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CAHBON ANI) GRAPHITE MATRICES IN CARBON-CARBON COM POSITE S:
AN OVERVIEW OF THEIR FORMATION, STRUCTURE, AND PROPERTIES
Gerald Rellick
Mechanics and Ma terials Technology Cente r. Technology
0
erations
The Aerospace Corporation,
PO.
Box 92957,
Los
Angeles, E A 90009
Keywords: Carbon-carbon composites. Graphitization. Composite properties
WHY CARBON-CARBON COMPOSITES?
Carbo n-carbo n (CIC) composites. so-called because they combine carbon-fiber reinforcem ent in an all-
carbon matrix. can best be viewed as part of the broader category of carbon-fiber-based composites,
a l l
of
which seek
to
utilize the light weight and exceptional streng th and stiff ness of ca rb on libers. However, in CIC.
the structural benefits
of
carbon-fiber reinforcement a re combined with the refractoriness
of
an all-carbon
materials system. mak ing CIC com posites the material
of
choice
for
severe-environment applications. such as
atmo spheric reentry, solid rocke t motor exhausts. and d isc brakes in high -perfor manc e military an d comm er-
cial aircraft. Their dimensional stability. laser hardness, and low outga ssing
also
make them ideal candidates
for various space stru ctural applications.
Such mechanical and refractory properties are not met by the various bulk graphites for t w o reasons:
(1)
graph ites arever y flaw sensitive and . therefore. brittle: and
(2)
graph ites ar e difficult to fabricate in to large
sizes and complex shapes. The se difficulties are hrgely overcome by taking adva ntage of the "two phase prin -
ciple
of
material structure and strength
[I]."
In the classical two-phase mat erials system. or compos ite. a high-s trength . high-m odulu s, discon tinuou s-
reinforcement phase is carried in a low-modulus. contin uous- matrix phase: e.g.. grap hite fibers
i n
a thermo-
plastic-resin matrix. The stre ss in a composite structure having fiber reinforcement that is continuous in
length. is carried in proportio n to the moduli of the constitu ent phases. weighted by their respective volume
fractions. 'merefo re. the much stiffer (highe r-mod ulus) fibers will
y
tlie princip al load bearer s. and the ma-
trix. in addition to havingthetaskof binding togetherthccomposite.will deform under load and distribute the
majority
of
stress
to
the fibers. At the same time. becau se the brittle carbon fibers arc rso/&4 tlie poss ibility
that an individual fiber failure will lead to propagation and catastropbic failure is practically eliminated.
Ano ther major benefit of com posites is that they permit the construc tion of complex geom etries. and
in
such a way that different amou nts of the load-carrying fibers can b e oriente d in specific directions t o accom-
modate the design loads
of
the final structure. Closely assoc iated with this "tailoring" f eatu re of com posites is
that carbon-fiber technology enables exploitation of the exceptional basal-plane stiffness (and strength.
i n
principle, although this is still much far ther from realization) of sp2 bonded carbon atoms-Le., the fibers are
not
isotropic. but rath er have their g raphite bass1 planes oriented preferentially in the fiber axial direction.
For
very-high-temperature carbo n-fiber -com posite applic ations. say. abov e 2000C. even for brief peri-
od s of time, it is necessary to employ a carb on m atrix: however, like the fiber. the carbo n matrix is also brittle.
When fiber-matrix bond ing is very stron g
i n
CIC. brittle fracture is frequently obs erved. The explanation is
that strong bonding permits the development of high crack tip stresses at the fiber-matrix interfac e; cracks
that i nitiate in either fiber or matrix can then propaga te through the com posite. However. if the matrix
or
the
fiber-matrix interface is very weak. or microcracked. then the primary advancing crack can be deflected at
such weakened interfaces
or
cracks . This is the Cook-Gordon theory [2] for streng thening of b rittle solids.
which states. more specifically. that
if
the ratio ofth e adhesive strength of the interface to the general cohesive
strength
of
the solid is in the right range. large increases in the strength and toughness
of
otherwise brittle
solids may re sult. There fore, good fiber strength utilization in a brittle-matrixcomposite like CIC depends on
control of the matrix and interfacial structures.
The objective of this pa per is to provide a brief overview of carb on an d g raph ite matrices in CIC. with an
emphasis
on
recent research
on
some of the more fun dam ental m aterials issues involved
13-71,
Muc h of what
is presented is tak en from our own published work, which has focused
on
understanding how the structure
of
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In
addition to orientation, another important feature of carbon matrices is their g raphitizab ility, which is
a measure of the eas e in converting the pyrolyzed carbon matrix product into crystalline graphite through
high-temperature heat treatments in the
-2000-3000'C
range. Th e state of graphitiz ation can b e assessed by a
number of techniqu es, the most co mm on of which is X-ray diffraction (XR D). However, in C/ C it is usually
very difficult to resolve th e result ant compo site diffraction response in to the respectiv e fiber and m atrix re-
sponses, because both phases are carbon. A technique to circumvent the samp le volume problem is laser
Raman microprobe spectroscopy (LRMS). Although the interpretation of the Ram an spectra is more ambig-
uous
than with XRD , LRM S permits focusingof avisible-light beam, as small as 1 im in diameter,
on
aregion
of th e specimen while recording the Ram an spec trum, which is active in carbon
[u)] .
Useful struc tural infor-
mation on a local scale ca n b e obtain ed in principle.
O ne majo r difficulty with applyin g LRM S to composite s is that the size of constitue nt phas es is of the
order o f microns. making it necessary to prepare the specim ens for examination using st and ard optical polish-
ing techniques. Such polishing tends to damage the near-surface structures and leaves behind a thin layer of
polishing debris. Since the p robe depth of the optical beam is only abo ut 50 nm
[ZO],
he Raman spectrum
unfortunately becomes a function
of
the preparation technique
[21-231.
A techniqu e we have employed extensively and with good success, and which is an outgrow th of early
work performed at Los Alamos L aboratories [24.25], involves SEM examination of specim ens that have been
polished an d then cathodica lly etched with xenon. When th e carbon struc ture is graphitic, and when the
graphite layer planes are oriented perpendicular to the plane
of
sect ion , we see, typically, a pronounced lamel-
lar
texture,
as
revealed for the inner- and outerm ost CVI layers in the C /C
of
Fig.
3.
Th e lamella r texture is the
result of differential etching rat es of the various microstructural units, the exact nature of which is still not
clear. Th e most likely mechanism is preferential removal a t lower-density, less-o rdered intercrystalline-type
bounda ries th at sep arate regions of good crystalline registry; this is seen very dram atically in highly oriented
pyrolytic graphit es reac ted in oxygen [26 27]. he techniq ue is effective, principally, in distinguis hing broadly
between graphitic and nongraphitic carbon
on
a scale of microns.
Returning to Fig. 3, this particular specimen has the CVI deposition sequence R U S U R L (as determined
separate ly from polarized-light microscopy) and has been heat- treated to 2500C for
1
hr.
The
lamella r tex-
ture of the RL zones indicates their graphitized structure. whereas the absence
of
significant texture in the SL
zone indicates that the SL struc ture is essentially glassy carb on. This observation was confirmed by XRD ,
LRMS , and by selected physical-p roperty meas urem ents [B] he effect of having
a
graphitic and well-ori-
ented matrix is illustrated by the higher therm al conductivities for heat-treated R L composites shown in Fig. 4.
Modulus enhancement is another interesting effect of
a
well-oriented, graphitic matrix (Fig.
5).
For the
particular pseudo-3-D. fe lt-based C/C co mpo site of the figure. there were two CV I densifications . Following
the first, the composite stru ctu rew as heat-treated to 2500'C; the second CV I was left int he as-deposited state
( - loOo-1200"C). Th e relative proportions of the first and secon d CV I varied with each sp ecim en, bu t the total
CV I weights were approximately the sam e. The fiber volume (and weight) fraction was cons tant ( -u) ) for
each composite.
The strong dependence
of
the modulus on the relative proportion
of
heat-treated C VI indicates that the
carbon m atrix can carry a significant fraction of the load, particular ly, in this case, if the structure is heat-
treated to typical graphitization temperatures. Th e modulus-enhancement effect by the m atrix
is
especially
striking in this composite beca use of the use
of low-modulus fibers at fairly low volume fractio ns, However, as
will be seen, this effect is
an
importan t m aterials and processing consideration in all CIC composites.
Coal-Tar
and
Petro
eum
Pitches
Th e second method for CIC densification
is
the use of coal-tar and petroleum pitches. Because they are
thermo plastic, pitches a re used mostly for redensification: Le., furt her densifying of
a
C/C structure that has
been "rigidized" by
an
earlier impregnationldensificationste p (e.g., a resinimpre gnated fabric preform) or
that has sufficient rigidity from the friction between the elements of the woven structure ( e g , 3-D braided
preform).
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Pitches a re unique i n passing through a liquid-crystalline ransformation at tempe ratures between about
350 an d 55OC [29]. In this transform ation, large lamel lar molecules form ed by the reactions of thermal crack-
ing and aromatic polymerization are aligned parallel to form
a n
optically anisotrop ic liquid crystal known as
the carbonaceous mesophase
[30].
T he alignment of the lamellar molecules is the basis for easy thermal gra-
phitizability of the carbonized product. On e
of
the features
of a
meso phase-b ased matrix is high bulk density,
which is achievable because the matrix density can a ppro ach the value for single-crystal graphite , 2.26 g/cm3.
The topic of pitch impregnation and densification of C IC intro duces the subject of densification efficien-
cy, th e most meaningful me asur e of which is volumetric densifica tion efficiency [31]. It is the ratio of thevol-
ume of carbon matrix in a process cycle to the volume of porosity available for densification.
For pitches carbonized a t atmosp heric pressure, coke yields are of the order of
50-60%,
impregnant den-
sities
are
- 1.35 g/cm3. and , as we have noted, densities for pitch-derived matrices ar e -2.2 g/cm3. From these
values we calculate volum etric densificatio n efficiencies
of
only 3 0 4 0 % at atmosp heric pressure [31]. By re-
sorting to so-called hot isostatic-pressure-impregnation-carbonization (HIPIC), to pressures of about
15,000psi, carbon yields of pitches can be increased to almost
90%
[lo]. But even with HIPIC, volumetric
filling is only
55 .
There fore, given
a
preform with initial porosity of 45%, typical for man y 3-D woven struc-
tures, three cycles at maximum
densifico/ion efficiency
would be required to reduce the porosity to 4%. With
curr ent HIP IC pro cedure s, however, it is found tha t at least five cycles at
15,000
psi are requi red to achieve this
sa m e level of porosity. Su ch reduce d efficiency in real systems is the result of forced expulsion of pitch from
the preform as a result of the gas-forming pyrolysis reactions accom panying carbonization.
Clearly, one way
to
incr ease efficiency, for a given weight-based carb on yield, is to select either
an
impreg-
nant o r an HTT that will lower the final matrix den sity. As will be seen in the next sub section, lower-density
carbo n matrices
can
be achieved by using resin precursors that form
a
glassy-carbon-type structure. But,
although this approach fills more of the available space,
it
does
so
with a lower-density carbon matrix, which is
different in struc ture from the higher-density graphitic matrix. Th e trade-offs in pro perties, particularly me-
chanica l, ar e not well under stoo d. We will touch on this topic again in the next subsection.
Appro aches to improving densification efficiency
of
pitch-based matrices without resorting to HIPIC
processing include the use of heat-treated and solvent-extracted pitches [32] and partially transformed (to
meso phase) pitc hes [33,34]. A novel app roa ch, dev eloped by W hite a nd Sheaffer 1351, is to oxidatively stabilize
the m esophase following impregnation and transformation,
an
approach similar to that employed in meso-
phase-fiber stabilization. Th e result is a hardene d mesophase that is resistant to the bloating effects of
pyrolysis gases but that, upon further heat treatment, yields
a
dense, graphitic carbon.
Th e strong orienting effect of the fiber surface
o n
the large lamellar m esopha se molecules is
an
interest-
ing feature ofmesophase formation in C /C composites. This effect was demonstrated by the work of Zimmer
an d Weitz [36], who used p olarized-lig ht microscopy to show that m esopha se molecules near
a
fiber sur face in
a
close-packed fiber bund le always aligned parallel t o the fiber surfac es, even in the pres ence of stron g magnet-
ic fields. S inger and Lewis dem onstrated earlier that m agnetic fields would orient mesophase molecules in
bulk mesophase [37]. Zim mer
and
Weitz showed that mesophase would also orient in matrix-rich regions
within the fiber bundles-i.e.,
at
points far removed from fiber surfaces
[36].
They calculated a magnetic
coherence length of 7 Fm, which correspon ds roughly to the distan ce over which the orien tation effect acts.
Such localized o rientation in the liquid-crystalline state w ould lead o ne to expect the final, graphitized
matrix also to be well oriented in the imm ediate vicinity of he fiber. First observed by Evangelides
[38]
using
SEM
in conjun ction with xenon-ion-etching, suc h
a
matrix sheath effec t is depicted in Fig.
6
in
a
coal-tar-
pitch-densified C/C.
Modulus enhancement in pitch-based C/ C has been widely reported, b ut whether the effect is due to the
matrix or t o an increase in the fiber m odulus, resulting from high-temperature heat-treatment-induced struc-
tura l changes in the fiber, has not been clarified [39]. Th e sheath effect is
also
pronounced in resin-based
car bon matrices, but for d ifferent reasons, which we will examine in the next subsection.
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Matrix microcracking is characteristic of all CICs, bu t it is particularly prevalent in graphitic matrices
because of the comb ination of weak shear planes in polycrystalline graphite an d the the rma l stresses gener-
ated during heat treatm ent (Fig.
7)
[40,41]. Microcracking also has importan t effects
on
the engineering pr op
erties ofC /C materials-particularly the matrix-dominated properties in the unreinforced directions, such as
the interlam inar shear strength and perpendicular-to-ply tensile strength in 2-D C/C lam inates. However, as
mentioned above, such m icrocrack ing app ears
to
improve in- plane flexural and tensile strength, by way of a
Cook-Gordon mechanism [42 45] .
Th e third, and last, class of C/ C impregnant to be discussed is thermoset resins, which are th e basis for
prepreg fabric and tapes, a s noted above; resin systems c an also be used for reimpregnation.
In
addition to
their easy fabricability. therm osets have the a dvantag e of charring-in-place; th at is, although they softe n and
deform on heating, they d o not fuse o r iquefy, and, therefore, no special tools or techniqu es must be employed
to retain the matrix in the composite during pyrolysis.
Therm oset resins are usually highly crosslinked, which makes them resistan t to therm al graphitiz ation n
bukform, even to temperatures of 3000C [5,46]. P henolic resins are currently m ost commonly used for pre-
preg operations, whereas furan -bas ed resins ar e used mo re for reimpregn ating. B oth have char yields typically
in the
5040%
range.
Th e development
of
ultra-high-char-yield resins derived from polymerization of diethy nylbenz ene DEB)
[47-511, usually term ed polyarylacetylenes (PAA) [47], has received mu ch focu s in recent years. T h e stru ctur e
of DEB is illustrated in Fig. 8. along with a synthesis route tha t involves a catalytic cyclotrimer izationprepoly-
merization in methyl ethyl ketone solvent [48.49]. T he cyclotrimerization l iber ates m uch
of
the exothermic
heat of polymerization, thereby allowing safe, controllab le curing. T he principa l appea l of PAAs is th eir ex-
tremely high char yield. From th e average structure, we calculate a theoretical carb on yield of ab out 95 ; in
practice, PAA s can have carbon yields of 90% to 700C. although mor e practical formu lations employing
monofunctional chain terminators to improve flow properties reduce this yield to about
85
[48,49].
Similar to other crosslinked therm osets, PAAs prod uce largely nongra phitizable carbo ns. To extend the
range of matrix struc tures for this fabricable resin system, we have been exploring approa ches t o in siru matrix
catalytic graphitization in C /C in
our
laboratory On e promising approach, by Za ld iv ar et d. [52], has been the
use of boron in the form of a carb oran e compo und. Figure 9a is a plot of room-temperature tensile strength
of
undoped and boron-doped unidirectional C/C s versus HTI ; the strengths a re calculated relative to the fiber
cross-sectional areas on the assum ption that the matrix carries negligible load relative to the fibers. The
strength of the fibe rs in the cured- resin comp osite is taken to be th e value for full stren gth utilization. T he plot
illustrates a number of im port ant features. First, for the und oped syste m, strength exhibits a large decrease as
the comp osite proceed s from cu re to carbonization , owing to conversion of the co mp liant polymer matrix into
a well-bonded, low-strain-to-failure carbo n matrix. Increasin g boron levels lead to incre ased strength utiliza-
tion for the 1100C
H T I
samples: T he undoped specim en behaves a s a monolithic solid and fractures in a
planar-catastrophic mode (Fig. loa); the
5%
B-do ped samples exhibit extensive fiber pullout (Fig. lob), which
indicates a weakened interface. Th e reasons for the weakened interfac e ar e unclear, since X-ray diffraction
revealed no significant difference in graphitization between doped and undoped specimens after this
HIT
At higher H T Is , the use of higher boron levels leads to a red uction i n strength utilization(and an increase
in modulus; Fig. 9b), du e tocatalytic graphitiza tion of the fiber. Fur ther H To f th e undop ed spec imen s beyond
1100C reclaims much of the lost fiber strength, for the reaso ns discussed above. Mo re work is needed to
define the mechanisms by which catalytic graphitiz ation of the matrix affects the properties of C/C.
We recently reported a striking modulus enhancem ent for the same type of 1-D composite studies (Zaldi-
var
el
a/. [7]), using four m esopha se-based fibe rs from D uPon t and PAA resin (Fig. 11). The number in the
fiber designation is the axial tensile mo dulus, i n Mpsi. For
HT
to 2750C. all the composites exhibit sharp
increases in f iber moduli, t o values exceeding 150 Mps i, which is the theoretica l limit
of
the graphite basal-
plane modulus. Since the moduli are calc ulated relative to the original fiber cross -sectional areas, su ch values
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indicate that the com posite mod ulus m ust have significant contributions from the matrix. An example of
a
matrix sheath that may be contributing to the composite modulus is shown in Fig. 12.
This figure brings us to the su bject of stress-induced orientation an d graphitization in otherwise nongra-
ng carbo n m atrices in C / C . While the phenomen on of stress graphitization of hard carb ons ha s been
noted for some time [24,25,53,54], only recently have serious efforts have been made to understand the
physical-mechanical mechan isms involved in matrix orientation and graphitization in C / C [5,6,55]. This topic
isof more than academic interest, because the formation
of
a two-phase m atrix of graphitic and nongraphitic.
oriented and unoriented, zones ca n have a major influence
on
the m echanical properties
of
CIC composites.
To examine preferred orientation in thermosetting-resin-impregnated matrices, cross sections
of
C/C
tows fabricated from Amoco T50 PAN-based fibers and a
PAA
resin were polished, then heat-treated to
2900'C for 1h r and xenon-ion-etched. T he polarized-light microg raph of Fig.
13a
reveals that in addition to
the pronounced lamellar zones, the smoo th-appearing zones-which, by definition, have formed no observ-
able texture with etching-are nevertheless oriented,
as
evidenced by the polarized-light extinction contours
sweeping across the surface
of
the sample
as
the analyzer is rotated.
We conclude tha t even the thickest
(> -20 am matrix regions in this specimen a re oriented. Pronounced optical anisotropy in the matrix for the
same composite heat-treated to onIy 1200 "C is revealed by Fig. 13b.
As
expected, etching produced
no
lamel-
lar texture for this low HTT
The highly localized na ture of the com bination of stress-induced orie ntation an d graphitization is one of
its mo re intere sting features; Le.. all
of
the carbon matrix in the specimen of Fig.
13a
is oriented to some de-
gree, yet only certain discrete reg ions becom e lamellar graphite u pon
HT
to 2900'C.
SEMs
of ion-etched
specimens reveal this localized graphitization more clearly (Fig.
14);
particularly striking
is
the shrinkage of
the matrix away from the fiber, which is
a
result of the volume decrease accomp anying graphitization.
TEM is an extremely effective technique for studying the
local
structure on an even finer scale.
In
the
transverse (Fig. 15a) and longitudinal (Fig. 15b) bright-field images of thin s ections of a T50 fiberlresin-
derived C / C heated to 2750 C. crystallite formation and orienta tion a re evident, particularly in the transverse
section (com pare with Fig. 13a). Selected a rea electron diffraction confirm ed the highly crystalline struc tureof
the interfilament matrix regions [56].
In the SEM of Fig. 16a, we observe a n interesting effect: A t the interstice of five continguous fibers there
is no lamellar formation in the m atrix pocket, except perhaps imme diately adjacent to th e filament surfaces.
This effec t was typically observed in close-packed groups of three t o five fibers. In contr ast, in m ore extensive
matrix regions-for example. those tha t bound two relatively fiber-rich areas, and where the matrix bound-
aries are fairly straight-we observe relatively unimped ed developm ent of lamellar structur e over a dista nce of
several microns (Fig. 16b). Such lamellar development is particularly striking at th e extreme outsid e of the
single-tow specimen s where quite thick
( -
1-2 fiber diame ters) lamell ar zones form (Fig. 16c). In Fig . 16d, we
see that an interruption in the uniformity of the interface between this outer m atrix cru st and the composite
leads to
a
transition from the lamellar
to
nonlamellar structu re.
Further microstructural features not seen in polished specimens are revealed in the SEM of a tensile-
fractu re surfa ce (Fig. 17).
The
lamellar regions in the matrix a re still evident, and th e PAN-based fibers show
their typical fibrillar struct ure. But we now obselve in the matrix both lam ellar and fibrillar textures, the latter
resembling that seen in the T 50 fibers, which a re generally considered to be oriented glassy carbon [46].
Two observations sugges ted to
us
that the key factor in determ ining lamellar-structu re formation in
a C/C
compo site matrix is
a
multiarial deformationof the resin du ring its pyrolysis to carbon. First, consider that, in
normal PAN-fiber manufacture, which leads to a fibrillar structure, the filaments ar e subjected to a uniaxial
tensile Stress durin g oxidation stabilization. However, when carbonized w ithout prior oxidation stabilization,
but in
very
th in
sections, such
as
between the layers of montm orillonite clay, PAN ha s been sh own to yield a
single-crystal structu re following subsequ ent graphitizatio n heat tre atm ent [57 econd, in partially oxidized
(through-the-th ickness) PAN fibers. the unoxidized, fusible core can form lamellar carb on [58].
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In both examples, the m echanical restraints imposed on the PAN during its pyrolysis would be expected
to produce multiaxial deformation. In this critical regime, a number of stresses act at the fiber-matrix inter-
face, assum ing good fiber-matrix adhesio n: First, there is an axial tensile stres s that acts
o n
the matrix; it is
a
consequence of the large matrix pyrolysis shrinka ge, and the high axial modulus an d low axial therm al expa n-
sion of the fiber. This matrix shrinkag e alsogenerates two additional matrix stresses in the plane perpendicu-
lar to the axial direction -a com pressiv e stress, which acts radially, and a tensile stress, which acts circumfer-
entially.
We tested this hypothesis by performing
a
linear elastic plane-strain thermal stress analysis for three
different local fiber-matrix com posite configurations: a clustered arrangement of three fibers and four fibers,
sketche d in Figs. 18a and b, respectively, and a matrix with free boundaries. These three cases correspond
closely to those seen in Figs. 16a-d. Th e material prope rties used for the
PAN
fiber and phenolic-resin matrix
ar e typical values obtaine d from
a
variety
of
source s. Th e mechanical prop erties of the pyrolyzing matrix are
those reported by Fitzer and B urger [59]. The therm al environm ent was a heatup from room tempe rature to
1M)O C.
In the analysis we ar e concerned only with the stresses in th e matrix in the plane p erpendicu lar to the fiber
axis, because th e tensile s tress of the matrix in the fiber directio n at any point i n the matrix is clearly m ore or
less constant at a given temperature owing to the plane-strain consideration. The stresses in the radial-tangen-
tial plane may vary signficantly, dependin g
on
their relative location
to
the fiber. At any point in the matrix,
therefore, we have a state of triaxial stress.
The development of lamellar structure in the matrix was postulated to be favored by two factors:
1)
large value of the maximum tensile stress in the plane, an d (2) a small value of th e ratio of minimum-to-maxi-
mum p rincipal stress in this same plane. Th at is. for a given value of m aximum tens ile stress in the matrix,
lamellar formation
is
favored more when the minimum-to-maximum stress ratio at any location is either small
or negative (Le., compressive). These two parameters may vary with the fiber spacing and boundary condi-
tions, e.g., constrained
or
free edge.
Figure 19a, a plot of principa l stress orientation and relative stress magnitudes, in dicates th at the maxi-
mum stress adjacent to the outside diameter of each fiber is dominate d by hoop tension with a very low leve l of
radial tensile stress; by contr ast, the maximum stress i n the center of the pocket is equal to abo ut one-third that
at the fiber surface, and the minimum (tensile) stress is now significant. From our hypothesis, these two fac-
tors will work in the direction of reduced lamellar formation relative to that at the fiber surface.
The effects of an increa se in the a/r ratio (Fig.
18)
are to decrease the maximum hoop stress
at
the fiber-
matrix bound ary and increase the stress ratio in the pocket region.
In
other words, when the th ree fibers are
mo re closely packe d, the formation of lamellar struc ture at the fiber surface is more favored t han when they
ar e loosely packed; however, within the pocke t,
it
is less favored. Sim ilar results were found for th e four-fiber
case.
We used th e model
of
Fig. 18b
to
make the calculation for the free-boundary condition occurring along
a
straight, resin-rich area; the stress in the matrix along the free boundary is primarily unidirectional. Fig.
ure 19b illustrates that the relative stress magnitude and orientation correlate with the location
of
form atio n of
lamellar structure depicted in Fig. 16c.
In
conclusion, it is seen that th e magn itude and orientatio n
of
the matrix shrinkage stresses during pyrol-
ysis,
as
estim ated by this analysis, ar e consistent with the proposed m odel for stres s orientation and
graphitization.
Much still remains t o be learned about matrix stress graphitization in
C/C:
e.g., the effects
of
fiber type,
fiber volume, matrix precursor. an d high-temperature creep deformation. Equally intriguing is
the possibility
of being able to control C/C properties by controlling the matrix orientation and graphitization behavior.
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SUMMARY
Carbo n-carbon com posites are an exceptional class of high-strength, low-weight refractory materials;
however, effective utilization of th e ca rbo n fiber properties requires a ppropriate selection of th e carb on
or
graphite matrix and processing conditions. Th e matrices may be derived from hydrocarbon gases, coal-tar
and petroleum pitches, and thermosetting resins, and represent a range of structures and properties. Curren t
research is beginning to elucidate how C /C com posite properties may be controlled by controlling the struc-
tures of th e matrix, both in bu lk matrix regions and, more sensitively, at th e crucial fiber-matrix inte rpha se
region.
ACKNOWLEDGMENTS
This pape r reviews selected asp ects of work funded by the Air Force Space Systems Division under C on-
tract No.
FO4701-88-C-0089,
and by the Aerospace Sponsored Research Program. The author wishes to hank
a number
of
coworkers: Rafael Zaldivar. particularly, who, as an A erospace MS and P hD Fellow, is responsi-
ble for much of the work discussed here; Dr. Dick Chang for his effective collaboration
on
portions of this
work; and Paul Adams , Jim Noblet. Ross Kobayashi. Joe Uht. Dick Brose, and C a
Su,
all
of
whom contributed
in significant ways.
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1102
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Figure
1. Schematic of (a) 3-D block construction and (b)
2-D
plain-weave
fabric (McAllister an d L achm an [lo]).
Figure
2. Polarized-light micrographs showing as-deposited CVI carbon microstructures
of
two specimens. Depo sition seque nce: (a) RUSL; (b) S U R L (Rellick [B] .
5v
Figure 3. Scanning electron micrograph of speci-
men after heat treatment at 25OO'C for 1 hr
[a].
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8/11/2019 CAHBON ANI) GRAPHITE MATRICES IN CARBON-CARBON COMPOSITES
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0.6 0.7
0.8
WEIGHT FRACTIONOF
FIBER
PLUS
FIRST
CVI
Figure
4.
Through-thickness thermal cond uctivity Figure
5.
Composite tensile modulus versus
(at KT) for composite specimens of different CVI
weight fraction o f fiber plus heat-treated CVI
128).
structures a nd processing stages [28].
FIBER
MATRIX
FIBER
2 pm
Figure 6. SEM showing highly aligned coal-tar-
pitch-derived graphite matrix in the interfilament
region of a
CIC
comp osite. Fibers are
AmwoT50
from PAN.
1104
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0.2
mm
Figure
7. Optical micrograph ofcross section of 3-D CIC
compo site densified
with
both pitch and resin and heat-
treated t o 2750C. Note extensive matrix microcrack ing.
DJETHYNYLBENZENE
MONOMER
POLYARYLA CETYLENE (PAA)
PREPOLYMER
A
CARBONIZATION
GRAPHITIZATION COMPOSITES
ARBON-CARBON
looo'c)
(heat, pressure)
PAA POLYMER
COMPOSITES
Figure
8.
Chemical structure and processing
of
PAA -based composites
(Barry el
01 [48]
and Katzman
[49]).
1105
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400
W
LL
z
1000 2000 3000
HEAT.TREATMENT TEMPERATURE
( C)
Y
m
0 1000 2000
3000
50
HEAT-TREATMENT TEMPERATURE
( C)
a) b)
Figure
9.
Plots of tensile (a) strength and (b) fiber modulus of undoped and B-doped
P M 5 0 C/C composites (Zaldivar et
nl. [52] ) .
00
pm
4
b)
Figure 10. Micrographs of fracture surfaces of (a)undoped an d (b) B-doped
PAA-derived
CIC
composites heat-treated to
11OOC
[52 ] .
300
E130
1000 2000
3
HTT
( C)
Figure 11. Moduli of composites
HTI (Zaldivar et a/ [7]).
1106
)o
versus
8/11/2019 CAHBON ANI) GRAPHITE MATRICES IN CARBON-CARBON COMPOSITES
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a m
Figure
13
Polarized-hght micrographs
of
unidirectlonal C/C com-
posite heat-treated to
(a)
2900'C and (b) 1100C (Rellick et
a/
[6])
0
l m
Figure
12.
Fracture sur-
face
of
E105 composite to
2750'C
H7T.
showing ma-
trix sheath tube (Zaldivar
et n l
[7]).
U
1 I rm
Figure
14.
SEM micrograph
of
PX-7 filament em -
bedded in PAA-derived carbon matrix heat-
treated t o 2750C.
2 r m a) 0.4 pm Ib)
Figure
15. T E M
bright-field images
of
C/C resin-matrix-derived unidirectional composite:
a ) ransverse and (b) longitudinal sections (Rellick and Ad am s [56]).
1107
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.I
I
Figure
16.
SEMs of ion-etched unidirectional C/Cs heat-treated to 2900C
(Rellick et
a/ . [ 6 ] ) .
U
4w
Figure 17. SEM fracture surface
of
T50/SC1008 heat-treated to2900'C [ 6 ] .
1108
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k
h
Figure
18. Schematic
of
the local packing arrangement of (a) three and (b) four fibers.
Shaded ar ea denotes region for which stresses are calculated [6].
e
u l
z 1.0
w
2 0.5
6
Y
0
0 0.5 1.0 1.5 2.0 2.5
r.AXIS
RELATIVE DISTANCE FROM CENTER OF FIBER
a)
r-MIS
RELATIVE DISTANCE FROM CENTER OF FIBER
b)
5
Figure 19. Com puter p lot of the directions and relative magnitudes of the matrix stresse s in
the plane
of
the fiber of various points, relative to Figs. 18a and b
[6].
1109