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U-M-I

THE BONDING MECHANISM OF
ARAMID FIBERS TO EPOXY MATRICES

By

Javad Kalantar

A THESIS

Submitted to
Michigan State University
in partial fulfillment of the requirements
for the degree of

MASTER OF SCIENCE

Department of Chemical Engineering

Winter 1988

ABSTRACT

THE BONDING MECHANISM OF
ARAMID FIBERS TO EPOXY MATRICES

By

J AVAD KALANTAR

The interface between aramid fibers and epoxy matrices lacks the level of
adhesion attained by other reinforcing fibers. For aramid composites, the interfacial —~~-~
properties attained to date are acceptable, but less than optimal for some -—
applications. By developing a basic understanding of the interfacial interactions
between aramid fibers and epoxy matrices, fundamental approaches for improving the
adhesion can be identified.

W ”Three types of interfacial interactions have been examined: (mechanical)
’W Wthat include thermal strain and Poisson’s ratio differences between fiber
and matrix, (chemical interactions that include covalent bonding and fiber-matrix
wetting, and physrcochemical weak boundary layers. Each of these interactions has
been evaluated by manipulating the interface and the curing conditions. Both aramid
and carbon fibers have been examined in order to access the interfacial interaction by

comparing the behavior of these two fibers.

Our results indicate that the adhesion of aramid fibers to epoxy matrices lacks
the mechanical and chemical interactions present in carbon-epoxy adhesion. Aramid
fibers exhibit an interfacial shear strength as much as four times lower than the
expected theoretical value. Direct observation of the aramid-epoxy interface, by
transmission electron microscopy, shows fibrillar separations within the fiber surface.
This type of interfacial failure suggests that Mamifippxyhadhesion could be due
ejectflrfaqsiailw W W W W W

TO MOHAMMAD J AHAN

iii

Acknowledgement

This work was accomplished at the Composite Materials and Structures
Center (CMSC), at Michigan State University. I want to thank Dr. L. T. Drzal for

the opportunity to be exposed to the project and his support, assistance, and patience
during my work.

I wish to gratefully acknowledge E. I. du Pont de Nemours & Co. for its
financial and research support.

Finally, I would like to express my appreciation to the Staff of Center for
Electron Optics for their help and informative discussions.

iv

Table of Contents

List of Tables .................................................................................................... vii
List of Figures .................................................................................................. viii
Nomenclature .................................................................................................... ix
Introduction ....................................................................................................... 1
Background ........................................................................................................ 4
Composite Interface ........................................................................................... 4
Polyaramid Fibers .............................................................................................. 5
Mechanical Properties of Aramid Fibers .......................................................... 12
Carbon Fibers ..................................................................................................... 15
Epoxy Matrices .................................................................................................. 16
Aramid-Epoxy Adhesion ................................................................................... 21
Aramid-Epoxy Interface .................................................................................... 23
Microstructural Approach ................................................................................. 23

1. Mechanical Interlocking ............................................................... 24

2. Adsorption Interactions ................................................................ 24

3. Electrostatic Attractions ................................................................ 25

4. Polymer Interdiffusion .................................................................. 25
Macrostructural Approach ................................................................................ 26

1. Mechanical Interactions ................................................................ 26

Surface Topography ............................................................... 26

Thermal Stresses .................................................................... 28

Poisson’s Ratio Difference .................................................... 29

2. Wetting .......................................................................................... 30

3. Weak Boundary Layers ................................................................ 30

Theory .................................................................................................................... 33
Experimental ..................................................................................................... 37
Results and Discussion ................................................................................ 42
Thermal Stresses ................................................................................................ 42
Chemical Bonding ............................................................................................. 46
Poisson Contraction ........................................................................................... 48
Fiber Wetting ..................................................................................................... 50
Three dimensional stress model ........................................................................ 50
Electron Microscopy .......................................................................................... 60
Conclusions and Recommendations ................................................... 69
Appendix A: Three-Dimensional Stress Model ....................................... 72
Appendix B: Critical Length distributions ............................................... 76
Appendix C: Thermal Expansion Data ..................................................... 101
Appendix D: Microtoming Technique ........................................................ 108
References ........................................................................................................... 113

vi

List of Tables

Table 1 Material properties of Kevalr 49 and As4 fibers. ......................................... 13

vii

List of Figures

\i

Figure l The sample conditions and the interactions present at each condition. .............
Figure 2 Structure of polyaramid PP'I‘A monomer. ...................................................
Figure 3 Dobb and Johnson model for the aramid pleated structure. .............................
Figure 4 Morgan’s model for the aramid fibrillar morphology. ...................................
Figure 5 Model of carbon fiber ribbon morphology. ..................................................
Figure 6 Structure of DGEBA epoxy. .....................................................................
Figure 7 Structure of DETA, MPDA, and DETDA curing agents. ...............................
Figure 8 SEM micrographs of fracture surfaces of unidirectional AS -4/epoxy and

Kevlar 49/epoxy composites. ....................................................................
Figure 9 Tensile jig for the critical length technique and a sample mold. ......................
Figure 10 Plot of % thermal shrinkage of the epoxy matrices. .....................................
Figure 11 Plot of experimental interfacial shear strength of untreated carbon and

aramid fibers with epoxy matrices cured at different temperatures. .................
Figure 12 Plot of experimental interfacial shear strength of untreated and gold coated

carbon and aramid samples made with 75°C and 125°C cured epoxy. .............
Figure 13 Plot of interfacial shear strength of untreated, gold coated, and silicone

coated carbon and aramid samples made with the room cured epoxy. .............
Figure 14 Plot of interfacial shear strength of untreated, gold coated, and silicone

coated aramid with epoxy matrices cured at different temperatures. ................
Figure 15 Plot of epoxy elastic modulus cured at different temperatures. .......................
Figure 16 Plot of theoretical and experimental critical lengths for aramid and

carbon fibers with epoxy matrices cured at different temperatures. .................
Figure 17 Plot of theoretical and experimental average interfacial shear stress over

the critical lengths for aramid and carbon samples. .....................................
Figure 18 Optical micrographs of fiber fragment at their critical lengths with

bright-field and cross-polarized light. .........................................................
Figure 19 Plot of theoretical average radial stresses for aramid and carbon fibers

with epoxy matrices cured at different temperatures. ....................................
Figure 20 TEM micrographs of Kevlar 49 and (50/50) epoxy matrices cured at

175°C (lognitudinal cut). .........................................................................
Figure 21 TEM micrographs of Kevlar 49 and room-cured epoxy (radial cut). ................
Figure 22 SEM micrographs of a single Kevlar 49 fiber. ............................................
Figure 23 TEM micrographs of a surface damaged Kevlar 49 in a epoxy matrix

cured at 175°C (radial cut). .....................................................................
Figure 24 TEM micrographs of a surface damaged Kevlar 49 in an epoxy matrix

crued at 175°C (radial cut). The section is stained with 0504. ....................
Figure 25 TEM micrographs of a shear damaged Kevlar 49 in an epoxy matrix

cured at 75°C (radial cut). ........................................................................
Figure 26 TEM micrographs of a shear damaged Kevlar 49 in an epoxy matrix

cured at 75°C (radial cut). .......................................................................
Figure 27 Diagram of stress distributions around a fiber fragment. ...............................

viii

22
39
43

45

47

49

51
52

54

55

57

59

62
63

65

67

68
75

Nomenclature

Elastic constant for fiber longitudinal stress, MPa
Elastic constant for fiber radial stress, MPa
Elastic constant for fiber radial displacement
First elastic constant for matrix radial displacement
Second elastic constant for matrix radial displacement
Axial fiber elastic modulus, MPa

Radial fiber elastic modulus, MPa

Matrix elastic modulus, MPa

Radial fiber elastic shear modulus, MPa

Matrix elastic shear modulus, MPa

Plane strain bulk modulus, MPa

Critical length, um

Radial coordinate, urn

Fiber radius, tun

Curing temperature, °C

Glass transition temperature, °C

Longitudinal fiber displacement

Longitudinal matrix displacement

Radial fiber displacement

Radial matrix displacement

Longitudinal coordinate, um

Dimensionless longitudinal coordinate

Weibull shape parameter

Matrix coefficient of thermal expansion, ppm/°C
Axial fiber coefficient of thermal expansion, ppm/°C
Radial fiber coefficient of thermal expansion, ppm/°C
Weibull scale parameter, um

Temperature difference between ambient and oven condition, °C
Far field axial strain

Fiber fracture strain

Matrix thermal strain

Longitudinal fiber thermal strain

Radial fiber thermal strain

Matrix Poisson’s ratio

Axial fiber Poisson’s ratio

Radial fiber Poisson’s ratio, assume to be equal to vuf
Fiber tensile strength, MPa

Longitudinal fiber stress, MPa

Radial fiber stress, MPa

Longitudinal matrix stress, MPa

Radial matrix stress, MPa

Average interfacial radial stress, MPa

Interfacial shear strength, MPa

Fiber shear stress, MPa

Matrix shear stress, MPa

Average interfacial shear stress, MPa

Bulk constant

INTRODUCTION

Aramid fibers have a unique combination of stiffness, high strength, and low ~
density which rivals the properties of inorganic reinforcing fibers such as glass and ——
carbon. Advanced composites made of aramid fibers have excellent axial properties -
as compared to inorganic fibers, but their off-axis properties are less than optimum
for some applications. The off-axis properties of fiber-reinforced composites are

generally controlled by the level of fiber-matrix adhesion.

In this study, attempts are made to understand the adhesion interactions of
aramid fibers with epoxy resins. These adhesion interactions are : mechanical strains,
chemical interactions, and effects of physicochemical weak boundary layers.
Mechanical strains are caused by thermal shrinkage and Poisson’s ratio differences
between fiber and matrix. Chemical interactions include wetting and covalent
bonding. The wetting of the fiber with liquid epoxy insures intimate molecular
contact between them which is a prerequiste for covalent chemical bonding. Weak
boundary layers can reduce the efficency of load transfer at the interface and
significantly affect other interfacial interactions. Application of controlled weak

boundary layers is used to manipulate fiber-matrix interactions.

The approach of this study is to separate and qualitatively analyze the
adhesion interactions by modifying the curing conditions and the fiber surface.
Figure 1 illustrates the experimental plan.

 

 

 

 

 

 

Mechanical Chemical
Sample Interactions Interactions
Conditions
Thermal Poisson Covalent Wetting
Stress Contraction Bonding
Untreated Fibers
& + + + +
Elevated Temp. Cured
Untreated Fibers
& + + +
Room Temp. Cured
Gold Coated Fibers
& + + +

Elevated Temp. Cured
Silicone Coated Fibers

 

 

 

& + +
Elevated Temp. Cured
Gold Coated Fibers
& + +
Room Temp. Cured
Silicone Coated Fibers
& +

 

 

 

 

 

 

 

 

Room Temp. Cured

 

Figure 1 - The experimental sample conditions and the interactions present at

each condition.

As shown in Figure 1, various combinations of epoxy curing temperatures and
fiber surface modifications have allowed distinguishing different interaction

mechanisms.

Gold and silicone coating of the fibers are the two types of surface treatments
examined. Both treatments introduce physicochemical weak boundary layers which
effectively eliminated all the covalent bonds that might be present at the fiber-matrix
interface, but also modified the thermodynamics of the fiber surface. Gold coating
has produced an inert, yet wettable fiber surface. Silicone coating has produced an

inert, but non-wettable fiber surface.

Carbon fibers have physical and chemical properties that are distinctly
different from aramid fibers. Comparison of these two types of fibers in samples
made with the same epoxy has enabled the effects of different interactions to be
distinguished.

A three-dimensional linear elastic stress model is used to compare the
experimental data with their theoretical values. The model has been developed by
Whitney et al. [1]. For a single fiber fragment, stress distributions and critical

lengths are predicted by the model.

A direct observation of the aramid-epoxy interface is carried out by
microtoming the single-fiber samples and examining them by transmission electron

microscopy.

BACKGROUND

Composite Interphase

 

The mechanical behavior of composite materials reflects the interactions r— ”

between their various constituents. When a load is applied to a fiber reinforced W

/’

—.__.——

composite, the load is transferred between matrix and fiber through their interphase. ~-
A sgggg interphase promotes greater involvement of the fibers, thus contributing to “r
the. composite strength. The fiber-matrix interphase also determines the failure mode

of the composite. At high levels of adhesion, the failure would start with matrix
cracks, but at lower adhesion levels, failure occurs along the fiber-matrix interphase.

For continuous fibgiéinforced composites, the fiber-matrix interphase is symmetric ~—
along the fiber longitudinal axsis, and is commonly referred to as the fiber-matrix ~

"interface".

Composite behavior is significantly affected by the condition of its interface
[2]. In particular, off-axis properties such as interlaminar shear and transverse
strength can be improved by increasing interfacial bonding. Improved interfacial -—
adhesion also enhances the environmental stability of polymeric composites by —-
reducing the formation of weak boundary layers. However, for some applications --
such as fracture toughness, a low level of fiber-matrix adhesion is desirable. In .
general, the optimum condition for interfacial strength depends on the particular -—

application and its expected loads.

In elementary treatments of the composite tensile properties, the effects of the
interface are usually ignored [3]. In practice, the interfacial properties have moderate /'
to critical influences on many mechanical or thermal properties of the composite. /‘
Studies by Drzal et al. [4] and Owen [5] on carbon fibers with different surface
prOperties have demonstrated the significance of interfacial properties on fiber-
dominated composite properties. Peters et al. [6] have shown that the mechanical
properties of the composite are affected more by the interfacial condition of the
fiber-matrix than by the degree of the cure of the matrix. Such observations suggest
that the curing cycles of the resin should be optimized with respect to the desired

fiber-matrix adhesion rather than optimum matrix mechanical properties.

At low levels of fiber-matrix adhesion, fracture toughness and impact
resistance of the composites usually increase. A report by Chang et a1. [7] on
carbon-epoxy composites with controlled interfaces has shown an inverse relation
between interlaminar shear strength and impact resistance of the composite. Similar
works by Mai et al. [8], on fracture toughness of Kevlar-epoxy composites has /
demonstrated a 200% to 300% greater fracture toughness for Estapol-7008 coated /
fiber composites than uncoated fiber composites. The above studies has shown that a ’
very high level of fiber-matrix adhesion can be detrimental to fracture toughness and ’

impact resistance of the composites.

Polyararrrid Fibers

 

At molecular levels, the strength of organic polymers is related to the rupture /
of their carbon-carbon bonds. In theory, the material strength can be calculated from /
the carbon-carbon bond dissociation energy (~ 83 kcal/mole) and the packing of the
polymers [9]. However, for most solid materials, the measured strength of the bulk
is several orders of magnitude smaller than the theoretical values. The main reason

is the existence of flaws or defects in the structure of the material. Misalignment in

the orientation of the polymer chains, broken chain ends, and slippage of the chains
can lead to stress concentrations on a few bonds which cause chain rupture and
catastrophic failures. To reach high material strengths, certain highly ordered *“
polymer morphologies are required. Polymer chain packing, orientation, and /
extension significantly affect the material strength. The distribution of flaws and

cracks which are detrimental to the strength and must also be minimized.

During the past two decades, considerable progress has been made in the
production of high performance synthetic fibers [10,11]. These fibers have high
degrees of crystallinity and their ultimate properties approach their theoretical
maximums. The most successful high performance organic fibers have been prepared 1
from wholly aromatic polymers [12]. These fibers have high modulus, high strength,
and are not brittle. Preston [13] has reviewed the development of aromatic polymer

fibers.

To date, the most successful high-performance organic fibers have been ,
polyaramid fibers. EJ. du Pont is the major manufacturer of one type of aramid fiber ’
which is marketed under the trade name Kevlar®. Three types of Kevlar fibers are
available for specific applications : (1) High modulus Kevlar 49 for composite '
reinforcement, (2) intermediate modulus Kevlar 29 for ropes and fabrics, and (3) tire "
cord Kevlar. Since its introduction in 1971, Kevlar has become the major reinforcing

fiber for applications where toughness and impact resistance is required [14].

Kevlar fibers consist of extended chains of highly oriented rod-like molecules /"
[15,16] formed into fibers with a nominal diameter of 12 um. The aramid monomer
i\s_pa_ra;Phenylene Terephthalamide (PPTA) and its chemical structure is shown in
Figure 2. The polymer chains are oriented in the fiber longitudinal direction and are
hydrogen bonded to each other. The structure of Kevlar fibers is not well

documented, but some conclusions about their morphology can be made.

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Dobb et al. [17], using electron diffraction and electron microscope dark-field
image studies, have reported that the structure of Kevlar 49 fiber consists of sheets of
polymer chains radially arranged and held together by hydrogen bonding (Figure 3).
These sheets are regularly pleated along the axial direction of the fiber, with a pleat
angle of about 170°. Over small transitional sections between the pleated sheets, the
PPTA polymers are parallel to the axial plane; this feature eliminates the possibility
of rotational molecular orientation. Dobb has observed two main types of 500 nm
and 250 nm periodicities in the fibers, but near the edges of the fiber he has reported

evidence of marked changes in the spacing.

Morgan et al. [18] have studied the relation between Kevlar fibers failure
process and its structure. While not disputing the pleat morphology of the fiber, they
have suggested that the primary structural factor affecting the deformation and failure
of the fiber is the concentration and distribution of the supermolecular chain ends
within the fiber. Based on the fiber fabrication procedure, structure of the PPTA
crystals, microscopic deformation, and fracture topography studies, they have
proposed a model of chain-end distribution. In this model shown in Figure 4, chain-
end distributions are random in the fiber exterior, but progressively more aligned and
clustered in the interior. Morgan’s model suggests a skin-core morphology for the
fiber with random chain distribution at the skin and periodic weak planes at the core.
The periodicity of the weak planes is about 200 nm, which is the suggested average
length of a PPTA rrricromolecule rod. Morgan further suggests that PPTA
macromolecules are clustered into cylindrical crystals with 60 nm diameter and 200
run length. A large percentage of macromolecules transverse the weak plane, keeping
the continuity of the crystals in the axial fiber direction. It can be inferred from ,
Morgan’s model that crack propagation can readily occur parallel to the rods and ,
across the weak planes, leading to fiber fibrillations.

 

Rm 3 - Dobb and Johnson model for the aramid fiber structure. The
diagram shows a system of radially pleated polymer planes.
A small vertical section is located between each pleat.

10

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

]
100 to 1000 nm skin
Path of a break

Figure 4 - An exaggerated model of aramid fiber morphology proposed by
Morgan. The PPTA chains are ramdomly distributed in the fiber
exterior and progressively more clustered at the interior. Such
a sturcture results in skin-core difference in the fiber.

11

Morgan’s suggested skin-core morphology and chain-end distribution model
has been implied by other workers. Skin-core morphology of the fibers has been
suggested in studies by Chatzi et al. [19]. Using' the Dignam-Roth theory, their
photoacoustic FI‘IR spectroscopic studies have determined a difference in the chain
orientation between the exterior and the interior of the Kevlar 49 fibers. Brown et al.
[20], using electron paramagnetic resonance studies, have determined that the
concentration of the stress-induced free radicals is more than estimated concentration
for the fracture surface. They have suggested that the excess free radicals are
produced by polymer chain scission at weak planes within the fiber.

A fiber fabrication process for the aramid fibers has been reported by Morgan
[21]. Aramid fibers are produced by the condensation polymerization of
terephthaloyl chloride and p-phenylene diamine [22]. The PPTA is polymerized
using a stoichiometric ratio of the reactants. The polymer solution is washed with
NaOH to neutralize the HCl formed during polymerization. The solution is then
extruded into hot walled cylinders, whereupon the solvent is removed and the shear
forces cause the PPTA liquid crystals to orient in the direction of the shear. The
resulting yarns are washed and the subsequent stretching and drawing Wtr'eatrnents. -
amuse stifmsssarrd “$9.31!:

In another report by Morgan et al. [23], the chemical impurities in Kevlar 49
fibers have been investigated. They have determined that there are ~ 0.7% of
Na2804 impurities within the fiber, which are the result of sulfuric acid neutralization
step. Similar impurity concentrations have been reported by Penn et al. [24].
Morgan has suggested that NaZSO4 residues in the interfibrillar regions are paths for
moisture diffusion, which during fiber fabrication can generate microvoids in the
fiber. Ashbee et al. [25] have proposed a chemical volume expansion model to
describe the hydration expansion which can result in fiber fracture. Based on "salt-

weathering" mechanisms in geology, Whalley et al. [26] have confirmed the

12

possibility of a NaZSO4 hydration fracture mechanism. Small angle X-ray scattering
studies by Lee et al. [27] have suggested that failure of Kevlar 49 fibers is due to
increases in the volume fraction of microvoids and their enlargement along the fiber

axis direction.

Mechanical Properties of Aramid Fibers

 

Chiao et al. have documented the common properties of aramid fibers and
their composites [28]. Table 1 lists the mechanical properties of Kevlar 49 relevant
to this study.

Kompaniets et al. [29] have examined the statistical aspects of aramid fiber
tensile strength. They have reported that for both monofilament and yarns, the tensile
strength decreases with an increase in gage length, but the tensile strength of the
unidirectional composites was found to be unaffected by the gage length. Their
tensile strength data were also independent of the deformation rate (1 to 20
rum/min). Many investigators have attempted to model the strength and modulus
behavior of the Kevlar fibers. Knoff [30] has proposed a modified weakest-link
model to describe the tensile strength of the fibers as a function of test length. His
experimental results have suggested that for aramid fibers below 1 cm gage length,
the tensile strength only slightly increases with decreasing length, but above 1 cm the
tensile strength is strongly dependent on the gage length. Their model can be
considered to describe the aramid fibers as a series of approximately 1 cm uniform

strength links.

13

Table 1 - Material properties of Kevlar 49 and AS—4 fibers

 

 

 

Property Ref. Kevlar 49 AS—4
E1f (GPa) [83] 119 231
E2f (GPa) [1] 6.9 21
v12f [1] 0.35 0.25
G2f (GPa) [1] 2.6 8.3
(11f ppm/°C [791,[1] -5.72 -2
(12f ppm/°C [791,[1] 65 8.5
ouf (GPa) [83] 3.31 5.86
euf (%) [83] 2.5 1.4

 

 

l4

Compressive buckling of aramid fibers has been modeled by DeTeresa et al.
[31]. They have suggested that the compressive strength of the fibers should be /
proportional to either the shear modulus or the shear strength of the fibers. Another / -'
report by DeTeresa et al. [32], describes a new technique to determine the
compressive and torsional behavior of the Kevlar 49 fibers. A torsional pendulum
has been developed and ratios of (5:1) tensile-to-compressive strength, (17:1) tensile-
to-shear strength, and (70:1) tensile-to-shear moduli have been reported. White et al.
[33] have proposed an interesting mechanical model to describe compressive buckling
of the aramid fiber. Their model involves the classical mechanics spring, dashpot,
and rigid rods elements. This model can describe both the compressive buckling and

the stress-strain characteristic under tensile loading after the buckling.

Fatigue, creep, and failure behavior of aramid fibers have been studied by
several workers. Wagner et al. [34] have investigated the creep-rupture statistics of
the Kevlar 49 fibers. They have observed that for a given stress level, fibers with 7%
lower tensile strength show an order of magnitude lower rupture lifetime. Wagner et
al. have suggested a power law dependence for the creep behavior, with different
regions of power-law exponents. Lafitte et al. [35] have reported similar multi-region
power law creep behavior for Kevlar 29 fibers. In another fatigue study by Lafitte et
al. [36] the tensile strength variability of Kevlar 29 fibers has been attributed to the
distribution of defects in the fibers. Cook [37] has suggested a kinetic model for

Kevlar 49 creep-failure behavior that allows prediction of its rupture lifetime.

Several techniques for determining the failure mode of aramid fibers has been
reported. Fracto—emission studies by Dickinson et al. [38] have indicated a
correlation between total emission of electron and positive ions with the extend of
fibrillation in a Kevlar fiber. Acoustic emission (AB) techniques on Kevlar 49 fibers

by Hamstad et al. [39] have been less fruitful. The load at which the failure of a

15

single fiber occurred was not directly correlatable with the AE event peak amplitude.
An interesting model for longitudinal splitting from surface defects has been set forth
by Wagner [40] which allows the prediction of onset and growth of the cracks from

surface flaws.

Carbon Fibers

 

Carbon fibers are the most widely used high-performance reinforcing fibers. /
Carbon fibers have distinctly different properties from aramid fibers. This study has /
concentrated on the adhesion properties of the aramid fibers, but comparing aramid
and carbon fibers is helpful in highlighting the interfacial interactions. The
morphology and the surface properties of carbon fibers will be described briefly.
Additional details and more extended discussions on carbon fibers and composites are
provided in a review article by Riggs et al. [41] and the reference book by Donnet
and Bansal [42].

In a carbon fiber, the carbon atoms form hexagonal rings which are extended
by covalent bonds to form a plane of carbon rings called "basal planes". These
planes stack and form layers which are held together by weak van der Waal forces;
however, there is little alignment between equivalent carbon atoms in adjacent
planes. This crystallographic structure of carbon is referred to as "turbostratic
graphite". In the plane of the rings, the tensile moduli is very high (910 GPa), but
normal to the rings, the moduli is considerably reduced (30 Gpa). The graphite
crystals form ribbon-like structures along their basal planes. The general orientation
of the ribbons is along the fiber axis. There is a large number of microvoids existing
between the ribbons. As the number of microvoids decreases and the ribbons are
more aligned, the tensile modulus and strength of the fiber increases. Diefendorf and
Tokarsky [43] have shown that the orientation of the ribbons varies from the surface

to the center; with more alignment closest to the surface. This morphology shown in

16

Figure 5, results in a fiber with more load-carrying capacity on the surface than in its

COI'C.

Carbon fibers are usually surface treated in order to enhance their adhesion
properties. Comparison of composite properties of surface treated A54 and untreated
AU -4 carbon fibers by Drzal et al. [4] has demonstrated the improvement of interface
sensitive properties. Most surface treatments tend to modify the interphase by
creating fiber surface porosity and increasing the fiber-matrix contact areas. These
treatments also improve the reactivity of the surface by forrnin g reactive functional
groups, mostly carbonyl and carboxylate. Drzal et al. [44] have demonstrated that
another effect of surface treatment is the removal of weak structural layers from the
fiber surface which can be a source of interfacial failure. It is usually difficult to
assess the relative importance of each mechanism on the improvement of the carbon

fiber adhesion.

Mechanical properties of carbon fibers vary significantly among different
brands due to different precursors and production techniques available. Table 1
includes some of the mechanical properties of the AS-4 carbon fibers utilized in this
study. More detailed discussions of the mechanical properties of carbon fibers are
available elsewhere [41,42]. A discussion of statistical aspects of carbon fiber
strength distributions has been presented by Phani [45].

Epoxy Matrices

 

Epoxy resins are the most common resins used in high-performance
composites. Epoxies are therrnoset resins with a wide range of viscosities and can be
reacted with a variety of curing agents to obtain a spectrum of mechanical
properties. The chemical reaction between epoxy and the curing agent is referred to
as "curing". The reaction forms a three-dimensional crosslinked solid structure. The

curing mechanisms and kinetics of epoxy systems have been discussed extensively in

 

 

A carbon ribbon

Figure 5 - Arr exaggerated model of carbon fiber ribbon structure.
The amplitude and wavelength of the ribbons vary
from the core to the surface with increasing alignment
near the surface.

18

the literature [46,47]. The mechanical properties of various epoxy systems have been
documented by Lee and Neville [48]. ‘(hrring reactions of epoxies do not release any
volatiles or excess by-products, thus, they usually have low volume shrinkage.
Epoxy resins exhibit a high level of adhesion among polymers and cured epoxy

resins have excellent chemical resistance and provide good electrical insulation.

The epoxy system used in the present study is Diglycidyl Ether of Bisphenol
A (DGEBA) shown in Figure 6. This epoxy has been cured with three types of
primary amine curing agents: DiEthyleneTriAmine (DETA); M—PhenyleneDiAmine
(MPDA); and DiEthleolueneDiAmine (DETDA) shown in Figure 7. These three
curing agents all contain two similar primary amines at each end. The chemistry of
the epoxy-amine crosslinking is identical in all three systems, but due to steric

hindrances, each system has different kinetics.

Cure kinetics of the DEGBA epoxy with other aromatic diamines has been
investigated by Moroni et al. [49]. In their study, using thermal analysis data, they
have distinguished between different kinetic mechanisms. A report by Wiggins [50]
compares the curing behavior of DGEBA and DETDA curing agent with other

aromatic curing agents. \

Mechanical properties of DGEBA systems cured with amine curing agents
have been investigated by several workers. For the DGEBA/MPDA system, Gupta et
al. [51] have reported the dependence of the mechanical properties on temperature,
composition, and the relaxation process. Their results have suggested that, in the
glassy state, the tensile modulus is dependent on the intermolecular packing, but in
the rubbery state, the crosslinking density is the important factor. The effect of
crosslinking density on the glass transition of several amine-cured epoxies has also
been examined by Bellenger er al. [52]. Another report by Gupta et al. [53] have
examined the free volume and its effect on moisture transport in DGEBA/MPDA

system. They have determined that below the glass transition, there is no covalent

19

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o

.583 <mmOQ Co ouaosbm - o oSwE

 

\ /
£0
_
mmo

/ \

lame: no I own
/ \

 

Emmoe < schism .8 55m 32965

20

DiEthyleneTriAmine (DETA)
H

HN—CH —CH —N—CH —CH—NH
2 2 2 2 2 2

m-PhenyleneDiAmine (MPDA)

NH2

NH

DiEthleolueneDiAmine (DETDA)

CH

3
NH 2
H3CH2C CHZCH3
NH 2

Figure 7 - Structure of DETA, MPDA, and DETDA curing agents.

21

bond formation with water, but at higher temperatures there would be reactions that
can have adverse effects on the modulus of the matrix. Jean et al. [54] have
examined the free volume of four epoxy matrices with different crosslinking densities
using positron annihilation spectroscopy. Their study has shown that the transitions
in free volume coincide with the glass transitions and crosslink density of the
polymers. Lee [55] has examined a molecular model to described the plastic
deformation of various amine and anhydride cured DGEBA resins and has related his

model parameters to the known chemical structures of the resin.

Aramid-Epoxy Adhesion

._.__’

 

Separation of aramid fibers from an epoxy matrix is characterized by bare
fibers with little resin adhering to them. Figure 8 illustrates the fracture surface of
unidirectional Kevlar 49-epoxy and AS-4-epoxy composites. The failure modes of
the aramid-epoxy composites are generally interfacial [56,57], with some surface
fibrillation. Significant amounts of research have been devoted to the interfacial
properties of glass and carbon fibers and many surface treatment techniques and
coupling agents have been developed. For glass and carbon fibers these surface /
treatment techniques can increase the interface strength two or three times /
[41,58,59]. Development of similar improvement with interfacial treatment methods /

/"

for the aramid reinforced composites has been difficult.

Cooke [60], Morgan and Allred [21] have documented various attempts on the
development of surface treatment techniques for the aramid fibers. Despite many
efforts, promising coupling agents have not been developed [61]. Surface treatments
such as surface oxidation techniques can improve adhesion but are usually
accompanied by loses in fiber tensile strength [62]. More promising approaches have
suggested forming chemical active groups on the fiber surface which can then react

with epoxy to produce covalent chemical bonds at the fiber-matrix interface.

22

Figure 8 - SEM micrographs of fracture surfaces of unidirectional composites.
(A) AS—4/epoxy composite.
(B) Kevlar 49/epoxy composite.

bar =10 um

 

 

 

 

 

 

 

 

 

 

 

 

 

 

23

Two main types of surface chemical modifications have been attempted;
chemical plasma treatments and wet chemical reactions. Kevlar 49 plasma treatment
has been reported by Allred et al. [63]. Using a RF plasma in the presence of
ammonia gas, Allred has reported a two-fold increase in the interlaminar peel
strength of treated Kevlar 49-epoxy composites and failure mode changes from
interface failure to mixtures of fiber and matrix failure. In the wet chemical reaction
approach, Wu et al. [64] have been able to incorporate amine functional groups on
the fiber surface by brornination followed by ammonolysis, nitration, and reduction.

They have reported results similar to those of Allred.

Aramid-Epoxy Interface

 

An understanding of the interfacial interactions between fiber and matrix is
critical for improving their adhesion. There are two approaches to the investigation
of fiber-resin interface. One approach deals with the microstructural aspects of bond
forming and concentrates on the chemistry and physics of the interface. The other
approach deals with the macrostructural aspects and mechanical analysis of the
interface. The two approaches must be combined for a viable explanation of the
aramid adhesion. The present study consider concepts from both approaches to
develop a practical understanding of the aramid-epoxy adhesion. A brief review of
the main concepts of both approaches and their implications on the aramid-epoxy

interface follows.

Microstructural aspects of polymer-polymer adhesion have been discussed
by Allen [65]. He has explained several mechanisms for polymer-polymer bonding,
each involving combinations of physical and chemical interactions. These
mechanisms are: mechanical interlocking, adsorption interactions, electrostatic

interactions, and interdiffusion of the polymer chain segments.

24

1. Mechanical Interlocking

Mechanical interlocking of two adhering surfaces is a result of surface
irregularities of the two phases. These surface irregularities can act as mechanical
anchors, and high bond strengths can be obtained even though other interactions may
be weak. This mechanism has a profound influence on the adhesion of porous
materials such as wood, paper, and textiles, but for the relatively smooth surfaces of

aramid fibers, mechanical interlocking is not expected to be significant.

2. Adsorption Interactions

Adsorption interactions refer to processes whereby the molecules of one phase
are attracted to specific sites on the other phase. These attractive forces originate
from basic chemical interactions, such as covalent chemical bonds and secondary
chemical interactions. Covalent bonds involve sharing of electrons among atoms, and
are the primary form of chemical interactions (39-1 11 kcal/mole). Secondary
interactions involve electron correlation between molecules and are much weaker
chemical interactions (2—6 kcal/mole). Secondary interactions include nonpolar
dispersion forces (Van der Waals forces), polar dipole interactions, and polar Lewis
acid/base interactions (including hydrogen bonding). All types of adsorption
interactions require intimate physical contact between the molecules of the two
phases [66], but secondary interactions are effective over greater atomic distances /
than covalent interactions and are a prerequisite for covalent interactions. To achieve /

a stable interface, formation of covalent bonds at the interface is very desirable. ,-

The reported results by Allred et al. [63] and Wu et al. [64] on a two fold
increase in interlaminar peel strength of aramid composites by forming chemical
active groups on the fiber surface can be attributed to the formation of covalent .
bonds. Timm et al. [67] have put forth the possibility of chemical reactivity of the

secondary aromatic amines in Kevlar 49, especially in anhydride cured epoxies.

25

. Using FT-IR techniques, Garton [68] has examined the interface of a model PPT A
surface with both anhydride and amine cured DGEBA epoxy. For the aromatic
amine cured epoxy, no significant effect from PPTA on the crosslinking was
observed; however, for the anhydride cured epoxy the effect of the aramid surface
was significant. Garton determined that the changes in the anhydride system were
due to reactions with the adsorbed water on the aramid coating. He has not rejected
the possibility that the amide functionality of the aramid may play a more significant
role at high temperatures. Using photoacoustic FTIR spectroscopy, Chatzi [69] et al.
have determined that 30% of the N-H groups in Kevlar 49 fibers may be accessible

for possible reactions, but the other 70% are sterically inaccessible.

3. Electrostatic Attractions

Electrostatic attraction can result from the transfer of electrons across the
interface, which creates positive and negative charges that attract each other.
Electrostatic interactions are usually significant in metal-polymer and fine particle
adhesion, but for the polymers at close distances the electrostatic forces are usually
small compared to other types of interactions. Inverse gas chromatography studies
by Chappell [70], have indicated that there is ~ 62% increase in the dispersive
component of Kevlar 49 surface free energy when the anti-static coating of the fibers
is extracted. In the absence of perturbation from other types of interactions, the

electrostatic interactions could become significant at the aramid-epoxy interface.

4. Polymer Interdifi’usion

The contribution of interdiffusion in polymer-polymer adhesion has long been
recognized [71]. The extent of diffusion of one polymer phase into another phase
depends on their mutual molecular affinities. For aramid-epoxy adhesion, polymer

interdiffusion is not possible, but macroscopic diffusion of the liquid epoxy polymers

26

into the fiber skin and between or within fibrils is possible.

Macrostructural properties of the adhering phases have pronounced effects
on the bond strength of their interface. Mechanisms of adhesion are influenced by
bulk and surface properties such as mechanical interactions, wetting, and weak
boundary layers. These properties do not directly create interfacial bonding, but they

can enhance or weaken any of the possible adhesion mechanisms.

I . Mechanical Interactions

Mechanical interactions have several origins including surface topography,

thermal stresses, and Poisson contraction.

Surface topography of contacting solids is very important to their adhesion.

 

Solid surfaces are generally rough on a microscale. When two solid surfaces are in
direct contact, the actual area of molecular contact is limited a to relatively few high
points on each surface. A low contact area results in limited interactions and a weak
adhesion. In liquid-solid adhesion, for low-viscosity liquids that can conform to the
surface of the solid, the increase in the surface area due to the surface roughness
improve all adhesion mechanisms. Conversely, when the liquid has high viscosity or
high surface tension, it can form bridges over the rough surfaces and create voids.

The presence of voids or bubbles at the fiber matrix interface is generally detrimental

to good adhesion.

The effects of surface topography can be analyzed in terms of fiictional
forces. Both adhesion and deformation contribute to the friction. The adhesion
contribution is due to rupture of molecular bonds that takes place during friction and
involves local motions of order of 1 nm. The deformation contribution is due to
mechanical interactions of the two surfaces and involves motions exceeding 1 pm.

Friction from surface deformations is always present even when there is no

27

adhesion. For instance, when an inert lubricant is interposed between the two
surfaces, the adhesion friction is eliminated, but there are still deformation frictions.

Additional details and discussions on friction are presented by Cherry [72].

Pull—out techniques to determine the interfacial adhesion either explicitly or
implicitly include the effects of friction. Piggott et al. [73] have reported on pull-out
experiments with carbon and glass fibers. Their results suggest that the normal
pressures and coefficients of interfacial friction are more dependent on the state of
cure than the type of resin used. Their reported values for the coefficient of friction
of carbon fibers vary from 0.42-0.58, with the higher value corresponding to the
higher curing temperature. The reported values for aramid fibers [74] vary from 0.41
to 0.46, which suggests the similarity of its surface topography to carbon. The
presence of coating on fiber can significantly affect the fiiction by modifying the
surface topography. The comparison of pull-out properties of coated Kevlar fibers by
Mai et al. [75] has demonstrated that shearing rate and interfacial viscosity can

significantly affect the interfacial friction of the coated fibers. /

Reedy [76] has reported on a finite element analysis of stress concentrations at
the Kevlar-epoxy interface. Comparing computed frictional stresses with the
experimental data, Reedy has suggested that for Kevlar 49 fiber the friction due to
deformation has a magnitude of roughly 50% of that due to adhesion. Reedy has
also suggested that when debonding occurs, the results of a linear elastic, perfectly

bonded fiber-matrix model are no longer applicable.

Shih et al. [77] have presented a theoretical model to describe the effects of
the interface on the tensile strength of unidirectional composites. Their model
suggests that for an interface strong in adhesive interactions, an increase in bonding
has only a marginal effect on the tensile strength of the composite, but for an
interface with low fiiction an increase in adhesive would significantly enhance the

tensile strength.

28

Thermal stresses are caused by dimensional shrinkage of the resin around the

 

fiber. Cooling, solvent removal, and chemical reactions can cause shrinkage.
Thermal stresses can increase surface contact and enhance frictional interactions, but
they can also cause elastic strains, which upon debonding, act as a locus of failure.
For composites, the thermal stresses are mainly due to the difference in the thermal
expansion properties of the fibers and the matrix. Anisotropic aramid and carbon
fibers have different thermal expansion coefficients in their axial and radial
directions. During cool-down from high curing temperatures, the mismatch between
the matrix and fiber shrinkage can result in radial and axial stresses. Nairn et al. [78]
have discussed the effects of thermal stresses on Kevlar and carbon composites made
with epoxies, amorphous thermoplastics, and semi-crystalline thermoplastics. They
suggest that generally, any matrix—dependent composite property is affected by the
thermal stresses.

Rojstaczer et al. [79] have reported on the thermal expansion properties of
aramid fibers. For Kevlar 49 fibers, they have reported an axial thermal expansion
coefficient of -5.7 ppm/°C, which is different from the -2 ppml°C value reported by '—
the manufacturer [74]. A negative thermal expansion coefficient indicates shrinkage
with increasing temperature. The off-axis mechanical properties of the fibers are
very difficult to determine and they are often back-calculated from the composite
properties. These calculations usually ignore the effects of the fiber-matrix interface
on the thermomechanical properties. Based on a matrix thermal expansion of 65
ppml°C and a Poisson’s ratio of 0.35 for both matrix and fibers, Rojstaczer et al.
have reported a value of 66.3 ppm/°C for the radial coefficient of thermal expansion

of Kevlar.

For carbon fibers, the reported thermal expansion coefficients are -0.1 to -0.5
ppml°C for axial and 7 to 12 ppml°C for the radial expansions [80]. We can infer
from the values of the thermal expansion coefficients that along the radial directions,

29

the thermal shrinkage mismatch between fiber and matrix is much larger for carbon
fibers than for aramid fibers. Comparing carbon and aramid fibers, Penn et al. [81]
have suggested the lower level of thermal stresses in aramid composites as an

important reason for their relatively weaker interfacial adhesion.

Poisson’s ratio differences between fibers and matrix can result in conditions

 

similar to thermal stresses. When a fiber has a lower axial Poisson’s ratio than the
matrix, upon application of axial tension to the composite the matrix shrinks to a
greater extent than the fiber, resulting in radial compressive strains referred to as
"Poisson contraction". This compressive load can increase the interfacial surface
contacts and bonding. Conversely, upon compression, the Poisson contraction can

contribute to the debonding of the fiber matrix.

Axial Poisson’s ratio of the polyaramid fibers has been back-calculated from
their composite properties [82] and is determined to be the close to neat resin values
(0.33 to 0.35). Composite measurements do not yield accurate determination of fiber
Poisson’s ratio. Direct measurement of the fiber Poisson’s ratio are not reported in
literature. For carbon fibers, the reported values of axial Poisson’s ratios is 0.22 to
0.25 [82]. The mismatch of Poisson’s ratio between fiber and matrix is much greater
for carbon fibers than aramid fibers, suggesting greater normal compressive stresses
for the carbon composite under tension. Drzal [83], in comparing the interfacial
behavior of aramid and carbon composite, has suggested that both the lower thermal

and Poisson’s ratio mismatch of aramid—epoxy are adversely affecting their adhesion.

30

2. Wetting

Adequate wetting of the fiber surface by liquid resin is a prerequisite for good
bonding. The importance of wetting in adhesion has been dramatically demonstrated
by Sharpe [84]. When liquid epoxy was cured on solid polyethylene, the adhesion

I was low, however, solidification of molten polyethylene on the cured epoxy produced
a much stronger adhesion. In the first case, the liquid epoxy with high surface
tension did not wet the solid polyethylene with low surface tension; but in the second
case, liquid polyethylene with low surface tension could spread over high surface

tension solid epoxy.

Using contact angle analysis, Penn et al. [81] have determined that the surface
energies of carbon and aramid fibers are very similar. Li et al. [85] have studied the
wettability of carbon and aramid fibers by both the Wihelmy and the solidification
front techniques. Wihelmy technique yielded 42.4 mN/m and 43.7 mN/m values for
carbon and aramid fibers surface energies, respectively. Solidification front
measurement resulted in 41.8 mN/m and 46.4 mN/m for carbon and aramid fibers,
respectively. A report by Wesson et al. [86] has shown that surface energies of
aramid fibers and liquid DGEBA resin are similar. From these results, both carbon
and aramid fibers are expected to have similar "good" wetting with liquid epoxies.

3. Weak Boundary Layers

Weak boundary layers refer to surface layers at the interphase with lower
cohesive or adhesive properties than their bulk substrates. Weak boundary layers can
be due to entrapped gas, contaminants, and structural anomalies of the substrates.
They are usually detrimental to adhesion and can prevent the formation of strong

adhesion even though extensive interfacial contact might be present.

31

There are two types of weak boundary layers, adhesive and cohesive. In an
adhesive weak boundary layer only the interface is affected and there is little damage
in either fiber or matrix, but in a cohesive weak boundary layer, one or both are
damaged. In an fiber-matrix adhesive failure there is clean fiber and matrix
separation, but in a cohesive failure there is substantial structural damage to either
fiber or matrix. Both types of weak boundary layers could coexist, but depending on

the mechanical state of stresses, the effect of one can mask the other.

Existence of cohesive weak boundary layers for untreated carbon fibers has
been demonstrated by Drzal et al. [44]. Their study has shown that untreated carbon
fibers possess a weak structural layer that can not support high shear loads. Surface

treatment of the fiber removes the defect layer and substantially improves adhesion.

For aramid fibers, the possibility of cohesive weak boundary layers exists.
Examination of Morgan’s model of aramid’s skin-core morphology indicates that the
PPTA monomers are more randomly dispersed on the fiber exterior but are
progressively more clustered towards the core. A reduced degree of crystallinity on
the skin could result in weaker properties in the surface region. Upon application of
shear to the fiber surface, cracks could be generated in the skin region producing an

inefficient fiber-matrix load transfer.

There are numerous indications that surface treatments that do not affect the
aramid surface morphology are ineffective in improving its epoxy adhesion. Fiber
pull—out experiments by Miller et al. [87] do not suggest any difference between
bond strength of "as received" and "acetone washed" aramid fibers. As received
fibers have a sizing on them which should introduce an adhesive weak boundary
layer, yet the fiber—matrix bond strength is unaffected. Penn et al. [61], have applied
various coupling agents to promote aramid-epoxy adhesion and have reported no
significant changes in the fiber-matrix bond strength. In a more recent study by Penn

et al. [88], flexible reactive pendent groups have been covalently attached to the

32

aramid fiber surface, but no mechanical improvement of interfacial bonding was
observed. Penn has attributed this observation to a lack of bonding between the
pendent groups and the matrix; however, failure in a cohesive weak boundary layer
could be another explanation. Data by Kompaniets et al. [29] have demonstrated that
for the aramid fibers the fiber tensile strength is the gage length dependent, but for
the unidirectional composites the tensile strength is independent of the gage length.
This observation suggests a different failure mode for the embedded fibers which

experience shear load transfer.

Theory

Interfacial Shear Characterization

 

Single fiber techniques to examine fiber-matrix interaction are inherently less
complex than multi-fiber composite techniques. Coupling of various fiber-matrix
interactions in a composite, complicates the examination of it fiber—matrix interface.
In this study, a single fiber testing method has been utilized. The technique involves
embedding a single fiber in a resin matrix tensile dogbone coupon. The sample is
subjected to a tensile load and the transfer of the load through shear at the interface
causes fiber fragmentation. The fiber fragmentation process continues until the
remaining fragments are too small to transfer enough load to cause another fiber
fracture. This final length of fiber is twice a dimensional limit called the critical
length (lc). Critical length is characteristic of the level of fiber-matrix adhesion as
well as the fiber and matrix properties. For a system consisting of a fiber fragment
surrounded by an unbounded matrix, a relation between (lo) and interfacial shear
strength (1],) was first demonstrated by Kelly and Tyson [89]:

r. = 32%?) (1)

C

where ouf is the fiber tensile strength and R is the fiber radius.

33

34

For this system, load transfer to the fiber takes place by shear forces,
therefore, T}, by equation (1) provides a good characterization of the shear strength of
the interface. The value of In obtained by equation (1) is the average shear over the

fragment.

The shear force at the fiber-matrix interface is the combined result of all the
stresses acting at the interface, however, equation (1) does not explicitly specify these
forces. To analyze the fiber-matrix interaction a more rigorous model is required.
An analytical model of the three-dimensional stress distribution around an isolated
fiber fragment has been proposed by Whitney et al. [1]. Whitney’s model proposes
an approximate solution using linear elasticity equations. The main assumptions of
the Whitney’s model are:

a) The system is axisymmetric.
b) Both fiber and matrix undergo elastic deformations.
c) Fiber is transversely isotropic.

(1) Fiber and matrix strains match at their interface, i.e. perfect bonding.

Complete solutions of Whitney’s model are presented in Appendix A, but of
particular concerns to this study are the shear (tn) and normal stresses (0,) produced

at the interface. The interfacial shear and the normal stresses are given by following

equations:
1,, (r3) = 4.75 a A1 e0 reef-75f (2)
o, (:12) = [A2 - 11,2 A1 (1 - 4.75;?) [MSW] so (3)
where 5:“: W 's a dimensionless length along the fiber axis and so is the far field
.-

axial strain. Whitney has described so as the strain at the onset of observed fiber

breakage. Since the testing jig is not very accurate for strain measurements, we

35

propose another definition:

80 = euf "' GMAT (4)
where euf is fracture strain of the fiber, or," is thermal expansion coefficient of the
matrix, and AT is temperature difference between ambient temperature (25°C) and

curing temperature. This definition should be close to actual strain at the critical

length. Other material property constants tti, A1, A2, and Kf are define by:

l
G _
¢=[ " ]2 (a
Elf — 4vam

 

a 4K v
Al: E1/(1-—lf-)+—JW-T—12£

 

 

 

Kf'i'Gm

1+ - - -

x[v12f'v,..+( v") WeoW VlWlf] (6)
E

KfW ' m (7)

2 2__E_2f__22’_21£21

202, 13,,
_fl(&_ (1+Vflfim-F’2f'vlzzlf

Az- Kf'i-Gm [Vlzf—Vm'f' £0 (8)

Whitney has defined the critical length 1,. such that 95% of the applied axial
load is transferred to the fragment. The model permits the calculation of the critical

length:

2.37519
4

which depends explicitly on fiber and matrix material properties, but not on

 

lo:

(9)

temperature.

36

Average values of shear and normal stresses are obtained by integrating

equations (2) and (3) with respect to f and then dividing by the fragment length i.e.

 

 

jigorx) d'x—
WI? = O _
x
[exam at?
e, = 0

After some manipulation the following results are obtained

¢Aleo [1 - est-75f (4.75r+ 1) ]
TX, =
4.75 r

 

“o“. = (42 — Ali? 84.75;) 60

The theoretical average values are evaluated at at f: 1 resulting:

r,,(iheo) = 0.200 a Al .20

e,(iheo) = (A2 - 0.00865 A102 ) £0

(10)

(11)

(12)

(13)

(14)

(15)

Experimental

Aramid fibers utilized in this study were Kevlar 49 supplied by E J. du Pont
(lot # 123481, 7200 denier). The fibers were supplied with a sizing. To eliminate
possible interferences by fiber sizing, the fibers were washed by absolute ethanol.
The washing procedure involved the following steps:
1) Removal of the exposed surface layer of the fibers from the spool
and a careful selection of a ~ 10 inch fiber tow.
2) Three successive washes in absolute ethanol with six hours soaks
between each wash.
3) Two hours drying of the fibers in a vented oven at 125°C.
4) Storage of the tow in clean aluminum foils inside a dedicated

desiccation chamber.

Carbon fibers in this study were AS-4 fiber, supplied by Hercules (lot # 708-
4C, 12000 denier). These carbon fibers had no sizing and did not require any

washing. They were carefully selected and stored in clean aluminum foils.

37

38

The selected epoxy resin was D.E.R. 331 supplied by Dow Chemical U.SA..
This resin is a Diglycidyl Ether of Bisphenol A (DGEBA) epoxy. The curing agents
for this study were DiEthyleneTriAmine (DETA) and a mixture of two curing agents
M-PhenyleneDiAmine (MPDA) and DiEthleolueneDiAmine (DETDA). The
D.E.R.33l/DETA system contained :1 110/100 mass ratio of curing agent to epoxy.
The mixture was degassed in a vacuum oven for 15 minutes at —29 in.Hg (gauge
pressure). DETA is highly reactive and its epoxy mixture was degassed only at room
temperature to avoid gelling. For D.E.R.33l/MPDA/DETDA system, a 7.25/100
mass ratio of MPDA and a 11.75/100 mass ratio of DETDA were combined; this
system is referred as (50/50) mixture. The (50/50) mixture was degassed in a
vacuum oven at 75°C for 5 minutes at -29 in.Hg to reduce the viscosity of the

solution as well as removing entrapped bubbles.

Characterization of the interfacial adhesion was obtained by a critical length
measurement technique. The sample was a cured resin dogbone with a single fiber
embedded along its center. For the fabrication of the dogbone specimen, a single
fiber was carefully hand separated from a tow. The fiber was then mounted in a
silicone mold and held in place with a small amount of rubber cement. The mold
was left exposed to air for 30 minutes permitting the cement to dry. Subsequently,
the mold was placed inside a ventilating oven at 150°C for 1 hour to allow the fibers
to desorbe moisture gained during the air exposure. An epoxy resin was then poured
into the mold and the fibers were slightly pulled to straighten them. The mold was

placed in an oven and the appropriate curing schedules were executed.

A cured dogbone sample was subjected to a tensile load using a tensile testing
jig (Figure 9). The load transfer to the fiber through shear at the interface causes the
fiber to fragment. The fiber fragmentation process was continued until the fragments
reached their critical length (lc). The resin matrix was transparent and the

fragmentation process was monitored under an optical microscope at 150x

39

Figure 9 - A sample mold and the tensile jig for the critical length technique.

 

    

Strain Gage

 

 

 

4o

magnification. For the carbon fibers the critical lengths were directly measured using
a motorized stage with a displacement readout. For aramid fibers, the failure process
was by fibrillation and did not have well defined edges. The average critical lengths
of the aramid fibers were obtained by counting number of failed regions within a 22
mm fiber length.

Four ranges of curing temperatures (Tc) were selected for this study, 25°C,
75°C, 125°C, and 175°C. For the 25°C and 75°C curing temperatures, the DETA
curing agent was chosen. DETA is one of the most reactive amine curing agents
available, however, at low curing temperatures its epoxy system was still too brittle
for the critical length testing. Subsequent post-curing of the DETA systems was
required to increase its fracture strain. The post-curing time and temperatures were
determined by the glass transition (T3) of the matrix. During the post-cure, the oven
temperature was maintained below the Ts of the matrix. This was to avoid building
up thermal stresses. At each post-curing temperature, initially the T8 was only a few
degrees above the oven temperature. After a certain time the glass transition
temperature was increased allowing the oven temperature to be raised in steps of

10°C. For the DETA systems, the following curing schedules were selected:

25°C for 48 hours, followed by 4 hours post-curing at 40°C, 50°C,
60°C, 70°C, and 80°C.

75°C for 8 hours, followed by 4 hours post-curing at 85°C.

41

For 125°C and 175°C curing temperatures, the mixture of MPDA/DETDA
(50/50) curing agents was utilized. This system required no post-curing
and the following curing schedules were selected :

125°C for 24 hours

175°C, for 3 hours

Thermal expansion and glass transition of the matrix were determined by a du
Pont thermal mechanical analyzer (TMA model 943). A 18x4x4 mm sample mold
was used to ensure consistent sample geometry. Appendix C includes TMA plots for
the thermal expansions and glass transitions of the resin systems. Initial curing times
were estimated by differential scanning calorimetry (DS C) performed at isothermal

temperatures.

The elastic modulus of the epoxy matrices at the ambient temperatures were
determined by a dynamic mechanical analyzer (DMA), ran isothermally at 1 Hz fixed
frequency for 5 minutes. To obtain reproducibility in the results, five samples for

each curing temperature were tested.

Direct observation of fiber-matrix interface was obtained by transmission
electron microscopy (TEM). Ultrathin sections of fiber embedded matrix were
microtomed using a diamond knife. The microtoming technique is described in
details in appendix D. The sections were examined with a JEOL CX100 TEM with
magnifications up to 320,000x. Some sections were stained with 0s04 to enhance

some particular features.

Results and Discussion

Thernurl Stresses

 

Cooling of a composite from oven temperature to ambient temperatures results
in thermal, shrinkage of the, fiber and the matrix. ~_'I_’o examine effects of thermal
sgessesgonfibervmauix adhesion, samples of epoxy dogbones were cured at ambient
(nominally 25°C), 75°C, 125°C, and 175°C curing temperatures. Figure 10 illustrates
the TMA results for the thermal expansion of these epoxy systems which corresponds
to 70 :l: 6 ppm/°C thermal expansion coefficient. The epoxy system cured at 175°C
has the Ts: 160°C and for this system the thermal stresses are expected to initiate
from 160°C. 1113929}! system cured” at the room temperature“ was, subjected to an
extensive p9§t7gufinggwhicjli resulted in;- 0.10 % post-curing shrinkage. Jhe._75°C
cured-epoxy was also post-cured to. increase its fracture strain Post-curing
temperatures were selected so that oven temperatures was always lower than glass
transition of the resin. Such a post-curing scheme has minimized the alteration of the
initial thermal stresses. Appendix C contains the TMA plots for the thermal
expansions of the epoxies and the post-curing schedules for the 75°C and room cured

epoxies.

42

43

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44

The experimental carbon fiber data has been analyzed by a Weibull model,
but a normal distribution has been used for the aramid data. The brittle fracture of
the carbon fiber permits the measurement of the individual fragments and with its
short critical lengths, a large number of fragments were measured for each sample
dogbone. The Weibull model is a two parameter model that is often used to describe
unsymmetric failure distributions. The large data population of the carbon samples
should provide a good estimation of the Weibull constants, however, the carbon fiber
is very flaw sensitive and has a large tensile properties variations which results in
large intrinsic variations in the critical length distribution. For the aramid fibers, the
indindual,__fr_agmcnts could not , be accurately. measured because of the fibrillation that
occurs at.it§-£ail.§.§.w!981;9ns. The average aramid critical lengths \yt‘ft'gfdflflmincd by
4311113529-}an Of the. fiber -by the total number, 0f".0b§medhfailed..rcgions. The
experimental aramid critical lengths are much longer than the carbon critical lengths,
but the aramid data have relatively small variations. With a smaller aramid data
population the Weibull constants can not be accurately estimated so a normal

distribution is used instead.

In the following results, the carbon fiber data indicate much larger error bars
than the aramid fiber data. The larger carbon data variation is due the intrinsic
variation in the carbon fiber tensile properties rather than the experimental
variations. For example, 500 and 1000 data populations for the gold coated 125°C
cured carbon samples both have similar mean and variations. The mean values of
the carbon data populations are very consistant and the large error bars should not

prevent us from making valid judgements on the possible correlations.

Expenmeng interfacial shear strength (“r”) values of untreated aramid-epoxy
amjagbon-epoxysystems at different curing temperatures are presented ,in Figure
11. For the examined curing temperatures, the aramid samples exhibit 3.5 to 5 times

lower Tu than the carbon samples. The mearbon, samples indicate increased fiber-

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46

m@wstm__m increasing _091,i!}a1¢rr_lneratu_rc§. but th¢ Wflafimlssfiseems
to be, unaffected ,by the curing temperature. During cool down from oven
temperature to ambient conditions the isotropic matrix shrinks uniformly, but the
orthotropic fiber undergoes longitudinal expansion and radial shrinkage. The overall
radial thermal stress at the fiber-matrix interface is due to both radial thermal

shrinkage of fiber and matrix and the Poisson contraction resulted from the fiber

longitudinal strains. Warm .independchQDmeP .3? 3.3347913?“ adhesion "

can be attributed to the close match between the Poisson’s ratio and the radial "I

thermal expansion coefficient of the aramid fiber and the epoxy matrices.

Chemical Bondig

 

Coatingge fibers by an inert material such as gold eliminates all [possible /’

coyalentfbonds at the fiber-matrix interface/The coating also introduces a k weak

/

boundary“ layer, but for relative comparisons the effect of the layer should be /-

constant/Figure 12 demonstrates the effect of the two different epoxy systems on
the interfacial shear strengths of gold coated and uncoated fibers. The uncoated
carbon samples exhibit smaller changes between the 75°C cured DETA system and
125°C cured MPDA/DETDA system than the corresponding gold coated carbon
fibers. For the aramid adhesion no trends with the curing temperature or epoxy

system for either coated and uncoated fibers are observed.

At the two curing temperatures, the difference in the gold coated samples
should be mainly due to thermal stresses and Poisson contractions. From 125°C to
75°C curing temperatures, the uncoated carbon fibers show a lower reduction in
adhesion than the gold coated samples. This observation can be attributed- to. the
compensation of the lower thermal stress at 75°C curing temperature by the greater
chemical reactivity of its resin. The DETA curing agent is chemically more active

an the MPDA/DEEA__enrin_g agents which increases, the ”extent . of possible

47

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48

challicalwhondings at the fiberematrix interface. For the . aramid fibers, both coated
and uncoated fibers exhibit the same levels of adhesion at the two curing
tcmperanuese The relative adhesion reduction for the gold coated aramid samples is
less than the corresponding carbon samples. The aramid. observations. assesses
@9336??? 113129. shemical reactivity of the, resins does net, effws the aramid-

eppxy adhesion.

Poisson Contraction

 

The Poisson contraction refers to the stresses due to the contraction of the
mgtmg thena. load. is applied to the composite.” _Figure 13 compares the interfacial
shear strength of gold coated, silicone coated, and uncoated aramid and carbon
samples made with the room temperature cured epoxy. The carbon samples
generally exhibit higher 'Tu values than the aramid fibers for both coated and uncoated

film‘- Thc room cured matrixrhasr gym .loyv. thermal“omisfldmtlrergatcdfihors

 

 

hale no covalent ehernical bonding at their interface. EQI‘flIQHQQQtQflngS, the main

interfacial interactions, present are the Poisson contraction, fibererrratrix wetting, and

 

u... ”Wm

the silicone negated carbon fibers have i, values two and half times the values of
corresponding aramid fibers. Since the coated aramid and carbon fibers have the,
serge sflurfaacempmperties, the higher adhesion of the coated carbon fibers indicates the
effect of Poisson contraction. Note that for the uncoated carbon fiber in is about four
times higher than the uncoated aramid fiber which can be attributed to the effect of

the chemical bonding and removal of the coating weak boundary layer.

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Fiber Wetting

 

Geeéflfine .95 the. Why. the. Bettie ,‘Ziii’llf‘crfasii9191¢921¥.,i9!eteet1998
wimfifi§9¢s.-flhich-.,m., turn. enhances the load transfer .t9.,Fh€H§§e_£ Figure 14
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giyee “1-15%.reductien inj,.while. siliconescoeted fiber?» £99399: ,.7§0./‘39!EC19_99LQD-
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with very poor wetting properties. The drastic reduction of the silicone coated fibers
clearly mpgegrgeewme--.mpormce of achieving thermodynamic wetting as a

Ewen: foredhesions

Three Dimensional Stress Model

 

The experimental values of the interfacial shear strengths have been
determined by equation (1), however, the effects of the thermal stresses and Poisson
contractions are not explicitly present in the equation. A three dimensional stress
model proposed by Whitney was examined to acquire additional insights to the
interfacial interactions. The material properties for fibers are presented in the Tables
1. The Poisson’s ratio of 0.35 is assumed for the matrices. The ambient matrix
elastic modulus have been determined by DMA and are plotted in Figure 15. The

matrix thermal expansions are given in the Figure 10.

For the aramid and the carbon fibers, Figure 16 compares the theoretical
critical lengths calculated by equation (9) with the experimental critical lengths for
the uncoated fibers. The data indicate that the theoretical results for the carbon fibers
are close to the experimental results, but for the aramid fibers there are factors of
four differences. Due to model assumptions, the theoretical critical lengths do not
show the same temperature trends as the experimental results. The model defines the
critical length as the length required to recover 95% of the applied longitudinal stress

on the composite and the resulting critical length equation is:

51

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which does not include any temperature term.

The model permits the evaluation of stress and displacement components.
Figure 17 compares the experimental and the theoretical interfacial shear strengths
calculated by equations (1) and (14). The experimental and the theoretical '17,, values
for the carbon fiber show much smaller differences than the result for the aramid
fiber. For both fibers, the theoretical results indicate increasing in values for higher
curing temperatures, although to a greater extent for the carbon fibers than for the
aramid fibers. The model predicts good interfacial shear strength for both types of
fibers.

To understand the discrepancies between the experimental and the theoretical
results, the assumptions of the model needs to be closely examined. The model
assume that both fiber and matrix undergo elastic deformations, however, the fiber
fragments or matrix may actually behave nonelastically. Reedy [7 6] has suggested
that when debonding occurs, the results of linear elastic, perfectly bonded fiber-
matrix models are no longer applicable. Figures 18 show the optical micrographs of
aramid and carbon fiber fragments under bright-field and cross-polarized lights. For
the carbon fiber, the fiber-matrix failure is by a combination of interface slippage and
matrix cracking, but the fragment and its surrounding matrix should behave elasticly.
The carbon fragment shows no longitudinal failures and the cracked matrix may only
display inelastic behavior near the crack regions which is only an small portion of the
fragment length. Another assumption of the model, the matching strains at the fiber-
matrix interface may also not be valid for the fibers. Sometimes small gaps between
carbon fi'agments are observed indicating interface slippage. For the aramid

fragments it is difficult to assess the interfacial slippage. The low intensity and

54

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56

dispersed stress fringe patterns of the aramid-epoxy interface may be due to either
interface slippage or internal fiber realignments due to surface to bulk fractures.
Presence of a strong chemical covalent bond at the fiber-matrix interface should
reduce the fiber-matrix slippage. Finally, the critical length equation which does not
take into account the effects of thermal stresses may introduce some additional errors

in the theoretical results.

The failure of the Whitney’s model to reasonably predict the aramid
interfacial interactions suggests the severity of the violation of model’s assumptions.
The inelastic longitudinal failure mode of the aramid fiber may be the principal
reason for the model’s deficiency. Upon the onset of the longitudinal cracks the fiber
no longer behave elastically and in fact a substantial decrease in the fiber modulus
are expected. Other fiber properties such as Poisson’s ratio or thermal strains may
also be altered by the fiber longitudinal fracture. The matrix is expected to maintain
an elastic behavior since no matrix crack is observed. The failure of the perfect
bonding assumption may also skew the result. Aramid fragment separation has not
been observed, but internal fibril slippage is possible. The temperature independence
of the critical length equation is not expected to introduce severe errors to the aramid

results since the experimental results are also temperature independent.

The interfacial radial stress is due to radial thermal strains and Poisson
contractions. Since fiber and matrix are in series, these radial strains should be fully
transferred despite any possible interfacial slippage. The fiber is embedded in a large
matrix environment, thus, the matrix has greater displacements and its properties
should dominate the fiber properties. The interfacial radial stress is therefore
expected to be much less dependent on the extent of fiber-matrix bonding and moduli
of the fiber than other stresses. Equation (15) represents the theoretical average
radial stress:

6',(theo) =- (A2 — 0.00865 A1¢2 ) so (15)

57

Figure 18 - Optical micrographs of fiber fragments at their critical lengths.
(A) Kevlar 49 under bright-field light.
(B) Kevlar 49 under cross-polarized light.
(C) AS-4 under bright-field light.
(D) AS-4 under cross-polarized light.

bar=100tt

 

 

 

58

An inspection of the equation suggests that the equation is more dependent on
the value of the A2 constant than the A1 constant. A numerical examination of the
equation (15) shows that for the aramid fiber A1 term has at most 20% contribution
to the 6', value and for the carbon fiber at most 4% contribution to the 6', value.
Inspection of constant A2 reveals that it is mainly determined by the matrix elastic
and shear modulus. The matrix modulus are expected to conform reasonably to
elastic behavior, therefore, the magnitude of theoretical radial stresses should be

approximately valid for both fiber systems.

Figure 19 compares the theoretical average radial stresses for both the aramid
and the carbon systems. The carbon systems show three to five times higher radial
stress values than the aramid system. The carbon system also shows higher 6‘,
increases with curing temperature than the aramid system. The higher 6, value of
the carbon system can be attribute to its higher thermal stresses and Poisson
contractions. The temperature trend of the aramid system is mostly due to the

Poisson contractions caused by its longitudinal thermal stresses.

The theoretical analysis for the aramid fibers indicates increasing radial
stresses with increasing curing temperature, however, the experimental interfacial
shear strengths are found to be independent of the curing temperatures. The
theoretical interfacial shear stresses are also much greater than the experimental
values. These observations suggests that the aramid-epoxy adhesion is substantially
limited by a failure process which prevent the interfacial stresses from reaching their
theoretical maximums. Altering and improving the aramid-epoxy failure process has

the potential of increasing adhesion to levels comparable to carbon-epoxy adhesion.

59

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Electron Microscopy

 

Direct observations of the fiber-matrix interface is possible through
transmission electron microscopy (TEM). The technique involves ultrathin sectioning
of the composite and high-magnification transmission electron imaging of the fiber-
matrix interface. The details of the sectioning technique are described in appendix
D. Care must be taken in the interpretation of the images to avoid mistaking the
sectioning artifacts for actual fiber-matrix interactions. During sectioning, the cutting
knife exerts bending and compression on the sections which can result in fracture,
thickness variation, and folding of the fibers at normal to the cutting direction. In the
following micrographs, the direction of the sectioning is indicated by an arrow, fiber
and matrix are identified by letters F and M respectively, and the designated

magnifications are shown by scale bars with better than 10% accuracy.

An ultra-thin section of a aramid-epoxy composite cut parallel to the long axis
of the fiber is shown in Figures 20. The aramid fiber ribbons generally show
interfacial separations with some fibrils remain attached to the matrix. The aramid
ribbons easily separated from the epoxy matrix and it was difficult to find sections
where the fiber ribbon was still connected to the matrix. The ribbon shows repeated
bands of knife marks along its width.

The aramid fiber was also sectioned normal to the fiber longitudinal
direction. This sectioning direction exerts both tension and compression loads on the
fiber-matrix interface which provides more insight to the fiber-matrix interactions.
Figures 21 and 23 to 26 illustrate aramid-epoxy composite sections cut along the
fiber radius. Figure 21 illustrates the fibrillation that takes place upon tearing of the
fiber. Similar fibrillation are also seen at the aramid-epoxy interface (Figures 23 to
26). Aramid fibers sometimes display a torn skin in the form of helical ribbons as
shown in Figure 22; radial cross—sections of such fibers are shown in Figures 23 and

24. The fiber-matrix separation again takes place through the fibrillation of the

61

aramid fiber. Figure 24 has concentrated on a portion of a skin-tom fiber that is
suspected to be the region where the skin was originally located. The section in
Figure 24 has been stained with 0.904 which reacts with the matrix and increases its
contrast. A high magnification of this portion, shows some aramid fibrils which are
still connected to the matrix. Finally, aramid fibers which have been damaged by the
interfacial shear test have been examined. The cross-section of the longitudinal
cracks which are the typical form of the aramid fracture can be seen in Figure 25.
The longitudinal cracks had propagated to the fiber core dividing the fiber into
pieces. Another damaged aramid section is shown in figure 26. A high
magnification observation of the interface illustrates the uniformity of the aramid-

epoxy interface suggesting a good fiber-matrix surface interactions.

The morphology of the aramid fibers can describe their observed mode of
interfacial failure. The aramid fibers are constituted of large polymer macro-
molecules which only form hydrogen bonding with their adjacent molecules. Such
structure tends to initially fail across the weak hydrogen bondings resulting in a
failure by fibrillation. When the load transfer is by shear forces, the fiber exterior
receives the load and distributes it to the rest of the fiber. For the aramid fibers, it is
conceivable that the fiber exterior could be of lower strength and modulus than the
fiber core and long critical lengths are required before the shear load is fully
transferred to its core. Furthermore, the aramid fiber failure could be initiated by

surface cracks which then propagate towards the fiber core.

62

Figure 20 - TEM micrographs of Kevlar 49 and MPDADETDA epoxy system cured
at 175°C. The section is cut parallel to the long fiber axis.
(A) bar = 2 pm
(B) bar = 50 nm

 

,.
H4.

"a

63

Figure 21 - TEM micrographs of Kevlar 49 and room-cured epoxy system.
The section is cut along the fiber radial direction.
(A) bar = l um
(B) bar = 500 nm

 

Figure 22 - SEM micrographs of a single Kevlar 49 fiber.
(A) bar = 10 um
(B) bar = 5 um

 

 

 

65

Figure 23 - TEM micrographs of a surface damaged Kevlar 49 and MPDADETDA epoxy
system cured at 175°C. The section is out along the fiber radial
direction.
(A) bar = 2 um
(B) bar = 500 nm

 

 

 

#—

 

Figure 24 - TEM micrographs of a surface damaged Kevlar 49 and MPDADETDA epoxy
system cured at 175°C. The section is stained with 0s04. The section
is cut along the fiber radial direction.
(A) bar = 1 mm
(B) bar = 100 nm

 

67

Figure 25 - TEM micrographs of a shear damaged Kevlar 49 and DETA epoxy
system cured at 75°C. The section is cut along the fiber radial
direction.

(A) bar = 1 pm
(B) bar = 500 nm

 

68

Figure 26 - TEM micrographs of a shear damaged Kevlar 49 and DETA epoxy
system cured at 75°C. The section is out along the fiber radial
direction.

(A) bar = 250 nm
(B) bar = 25 nm

 

Conclusions & Recommendations

The results of this study lead to some important conclusions on the aramid-
epoxy bonding mechanisms. The results also point out the potential approaches to
improving aramid-epoxy adhesion. Some recommendations for further studies on

aramid-epoxy bonding are presented.

Conclusions

 

Resin shrinkage does not affectuthe adhesion *ofthe aramid fibers to, epoxy
matrices. The predicted higher radial compressive thermal stress at higher curing
temperatures is not observed experimentally, reflecting a curing temperature

independent failure mode.

The close matching of aramid radial thermal expansion and its Poisson ratio
{9-131089 of the matrix results in lower mechanical fiber-matrix interactions for the
aramid-epoxy interface than for the carbon-epoxy interface? However, theoretically if
the aramid-epoxy interfacial failure mode is altered, the aramid-epoxy adhesion has

the potential of up to four times higher adhesion levels.

69

70

High magnification TEM aramid-epoxy interface observation along with
coated fiber experiments have shown that the fiber is thermodynamically "wet" by the
matrix. Increasing the aramid surface tension is not expected to improve its wetting

since the interfacial contacts are already very efficient.

The aramid interfacial failure is made by fiber fibrillation of the outer
surface. This observation suggests the presence of a cohesive fibrillar layer on- the
fiberflexterior which fails at low shear levels resulting in inefficient fiber-matrix load

and low values of interfacial shear strength.

Aramid fibers lack the interfacial chemical bonding present in carbon fibers,
but formation of chemical covalent bonds that only affect the aramid surface are not

expected to overcome the fiber surface structural limitations.

Recommendations

 

These conclusions explain the failure of some of previous attempts to improve
the aramid-epoxy adhesion and suggest other potential approaches for improving the

adhesion. Recommendations based on the results of the present study are suggested.

First, the aramid fiber exterior needs to be modified before significant /“
adhesion improvements can be achieved. Although, fiber surface removal by etching /
techniques has been shown to improve the adhesion, but they are usually
accompanied by losses in fiber tensile strength, possibly by inducing surface flaws.
The fiber surface modification should not result in excessive losses of tensile
strength. The fiber skin alteration can be done by changing of the fiber morphology
during the manufacturing or through chemical crosslinking of the aramid polymer
chains after the manufacturing. Morphological modifications during the fiber
manufacturing are beyond the scope of the present study, but several ways to produce

chain crosslinkings can be suggested.

71

The aramid polymers have amine functionalities available which could act as
sites for chain crosslinking. Works by Chatzi [69] et al. has determined that 30% of
the N-H groups in Kevlar 49 may be accessible for possible reactions. The
hydrophilic nature of the aramid fibers permits penetration of water soluble reactants
and wet chemical reactions are good possibilities. Plasma ion treatments can deposit
chemical active sites within the fiber permitting epoxy crosslinking within the fiber
exterior. Such matrix penetration could increase the efficiency of the shear load
transfer to fiber. A novel chemical kinetic approach would be ion implantation.

Metallic ions could act as catalyst for any number of chemical reactions.

Coupling agents that only react with the aramid surface are not expected to
improve the aramid-epoxy adhesion since the cohesive weak boundary layer is not
affected by the coupling agents. However, if the fiber skin failure is moderately

improved then a coupling agent approach could have potentials.

Appendix A

Appendix A

Three-Dimensional Stress Model

Whitney et al. [1] has proposed a solutions for three-dimensional stress
distribution around an isolated fiber fragment surrounded by an unbounded matrix.
Main assumptions of the Whitney’s model are :

a) The system is axisymmetric.

b) Both fiber and matrix undergo elastic deformations.
c) Fiber is transversely isotr‘Opic.

d) Fiber and matrix strains match at their interface,

i.e. perfect bonding is assumed.

Complete solutions of the Whitney’s model are followings :

04m = [1 — (1 + 4.75)?) 64-75" A180 (A-l)
"- - A—A 22-i1—475 “”5? A2
04(x.r)- 2 NH RZX . fie so (-)
Inf-(5:7) = 4.75 0A1 £0 (%)x‘e'4'75f (A-3)

2
u,,, (zr) = { 2.375 Elfn A, [1 — (%)

_ R
+ 2.375 f(l + 4V12fl2) ] 84'75x} Tb—ET— (A-4)
1f

72

73 Appendix A

 

v4 (r) = A3 r so (A-S)
Em R _
0m (it) = [Ema - a) - A1(7)(1 - 4.752) 5475*] so (A-6)
-— 2 —475— f 2
5m; 0"") = [A2 — A1¢ (1 — 4.753 8 ‘ J‘](7€) 80 (A-7)
r 3 — 5
cm can = 4.75 (pt-1;) 1411750 e475" (A-8) t3
it “
R 2 3
uxrm (1:7) = 2.375 Es}: [El‘fY—‘i' A1 (1 + 4V12j¢2) (%) x—e—4.752] (A-9) _
j
A5 y
Vrm (f) = (A4r + T) 80 (A-IO) '

Constants of these equations are :

El 4K V
f m

 

 

 

 

 

 

1 _— _
x [vuf—vm+( +V"‘)€’" 81” Vlz’glf] (A-ll)
£0
2K0," (1+vm)€ -EZf-vzj€
_ f _ m 1 1f
Az— Kiwi-Gm [VIZf Vm+ £0 ] (A-12)
_ [Kf’VIZf-l- Vme — Kf (62f+ “2151i: Gm(1 + VIM-Em]
A3 = (Kf+ Gm) (A-13)
1+ -
A4: ( v") 8’" —vm (A-14)

30

74 Appendix A

 

 

 

RZK €+v ——(1+v,,,)
_ f 2f rzfilf
A5 — 79—476,; [Vm — V12f+ £0 ] (A-15)
E
Kf— ”' (A-16)
2 [ _ 132, _ 2v2115y]
262, Elf
G .1.
2
¢ = ’” (A-17)
Elf — 4"]me
Em E amAT (A-18)
€le aleT (A-19)
Elf; ayAT (A-20)
80 = 8.4- E... (A-21)
r: l (A-22)
10

Note that AT 5 0.

Whitney has defined the critical length [C such that 95% of the applied axial
load is transferred to the fiber i.e.

oxf (16) = 0.95 Also

The interfacial stress distribution obtained from the model are illustrated in

the Figure 27.

Appendix A

   
     

 

 

Appendix B

Appendix B

Critical Length Distributions

The following diagrams display the experimental critical length distributions
for the coated and uncoated aramid (Kevlar 49) and carbon (AS-4) samples cured at
differ€=nt curing temperatures. The carbon fiber data have been modeled by Weibull
diStl‘ibutions and their plots include the Weibull constants as well as their Weibull
mean and standard deviations. For the aramid data a normal distribution has been

utililtzd and their corresponding mean and standard distributions are listed.

76

Appendix B

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Appendix C

Appendix C

Thermal Expansion & Tg Data

TMA plots of % thermal expansion of the 75°C, 125°C, and 175°C cured
epoxies are presented. The plots also show the T8 of these epoxy systems. Plots of

post-curing schedules for the room temperature and 75°C cured epoxy systems are

also included.

101

Appendix C

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Appendix D

Appendix D

Ultra-Thin Microtoming of Kevlar® 49 Fibers

In this appendix the techniques for microtoming the Kevlar fibers embedded
in epoxy resins are discussed. For the benefit of others who may want to follow
these technique, some of the encountered problems and ways to get around them are

also mentioned.

Introduction

An essential part of the TEM microscopy is the ultrathin microtoming of its
samples. The TEM samples must be thin enough to transmit sufficient electrons to
form an) image. The samples must also be stable under the electron beam and in a
high vacuum. TEM images are formed by combination of elastic, inelastic, and
absorption interactions of the beam with the sample. Increasing the sample thickness
increases the beam absorptions and reduces the image resolution and the beam
stability. For a 100 kV electron beam the practical specimen thickness is limited to
100 nm.

108

109 Appendix D

Procedure
Kevlar fibers are embedded into epoxy dogbones as described in the dogbone

preparation procedure. The gauge length of a dogbone is cut into -;- inch sections.

The cuts are made with a razor blade, normally positioned and then tapped to
fracture the section. The two halves of the first cut are examined to see which one
has the pulled out fiber end, that half is chosen for the second cut. This inspection is
necessary to avoid having samples with no embedded fibers. A sample is then

placed into a sample holder and is clamped tightly. The sample is hand trimmed into
a trapezoid of dimensions no longer than 3:- mm. All the razor blades used in the

above process are new blades and their edges are cleaned with ethanol soaked cotton

swabs to avoid surface contamination.

The hand trimming of the trapezoid blocks is a very important step and needs
further elaboration. The sample holder is placed under a sectioning microscope and
the t0p face of the epoxy block is viewed at 40x magnification. The exposed pulled
out fiber on the face helps to locate the position of the fiber. Four large blade marks
of about 2 mm lengths are made to isolate the location of the fiber. With the fiber at

the center of the blade marks, the epoxy outside the marks are trimmed at 45°
angles. Next, two parallel cuts ~ 71- mm apart are made above and below the fiber.

It is very important to have these two cuts parallel in order to have a good ribbon
formation during the microtoming. Two more cuts at 45° to the two parallel cuts are

made in order to form the final trapezoid block. All of these four cuts are normal to
the face and about 1 mm deep. Finally, the top -;- mm of the trapezoid face is cut in

order to remove the excess fiber and have a clean undamaged face for the

microtomin g.

110 Appendix D

Ultra-thin microtoming of the Kevlar samples requires a diamond knife. A
55° angle knife with 3 mm edge width has been purchased (Micro Star Co.). The

right most % of the knife edge is dedicated to initial thick sectioning of the block.
The next 7‘;- of the knife edge is dedicated for the smoothing of the block face and
experimenting with a new face for its ribbon formation behavior. The next 33‘- of the

knife edge is dedicated for the final thin sectioning of the sample. The left % of the

knife edge has not yet been used and will be used gradually as the rest of the edge
becomes dull.

The sample holder is mounted into the microtoming arm and is fastened to
avoid vibration problems. Attention was paid to make sure the longest edge of the
sample trapezoid is facing down. The diamond knife is mounted at a 3° angle with
respect to the block face. To wet and clean the diamond edge, the tip of a wooden
applicator stick is shaved into a thin plate and is soaked in de—ionized water for at
least 15 minutes. The boat of the knife is filled with de-ionized water and the
diamond edge is wetted. The excess water is drawn off with a filter paper so that the
surface of the fluid is concave behind the diamond edge. The knife is advanced

manually toward the block face. Before the knife reaches the block, it is adjusted so
that the right most 3!;— of the edge would touch the block first. The knife is then

slowly advanced toward the block at 1 pm steps and block is manually moved. Once

the cutting begins, several 1 um sections are cut in order to prepare a clean smooth
face for the thin sectioning. The knife is then receded and is shifted % of the edge to

the left. Again the knife is slowly advanced until thin sections of gold or purple
color are cut. Motorized sample advancing is turned on at 0.45 mm/sec speed and
the thickness setting is adjusted to obtain gold colored sections. The ribbon

formation of the sample is examined. If the ribbons do not form or badly tilted to a

111 Appendix D

side or the sections have cut marks, these indicate that the block face is not ready for
the thin sectioning. If the ribbons do form but the sections are compressed and vary
in thickness (i.e. color), that indicates that the water level is too high and then the

water is drained until the problem is corrected. Once everything is ready for the thin
sectioning, the knife is shifted to the middle % of its edge and the gold color sections
are prepared at 0.33 mm/sec cutting speed.

Sections form ribbons which float on the water. To manipulate ribbons, an
eyelash applicator is prepared. An eyelash applicator is simply an eyelash mounted
on a wooden applicator using a drop of nail-polish. With the eyelash applicator, the
ribbons are assembled for the grid pick up. First the knife is receded and water is
added to the boat. The ribbon sections are then arranged away from the diamond
edge. The grids used are 200 mesh fine wire copper grids which have small end tabs
for easy pick up. Under the microscope of the microtomin g machine a grid is
submerged in water and is approached underneath the assembled ribbons. The
ribbons are picked up and the grid is drained on a filter paper. The grids are stored

in a dessicator or treated for sample staining and carbon coating.

Special Comments

The following statements summarize the important aspects of the described

microtoming method plus a few other trouble-shooting suggestions.

-- Cutting speed of 0.33 mm/sec produces compression free sections.

112 Appendix D

-- For untreated Kevlar samples with good epoxy adhesion, obtain gold coated
sections, however, for poorly bonded samples such as silicone coated Kevlar samples

obtain thicker gold-brown sections.

-- Use of 3° knife angle produces compression free and uniform color ribbons and is

recommended over the 2° and 4° knife angle settings.

-- Always maintain the minimum water level possible without losing the meniscus.
Sometimes it is hard to wet the diamond edge, to eliminate the problem, wet the

eyelash applicator with your thong and then brush it against the diamond edge.

-- Never breathe on the samples or use table light during the sectioning since they

alter block dimensions.

-- Fill up the boat with water before immersing the grid. Never pick up the ribbons

from top because this can crumble your sections.

-- 200 mesh fine wire grids are the best grids for our sample dimensions and provide
good support with large exposed sample areas.

References

10.

11.

12.

13.

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L2
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