The dependence of the fracture toughness of thermoplastic composite laminates on interfacial interaction

$Page 1: The dependence of the fracture toughness of thermoplastic composite laminates on interfacial interaction$
The dependence of the fracture toughness of thermoplasticcomposite laminates on interfacial interaction

Tie Qi Lia,b,c, Ming Qiu Zhanga,*, Kun Zhanga, Han Min Zenga

aMaterials Science Institute, Zhongshan University, Guangzhou 510275, People's Republic of ChinabDepartment of Materials Science and Engineering, Guangdong University of Technology, Guangzhou 510090, People's Republic of China

cLaboratory of Polymeric Composite and Functional Materials, The Ministry of Education of China, Zhongshan University, Guangzhou 510275,

People's Republic of China

Received 18 November 1998; received in revised form 4 August 1999; accepted 12 October 1999

Abstract

The paper investigates the in¯uence of microstructure of the matrix resin and polymer/reinforcement interface on the fractureperformance of thermoplastic composites in relation to processing conditions. Solution-pre-impregnated, unidirectional, carbon-®ber-reinforced poly(ether ether ketone) laminates prepared at di�erent melt residual times were used as the experimental materialsto obtain di�erent interfacial e�ects and matrix morphologies. Composite crystalline structures were analyzed by X-ray di�raction

and infrared spectroscopy. The stress-intensity factor was measured by means of the single-edge-notched bending method along thetwo orthogonal directions of the laminates. The results indicate that the Mode-I fracture toughness along the ®ber direction, KIC

00,increases with a rise in melt residual time, while the stress-intensity factor at failure perpendicular to the ®ber direction, KIC

? ,

changes with the melt residual time in a more complicated way. Through a careful comparison of the two stress-intensity factorswith composite microstructure, the dependence of the fracture toughness on the crystalline morphology of the matrix resin has beenfound. That is, KIC

? increases with increasing matrix crystallinity and KIC00 increases with either a decrease in the heterogeneous

crystallization or an increase in the orientation of matrix crystallites. It is, thus, proved that the fracture toughness of the laminatesdepends on crystallinity and crystalline morphology of the matrix as well as the interfacial interaction, but not merely on matrixcrystallinity. # 2000 Elsevier Science Ltd. All rights reserved.

Keywords: B. Interface; B. Fracture toughness; B. Microstructure; Polymer composites; Melt residual time

1. Introduction

Carbon-®ber-reinforced composite laminates withthermoplastic poly(ether ether ketone) (PEEK) matrixhave attracted considerable attention for more than adecade [1±3]. In addition to having a better processa-bility, the unique advantage of the material lies in thehigher delaminating resistance compared to its thermo-setting counterparts [4,5]. To understand this improve-ment in toughness, e�orts have been made to probe intothe nature of interlaminar fracture behavior in terms oflinear-elastic fracture mechanics (LEFM) [4±9] as wellas into the e�ect of the composite matrix structure andmorphology [10±12]. It was shown that ®ber bridging

ahead of a crack, substantial matrix deformation andstrong ®ber/matrix adhesion play leading roles. On thebasis of these fundamental researches, dependence ofthe composite fracture toughness on processing condi-tions seems to be predictable.However, the results are factually a little ambiguous

because interfacial bonding and matrix morphology,which depend on polymer microstructure and in turnrely on the type of prepregs and processing conditions,were not taken into overall consideration. For example,Talbot and co-workers [10] found that the fracturetoughness of APC-2 laminates (carbon-®ber/PEEKcross-ply laminates with a ®ber content of 60% byvolume produced by ICI) was mainly controlled bymatrix crystallinity and was rather insensitive to ther-mal history, but a later report suggested the crystal-lization mechanism should be the main in¯uencingfactor [12]. Therefore, further work is needed to see

0266-3538/00/$ - see front matter # 2000 Elsevier Science Ltd. All rights reserved.

PI I : S0266-3538(99 )00147-5

Composites Science and Technology 60 (2000) 465±476

* Corresponding author. Tel.: +86-20-84036576; fax: +86-20-

84036564.

E-mail address: [email protected] (M.Q. Zhang).

whether the dependence on matrix crystallinity is gen-erally applicable to the composites and how the proces-sing variables a�ect the fracture toughness by changingthe interfacial and matrix structure. This is also the aimof the present study.For this reason, carbon-®ber-reinforced PEEK lami-

nates prepared from the solution-pre-impregnated pre-forms is investigated, ensuring a uniform dispersion ofresins in and between ®ber bundles. The melt residualtime (MRT) at 400�C is selected as the processing vari-able so as to adjust interfacial interaction [13]. Polymermorphology is characterized by using di�erential scan-ning calorimetry (DSC), wide angle X-ray di�raction(WAXD) and Fourier transform infrared spectroscopy(FTIR). Laminate fracture resistance is measured bymeans of single-edge notched bending (SENB) tests andthen compared with the determined interfacial andmatrix microstructure.

2. Experimental

2.1. Materials

The polymer employed in the present work is pow-dered PEEK, kindly supplied by Jilin University, China.It is claimed to be similar to ICI 150P PEEK resin. Toprepare carbon-®ber/PEEK prepregs, eight layers of thesize-free Hercules AS4 carbon ®ber, which had beenhand-wound on a polytetra¯uoroethylene (PTFE) ®x-ture, were immersed in a 2 g/100 ml chlorophenol solu-tion of PEEK [14] and then degassed at 60 mmHgvacuum. The amount of the solution was controlled toensure the highest possible ®ber volume content, calcu-lated to be 66% by assuming random dense packing ofthe ®bers in the unidirectional ®ber preforms. Solvent inthe system was subsequently evaporated in vacuum at150�C.Prior to laminates manufacturing, the above preforms

were cut into the size ®t for a matched mold. Two piecesof the preforms were then stacked in the mold andcompressed into a unidirectionally laminated composite2 mm thick at the desired melt residual time under400�C in a hydraulic press equipped with a cooling sys-tem. The cooling rate of the mold was kept at 20�C/minunless otherwise speci®ed. In order to conduct the sub-sequent SENB measurement, the laminates weremachined with the wire electrical discharge method toavoid possible damage due to machining [15].

2.2. Structure characterization

FTIR spectra were collected on a Bio-Rad FTS 600spectrometer with a re¯ective geometry by using gold®lm as background. Each spectrum was the average of200 scanning results with a resolution of 4 cmÿ1. Prior

to the quantitative analysis, the recorded spectra weretransformed to absorbency. Ordering of the polymerwas calculated based either on the integrated intensityratio, A966/A952, or on the peak intensity ratio, I966/I952,of the 966 and 952 cmÿ1 bands.WAXD experiments were performed by using a Y-4Q

di�ractometer. A monochromatized copper CuKa raywas used with the generator operating at 40 Kv and 30mA. The slit arrangement was set at 1�, 1� and 0.4 mm,allowing a 2� resolution to be better than 0.05�. WAXDspectra were recorded in scanning mode at a scanningrate of 0.05�/s and an interval of 0.03�. For analyzingthe crystallinity of the matrix resin, the measured inten-sity versus 2� curve from 11 to 36� was normalized andthen the scattering of carbon ®ber, which had also beennormalized, was subtracted from it. The resultant dif-ference spectrum was then de-convoluted to four Lor-entzian components corresponding to the re¯ections of(110), (111), (200) and the amorphous phase, respec-tively. The center of the amorphous background waslocated at 2� � 20� 0:2� according to the di�erencespectrum of a quenched laminate. The method is thesame in physical nature as the traditional one [10,16]and the arbitrariness in background subtraction can,thus, be prevented. Lorentzian polarization correctionwas made prior to the quantitative analysis.With respect to DSC characterization, Perkin±Elmer

DSC-2 equipment was used. The specimens were heatedto 400�C at 10�C/min, and then cooled at a rate ofÿ10�C/min and heated up again to 400�C at 10�C/min.The exothermic enthalpy on the cooling trace and theendothermic enthalpy of the second heating scan wererecorded. In the case of crystallinity calculation, theheat of fusion of fully crystalline PEEK was assumed tobe 130 J/g [16].

2.3. Mechanical properties testing

SENB experiments were conducted on 12.7�56�2mm3 specimens according to the method designated forrigid plastics [17]. The notch was made by cutting agroove 3mmwide with a hand saw and then tapering to thedesired length with a new razor blade. Great care was takenfor the specimens with ®ber axis parallel to the notch.As revealed by the load versus displacement plots in

Fig. 1, it is found that the experiments with crackgrowth direction parallel to the ®ber axis satisfy the rule

Pmax=P5% < 1:1 �1�

whereas those with crack growth direction perpendi-cular to the ®ber axis are not the general case (Pmax andP5% stand for the maximum load and the load allowing5% non-linearity, respectively). Stress intensity factor,KQ, is calculated from the load, PQ, which is the lowerone of Pmax and P5%:

466 T.Q. Li et al. / Composites Science and Technology 60 (2000) 465±476

KQ � fPQ=�BW1=2� �2�

where B and W denote specimen thickness and width,and f is an orthotropic correction factor with its valuedepending on the elastic properties and the crackgrowth direction [18±20]. When a specimen was testedwith the ®ber axis parallel to the crack propagatingdirection, the great sti�ness along the ®ber directionmade it unnecessary to check further the linearity of theresults. The corresponding KQ is thus taken as theMode-I fracture toughness of the composites along thisdirection and is denoted as KIC

00. On the other hand, KQ

values along the other direction never satisfy the linearelastic requirements and the crack initiation processinvolved substantial ®ber full-out and matrix deforma-tion [15]. The stress intensity factor at fracture initiationis written as KIC

? in the following text for convenience,noting that it does not describe the crack growth resis-tance transverse to the ®ber direction. Accordingly, Eq.(2) is rewritten as:

K00IC � �Eo=E22�1=2PQ=�BW1=2� �3�K?IC � �Eo=E11�1=2PQ=�BW1=2� �4�

where E11 and E22 denote the Young's moduli of thelaminates, Eo is the orthotropic modulus and is given asfollows for the plane stress situation [18±20]:

Eo � �2E11E22�1=2��E11=E22�1=2 � �ÿ2�12=E11

� 1=G12�E11=2�ÿ1=2 �5�

where �12 and G12 stand for Poisson's ratio and shearmodulus, respectively, and the subscripts 1, 2 and 3denote x, y, and z axes, respectively.

For the purposes of providing elastic properties of thecomposites, the ultrasonic immersion technique, a non-destructive technique by which the ®ve independentelastic constants of an uniaxial anisotropic material canbe obtained [21], was used. During the measurements,the specimen, the transmitting and the receiving trans-ducers were immersed in a tank ®lled with water. Bychanging the angle of the incident ultrasonic beam, aseries of refracted wave velocities were measured. Theelastic constants can, thus, be determined and trans-formed into moduli and Poisson's ratios.

3. Results and discussion

3.1. Crystallinity

Fig. 2 illustrates the crystallinity of the laminates cal-culated from the melting enthalpy measured with thecalorimeter. It can be seen that the DSC crystallinity,XDSC

c , increases with a rise in MRT from 4 to 16 min. Inthe case of short and long MRT, the laminates have arelatively low crystallinity of 16�1%.By means of the WAXD technique, however, a dif-

ferent dependence of crystallinity on MRT can beobserved (Fig. 3). Except for the lower value found forthe composite with MRT of 4 min, the crystallinitytends to increase with increasing MRT and thendecrease. A maximum value of 28% is again found at 16min of MRT. The unique S-shape dependence of crys-tallinity can be further proved by FTIR measurement,as shown in the same ®gure.The apparent inconsistency between the results given

by DSC, WAXD and FTIR methods should be rootedin their working principles. Both WAXD and FTIR

Fig. 1. Typical recorrected load (P)/displacement (u) plots of the laminates obtained from SENB tests.

T.Q. Li et al. / Composites Science and Technology 60 (2000) 465±476 467

characterization are non-destructive in nature, while theDSC method is not. Crystallites in a specimen wouldmelt and recrystallize in the heating process of a DSCexperiment. For example, it has long been known thatPEEK tends to be annealed during a DSC scan [22]. Themeasured melting enthalpy, therefore, is in¯uenced bythe in situ morphology change in the case of DSC mea-surement and is not a proper description of real crys-tallinity of the specimen for the current study. In otherwords, crystallinity determined by the WAXD andFTIR methods could re¯ect the e�ect of MRT in amore objective way.According to our recent report on the bulk crystal-

lization in AS4/PEEK laminates [13,15], prolongedMRT a�ects the crystallization of PEEK through twomechanisms. On the one hand, prolongedMRT facilitates

short-range interaction by promoting absorption ofmacromolecular chains onto the carbon ®ber surface.As a result, heterogeneous nucleation is enhanced,leading to an increase in primary crystallization (the®rst stage of crystallization, consisting of outwardgrowth until impingement) and secondary crystal-lization (corresponding to the growth ®lling in theinterstices left by the primary one) as well. Because thetime required to form the short range interaction isabout a few minutes for the system [15,23], the increasein the contribution of secondary growth accounts forthe decrease in bulk crystallinity when MRT increasesfrom 2 to 4 min.On the other hand, a rise in MRT also strengthens the

long-range e�ect of carbon ®ber due to the appearance ofthermal strain in bulk polymer induced by the mismatch

Fig. 2. DSC crystallinity, XDSCc , of the laminates as a function of MRT.

Fig. 3. Crystallinity of the laminates determined with WAXD and FTIR methods, XWAXDc and XFTIR

c (A966=A952), as a function of MRT.


in thermal properties between reinforcements andmatrix resins. With a predetermined short-range inter-facial interaction as the prerequisite, bulk nucleation isimproved but the growth processes tend to be restricted.Bulk crystallization rate is, thus, reduced, especiallywhen a strong short-range interfacial interaction hasbeen formed. In summary, the aforesaid MRT depen-dent variation in composite crystallinity is the result of acompetition between nucleation and restriction on crys-tallite growing, which brings about the maximum crys-tallinity at a moderate MRT of 16 min when MRT islonger than 4 min (Fig. 3).In fact, the two-stage intensi®cation of long-range

interfacial interaction can also be perceived with theDSC method. Fig. 4 shows the normalized enthalpydi�erence ��Hm ÿ�Hc�=�Hm with MRT, where �Hm

is the endothermic enthalpy on a DSC heating trace and�Hc the exothermic enthalpy on a cooling trace. As ameasure of the extent of in situ annealing in the DSCcell, ��Hm ÿ�Hc�=�Hm allows a description of therestriction e�ect on crystalline growth during cooling.The higher the parameter, the larger the restrictionplaced by the reinforcement [24]. It is seen from Fig. 4that ��Hm ÿ�Hc�=�Hm increases with MRT over tworanges of MRT. It means the restriction on growthprocesses due to the wetting and absorption of thepolymer on the ®ber surface is enhanced mainly within 8min, while that resulting from thermal strain becomes moreand more signi®cant when MRT is longer than 16 min.

3.2. Matrix morphology

It should be noted that the crystallinity discussed inthe last section characterizes the overall crystalline por-tion of the polymer matrix but is not capable of

describing the details dealing with ®ne structure. Acareful analysis of the individual re¯ections on theWAXD spectra should obtain more useful informationin this direction.As the re¯ections of (111), (110) and (200) have dif-

ferent crystallographic originations, their crystallinitydependences on MRT are given separately in Fig. 5. Asexpected, they are di�erent from each other. Fig. 5illustrates that the partial crystallinity corresponding tothe (111) re¯ection, XWAXD

c;�111� , reaches the maximum whenMRT is 16 min, just like the overall WAXD and FTIRcrystallinity (see Fig. 3). However, the minimum at 4min of MRT appearing in Fig. 3 is no longer found. TheMRT dependence of XWAXD

c;�111� factually resembles that ofthe FTIR peak height ratio of the 966 and 952 cmÿ1

bands (Fig. 6). Generally, the height of a FTIR bandcharacterizes the most probable conformation of thecorresponding conformer and, an integrated intensity ofa FTIR band characterizes content of the conformer. Inother words, the FTIR peak intensity ratio measures theordering of the matrix, rather than the overall crystal-linity described by the integrated intensity ratio [13].Therefore, XWAXD

c;�111� can only be related to the matrixordering, as indicated in [25] the (111) re¯ection has nospeci®c crystallographic meaning. The crystalline fea-tures of PEEK in interphase could not be observed fromthe plot of XWAXD

c;�111� against MRT. On the other hand, thetwo-stage variation of XWAXD

c;�111� with MRT in Fig. 5further demonstrates the long-range e�ect of carbon®ber on matrix crystallization as follows [13]. In the caseof shorter MRT, an increase in MRT tends to acceleratematrix crystallization. In the case of longer MRT,matrix crystallization is slowed down owing to stricterrestriction on crystalline growth imposed by carbon®ber.

Fig. 4. Normalized enthalpy di�erence, ��Hm ÿ�Hc�=�Hm, of the laminates as a function of MRT.


In contrast, the intensities of (110) and (200) re¯ec-tions vary with MRT in di�erent ways (Fig. 5). XWAXD

c;�110�increases with a rise in MRT except in the shorter andlonger MRT regions. With increasing MRT, on theother hand, XWAXD

c;�200� tends to decrease from 2 to 4min, but increases to a higher level when MRTreaches 8 min and decreases again with a furtherincrease in MRT.In consideration of the fact that the (110) re¯ection

reveals the level of chain orientation in the specimen, asevidenced in many neat thermoplastic crystalline poly-mers, the quantity XWAXD

c;�110� can act as a measure of chainorientation to some extent. Hence, the above observedtrend of XWAXD

c;�110� supports our deduction that prolonged

MRT promotes long-range e�ect of carbon ®ber onchain conformation.It is worth noting that the (200) re¯ection is strength-

ened when MRT increases from 4 to 8 min (Fig. 5).Based on the heterogeneous nucleation model proposedby Lovinger and co-workers [25], it was proved that themelt-born, heterogeneous nuclei in the AS4/PEEK sys-tem should also be in the form of face-on lamellae [26].A quasi-epitaxial growth from the (102) fronts of theseface-on lamellae would have a (200) lattice that formsan angle to ®ber surface as low as 5.6�, though thegrowth of the face-on lamellae itself does not favor the(200) re¯ection. Therefore, the rise in (200) re¯ection inthis MRT range is attributed to a signi®cant increase in

Fig. 5. Partial crystallinity obtained from the individual WAXD re¯ections (111), (110) and (200) of the laminates, XWAXDc;�111� , X

WAXDc;�110� and XWAXD

c;�200� , as afunction of MRT.

Fig. 6. Partial crystallinity calculated from the (111) WAXD re¯ection, XWAXDc;�111� , and the FTIR ordering calculated from the peak intensity ratio,

XFTIRc (I966=I952), of the laminates as a function of MRT.


heterogeneous nucleation consequent on the wettingand absorption of chains on the ®ber surface. The par-tial crystallinity of the (200) re¯ection, XWAXD

c;�200� , is thuscorrelated to the content of heterogeneously nucleatedinterfacial crystallites.

3.3. Fracture toughness of the laminates

The critical stress intensity factors of the laminateswith crack propagating perpendicular to and along the®ber axis are shown in Figs. 7 and 8, respectively. FromFig. 7 it can be seen that the stress intensity factor of theformer case, K?IC, varies with MRT in a similar way tothe bulk crystallinity (Fig. 3), while Fig. 8 indicates thefracture toughness, K00IC, increases with increasing MRTup to 32 min. The general trend for all the specimenscan be described with a semi-logarithmic relationship:

K00IC � 0:06��0:15� � 0:48��0:13� log MRT �6�

where K00IC is in MN/m3=2 and MRT is in min.By examining the stress intensity factor in relation to

crystallinity (Fig. 9), the increasing trend with XWAXDc

can be found for K?IC:

K?IC � ÿ0:42��0:50� � 0:06��0:02�XWAXDc �7�

where K?IC is in MN/m3=2 and XWAXDc is in percentage.

In fact, a better linearity can be observed in the specimenswhen the longest MRT is excluded from consideration(refer to the dashed regression line in Fig. 9):

K?IC � ÿ2:01��:0:01� � 0:11��0:01�XWAXDc �8�

though the negative truncation has no clear physicalmeaning. Because higher crystallinity leads to higher

shear modulus and shear strength but lower Mode-Ifracture toughness of neat PEEK [10], the increase inK?IC, therefore, further indicates that the parameter hasmuch to do with the shear deformation of the matrixpolymer and may be comparable to the Mode-II frac-ture toughness of the composites. On the other hand,the fracture toughness, K00IC, does not show any simpledependence on crystallinity (Fig. 9). According to theresults of Ref.[10] and the additive law of compositematerials, an increase in crystallinity of the currentlaminates should be accompanied by a decrease in K00IC.However, the data in Fig. 9 show this trend with poorcon®dence:

K00IC � 0:91��1:16� ÿ 0:01��0:04�XWAXDc �9�

It strongly suggested that other factors such as inter-facial adhesion and crystallites orientation must play animportant role in determining K00IC, which should bediscussed hereinafter.

3.4. E�ect of interfacial structure and matrix morphologyon the fracture resistance

As mentioned above, the evident deviation of thecrystallinity dependence of Mode-I toughness from thedesired linearity puts a question to the deformationmechanism during the fracture process. Due to the highstress concentration and the heterogeneity in the lami-nates, localized yielding of the tough PEEK is unavoid-able [27]. Correlation of stress intensity factor withindividual WAXD re¯ections would help to reveal thedependence of fracture toughness on morphology of thepolymer. From Fig. 10, it can be seen that K?IC varieswith XWAXD

c;�111� linearly except for the laminates with aMRT of 4 min:

Fig. 7. Stress intensity factor K?IC of the laminates as a function of MRT.


K?IC � ÿ0:49��0:84� � 0:29��0:02�XWAXDc;�111� �10�

where the partial crystallinity XWAXDc;�111� is also in percen-

tage. The result agrees well with Eq. (7). However, theinversely proportional relationship between the fracturetoughness K00IC and XWAXD

c;�111� is tenable only in the case ofconsiderable deviation:

K00IC � 1:47��1:42� ÿ 0:16��0:26�XWAXDc;�111� �11�

K00IC factually increases with increasing XWAXDc;�111� in the

MRT ranges 2±16 min and 32±64 min, but decreases witha rise in XWAXD

c;�111� in the MRT range 16±32 min (Fig. 10).The decrease of matrix Mode-I toughness with a rise in

its crystallinity as found in Ref. [10] cannot account forthe current increasing trends of the laminates Mode-Ifracture toughness within the two ranges of crystallinity.Obviously, an increase in interfacial interactions plays amore important role than crystallinity.The in¯uence of interfacial structure and matrix mor-

phology becomes more evident when K?IC and K00IC arecompared with the partial crystallinity calculated fromthe (200) re¯ection, XWAXD

c;�200� , a parameter relating closelyto crystallites growing under strict geometric constraintsof the ®bers [28,29] or to crystalline growth indirectlyresulting from heterogeneous nuclei [26]. As shown inFig. 11, agreement with the rule of mixture can be foundfor both K?IC and K00IC, though the scattering is relativelysigni®cant:

Fig. 8. Fracture toughness K00IC of the laminates as a function of MRT.

Fig. 9. Stress intensity factors, K?IC and K00IC, versus crystallinity determined by WAXD, XWAXDc (the numerals stand for MRT in minutes).


K?IC � 0:39��0:35� � 0:07��0:04�XWAXDc;�200� �12�

K00IC � 1:21��0:56� ÿ 0:07��0:06�XWAXDc;�200� �13�

where XWAXDc;�200� is in percentage. In fact, the linear rela-

tionship between K?IC and XWAXDc;�200� falls into two cate-

gories according to di�erent MRT (refer to the dashedand dotted regression lines in Fig. 11). The slope corre-sponding to the specimens with shorter MRT (2±16min) is higher than that corresponding to the specimenswith longer MRT (8±64 min), coinciding with the fact

that matrix morphology is controlled by di�erentnucleation and growth mechanisms at the two ranges ofMRT [13]. Sti�ness of polymer, indeed, determines K?ICprovided the nucleation and growth mechanisms ofcrystallites are the same.On the other hand, according to the data obtained at

a similar loading rate by Talbot and co-workers [10], theMode-I stress intensity factor, KQ, of the ICI 150Ppolymer was found to decrease with increasing XWAXD

c

KQ � 16:68��1:63� ÿ 0:38��0:06�XWAXDc �14�

Fig. 10. Stress intensity factors, K?IC and K00IC, versus partial crystallinity calculated from the (111) WAXD re¯ection, XWAXDc;�111� (the numerals stand for

MRT in minutes).


MRT in minutes). The linear regression for K?IC � XWAXDc;�200� in the MRT range 2±16 min (dashed line) is: K?IC � ÿ1:11��0:75� � 0:21��0:07�XWAXD

c;�200� ,and that in the MRT range 8±64 min (dotted line) is: K?IC � 0:43��0:14� � 0:08��0:01�XWAXD

c;�200� .


where KQ is still in MN/m3=2 and XWAXDc varies between

15 and 40%. It is seen that the dependence of K00IC oncrystallinity, as re¯ected in Eq. (13), is quite similar tothis. Since the KQ value determined in Ref.[10] is linearelastic in nature only when the specimen has a crystallinityhigher than 30%, and the overall WAXD crystallinity ofthe present laminates ranges from 20 to 30%, the rela-tively large deviation of the inverse K00IC dependence onXWAXD

c;�111� and XWAXDc;�200� (Figs. 10 and 11) manifests that the

deformation of PEEK in the macroscopicMode-I fracture,is not necessarily linear elastic in nature on microscopicscale, as deduced at the beginning of this section. Infact, an increase in K00IC with a rise in the individualpartial crystallinity can be perceived in the case of all the

three re¯ections at a certainMRT range (Figs. 10±12), butthe increase with XWAXD

c;�110� is the most obvious.As discussed in the previous section, the physical

meaning of the (111), (200) and (110) re¯ections is notthe same. XWAXD

c;�111� and XWAXDc;�200� re¯ect the bulk matrix

crystallinity and the interfacial crystallinity, respec-tively, while XWAXD

c;�110� indicates the level of main chainorientation. Comparing the MRT dependence ofXWAXD

c;�110� (Fig. 5) with the XWAXDc;�110� dependence of K00IC

(Fig. 12), it can be concluded that chain orientationmakes greater contributions to the increase in K00IC.Although further crystallographic study, for example, apole plot examination [30], is required to discuss theexact nature of this phenomenon, the dependence of K00IC


MRT in minutes).

Fig. 13. E�ect of MRT on stress intensity factor ratio K00IC=K?IC.


on chain orientation is readily understood according tothe Mode-I fracture mechanism of the unidirectionallaminates. That is, the chain breakage is inevitablyinvolved in the macroscopic Mode-I fracture. Conse-quently the orientation of macromolecular chainsimproves the fracture resistance of the laminates. TheMode-I toughness of the laminates, therefore, also relieson chain orientation.So far the chain orientation plays an important role in

determining K00IC, it is worth comparing K00IC with K?IC.From Fig. 13, it can be seen that the ratio K00IC=K

?IC

increases with a rise in MRT and changes from the ratioof Young's moduli to that of shear moduli. Clearly,matrix shearing makes more and more contribution toK00IC when MRT is increased. Fig. 14 illustrates thatK00IC=K

?IC decreases linearly with the integrated intensity

ratio of the (200) and (110) WAXD bands, A200=A110:

K00IC=K?IC � 1:79��0:18� ÿ 1:58��0:23�A200=A110 �15�

It means that the fracture resistance along the ®berdirection tends to increase with either a decrease in het-erogeneous nucleation or an increase in crystalliteorientation. In summary, fracture resistance of thelaminates depends strongly on the matrix morphology,but not merely on polymer crystallinity.

4. Conclusions

Microstructure and fracture toughness of the solutionpre-impregnated AS4/PEEK laminates are character-ized in this paper. Through a careful analysis of thestress intensity factors and crystalline details of matrixresin, it can be concluded that the fracture toughness

relies on both interfacial interactions and matrix mor-phology.The SENB K?IC is closely related to bulk crystallinity

of the composites, which is in turn determined by theshort-range and long-range e�ects of the ®ber duringprocessing. A linear increase in the stress intensity fac-tor with increasing crystallinity is a property that can beexpected for Mode-II experiments, rather than for pre-sent Mode-I tests. Further studies are needed to under-stand the meaning of the quantity.The fracture toughness along the ®ber direction, K

00IC,

depends more signi®cantly on interfacial adhesion andpolymeric chain orientation. It has been proved that K00ICwould increase either with a decrease in heterogeneouslynucleated crystallization or with an increase in theorientation of matrix crystallites.In short, fracture toughness of the present composite

system depends on both crystallinity and crystallitestexture of the matrix polymer in which interfacial interac-tion plays an important role, but not merely on crystallinityas suggested for APC-2 laminates [10]. It is, thus, believedthis complex dependence on morphology leaves spacefor optimizing both the elastic properties and the frac-ture toughness in the processing of the thermoplasticcomposites.

Acknowledgements

The support of the National Natural Science Foun-dation of China (Grant No. 59725307), the Trans-Cen-tury Training Program Foundation for the Talents theMinistry of Education of China are gratefully acknowl-edged. Dr. T. Q. Li is grateful to the support of theLaboratory of Polymeric Composite and Functional

Fig. 14. Stress intensity factor ratio, K00IC=K?IC, versus integrated intensity ratio of the (200) and (110) WAXD re¯ections, A200=A110 (the numerals

stand for MRT in minutes).


Materials, the Ministry of Education of China (GrantNo.19703). The authors wish to thank Prof. Yiu-WingMai for his valuable comments, as well as Mr. HaoZhang and Prof. Yu-Long Liu for their helps in IRexperiments.

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The dependence of the fracture toughness of thermoplastic composite laminates on interfacial interaction