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Matrix effect on the static and dynamic interlaminar fracture toughness of glass-fibre marine composites
Composites Part B 29B (1998) 505-516
PIh S1359-8368(98)00004-3
ELSEVIER
© 1998 Elsevier Science Limited Printed in Great Britain. All rights reserved 1359-8368/98/$19.00
Matrix effect on the static and dynamic interlaminar fracture toughness of glass-fibre marine composites
P. C o m p s t o n a'*, P.-Y.B. Jar a and P. Davies b aDepartment of Engineering, The Australian National University, Canberra 0200, Australia bMarine Materials Laboratory, IFREMER- Centre de Brest, BP 70, 29280 Plouzane, France (Received 20 November 1996; accepted I September 1997)
The interlaminar fracture toughness of four marine composites has been investigated. The matrix material was different in each composite. Mode I and Mode II tests were performed under static loading and the focus was placed on the transfer of matrix toughness to the composite. A coefficient of toughness transfer was used to quantify the synergy experienced in the Mode I results. Mode II tests were also performed under dynamic loading with the focus placed on the order of results in relation to the order of matrix toughness. The study concludes that synergy in Mode I results for crack initiation and steady-state propagation is greatest in the composites with superior interracial bonding. The Mode I results do not correspond directly to matrix toughness owing to interracial strength limitations, relative to matrix strength, in the composites with the toughest matrices. A similar pattern is evident in the results for crack initiation in static mode II tests. At the maximum load point in static Mode II tests, where fast fracture occurred, and in dynamic Mode II tests, the results are consistent with the order of matrix toughness. For these results, differences in interfacial bonding seem to have an insignificant effect. © 1998 Elsevier Science Limited. All rights reserved
(Keywords: interlaminar fracture; matrix toughness; glass-fibre marine composites)
INTRODUCTION There are many examples of the increasing use of fibrereinforced polymer-matrix composites for structural applications in marine environments. ~-7 While composites can offer high strength and stiffness in conjunction with weight reduction, polymer matrices in particular improve sonic characteristics, corrosion resistance, and aesthetic qualities. For economic reasons, glass fibre is the preferred reinforcement in marine composites. Therefore, cost effective improvement of any composite property will be achieved through the use of matrix materials with improved properties. Polyester resins have been the most common matrix material in the past, but attention is now focused on replacing them with resins such as epoxy and vinyl ester. The vinyl ester is of particular interest owing to its ageing resistance, s As marine composites are susceptible to delamination damage, 9 the effect of different matrices on the interlaminar fracture toughness requires investigation, t° Static tests are useful for characterisation of mechanical properties, however they do not completely reflect practical situations. * Corresponding author.
Delamination damage in marine composites can be caused by low-speed impact, ~1 and this can adversely affect the residual strength and stiffness in the composite. 12 Composite interlaminar fracture toughness under Mode I and Mode II static loading is characterised in terms of critical strain energy release rate, Go. 13-15 Many different composite systems have been characterised in this way, as indicated by Refs 16A7. It has been shown that matrix toughness, fibre-matrix bond strength and, especially in Mode I, fibre bridging 18 can all influence the value for composite Gc. Hibbs e t al. 19 showed that increased matrix toughness will increase composite interlaminar fracture toughness. However, along with other published work, 2°'21 Hibbs e t al. also emphasised that good interfacial bonding is necessary to promote cohesive failure through the matrix and, thus, ensure that matrix toughness is transferred to the composite. In a systematic study of the interface effect, Albertsen e t al. 22 showed that as adhesion between the fibre and the matrix increases, Gc for crack initiation in both Mode I and Mode II also increases. For a carbon-fibre/epoxy composite, they also found that an optimum level of adhesion between the fibres and the matrix, which promotes a combination of cohesive and interfacial failure, will
505
Fracture toughness of glass-fibre marine composites: P. Compston et al. maximise the Gc value for Mode I steady-state crack propagation. The interfacial failure is necessary for the formation of bridging fibres, failure of which is the major energy absorbing mechanism. As discussed by Bradley, 23 much of the published work from Mode I testing of brittle-matrix composites shows a synergistic effect, whereby composite Gc is higher than matrix Go. This was attributed to increased fracture area created in the matrix by crack propagation, the effect of fibre bridging and the slower rate of decay of the stress field ahead of the crack tip. Strong interfacial bonding would also contribute to the synergistic effect by promoting cohesive failure. Research on interlaminar fracture toughness under impact loading is less extensive and test methods are not standardised. However, Maikuma et al. 24 successfully induced Mode II dominated delamination in carbon-fibre composites using a drop weight impact tester. Two composites were tested, one with an epoxy matrix and the other with a tough PEEK matrix. The results showed interlaminar fracture toughness under low-speed impact to be dominated by matrix toughness. This paper reports on a study to characterise the interlaminar fracture toughness of marine composites. They were subjected to Mode I and Mode II static loading, and Mode II low-speed impact loading. The only variant among the composites was the matrix material. Discussion of the results from the static tests will focus on the transfer of matrix toughness to composite interlaminar fracture toughness. Under impact loading, the focus will be placed on the order of the results in relation to the order of matrix toughness.
Table 1 Matrix resin details Matrix resin
Supplier
Trade name
Epoxy (EP) Vinyl ester (VE) Isophthalic polyester (PI) Orthophthalic polyester (PO)
Sicomin DSM Scott Bader Scott Bader
SRI500 Atlac 580 Crystic 491 PA Crystic R115 PA
Table 2
Fibre volume fraction in UD region
Composite
Fibre volume fraction in UD region (%)
Coefficient of variation (%)
GF/EP GF/VE GF/PI GF/PO
40 46 39 39
6 5 4 7
water-cooled diamond cutting wheel. They were then dried under vacuum at room temperature for 24 h. Using the method proposed by Waterbury and Drzal, 25 the fibre volume fraction in the UD region of each composite was determined by calculating fibre area fraction from images of cross-sectioned samples. Each cross-section was 2 0 mm wide and 6 mm thick. Images of the UD region were obtained using a scanning electron microscope. The magnification was such that the total thickness of the UD region was contained within each image. This equated to 2000-3000 fibres per image. Using the public domain NIH Image program (version 1.57), the fibre area fraction, and hence fibre volume fraction, were calculated. For each composite, two images were taken from three samples. Therefore, each value for fibre volume fraction given in Table 2 is an average of six results. It should be noted that the fibre volume fraction is significantly higher for GF/VE than for the other composites.
MATERIALS E-glass fibres with a multi-compatible sizing, made by Vetrotex, were used for this study. Four thermoset resins were chosen as the matrix materials. They are detailed in Table 1. According to Vetrotex, the sizing (P177) is compatible for bonding with each resin. Table 1 also introduces a reference for each matrix which will be used in an abbreviated reference for each composite throughout the rest of the text. For example, the glass-fibre reinforced composite with the epoxy matrix will be referred to as the GF/EP composite. Four 12-ply composite laminates, one for each matrix resin, were produced by hand lay-up. The central four plies of each laminate contained unidirectional (UD) fibres of 250 g/m 2. The outer layers consisted of 0/90 ° stitched fabric plies of 590g/m z. To introduce a starter crack, a polypropylene film, 8/zm thick and 75 mm long, was placed between the central plies in the UD region. As recommended by the resin suppliers, the laminates with the polyester matrices were stabilised at room temperature for 1 month before testing and the laminates with the epoxy and vinyl ester matrices were post-cured at 90°C for 6 h. Test specimens were cut from these laminates using a
506
EXPERIMENTAL PROCEDURES The mechanical properties of each matrix resin were established through fracture toughness and tensile testing. The fracture toughness tests were conducted in accordance with the draft international standard ISO/DIS 13586(E) for plastics--determination of fracture toughness (Gc and Kc)--linear elastic fracture mechanics (LEFM) approach, 1996. Mode I double cantilever beam (DCB) and Mode II end notch flexure (ENF) tests were used to establish composite delamination resistance under static loading. These tests were conducted in accordance with the European Structural Integrity Society (ESIS) Protocols for Interlaminar Fracture Testing of Composites. 26 Mode II central notch flexure (CNF) tests were used to determine delamination resistance under dynamic loading. For this test, the method described by Maikuma et al. 24 was followed. In each test, the delamination crack was propagated parallel to the fibre direction in the UD region of the test specimen, and a minimum of five specimens from each composite were used. The specimens were not pre-cracked.
Fracture toughness of glass-fibre marine composites: P. Compston et al. RESIN FRACTURE TOUGHNESS AND TENSILE TESTING
in each test. Owing to material availability constraints, only one or two specimens could be tested for each resin.
The geometry of the single edge notched bend (SENB) specimen used for the resin fracture toughness, Ko tests is shown in Figure 1. Specimen dimensions were width, W, 14 mm, nominal thickness, B, 4 mm, and span distance, 2L, 56 mm. A natural crack was introduced by tapping a new razor blade placed in a machined notch, to produce a crack length in the range 0.45 < a/W < 0.55. Specimens were loaded in a three-point bend fixture at 10 mm/min. The load-displacement plots recorded were linear elastic and the maximum load Pmax w a s used to determine Kc from the equation: f(emax) g c -- - I/2(BW)
(1)
The calibration factor f is given in the draft ISO standard. Tests were valid according to the specimen size criterion given in that document. Between four and eight valid tests were performed for each material. Scatter in Kc was typically + 10%. From the Kc results, the Mode I critical strain energy release rate for each matrix resin, Gic_matrix, w a s calculated as f o l l o w s : 27
K2 Glc = ~-(1 - p 2)
(2)
where E is tensile modulus and v Poisson's ratio. Young's modulus, maximum strength and maximum strain were established using miniature tensile test specimens. They had a width of 10 mm, thickness 5 mm and gauge length 25 mm. Testing was conducted in displacement control on an Instron 4505 Universal Testing Machine (UTM) at a crosshead speed of 1 mm/min. An extensometer was used
Mode I DCB testing
The DCB specimen geometry used for static Mode I testing is shown in Figure 2. The thickness was approximately 6 mm, width 20 mm and length 160 mm. The latter dimension allowed for 70 mm of crack growth. For load introduction, aluminium blocks were attached to the end containing the crack starter film. A thin layer of white correction fluid was applied to one edge of the specimen, starting at the tip of the crack starter film. This edge was marked at 1 mm intervals for the first 5 mm and then at 5 mm intervals up to the 70 mm mark. This allowed the crack to be identified and measured as the test progressed. Testing was conducted in displacement control on the Instron UTM at a crosshead speed of 2 mm/min. A load-displacement plot was produced for each test, and load-displacement data were recorded for each individual crack length measurement. The Mode I critical strain energy release rate, Gic, was calculated using the corrected beam theory method, z8 It is derived from the Irwin and Kies expression for fracture energy: 29 p2 d C
Gc -- 2B da
(3)
where P is load, B specimen width, C compliance and a crack length. This expression accounts for the change in compliance with crack length. Using the simple beam theory expression for compliance: 6 2a 3 C. . . . P 3EI
(4)
where E is flexural modulus and I moment of inertia, eqn (3) may be rewritten as: 3P6 Glc = 2Ba
(3. / I~
-I
ZL
Figure 1 Single edge notched bend (SENB) test specimen geometry
where 6 is crosshead displacement. However, a correction factor IAL needs to be added to crack length to allow for crack tip rotation at the root of the cantilever beam. IAI is the x-axis intercept when the cube root of compliance, C m, is plotted as a function of crack length, a. Therefore, the corrected beam theory expression for Glc is: 3P3 GIc -- 2B(a + IAI)
crack starer
ao
I I I I I
(5)
(6)
film
LID region
Figure 2 Mode I DCB test specimen geometry with initial crack length a o
Values for GIc will be plotted as a function of crack length to produce a resistance (R) curve. Results will be presented for crack initiation, Glc-init, defined as the first deviation from linearity on the loaddisplacement plot, and steady state crack propagation, Gtc-prop, defined by a plateau on the R curve. The extent to which matrix toughness has been transferred to interlaminar fracture toughness in each composite, and the degree of synergy experienced for the
507
Fracture toughness of g/ass-fibre marine composites: P. Compston et al. Results will be presented for crack initiation, Gllc.init, again defined as the first deviation from linearity on the loaddisplacement plot, and for the maximum load point, Gllc-ma×. I
Mode H CNF impact testing 2L
The geometry of the CNF specimen used for the Mode II impact tests is shown in Figure 4. The thickness was 6 mm, width 20 mm and length 180 mm. The crack starter film is located in the centre of the specimen length. The CNF specimen was placed in a three-point bending fixture. Span length was set at 120 mm so that the specimens were sufficiently rigid to prevent excessive deflection during the impact. The tests were conducted using a Mitsubishi Yuka Co. instrumented drop weight impact tester. The striker width was 25 mm, sufficient to cover the specimen width and provide a contact condition similar to that in the Mode II ENF test fixture. The striker weight was constant at 2.6 kg and impact speed was varied between 1.5 and 3 rn/s in order to obtain a range of impact energies. The force during the impact was recorded by deformation of a strain gauge attached to the striker, and sampled every 20 #s. A forcetime plot was produced for each test and reduced into force and displacement data. The total absorbed energy was the area under the force-displacement curve. The total delamination fracture area in each specimen was measured after impact. A plot of absorbed energy versus delamination fracture area for a number of tests showed a linear relationship, therefore the average G~ value for each composite was taken as the slope of the line fitted to these points. A detailed description of how this test was conducted on these composites is available elsewhere. 31
I
Figure 3 Mode II ENF test specimen geometry
impact
Figure 4 Mode II CNF impact test specimen geometry
Glc.init and Glc.prop results,
will be expressed as follows:
GIc - composite ~
CTTGIc
-
matrix
(7)
where CTT is the coefficient of toughness transfer.
Mode H ENF testing The ENF specimen geometry used for static Mode II testing is shown in Figure 3. The thickness was approximately 6 m m , width 2 0 m m and length 160mm. The specimen was placed in a three-point bending fixture with half span length, L, set at 50 mm and the ratio of crack length to half span length, a/L, maintained at 0.5. One edge of the specimen was covered with white correction fluid, allowing the crack to be identified and measured at the end of the test. This test was also conducted in displacement control on the Instron UTM at a crosshead speed of 1 mm/min. It is stable until a maximum load point is reached, after which unstable fracture occurs. A load-displacement plot provided the raw data for each test. The Mode II critical strain energy release rate, Gtic, w a s calculated using direct beam theory. It is also based upon eqn (3). Using the ENF compliance expression given by Russell and Street: 3°
Scanning electron microscopy Fracture surface micrographs were obtained using a Cambridge $360 scanning electron microscope (SEM). Specimens were coated with a thin layer of gold prior to examination. The micrographs were used to characterise the dominant failure modes in each composite for each test. All micrographs presented in this paper show regions that are close to the crack initiation point.
2L 3 -t- 3a 3
C=
RESULTS AND DISCUSSION
8EBh3
(8)
where h is specimen thickness, eqn (3) may be rewritten as:
9a2P6 GIIc - 2B(2L3 + 3a3) Table 3
Resin properties The tensile properties of each matrix resin are given in Table 3. While the modulus is similar for each resin, the
(9)
Tensile and fracture toughness, Kc, properties of the matrix resins
Matrix resin
Young' s modulus (GPa)
Maximum strength (MPa)
Maximum strain (%)
Fracture toughness, Kc (MPa ,v/~)
Epoxy (EP) Vinyl ester (VE) Isophthalic polyester (PI) Orthophthalic polyester (PO)
3.4 3.5 3.5 3.7
70 85 68 58
5.4 3.2 2.9 2.1
1.125 0.694 0.451 0.419
508
Fracture toughness of glass-fibre marine composites: P. Compston et al. et al., 32 which showed that 50 mm of crack growth was required before a plateau region develops. Therefore, in the present study, Glc values for 50-70 mm of crack growth were averaged to give Glc-prop. The results for Gic_init are shown in Figure 7. They can be placed on two distinct levels. The GF/EP and GF/PI composites are on the higher level, and the GF/VE and GF/PO composites comprise the lower level. The results for Gic.prop a r e shown in Figure 8. The GF/PI and GF/VE composites have the highest values followed by GF/EP then GF/PO. The GF/VE composite, however, has a higher fibre volume fraction in the UD region (see Table 2) and has the potential to form more bridging fibres. Therefore, the value of Gk-prop for GF/VE cannot be compared with the other composites. Nevertheless, it is clear that the trend in the GIcinit and Glc_prop results is not consistent with the trend in matrix toughness as shown in Figure 5. In addition, the values of CTT for Gtc-init and Glc.prop in Table 4 show that greater synergy is achieved in the GF/PI and GF/PO composites. The CTT is discussed in terms of fibre-matrix adhesion, but other factors may also influence the toughness
values for maximum strength and maximum strain are higher for epoxy and vinyl ester. Therefore, the use of these two resins is consistent with the aim of improving composite properties through improved matrix properties. The K~ values, also given in Table 3, show epoxy to be the toughest matrix resin, followed in descending order by vinyl ester, isophthalic polyester and orthophthalic polyester. The Mode I critical strain energy release rate of each matrix resin, Glc-matrix, is depicted in Figure 5. The trend shown confirms the order of matrix toughness given by the K~ results, and will be compared with the trend in the results for composite Mode I interlaminar fracture toughness.
Mode I DCB fracture toughness Typical R curves for each composite are shown in Figure 6. The Glc_init value is the first value in the curve and corresponds to the onset of crack growth. In each R curve, the Glc values continue to rise until approximately 50 mm of crack growth is achieved. This is consistent with the R curves for glass-fibre composites produced by Kawada
0.35 0.3 0.25 0.2
0.15 ~,~
0.1 0.05
EP
VE
Figure 5 Mode I critical strain energy release rate for each matrix resin,
PI
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PO
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Crack growth (ram) Figure 6
Typical Mode I R curves for each composite: GF/EP (D); GF/VE (©); GF/PI (A); and GF/PO (V)
509
Fracture toughness of glass-fibre marine composites: P. Compston et al. 0.35 0.3 0.25 0.2 ,m ,~
~,~
0.15 0.1 0.05
GF/EP
GF/VE
GF/PI
GF/PO
Figure 7 Mode I DCB fracture toughness for crack initiation, alc.init. The error bar signifies _ 1 standard deviation
1.5
eL
1
O L
0.5
GF/EP Figure 8
GF/VE
Coefficient of toughness transfer, C/T, for Mode I results
Composite GF/EP GF/VE GF/PI GF/PO
C/T~initiation
CTT--propogation
( G lc.init/G lc-matrix)
( G lc_prol~G Ic-matrix)
0.9 1.3 5.1 3.9
2.9 11.5 a 25.6 18.4
~Comparitively high fibre volume fraction.
510
GF/PO
Mode I DCB fracture toughness for steady-state crack propagation, Glc-prop-The error bar signifies + 1 standard deviation
transfer, such as fibre constraint and shrinkage stresses. The contributions of such factors are not considered here. The fracture surfaces of the Mode I DCB specimens in Figure 9 show an important variation in the level of fibrematrix adhesion. The fibre surfaces in the GF/EP and GF/ VE composites are completely devoid of matrix, indicating adhesive failure along the interface. In contrast, the fibre surfaces in the GF/PI and GF/PO composites are partly covered with matrix, indicating a degree of cohesive failure through the matrix. Therefore, there is greater utilisation of matrix toughness in GF/PI and GF/PO, and this is reflected in the values for CTT. In the case of GF/EP, the value for CT/" suggests that the superior matrix toughness is negated Table 4
GF/PI
by interfacial strength limitations, relative to the matrix strength, resulting in a GIc-init which is similar to GF/PI. For the same reason, the value of GIc-init for GF/VE is similar to GF/PO. The difference in fibre-matrix adhesion between GF/PI and GF/EP has also been sufficient to produce a large difference in the value of Glc-prop. The increased cohesive failure in GF/PI has served to utilise matrix toughness and increase bridging fibre lifetime. Thus, the GIc-prop value for GF/PI is higher than GF/EP. However, GF/PO has the lowest GIc-prop value despite good fibre-matrix adhesion and the comparatively high CTT value. In this case, the failure of the improved bonding to increase Gic_prop is attributed to the inherent brittleness of the matrix. These propagation values are strongly influenced by fibre bridging, but may be important in areas where unidirectional reinforcements are used to reinforce critical areas of structures.
Mode H ENF fracture toughness The results for Gllc.init a r e shown in Figure 10. The GF/EP composite has the highest value, and this is consistent with
Fracture toughness of glass-fibre marine composites: P. Compston et al.
50/~m Figure 9
Fracture surfaces of Mode I DCB specimens: (a) GF/EP; (b) GF/VE; (c) GF/PI; and (d) GF/PO
0.4
0.3
°~
•-
0.2
= 0.1
GFfEP
GF/VE
GF/PI
GF/PO
Figure 10 Mode II ENF fracture toughness for crack initiation, Guc-init. The error bar signifies -+ 1 standard deviation matrix toughness. However, the value of Gllc.init for GF/VE is similar to GF/PI and GF/PO. The results for GIlc-max, shown in Figure 11, are consistent with matrix toughness. The fracture surfaces of the Mode II ENF specimens are shown in Figure 12. As seen for Mode I DCB specimens, the fibre surfaces in GF/EP and GF/VE are devoid of matrix whereas in GF/PI and GF/PO they are still partly covered with matrix. The cohesive failure in GF/PI and GF/PO and
the adhesive failure in GF/VE has resulted in a similar value of Gtlc_init for these three composites. However, the fact that GF/EP has the highest value suggests that fibre-matrix adhesion has been less influential at crack initiation in the Mode II test than in the Mode I test. Matrix deformation, in the form of hackle marks, can also be seen in Figure 12 and is typical for Mode II specimens. ~9 This observation is consistent with a matrix dominated failure process and supports the Gllc-max results.
511
Fracture toughness of glass-fibre marine composites: P. Compston et al.
3.5
2.5
2 1.5
0.5
GF/EP Figure 11
GF/VE
GF/PI
GFfPO
M o d e II E N F fracture toughness at m a x i m u m load point, GIIc-max.The error bar signifies _+ 1 standard deviation
50 #m Figure 12 Fracture surfaces o f M o d e II E N F specimens: (a) GF/EP; (b) GF/VE; (c) GF/PI; a n d (d) G F / P O
Mode H CNF impact fracture toughness The plots of absorbed energy versus delamination fracture area for each composite are shown in Figure 13. The G~ value is the slope of the line fitted to the points in these plots. The equations for each line are as follows:
512
GF/EP:
GF/VE: GF/PI: GF/PO:
y y y y
= 0.006985 + 2.0405x = 0 . 0 7 2 1 8 4 + 1.7657x = - 0.21507 -f- 1.5346x = 0.3754 + 1.4103x
R R R R
= = = =
0.95067 0.95932 0.8158 0.88041
where y is absorbed energy loss, x fracture area and R goodness-of-fit for each line (where R = 1 is a perfect fit).
Fracture toughness of glass-fibre marine composites: P. Compston et al. .
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Fracture area (103m 2) Figure
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Plots of absorbed energy Eah versus fracture area AA from Mode II CNF impact tests: (a) GF/EP; (b) GF/VE; (c) GF/PI; and (d) GF/PO
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GF/VE
GF/PI
GF/PO
Mode II CNF impact fracture toughness, G~, For comparison, Gu~ m,~ results from Mode II ENF tests are included
The G~ results, shown in Figure 14, correspond directly to matrix toughness. The fracture surfaces of the CNF specimens are shown in Figure 15. There is greater fibre-matrix adhesion in the GF/PI composites after fracture, which is consistent with Figures 9 and 12. Hackle mark deformation can be seen in each matrix,
consistent with Figure 12, with increased deformation apparent in the tougher EP and VE matrices. In conjunction with the GI~ values, observation of such matrix deformation characteristics supports the conclusion that energy absorption has been matrix dominated, and that the variations in fibre-matrix adhesion which
513
Fracture toughness of glass-fibre marine composites: P. Compston et al.
a
Figure 15 Fracture surfaces of Mode II CNF impact specimens: (a) GF/EF; (b) GF/VE; (c) GF/PI; and (d) GF/PO
influenced Glc.init, Gic_prop and Giic_init results have had a less significant effect. The Gllc-max values are also included in Figure 14 for comparison with G D. While the trend matches well, the absolute values for G D are lower than for Gn ..... even though both results represent energy absorbed in resisting fast fracture. This suggests a loading rate effect. The difference in specimen geometries for the quasi-static and dynamic tests may contribute to this difference. Published values for carbon-fibre reinforced composites indicate that the ENF and CNF specimens yield similar values. 24 To examine this further for the composites used here, a series of static tests was performed on CNF and ENF specimens taken from the same glass-fibre/epoxy laminate. The expression used to determine Glic was:
9PZa2
(10)
GIIc = 64EB2h 3 for a specimen of thickness 2h. The results were analysed using the same modulus values. Thus, five tests on CNF specimens yielded mean (standard deviation) GHc.... values of 3.72 (0.26) kJ/m 2, compared with 3.50 (0.35) from the ENF specimens. This suggests that the specimen geometry does not explain the rate effect. Further work is being carried out to verify the effect of loading rate on glassfibre reinforced composites.
514
CONCLUSIONS The interlaminar fracture toughness of four glass-fibre marine composites has been investigated. Mode I (DCB) and Mode II (ENF) tests were conducted under static loading. The focus of the discussion on results from these tests was placed on the transfer of matrix toughness to composite interlaminar fracture toughness. The synergy experienced in the Mode I results, expressed in terms of a coefficient of toughness transfer, CTT, was also discussed. In addition, Mode II (CNF) impact tests were conducted under low-speed impact loading, with the focus placed on the order of results in relation to the order of matrix toughness. The following conclusions are made: • In static Mode I tests, the alc.init results show that the transfer of matrix toughness and the synergistic effect, reflected in values of C/T, is greatest in the composites with the brittle polyester matrices. That is, in GF/PI and GF/PO. This is attributed to superior adhesion between the fibres and the matrix, which has promoted greater cohesive failure and utilisation of matrix toughness. The GF/EP composite, which has the toughest matrix, has the highest absolute value for Glc-init, but the lowest C / T value. This indicates, in conjunction with the interfacial failure observed on the fracture surface, that interfacial strength limitations have reduced the effect of superior
Fracture toughness of glass-fibre marine composites: P. Compston et al. matrix toughness. The same reason is given for the comparatively poor performance of the GF/VE composite which has the vinyl ester matrix. The highest value of GIc-prop, and corresponding CTT, was obtained for GF/PI. As for Glc.init, this is attributed to the superior fibre-matrix adhesion. It has enhanced the effect of fibre bridging and placed GF/PI closer than GF/EP to the optimum level of bonding required to maximise Gic-prop.The GF/PO composite also has a high CTT value, but the lowest value for G~c-prop,and this is attributed to the inherent brittleness of the matrix. The GF/VE composite was omitted from this comparison owing to its comparatively high fibre volume fraction. • In static Mode II tests, the value of Giic.init is highest for GF/EP. This is consistent with matrix toughness. However, the value for GF/VE is similar to GF/PI and GF/PO. This is attributed to the superior bonding and greater cohesive failure in GF/PI and GF/PO which is evident on the fracture surfaces. The GII..... results, which reflect energy absorbed in resisting fast fracture, are consistent with matrix toughness. The fracture surfaces show the areas of pure matrix resin to be highly deformed, indicating extensive participation of the matrix in the fracture process. • The G~ results from the CNF impact tests are also consistent with matrix toughness. Fracture surface examination revealed hackle mark deformation in the matrix. A greater amount of deformation was observed in the tougher matrices. This supports the conclusion that, for these composites, the fracture process under lowspeed impact is matrix dominated. • From a practical point of view, the use of a sizing which enhances compatibility of glass fibre with epoxy and vinyl ester resins is recommended. This would enhance the static interlaminar fracture toughness in glass-fibre marine composites which use these resins as the matrix material.
4.
5. 6.
7.
8.
9.
10.
11.
12.
13. 14.
15.
16.
17.
18.
19.
ACKNOWLEDGEMENTS This study was supported by the Australian Research Council (ARC) under the Large Grant Scheme. The technical support of staff in the ANU Electron Microscopy Unit is also appreciated.
20. 21.
22.
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Maccan, A. and Farolfi, F., Innovations and technical-economic aspects of structural and non-structural applications of composite materials on large passenger vessels and merchant ships. In Int. Conf. Nautical Construction With Composite Materials, eds P. Davies and L. Lemoine, IFREMER. Paris, 1992, pp. 39-47. Saman, A. B., Kader, A. B. D., Yaakob, O. and Samian, Y., Design and construction of fibreglass (GRP) assault boat. In Int. Conf. Nautical Construction With Composite Materials, eds P. Davies and L. Lemoine, IFREMER. Paris, 1992, pp. 67-74. Lemiere, Y., The evolution of composite materials in submarine
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structures. In Int. Conf. Nautical Construction With Composite Materials, eds P. Davies and L. Lemoine, IFREMER. Paris, 1992, pp. 441-449. Hihara, L. H., Bergman, R. and Takahashi, P. K., Marine applications for advanced composites. In Int. Conf. on Advanced Composites, eds T. Chandra and A. K. Dhingra. Wollongong, Australia, 1993, pp. 95-100. Leslie, J. C., Development of a composite submarine radome. ICCM-IO, Whistler, Vol. III, August 1995, pp. 669-676. Hollingsworth, J. P., Comparison of alternative laminating materials for the navy standard 22 foot utility boat. In Proc. of 6th Int. Conf. on Marine Applications of Composite Materials (MACM), Florida, March 1996. Wyman, D., Composites in the navy's new SDV. In Proc. of 6th Int. Conf. on Marine Applications of Composite Materials (MACM), Florida, March, 1996. Gutierrez, J., Le Ley, F. and Hoarau, P., A study of the aging of glass fibre resin composites in a marine environment. In Int. Conf. Nautical Construction with Composite Materials, eds P. Davies and L. Lemoine. Paris, 1992, pp. 338-346. Barnoum, B. and Renaud, P., Damage observed during inspection of composite pleasure boats. In Int. Conf. Nautical Construction With Composite Materials, eds P. Davies and L. Lemoine, IFREMER. Paris, 1992, pp. 389-393. Davies, P., Fracture of marine composites. In Fracture of Composites, ed. E. Armaois. Trans Tech Publications, 1996, pp. 583-596. Davis, G. T., In-service failures from the marine surveyor's point of view: an update. In Proc. of 6th Int. Conf. on Marine Applications of Composite Materials (MACM), Florida, March, 1996. Benzeggagh, M. L. and Benmedakhene, S., Residual strength of glass/polypropylene composite material subjected to impact. Composite Sci. Technol., 1995, 55, 1-11. Griffiths, A. A., The phenomena of rupture and flow in solids. Phil. Trans. R. Soc. Lond., 1921, A-221, 163. Davies, P. and Benzeggagh, M. L., Interlaminar Mode-I fracture testing. In Application of Fracture Mechanics to Composite Materials, ed. K. Friedrich. Elsevier, Amsterdam, 1989, pp. 81112. Carisson, L. A. and Gillespie, J. W., Mode-II interlaminar fracture of composites. In Application of Fracture Mechanics to Composite Materials, ed. K. Friedrich. Elsevier, Amsterdam, 1989, pp. 113157. Sela, N. and Ishai, O., Interlaminar fracture toughness and toughening of laminated composite materials: a review. Composites, 1989, 20(5), 423-435. O'Brien, T. K. and Martin, R. H., Round robin testing for Mode I interlaminar fracture toughness of composite materials. J. Composite Technol. Res., 1993, 15(14), 269-281. Johnson, W. S. and Mangalgiri, P. D., Investigation of fiber bridging in double cantilever beam specimens. J. Composite Technol. Res., 1987, 9, 10-13. Hibbs, M. F., Tse, M. K. and Bradley, W. L., lnterlaminar fracture toughness and real-time fracture mechanism of some toughened graphite/epoxy composite, Toughened Composites ASTM STP, 1987, 937, 115-130. Hunston, D. L., Composite interlaminar fracture: effect of matrix energy. Composite Technol. Rev., 1984, 6, 176-180. Madhukar, M. S. and Drzal, L. T., Fiber-matrix adhesion and its effects on composite mechanical properties: IV. Mode I and Mode II fracture toughness of graphite/epoxy composites. J. Composite Mater., 1992, 26(7), 936-968. Albertsen, H., Ivens, J., Peters, P., Wevers, M. and Verpoest, I., Interlaminar fracture toughness of CFRP influenced by fibre surface treatment: part I. experimental studies. Composite Sci. Technol., 1995, 54, 133-145. Bradley, W. L., Relationship of matrix toughness to interlaminar fracture toughness. In Application of Fracture Mechanics to Composite Materials, ed. K. Friedrich. Elsevier, Amsterdam, 1989, pp. 159-187. Maikuma, H., Gillespie, J. W. and Wilkins, D. J., Mode II interlaminar fracture of the center notch flexural specimen under impact loading. J. Composite Mater., 1990, 24, 124-149. Waterbury, M. C. and Drzal, L. T., Determination of fiber volume fractions by optical numeric volume fraction analysis. J. Reinforced Plastics Composites, 1989, 8, 627-636. European Structural Integrity Society (ESIS), Protocols for lnterlaminar Fracture Testing of Composites, 1993.
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Hertzberg, R. W., Deformation and Fracture Mechanics of Engineering Materials, 4th edn. Wiley, New York, 1996. Hashemi, S., Kinloch, A. J. and Williams, J. G., Corrections needed in double cantilever beam tests for assessing the interlaminar failure of composites. J. Mater. Sci. Lett., 1989, 8, 125-129. Irwin, G. R. and Kies, J. A., J. Welding, 1954, 33, 193. Russell, A. J. and Street, K. N., Factors affecting the interlaminar fracture energy of graphite/epoxy laminates. ICCM-IV Tokyo, 1982, p. 279.
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PIh S1359-8368(98)00004-3
ELSEVIER
© 1998 Elsevier Science Limited Printed in Great Britain. All rights reserved 1359-8368/98/$19.00
Matrix effect on the static and dynamic interlaminar fracture toughness of glass-fibre marine composites
P. C o m p s t o n a'*, P.-Y.B. Jar a and P. Davies b aDepartment of Engineering, The Australian National University, Canberra 0200, Australia bMarine Materials Laboratory, IFREMER- Centre de Brest, BP 70, 29280 Plouzane, France (Received 20 November 1996; accepted I September 1997)
The interlaminar fracture toughness of four marine composites has been investigated. The matrix material was different in each composite. Mode I and Mode II tests were performed under static loading and the focus was placed on the transfer of matrix toughness to the composite. A coefficient of toughness transfer was used to quantify the synergy experienced in the Mode I results. Mode II tests were also performed under dynamic loading with the focus placed on the order of results in relation to the order of matrix toughness. The study concludes that synergy in Mode I results for crack initiation and steady-state propagation is greatest in the composites with superior interracial bonding. The Mode I results do not correspond directly to matrix toughness owing to interracial strength limitations, relative to matrix strength, in the composites with the toughest matrices. A similar pattern is evident in the results for crack initiation in static mode II tests. At the maximum load point in static Mode II tests, where fast fracture occurred, and in dynamic Mode II tests, the results are consistent with the order of matrix toughness. For these results, differences in interfacial bonding seem to have an insignificant effect. © 1998 Elsevier Science Limited. All rights reserved
(Keywords: interlaminar fracture; matrix toughness; glass-fibre marine composites)
INTRODUCTION There are many examples of the increasing use of fibrereinforced polymer-matrix composites for structural applications in marine environments. ~-7 While composites can offer high strength and stiffness in conjunction with weight reduction, polymer matrices in particular improve sonic characteristics, corrosion resistance, and aesthetic qualities. For economic reasons, glass fibre is the preferred reinforcement in marine composites. Therefore, cost effective improvement of any composite property will be achieved through the use of matrix materials with improved properties. Polyester resins have been the most common matrix material in the past, but attention is now focused on replacing them with resins such as epoxy and vinyl ester. The vinyl ester is of particular interest owing to its ageing resistance, s As marine composites are susceptible to delamination damage, 9 the effect of different matrices on the interlaminar fracture toughness requires investigation, t° Static tests are useful for characterisation of mechanical properties, however they do not completely reflect practical situations. * Corresponding author.
Delamination damage in marine composites can be caused by low-speed impact, ~1 and this can adversely affect the residual strength and stiffness in the composite. 12 Composite interlaminar fracture toughness under Mode I and Mode II static loading is characterised in terms of critical strain energy release rate, Go. 13-15 Many different composite systems have been characterised in this way, as indicated by Refs 16A7. It has been shown that matrix toughness, fibre-matrix bond strength and, especially in Mode I, fibre bridging 18 can all influence the value for composite Gc. Hibbs e t al. 19 showed that increased matrix toughness will increase composite interlaminar fracture toughness. However, along with other published work, 2°'21 Hibbs e t al. also emphasised that good interfacial bonding is necessary to promote cohesive failure through the matrix and, thus, ensure that matrix toughness is transferred to the composite. In a systematic study of the interface effect, Albertsen e t al. 22 showed that as adhesion between the fibre and the matrix increases, Gc for crack initiation in both Mode I and Mode II also increases. For a carbon-fibre/epoxy composite, they also found that an optimum level of adhesion between the fibres and the matrix, which promotes a combination of cohesive and interfacial failure, will
505
Fracture toughness of glass-fibre marine composites: P. Compston et al. maximise the Gc value for Mode I steady-state crack propagation. The interfacial failure is necessary for the formation of bridging fibres, failure of which is the major energy absorbing mechanism. As discussed by Bradley, 23 much of the published work from Mode I testing of brittle-matrix composites shows a synergistic effect, whereby composite Gc is higher than matrix Go. This was attributed to increased fracture area created in the matrix by crack propagation, the effect of fibre bridging and the slower rate of decay of the stress field ahead of the crack tip. Strong interfacial bonding would also contribute to the synergistic effect by promoting cohesive failure. Research on interlaminar fracture toughness under impact loading is less extensive and test methods are not standardised. However, Maikuma et al. 24 successfully induced Mode II dominated delamination in carbon-fibre composites using a drop weight impact tester. Two composites were tested, one with an epoxy matrix and the other with a tough PEEK matrix. The results showed interlaminar fracture toughness under low-speed impact to be dominated by matrix toughness. This paper reports on a study to characterise the interlaminar fracture toughness of marine composites. They were subjected to Mode I and Mode II static loading, and Mode II low-speed impact loading. The only variant among the composites was the matrix material. Discussion of the results from the static tests will focus on the transfer of matrix toughness to composite interlaminar fracture toughness. Under impact loading, the focus will be placed on the order of the results in relation to the order of matrix toughness.
Table 1 Matrix resin details Matrix resin
Supplier
Trade name
Epoxy (EP) Vinyl ester (VE) Isophthalic polyester (PI) Orthophthalic polyester (PO)
Sicomin DSM Scott Bader Scott Bader
SRI500 Atlac 580 Crystic 491 PA Crystic R115 PA
Table 2
Fibre volume fraction in UD region
Composite
Fibre volume fraction in UD region (%)
Coefficient of variation (%)
GF/EP GF/VE GF/PI GF/PO
40 46 39 39
6 5 4 7
water-cooled diamond cutting wheel. They were then dried under vacuum at room temperature for 24 h. Using the method proposed by Waterbury and Drzal, 25 the fibre volume fraction in the UD region of each composite was determined by calculating fibre area fraction from images of cross-sectioned samples. Each cross-section was 2 0 mm wide and 6 mm thick. Images of the UD region were obtained using a scanning electron microscope. The magnification was such that the total thickness of the UD region was contained within each image. This equated to 2000-3000 fibres per image. Using the public domain NIH Image program (version 1.57), the fibre area fraction, and hence fibre volume fraction, were calculated. For each composite, two images were taken from three samples. Therefore, each value for fibre volume fraction given in Table 2 is an average of six results. It should be noted that the fibre volume fraction is significantly higher for GF/VE than for the other composites.
MATERIALS E-glass fibres with a multi-compatible sizing, made by Vetrotex, were used for this study. Four thermoset resins were chosen as the matrix materials. They are detailed in Table 1. According to Vetrotex, the sizing (P177) is compatible for bonding with each resin. Table 1 also introduces a reference for each matrix which will be used in an abbreviated reference for each composite throughout the rest of the text. For example, the glass-fibre reinforced composite with the epoxy matrix will be referred to as the GF/EP composite. Four 12-ply composite laminates, one for each matrix resin, were produced by hand lay-up. The central four plies of each laminate contained unidirectional (UD) fibres of 250 g/m 2. The outer layers consisted of 0/90 ° stitched fabric plies of 590g/m z. To introduce a starter crack, a polypropylene film, 8/zm thick and 75 mm long, was placed between the central plies in the UD region. As recommended by the resin suppliers, the laminates with the polyester matrices were stabilised at room temperature for 1 month before testing and the laminates with the epoxy and vinyl ester matrices were post-cured at 90°C for 6 h. Test specimens were cut from these laminates using a
506
EXPERIMENTAL PROCEDURES The mechanical properties of each matrix resin were established through fracture toughness and tensile testing. The fracture toughness tests were conducted in accordance with the draft international standard ISO/DIS 13586(E) for plastics--determination of fracture toughness (Gc and Kc)--linear elastic fracture mechanics (LEFM) approach, 1996. Mode I double cantilever beam (DCB) and Mode II end notch flexure (ENF) tests were used to establish composite delamination resistance under static loading. These tests were conducted in accordance with the European Structural Integrity Society (ESIS) Protocols for Interlaminar Fracture Testing of Composites. 26 Mode II central notch flexure (CNF) tests were used to determine delamination resistance under dynamic loading. For this test, the method described by Maikuma et al. 24 was followed. In each test, the delamination crack was propagated parallel to the fibre direction in the UD region of the test specimen, and a minimum of five specimens from each composite were used. The specimens were not pre-cracked.
Fracture toughness of glass-fibre marine composites: P. Compston et al. RESIN FRACTURE TOUGHNESS AND TENSILE TESTING
in each test. Owing to material availability constraints, only one or two specimens could be tested for each resin.
The geometry of the single edge notched bend (SENB) specimen used for the resin fracture toughness, Ko tests is shown in Figure 1. Specimen dimensions were width, W, 14 mm, nominal thickness, B, 4 mm, and span distance, 2L, 56 mm. A natural crack was introduced by tapping a new razor blade placed in a machined notch, to produce a crack length in the range 0.45 < a/W < 0.55. Specimens were loaded in a three-point bend fixture at 10 mm/min. The load-displacement plots recorded were linear elastic and the maximum load Pmax w a s used to determine Kc from the equation: f(emax) g c -- - I/2(BW)
(1)
The calibration factor f is given in the draft ISO standard. Tests were valid according to the specimen size criterion given in that document. Between four and eight valid tests were performed for each material. Scatter in Kc was typically + 10%. From the Kc results, the Mode I critical strain energy release rate for each matrix resin, Gic_matrix, w a s calculated as f o l l o w s : 27
K2 Glc = ~-(1 - p 2)
(2)
where E is tensile modulus and v Poisson's ratio. Young's modulus, maximum strength and maximum strain were established using miniature tensile test specimens. They had a width of 10 mm, thickness 5 mm and gauge length 25 mm. Testing was conducted in displacement control on an Instron 4505 Universal Testing Machine (UTM) at a crosshead speed of 1 mm/min. An extensometer was used
Mode I DCB testing
The DCB specimen geometry used for static Mode I testing is shown in Figure 2. The thickness was approximately 6 mm, width 20 mm and length 160 mm. The latter dimension allowed for 70 mm of crack growth. For load introduction, aluminium blocks were attached to the end containing the crack starter film. A thin layer of white correction fluid was applied to one edge of the specimen, starting at the tip of the crack starter film. This edge was marked at 1 mm intervals for the first 5 mm and then at 5 mm intervals up to the 70 mm mark. This allowed the crack to be identified and measured as the test progressed. Testing was conducted in displacement control on the Instron UTM at a crosshead speed of 2 mm/min. A load-displacement plot was produced for each test, and load-displacement data were recorded for each individual crack length measurement. The Mode I critical strain energy release rate, Gic, was calculated using the corrected beam theory method, z8 It is derived from the Irwin and Kies expression for fracture energy: 29 p2 d C
Gc -- 2B da
(3)
where P is load, B specimen width, C compliance and a crack length. This expression accounts for the change in compliance with crack length. Using the simple beam theory expression for compliance: 6 2a 3 C. . . . P 3EI
(4)
where E is flexural modulus and I moment of inertia, eqn (3) may be rewritten as: 3P6 Glc = 2Ba
(3. / I~
-I
ZL
Figure 1 Single edge notched bend (SENB) test specimen geometry
where 6 is crosshead displacement. However, a correction factor IAL needs to be added to crack length to allow for crack tip rotation at the root of the cantilever beam. IAI is the x-axis intercept when the cube root of compliance, C m, is plotted as a function of crack length, a. Therefore, the corrected beam theory expression for Glc is: 3P3 GIc -- 2B(a + IAI)
crack starer
ao
I I I I I
(5)
(6)
film
LID region
Figure 2 Mode I DCB test specimen geometry with initial crack length a o
Values for GIc will be plotted as a function of crack length to produce a resistance (R) curve. Results will be presented for crack initiation, Glc-init, defined as the first deviation from linearity on the loaddisplacement plot, and steady state crack propagation, Gtc-prop, defined by a plateau on the R curve. The extent to which matrix toughness has been transferred to interlaminar fracture toughness in each composite, and the degree of synergy experienced for the
507
Fracture toughness of g/ass-fibre marine composites: P. Compston et al. Results will be presented for crack initiation, Gllc.init, again defined as the first deviation from linearity on the loaddisplacement plot, and for the maximum load point, Gllc-ma×. I
Mode H CNF impact testing 2L
The geometry of the CNF specimen used for the Mode II impact tests is shown in Figure 4. The thickness was 6 mm, width 20 mm and length 180 mm. The crack starter film is located in the centre of the specimen length. The CNF specimen was placed in a three-point bending fixture. Span length was set at 120 mm so that the specimens were sufficiently rigid to prevent excessive deflection during the impact. The tests were conducted using a Mitsubishi Yuka Co. instrumented drop weight impact tester. The striker width was 25 mm, sufficient to cover the specimen width and provide a contact condition similar to that in the Mode II ENF test fixture. The striker weight was constant at 2.6 kg and impact speed was varied between 1.5 and 3 rn/s in order to obtain a range of impact energies. The force during the impact was recorded by deformation of a strain gauge attached to the striker, and sampled every 20 #s. A forcetime plot was produced for each test and reduced into force and displacement data. The total absorbed energy was the area under the force-displacement curve. The total delamination fracture area in each specimen was measured after impact. A plot of absorbed energy versus delamination fracture area for a number of tests showed a linear relationship, therefore the average G~ value for each composite was taken as the slope of the line fitted to these points. A detailed description of how this test was conducted on these composites is available elsewhere. 31
I
Figure 3 Mode II ENF test specimen geometry
impact
Figure 4 Mode II CNF impact test specimen geometry
Glc.init and Glc.prop results,
will be expressed as follows:
GIc - composite ~
CTTGIc
-
matrix
(7)
where CTT is the coefficient of toughness transfer.
Mode H ENF testing The ENF specimen geometry used for static Mode II testing is shown in Figure 3. The thickness was approximately 6 m m , width 2 0 m m and length 160mm. The specimen was placed in a three-point bending fixture with half span length, L, set at 50 mm and the ratio of crack length to half span length, a/L, maintained at 0.5. One edge of the specimen was covered with white correction fluid, allowing the crack to be identified and measured at the end of the test. This test was also conducted in displacement control on the Instron UTM at a crosshead speed of 1 mm/min. It is stable until a maximum load point is reached, after which unstable fracture occurs. A load-displacement plot provided the raw data for each test. The Mode II critical strain energy release rate, Gtic, w a s calculated using direct beam theory. It is also based upon eqn (3). Using the ENF compliance expression given by Russell and Street: 3°
Scanning electron microscopy Fracture surface micrographs were obtained using a Cambridge $360 scanning electron microscope (SEM). Specimens were coated with a thin layer of gold prior to examination. The micrographs were used to characterise the dominant failure modes in each composite for each test. All micrographs presented in this paper show regions that are close to the crack initiation point.
2L 3 -t- 3a 3
C=
RESULTS AND DISCUSSION
8EBh3
(8)
where h is specimen thickness, eqn (3) may be rewritten as:
9a2P6 GIIc - 2B(2L3 + 3a3) Table 3
Resin properties The tensile properties of each matrix resin are given in Table 3. While the modulus is similar for each resin, the
(9)
Tensile and fracture toughness, Kc, properties of the matrix resins
Matrix resin
Young' s modulus (GPa)
Maximum strength (MPa)
Maximum strain (%)
Fracture toughness, Kc (MPa ,v/~)
Epoxy (EP) Vinyl ester (VE) Isophthalic polyester (PI) Orthophthalic polyester (PO)
3.4 3.5 3.5 3.7
70 85 68 58
5.4 3.2 2.9 2.1
1.125 0.694 0.451 0.419
508
Fracture toughness of glass-fibre marine composites: P. Compston et al. et al., 32 which showed that 50 mm of crack growth was required before a plateau region develops. Therefore, in the present study, Glc values for 50-70 mm of crack growth were averaged to give Glc-prop. The results for Gic_init are shown in Figure 7. They can be placed on two distinct levels. The GF/EP and GF/PI composites are on the higher level, and the GF/VE and GF/PO composites comprise the lower level. The results for Gic.prop a r e shown in Figure 8. The GF/PI and GF/VE composites have the highest values followed by GF/EP then GF/PO. The GF/VE composite, however, has a higher fibre volume fraction in the UD region (see Table 2) and has the potential to form more bridging fibres. Therefore, the value of Gk-prop for GF/VE cannot be compared with the other composites. Nevertheless, it is clear that the trend in the GIcinit and Glc_prop results is not consistent with the trend in matrix toughness as shown in Figure 5. In addition, the values of CTT for Gtc-init and Glc.prop in Table 4 show that greater synergy is achieved in the GF/PI and GF/PO composites. The CTT is discussed in terms of fibre-matrix adhesion, but other factors may also influence the toughness
values for maximum strength and maximum strain are higher for epoxy and vinyl ester. Therefore, the use of these two resins is consistent with the aim of improving composite properties through improved matrix properties. The K~ values, also given in Table 3, show epoxy to be the toughest matrix resin, followed in descending order by vinyl ester, isophthalic polyester and orthophthalic polyester. The Mode I critical strain energy release rate of each matrix resin, Glc-matrix, is depicted in Figure 5. The trend shown confirms the order of matrix toughness given by the K~ results, and will be compared with the trend in the results for composite Mode I interlaminar fracture toughness.
Mode I DCB fracture toughness Typical R curves for each composite are shown in Figure 6. The Glc_init value is the first value in the curve and corresponds to the onset of crack growth. In each R curve, the Glc values continue to rise until approximately 50 mm of crack growth is achieved. This is consistent with the R curves for glass-fibre composites produced by Kawada
0.35 0.3 0.25 0.2
0.15 ~,~
0.1 0.05
EP
VE
Figure 5 Mode I critical strain energy release rate for each matrix resin,
PI
alc_matri
PO
x
2000
ea
1500
~
0
0
D
V
V
~ 0
0
0 0 [] []
I000
500
~ V
I
I
0
i
i
,
i
I
10
i
i
,
J
I
20
[] V
I
,
V
i
i
I
30
~ v
V
i
i
i
b
I
40
i
i
'
'
I
. . . .
50
I
. . . .
60
I
,
70
Crack growth (ram) Figure 6
Typical Mode I R curves for each composite: GF/EP (D); GF/VE (©); GF/PI (A); and GF/PO (V)
509
Fracture toughness of glass-fibre marine composites: P. Compston et al. 0.35 0.3 0.25 0.2 ,m ,~
~,~
0.15 0.1 0.05
GF/EP
GF/VE
GF/PI
GF/PO
Figure 7 Mode I DCB fracture toughness for crack initiation, alc.init. The error bar signifies _ 1 standard deviation
1.5
eL
1
O L
0.5
GF/EP Figure 8
GF/VE
Coefficient of toughness transfer, C/T, for Mode I results
Composite GF/EP GF/VE GF/PI GF/PO
C/T~initiation
CTT--propogation
( G lc.init/G lc-matrix)
( G lc_prol~G Ic-matrix)
0.9 1.3 5.1 3.9
2.9 11.5 a 25.6 18.4
~Comparitively high fibre volume fraction.
510
GF/PO
Mode I DCB fracture toughness for steady-state crack propagation, Glc-prop-The error bar signifies + 1 standard deviation
transfer, such as fibre constraint and shrinkage stresses. The contributions of such factors are not considered here. The fracture surfaces of the Mode I DCB specimens in Figure 9 show an important variation in the level of fibrematrix adhesion. The fibre surfaces in the GF/EP and GF/ VE composites are completely devoid of matrix, indicating adhesive failure along the interface. In contrast, the fibre surfaces in the GF/PI and GF/PO composites are partly covered with matrix, indicating a degree of cohesive failure through the matrix. Therefore, there is greater utilisation of matrix toughness in GF/PI and GF/PO, and this is reflected in the values for CTT. In the case of GF/EP, the value for CT/" suggests that the superior matrix toughness is negated Table 4
GF/PI
by interfacial strength limitations, relative to the matrix strength, resulting in a GIc-init which is similar to GF/PI. For the same reason, the value of GIc-init for GF/VE is similar to GF/PO. The difference in fibre-matrix adhesion between GF/PI and GF/EP has also been sufficient to produce a large difference in the value of Glc-prop. The increased cohesive failure in GF/PI has served to utilise matrix toughness and increase bridging fibre lifetime. Thus, the GIc-prop value for GF/PI is higher than GF/EP. However, GF/PO has the lowest GIc-prop value despite good fibre-matrix adhesion and the comparatively high CTT value. In this case, the failure of the improved bonding to increase Gic_prop is attributed to the inherent brittleness of the matrix. These propagation values are strongly influenced by fibre bridging, but may be important in areas where unidirectional reinforcements are used to reinforce critical areas of structures.
Mode H ENF fracture toughness The results for Gllc.init a r e shown in Figure 10. The GF/EP composite has the highest value, and this is consistent with
Fracture toughness of glass-fibre marine composites: P. Compston et al.
50/~m Figure 9
Fracture surfaces of Mode I DCB specimens: (a) GF/EP; (b) GF/VE; (c) GF/PI; and (d) GF/PO
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Figure 10 Mode II ENF fracture toughness for crack initiation, Guc-init. The error bar signifies -+ 1 standard deviation matrix toughness. However, the value of Gllc.init for GF/VE is similar to GF/PI and GF/PO. The results for GIlc-max, shown in Figure 11, are consistent with matrix toughness. The fracture surfaces of the Mode II ENF specimens are shown in Figure 12. As seen for Mode I DCB specimens, the fibre surfaces in GF/EP and GF/VE are devoid of matrix whereas in GF/PI and GF/PO they are still partly covered with matrix. The cohesive failure in GF/PI and GF/PO and
the adhesive failure in GF/VE has resulted in a similar value of Gtlc_init for these three composites. However, the fact that GF/EP has the highest value suggests that fibre-matrix adhesion has been less influential at crack initiation in the Mode II test than in the Mode I test. Matrix deformation, in the form of hackle marks, can also be seen in Figure 12 and is typical for Mode II specimens. ~9 This observation is consistent with a matrix dominated failure process and supports the Gllc-max results.
511
Fracture toughness of glass-fibre marine composites: P. Compston et al.
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GF/VE
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M o d e II E N F fracture toughness at m a x i m u m load point, GIIc-max.The error bar signifies _+ 1 standard deviation
50 #m Figure 12 Fracture surfaces o f M o d e II E N F specimens: (a) GF/EP; (b) GF/VE; (c) GF/PI; a n d (d) G F / P O
Mode H CNF impact fracture toughness The plots of absorbed energy versus delamination fracture area for each composite are shown in Figure 13. The G~ value is the slope of the line fitted to the points in these plots. The equations for each line are as follows:
512
GF/EP:
GF/VE: GF/PI: GF/PO:
y y y y
= 0.006985 + 2.0405x = 0 . 0 7 2 1 8 4 + 1.7657x = - 0.21507 -f- 1.5346x = 0.3754 + 1.4103x
R R R R
= = = =
0.95067 0.95932 0.8158 0.88041
where y is absorbed energy loss, x fracture area and R goodness-of-fit for each line (where R = 1 is a perfect fit).
Fracture toughness of glass-fibre marine composites: P. Compston et al. .
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Plots of absorbed energy Eah versus fracture area AA from Mode II CNF impact tests: (a) GF/EP; (b) GF/VE; (c) GF/PI; and (d) GF/PO
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GF/PO
Mode II CNF impact fracture toughness, G~, For comparison, Gu~ m,~ results from Mode II ENF tests are included
The G~ results, shown in Figure 14, correspond directly to matrix toughness. The fracture surfaces of the CNF specimens are shown in Figure 15. There is greater fibre-matrix adhesion in the GF/PI composites after fracture, which is consistent with Figures 9 and 12. Hackle mark deformation can be seen in each matrix,
consistent with Figure 12, with increased deformation apparent in the tougher EP and VE matrices. In conjunction with the GI~ values, observation of such matrix deformation characteristics supports the conclusion that energy absorption has been matrix dominated, and that the variations in fibre-matrix adhesion which
513
Fracture toughness of glass-fibre marine composites: P. Compston et al.
a
Figure 15 Fracture surfaces of Mode II CNF impact specimens: (a) GF/EF; (b) GF/VE; (c) GF/PI; and (d) GF/PO
influenced Glc.init, Gic_prop and Giic_init results have had a less significant effect. The Gllc-max values are also included in Figure 14 for comparison with G D. While the trend matches well, the absolute values for G D are lower than for Gn ..... even though both results represent energy absorbed in resisting fast fracture. This suggests a loading rate effect. The difference in specimen geometries for the quasi-static and dynamic tests may contribute to this difference. Published values for carbon-fibre reinforced composites indicate that the ENF and CNF specimens yield similar values. 24 To examine this further for the composites used here, a series of static tests was performed on CNF and ENF specimens taken from the same glass-fibre/epoxy laminate. The expression used to determine Glic was:
9PZa2
(10)
GIIc = 64EB2h 3 for a specimen of thickness 2h. The results were analysed using the same modulus values. Thus, five tests on CNF specimens yielded mean (standard deviation) GHc.... values of 3.72 (0.26) kJ/m 2, compared with 3.50 (0.35) from the ENF specimens. This suggests that the specimen geometry does not explain the rate effect. Further work is being carried out to verify the effect of loading rate on glassfibre reinforced composites.
514
CONCLUSIONS The interlaminar fracture toughness of four glass-fibre marine composites has been investigated. Mode I (DCB) and Mode II (ENF) tests were conducted under static loading. The focus of the discussion on results from these tests was placed on the transfer of matrix toughness to composite interlaminar fracture toughness. The synergy experienced in the Mode I results, expressed in terms of a coefficient of toughness transfer, CTT, was also discussed. In addition, Mode II (CNF) impact tests were conducted under low-speed impact loading, with the focus placed on the order of results in relation to the order of matrix toughness. The following conclusions are made: • In static Mode I tests, the alc.init results show that the transfer of matrix toughness and the synergistic effect, reflected in values of C/T, is greatest in the composites with the brittle polyester matrices. That is, in GF/PI and GF/PO. This is attributed to superior adhesion between the fibres and the matrix, which has promoted greater cohesive failure and utilisation of matrix toughness. The GF/EP composite, which has the toughest matrix, has the highest absolute value for Glc-init, but the lowest C / T value. This indicates, in conjunction with the interfacial failure observed on the fracture surface, that interfacial strength limitations have reduced the effect of superior
Fracture toughness of glass-fibre marine composites: P. Compston et al. matrix toughness. The same reason is given for the comparatively poor performance of the GF/VE composite which has the vinyl ester matrix. The highest value of GIc-prop, and corresponding CTT, was obtained for GF/PI. As for Glc.init, this is attributed to the superior fibre-matrix adhesion. It has enhanced the effect of fibre bridging and placed GF/PI closer than GF/EP to the optimum level of bonding required to maximise Gic-prop.The GF/PO composite also has a high CTT value, but the lowest value for G~c-prop,and this is attributed to the inherent brittleness of the matrix. The GF/VE composite was omitted from this comparison owing to its comparatively high fibre volume fraction. • In static Mode II tests, the value of Giic.init is highest for GF/EP. This is consistent with matrix toughness. However, the value for GF/VE is similar to GF/PI and GF/PO. This is attributed to the superior bonding and greater cohesive failure in GF/PI and GF/PO which is evident on the fracture surfaces. The GII..... results, which reflect energy absorbed in resisting fast fracture, are consistent with matrix toughness. The fracture surfaces show the areas of pure matrix resin to be highly deformed, indicating extensive participation of the matrix in the fracture process. • The G~ results from the CNF impact tests are also consistent with matrix toughness. Fracture surface examination revealed hackle mark deformation in the matrix. A greater amount of deformation was observed in the tougher matrices. This supports the conclusion that, for these composites, the fracture process under lowspeed impact is matrix dominated. • From a practical point of view, the use of a sizing which enhances compatibility of glass fibre with epoxy and vinyl ester resins is recommended. This would enhance the static interlaminar fracture toughness in glass-fibre marine composites which use these resins as the matrix material.
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5. 6.
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11.
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13. 14.
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ACKNOWLEDGEMENTS This study was supported by the Australian Research Council (ARC) under the Large Grant Scheme. The technical support of staff in the ANU Electron Microscopy Unit is also appreciated.
20. 21.
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