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Modes I and II interlaminar fracture toughness and fatiguedelamination of CF/epoxy laminates with self-same epoxy interleaf

Masaki Hojo a,*, Tadashi Ando a, Mototsugu Tanaka a, Taiji Adachi a,Shojiro Ochiai b, Yoshihiro Endo c

a Department of Mechanical Engineering and Science, Kyoto University, Kyoto 606-8501, Japanb International Innovation Center, Kyoto University, Kyoto 606-8501, Japan

c Toho Tenax Co., Ltd, Nagaizumi, Shizuoka 411-8720, Japan

Available online 20 March 2006

Abstract

Interlaminar fracture toughness and delamination fatigue crack growth behavior were investigated for carbon fiber (CF)/epoxy lam-inates with the self-same epoxy interleaf. The matrix epoxy with a thickness of 50 lm was chosen as the interleaf material in order toclarify the effect of resin-rich layer thickness on the delamination fatigue crack growth behavior. Tests under mode I loading were carriedout using double cantilever beam specimens. For tests under mode II loading, three-point end notched flexure specimens were used forinterlaminar fracture toughness tests, while four-point end notched flexure specimens were used for delamination fatigue tests. The modeI properties (interlaminar fracture toughness and delamination fatigue threshold) of these epoxy-interleaved CFRP laminates werealmost identical to those of the laminates without interleaf (base CFRP laminates). The effect of epoxy interleaf was completely differentunder mode II loading. The mode II interlaminar fracture toughness for the epoxy-interleaved laminates was 1.6 (initial value) and 3.4(propagation value) times higher than that for the base CFRP laminates. The mode II delamination fatigue threshold of the epoxy-inter-

leaved laminates was 2–2.3 times higher than those of the base CFRP laminates. While the toughness of the interleaf is the key factorunder mode I, the thickness of interlayer is the key factor under mode II. The difference in the effect of the self-same epoxy interlayer onthe interlaminar fracture properties under modes I and II loadings was discussed on the bases of the fractographic observations andmechanism considerations. 2006 Elsevier Ltd. All rights reserved.

Keywords: Polymer–matrix composites; Fatigue crack growth; Delamination; Self-same epoxy interleaf; Mode I; Mode II

1. Introduction

Although advanced composite materials such as carbonfiber reinforced plastics (CFRP) have attained a certain

position as structural materials for aircraft, demands forhigher fuel efficiency require further increasing the struc-tural weight ratio for composite materials. The Boeing 787project is reported to have overcome difficulties in replacingmain wing and fuselage with CFRP, and thus the expectedweight ratio of composite materials is 50% [1]. AirbusA380, the largest commercial aircraft with increased weightratio of composite materials to 25%, is almost ready [2].

Since poor interlaminar strength is one of the design limitingfactors in CFRP laminated structures [3], improving inter-laminar fracture toughness and resistance to fatigue delam-ination is still an important topic [4–6].

Attempts to improve the interlaminar fracture tough-ness and the delamination fatigue properties of CFRP lam-inates have so far shown various results. Some levels of toughening have already been achieved, both under staticand fatigue loading, by replacing matrix resin with atougher thermoplastic system [4,7,8]. However, the costand the processing difficulties lower the applicability.Another way to increase the interlaminar properties is toreplace only the resin layer at the prepreg interface with atougher system [9–11]. This is often referred to as ‘‘inter-leaf’’ or ‘‘interlayer’’ method. In particular, T800H/3900-2

International Journal of Fatigue 28 (2006) 1154–1165

www.elsevier.com/locate/ijfatigue

International Journalof

Fatigue

0142-1123/$ - see front matter 2006 Elsevier Ltd. All rights reserved.

doi:10.1016/j.ijfatigue.2006.02.004

* Corresponding author. Tel.: +81 75 753 4836; fax: +81 75 771 7286.E-mail address: [email protected] (M. Hojo).

mailto:[email protected]:[email protected]

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with a heterogeneous interlayer consisting of fine thermo-plastic particles has shown both excellent compressivestrength after impact (CAI strength) and hot-wet charac-teristics [12–14]. T800H/3900-2 has already been appliedto primary structures of Boeing 777 [13]. Ionomer inter-leaved CFRP laminates have shown higher toughness

under both modes I and II loadings [15–17].In these interleaf or interlayer methods, there are twofactors involved in the improvement of the interlaminarproperties. One is that interleaf or interlayer has highertoughness than the resin-rich layer of the base laminates.Another is the increase in resin-rich layer thickness. How-ever, inserting a tougher resin layer by the interleaf or inter-layer method usually makes resin-rich layer thicker [10,13].Thus, for the better design of interleaved or interlayertoughened laminates, it is necessary to clarify separatelythe effect of interleaf/interlayer toughness and that of theresin-rich layer thickness on the improvement of the inter-laminar properties by controlling the mesoscopic struc-

tures. Though Partridge and coauthors have already triedthis subject using laminates with self-same epoxy matrixlayer [11,18], the obtained information was limited.

In the present study, unidirectional CF/epoxy laminateswith self-same epoxy interleaf were specially prepared. Theeffect of self-same epoxy interlayer on the modes I and IIinterlaminar fracture toughness and delamination fatiguecrack growth behavior were investigated. The differencesin fracture mechanisms under modes I and II loadings werediscussed on the bases of the mechanism consideration andmicroscopic observation.

2. Experimental procedure

2.1. Materials and specimens

Unidirectional laminates with a nominal thickness of 3 mm (Base laminates: (0)24) were fabricated from carbon

fiber/epoxy prepregs of Toho UT500/111 using an auto-clave. The epoxy #111 used in this study is rather a conven-tional one. It has a curing temperature of 120 C, and isused as matrix resin for general purposes. Polyimide filmsof 13 lm in thickness were inserted during molding at themidplane as initial cracks for these laminates.

The matrix epoxy #111 with a thickness of 50 lm wasalso used as interleaf material. The epoxy-interleaved lam-inates (012/111/012) were fabricated using the followingprocess in order to maintain the epoxy interleaf thickness:(1) The upper and lower part of the base lamina whichconsists of 12 layer prepregs, (0)12, was half-cured byapplying the normal cure pressure. (2) The two 50 lm

thick epoxy films were half-cured. (3) Then, the upperpart of the base lamina (0)12, the two 50 lm thick epoxyfilms and the lower part of the base lamina (0)12 were laidup. (4) This was followed by the molding process of applying only the vacuum pressure. Two 50 lm-thickepoxy #111 films were inserted at the mid plane of thelaminates. Polyimide films of 13 lm in thickness wereinserted between these two epoxy films as initial cracksfor the epoxy-interleaved laminates.

Fig. 1 shows the optical micrographs of the transversesection for the base laminates and the epoxy-interleaved

Fig. 1. Optical micrographs of transverse sections for base laminates and epoxy-interleaved laminates, and method for measurement of resin-rich layer

thickness. (a) Method for measurement of resin-rich layer thickness. (b) Optical micrographs.

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laminates, and the definition of the resin-rich layer thick-ness. The resin-rich layer thickness of the base laminates

and that of the epoxy-interleaved laminates were measuredas 15 and 50 lm, respectively. Although two 50 lm-thickepoxy films were inserted, the resulting total interlayerthickness of the self-same epoxy matrix was 50 lm. Thisepoxy-interleaved laminates is named as 50 lm-epoxy-interleaved laminates in this study. The fiber volumefractions for the base and the 50 lm-epoxy-interleavedlaminates were 60% and 56%, respectively.

Double cantilever beam (DCB) specimens (widthB = 20 mm, length L = 140 mm, and nominal thickness2h = 3 mm) were used for tests under mode I loading.Fig. 2 shows the DCB specimen and aluminum blocks

for load introduction [8,19]. Special loading apparatus withuniversal joints were used in these tests [7,8]. The side sur-faces of the specimens were polished with abrasive papersand diamond paste (1 and 6 lm). Then, white brittle paintwas coated on both sides of the test specimen to enable thecrack length measurement. Mode I precracks of 2–5 mm inlength were introduced into the specimens by clamping thespecimens across the entire width at approximately the endof the starter slits and then manually wedging open thespecimens [20]. Crack length, a, was measured on bothsides of specimens with traveling microscopes. Initial cracklength was 25 mm.

End notched flexure (ENF) specimens were used fortests under mode II loading. Four-point ENF tests werecarried out for static fracture toughness in order to stabilizecrack growth [21,22]. Specimens of width, B = 20 mm,length between the supports, 2L = 100 mm, totallength = 160 mm and nominal thickness, 2h = 3 mm, wereused for four-point ENF tests. Conventional three-pointENF tests were carried out for tests under fatigue loading.The specimens of width, B = 10 mm, length between thesupports, 2L = 100 mm, total length = 160 mm and nomi-nal thickness, 2h = 3 mm were used for three-point ENFtests. Fig. 3 shows the ENF specimens. Mode I precracksof 2–5 mm in length were introduced in all specimens. To

avoid the friction between fracture surfaces, a polytetraflu-

oroethylene (PTFE) film (length in specimen longitudinaldirection: 5 mm, thickness: 0.1 mm) was inserted in the ini-tial crack at the position above the left support for thethree-point ENF specimens, and above the left support

and below the left loading nose for the four-point ENF

10

Aluminum block

Polyimide film

1 0

5

5

25

140

3

2 0

Fiber direction

5

10

Aluminum block

Polyimide film

1 0

5

5

25

140

3

2 0

5

Fig. 2. DCB specimen with aluminum blocks (dimensions are in mm).

Fig. 3. ENF specimen for mode II test (Dimensions are in mm). (a) Four-point ENF specimen for mode II static test. (b) Three-point ENFspecimen for mode II fatigue test.

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specimens [23]. Special loading apparatus with universal joints were used for tests [19,24].

2.2. Fracture mechanics parameter

The energy release rate under mode I loading was calcu-

lated using the modified compliance calibration method.The exact equations used in this study are presented inRef. [20]. Since the compliance calibration curves for fati-gue loading agreed well with those for static loading, thestatic curves were also used for the calculation of theenergy release rate under fatigue loading. The energyrelease rate under mode II was calculated using experimen-tally obtained compliance curves for each specimen. Herethe equations were designed to avoid the effect of the spec-imen thickness scatter [19].

Assuming the materials as homogeneous and aniso-tropic, the stress intensity factor K i (i = I, II) was calcu-lated using the relation between the energy release rate G i (i = I, II) and K i (i = I, II) as follows:

G i ¼ H i K 2i ; ð1Þ

where H i is a function of elastic moduli (E ij , mij ) [25]. H I wascalculated as 6.74 · 1011 Pa1 for the base laminates and7.28 · 1011 Pa1 for the 50 lm-epoxy-interleaved lami-nates. H II was calculated as 2.35 · 10

11 Pa1 for the baselaminates and 2.59 · 1011 Pa1 for the 50 lm-epoxy-inter-leaved laminates.

For fatigue tests, the stress intensity range, DK , isdefined as DK = K max K min = K max(1 R), where K maxand K min are the maximum and minimum values of the

stress intensity factor, and R is the stress ratio of theminimum load to the maximum load. The maximum andminimum energy release rates, G max and G min, are definedcorresponding to the maximum and minimum loads. Wehave shown the following equivalent stress intensity range,DK eq, as a controlling parameters under various R values inorder to discuss the stress ratio dependency of fatiguedelamination [26]:

D K eq ¼ D K 1 Rð Þc ¼ D K 1c K cmax ð2Þ

where the stress-ratio effect parameter, c (0 < c < 1), indi-cates the relative contribution of the maximum load tothe cyclic load in determining the crack propagation rate.The c values are obtained from the experimental relation-ship between DK and (1 R) [26].

2.3. Fracture toughness test and fatigue test

The tests were carried out in a computer-controlled ser-vohydraulic testing system (Shimadzu 4880, 9.8 kN) withoriginal software [8,26]. Load cells of 490 N and 2.45 kNin capacity were attached for tests under modes I and IIloading, respectively. For static tests, the cross head speedwas controlled at 0.5 mm/min in DCB tests [20] and0.2 mm/min in four-point ENF tests [21]. Initial values of

the fracture toughness, G Ic, were determined from the onset

of the nonlinearity in the initial load–loadline displacementcurves (NL point) [20].

For fatigue tests, the stress ratio, R, of the minimumload to the maximum load was kept constant at 0.1 and0.5. The frequency of stress cycling was 10 Hz. The cracklength was computed from the measurement of the compli-

ance calibration curves [8,20].For fatigue delamination tests under mode I loading, thefatigue crack growth resistance increases with the incre-ment of the crack length, Da, similar to the rising R-curvesunder static loading. Fig. 4 schematically indicates thiseffect. G max-constant test has been proposed where the datapoints for Da = 0 give the exact da/dN G max relation andthe threshold [27,28]. However, this method is time con-suming and requires a large number of specimens. Anotherpossible way is the iteration of the load-shedding ( G max-decreasing) tests developed by the authors (Fig. 5) [29].

a=0

a : Large

Gmax-constant test

Gmax-decreasing test

L o g ( d a / d N )

Log(Gmax)

a=0

a : Large

Gmax-constant test


L o g ( d a / d N )

Log(Gmax)

Fig. 4. Effect of Da on fatigue crack growth behavior for G max-constanttest and G max-decreasing test.


Gmax-constant test

at a=0

1st test

2nd test

3rd test

Log(Gmax)

L o g ( d a / d N

)


Gmax-constant test

at a=0

1st test

2nd test

3rd test

Log(Gmax)

L o g ( d a / d N

)

Fig. 5. Fatigue crack growth behavior and saturation of growth threshold

for iteration of several G max-decreasing tests.


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The procedure is as follows: (1) Carry out the first G max-decreasing test to obtain a tentative threshold. (2) Startthe second G max-decreasing test by using new specimen atthe tentative threshold value of the former test. Tests arerepeated until the saturation of the tentative threshold val-ues. A normalized gradient of the energy release rate (1/G )

dG /da, was controlled at 0.1 mm1

during load shedding.Since the fatigue crack growth resistance is often notaffected by the increment of the crack length, normalG max-decreasing tests were carried out under mode IIloading.

3. Results and discussion

3.1. Mode I interlaminar fracture toughness tests

Fig. 6 indicates the relation between the load and thecrack opening displacement for the mode I interlaminarfracture toughness tests. The load–displacement curves

for both laminates were almost linear up to the maximumpoint (P max). Stick-slip behavior was observed only for the50 lm-epoxy-interleaved laminates.

Fig. 7 shows the relation between the mode I fracturetoughness and the increment of crack length (R-curves)for both laminates. The propagation values of the fracturetoughness, G IR, for the base laminates gradually increasedwith the increment of the crack length, Da. The G IR valuesfor the 50 lm-epoxy-interleaved laminates rapidlyincreased, and then leveled off where Da is larger than5 mm. While the initial values, GIc, for the 50 lm-epoxy-interleaved laminates (170 J/m2) were almost identical with

those for the base laminates (160 J/m2

), the averaged G IRvalues (Da > 10 mm) for the 50 lm-epoxy-interleaved lam-inates (280 J/m2) were slightly higher than those for thebase laminates (230 J/m2). This tendency is similar to thosereported by Partridge [18]. She reported that the mode I

interlaminar fracture toughness with self-same epoxymatrix interleaf only increased when toughened epoxywas used as matrix. In the present study, the G IR valuesfor the 50 lm-epoxy-interleaved laminates were calculatedfrom the maximum values during stick-slip behavior whichcorrespond to stable crack growth. Table 1 summarizes theresults of mode I fracture toughness tests for bothlaminates.

3.2. Mode I delamination fatigue test

Fig. 8 shows the relation between the crack propagationrate and the maximum energy release rate in G Imax-decreas-ing tests. Though only two tests were carried out for eachlaminates under each stress ratio, these two tests gavealmost identical threshold values of the maximum energyrelease rate, G Imaxth. Thus, these obtained G Imaxth valuesare acceptable for crude evaluation. The data points withthe highest da/dN values for each specimen (marks withdash) indicate the data obtained just after the G Imax-decreasing tests were started (Da = 0). Then, the solid linesconnecting these points depict the da/dN G Imax relationswithout the effect of crack length.

G Imaxth

for the 50 lm-epoxy-interleaved laminates (28J/m2 for R = 0.1, 56 J/m2 for R = 0.5) were almost identi-cal with those for the base laminates (31 J/m2 for R = 0.1,53 J/m2 for R = 0.5). The fatigue crack growth resistance

Base laminates

50 m-epoxy-interleaved laminates

0

10

20

30

40

50

60

70

80

0 2 4 6 8 10

L o a d ,

P ( N )

Crack opening displacement, (mm)

Mode IUT500/111

NL(=Pmax)

Fig. 6. Relation between load and crack opening displacement for mode I

fracture toughness tests of base and 50 lm-epoxy-interleaved laminates.

0

0.2

0.4

0.6

0.8

0 10 20 30 40 50 60 M o d e I f r a c t u r e t o u g h n e s

s ,

G I c ,

G I R

( k J / m 2 )

Increment of crack length, a(mm)

Mode IUT500/111

Base laminates


Fig. 7. Relation between mode I fracture toughness and increment of crack length for base and 50 lm-epoxy-interleaved laminates.

Table 1Interlaminar fracture properties under mode I loading

Laminates Base 50 lm-epoxy-interleaved

G Ic (J/m2) 160 170

G IR (J/m2) 230 280

G Imaxth (J/m2)

R = 0.1 31 28

R = 0.5 53 56


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for the 50 lm-epoxy-interleaved laminates was slightlyhigher than that for the base laminates in the power-lawrelation region. The exponents of the power function forthe 50 lm-epoxy-interleaved laminates (4.7 for R = 0.1,6.0 for R = 0.5) were almost the same as those for the baselaminates (4.6 for R = 0.1, 5.0 for R = 0.5). Table 1 alsosummarizes the results of mode I delamination fatiguecrack growth tests.

Stress-ratio effect parameter, c [26], was calculated foreach crack propagation rate, da/dN , in order to discussthe stress ratio dependency on fatigue crack growth(Fig. 9). The c values for the base laminates were around0.6 without respect to da/dN . Though the c values for the50 lm-epoxy-interleaved laminates slightly increases withincrease in da/dN , the difference in the c values betweenthese laminates is small. The rather higher values of c indi-

cate that the contribution of maximum stress is relatively

large in the mode I delamination fatigue for bothlaminates.

3.3. Consideration of fracture mechanism under mode I

loading

Fig. 10 shows the scanning electron micrographs(SEMs) of the fracture surfaces under mode I loading forboth laminates. Fig. 10(a-1) and (a-2) compare the static

fracture surfaces. Here, the region of stable crack growthwas selected for the observation of the fracture surface of the 50 lm-epoxy interleaved laminates, Fig. 10(a-2). Forthe base laminates (Fig. 10(a-1)), a large number of carbonfibers and their pullout traces were observed. The arealratio of the interfacial fracture was about 30%. On theother hand, about 90% of the fracture surface was coveredwith resin in the case of 50 lm-epoxy-interleaved laminates(Fig. 10(a-2)). The morphology of the resin part shows brit-tle fracture with little trace of the plastic deformation.Fig. 10(b-1) and (b-2) show the fracture surfaces under fati-gue loading. Since no significant effect of the stress ratio onthe fracture surfaces was observed under fatigue loading,only surfaces for R = 0.5 are presented. Moreover, theeffect of da/dN was also minimal. While the interfacial frac-ture was dominant for the base laminates (Fig. 10(b-1)), theresin fracture was dominant for the 50 lm-epoxy-inter-leaved laminates (Fig. 10(b-2)). Plastic deformation of thematrix epoxy was hardly observed on all of the fatigue frac-ture surfaces in both laminates. Thus, the morphology of the fatigue fracture surfaces were similar to those of thestatic fracture ones. This similarity is supported by thehigher c values which indicate that the contribution of the maximum load is high in the mechanism of fatiguecrack growth. The only noticeable difference is that while

a rough surface was observed on the static fracture surface,

10-11

10-10

10-9

10-8

10-7

C r a c k p r o p a g a t i o n r a t e ,

d a / d N ( m / c y c l e )

Maximum energy release rate, GImax

(J/m2)

Base laminates Mode I

1003020 40 2007050

5.04.7

R

1st GImax

-decreasing test

2nd GImax

-decreasing test

0.5 0.1

10-11

10-10

10-9

10-8

10-7

C

r a c k p r o p a g a t i o n r a t e ,

d a / d N ( m / c y c l e

)

Maximum energy release rate, GImax

(J/m2)

50 m-epoxy-interleaved laminates Mode I

1003020 40 2007050

6.0

4.6

R

1st GImax

-decreasing test

2nd GImax

-decreasing test

0.5 0.1

- - - - - Base laminates

4.75.0

(a)

(b)

Fig. 8. Relation between crack propagation rate and maximum energyrelease rate under mode I loading for base and 50 lm-epoxy-interleavedlaminates.

50 m-epoxy-interleaved laminatesBase laminates

0

0.2

0.4

0.6

0.8

1

Crack propagation rate, da/dN (m/cycle)

10-9

10-7

UT500/111 Mode I

10-8

Fig. 9. Change of stress-ratio effect parameter at various crack propaga-tion rate for mode I fatigue tests of base and 50 lm-epoxy-interleaved

laminates.


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the fatigue fracture surface is rather smooth for the 50 lm-epoxy-interleaved laminates.

The 50 lm-epoxy-interleaved laminates did not showsignificant improvement of mode I interlaminar propertieswith the increase of the resin-rich layer thickness. The SEMobservations indicate that the fracture mechanism of the50 lm-epoxy-interleaved laminates was controlled by thebrittle fracture of matrix epoxy. Thus, the thicker resin-richlayer did not contribute directly to the increased propertiesbecause the damaged zone size was expected to be rathersmall. However, the G IR values and the fatigue crackgrowth resistance in the high crack propagation rate regionfor the 50 lm-epoxy-interleaved laminates were a littlehigher than those for the base laminates. One of the possi-

ble explanations for these increases is that the damage zonewas slightly large even though thicker resin-rich layer did

not contribute fully to the increase of the damage zone(Fig. 11). These conclusive results indicate that the tough-ness, and not the thickness, of the interleaved materials isthe key factor to the increase in mode I interlaminar prop-erties both under static and fatigue loading [30].

3.4. Mode II interlaminar fracture toughness tests

The load versus load–line displacement diagram in themode II interlaminar fracture toughness test is shown inFig. 12. The arrows in Fig. 12 indicate NL points and5% points. Nonlinear behaviors were observed before the

Fig. 10. Scanning electron micrographs of fracture surfaces under mode I loading. (a-1) Base laminates, static fracture. (a-2) 50 lm-epoxy-interleavedlaminates, static fracture. (b-1) Base laminates, fatigue fracture, da/dN = 1 · 1010 m/cycle. (b-2) 50 lm-epoxy-interleaved laminates, fatigue fracture,da/dN = 1 · 1010 m/cycle.

(a) (b)

Carbon fiber

Epoxy Crack path

Damage zoneCarbon fiber

Epoxy Crack path

Carbon fiber

Epoxy

Carbon fiber

Epoxy Crack path

Damage zone

Fig. 11. Schematic explanation of crack growth behavior under mode I loading. (a) Base laminates. (b) 50 lm-epoxy-interleaved laminates.


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load reached peak values for either of the laminates. Afterthe peak values, the load increased slightly and the crackgrowth behavior became stable.

Fig. 13 shows the relation between the mode II fracturetoughness and the increment of the crack length for bothlaminates. Solid symbols indicate the results for the baselaminates, and open symbols indicate those for the50 lm-epoxy-interleaved laminates. The data points atDa = 0 indicate the initial values of the fracture toughness,G IIc, at the NL points. The propagation values of the frac-

ture toughness, G IIR, for the base laminates graduallyincreased with the increment of the crack length, and thenleveled off where Da was larger than 3 or 5 mm. On the con-trary, G IIR values for the 50 lm-epoxy-interleaved lami-nates increased sharply at the initial stage of the crack

growth, and then leveled off where Da was larger than10 mm. Thus, the increase of the fracture toughness fromG IIc to G IIR for the 50 lm-epoxy-interleaved laminates(740–2100 J/m2) is much larger than that for the base lam-inates (470–610 J/m2). It is interesting to note that theincrease in fracture toughness achieved by inserting the

epoxy interleaf is much larger for G IIR (3.4 times) thanfor G IIc (1.6 times). This fact also agrees with the resultsof our separate research on the toughened interleaf [14].Thus, although the G IIR values for the 50 lm-epoxy-inter-leaved laminates are comparable to those for the tough-ened laminates such as AS4/PEEK and T800H/3900-2,the G IIc values are smaller than those for the toughenedlaminates [14,29]. Table 2 summarizes the results of modeII fracture toughness tests for both laminates.

3.5. Mode II delamination fatigue test

Fig. 14 shows the relation between the crack propaga-

tion rate, da/dN , and the maximum energy release rate,G IImax, for both laminates. Solid and open symbols indicate

Base laminates


0

500

1000

1500

2000

0 1 2 3 4 5

L o a d ,

P ( N )

Load-line displacement, v(mm)

Mode IIUT500/111

NL

5%

Fig. 12. Relation between load and load–line displacement for mode II

fracture toughness test of base and 50 lm-epoxy-interleaved laminates.


Base laminates

0

0.5

1

1.5

2

2.5

0 5 10 15 20 25 30 M o d e I I f r a c t u r e t o u g h n e s s ,

G I I c

, G

I I R

( k J / m 2 )

Increment of crack length, a (mm)

UT500/111 Mode II

Fig. 13. Relation between mode II fracture toughness and increment of

crack length for base and 50 lm-epoxy-interleaved laminates.

Table 2Effect of interleaf under mode II loading for UT500/111

Laminates Base 50 lm-epoxy-interleaved

50 lm-epoxy-interleaved/base

Interlayer thickness (lm) 15 50 3.3G IIc (J/m

2) 470 740 1.6

G IIR (J/m2) 610 2060 3.4

G IImaxth (J/m2)

R = 0.1 70 140 2.0

R = 0.5 170 390 2.3

Base 50 m-epoxy-interleaved

R=0.5

R=0.1

10-11

10-10

10-9

10-8

10-7

C r a c k p r o p a g a t i o n r a t e

, d a / d N ( m / c y c l e )

Maximum energy release rate, Gmax

(J/m2)

50 70 100 200 300 500

25152013

UT500/111 Mode II

Fig. 14. Relation between crack propagation rate and maximum energyrelease rate under mode II loading for base and 50 lm-epoxy-interleaved

laminates.


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the data for the base and for the 50 lm-epoxy-interleavedlaminates, respectively. The threshold values of the maxi-mum energy release rate, G IImaxth, for the 50 lm-epoxy-interleaved laminates (140 J/m2 for R = 0.1, 390 J/m2 forR = 0.5) were 2.0–2.3 times larger than those of the baselaminates (70 J/m2 for R = 0.1, 170 J/m2 for R = 0.5).

Also, it is interesting to note that these increases under fati-gue loading is larger than those for T800H/3900-2 asshown in Table 3. The exponents of the power functionfor the 50 lm-epoxy-interleaved laminates (20 forR = 0.1, 25 for R = 0.5) were slightly higher than thosefor the base laminates (13 for R = 0.1, 15 for R = 0.5).The threshold values for the 50 lm-epoxy-interleaved lam-inates are almost identical to those for the toughened lam-inates such as AS4/PEEK and T800H/3900-2 [14,29].Table 2 also summarizes the mode II delamination fatiguethreshold values.

Previously reported experimental results showed that

the relative increase in fatigue crack growth resistancewas lower than that of the fracture toughness for interlayertoughened laminates such as T800H/3900-2 [14,31]. TheG IImaxth/G IIc values under R = 0.5 for the base laminates,the 50 lm-epoxy-interleaved laminates and the T800H/3900-2 were 0.36, 0.52 and 0.33, respectively. It is interest-ing to note that the G IImaxth/G IIc value was highest for the50 lm-epoxy-interleaved laminates. The stress-ratio effectparameters, c, for both laminates were relatively low asindicated in Fig. 15, indicating that the contribution of DK is large in the fatigue fracture behavior under modeII loading for the base and also for the 50 lm-epoxy-inter-leaved laminates.

3.6. Consideration of fracture mechanism under mode II

loading

Figs. 16(a-1) and (a-2) show the SEMs of the static frac-ture surfaces of the base laminates and those of the 50 lm-epoxy-interleaved laminates, respectively. Typical hacklemarkings were observed for both laminates. The heightand interval of the hackle markings for the 50 lm-epoxy-interleaved laminates are much larger than those for thebase laminates. The intervals of the hackle marking forthe base and the 50 lm-epoxy-interleaved laminates were

about 2–5 and 10 lm, respectively. Moreover, the plastic

deformation of the hackle markings for the 50 lm-epoxy-interleaved laminates was much larger than those for thebase laminates.

Figs. 16(b-1) and (b-2) show the SEMs of the fatiguefracture surfaces for the base laminates and the 50 lm-epoxy-interleaved laminates, respectively. The areal ratioof the resin fracture was larger than that observed understatic loading. The hackle markings were observed on thefatigue fracture surfaces of both laminates as well as onthe static fracture surfaces. However, the roughness and

the height of the hackle markings were much smaller thanthose of the static fracture surfaces for both laminates. Thisdifference in roughness can be supported by the relativelylow c values, indicating that the contribution of DK is largein the fracture behavior under mode II fatigue loading forboth laminates. This can be explained as the main fracturemechanism by the reversed shear plastic deformation[14,31]. The size of the hackle markings for the 50 lm-epoxy-interleaved laminates was much larger than thatfor the base laminates. This indicates that the damagedzone for the 50 lm-epoxy-interleaved laminates was muchlarger than that for the base laminates. Even though theshear plastic deformation of epoxy resin is large, the finalfracture is often normal to the principal tensile stress [32].This is why we observed hackles although the reversedshear plastic deformation is the principal fracture mecha-nism as indicated by low c values.

Contrary to the results under mode I loading, the50 lm-epoxy-interleaved laminates showed higher fracturetoughness than the base laminates under mode II loading.The plastic deformation of the hackle markings for the50 lm-epoxy-interleaved laminates was much larger thanthat for the base laminates as shown in Fig. 16(a). This lar-ger deformation is the main reason for the higher fracturetoughness of the 50 lm-epoxy-interleaved laminates. It is

interesting to note that the increase in G IIR value from

Table 3Effect of heterogeneous interlayer under mode II loading for T800H/3900-2

Laminates T800H/3631 T800H/3900-2 (T800H/3900-2)/(T800H/3631)

Interlayerthickness (lm)

7–15 30 2–4

G IIc (J/m2) 500 1200 2.4G IIR (J/m

2) 540 2100 3.9G IImaxth (J/m

2)

R = 0.1 80 120 1.5

R = 0.5 250 400 1.6

50 m-epoxy-interleaved laminatesBase laminates

0

0.2

0.4

0.6

0.8

1

Crack propagation rate, da/dN (m/cycle)10

-910

-7

UT500/111 Mode II

10-8

Fig. 15. Change of stress-ratio effect parameter at various crack propa-

gation rate for mode II delamination fatigue test of base and 50 lm-epoxy-interleaved laminates.


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the base laminates to the 50 lm-epoxy-interleaved lami-nates (about 3.4 times) agrees with the increase in theresin-rich layer thickness (about 3.3 times) (see Table 2).This was suggested by Partridge that the increase of thefracture toughness was linearly correlated to the interleaf thickness [18]. Since fibers and traces of pullout fibers wereoften observed for the fracture surfaces under static load-ing, the hackle markings were located throughout theresin-rich layer. Thus, the size of the hackle markingsand the size of the damage zone should be in proportion

to the resin-rich layer thickness (Fig. 17(a)). This is the rea-

son why the increase in G IIR value from the base laminatesto the 50 lm-epoxy-interleaved laminates agrees with theincrease in resin-rich layer thickness. On the other hand,the increase in G IIc value from the base laminates to the50 lm-epoxy-interleaved laminates (about 1.6 times) wassmaller than the corresponding increase in resin-rich layerthickness (about 3.3 times). The mechanism of the crackgrowth behavior at the initial stage of the crack propaga-tion under mode II loading is still unknown. However, itis clear that the onset and the growth of microcracks play

a major role in determining the fracture behavior. This is

Fig. 16. Scanning electron micrographs of fracture surfaces under mode II loading. (a-1) Base laminates, static fracture. (a-2) 50 lm-epoxy-interleavedlaminates, static fracture. (b-1) Base laminates, fatigue fracture, da/dN = 1 · 107 m/cycle. (b-2) 50 lm-epoxy-interleaved laminates, fatigue fracture,

da/dN = 1 ·

10

7

m/cycle.


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probably why the increase in resin-rich layer thickness wasnot directly translated into increase in G IIc value.

In the mode II delamination fatigue tests, the 50 lm-epoxy-interleaved laminates showed higher fatigue crackgrowth resistance than the base laminates. The largerhackle markings for the 50 lm-epoxy-interleaved lami-

nates indicate larger damaged zone under fatigue loading.However, the hackle markings and the damage zone werenot expanded throughout the resin-rich layer (Fig. 17(b)).This is why increase in resin-rich layer thickness was notdirectly translated into increase in G IImaxth value. Compar-ison between Tables 2 and 3, and also between Table 2 andthe result of Ref. [18] indicates that while increase in fati-gue threshold (about two times) for brittle #111 self-sameepoxy interleaf is larger than that for toughened epoxy,increase in fracture toughness is smaller than that fortoughened epoxy. Thus, the toughness of the matrix resinhas no direct contribution to higher fatigue threshold. The

conclusive results are that it is possible to increase themode II fracture toughness and fatigue threshold only byinserting the self-same epoxy interleaf without toughening.

4. Conclusions

The effect of the self-same epoxy interleaf on the inter-laminar fracture properties of unidirectional CF/conven-tional epoxy laminates were investigated for modes I andII under both static and fatigue loadings. The results weresummarized as follows. The mode I interlaminar fracturetoughness values for the 50 lm-epoxy-interleaved lami-

nates were almost identical to those for the base laminates.

The fatigue threshold values for the 50 lm-epoxy-inter-leaved laminates were again almost the same as those forthe base laminates. The analysis of the stress ratio depen-dency indicated that the contribution of the maximum loadwas rather large in the mechanism of fatigue delamination.The fracture surface morphology indicated rather brittle

fracture under both static and fatigue loadings.Whereas the initial values of the mode II interlaminar

fracture toughness for the 50 lm-epoxy-interleaved lami-nates were 1.6 times higher than those for the base laminates,the propagation values for the 50 lm-epoxy-interleavedlaminates were 3.4 times higher than those for the baselaminates. The fatigue threshold values for the 50 lm-epoxy-interleaved laminates were 2.0 times (R = 0.1) and2.3 times (R = 0.5) higher than those for the base lami-nates. The analysis of the stress ratio dependency indicatedthat the contribution of the cyclic load was rather large inthe mechanism of fatigue delamination. The fracture mor-

phology indicated that the damaged zone size for the50 lm-epoxy-interleaved laminates was much larger thanthat for the base laminates.

The contribution of the self-same epoxy interleaf is com-pletely different for modes I and II loadings. Self-samematrix interleaf with conventional brittle epoxy canincrease mode II interlaminar properties comparable toadvanced systems with toughened interlayer or thermo-plastic matrix.

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PrincipalstressPrincipalstress

(a-1)

(b-1) (b-2)

(a-2)

Carbon fiber Epoxy

Microcrack (Tensile fracture)

Carbon fiber Epoxy

Microcrack (Tensile fracture)

MicrocrackMicrocrack

Damage zoneDamage zone

Fig. 17. Schematic explanation of crack growth behavior under mode II loading. (a-1) Static fracture, base laminates. (a-2) Static fracture, 50 lm-epoxy-interleaved laminates. (b-1) Fatigue fracture, base laminates. (b-2) Fatigue fracture, 50 lm-epoxy-interleaved laminates.


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