bridging effect. The fiber breakage

seen in Fig. 3(a) also acts as an energy absorbing

mechanism, as reported in a literature [16]. Therefore,

the crack front absorbs higher energy, which can then be

released on propagating the crack at a faster speed and

possibly in a more brittle mode. In the filament C/P

specimens, the relatively clean fracture surface can beobserved. Here, the reinforcing fibers are separated from

the matrix with a furrow-like pattern. Also, broken fi-

bers and a number of fiber hackles can be found in Fig.

3(b). During the formation of fiber hackles, the crack is

normally initiated and propagated along the weak point

of the matrix when shear force is applied to the matrix.

The extended propagation of crack is continued pro-

ducing an irregular shape of the matrix fracture‘‘hackle’’ coincident with or opposite to the direction of

propagation. In general, the hackle pattern may be fre-

quently found in a tough matrix undergoing ductile

deformation [19]. However, in carbon fiber reinforced

Fig. 4. Scanning electron microphotographs showing the crack growth propa

C/P composites after the DCB test.

phenolic composites, it is also observed that there is a

sort of small hackle pattern due to brittle polymer ma-

trix surrounding the fibers, as found in this work.

Fig. 4 shows scanning electron microphotographs

observed at the cross-section of the crack path propa-gated along the laminar direction for spun C/P and fil-

ament C/P composites after the DCB test. The cracks in

gated along the laminar direction for (a,b) spun C/P and (c,d) filament

800 J.K. Park et al. / Carbon 42 (2004) 795–804

all the specimens are propagated along the intralaminar

(Fig. 4(a) and (b)) or interlaminar (Fig. 4(c) and (d))

direction upon fracture. In the spun C/P composite

specimen, the fracture behavior such as fiber pullout andfiber breakage has been clearly observed. Also, occa-

sional crack-deflection and crack-bifurcation have been

found. It indicates that the protruded fibers apparently

contribute to restricting the crack path formation and

consequently possible delamination by both fiber rein-

forcing and fiber bridging effects. In the filament C/P

composite specimen, the relatively straight crack path

has been observed. It indicates that the fracture patternreflects the occurrence of matrix debonding. Therefore,

it is concluded that the interfacial properties of the spun

C/P composite in terms of interlaminar shear strength

and crack growth energy are significantly greater than

those of the filament C/P composite, with qualitative

support of fracture surface observations.

Fig. 5. Optical photographs showing the cross-section of (a) spun C/P

and (b) filament C/P composites (·100).

3.3. Thermal properties

Table 3 summarizes the longitudinal and transverse

thermal conductivities and the erosion rate measured for

spun C/P and filament C/P composites. Here, the lon-gitudinal and transverse thermal conductivities are de-

fined as the thermal conductivity measured along the

directions parallel and perpendicular to the laminar

plane of each composite, respectively. As tabulated, the

longitudinal thermal conductivity of the spun C/P

composite is about 7% lower than that of filament C/P

composite. This turns out that the heat transferred from

spun yarn structure may be more or less discrete. Theresult implies that some structural modification of

reinforcing spun yarn may be desirable for decreasing

the thermal conductivity of a phenolic composite. On

the other hand, the transverse thermal conductivity

of the spun C/P composite is same as that of the filament

C/P composite. One possible explanation on the trans-

verse thermal conductivity behavior is as follows. A

microscopic observation on the distribution of the fibersand the matrix in the composites provides an evidence

for understanding the difference in the transverse ther-

mal conductivity. Fig. 5 shows that the resin-rich region

of the spun C/P composite is significantly lesser than

that of the filament C/P composite as well as a number

of the protruded fibers exist in the resin-rich region of

Table 3

A summary of the thermal conductivity and the erosion rate measured

for spun C/P and filament C/P composites

Spun C/P Filament C/P

Thermal conductivity (W/mK)

Longitudinal 1.75 1.88

Transverse 0.76 0.76

Erosion rate (mm/s) 0.089 0.103

the spun C/P composite. It has been known that the

thermal conductivity of a polymer composite increases

with increasing the quantity of carbon black in the

matrix as reported in a literature [20]. A number of theprotruded fibers of spun yarn may somewhat contribute

to transferring the phenolic matrix from an insulation

material to a semi-conductive material, as similarly

found in the carbon black case.

Fig. 6 shows the variations of the coefficient of ther-

mal expansion (CTE) measured along the parallel

(longitudinal) and perpendicular (transverse) directions

to the laminar plane of spun C/P and filament C/Pcomposites as a function of temperature. The spun C/P

composite measured along the longitudinal direction

gradually expands up to 450 �C, contracts in the range

of 450–650 �C, and then expands again above 650 �C.The thermomechanical behavior is predominantly gov-

erned by the phenolic matrix in the composite. The

contraction is due to thermal degradation of the phe-

nolic matrix. In the case of spun C/P composite, theshrinkage of matrix resin occurs predominantly since the

resistance against the shrinking force of phenolic matrix

occurring upon thermal degradation is relatively weak.

0 100 200 300 400 500 600 700 800

0.0

0.2

0.4

0.6

0.8 Spun C/PFilament C/P

CTE

(x10

-6/o C

)

Temperature(oC)

0 100 200 300 400 500 600 700 800-200

-150

-100

-50

0

50

Spun C/PFilament C/P

CTE

(x10

-6/o C

)

Temperature (oC)

(a)

(b)

Fig. 6. Variations of the coefficient of thermal expansion (CTE) for

spun C/P and filament C/P composites: (a) longitudinal and (b)

transverse.

100 200 300 400 500 600 700 800

123456789

1011

Insu

latio

n In

dex

(s/m

m)

Back-face Temperature of Specimen (oC)

Spun C/P Filament C/P

Fig. 7. A comparison of the insulation index as a function of back-face

temperature for spun C/P and filament C/P composites measured

during the ablation test using an arc plasma torch.

J.K. Park et al. / Carbon 42 (2004) 795–804 801

The transverse thermal behavior of both composites isquite different from the longitudinal cases, showing a

large variation of negative CTE value in the range of

450–650 �C. This is definitely explained by the shrinkage

of the phenolic matrix due to thermal degradation

without significant restraint of reinforcement in the

transverse direction. The spun C/P composite exhibits

lower shrinkage behavior in the transverse direction

than the filament C/P composite due to a bridging effectby the protruded fibers, as described earlier.

Front-face

Back-face

Front-face(1) (2)(3)

(4)Specimenthickness

Cutting Plane

Burn-through hole

100mm

100mm

Fig. 8. Schematic illustration of a cross-section of an eroded specimen

indicating four different sections for SEM observations.

3.4. Ablation properties

The erosion rates obtained for spun C/P and filament

C/P composites after the ablation test are summarized in

Table 3. The spun C/P composite exhibits about 14%

lower erosion rate than the filament C/P counterpart. As

shown in Fig. 7, the insulation index of the spun C/P

composite is higher than that of the filament C/P

counterpart over the whole range of the back-face

temperature of each specimen. Here, the insulationindex is defined as the time required to achieve a certain

back-face temperature from the surface temperature

prior to ablation test divided by the average thickness of

each specimen. Therefore, it is expected that the higher

the insulation index the greater the ablation resistance.

The result indicates that the spun C/P composite re-

quires about 54% longer time in second to achieve 100

�C than the filament C/P composite, showing the similar

trend at higher temperatures. This may be explained onthe basis of thermal diffusivity, which can be expressed

by K=q � Cp [21,22]. Here, K is the thermal conductivity,

q is the density, and Cp is the specific heat capacity of a

composite. The expression reflects that the smaller the

thermal diffusivity the better the heat insulation. When

specific heat capacities of two composites are constant, a

decrease of the thermal conductivity of a composite

material leads to a lower value of thermal diffusivity.The insulation index on the basis of thermal diffusivity

may result in different thermal conductivities of two

composites.

On the other hand, the insulation index may impor-

tantly influence the difference in the erosion rate of

composites. The following descriptions also support the

result. The times needed for reaching the back-face

temperature of 300 �C are approximately 74 s for thespun C/P composite and 58 s for the filament C/P

composite, respectively. The erosion depth is calculated

to be 3.4 mm for the spun C/P composite and 5.9 mm

for the filament C/P composite from the erosion rate

given in Table 3. The thickness of each specimen mea-

sured before the ablation test is identically 10 mm in

average. The sample thickness remaining after the back-

802 J.K. Park et al. / Carbon 42 (2004) 795–804

face temperature of 300 �C is reached is 6.6 mm for

the spun C/P composite and 4.1 mm for the filament C/P

composite, respectively. This may be caused by the

difference in the rate of back-face temperature change.Therefore, it may be concluded that the insulation index

of the spun C/P composite is greater than that of the

filament C/P composite because the erosion rate of the

former is lower than that of the later.

Fig. 9. Scanning electron micrographs showing ablation patterns of the fiber a

of spun C/P and filament C/P composites. Numbers (1), (2), (3), and (4) des

designate spun C/P composite and filament C/P composite, respectively.

After the ablation test, the eroded specimens are di-

vided into four different sections depending on the dis-

tance from the central region of the test specimen of

100 mm · 100 mm where the flame passed through itmaking a burnt-through pinhole to the outer region of

the specimen where the flame is less effective. Here, the

difference in the distance reflects the difference in the

degree of ablation. As depicted in Fig. 8, the sections (4)

nd matrix at the eroded surface from the center of specimen to the edge

ignate different locations in the eroded specimen. The letters S and C

J.K. Park et al. / Carbon 42 (2004) 795–804 803

and (1) designate the central and outer regions in a

tested specimen, respectively. The flame is expected to be

most intensive near the section (4) and least intensive

near the section (1). Even after the ablation test, nodelamination has been observed in all the specimens.

The appearance of each tested specimen shows a typical

ablation pattern formed by flame of high temperature,

high velocity, and high pressure simultaneously.

Through scanning electron microscopic observations

on the eroded faces of spun C/P and filament C/P

composites after the ablation test, the ablation pattern

has been closely examined, as shown in Fig. 9. It hasbeen found that the ablation of all the reinforcing fibers

in the composites has a needle- or an icicle-shaped

pattern formed by simultaneous combination of ther-

momechanical, thermochemical, and thermophysical

effects [8]. The fibers between the cross-plies exhibit a

necking pattern of a single filament. Also, some fibers

are broken away, as seen in S(1), S(2) and C(1) in Fig. 9.

In these sections, a number of fibers exist with a charredmatrix. The fibers in a discontinuous matrix region are

already broken away or almost broken showing an icicle

shape. This is probably ascribed to the thermochemical

erosion like oxidation and sublimation. It is also found

that the alignment of the eroded fibers is not regular in

the spun C/P composite but relatively regular in the

filament C/P composite. The matrix part between the

cross-plies is also broken away showing deep cracks andprogressive erosion with the fibers, as seen in C(1) and

C(3) in Fig. 9. The fibers in the lost matrix may be

broken away upon removal of the weak matrix. Such the

ablation behavior has not been found in the spun C/P

composite specimens.

Severe microstructural damages in carbon fibers as

pitting and digging are rarely observed in the both

composites, but only sharp and smooth fiber surface canbe found. In the section (4), which is most severely ex-

posed to the plasma flame, the matrix layer is rarely

found, and fibers are uniformly eroded forming the

needle shape. In the spun C/P composite, it has been

found that the protruded fibers exist with or without the

charred matrix between the cross-plies. The results

suggest that the protruded fibers in the spun C/P com-

posite somewhat affect not only on the thermal insula-tion but also on the ablation behavior with reduction of

abnormal erosion pattern such as spalling, pocketing

and/or ply-lifting.

4. Conclusions

The interfacial, thermal, and ablative properties of

phenolic composites reinforced with spun yarn type

carbon fabrics and filament yarn type carbon fabrics

heat-treated at 1100 �C have been compared in terms of

interlaminar shear strength, crack growth rate, fracture

pattern, thermal expansion behavior, insulation index,

erosion rate, and ablation pattern.

It has been concluded that the interfacial properties

of the spun C/P composite are significantly greater thanthose of the filament C/P composite, with qualitative

support of fracture surface observations. The improved

interfacial property may be explained to be due to a fiber

bridging effect by the protruded fibers located in the

phenolic matrix of the spun composite. The longitudinal

thermal conductivity of the spun C/P composite is about

7% lower than that of the filament C/P counterpart. The

result implies that some structural modification ofreinforcing spun yarn may be desirable for reducing the

thermal conductivity of a phenolic composite for ther-

mal insulation applications. The spun C/P composite

exhibits lower shrinkage behavior in the transverse

direction than the filament C/P composite due to a

bridging effect by the protruded fibers. The ablation

results indicate that the insulation index of the spun C/P

composite is greater than that of the filament C/Pcomposite because the erosion rate of the former

is lower than that of the later. The qualitative infor-

mation from microscopic observations also supports the

quantitative results obtained in the present study.

Consequently, all the results obtained from this work

suggest that use of spun yarn type carbon fabrics

as reinforcement in a phenolic composite may signifi-

cantly contribute to improving the interfacial, thermal,insulative and ablative properties of C/P composite as

well.

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