mode i delamination behaviour of short fibre reinforced carbon fibre/epoxy composites following...

Mode I delamination behaviour of short fibre reinforced carbonfibre/epoxy composites following environmental conditioning

Laurence Walker*, Xiao Zhi Hu

Department of Mechanical and Materials Engineering, University of Western Australia,

35 Stirling Highway, Nedlands, Western Australia 6009, Australia

Received 11 June 2001; accepted 19 March 2002

Abstract

Carbon fibre laminates based on Newport Adhesives NCT-301 uniaxial pre preg were prepared for Mode I DCB testing. Poly-amide and Polyolefin web materials supplied by Spunfab Ltd and Toyobos’ Zylon HM fibres were applied to the interlaminar

region to enhance fracture toughness. Three of each specimen type were ‘‘treated’’, subjected to immersion in 95 �C water for aperiod of 600 h. Two ‘‘untreated’’ control specimens of each type were aged at 21 �C and 40-95% relative humidity. Specimenweight gains in the order of 2% were found to occur in all immersed samples. Fracture toughness increases were evidenced in alltreated samples with the exception of the Polyamide Web sample. Average GIc increases in treated samples were 250–300%.

Untreated samples were found to have exclusively interlaminar crack propagation. Untreated samples containing Polyamide orPolyolefin Web reinforcement displayed increased fracture toughnesses. These toughness increases were attributed to improvedinterlaminar bonding.

# 2002 Published by Elsevier Science Ltd.

Keywords: A. Carbon fibres; B. Environmental degradation; B. Fracture toughness; C. Delamination; D. Scanning electron microscopy (SEM)

1. Introduction

The use of Short Fibre Interlaminar Reinforcement(SFR) is a simple and low cost method for the tough-ening of laminar composites. The technique uses a verylow density of tough fibres as reinforcement within theinterlaminar region. SFR systems incorporating Kevlaror Zylon fibres have been shown to increase the quasi-static toughness of double cantilever beam (DCB) spe-cimens [1,2]. Development work with these fibres andalso Spunfabs Polyamide Web system have reduced thearea and intensity of damage in low velocity impactspecimens [3]. In addition, the SFR process has littleeffect on the existing laminating processes and can beincorporated into current pre preg manufacture withlittle effect on production costs.The aerospace industry pays particular attention to high

temperature/humidity environments due to their deleter-ious effects on material properties. High temperature

environments, in excess of a systems current Tg,decreases the mechanical properties of the matrix forshort exposure periods. Long term exposure to elevatedtemperatures however usually enhances crosslinking viaa post cure mechanism [4]. This enhanced crosslinkingincreases laminate stiffness and consequently reducesfracture toughness. The effects of long term exposure inhigh humidity environments however can be con-siderably more damaging. The absorption of water inhot/wet environments is found in most polymeric com-posites. The degree of absorption is both matrix andfibre dependent with many properties decreasing sub-stantially with increasing moisture levels. Early work inthe field suggested that property reductions were linkedto moisture content. Recent work however has deter-mined that the exposure cycle is the determinant ofmechanical property reductions [5]. Outcomes arisingfrom moisture ingress include reduced matrix Tg andreduced bonding at the matrix/fibre interface [6]. Akay[7] observed reductions in interlaminar shear, compres-sive and flexural strengths of 17, 10 and 7% in uni-directional carbon fibre/epoxy laminates associated witha 1.5% by mass laminate moisture gain. These results

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Composites Science and Technology 63 (2003) 531–537

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* Corresponding author. Tel.: +61-8-9339-0111; fax: +61-8-9438-

1203.

E-mail address: [email protected] (L. Walker).

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were obtained by immersing T300 3K/Cycom 985 epoxysamples in water for 21 days at 70 �C.The range of test conditions available for environ-

mental exposure testing is considerable. Birger et al. [8]found the absorption rate for laminates immersed at100 �C was approximately 5 times faster than that oflaminates exposed to 50 �C/95% RH. Laminates agedfor short periods at 100 �C produced substantiallyreduced strengths that decreased further with time.Glass transition temperatures were observed to decreasefrom 135 to 95 �C after 24 h immersion at 100 �C. SEMimages taken from Birger et al. [8] show that even shortexposures to boiling water can destroy the fibre/matrixadhesion and create large voids within the matrix. Theselarge voids were attributed to the dissolution of theuncured DICY curing agent at elevated temperaturesand humidities.The work of Biro et al. [9] utilised the microbond

technique. Following debonding, the frictional coeffi-cient of pullout was found to increase by at least 85% inboth AS4 and T300 fibres. Resin Tg was found todecrease by approximately 11 �C with moisture weightgains of 1.54–1.97%. The reduced fibre/matrix bondingwas attributed to a swelling of the epoxy due to waterabsorption reducing the radial pressure on the fibre andweakening interfacial interactions.The use of SFR materials for interlaminar toughening

relies on combining tough fibres with carbon/epoxypre preg. It has been noted previously that Kevlar hasa tendency to absorb larger quantities of water thancarbon and glass fibres in high humidity/high tem-perature environments. Further, it has been found thatKevlar fibres provide a faster path for the transfer ofmoisture than purely matrix diffusion [10]. Environ-mental conditioning is therefore necessary to examinethe effects of hot/wet environments on SFR modifiedlaminates.

2. Materials and methods

A series of 210 mm � 105 mm test panels containingthree SFR materials as listed in Table 1 were manu-factured. High modulus Poly (p-phenylene bensobisox-azole) (PBO HM) fibres were supplied by Toyobo Co

under the trade name Zylon for use in the study. SpunfabLtd supplied Polyamide (12 g/m2) and Polyolefin Web(10, 30 g/m2) materials for use as interlaminar reinfor-cements. Specimens utilised Newport AdhesivesNCT-301 300 g/m2 uniaxial pre preg tape. Panels werecomprised of 12 plies at a nominal ply thickness of 0.2mm/ply. The mid plane of each panel contained SFRfibres and an initial DCB starter crack was generatedthrough the addition of a folded piece of aluminium filmas suggested by Davies et al. [11]. Specimens of 18 mmwidth were cut from the panel, ground to a tolerance of0.05 mm and polished to 1200 grit.Immersion tests consisted of 600 h immersion at 95 �C.

Three samples were environmentally treated whilst twosamples were aged in air at 22 �C with humidity varyingfrom 40 to 90%. Following immersion, all specimenswere placed in a 90 �C oven for 2 h to remove freemoisture. Specimen sides were cleaned of calciumdeposits to remove the possibility of mass related errorsassociated with the deposition of water contaminants.Cleaning was performed with mild abrasion from ascotch pad and water.Each specimen was analytically weighed prior to and

at the completion of the drying period to determine thepercentage of water uptake. Immersed specimens arereferred to as treated whilst room aged specimens areuntreated.For use in DCB testing, specimen sides were painted

with white correction fluid to facilitate easy crack iden-tification. Specimens were tested using the establishedloading/unloading technique with measurements readfrom the laminate with the aid of a magnifying glass andvernier calipers. Testing took place on an Instron 4301at 2 mm/min cross head speed and a Lloyds PL3 chartrecorder.As this test programme was purely comparative and

there is little agreement for the use of correcting equa-tions in fibre bridged DCB specimens, this testing uti-lised an uncorrected form of DCB equations. Thisuncorrected form has been used in previous papersexamining SFR laminates due to its applicability tospecimens demonstrating fibre bridging.The uncorrected form was developed as follows:

1. Values of Load (P), Deflection (�) and crackextension (a) were taken at intervals determinedby the generation of crack propagation. Theseintervals were approximately 5 mm however datawas only recorded if the load was decreasing.

2. Initial compliance values were calculated for theelastic load displacement curves.

3. Youngs modulus was calculated using Eq. (1)and the compliance determined in step 2. Thecompliance crack extension was taken to be theinitial crack extension, as determined during thefracture surface examination.

Table 1

Specimen composition and mass changes due to hot/wet conditioning

SFR Length Density

(%)

Mass increase g (%) by specimen

1 2 3

POW4 As supplied 4 0.35 (2.11) 0.34 (2.06) 0.34 (2.06)

None N/A N/A 0.35 (2.12) 0.36 (2.16) 0.37 (2.19)

PAW As supplied 4 0.35 (2.09) 0.32 (1.89) 0.35 (2.06)

ZAS 6 mm 1 0.41 (2.39) 0.39 (2.30) 0.33 (1.96)

532 L. Walker, X. Zhi Hu /Composites Science and Technology 63 (2003) 531–537

4. Youngs modulus was averaged across the fivesamples of each SFR type.

5. Subsequent values of P and � were coupled withthe value of a generated by the previous unstablecrack extension and substituted into the uncor-rected Eq. (2).

�

P¼8a3

BEh3ð1Þ

G1c ¼P2a2

BEavgIð2Þ

Fracture analysis involved the use of two methods;visual observation and scanning electron microscopy(SEM). A visual analysis of the fracture surface is a cri-tical part of the work required when determining SFRtoughening mechanisms, as the presence of short fibresand their distribution are clearly visible. The globalmechanisms identifiable via surface analysis can be fur-ther analysed using SEM to determine the processesinvolved in originating and/or developing improvedtoughness. For the SEM study, 10 mm long, full widthsamples were cut from both treated and untreated spe-cimens of each SFR type. Specimens were gold andcarbon coated and analysed using a Phillips 505 SEM at50 nm and 15 kV excitation voltage.

3. Results and observations

3.1. Quantitative assessment

Examining Table 1 it can be seen that specimens dis-played consistent mass increases after immersion.Increases ranged from 0.32 g (1.89%) to 0.41 g (2.39%)with a mean of 2.11% and a standard deviation of0.04%.Figs. 1–4 display the fracture toughness of specimens

featuring varying interlaminar reinforcements. Solidsymbols represent the environmentally treated samples

whilst outlined symbols were room aged. Most of thesamples shown in the figures display relatively con-sistent fracture toughness curves. Both treated anduntreated samples display trends relative to each SFRtype. It must be noted that the scales for Figs. 2 and 4differ from those of Figs. 1 and 3.

Fig. 1. Fracture toughness of environmentally treated (solid symbols)

and untreated samples featuring unreinforced interlaminar regions.


and untreated samples featuring interlaminar regions reinforced by

Polyamide web with a density of 4% by mass.


and untreated samples featuring interlaminar regions reinforced by

Zylon HM with a density of 1% by mass.


and untreated samples featuring Polyolefin web with a density of 4%

by mass.

L. Walker, X. Zhi Hu /Composites Science and Technology 63 (2003) 531–537 533

The high modulus Zylon (ZHM) sample of Fig. 3shows a clear delineation between the treated anduntreated samples. Toughness values for treated sam-ples are approximately 200% higher than those foruntreated samples. Both curves are consistent andincrease with increasing crack length. Similar results,consistent outcomes, delineated behaviour and increas-ing toughness with crack length, are shown by theunreinforced specimen.The effect of environmental treatment on polyolefin

web samples is less pronounced than on SFR samples.The 4% PO Web sample shows relatively linear tough-ness curves, increasing with increasing crack length. Thetwo untreated samples display widely varying initiationtoughnesses and different toughness slopes. The treatedsamples show similar results for all three samples andhigher toughness values than the untreated samples.The PA Web specimen displays an opposite form of

behaviour to all of the other specimens. Treated samplesproduce relatively poor toughnesses, higher than mostuntreated samples but lower than all of the treated spe-cimens. The initiation toughness of the untreated sam-ples is very high relative to all of the other samples.Crack extension in these samples tends to be very highfor each load cycle. Scatter for the untreated samplesalso appeared very high.This scatter may be attributable to the limited number

of data points that could be obtained for each samplemaking continuity between points difficult to define.Testing showed that crack propagation occurred in anunstable manner. In one case the crack extendingapproximately 30 mm from its initiation point in a sin-gle cycle.

4. Fracture analysis

Fracture surface comparison is the most obvious wayof determining the global mechanics behind laminateperformance differences. Images of visual fracture sur-faces are very difficult to display and interpret due toproblems associated with displaying 3 dimensional var-iations in a 2 dimensional image. SEM images can moreaccurately portray the global outcomes due to theirresolution at varying depths. SEM images also detaillocal fracture mechanisms that can be linked to globalfracture mechanics. For example, the image of Fig. 7displays relatively flat topography and SFR can beobserved in the surface. This translates to a globalinterpretation that the fracture remained interlaminarrather than travelling to the intra-ply region. SEMimages shown below were chosen to be indicative of thefracture surface of the whole sample and therefore ofthe global mechanisms involved in each type of fracture.A comparison between Figs. 5 and 6 shows the effect

of environmental treatment. Fig. 5 shows continuous

carbon fibres covered in a thin layer of epoxy resin.Hackle marks are evident on the surface and none of thecontinuous fibres have been dislocated from the ply.

Fig. 5. SEM image showing the fracture surface of an unreinforced

and untreated DCB sample.

Fig. 6. SEM image showing the fracture surface of an unreinforced

and treated DCB sample.

Fig. 7. SEM image showing the fracture surface of a untreated sample

using Polyolefin web as an interlaminar reinforcement.


Conversely, Fig. 6 displays substantial quantities ofdislocated carbon fibres often bundled into smallgroups. Shards of matrix can be seen throughout thespecimen and substantial quantities of matrix areattached to the fibres. The topography is highly variedand large quantities of matrix damage and voidage arein evidence.The images of Figs. 7 and 8 show the difference

between the fracture surfaces of polyolefin web with andwithout treatment. Central to the image of Fig. 7 is anarea of polyolefin web ‘‘fibre’’, passing under the con-tinuous carbon fibres on the left and continuing off theimage to the right. The polyolefin web is originallycomprised of a thin sheet of polyolefin fibres similar inappearance to a low-density non-woven veil. In thisimage the fibre has melted during processing to form awidened area of adhesive. The fibre has maintained itsexternal shell however this has fractured leaving acylindrical valley in the matrix material. Around theweb, the matrix has fractured in the interlaminar regionwith little fibre fracture or interfacial damage. Fig. 8displays very similar fracture mechanisms to the imagesof Fig. 6. There are a substantial number of dislocatedcontinuous fibres, matrix shards and matrix attached tocontinuous fibres, topography is varied and voidage isapparent. The large object to the right of the image isPO Web material. Similar to the PO Web fibre of Fig. 7,the difference however arises in the continuity of theweb material and its attachment to the matrix materialbelow it. Had the background been consistent andunbroken as in Fig. 7 their appearance would be verysimilar. As the matrix has fractured, the web materialhas been separated by the matrix or obscured as thecrack has passed into an intra-ply plane.The treated ZHM sample of Fig. 9 shows limited evi-

dence of ZHM fibres. The surface is comprised pri-marily of exposed continuous carbon fibres withevidence of pulled out groups or bundles of carbon

fibres. Most carbon bundles have matrix attached to thesurface and shards of matrix are seen across the image.Darker areas on the image are evidence of topo-graphical low points. These are areas where carbonbundles have been separated from the plies. A couple offibrillated fibres can be seen at the top and middle left ofthe picture.Fig. 10 is a treated sample of Polyamide Web. The

figure displays exposed carbon fibres, shards of matrixmaterial, matrix material attached to separated fibrebundles and evidence of bundle pullout. Central to theimage of Fig. 10 is the plastic fracture appearance ofPolyamide Web. A couple of fractured carbon fibres arein evidence on the specimen. The web material has theimprint of continuous fibres on its surface. It appearsthat the fibres have separated from the material fairlyeasily as the fibres and the web are both fairly clean withno remnants of either material found on the other. Notobvious from the SEM image but readily seen on thefracture surface is the white appearance of the PAW

Fig. 10. SEM image showing the fracture surface of a treated sample

using PA web as an interlaminar reinforcement.


using Polyolefin web as an interlaminar reinforcement.


using Zylon HM as an interlaminar reinforcement.


after treatment. This is in contrast to a dark appearanceseen on untreated PAW specimens. The untreatedPolyamide Web system displayed unique behaviour inthat the surface consisted entirely of a resin rich inter-laminar region. Limited surface hackling was in evi-dence and the continuous fibres were completelycovered by a thin layer of PA Web and matrix materiali.e. no exposed fibres.SEM images have not been included for the untreated

ZHM and PA Web systems due to space restrictions.Both of these systems showed similar fracture surfacesto some of those described above.The surface of the untreated ZHM sample is most

readily represented by the top left corner of Fig. 5 withthe exception of an occasional visible ZHM fibre. Itdisplays considerable quantities of continuous fibrescovered in a thin but continuous coating of epoxy resin.Some carbon fibres have been dislocated from thematrix and occasionally a fibrillated Zylon fibre can beseen. There is limited SFR in evidence and limitedfibrillation. The untreated Polyamide Web system dis-played unique behaviour in that the surface consistedentirely of a resin rich interlaminar region. A visibledarkening due to the addition of the PA Web materialwas seen across the entire surface. Limited surfacehackling was in evidence and the continuous fibres werecompletely covered by a thin layer of PA Web andmatrix material i.e. no exposed fibres.

5. Discussion

Mass increases occurred across all of the samples tes-ted. This is indicative of water penetration into thecomposite material. The largest area of the materialexposed to the water was the top and bottom of thespecimens. This area was consistent across all samplesand therefore should reflect fairly similar moisturegains. The second largest surface area was the inter-laminar region associated with the initiation crack.Moisture gain within this area is related to interlaminarmaterial, however it must be noted that only one surfaceof the initiation crack was covered in short fibre rein-forcement. This limited difference in area between spe-cimens likely accounts for the similarity in masspercentage increases. No specimen displayed significantvariation in mass gain to warrant further investigation.The propensity for increased Mode I fracture tough-

ness with environmental treatment carries across almostall of the specimens tested. The only specimen notexperiencing toughening was that of the PAW specimen.Polyamide hydrolyses in water to form an ammoniumsalt [12]. The combination of high temperature, longduration and a large surface area relative to materialthickness lead to substantial degradation of the PAW.The PAW material covering the whole specimen surface

had hydrolysed from a dark translucent film to a whitesalt like deposit. This transition is evidence of the diffu-sion of moisture throughout the entire specimen.Toughness results show the treated PAW specimenaveraging 0.8 kJ/m2 whilst the untreated samples aver-aged around 0.5 kJ/m2. These figures, coupled with thepropensity for interlaminar cracking suggest thatdespite the material degrading extensively, the outcomeremained tougher than the untreated samples for thisduration and intensity of exposure. The performance ofthe untreated PAW sample demonstrates the hightoughening potential of the material as an interlaminarreinforcement adhesive. Not obvious from the graph ofFig. 6 however is the catastrophic crack extensionsexhibited by the sample during testing. This leads tolimited results being available for the material. Thefracture surface remained entirely interlaminar andappears to be highly affected by strain rate. Conclusionsfor the PA Web material suggest that it provides a goodadhesive bond in the interlaminar region. At the level ofenvironmental treatment provided here the PA Web stillprovides Mode I toughening however this cannot beguaranteed at higher exposure levels due to hydrolysis.The untreated Polyolefin samples displayed significant

toughness increases. Using Fig. 7 as an indicative guideto the entire fracture surface it can be seen that thefracture remained interlaminar and therefore toughnesswas enhanced by the adhesive action of the web mate-rial. Melting of the web material during processingincreased the area covered by the web to create a moreuniform distribution of adhesive rather than a dis-continuous fibre interlayer as is usually found in SFRsamples.Polyolefin materials display the lowest moisture gains

of any of the man made fibres and are highly stable inhigh temperature/high humidity environments [13]. Thisstability is evident in the maintenance of fracturetoughness after environmental treatment. Treated sam-ples displayed toughnesses that were significantly higherthan the other SFR materials. The results of the treatedsample displayed a combination of both adhesivetoughening and crack deviation into the intra-ply regionresulting in the highest toughness of the systems exam-ined. Considerable amounts of continuous fibres havebeen dislocated yet the presence of the interlaminarreinforcement generated by the adhesive web is stillobvious. Also notable and seen in all of the treatedsamples is the quantity of exposed continuous fibres,due to matrix fracture and interfacial failure.The trend for crack propagation to remain inter-

laminar in untreated samples is reinforced in both thePO Web sample, Fig. 7, and the unreinforced sample ofFig. 5. In the PO Web sample, the interlaminar region ismarked by the presence of SFR and the continuouscarbon fibres are covered by matrix material. Theunreinforced sample of Fig. 5 displays hackling and


continuous fibres coated in a thin layer of matrix mate-rial both strong indications of interlaminar failure.In Fig. 9, the treated ZHM sample displays a combi-

nation of interlaminar and intra-ply cracking withresulting fibrillation fracture and fibre pullout. The sur-face also features a number of broken carbon fibres dueto fracture occurring in continuous bridging fibres.Similarly, the unreinforced sample of Fig. 5 and the POWeb sample of Fig. 8 contain large portions of pulledout and fractured carbon fibres. There are limitedobservable fibrillated fibres in the case of the ZHMsample due to the crack travelling primarily within theintra-ply region.Common to all of the treated samples is increased

fracture toughness and a transition from interlaminar tointra-ply cracking. This propensity for treated samplesto display intra-ply cracking has been described beforeby Birger et al. [8]. The changed fracture mode occurs asa result of leaching occurring in the partially reacteddicyandiamide curing agent. These leached particlescause the generation of voids and weaknesses to aid inthe deviation of the crack from the interlaminar regionto the intra-ply region.

6. Conclusions

The addition of SFR to the interlaminar region ofenvironmentally treated composites has not producedadverse effects. Specimen weight gains in the order of2% were found to occur in all immersed samples. Thesimilarity in weight gain is attributed to the small pro-portion of SFR or web modified material exposed to theenvironment. Crack propagation was found to be inter-laminar in all untreated samples. Polyamide and Poly-olefin web samples displayed increased fracturetoughnesses attributable to improved interlaminaradhesion Treated samples displayed a combination ofinterlaminar and intra-ply crack propagation. Intra-plycrack propagation being attributed to void creation, dueto dissolution of uncured DICY curing agent, creatingcrack deviation. Fracture toughness values in SFRsamples increased by 250–300% as a result of environ-mental treatment and were fairly consistent. This levelof toughness arose through a combination of intra-plytoughness, extensive continuous carbon fibre bridgingand limited interlaminar crack propagation. Some smalladditional increases occurred as a result of SFR fibril-lation fracture however they are sufficiently small to bealmost negligible.

Polyolefin web reinforced materials also showed250–300% toughness gains arising from environmentaltreatment. The ingress of moisture into the laminatesdid not significantly affect the bond strength of thepolyolefin web. PO Web toughness was comprised ofintra-ply crack propagation, continuous fibre bridgingand increased interlaminar bonding. Polyamide webmaterials hydrolysed during environmental treatmentcausing a weakening of the interlaminar region. Thesesamples displayed purely interlaminar crack propaga-tion and the lowest toughness of the treated samples.

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