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: [email protected] (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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