Mode I fracture toughness of fiber-reinforced polymer composites: A review

Abstract

Composite materials are known for their high stiffness and strength at lower weight as compared to conventional structural materials. Recently, there has been a growing interest in finding the new ways to decrease delamination failure, which is a life limiting factor of laminated composites. This review paper emphasizes on the effects of different reinforcement structures on mode I fracture toughness and possible ways to improve fracture toughness. A brief description on intrinsic and extrinsic mechanisms of crack growth has been discussed along with the earlier investigations and recent developments for mode I fracture toughness testing. Factors that affect the fracture toughness are also discussed. A brief knowledge of mode I fracture toughness of traditional and advanced fiber-reinforced composites is given, which could help researchers to understand fracture behaviors of composites and thus, it can help engineers to design composites with higher interlaminar strength.

Keywords

Fiber-reinforced composites fracture toughness delamination mode I interlaminar 2D composites 3D composites

Introduction

The high specific modulus, also known as stiffness to weight ratio, of fiber-reinforced polymer composites (FRPCs) has resulted in the widespread use of these composites as key structural engineering materials in automotive, aerospace [1 –3], marine, transportation, infrastructure, civil engineering applications [4 –8] and motorsport industries [4,9,10]. Since, there has been an ever-increasing demand for composite materials and new materials are being introduced in the market frequently, so it is challenging to assess their properties in order to forecast service life and failure behaviors. The performance evaluation of the advanced reinforcing fibers (glass [11], carbon [12,13], Kevlar [14,15] and PBO (poly(p-phenylene-2, 6-benzobisoxazole) [16,17] fibers) and resins (Epoxy [18 –21], vinyl ester [11], polyimide (PI) [22] and cyanate ester [23]) in final composite is necessary for their safe application.

Compared to other damage categories, delamination evolution is a dominant life-limiting failure mode for composite structures [13,24 –28]. It may be introduced either during manufacturing or caused by damaging events during service, for example, impact damage [29,30], large hailstones [31,32] and bird strike events [33,34], thus, hampering structural integrity and durability [35,36]. Keeping in view the adverse impact of delamination on the composite structures, it is necessary to consider in structural design process as well as in verification testing, to ensure safe application with prolonged life of these composites.

Property which defines the ability of a material to resist fracture is known as fracture toughness [37,38]. It indicates the amount of stress required to propagate a pre-existing thin crack. Damage tolerance is the desired basic property for various structures depending upon the end application [39]. Durability is an economic life-cycle design concern, whereas, damage tolerance tells about the structure's ability to securely withstand flaws until these flaws are removed or fixed. Fracture toughness characterization of composites is still on the way of growth as compared to metals. Fracture toughness depends on many factors. Many researchers reported that in FRPCs, the performance of fibers depends on the reinforcement's architecture [13,28,40], which is discussed in the next sections. Hence, structural effects of reinforcement as well as effects of other factors (manufacturing process, loading type, loading rate and temperature) on interlaminar fracture toughness are worth studying. Furthermore, the knowledge of damage and its mechanism should be extended to design composites with higher interlaminar strength for longer life span.

Fracture toughness

Fracture toughness is a distinguished property for many structural designs and applications to ensure reliability [41,42]. The value of this property is typically denoted by the strain energy release rate ( $G_{C}$ ).

In composites, strain energy release rate ( $G_{C}$ ) denotes the material resistance to interlaminar fracture. The value of strain energy release rate is denoted as $G_{IC}$ , $G_{IIC}$ , or $G_{IIIC}$ with respect to cracking mode being tested (Figure 1). The $G_{C}$ value at which the delamination essentially starts to spread differs largely depending on the mode of loading [5]. As the material is being tested and the crack begins to propagate, the stiffness and force on the material begin to decrease. The decrease in the load means that the strain energy stored in the material is also reducing or being released.

Figure 1.

Various cracking modes for fracture toughness classification.

Modes of fracture

The manner in which a crack propagates through a material gives insight into the mode of fracture. Thus, there are different modes of fracture for these composite materials. Cracking mode I is considered to be the crack opening or tensile mode of delamination. It is the most common form of fracture failure as its motion is like pulling plies of material away from each other. The crack faces undergo opening displacements relative to one another as it propagates. The corresponding material property $G_{IC}$ (fracture toughness), under mode I loading, is commonly obtained by the double cantilever beam (DCB) test method [43].

Cracking mode II is the in-plane shear mode of delamination. A shear stress acts parallel to the plane of the crack and perpendicular to the crack front. This is classified by two separated plies of material sliding above each other in the path of crack growth. This cracking mode is less common than the crack opening tensile mode, but it is still relevant to designs where force is not particularly down the center of a structural component. In cracking mode III, known as tearing mode, a shear stress acts parallel to the plane of the crack and parallel to the crack front. This is also termed as out of plane shear mode.

Types of preform used in textile-reinforced polymer composites and their impact on fracture toughness

The high specific stiffness and strength of composites have been of great deal for their utilization in structures where higher weight is a life-threatening variable. Hence, the development of low weight yet high strength composites are used in aerospace [44,45] and defense fields [46]. The advanced textile fibers such as Dyneema, Kevlar, glass, carbon and Zylon are known for their outstanding performance in protective applications. The performance of textile fibers varies depending on their architecture when used as reinforcement in composites. Hence, mechanical characterization in different structural forms [13,47 –49], under several static and dynamic loading situations, is the area of interest in recent years.

Laminated composites

Most of the FRPCs are made from laminates. Interlaminar delamination presents one of the most important life limiting failure mode that restricts their applications [5,8 –10,28,50,51] especially in primary aircraft structures [4,52]. As reported by Tamuzs et al. [53], the critical energy release rate of laminated composites does not usually exceed from 0.2 to 0.4 kJ/m². These composite materials show a worrying weakness towards the presence and development of cracks between the layers [10] especially in mode I loading. [40]

Delamination can be initiated due to the occurrence of flaws during fabrication process or produced in structural components during service due to interlaminar tension and shear. It may also develop due to a multiple factors such as impact of falling objects, structural discontinuities, free edge effects, differences in moisture and temperature [7,10]. The growth of delamination results in stiffness loss and could ultimately result in catastrophic failure [50,54,55]. Crack growth is a result of competition among micro structural intrinsic and extrinsic mechanisms. Intrinsic mechanisms perform in front of crack tip and promote crack evolution. Extrinsic mechanisms perform behind crack tip and impede crack evolution [56].

In fibre-reinforced composites, the intrinsic mechanisms are structural voids [57], material dislocation [58] or debonding [59,60] generated during manufacture or in-service. The shielding mechanism, also known as external mechanism, results from the closure traction between crack faces in the crack opening due to unbroken fibres or yarns and by deflection of crack path [13,61]. Extrinsic mechanisms give rise to resistive curve (R-curve) behaviour and thus, increase the force required for crack growth [13,28,62]. In this regard, various methods have been introduced to improve interlaminar strength [63]. Z-pinning [64 –68], stitching [39,68 –70], 3D weaving [71] and braiding [9,68] are popular methods to improve fracture toughness of textile composites. Advanced textile composites prepared by these technologies have shown good impact damage tolerance [72 –74].

Stitched laminated composites

Some experimental studies have been conducted on stitched unidirectional or woven FRPC composites to find interlaminar delamination. It was found that the stitched FRPC laminates present significantly larger interlaminar toughness as compared to un-stitched laminates [4,74 –79], and that an increase in stitch density caused an increase in mode I strain energy release rate [4,74 –76]. A recent study [80] has been carried out on flax fiber-epoxy composite laminates to find the effect of stitching in-through-thickness direction on mode I fracture toughness. At least 10% improvement in fracture toughness at the lowest fraction of stitch fiber was reported.

But, it is problematic to stitch big and complex structures by using existing stitching machines. It requires automated, intricate multi-needle stitching machines, which are costly for composite manufacturers. Meanwhile, the in-plane yarns can be distorted or damaged because of stitching leaving a negative impact on the in-plane mechanical properties [4,75].

Knitted fabric composites

Mode I fracture studies performed on knitted composites reported that these composites have extraordinary fracture toughness properties, being their $G_{IC}$ value 2 to 10 times greater than traditional composites [74,81,82]. The origin of the high $G_{IC}$ values, which are superior to woven or UD laminates, lies in the very complex fabric architecture [82]. Three-dimensional loop structure induces various energy-consuming mechanisms such as crack branching, friction, yarn bridging and breakage, which increase their fracture resistance.

Z-anchor-reinforced composites

Z-anchor (z-pinning) is a novel through-the-thickness reinforcement, which is used to improve the interlaminar strength of composite laminates [83 –88]. In the z-anchor process, in-plane yarns are entangled each other by sticking with special needles as shown in Figure 2. Z-pinned composites have significantly higher mode I delamination resistance than two-dimensional (2D) equivalents [9,88 –91]. Kusaka et al. [88] experimentally characterized the mode I interlaminar fracture behavior of z-anchor-reinforced composite laminates with different z-anchor densities. Mode I interlaminar fracture behavior of z-anchor-reinforced composite laminates increased almost linearly with the z-anchor density. Pingkarawat et al. [92] studied the effects of length, diameter and the volume content of z-pins on interlaminar fracture toughness. It was reported that fracture toughness increased with an increase in volume content and length of z-pins and by decrease in diameter. During crack propagation process, z-anchor reinforcement induced a large amount of fibers bridging resulting in the improvement of the mode I fracture toughness of composite laminate [88,92].

Figure 2.

Z-pinned reinforced composite [87].

Interlaminar z-directional flock-fibers-reinforced composites

Another technique of manufacturing composites with greater delamination resistance is to introduce interlaminar z-directional flock-fibers [93]. Jia et al. carried out a study to compare the mechanical properties, specifically, the mode I fracture toughness of z-directional flock fiber-reinforced composite with 3D braided composites. It was reported that the $G_{IC}$ value of 3D braided composite was higher than z-directional flock fiber-reinforced composite due to difference in toughening mechanisms as shown in Figure 3. But, 3D braiding has low productivity as compared to flocking technology. Moreover, z-reinforced flocking composite technology is found to be more cost effective [50].

Figure 3.

(a) 3D braiding composite (b) z-directional flocked fibers-reinforced composite.

Chemical methods to improve fracture toughness of laminated composites

A common technique that has been adopted to improve the fracture toughness of laminated composites is to use toughened matrix material [39,94 –97]. Later it was found that at higher temperatures, these resins show instability [50].

Another method to increase the interlaminar toughness of laminated composites is the addition of whiskers. In this regard, Yamashita et al. [98] and Sohn et al. [99] conducted experimental tests and found significant improvement in mode I fracture toughness of the composites. Wang et al. [4] also reported a significant improvement in mode I and mode II fracture toughness of CFRP laminates by dispensing whiskers at interface during the lay-up procedure. However, it was found that complications in the fabrication process and high cost are major hurdles in the practical application of this method [4].

Each of the above-mentioned methods proposed a way that can result in improved interlaminar fracture toughness and damage resistance. However, three-dimensional (3D) textile technologies are supposed to be the most effective ones. Numerous varieties of 3D fabrics have been developed which contribute distinctive features [5,9,13,100].

3D Woven textile composites

3D woven composites offer two major benefits compared to the 2D conventional laminated composites. Firstly, by altering warp and weft yarns, weave design, amount of binder, their mechanical properties can be customized to fulfill particular requirements without laminating [101,102]. Secondly and possibly more essentially, the z-yarns in through-thickness direction are interweaved with warp and weft yarns which prevent the growth of delamination [6] and increase post-impact mechanical properties [103]. In consideration of the angle of binder yarns, angle-interlock and orthogonal weaves can be identified as shown in Figure 4. In angle-interlock weave, binder tows are placed at an angle of θ < 90°, which is called undulation angle. In orthogonal-interlock weaves, the tows are placed at θ = 90°. When z-yarns are limited to an adjacent layer or a few layers, we call it layer-to-layer (or ply-to-ply) interlock weave [13], and when they go all the way through-the-thickness and interlace with all the layers, we call it through-thickness weave [105].

Figure 4.

3D woven fabrics [104] (a) Orthogonal woven fabric. (b) Layer by layer angle interlock. (c) Through thickness angle interlock.

3D woven fabric composites have the ability to provide superior interlaminar fracture toughness and greater impact damage resistance than 2D conventional laminated composites [28,72,101,102]. Other desirable attributes include weight reduction, improved fatigue and fracture resistance and fully integrated assemblies.

3D woven composites are used for highly engineered applications such as I-Beams for civil infrastructure [106], multi-hit composite armor systems [107], aircraft frame work [105], leading edges of air craft wings [106] and engine fan blades of aircrafts [105,108]. For these applications, high through-thickness mechanical properties and damage tolerance are inevitable. Numerous experimental investigations on interlaminar toughness have been reported for 3D-reinforced composites [5,6,100]. Enhanced damage tolerance is attained due to bridging formation by z-binders in the crack wake. It impedes the growth of impact created damage, and thus minimizes the depletion of the post-impact properties [109]. It was found that a small volume content of binder yarns results in a big improvement to the damage tolerance [109,110].

3D Orthogonal inter-lock woven composites

Watanabe et al. [5] studied the interlaminar fracture toughness of 3D orthogonal interlocked fabric composites. They developed a numerical model for the characterization of delamination. The influence of through-thickness yarns on delamination growth was quantified using that model. Guénon [100] carried out a study on 3D carbon/epoxy orthogonal fabric composite to find the effect of 3D fabric geometery on interlaminr and in-plane fractuer behaviour of composite. The interlaminar fracture toughness results of 3D-reinforced composite were compared with 2D laminates having a compareable in-plane stacking arrangement and material system. A 10 times gain of interlaminar toughness and 25% increase in the in-plane toughness were attributed to the through-the-thickness fibers. Energy dissipitating processes associated to the existence of z-direction fibers were: fracture debonding, crack brancing and deviating, crack bridging caused by in-plane fibers and pull out of the z-direction fibers.

Iwahori et al. [6] developed a mechanical model of DCB for 3D orthogonal interlock fiber composite using a 2D finite element method (FEM). It was to simulate the DCB test results. They found a good agreement between simulation and experimental test results. It was verified that the $G_{IC}$ usually achieved from DCB tests can be assessed by interlaminar tension test and 2D FEM simulation. They also reported the strong influence of the z-fibers on $G_{IC}$ .

3D angle interlock woven fabric composites

Among different 3D woven fabric composites, interlock fabric composites provide a fibre architecture intended to improve interlaminar toughness, while, maintaining in-plane performance by increasing only a small volume of through thickness reinforcement [9,72,102,111,112]. According to Siddique et al. [13,28], fracture toughness of 3D angle interlocked composites is supplemented by energy-absorbing mechanisms, such as matrix breakage, yarns bridging and pull out and occurrences of secondary cracks as well as their propagation.

Comparisons of mode 1 fracture toughness properties of traditional composites and advanced textile composites

The comparison of mode I interlaminar toughness of traditional composites made from unidirectional or woven roving and composites manufactured from innovative textile technologies of stitching, knitting, braiding, and through-the-thickness weaving is given by Table 1. Their $G_{IC}$ values are at least twice as high as traditional composites. The magnitude of the improvement in $G_{IC}$ depends upon the textile processing factors, such as density and thickness of the through-thickness stitch/binder, the knit yarn density and braiding angle [74]. The toughening of knitted and braided composites is attributed to widespread crack splitting combined with the zig-zag delamination path. The better interlaminar toughness in the stitched and through-the thickness woven composites is attributed to crack bridging, though crack splitting and stitch/binder pull out effects [74]. Figure 5 illustrates the dominating mode I fracture mechanisms of traditional and advanced textile composites.

Figure 5.

Fracture mechanisms (a) Plies separation (b) matrix cracking (c) yarns bridging (d) crack deflection (e) secondary crack formation.

Table 1.

Overview of the mode I fracture toughness of composites.

Fiber/matrix	Weave	G_IC (J/m²)	References
Carbon/epoxy	Unidirectional	629	[113]
Carbon/epoxy	Unidirectional	820	[114]
Carbon/epoxy	Unidirectional Cross ply [0°/90°]	751 1277	[115]
Carbon/epoxy	UD tape satin fabric	280 780	[116]
Carbon/epoxy	UD woven fabric	890	[113]
Carbon/epoxy	UD laminate	300	[118]
Carbon/epoxy	[0°/90°]	1295	[119]
Carbon/epoxy	[0°/90°]	282	[120]
E-glass/vinyl ester	Plain woven	450	[11]
Carbon/epoxy	Plain woven	508	[45]
Carbon/epoxy	Laminate	216,000	[121]
Carbon/epoxy	Multidirectional laminates [0°/5°] [−45°/+45°] [−90°/+90°]	825 835 1000	[122]
Carbon/epoxy	Woven laminate [0°]	728	[123]
Carbon/epoxy	3D Angle interlock 3D Orthogonal	704 1569	[9]

The analysis of the previous mode I interlaminar fracture studies shows that advanced textile manufacturing methods are proficient to fabricate composites with significantly increased delamination resistance.

Testing of fracture toughness

Delamination is a crucial failure mode of fiber-reinforced composite structures. For the evaluation of delamination resistance, researchers have made attempts, which resulted in the improvement of various test methods [124].

ASTM 5528 test is recommended for mode I fracture toughness characterization ( $G_{IC}$ ) of laminated reinforced polymer composites, using the DCB specimen [10,125].For mode II fracture toughness (G_IIC), the end notch flexure test is commonly used [126,127]. For pure mode III fracture toughness (G_IIIC), the edge crack torsion test was proposed by Ratcliffe [128]. ASTM D6671 recommends mixed-mode bending test to measure the fracture toughness under various mode I and mode II loading combinations [127]. There is still a lack of standard test methods to measure fracture toughness under variety of loading conditions and materials.

Most of the mode I fracture tests use a split cantilever beam specimen as shown in Figure 6(a). However, in case of materials with very high toughness (e.g. due to through-the-thickness reinforcement) or with very low thickness, the failure of the specimen arms has been observed prior delamination growth, thus, making a delamination fracture toughness measurement almost impossible [124]. The failure of beam arms normally starts on the outer surface of the specimen due to the concentration of compressive stress [124]. Normally, the fiber-reinforced composites have lower strength under compression as compared to tension, and the lowest compressive strength is particularly observed at the outer surface of the material where the fibers have less support [124]. Hence, it leads to failure of the specimen arm at surface, before delamination extension occurs [70,124]. In this situation, tabbed DCB specimen is better choice.

Figure 6.

Geometery of specimens for mode I testing (a) Double cantilever beam [145]. (b) Single edge notch beam tension specimen. (c) Compact tension specimen [142].

Types of specimens

Single edge notch tension specimen

The geometry of single edge notch tension specimen for Mode I tests [129] is shown in Figure 6(b). It has been used under static [129 –131] and impact [73] loading conditions. This type of specimen offers lower stresses in front of crack tip, and thus results in self similar crack propagation and lower failure loads [129,132].

Compact tension specimen

Another specimen type used for mode I testing is compact tension specimen [133 –144]. Figure 6(c) shows configuration of compact tension specimen. It is preferred for materials having limited orthotropic behavior such as: particulates, short fibers, nanoparticles and whiskers-containing composites. Moreover, it is used to get fracture toughness just at damage initiation because of its limited ligament length to find fracture propagation. Rikards [125] used a modified compact tension specimen for fracture toughness testing of glass-epoxy-laminated composites.

DCB specimen

DCB is an ideal specimen type in mode I interlaminar fracture tests [4,9,41,45,114,120,123,127,146 –154]. It comprises of a rectangular composite specimen with uniform thickness as shown in Figure 6(a). A mechanical cut is machined [71], or a non-adhesive Teflon film is inserted in the middle plane during production, which acts as delamination starter. Two loading masses are attached on bottom and top surfaces of the end of DCB specimen arms. The delamination starter end of the DCB specimen is moved apart by quasi-static or dynamic loading. During the test, the delamination lengths are recorded. For more accurate delaminating length measurements, a travelling microscope is recommended by ASTM [127]. DCB specimen is used for quasi-static tests and has an advantage of moderate cost of manufacturing.

Wedge insert fracture specimen

Mode I testing using WIF type specimen is also an accepted method [155,156]. WIF is the modified form of DCB specimen and is mostly used for the dynamic fracture toughness tests. The only disadvantage is to measure the coefficient of friction between composite surface and wedge [157].

General experimental procedures

Rectangular shape specimen for mode I test can be manufactured from different types of materials and methods as mentioned above. An initial crack can be formed using non-adhesive film [8,115,120,158] or mechanical cutting using saw [13,71]. Screw-driven machines [117], servo-hydraulic [8,159] or tensile testing machines [115,158,160] can be used to perform quasi-static test at constant cross-head displacements [161]. According to the aim of investigation, the test materials and test conditions can be changed. During test, crack length can be found visually or by using travelling microscope that help to see crack growth path [117]. Fracture process can be recorded using camera to analysis crack initiation and propagation. Fracture toughness calculations can be done using schemes mentioned in the next section (Data reduction methods). Scanning electron micro-scope images can be used to determine the micro damages in the specimens [113,160,162].

Data reduction methods

To determine the mode I fracture toughness according to the ASTM D5528 standard, three methods are used [78]: the modified beam method (MBT) [163 –165], the compliance calibration (CC) method and the modified compliance calibration (MCC) method. The difference between them by means of fracture results is not more than 3.1%. Among these, the MBT gives the most conservative value of three methods as reported by O'Brien and Martin [7].

Figure 7.

Graphs of data reduction theories. (a) Modified beam theory. (b) Compliance calibration. (c) Modified compliance calibration.

Modified beam theory

G_{I} = \frac{3 P δ}{2 b (a + Δ)}

(1)

where P and δ represent the load applied and displacement. a, b and Δ represent crack length, specimen width and correction for crack length, respectively.

G_{I} = \frac{nP δ}{2 ba}

(2)

where n is the exponent form

\frac{Δ y}{Δ x}

, whereas, Δy and Δx are defined in Figure 7.

MCC

G_{I} = \frac{3 P^{2} C^{23}}{2 A_{1} bh}

(3)

where A₁ represents the slope of line and h is the thickness of specimen.

Factors effecting mode I fracture behaviour

Matrix toughness

Bradley [96] and Jordan et al. [97] found that toughness of matrix shows an important role in the interlaminar fracture behavior of composites. Decrease in the yield strength of the matrix results an increase in the delamination fracture energy by increasing the size of non-linear visco-elastic region or plastic deformation ahead crack tip [13,166,167]. Finally, it results in higher load redistribution far from the crack tip and therefore, more crack-tip blunting is seen. Crack tip blunting increases the radius of crack tip [168,169], and ultimately increases the fracture toughness by decreasing the stress intensity at crack tip. [13]

Weave type

Some investigations have been done to find the effect of weave pattern on fracture toughness ( $G_{IC}$ ) of laminated composites by mode I fracture toughness test, and it showed some significant influence [11,170]. Suppakul and Bandyopadhyay [11] compared the mode I fracture toughness of twill, satin and plain weave laminates and reported that twill weave showed the highest $G_{IC}$ value followed by satin and plain weave, respectively. Fibre bridging is reported as dominant mechanism which contributed to the fracture resistance of twill and satin woven laminates, which was absent in plain woven laminated composite. Fishpool et al. [9] compared the interlaminar fracture behaviour of three geometries of woven carbon fibre composites layer to layer, orthogonal and angle interlock using DCB specimen, and found significant effect of weave architecture. Orthogonal weave was found to be the most effective in resisting delamination propagation, whereas, angle interlock weaves were the least influential. In orthogonal weave, delamination resistance was contributed by through-the-thickness fiber tows bridging, fibers and resin pull out behind the crack front. In angle interlock, weave architecture as fiber tows was staggered across the width of specimen, and thus, it resulted in twisting of the specimen arms. Guénon [100] and Tanzawa et al. [5,61] also reported that in orthogonal interlocked composites, many complicated mechanisms such as debonding of z-fiber tows from in-plane layer, crack branching, pull-out of z-fiber tows contribute for their improved delamination resistance.

Effect of manufacturing process

Composites manufactured by different moulding methods present different fracture toughness values. Interlaminar fracture toughness can be affected by both composite manufacturing process and specimen thickness [171]. Specimens manufactured by hand lay-up method showed decrease in $G_{IC}$ with an increase in specimen thickness. Contrary to this, specimens manufactured by resin transfer moulding technique showed opposite trend [171]. Composites manufactured by resin transfer moulding technique presented higher fracture resistance as compared to composites manufactured by hand lay-up. During testing, specimens undergo mixed mode loading and twisting due to the lack of reinforcement symmetry caused by lay-up moulding [172]. Carraro et al. [173] studied the effect of two different resin infusion methods on mode I toughness and damage mechanisms of carbon-epoxy laminates. Schober et al. [174] studied the effect of process parameters (consolidation pressure and temperature) on glass fibre-reinforced thermoplastic tapes. Both studies [173,174] reported a negligible influence of manufacturing process on mode I toughness. Kim and White [153] also reported a negligible effect of staged curing cycles on mode I toughness of composite laminates. Contrary to this, Hunt et al. [154] reported cure path-dependent mode I fracture toughness of particle interleaf laminated composites. Zhang and Fox [175] reported the influence of manufacturing method (Quick step curing and autoclave curing) on mode I interlaminar toughness of carbon-epoxy laminated composite. They reported that composites manufactured by quick step cure process have greater mode I toughness as compared to those manufactured by autoclave curing.

Specimen thickness

Mode I fracture characterization using DCB test exhibits large scale fibers bridging phenomena (when bridging zone length becomes comparable to length to specimen) [176]. In this scenario, R-curve depends on specimen geometry. Thus, mode I fracture toughness dependence on specimen thickness has gained particular interest of many researchers [149,176 –179]. Hojo and Aoki [180] investigated the effects of DCB thickness for carbon-PEEK composites and measured initiation $G_{IC}$ values in the range of 1100 to 1300 J m⁻², essentially, independent of the specimen thickness. In contrast to the initiation values, the higher propagation $G_{IC}$ values were reported with an increase in the specimen's thickness. Manshadi et al. [179] and Farmand et al. [149] also investigated the effect of specimen's thickness on mode I toughness in a glass fiber-polyester and carbon fiber-epoxy laminated composites. There is consensus that an increase in specimen thickness results an increase in fracture toughness, due to enhanced bridging zone length and the energy release arte (ERR) at the plateau level by increasing of specimen scaling.

Loading rate

Hug et al. [181], Smiley and Pipes [182], and Gillespie et al. [183] conducted mode I interlaminar fracture toughness tests on carbon-epoxy composites to find the influence of the loading rates on the interlaminar fracture toughness. They observed stable crack propagation independent of testing rates with a little effect of rate on $G_{IC}$ . Contrary to these, several works [126,184–190] report dependence of $G_{IC}$ on loading rate. Kusaka et al. [155] observed that carbon-epoxy laminated composites showed a negative rate-dependent fracture behavior (i.e. increasing loading rate results decrease in fracture toughness). An increase in loading rate changed the crack propagation behavior from an unsteady manner to a steady continuous growth, and consequently decreased $G_{IC}$ . Aliyu and Daniel [191] reported that in graphite-epoxy composite, increasing the loading rate results the increment in fracture toughness within a loading rate of 0.5 to 500 mm min⁻¹. In contrast, Smiley and Pipes [192] and Gillespie et al. [183] observed that graphite-PEEK showed negative rate-dependent fracture behavior at loading rate of 0.25–40,000 mm min⁻¹. This was due to the transition of ductile to brittle crack growth behavior with an increase in load rate. Mall et al. [192] also observed the negative rate-dependent fracture response of carbon-PEEK woven composites with unstable crack growth behavior at a range of loading rate from 0.005 to 1000 mm min⁻¹. It was explained by the decrease in plastic deformation zone of matrix beneath surface by increase in loading rate.

Temperature

Several researchers [193 –196] reported that mode I fracture toughness increases with temperature. This response has been attributed to the ductility of the matrix at higher temperature. However, some researchers [197,198] reported a complicated behavior of composites at different temperatures. Coronado and his co-workers [199] also analyzed the effect of temperature on mode I delamination in a carbon-epoxy unidirectional composite under static and fatigue loadings. During static tests, the material performed better at room temperature. During fatigue loading at 90^oC, the ductility of the matrix increased, which enhanced resistance to delamination during dynamic propagation [193]. They carried out another similar study at low temperature (−60^oC to 20^oC) and found that toughness decreases with decreasing test temperature [7].

Generally, the temperature-dependent fracture toughness of a composite is attributed to mechanical properties of matrix, its structure and fiber-matrix interfacial strength. Thus, still there has been much dispute about the mechanisms of temperature dependence of composites [200,201].

Summary and future directions

Summarizing the papers published on mode I facture toughness, significant progress on fracture characterization is noticed in the last decade. For laminated composites, standard test methods, including specimen size and shape, are published for fracture characterization but, for 3D-reinforced composites, researchers are trying to find standards. Many articles have been published on the effect of reinforcement geometry on fracture toughness and proposed ways to improve fracture by reinforcement deign.

We discussed the two main classes of toughening mechanisms, intrinsic and extrinsic, which are more and less related to the microstructure of materials. Traditionally, a single value of $G_{IC}$ for crack initiation is considered as material fracture toughness, which may or may not adequately capture the toughening of material as in case of multidirectional and 3D-reinforced composites where a prominent R-curve behavior is observed. In order to improve fracture toughness, extrinsic shielding approach is found to be effective. This is attained by designing microstructure in such a way that can offer a complex crack propagation path, and thus can dissipate more strain energy and less crack growth. In this regard, multidirectional laminated composites and 3D-reinforced composites are very effective. Moreover, for fracture characterization of 3D-reinforced composites, lack of standard test methods needs attention of researchers.

The various factors and mechanisms by which these influence the mode I fracture toughness of FRPCs' are reviewed. Loading rate, weave type and temperature have dominating effect on toughness. The effect of loading rate and temperature on mode I toughness is mainly attributed to the properties of matrix, its micro-structure and fiber–matrix interfacial strength. However, there are discrepancies on the mechanisms of influence, that need further attention of researchers.

Footnotes

Declaration of conflicting interests

The author(s) declared no potential conflicts of interest with respect to the research, authorship, and/or publication of this article.

Funding

The author(s) disclosed receipt of the following financial support for the research, authorship, and/or publication of this article: The authors acknowledge the financial supports from National Science Foundation of China (Grant Number 11572085 and 51675095). The Fundamental Research Funds for the Central Universities of China (2232018G-02) are also gratefully acknowledged.

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