Strength Improvement of Short Graphite
Fiber-Re info rced PoIy propyIe ne
R . A. WEISS
Ctwporate Research-Science Laboratories
Erron Research and Engineering Company
Linden, New Jersey 07036
The improvement of the tensile strength of discontinuous
graphite fiber-reinforced polypropylene is described. Two
methods were employed: (1) modification of' the polymer
backbone by the introduction of' acrylic acid grafts and (2)
varying ths chemistry of the fiber sizing. The strength improvemen1 s attained are attributed to intermolecular interactions, such as hydrogen bonding, which improve the adhesion between the fiber and the polymer.
INTRODUCTION
n an earlier paper (I), the properties cipolypropylene
reinforced with discontinuous graphite fibers were
described. Little attention was given i I that investigation to the optimization of the strength of the resulting
composite, and in fact, the tensile strengths achieved
were relatively low. The failure mechanism of these
materials involved both local matrix yic:lding and shear
fracture of the fiber-polymer interface, the latter being
most likely a consequence of poor adhesion between
polypropylene and the graphite fibers
Whereas considerable work has been done on the
optimization of the fiber-polymer interface in order to
promote adhesion in glass fiber-polypropylene composites, most notably the introduction of coupling agents,
the author is not aware of any similar work for graphite
fiber-polypropylene composites. The objective of the
present investigation was to ascertz in whether the
strength of the interface can be imprcved by chemical
modification of the polymer and/or the fiber sizing in
order to promote stronger component interactions, such
as hydrogen bonding.
I
have circular cross-sections (3).The cross-sectional areas
of both types of fibers were of the order of 50-60 p m Z
which corresponds to a diameter of 8-9 p m for a circular
fiber.
The fibers used in these experiments had a very thin
polymer sizing in order to protect the fibers during
processing and to promote wetting of the fiber by the
polymer. The Fortafil fibers (CG-3) had a sizing containing epoxy resin, while the Magnamite fibers were sized
with either polycarbonate (AS-1805) or a polyamide
(AS-1800).
Fiber-reinforced copolymer or polypropylene compositions ranging from 10 to 40 weight percent fibers
were prepared by letdown of a composite masterbatch
containing 30 to 40 weight percent fibers. The masterbatches were prepared by d r y blending of 114 in.
chopped fibers with the appropriate polymer, followed
by melt extrusion. The acrylic acid concentration of the
polymer was varied by blending a polypropylene
Table 1. Polypropylene Starting Materials
Polypropylene
EXPERIMENTAL
Materials
The modified polypropylene was a a-opylene-acrylic
acid graft copolymer containing 5.3 weight percent
acrylic acid (2).The characteristics o f t his copolymer are
given in Table 1 , as are the properties of the propylene
homopolymers also used in this investigation. All the
polymers were commercial products of the Exxon
Chemical Company.
Three different graphite fibers we -e evaluated, one
obtained from the Great Lakes Cai%on Corporation
(Fortafil) and two from Hercules, Inc. (Magnamite).
These are described in Table 2 . The Fortafil fibers are
prepared from Orlon polyacrylonit rile (PAN) fibers
which yield a dogbone cross-section while the Magnamite fibers are based on a Courtelle PAN precursor and
POLYMER COMPOSITES, IULY, 1981, Vol. 2, Yo. 3
~~
Percent acrylic acid
Melt flow rate'
(dg/min)
Zero shear viscosity ci;
204°C (pascal . s)
Flexural modulus* (GPa)
Tensile modulus' (GPa)
Tensile strength3 (GPa)
Impact ~ t r e n g t h room
,~
temperature
notched ( J i m )
unnotched (J/crn)
Heat deflection temperature @ 455 kPa ("C)
~~
18.8
1100
1.6
1.7
31.3
0.34
8.2
101
Propyleneacrylic acid
copolymer
5.3
38.0
390
1.6
1.8
32.0
0.22
4.8
95
ASTY Standard 0-1238, condition L.
ASTM Standard D-790, narrow section cu1 from injection molded sample Type I
tensile bar.
a ASTM Standard D-638, injection molded Type I tensile bar.
ASTM Standard D-256.
I
89
R . A. Weiss
Table 2. Graphite Fibers
~~~~~~~
~
Fortafil CG-3
Manufacturer
Precursor
Cross section
Sizing
Starting length (in.)
Fiber diameter (pm)'
Cross sectional area (pm')
Specific gravity'
Tensile modulus (GPa)'
Tensile strength (GPa)'
Great Lakes Carbon
Orlon PAN
dogbone
epoxy
.25
(a)'
13.5
(b)'
5.3
55
1.8
207
2.5
~
~
~~
~~
Fiber designation
Magnamite
AS-1800
Hercules
Courtelle PAN
round
polyamide
.25
8
Magnamite
AS-1805
Hercules
Courtelle PAN
round
polycarbonate
.25
8
50
1.8
234
2.8
50
1.8
234
1.8
' Nominal valuea reported by manufacturer.
For Fortafll fibers (a) 11 the major diameter of the dogbone and (b) Is the mlnor dlameter.
homopolymer with the copolymer and was determined
using infrared spectroscopy.
The mechanical characterizations were carried out on
injection molded test specimens, and these techniques
have been described elsewhere (1).
RESULTS AND DISCUSSION
Theoretical Considerations
In order to quantify the strength of the interaction
between the fiber and the polymer in a short fiberreinforced plastic, it is necessary to measure the adhesion between the components. Because of the size of
fibers, generally of the order of 10 p m in diameter, it is
easy to appreciate the experimental difficulties involved
with such measurements. Although several techniques
have been proposed for measuring adhesion or the shear
strength of the bond, (4)experimental data are scarce
and the reliability of such data is questionable. An alternative method, for estimating the degree of interaction
between the composite components, albeit strictly qualitative, is to compare the mechanical strengths of wellcharacterized systems using some appropriate model or
criterion relating strength to the interfacial adhesion.
This technique, though experimentally simple, has
some very important limitations or inherent difficulties
which may complicate the interpretation of the experimental data.
One of the simplest and probably the most widely
quoted models for predicting the strength of the fiberpolymer interface is the shear lag analysis formulated by
Kelly and Tyson (5).An important concept introduced
by this analysis is that of an ineffective fiber length, I,, in
which stress is transferred from the matrix to the fiber
and in which the fiber normal stress increases from zero
at the fiber end to a maximum value at a distance of 1,/2
from the fiber end. For an ideal plastic material this
critical length, i.e., the length over which the fiber
stress is less than the maximum stress, is related to the
fiber-polymer interfacial shear stress by E q 1 ,
age fiber stress, (uJ, can be calculated by integrating
over the entire fiber length, E q 2 .
cr =
(ai)maJl -
W?
(2)
Therefore, the composite stress, which is a function of
the average fiber stress and the matrix stress, depends
upon (cf)mas
and ZJT). If composite failure is caused by
fibers breaking, (Ur)ma.ris simply the fiber strength, and
the shear lag model predicts that the composite strength
increases with increasing interfacial strength or adhesion.
The shear lag analysis does not consider stress concentrations in the matrix, which can be significant especially
near the fiber ends (6), and as a consequence it overpredicts the composite strength. Similarly, failure of a
composite by some mechanism other than fiber fracture
can lower the observed strength. Halpin and Kardos (7)
have developed an empirical relationship for predicting
the tensile strength of uniaxially oriented short fiberreinforced plastics by introducing a strength reduction
factor which accounts for the difference in strengths
observed experimentally and predicted by the shear lag
analysis. Their strength reduction factor is exponentially
related to the ineffective fiber length predicted by the
shear lag analysis, and like the shear lag analysis, their
relationship predicts an increase in the tensile strength
of a short fiber-reinforced plastic with an increase in the
interfacial shear strength. In other words, although
quantitatively incorrect, the shear lag analyses at least
qualitatively predicts the correct relationship between
the fiber-polymer adhesion and the composite strength.
The effect of differences in the fiber-matrix adhesion
on the composite modulus deserves special attention.
Implicity, the ineffective fiber length is a function of
strain, and at very small strains, 1, is small compared to I
and can be ignored. An important consequence of this
result is that the modulus is relatively insensitive to
differences in the fiber-polymer adhesion as long as no
slippage occurs at the fiber-polymer interface.
Varying the Polymer Matrix Composition
where (u&,,~
is the maximum fiber stress, d is the fiber
diameter. and 7 is the interfacial shear stress. The aver90
The engineering tensile and flexural properties of
injection molded composites prepared with both polypropylene and copolymer matrices are plotted vs
POLYMER COMPOSITES, JULY, 1981, Yo/. 2, No. 3
Strength Impro jement of Short Graphite Fiber-Reinforced Polypropylene
graphite fiber volume concentration in Figs 1 and 2 .
Whereas there is relatively little differrmce between the
moduli of the polypropylene and copolymer-based
composites, a significant increase in strength is attained
by changing from a polypropylene to a copolymer matrix. As was discussed in an earlier papix, the differences
in the flexural and tensile properties ate not unusual and
can be adequately explained by statistical strength
theories (8).
Differences in the strengths of the :wo types of composites can be a result of several factcrs, such as differences in the fiber lengths, the fiber orientations, and/or
E
iGPa)
.clI_
Fle ural Modulus
,
0
0.10
0.05
0.25
0.20
0.15
0.30
6
Fig. 1 . Tensile and flexural moduli of Fcrtufil graphite fiberpolypropylene (empty symbols) und prop!llene-acrylicacid copolymer (filled symbols) composites us oolunie fracture fibers.
the adhesion of the fiber-polymer interface. Because the
rheological properties of the two matrix polymers are
markedly different, as shown in Fig. 3 , the processing
shear histories of the two types of composites are expected to be different, afact which suggests the possibility of different fiber lengths and orientations in the two
composites. These dfierences would result, however,
in differences in the moduli as well as the strengths.
The fiber lengths in these materials were confirmed
to be similar by isolating the fibers from injection
molded test specimens and measuring the lengths of
300-700 fibers. The results given inTable 3 indicate that
the effect of processing on the fiber length attrition in
the two systems was not significantly different. Because
no significant differences were observed in the moduli
and the fiber lengths in these materials, it follows that a
reasonable assumption is that the fiber orientations were
also similar. This conclusion is supported by the results
of an investigation in our laboratory of the variation of
the mechanical properties of glass fiber-reinforced polypropylene with processing conditions, which indicated
that the fiber lengths and fiber orientations of specimens
molded with the Boy injection molding machine used in
the current investigation were independent of the melt
rheology of the material, but were determined solely by
the geometry of the nozzle and the gate of the Boy
machine.
If the fiber lengths and fiber orientations of the composites described in Fig. 4 are equivalent, then the
differences in strength must be due to differences in the
adhesion between the fiber and the polymer, that is, T .
Actually, because the fibers used in this investigation
were sized, it is probably the adhesion between the
sizing and the matrix which is changing. The composites described in Figs. 1 and 2 were prepared with
Fortafil CG-3 fibers which have an epoxy sizing, and the
improvement in the interfacial adhesion resulting from
the use of the propylene-acrylic acid copolymer in place
of polypropylene is most likely due to hydrogen bonding
between the ether linkages of the epoxy resin and the
carboxylic acid functionality of the copolymer.
The effects of acrylic acid concentration on the tensile
properties of graphite fiber-polypropylene composites
are shown in Figs. 4 and 5 ; each point represents the
Tensile Strength
1
10 0.1
0.05
0.10
0.15
0.20
0.25
0.M
d
Fig. 2 . Tensile andflexural strengths of Fortajil graphitefiberpolypropylene (empty symbols) and propylene-acrylic acid copolymer (filled symbols) composites os v d u m e fracturefibers.
POLYMER COMPOSITES, JULY, 1981, Vol. 2 No. 3
10
100
lo00
10.000
Shear Rate (rec-l)
Fig. 3. Apparent viscosity us shear rate f o r polypropylene and
propylene-acrylic acid cop01ymer at 204°C (capillary rheometer, Ild = 40).
91
R. A. Weiss
80
Table 3. Average Fiber Length in Injection Molded Tensile
Bars of Graphite Fiber-Reinforced Polypropylenes
Matrix
Fiber volume
fraction
Number average
fiber length, mm
0.047
0.220
0.047
0.218
0.16
0.18
0.13
0.13
Polypropylene
Copolymer
70
60
Qc
(MPal
average of five determinations. In general, the composite modulus was independent of the acrylic acid concentration, whereas, the strength appeared to be extremely
sensitive to the presence or absence of acrylic acid. All
the composites failed by interface failure at high loadings
and by matrix fracture at lower loadings. Micrographs of
the tensile fracture surfaces were similar to those given
in an earlier paper (I).
The data in Figs. 4 and 5 are consistent with the
conclusion of improved adhesion with the acrylic acid
graft and demonstrate that a significant increase in
strength can be realized with as little as one weight
percent graft. In fact, the greatest strength improvement occurs with the addition of 1-2 percent acrylic acid,
though further increases in the acrylic acid concentration result in modest improvements in tensile strength.
The acrylic acid concentrations of the materials described in Figs. 4 and 5 were prepared by blending the
propylene-acrylic acid copolymer and a propylene
homopolymer, and one might question whether any
differences would result from varying the acrylic acid
concentration by changing the actual copolymer composition. This raises the question of potential differences
18
l4
I
I
I
t
0-O
0
-
0
-
0
-
6
0-0-
I
0
1
I
2
I
3
1
I
I
0555
1
0
0.099
0.047
0-
I
1
I
4
5
6
Acryltc Acids Concentration ( W t . X)
Fig. 4 . Tensile modulus os acrylic acid concentration for Fort a i l graphite fiber-polypropylene composites.
92
50
40
@=0.047
-
30
20
0
1
2
3
5
4
ACRYLIC ACID CONCENTRATION (WT.
6
7.)
F i g . 5 Tensile strength os ucrylic acid concentrutionfor Fortufil
gruphite ~ b e r - p o l y p r o p y l e n ec o ~ n ~ i o s i t e s .
between the morphologies of a polymer-polymer blend
and a graft copolymer. This issue, albeit important, was
not considered in this investigation. O n e might
presume, however, that because of the similarities of the
two polymer backbones, they might co-crystallize. For
example, Pegoraro and his coworkers (9) have demonstrated that graft copolymers of propylene and acrylic
acid exhibit microphase separation; that is, the acrylic
acid exists as discrete domains dispersed in a semicrystalline polypropylene continuum. Therefore, it may
be reasonable to expect the morphologies of the continuous phase of a propylene-acrylic acid graft copolymer and of a blend of a propylene-acrylic acid graft
copolymer and polypropylene to be indistinguishable.
T h e heat deflection temperature ( H D T ) of t h e
graphite fiber composites is plotted in Fig. 6 as a function of fiber concentration and acrylic acid concentration. Increasing the H D T of polypropylene by the addition of graphite fibers was discussed in Ref. 1and can, for
the most part, be explained by a shift of the modulustemperature relationship towards higher temperature
with increasing fiber concentrations. This does not,
however, account for the additional increase in the H D T
resulting from the addition of grafted acrylic acid to the
matrix polymer. Whereas, for a single component system, for example polypropylene, the H D T reflects the
modulus-temperature relationship of the material, the
analysis becomes somewhat more complex for a multicomponent system where the H D T of the composite
occurs above that of at least one of the components. For
example, for the systems described here the HDT's are
above those of polypropylene and the copolymer, which
are approximately 100°C. In this situation, creep or viscous flow of the polymer component above 100°C may
become significant and have a noticeable influence on
the composite HDT. This flow influence, however,
should be afunction of the adhesion or shear strength of
the polymer-fiber interface. Therefore, an improve-
POLYMER COMPOSITES, JULY, 1981, Vol. 2, No. 3
Strength Improvement of Short Graphite Fiber-Reinforced Polypropylene
190
r-----l
/
c
0-
I5t
I
JU
Matrix
PP
PP+2%AA
AS-1800
PP
A S - 1 8 0 0 PP + 2% AA
AS -1 805
PP
A S - 1 8 0 5 PP + 2% AA
CG-3
PP
CG-3
PP+2%AA
-
0
90t
A
8
0
0
A
A
I
70
0.05
0.15
0.10
0.20
0.25
1
I
I
1
0.05
0.10
0.15
0.20
+
0
Fig. 6 . f l e at d e j e c t ion temperature us f i ber concen t ra tiori and
acrylic uc id co tice ntrut io t i fo r graphite fiber-polypropylene
coinposites.
J
0.25
Fig. 7 . lensile modulus us fiber concentration f o r 1;urioti.s
gruphite jibers and polypropylene matrices.
ment of the HDT of the composite resulting from the
improvement of the fiber-polymer adhesion by the addition of acrylic acid is consistent with thl? tensile strength
results discussed above.
I
/
I
Varying the Fiber Sizing
The choice of the sizing is in many instances one of
convenience, but it can have an impor:ant influence on
the degree of the fiber-matrix adhesion. In this investigation several sizings were evaluated in order to judge
their ability to promote improved fiber-matrix adhesion
or wetting. As detailed earlier in this report, differences
in adhesion were inferred from differences in the composite mechanical properties, assumin!; equivalent fiber
loadings, fiber lengths, and fiber alignments.
Given the above restrictions, two possible explanations can b e invoked to account far differences in
mechanical properties observed for the different composites described here: (1)differences in stress concentrations between the fibers or adhesion between the
components d u e to differences in the shapes of the fiber
cross-sections or (2)d a e r e n c e s in adhc sion between the
fiber and the polymer matrix resulting from differences
in the chemistry of the sizings. Hancock (10) has considered the first possibility for both r o m d and dogbone
graphite fibers in an epoxy matrix and found no differences in the mechanical properties a:tributable to the
fiber shape. It is recognized that be-ause of possible
differences in the crystalline morpholc,gy resulting from
variation of the fiber shape, the translation of this result
to composites with a semi-crystalline matrix may not b e
strictly valid. Nevertheless, it is assumed here that the
mechanical properties of the graphite fiber-reinforced
polypropylenes are not sensitive to the shape of the
fibers.
The effects of varying the fiber sizing on the tensile
properties of these composites are shown in Figs. 7 and
8 ; the matrix containing 2 weight percent acrylic acid
POLYMER CoMPosirEs,JULY, 1981,YO/. 2,140. 3
ocu
(MPa)
PP
0.05
0.15
0.10
0.20
0.25
9
Fig. 8. Tensile sfreiigth 1;s fiber concentrution f o r curious
graphite fibers and polyprvpylene 1natrice.s.
was a blend of the propylene-acrylic acid graft copolymer and polypropylene. The assumption of equivalent
fiber lengths and fiber alignments is supported by the
tensile moduli data, Fig. 7 , as discussed previously.
Interfacial adhesion, however, was particularly sensitive
to the choice of the fiber sizing and the polymer matrix as
is demonstrated by the tensile strength data in Fig. 8.
In regard to these results, several points are noteworthy. First, the tensile strength data indicate that a large
increase in the strength of a composite based on a polypropylene matrix can be attained by using a polycarbonate sizing (AS-1805) as opposed to an epoxy sizing
(CG-3) or a nylon sizing (AS-1800). In fact, the tensile
strengths of such composites are greater than that
achieved with an epoxy sizing and the copolymer matrix,
the materials described earlier in this paper. The implication here is that the polycarbonate-polypropylenesys93
R . A. Weiss
tem results in significantly better adhesion than that
achieved in these other systems.
An explanation of these results may be found by considering an investigation of Ford and Goettler (11)in
which they correlate the reinforcement efficiency for a
variety of fiberglass-reinforced thermoplastics with the
dielectric constant of the matrix. As the polarity of the
matrix is increased, the reidorcement efficiency, which
is related to interfacial adhesion, increases as a result of
stronger dipole interactions between the polymer and
the oxide surface of the glass.
In the compositions described in Fig. 8, the matrix
was held constant and fiber sizing was varied. Following
the reasoning of Ford and Goettler, one expects that the
adhesion should increase as the dielectric constant of the
sizing approaches that of polypropylene. This is demonstrated by the data in Table 4 . Of the sizings used,
polycarbonate comes closest to matching the dielectric
constant of polypropylene and, as a consequence, improves t h e tensile strength of the polypropylene
matrix-composites to the greatest extent.
In a similar manner, the sizing-copolymer results appear to be consistent with solubility parameter data, c.f.,
Table 4 . Solubility parameters were used here because
of the unavailability of the dielectric constant for acrylic
acid. In this case, the solubility parameters of the
acrylic acid and the nylon are the best match, and
accordingly, this system results in the highest tensile
strengths. T h e superiority of t h e polycarbonatecopolymer system over the epoxy-copolymer system is
probably d u e to the interaction of the polycarbonate
sizing with the polypropylene phase. The dielectric constant and solubility parameter of the epoxy resin are not
similar to either that of polypropylene or acrylic acid,
and, consequently, an epoxy sizing results in the poorest
mechanical properties.
CONCLUSIONS
It has been shown that the strength and the temperat u r e resistance of discontinuous g r a p h i t e fiberpolypropylene composites can be varied considerably
by varying the chemistry of the fiber sizing and the
polymer matrix. These differences are presumably the
result of differences in the fiber-polymer adhesion.
Where the polymer matrix contains acrylic acid, the
improvements in interfacial adhesion, i.e., strength, are
most likely d u e to hydrogen bonding between the carhoxylic acid groups and the sizing polymer. This effect is
94
Table 4. Dielectric Constants1and Solubility Parameters* of
Selected Materials
Material
Acrylic acid
Polypropylene
Polycarbonate
Epoxy resin
Nylon
E
2.2
-3
3.5
-4
6(cal/~rn~)~~~
12.0
11.9-14.5
9.2-9.4
10.9
9.5-10.6
‘ “Polymer Handbook,” J. Brandrup and E. Immargut, eds., Interscience Publ.
(1975).
“Polymer Handbook,” J. Brandrup and E. Immergut, eds., Interscience Publ.
(1966).
’
most dramatic in t h e case where t h e sizing is a
polyamide. For an unmodified polypropylene matrix,
improvements in strength can be attained by matching
the dielectric constants or solubility parameters of the
sizing and the polypropylene. This is best accomplished
with a polycarbonate sizing.
The morphology of the matrix polymer in the vicinity
of the graphite fibers has not been considered, but may
have an important influence on interfacial adhesion and
the properties reported in this paper. For example,
carboxylic acid salts are known nucleating agents for
polypropylene (12) and, in fact, the crystallization kinetics of the copolymer used in this study indicate that the
acrylic acid grafts nucleate the polypropylene phase.
Similarly, graphite fibers are also excellent nucleating
agents for polypropylene (12). As a consequence, the
matrix morphologies of the different composites described here are potentially different and complex and
deserve further investigation.
REFERENCES
1. R. A. Weiss, Polym. Compos., 2, 95 (1981).
2. R. A. SteinkanipandT. J. Grail, U.S. Patent 3,862,265(1975).
3. R. J . Diefendorfand E. Tokarsky, Polym. E n g . Sci., 15, 150
(1975).
4. L. J. Broutrnan in “Interfaces in Coniposites,” p. 27, ASThl
STP 452, ASThI (1969).
5. A. Kelly and W. R. Tyson, J. Mech. Phys. Solids, 13, 329
(1965).
6. R. M . Barker and T. F. ?vlacLaughlin,J. Coinpos. Muter., 5,
492 (1971).
7. J. C. Ilalpin and J. L. Kartlos, Polym. E n g . Sci.,l8, 496
i1978).
8. 8. E. Bullock, J. Compos. Muter., 8, 200 (1974).
9. M. Pegoraro, Pure A p p l . Chem., 30, 199 (1972).
10. N. L. Hancock, Fibre Sci. Tech., 10, 179 (1977).
11. E. A. Ford and L. A. Goettler, Polyni. Prepr., Am. C h i n .
Soc., Diu. P o l y m . Chem., 15, 451 (1973).
12. S. Y. Hobbs, Nature (London), Pkys. Sci., 234, 12 (1971).
POLYMER COMPOSITES, JULY, 1981, Vol. 2, No. 3