Fatigue damage behaviors of carbon fiber-reinforced epoxy composites containing nanoclay

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Fatigue damage behaviors of carbon fiber-reinforced epoxy composites

containing nanoclay

Shafi Ullah Khan a, Arshad Munir b, Rizwan Hussain b, Jang-Kyo Kim a,⇑

a Department of Mechanical Engineering, Hong Kong University of Science and Technology Clear Water Bay Kowloon, Hong Kong b National Engineering and Scientific Commission, P.O. Box 2801, Islamabad, Pakistan

a r t i c l e i n f o

Article history:

Received 13 February 2010

Received in revised form 19 June 2010

Accepted 5 August 2010

Available online 11 August 2010

Keywords:

A. Nanoclay

A. Carbon fiber

B. Fatigue

B. Interfacial strength

C. Damage tolerance

a b s t r a c t

The effects of nanoclay inclusion on cyclic fatigue behavior and residual properties of carbon fiber-rein-

forced composites (CFRPs) after fatigue have been studied. The tension–tension cyclic fatigue tests are

conducted at various load levels to establish the S-N curve. The residual strength and modulus are mea-

sured at different stages of fatigue cycles. The scanning electron microscopy (SEM) and scanning acoustic

microscopy (SAM) are employed to characterize the underlying fatigue damage mechanisms and progres-

sive damage growth. The incorporation of nanoclay into CFRP composites not only improves the mechan-

ical properties of the composite in static loading, but also the fatigue life for a given cyclic load level and

the residual mechanical properties after a given period of cyclic fatigue. The corresponding fatigue dam-

age area is significantly reduced due to nanoclay. Nanoclay serves to suppress and delay delamination

damage growth and eventual failure by improving the fiber/matrix interfacial bond and through the for-

mation of nanoclay-induced dimples.

2010 Elsevier Ltd. All rights reserved.

1. Introduction

Carbon fiber-reinforced composites (CFRPs) are widely used as

structural material in load bearing applications because of high

strength and stiffness, dimensional and thermal stability, and cor-

rosion resistance. Fatigue is known to be oneof the primaryreasons

for failure in many structural materials, including CFRPs [1–3].

When subject to cyclic loading, CFRPs exhibit gradual degradation

of the mechanical and structural performance as a result of damage

accumulation. The nature of fatigue damage in CFRPs is very com-

plicated and is quite different from those of isotropic materials.

The damage states are closely related to the anisotropy and heter-

ogeneity which leads to the formation of different stress levels

depending on the lay-up sequence and orientation of laminate.

The fatigue damage modes in CFRPs include combinations of inter-

facial debonding, matrix cracking, delamination, fiber breakage, etc.

Early work on unidirectional CFRP laminates under tensile fatigue

loading displayed a high degree of resistance before sudden cata-

strophic failure [4]. However, when the matrix was more highly

loaded such as laminates with off-axis fiber orientations, the re-

sponse was completely different: there were multiple mechanisms

of failure throughout the material involving combinations of fiber

and matrix damage interaction. The fatigue behavior of on-axis

specimens was influenced by the stochastic breakage of brittle fiber

bundles, whereas that of off-axis angle-ply was strongly affected by

the inelastic shear deformation and crack propagation of the ductile

polymer matrix [5]. Similar conclusions were drawn in a recent

study where the failure modes were as much related to the cyclic

stress as to the off-axis angle [6]. For on-axis specimens the failure

modes were fiber-dominated and matrix-dominated when high

and low cyclic stresses, respectively, were applied. In sharp con-

trast, for off-axis specimens the failure mode was always matrix-

dominated irrespective of the stress level.

Thermosetting epoxy resin systems are widely employed as

matrix materials for composites in many fields such as aerospace,

automotive and microelectronics. Toughening of epoxies has been

one of the topics most extensively studied because of the brittle

nature of epoxies and their widespread applications for engineer-

ing components. Understanding the fatigue crack propagation

behaviors of epoxy composites has been of great importance be-

cause such composites are often used for engineering components

that are subject to cyclic loading. Curtis [7] found that the tough-

ened resin system can improve the tensile fatigue response in

the low cycle fatigue regime, while in the high-cycle fatigue range

the fatigue performance of the toughened epoxy is inferior to that

of standard epoxy-based composites. Epoxy matrices with a high

ductility exhibited a higher compressive fatigue resistance [8].

The mode I delamination fatigue crack growth was studied of

interlayer/interleaf-toughened CFRP laminates [9]. The heteroge-

neous interlayer with fine polyamide particles increased the crack

growth resistance.

0266-3538/$ - see front matter 2010 Elsevier Ltd. All rights reserved.doi:10.1016/j.compscitech.2010.08.004

⇑ Corresponding author. Tel.: +852 23587207; fax: +852 23581543.

E-mail address: [email protected] (J.-K. Kim).

Composites Science and Technology 70 (2010) 2077–2085

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Although many research efforts have been directed toward

understanding the mechanisms of fatigue in polymer matrix com-

posites, the effects of nanoparticles on their fatigue performance is

still not fully understood. The addition of 1 wt% of carbon nano-

tubes (CNTs) to the matrix of glass fiber-epoxy composite lami-

nates improved their high-cycle fatigue life by a remarkable 60–

250% [10]. Even more impressively, the addition of 2 and 5 wt%

multi-walled CNTs enhanced the fatigue performance of physio-

logically maintained methyl methacrylate–styrene copolymer

(MMA-co-sty) by 565 and 593%, respectively [11]. Zhang et al.

[3] demonstrated an order of magnitude reduction in fatigue crack

propagation rate for an epoxy system with the addition of 0.5 wt%

of CNTs. The crack-tip bridging and frictional pull-out mechanisms

were responsible for the suppression of fatigue in the nanocom-

posite. Other types of nanofillers also gave rise to improved frac-

ture properties. For example, the introduction of SiO2 particles

increased both the initiation fracture toughness and the corre-

sponding cyclic fatigue behavior of epoxy [12]. Al2O3 and TIO2

nanoparticles improved the flexural strength, stiffness and fracture

toughness as well as the fatigue crack propagation resistance of the

epoxy [13]. The incorporation of organoclay in polyurethane elas-

tomers showed significantly improved fatigue life in addition to

more than 150% increase in static strength and failure strain [14].

Many studies have devoted to improving the mechanical prop-

erties of fiber-reinforced composites by adding nanoclay. In addi-

tion to mechanical properties, clay-epoxy nanocomposites have

shown wide array of property improvements with only very low

fractions of clay, including the enhanced thermal stability

[15,16], reduced moisture and gas permittivity [17] and superior

flame retardancy [18]. The nanoclay, in particular, exhibited ame-

liorating effects on fracture and fatigue resistance of carbon fiber

composites (CFRPs): e.g. increased mode 1 delamination resistance

[19], enhanced impact damage resistance and tolerance [20] and

better static and impact fracture toughness [21]. However, very

few studies have appeared in the open literature on fatigue perfor-

mance of hybrid CFRP composites containing nanoclay. As a con-

tinuation of our previous studies on clay-CFRP hybrid composites[19–21], this work specifically studies the fatigue performance of

CFRP composites affected by the incorporation of nanoclay. The

S-N curves and the residual properties of hybrid composites after

tension–tension cyclic loads of different levels were specifically

evaluated.

2. Experiments

2.1. Materials and fabrication of composite laminates

The laminate composites were fabricated from unidirectional

carbon fiber and organoclay filled epoxy resin. The epoxy resin sys-

tem was basically the same as that employed in our previous stud-ies [19–21]: a diglycidyl ether of bisphenol A (DGEBA) epoxy

(Epon828, supplied by Shell Corp) mixed with 1,3-phenylenedi-

amine (supplied by Aldrich) hardener at a ratio of 100:14.5 by

weight. Unidirectional carbon fabric (supplied by Taiwan electrical

insulators) with a unit weight of 200 g/m2 was used as the main

reinforcement for composite laminates. The organoclay, Nanomer

I30P (supplied by Nanocor), is an octadeclyamine modified mont-

morillonite suitable for dispersion in epoxy resins [17]. The

organoclay was dried overnight at 75 C in an oven prior to use.

The epoxy in a glass beaker was heated at 75 C to lower the vis-

cosity and the organoclay was added. The organoclay content

was varied between 0, 3, and 5 wt% of the epoxy resin-hardener

mixture. Mixing was conducted at a shear rate of 3000 rpm for

1 h using a high speed shear mixer (Ross Mixer). The mixturewas subjected to sonication using an ultrasonicator (Branson

2510) at an ultrahigh frequency for 3 h to further disperse the clay,

while maintaining the resin temperature at 75 C using a hot water

bath. After sonication, the translucent color of the epoxy/clay mix-

ture indicates uniform distribution of organoclays, partly confirm-

ing the efficiency of the sonication conditions used. The mixture

was degassed in a vacuum oven followed by addition of curing

agent, and the mixture was stirred while avoiding the formation

of bubbles. Twelve ply laminates of 30 cm square were prepared

by hand lay-up of carbon fabrics with a stacking sequence [0/

90]3S on a steel mould plate. To keep fabrics well aligned, necessary

precautions were taken during hand lay-up. The molded laminates

were wrapped with bleeders and peel plies within Teflon dam all

around, which was cured at 80 C for 2 h and at 150 C for 8 h, fol-

lowed by post-cure at 160 C for 2 h in a vacuum hot press (Tech-

nical Machine Product Corp). The high cure temperature

excursions for long durations were aimed at complete cure of the

resin. The cured composite laminates were cut, by a diamond

wheel, at 45 off-axis directions to obtain a resultant stacking se-

quence of [±45]3S. Introduction of clay into epoxy inevitably in-

creases the viscosity of the resin, which may result in composite

laminates thicker than those without clay. To lower the viscosity

and thus to avoid the thickness variation, the resin was heated to

75 C during the whole processing steps, including shear mixing,

sonication and degassing, as well as before hand lay-up after mix-

ing with the hardener. A uniform laminate thickness and a con-

stant fiber volume fraction were further assured through the use

of Teflon dams of required thickness and a constant pressure of

0.32 MPa during curing. The volume fraction of carbon fibers, V f ,

was consistently maintained at about 0.55 for both the composites

with and without nanoclay, which was determined from the

known weights and densities of the composite constituents.

2.2. Characterization, static and cyclic fatigue tests

The static tensile tests were conducted according to the specifi-

cation, ASTM D3039, on a universal testing machine (MTS Sintex

10/D) to determine the tensile strength and modulus. Rectangularspecimens of 230 mm long 20 mm wide 2.5 mm thick were

loaded at a crosshead speed of 2 mm/min. An extensometer with

gauge length of 25 mm was attached to the specimen to monitor

the strain during loading.

The tension–tension cyclic fatigue tests were conducted accord-

ing to the specification, ASTM D3479, on a universal testing ma-

chine (25 KN servo-hydraulic Instron 1300). The tests were

conducted at room temperature on a load control mode at a stress

ratio of 0.1, and with constant-amplitude sine-wave loading. To

determine the fatigue S-N curves, the maximum stress levels were

kept at 80, 70, 60 and 45% of the corresponding ultimate tensile

strength (UTS) of the composite. A test frequency of 2 Hz was used

which was low enough to minimize the effect of adiabatic heating.

Rectangular specimens, 230 mm long, 20 mm wide and 2.5 mmthick, were cut from the composite plates, and end tabs made of

glass fabrics and 40 mm long were bonded at both ends of the

specimen to avoid failure around the gripping device during the

tests. At least four specimens were tested for each set of loading

conditions. The residual properties of the composites were mea-

sured after different periods of fatigue loading at a maximum load

equivalent to 60% of the ultimate tensile strength of the composite.

Static tensile tests were conducted on the pre-cycled specimens to

measure the residual tensile strength and modulus. The tests fol-

lowed the same procedure as those for the static tensile properties

on virgin specimens.

The scanning electron microscopy (SEM) was used to examine

the surface morphologies of the static and fatigue fractured speci-

mens and thus to identify the different failure mechanisms in-volved in CFRPs with and without nanoclay. The scanning

2078 S.U. Khan et al. / Composites Science and Technology 70 (2010) 2077–2085



acoustic microscopy (SAM, Sonix Micro-Scan System) was em-

ployed to characterize the progressive fatigue damage growth at

different stages of fatigue cycles. A focused acoustic beam was

scanned over the damaged laminate using a transducer equipped

with a 35 MHz probe in a through-transmission mode. Both the

neat CFRP composites and the hybrid composites containing

5 wt% nanoclay were examined before loading and after 5 K,

10 K, 20 K, 25 K and 30 K fatigue cycles. For ease of understanding

only two colors, black and grey, were used to present the damage

state and a threshold value of 10% was used as the border line be-

tween the two colors.

3. Results and discussion

3.1. Static tensile properties

Fig. 1 presents a typical TEM image of nanocomposites with

5 wt% clay content, indicating a mixture of full intercalation and

partial exfoliation. Representative stress–strain curves obtained

fromthe static tensile tests are shown in Fig. 2. All materials exhib-

ited a typical bilinear stress–strain behavior before failure. As re-

ported previously [22–24], the tensile stress–strain curves for

angle-ply specimens are non-linear due to the significant contribu-

tion of the polymer matrix. It is clearly seen that both the yield

strength and the failure strain increased with increasing the clay

content. Fig. 3 summarizes the static tensile strength and modulusof clay-CFRP hybrid composites containing varying clay contents.

Both the tensile strength and modulus increase continuously with

increasing clay content, which is again a reflection of the compos-

ite property significantly affected by the matrix property. This

observation is generally consistent with the flexural properties re-

ported earlier [21] although the flexural strength tended to be mar-

ginally reduced at a high clay content due to the potential lack of

dispersion of clay.

The fracture surface morphologies as shown in Fig. 4 exhibited

sharp contrast between the composites without and with 5 wt%

nanoclay. Interfacial debonding between the fiber and matrix, as

well as limited deformation of matrix material are the major fail-

ure mechanisms observed in the composites without clay. The

fracture surface was generally smooth and featureless indicatingbrittle failure. Meanwhile, the clay-CFRP hybrid composites re-

vealed improved fiber–matrix interfacial bonding due to the pres-

ence of nanoclay in the matrix material that maximizes the stress

transfer between matrix and fiber. The modified epoxy adhered

well to the long carbon fibers and the fracture surface was rougher

and textured, quite similar to those observed from the interlaminar

fracture surfaces [19]. It is thought that the octadecylamine modi-

fier used for I.30P organoclay had alkyl and amine groups that are

functionally compatible with carbon fibers to give rise to strong

adhesion [20]. Similar amine groups have been extensively used

to functionalize carbon nanotubes/nanofibers for polymer compos-

Fig. 1. Typical TEM image of nanocomposite containing 5 wt.% clay, showing

dispersion state of nanoclay.

Fig. 2. Representative stress–strain curves of clay-CFRP hybrid composites.

Fig. 3. Tensile properties of clay-CFRP hybrid composites containing varying clay contents.

S.U. Khan et al. / Composites Science and Technology 70 (2010) 2077–2085 2079



ite applications, which may also be responsible for the improved

adhesion between the modified epoxy and ultra-high molecular

weight polyethylene fibers [25]. It is well known that the proper-

ties of the composites with [±45]S ply orientation are dominated

by the in-plane shear properties, which is further confirmed by

the general view of the failed specimen shown in Fig. 5a. The inter-

laminar shear and in-plane shear are considered to be the matrixdominated properties [26]. It was noted (Fig. 2) that the presence

of nanoclay in the matrix not only increased the apparent yield

stress but also the strain to failure, which is consistent with the

previous observations of improved interlaminar shear strength

(ILSS) [19,27] and in-plane shear strength of fiber composites [28].

3.2. Fatigue life and residual strength

Fig. 6 presents the S-N data of clay-CFRP hybrid composite at

varying clay contents. For the same level of maximum applied

stress, the clay-CFRP hybrid composite exhibited much longer fati-

gue life than the composite without clay at all the stress levels

tested. A maximum improvement of about 74% in fatigue lifewas achieved with 3 wt% clay when cyclic fatigue was carried

out at a load equivalent to 45% of the tensile strength of the spec-

imen. Results indicated that nanoclay modification produced more

improvement in fatigue life at a low stress or a high cycle regime

and was less potent in improving life at high stress levels. At high

Fig. 4. Fracture surface morphologies of CFRP composites: (a and b) without clay and (c and d) with 5 wt.% clay.

Fig. 5. Images of the failed specimens from (a) static tension and (b) cyclic fatigue.




stress levels the nanoparticles are less resistive in suppressing the

rapidly propagating cracks. This was explained on the basis of

stress intensities or strain densities [10,11]. At high stress levels

the fatigue crack grow at a rapid rate and at several fronts becauseof the high stress intensity and high strain density, respectively.

The nanoparticles become relatively ineffective in slowing down

the onset or subsequent growth of damage at high stress

intensities.

Fig. 7 shows the residual tensile strength and modulus mea-

sured after a given number of cycles. The residual properties exhib-

ited a gradual decrease with increase in number of cycles with

small variations between specimens. The nanoclay-modified com-

posite showed higher residual properties than those without

throughout the whole fatigue life. Judging from the typical images

of failed specimens in Fig. 5b, in-plane shear, instead of fiber break-

age, is the primary failure mode. Since no stochastic fiber failure

was involved, erratic changes in the residual strength and modulus

were not observed, further details of the latter will be discussed in

Section 3.4.

3.3. Fatigue damage index

Amongst several methods to characterize the extent of damage

arising from cyclic fatigue, the ‘fatigue damage index’ was em-

ployed in this study. The fatigue damage index, D, is defined as:

D ¼ 1 E r

E oð1Þ

where E is the modulus of specimen, with the subscripts o and r

referring to the undamaged state and residual value after a certain

fatigue life. D varies between 0 and I, and a low D value means little

modulus reduction due to fatigue. Thus, D is a macroscopic measure

of fatigue damage because the structural changes on the micro-

scopic scale (due to matrix cracks, fiber/matrix interfacial failure,

etc.) are characterized by a macroscopic reduction of the modulus[29–32]. Fig. 8 shows the damage index, D, plotted a function of fa-

tigue cycles for hybrid composites containing different clay con-

tents. It can be seen that at the early stage of fatigue (say, 0–

12.5 k cycles) the hybrid composites in general exhibited margin-

ally more damage than the neat composites. After the initial dam-

age period, the hybrid composite specimens sustained a relatively

longer stable period with low damage indices for the rest of fatigue

life. The final failure took place much earlier in the neat composite

than the clay-CFRP hybrid composites; and the higher was the clay

content, the longer was the fatigue life, with the exception of the

hybrid composites containing 5 wt% clay. The diminishing improve-

ment in fatigue life at clay contents higher than 3 wt% is attributed

to the higher possibility of forming unwanted agglomerates of a rel-

atively large size.

The early stage is generally considered as the crack initiation

stage. The hybrid composites have a large number of interfaces

due to the presence of nanoclay, and there are many weak inter-

faces between the clay galleries. The cleavage of clay tactoids or

the inter-gallery debonding might have occurred during this early

stage of fatigue generating micro or nanoscale cracks. It is likely,

however, that these micro or nanocracks took significantly longer

time to coalesce and propagate to form critical damage than in

the neat CFRP composites. Other important toughening mecha-

nisms responsible for the enhanced crack growth resistance of-

fered by nanoclay have been identified previously [18–

21,27,33,34]. Nanoclay can serve as the trigger for crack deflection,

Fig. 6. S-N Curves of clay-CFRP hybrid composites with varying clay contents.

Fig. 7. Residual fatigue properties: (a) strength and (b) modulus. Fig. 8. Fatigue damage variable, D, plotted as function of fatigue life.




pinning, as well as crack arrest mechanisms, which significantly in-

crease the fracture toughness. These mechanisms could also be

operative to retard fatigue failure by preventing or slowing the

damage buildup through fatigue crack growth, especially at the la-

ter stage of fatigue.

3.4. Fatigue damage mechanisms

Fig. 9 shows SEM micrographs of fractured specimens that were

fatigued at 80%, 60% and 45% of UTS of the CFRP composites with-

out and with 5 wt% clay. Both composites showed a number of

Fig. 9. SEM images showing fatigue fracture surfaces of the specimens failed at different maximum stress levels: (a) neat CFRP composites and (b) 5 wt.% clay-CFRP hybrid

composites.

Fig. 10. SEM micrographs of 5 wt.% clay-CFRP hybrid composite, showing improved fiber–matrix and particle–matrix interfacial bonding.




dimples on the fracture surface. The dimples are formed as results

of highly localized deformation around matrix defects, microvoids

or nanoparticles [35,36]. It is interesting to note that the size of

these dimples in general increased in both width and depth with

decreasing the stress level applied during fatigue or alternatively

with an increase in fatigue life. The mechanisms behind the forma-

tion and growth of dimples can be explained as follows: microcrac-

kes are formed during the early stage of fatigue loading at some

weak locations where there are high stress concentrations. Local-

ized plastic deformation occurs due to the strain concentration

around these microcracks, which in turn promotes the formation

of dimples. The lower the applied stress level, the longer this stage

lasts. Dimples grow in a stable fashion with increasing number of

fatigue cycles before they coalescence into major cracks and final

failure. At high stress levels, they tend to coalescence rapidly, lead-

ing to premature fracture without extensive growth of deep and

wide dimples. These observations suggest that the elongation or

stretching of the localized deformation zones or dimples is highly

dependent on the applied stress.

Indeed, the morphology of the nanoclay-induced dimples is

considerably different from those present in neat composite. The

neat composites had generally wider and shallower dimples than

the nanoclay hybrid composites. There was significant difference

in size of these dimples for a given maximum stress level, espe-

cially at a low stress or high cycle regime: e.g. at 45% of UTS level

these cavities are about 4–6 and 2–3 lm wide in the neat and hy-

brid composites, respectively. For a given fractured cross-sectional

area, there were more dimples of smaller sizes in the hybrid sys-

Fig. 11. SAM images showing progressive fatigue damage growth at different stages of fatigue life.




tem than in the neat system. Similar dimple patterns were also

found near the pre-cracked region of fractured PA/nano-TIO2 and

PA/nano-SIO2 composites [35,36] where the density of dimples

was much higher and the size was much smaller in the nanocom-

posites than in the neat PA66. In addition, there is significant anal-

ogy between our findings and those reported previously on glass

fiber composites (GFRPs) containing carbon nanotubes (CNTs)

[10] in terms of crack density: the addition of CNTs into the com-

posites resulted in a higher density of nanoscale cracks than the

composites without CNTs.

It is also worth noting that the nanoclay-induced dimples are

often divided into sub-partitions as evidenced from Figs. 9b and

10. There appears to be distinct partitioning lines emerging from

clay particles in the center to the periphery of the dimples. These

dividing walls significantly increased the density of dimples and

the total fracture surface area. The increased surface area implies

meandering crack tips during crack growth, consequently resulting

in high energy dissipation. The increased number of distinct parti-

tioning lines in hybrid composites containing nanoclay is again an

indication of simultaneous initiation and growth of small cracks in

a large density, in contrast to larger cracks in a low density in the

neat composite. Nevertheless, the static fracture surfaces in both

the hybrid and neat composites did not show any sign of these

dimples. The rapid crack propagation at a high stress intensity

and a high loading speed did not allow the matrix material with

enough time to stretch and plastically deform into these dimples.

The clay/matrix interfacial bond appears to be very strong even

when the clay was agglomerated to a microscale. The strong inter-

facial bond gave rise to magnification of the strain that the matrix

material between the nanoparticles could sustain.

3.5. Characterization of fatigue damage growth

The SAM technique was employed in a through-transmission

mode [37–39] to monitor the progressive damage growth due to

cyclic fatigue and the representative SAM images are shown inFig. 11. The corresponding damage area divided by the total area

within the gauge length is plotted as a function of life cycles in

Fig. 12. Black color (Fig. 11) represents the sound waves that are

less than the threshold value and are absorbed by the composite,

thus is an indication of damage. Grey color represents the sound

waves that pass through the composites and are received at the re-

ceiver unit, and thus is an indication of bonded region. The slight

variation in the grey color intensity occurred because of the den-

sity variation of fiber, matrix and the particle rich areas, which in

turn changed the extent of received waves.

Both materials exhibited quite uniform damage distributions

over the whole gauge area until damage became more concen-

trated at the later stage. It is seen that the hybrid composites

exhibited slightly more damage than the neat composites at the

very early stage of loading at below about 10 k cycles. This obser-

vation is functionally very similar to the damage index variation

discussed in Section 3.3 (Fig. 8). Except the early stage, the clay-

CFRP hybrid composite showed in general much less damage than

the neat composite throughout the whole fatigue life. The clay-

CFRP hybrid composites sustained about 15% more damage before

final failure than the neat composites, confirming substantially

higher damage-tolerant characteristics. Based on the above obser-

vations, the fatigue damage in the laminate composites studied

here can be divided into two stages: Stage I for damage initiation

and stable damage growth; and Stage II for rapid damage growth

to failure. It is seen that the hybrid composites had a longer Stage

I period than the neat composites (i.e. approximately 0–20 k cycles

vs 0–15 k cycles).

4. Conclusions

The tension–tension fatigue behavior of clay-CFRP hybrid com-posite was investigated and the prevalent toughening mechanisms

arising from nanoclay were identified. The following conclusions

can be highlighted from this study.

(1) The static tensile strength and modulus of CFRP composite

were significantly enhanced by the addition of nanoclay.

The presence of nanoclay in the matrix increased the appar-

ent yield strength and strain to failure.

(2) The clay-CFRP hybrid composites showed better perfor-

mance in terms of residual tensile strength and modulus

than the neat composite after a given fatigue cycle.

(3) Fatigue life was significantly extended with the incorpora-

tion of nanoclay to CFRP composite and the maximum

improvement was about 74% with 3 wt% clay content.(4) Nanoclay suppressed the fatigue damage growth of CFRP

composites in terms of damage area over the whole fatigue

life except the very early stage of loading.

(5) Improved fiber/matrix interfacial bond and nanoclay-

induced dimples were identified as underlying toughening

mechanisms responsible for the enhanced fatigue life of

clay-CFRP hybrid composites.

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Fatigue damage behaviors of carbon fiber-reinforced epoxy composites containing nanoclay

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