Mechanical Performance of Carbon-Fiber-Reinforced Composite Textile Laminates Integrated with Graphene Nanosheets

Abstract

This chapter discusses the fabrication and mechanical characterization of nano-engineered composite laminates fabricated using variable-thickness graphene sheets incorporated in non-crimp carbon fiber prepregs. The effect of graphene sheet thickness on interlaminar strength (Mode I fracture toughness) of the carbon fiber composites was evaluated. The graphene lattice structure used in the present research had linear and square grids. Linear grids were arranged parallel and perpendicular to the 0° fibers in the composite laminates and labeled as vertical and horizontal grid patterns, respectively. Mechanical characterization involved the study of the effects of sheet thickness and grid pattern with and without nanoengineered enhanced laminates at the midplane. The composite laminates fabricated using a lattice graphene structure had better interlaminar strength than those fabricated with straight graphene sheets. Nanoengineered sheets with minimal thickness showed better interlaminar strength than the thicker sheets. The polymer used to manufacture the graphene sheet could not bond with the epoxy used in the composite laminate. In the literature, when the graphene nanoparticles are dispersed in the epoxy, the challenge is a uniform distribution of the nanoparticles. To overcome this dispersion problem, sheets made using nanomaterials can be used to enhance the mechanical properties of the composite laminates.

1. Introduction

Due to superior specific properties over traditional materials, the use of textile composites for aerospace, automotive, and marine applications has increased dramatically. Metal failure depends on the yield strength of the material, whereas the applied stresses influence the failure of composite material, the interaction between the two phases, the stacking sequence, and voids or defects in the composite laminate. Composite laminates are bonded together by a thin layer of resin between them. The function of this interface layer is to transfer the displacement and the force from one layer to another. As this matrix is a weaker element in the composites, the matrix fails at the beginning and then the fibers. When these layers get damaged or weakened, they form a crack between adjacent plies, separating the adjacent layers. This results in a failure of the lamination, which is known as delamination. The weak interlaminar region of each ply is the primary reason for the interlaminar-induced delamination of composite laminates. Researchers have developed various methods from different perspectives in diverse fields to overcome this challenge and enhance interlaminar strengths. The next-generation reinforcements, non-crimp fibers (NCF), are being explored for various structural applications. Non-crimp fabric provides excellent laminate strength, and the fabrication cost is lower than that of traditional composite manufacturing due to its drape ability and lack of fiber crimp in the woven fabric.

Various techniques such as three-dimensional weaving, stitching, Z-pinning, edge notching, braiding, fiber hybridization, and edge capping have been explored. These techniques suppress delamination but degrade in-plane laminate properties and incur additional manufacturing costs due to the process modifications. Degradation of in-plane properties and an increase in the weight of the composite laminate also lead to macroscale defects in the composite laminate, which may act as a crack initiation point for failure.

Tsai and coworkers studied the effect of delamination and tensile strength in the thick and thin ply composite laminate. To achieve a thickness of less than one-third of conventional plies (thin ply), a constant airflow through a sagged fiber filament technique was used to spread thick tows without damaging the fibers. Smaller ply thickness provides more choice in optimizing the laminate composite structure. The authors found that the tensile strength was higher for the thin-plies than the thicker-plies composite laminates due to the less resin-rich area used, but manufacturing thinner plies is more challenging [1]. With the advanced technology and concept of Prof. Stephen W Tsai’s theory, known as the CHOMARAT, which was founded in 1998, four industrial sites in France, Tunisia, the USA, and China, manufactures non-crimp fabric (NCF) [2].

Nanomaterials, such as carbon nanotubes (CNTs)/graphene nanoplatelets (GNPs) incorporated polymer composites, are becoming popular due to their excellent mechanical, thermal, and electrical properties. Adding a small portion of nanomaterials improves the mechanical properties of the composite laminates, such as their strength, stiffness, impact toughness, and electrical and thermal conductivity by several orders of magnitudes, and allows them to work with inner plies However, due to strong intermolecular interactions based on dipole interactions and van der Waals forces, non-functionalized nanomaterials form bundles or agglomerates. Agglomeration causes the collection of bundles of these nanomaterials together, distributing them unevenly in the matrix.

Nanomaterials can be aligned and dispersed in the polymer composite by applying external force. Various techniques, such as mechanical stretching, electrical fields, spinning processes, and magnetic fields, are used to overcome the agglomeration issue. Mechanical methods, such as ultrasonication (bath or probe), three-roll milling (calendaring) ball milling techniques, or a combination of these methods in series or parallel, are used to achieve dispersion. Each method has its pros and cons. High energy created during the sonication or milling process often damages the nanomaterials. So, surface modification techniques became famous for dealing with agglomeration problems. In these methods, the nanomaterials are treated with chemicals such as nitric, phenolic, and carboxylic groups to improve the interaction with the solvent, which leads to uniform dispersion and avoid agglomeration.

Various tactics are used to integrate carbon nanotubes or graphene into fiber-reinforced plastic structural composites to improve the mechanical properties of the composite laminate. Micro/nanoscale materials incorporation into the matrix resin to enhance the interlaminar properties have become popular compared to mechanical insertion due to both in-plane and out-of-plane enhancement of the properties per unit mass. Furthermore, even though mechanical insertion improves the interlaminar strength, it degrades in-plane properties due to cutting of the fibers during the insertion and adding mass to the composite laminate. To overcome the agglomeration problem, researchers have used nanoengineered interleaving sheets typically inserted at various locations through the thickness of the laminate [3].

In the present research work, graphene sheets of various thicknesses were incorporated into composite laminate and the interlaminar strength investigated by performing a double cantilever beam test. Graphene sheet-embedded composite laminates showed degraded interlaminar strengths, due to low bonding between the resin and graphene sheets. To overcome this problem, grids were formed on the graphene sheet, which helped to form the thin layer of resin at the midplane by enhancing the interlaminar fracture toughness of the composite laminates. After the test to check the bonding at midplane, the tested specimens were opened at midplane for microscopic analysis.

2. Literature background

Nanomaterials have superior specific (per unit mass) mechanical properties as compared to resins and fibers. A small amount of nanomaterial added to the composites increases the mechanical properties by a few orders of magnitudes. Carbon nanotubes (CNTs)/graphene nanoplatelets (GnPs)-incorporated polymer composites have become popular in composite industries. The carbon nanotubes (CNTs) improve mechanical properties such as strength, stiffness, impact toughness, and electrical and thermal conductivity of the composites.

Toughening of the matrix resin, interleaving with short fibers or micro/nanoscale particles, became popular in recent years to overcome delamination issues by avoiding mechanical inserts (pins). The addition of nanofillers leads to additional toughening mechanisms such as fiber pull-out, peeling/pull-out hackle formation, and stick-slip mechanism, which are controlled by: (i) the method by which the laminate is manufactured, (ii) the even dispersion of the nanofiller, and (iii) the matrix mixing method. Table 1 presents the techniques used to improve the interlaminar using different nanofillers with various results/laminar improvements reported by researchers in the last few decades.

Improvement techniques	Key results/improvements	Reference
The adhesive layer of 0.008″	Load-carrying capacity was increased by 50%, with a loss of in-plane strength.	Lagace et al. [4]
Graphene platelets	Mode I fracture toughness enhanced by ~53% over that of neat epoxy.	Rafiee et al. [5]
Alumina nanoparticles	Mode I fracture toughness was increased by 51% for non-functionalized alumina nanoparticles, while 74% with functionalized silica nanoparticles.	Kelkar et al. [6]
Electrospun nanofibers	Mode I fracture toughness of composites with electrospun nanofiber interleaving increased by 80% as compared to neat composite.	Kelkar et al. [7]
Multi-wall carbon nanotubes (MWCNTs), graphene nanoplatelets, and carbon blanks polymer nanocomposite	The interlaminar shear strength of the carbon fiber reinforced composite increased marginally.	Shrivastava et al. [8]
0.2 wt.% of highly dispersed rGO (reduced graphene oxide)	A 52% improvement in fracture toughness (KIC) of the cured epoxy.	Long-Cheng Tang [9]
0–7 wt.% GO with an increment of 1 wt.% as a function of PVP content	Mode I fracture toughness improved by ~100% for 7 wt.% GO.	Vaidyanathan et al. [10]
POSS-PVP (polyhedral oligomeric silsesquioxanes-polyvinylpyrrolidone) 1, 3, 5, and 10 wt.% to PVP	For 5 wt.% of POSS loading 70% increase in interlaminar fracture toughness.	Mishra et al. [11]
CNT concentration of 0.02% to 5%	0.047 wt.% caused an increase in the fracture toughness by 22%.	Han Zhang et al. [12]
Aligned CNTs in unidirectional carbon tape composite	The Mode I fracture toughness increased by 1.5–2.5 times and Mode II by three times.	Enrique and coworkers [12]
Buckypaper intervals carbon nanofibers unidirectional prepregs [0]12 fabricated by hand layup	ILSS increased by 31%, whereas Mode II interlaminar fracture toughness increased by 104%.	Khan et al. [13]
CF/Biaxial NCF/Toray T1700Sc 50C) 0.5 and 1 wt.% CNT premixing, shear mixing/degassing	Mode I fracture toughness increased from 428 ± 27 (pristine) to 462 ± 25 (0.5 wt.%) and 537 ± 34 J/m^2 (1 wt.%).	Quan et al. [14]
Polyphenylene-sulfide (PPS) veils doped with multiwall carbon nanotubes and GNPs	Fracture toughness of MWCNTs veils improved while the incorporation of GNPs on the veils decreased adhesion.	Quan et al. [15]
0.5 wt.% GO-epoxy nanocomposites + use of high pressure & shear rates	Fracture toughness increased up to 220% with adequate process.	Guijosa [13]
Graphene nanoplatelets (GnPs)	Fracture toughness was found optimum at 1 wt.% for higher loading, the properties deteriorate due to agglomeration of GnPs.	Valorosi et al. [16]
Laminates interleaved by multiwall carbon nanotubes loadings of 5, 10, and 15%	Mode I interlaminar fracture toughness increased by 206% as compared to the controlled sample.	Liu et al. [17]

Table 1.

Summary of the techniques used to fabricate nanoengineered composite laminates and their effect on the mechanical properties.

A key method of incorporating nanomaterials in composite processing is based on the use of vacuum assisted resin transfer molding method (VARTM). However, the use of nanomaterials in composite processing using VARTM presents a significant challenge due to the narrow gaps between fibers as nanoparticles filter out during infusion as depicted in Figure 1. To overcome these challenge, researchers utilize advanced three-roll mills, centrifugal mixers, functionalization, sprayers, etc., for uniform dispersion of nanoparticles into the resin to produce nanomodified prepregs [19].

Figure 1.

Filtering of larger particles by fiber array and infiltration of smaller particles between fibers [18].

Furthermore, the use of nanofillers poses additional challenges of nonuniform dispersion resulting from nano-agglomeration, which degrades the properties of the laminate. The advanced method of inserting the carbon nanotube buckypaper was used by various researchers with promising results. Wang et al. [20], for example, studied glass fiber reinforced polymer composite with and without buckypaper. They reported 25% more shear strength due to incorporating carbon nanotubes (CNTs) buckypaper. Li et al. [21], on the other hand, studied the effect of curing pressure on glass fiber reinforced plastic (GFRPs) laminates fabricated by adding CNTs buckypaper to the midplane layer resulting in fracture toughness increments of 174.81% for initiation and 179.97% for propagation relative to baseline composite for 2 MPa curing pressure. Then, Chen et al. [22] investigated fracture toughness by incorporating buckypaper at the midplane of reinforced fiber and found the initiation interlaminar fracture toughness was increased by 29% compared to that of the one without the buckypaper composite laminate. These studies demonstrated that the interlaminar fracture toughness of fiber-reinforced plastics could be significantly enhanced by incorporating CNTs buckypaper interleaf. The researchers used nanomaterials with a small percentage of the composite and the nanoengineered micro thickness sheets were incorporated into composite laminates.

On the other hand, the use of graphene sheet interleaf is not found in the literature. So, in the present study, the graphene sheets with 50, 120, and 240 μm thickness were incorporated at midplane in order to study the initiation and propagation fracture toughness of quasi-isotropic layered non-crimp carbon fibers. Straight graphene sheets of 50, 120, and 240 μm thickness incorporated composite laminates degraded the initiation interlaminar fracture toughness by 75, 92, and 86% compared to the composite laminate without graphene sheet. To overcome this problem, uniform grids were formed on the graphene sheet, aligned parallel or perpendicular to 0^o fibers at the midplane, and labeled as either vertical or horizontal grids in the present research work.

3. Materials and the process

3.1 Procurement and characterization of graphene sheets

Graphene sheets of three different thicknesses: 50, 120, and 240 μm (PN40003, PN40008, and PN 40009) are supplied by XG Sciences Inc, Lansing, MI, USA, were used as interleaving material. SEM images of the graphene sheet, as shown in Figure 2, depict the two-dimensional graphene nanoplatelets (GnPs).

Figure 2.

SEM images of the graphene sheet.

3.2 Panel fabrication from NCF-MTM45-1 Prepreg

Almost all high-tech composite companies manufacturing aerospace and sporting goods use prepreg because of the high quality of fiber alignment and uniform fiber volume fraction. In the present study, NCF-MTM45-1 pre-impregnated prepregs supplied by SHD composites, NC, USA, were used to fabricate nanoengineered composite laminates. The graphene sheets were employed as interleaf at the midplane of the composite laminate to analyze the fracture toughness at the initiation and the propagation point. Other supporting materials, including glass mold, vacuum bag, release film, breather, and sealant tape, were assembled as shown in Figure 3.

Figure 3.

Fabrication set up for nanoengineered composite laminate.

The NCF composite laminates were fabricated using 48 layers (24 bilayers) of biaxial spread-tow carbon NCF fabric prepreg. Each NCF prepreg layer had [0/−45] orientation. These 24 bilayers were symmetrically stacked about midplane with the orientation of [0/−45/90/45] to achieve a balanced layup in the final cured laminate. Hand-cutting fibers produced the composite laminate with dimension 12″ × 12″, with layup [(0/−45/90/45)₆/GS/(45/90/−45/0)₆] [22, 23]. Figure 4a depicts the debulking, using vacuum pressure to remove the entrapped air in the prepregs. Then, the 12 layers of biaxial spread-tow carbon fibers were placed on the mold, and a Teflon sheet of 14″ × 3″ was inserted at one end such that the 14″ side is perpendicular to 90^o fibers in the laminate, as shown in Figure 4b and c. The graphene sheet (either 50, 120, or 240 μm) was incorporated next to the Teflon sheet at the midplane to produce nanoengineered composite laminate. The fabrication process included composite laminate, with and without nanoengineered graphene as depicted in Figure 4d [2, 23, 24].

Figure 4.

Fabrication—nanoengineered composite laminate. (a) Debulking of Prepregs. (b) NCF prepreg layer preparation. (c) NCF prepreg layer arrangement. (d) GS at midplane.

Weak bonding between the matrix and graphene resulted in degraded interlaminar resistance causing delamination and separation at the midplane when manual force was applied. This was then corrected by perforating the graphene sheets forming a lattice structure.

3.3 The curing cycle of the process

The curing cycle for the process is shown in Figure 5. After oven curing, the coupons were waterjet cut as shown in Figure 6a, and tested as per American Society for Testing and Materials (ASTM) D5528 standards [25]. The edges were made smooth and uniform using the grinding machine to achieve uniform width, as depicted in Figure 6b.

Figure 5.

Cure and post cure time-temperature cycle (MTM 45-1).

Figure 6.

(a–c) Preparation and DCB test. (a) Waterjet cutting, (b) edge grinding of coupons and (c) DCB Test.

3.4 Fracture toughness (G_IC)

Strain energy release rate G, expressed as Mode I interlaminar fracture toughness in ASTM D 5528 standards is given as:

$$G_{IC}=\frac{3P\delta }{2b\left(a+\text{l}\Delta \text{l}\right)}$$E1

where P = load, δ = load point displacement, b = specimen width, a = delamination length and Δ is constant, determined experimentally by generating the least-square plot of the cube root of compliance, $C^{1/3}$ as a function of delamination length. Compliance C is the ratio of crack opening displacement to the applied load, δ/P. For more details, readers are referred to appropriate literature sources [25]. Instron electromechanical system is used to record the load and displacement at every 10 s after test analysis was done to calculate fracture toughness and coupons were separated about midplane to observe the midplane bonding using microscopic images.

The present work uses the modified beam method to calculate fracture toughness as per ASTM 5528 standards.

4. Results and discussion

Strain energy release rate (G_IC) values for with- and without-graphene sheet reinforced-composite laminate were calculated using Eq. (1) and the results are shown in Figure 7. The initiation and propagation values of strain energy release rate (G_IC) with brittle graphene-epoxy for graphene sheet reinforced-composite reduced initiation fracture toughness by 75, 92, and 86% for 50, 120, and 240 μm thick graphene sheets, and similar trends observed for propagation fracture toughness [26].

Figure 7.

Initiation and propagation of GIC with and without graphene sheet.

Weak bonding between the matrix and graphene resulted in degraded interlaminar resistance. Because of this, the specimens were easily delaminated and separated at the midplane when manual force was applied, as shown in Figure 8. To overcome this challenge, the graphene sheet was perforated by forming a lattice structure.

Figure 8.

Testing of coupon manually. (a) Waterjet coupon, (b) manually applied force and (c and d) separation of the laminate at the midplane.

The lattice grids were developed with 5 mm spacing horizontally, vertically, and in square form, as shown in Figure 9. The gap between the graphene sheets helped to form the fiber-epoxy bonding.

Figure 9.

Lattice structure formed on a graphene sheet.

Figure 10 represents the delamination resistance curve (R curve) for the horizontal, vertical, and square grid embedded graphene sheet at midplane in the composite laminate. It then compares it with the composite laminate without a graphene sheet. Delamination resistance for composite laminate without nanoengineered graphene sheet was dominant as compared to the plain and lattice structure graphene sheet-embedded composite laminates. The lattice structure formed on the graphene sheet helped to improve the delamination resistance of the composite laminate.

Figure 10.

Delamination resistance curve (R curve) for all GS.

Figure 11 compares the fracture toughness of the 24 bilayers, NCF carbon fiber with and without incorporated lattice structure graphene sheet at the midplane with bare composites. Irrespective of the thickness, the vertical grids (grids parallel to 90° fibers in the laminate) had better performance than horizontal and square grids. The graph for the horizontal grids showed a sudden change in values due to graphene strips, because of the alternating resin patch.

Figure 11.

Fracture toughness for all samples.

The fracture toughness of the vertical grid 50 μm graphene sheet was the highest and was almost the same as the baseline composite laminate. The fracture toughness of the square grids was observed to be less than that of the other grid patterns. The fracture toughness of a straight 120 μm-thick graphene sheet was the lowest compared to those of similar groups.

Sample images for the above configurations are shown in Figure 12a and b.

Figure 12.

(a and b) Images of the coupons after the DCB test. (a) 50 μm straight GS and (b) 50 μm horizontal grid.

For the straight graphene sheet-incorporated panels, it was observed that there was very little binding between the graphene sheet and adjacent plies. This degraded its interlaminar fracture toughness compared to the other configurations. In straight graphene composite laminates, laminates fabricated using 50 μm-thick graphene sheet-embedded showed higher fracture toughness when compared to that of the 120 and 240 μm thick graphene sheet-embedded composite laminate. This could be due to the smaller thickness of the sheet. Also, the vertical lattice structure on 50 μm sheet showed a fracture toughness of 0.28 kJ/m² as compared to 0.18 and 0.24 kJ/m² when 120 and 240 μm vertical lattice structures graphene sheets were used, respectively, as the matrix formed a thin layer in between the lattice structure, which helped to improve the interlaminar resistance. As vertical grids were arranged parallel to 0° fibers in the composite laminates, which formed the thin alternate layer of epoxy along the length of the DCB test specimen, it helped to increase the interlaminar strength of the composite laminate. The grids are aligned horizontally perpendicular to the 0^o fibers in the composite laminate, which forms alternate thin layers of the matrix between two strips of graphene sheet along the length of the DCB test specimen, resulted in instability that was observed for delamination resistance (in the delamination curve) of horizontal grid pattern of graphene sheet embedded composite laminates.

To maintain uniformity in the fabrication process, horizontal-grid-formed sheets were arranged in a manner that a graphene sheet follows the same gap after the Teflon at midplane. Experimentally, maximum fracture toughness was observed for vertical grids. Furthermore, all the vertical grid samples exhibited fiber pullout, which showed bonding between the top and bottom layers at midplane, and this was not observed in the specimens with other grid patterns. Khan and Kim [13] used buckypaper intervals and reported a 31% increase in the Mode II interlaminar shear strength while reporting doubled fracture toughness, the cost of making CNT buckypaper is way too expensive compared to the commercial graphene sheets used in the present study. Also, Kelkar et al. [6] reported an 80% increase in interlaminar strength as compared to neat composite using electrospun fibers.

5. Conclusion

Graphene sheets-embedded composite laminate was used to incorporate the nanoparticles and avoid the nanoparticles’ uniform dispersion problem in the epoxy resin. Mode I fracture toughness (strain energy release rate) studies indicated that when graphene sheets were inserted in the midplane of the carbon fiber composite laminate, the results for fracture toughness were poorer as there was no bonding between the graphene sheet and adjacent prepregs. Therefore, the graphene sheets were converted into three different lattice configurations, including horizontal, vertical, and square structures, so as to alleviate this problem. Grids formed on the graphene sheet showed enhanced properties and vertical grids showed peak mechanical properties as compared to horizontal and square grids. The addition of grids helped to form a tiny epoxy layer and to improve the fracture toughness of the 50 μm lattice structure sheet, more than that of the composite without a graphene sheet. Vertical grids formed on 50, 120, and 240 μm thick graphene sheets showed improvement in the fracture toughness by 4.5, 10, and 7 times compared to the straight graphene sheet, but the values were always less compared to composite without graphene sheet. Overall use of two-dimensional nanoengineered graphene sheet showed degraded interlaminar strength as compared to the composite laminate without a graphene sheet.