CARBON 5 0 ( 2 0 1 2 ) 1 3 1 6 –1 3 3 1 Available at www.sciencedirect.com journal homepage: www.elsevier.com/locate/carbon Effects of carbon nanofiller functionalization and distribution on interlaminar fracture toughness of multi-scale reinforced polymer composites Ye Zhu a,* , Charles E. Bakis a, James H. Adair b a Department of Engineering Science and Mechanics, The Pennsylvania State University, 212 EES Building, University Park, PA 16802, United States b Department of Material Science and Engineering, The Pennsylvania State University, 108 Steidle Building, University Park, PA 16802, United States A R T I C L E I N F O A B S T R A C T Article history: Carbon nanofillers with different surface functional groups and aspect ratios, including Received 8 July 2011 carboxyl carbon nanotubes, un-functionalized carbon nanofibers (CNFs), glycidyloxypro- Accepted 1 November 2011 pyl-trimethoxysilane carbon nanotubes (GPS-CNTs) and nanofibers were evaluated for Available online 9 November 2011 their potential for increasing the interlaminar fracture toughness of an S2-glass fiber/epoxy composite. The fillers were added in the matrix of the fiber reinforced plies, in the resin interlayer between plies, or in both regions. Comparisons were made based on mode I and mode II interlaminar fracture toughness. For composites made with CNTs dispersed in the matrix, fracture toughness was largely unaffected except for a slight increase seen with long GPS-CNTs. However, adding a CNF or CNT modified resin interlayer significantly increased the fracture toughness, with the highest improvement over the baseline material achieved by adding long GPS-CNTs in the interlayer (79% and 91% for mode I and mode II onset toughness, respectively). Important material parameters identified for improving interlaminar fracture toughness are the nanofiller aspect ratio and concentration at the fracture plane. Based on microscopic evaluations of the fracture surfaces, a high density of high aspect ratio nanofillers causes the best entanglement between the filler and glass fibers and effectively obstructs interlaminar crack propagation. 2011 Elsevier Ltd. All rights reserved. 1. Introduction Continuous fiber reinforced polymer (FRP) laminates have numerous applications in high performance structures such as aircraft, boats, automobiles, wind turbines, and sporting goods. Attractive in-plane mechanical properties are obtained by using continuous in-plane fiber reinforcement. On the other hand, the out-of-plane mechanical properties of fiber reinforced composites are not comparably superior to competing materials due to the lack of continuous fiber reinforcement in this direction [1,2]. One approach for improving the out-of-plane properties of FRP composites focuses on modifications of the polymeric matrix. The out-of-plane properties of FRP composites are often characterized in terms of the interlaminar fracture toughness [3,4], which is related to the fracture toughness of the polymer matrix along with other factors such as fiber–matrix adhesion. The fracture toughness of epoxy resins used as matrix materials for FRP composites can be increased by the addition of micro-sized fillers of high [5,6] or low [7] modulus of elasticity. The toughening mechanisms of these fillers have been postulated to be crack pinning, crack deflection, crack bowing, * Corresponding author: Fax: +1 814 863 6031. E-mail address: yzhupsu@gmail.com (Y. Zhu). 0008-6223/$ - see front matter 2011 Elsevier Ltd. All rights reserved. doi:10.1016/j.carbon.2011.11.001 CARBON 5 0 ( 20 1 2 ) 1 3 1 6–13 3 1 crack front trapping and cavitation [5–10]. Fillers used as mechanical reinforcement in a polymer are usually loaded beyond the percolation threshold, at which concentration the fillers form an interconnected three-dimensional network occupying the whole volume of the matrix material. Carbon nanofibers (CNFs) and carbon nanotubes (CNTs) are particularly well suited for this purpose on account of their nanoscale diameter and high aspect ratio (>100), which promotes percolation at low volume fractions. With a combination of superior properties, including extremely high specific surface area, axial stiffness and strength, only a small amount of CNTs or CNFs are needed to improve the mechanical properties of polymers [10–16], including fracture toughness. It has been proposed that CNTs and CNFs are potentially ideal fillers for FRP composites, as well as neat polymers [17,18]. However, the addition of these nanofillers does not always improve the modulus and fracture toughness of FRP composites significantly [18–20]. For multi-scale reinforced composites, many processing and filler parameters exist and their individual roles in controlling mechanical properties are not well characterized, to-date. These parameters include, but are not limited to, the nanofiller dispersing method, distribution, filler aspect ratio, and surface functionality. Two possible reasons for the reduced reinforcing effect of nanofillers in multi-scale reinforced composites are as follows. Firstly, it is challenging to control the dispersion and distribution of nanofillers in FRP composites. The flow of resin causes the migration of nanofillers along and across micro-channels between microfibers, which promotes nanofiller aggregation, entanglement with microfibers, and shear rate induced segregation. Secondly, depending on the method of introducing nanofillers into an FRP composite, the nanofillers are not necessarily located where they are most needed to improve the mechanical properties, such as in between plies or at the fiber–matrix interface. On the other hand, the dispersion of CNTs in two-phase CNT/polymer composites is often improved by functionalization of the CNTs [21,22]. The mechanical properties of CNT/polymer composites were improved by CNT functionalization as well [23]. The effects of functionalization on filler dispersion and mechanical properties of multi-scale reinforced composites are still under investigation [14,19,24,25]. To-date, no investigation has compared the effects of filler functionality, aspect ratio, and distribution on interlaminar fracture toughness of an FRP material system made by a technique suitable for large scale structural applications. In the present investigation, an S2-glass fiber/epoxy laminated composite made by the filament winding method was used as the baseline material. The baseline material system was modified using nanofillers with different surface functional groups (carboxyl and glycidyloxypropyl-trimethoxysilane, GPS) and with different ranges of aspect ratio by using one or both of two placement approaches: (1) fillers uniformly dispersed in the matrix and (2) fillers added in the interlayer region between plies. The scientific rationale for surface functionalization of filler material with silane coupling agents is to match the surface energy of the filler material to the polymer matrix. Reed has recently reviewed this approach [26]. For example, the critical surface tensions for glycidylether and acid hydrolyzed glycidoether functional groups are 39 and 44.6 J/m2, 1317 respectively [27]. Thermal stability and resistance to chemical and mechanical degradation during processing are also important criteria for silane coupling agents [26]. However, the surface tension of the polymer matrix changes with processing conditions related to temperature and mechanical shear. Thus, the surface functionalization effects on processing and performance are usually experimentally determined. The goal of the present investigation is to shed light on criteria for selecting the nanofiller functionality, filler size or aspect ratio and filler placement method in order to achieve optimal improvement in the interlaminar fracture resistance of multi-scale reinforced composites. 2. Specimen preparation and testing 2.1. Materials 2.1.1. Baseline material Unidirectional fiber reinforced laminates were made by wet filament winding method using type 449AA-750 S-2 Glass fiber (AGY, Aiken, SC) and an epoxy resin matrix system. The baseline resin system is a bisphenol-A based epoxide diluted with alkyl glycidyl ether, EPONTM 8132 (Momentive Specialty Chemicals, Columbus, OH). The curing agent is a polyether amine, Jeffamine T403 (Huntsman Performance Products, The Woodlands, TX). The curing agent was added to the epoxide in a ratio of 40:100 by weight. An air-release additive, BYKA 501 (BYK Chemie, Wallingford, CT), was added at 0.5 wt.% to the catalyzed resin mixture to facilitate the release of bubbles during fiber impregnation. The tensile properties of the baseline resin without nanofillers were obtained following the ASTM standard D638 using dogbone specimens [28] (Table 1). 2.1.2. Carbon nanotube/nanofiber Three types of multi-walled CNTs were obtained from CheapTubes.com (Brattleboro, VT): short carboxyl-functionalized CNTs (COOH-CNTs), short hydroxyl-functionalized CNTs (OH-CNTs), and un-functionalized vertical grown CNTs (VGCNTs). Un-functionalized heat treated carbon nanofibers (UF-CNFs) were obtained from Applied Sciences, Inc. (Cedarville, OH). The vendor specifies that the short OHCNTs contain 3.06 wt.% hydroxyl groups and the COOHCNTs contain 2 wt.% carboxyl groups and 1 wt.% hydroxyl groups. The VGCNTs and UF-CNFs were of considerably higher as-received aspect ratio than the COOH-CNTs and OH-CNTs. Specifications for the nanofillers, obtained from the manufacturer, are provided in Table 2. VGCNTs, UF-CNFs, and OH-CNTs were used as raw material for the in-house production of silane-functionalized nanofillers. First, the VGCNTs and UF-CNFs were oxidized by refluxing 0.5 g of either filler in 100 ml 40% nitric acid for 4 h at 80 C. The oxidized fillers were washed with excessive distilled water and filtered. The silane coupling agent used to functionalize the oxidized VGCNTs and CNFs and the as-received OH-CNTs is (3-glycidyloxypropyl)trimethoxysilane, referred to as GPS. According the vendor (Sigma–Aldrich, St. Louis, MO), the GPS has a purity of >97% and has the chemical structure shown in Fig. 1. In the GPS treatment procedure, 15 g of GPS, 15 g of methanol, 0.15 ml glacial acetic acid, and 0.75 ml deionized water were first mixed together. Then, 4 ml of GPS mixture 1318 CARBON 5 0 ( 2 0 1 2 ) 1 3 1 6 –1 3 3 1 Table 1 – Mechanical properties of baseline resin system determined from dogbone tests based on the average value of three specimens. Values in parentheses are the coefficient of variation in percent. Young’s modulus, GPa Poisson’s ratio 2.43 (3.1) Ultimate strength, MPa 0.36 (0.3) Ultimate strain,% 47.6 (2.8) 5.43 (36) Table 2 – Name and manufacturer’s specifications of as-received carbon nanofillers. Name in this investigation Length, lm Outer diameter, nm Inner diameter, nm Initial aspect ratio, length/diameter Short COOH-CNT Short OH-CNT VGCNT UF-CNF 0.5–2 0.5–2 10–50 30–100 10–20 10–20 8–15 60–150 3–5 3–5 3–5 N/A 25–200 25–200 660–6250 200–1670 Fig. 1 – Chemical structure of (3-glycidyloxypropyl)trimethoxysilane (Sigma–Aldrich product information. (3-Glycidyloxypropyl)trimethoxysilane http://www.sigmaaldrich. com/). was added to 40 ml of toluene. About 0.1–0.2 g of oxidized CNTs or CNFs was added to 44 ml of toluene/GPS mixture. The nanoparticle mixture was stirred at room temperature for 60 h. The GPS functionalized CNTs (GPS-CNTs) or GPS functionalized CNFs (GPS-CNFs) were washed with excessive toluene and then excessive methanol, filtered, and finally dried in a vacuum oven at 80 C for 4 h. The GPS-CNTs obtained by the above method were subjected to chemical analysis to verify Manufacturer’s product name Short COOH CNTs Short OH CNTs MWNT Arrays Pyrograf-III PR-24 HT the successful attachment of desired functional groups. Short GPS-CNTs, long GPS-CNTs, and GPS-CNFs were used for manufacturing multi-scale reinforced composites. Additionally, multi-scale composites were also manufactured using COOH-CNTs and UF-CNFs in the as-received condition. All of the functionalized nanofillers used to manufacture the multi-scale composites are summarized in Table 3. 2.1.3. Carbon nanotube/nanofiber modified composite material It is known that the nanofiller loading will potentially have an effect on the fracture toughness and viscosity of the nanofilled resin. The filler loading used in this investigation was selected based on the information available in the literature and results from previous investigations by the authors. According to Ma et al. [23], the fracture toughness of a silane functionalized CNT-modified epoxy increases as the concentration of CNT increases. A saturation point is reached at a CNT loading of 0.5 wt.%. A silane-functionalized filler loading of 0.5 wt.% Table 3 – Functionalized nanofillers used for making multi-scale reinforced composites. Name in this investigation Short COOH-CNT Short GPS-CNT Long GPS-CNTs GPS-CNFs Functional group Raw material for functionalization COOH–, OH– 3-Glycidyloxypropyl silane (Fig. 1) 3-Glycidyloxypropyl silane (Fig. 1) 3-Glycidyloxypropyl silane (Fig. 1) Used as-received Short OH-CNTs (Table 2) VGCNTs (Table 2) UF-CNFs (Table 2) Table 4 – Compositions of resin systems used for manufacturing multi-scale composites and corresponding bath sonication times and temperatures. Compositiona Resin system Baseline epoxy 1 wt.% short COOH-CNT epoxy 0.5 wt.% short GPS-CNT epoxy 0.25 wt.% long GPS-CNT epoxy 0.5 wt.% UF-CNF epoxy 0.5 wt.% GPS-CNF epoxy a Bath sonication CNTs CNFs – 1 wt.% 0.5 wt.% 0.25 wt.% – – – – – – 0.5 wt.% 0.5 wt.% Weight percentages are in terms of the resin and curing agent mixture. Time (h) Temperature (C) – 5 8 8 4 2 – 80 60 60 60 60 CARBON 1319 5 0 ( 20 1 2 ) 1 3 1 6–13 3 1 was therefore selected for the current investigation. From a previous investigation by the authors [29], it is known that adding 1 wt.% short COOH-CNTs in a fiber reinforced composite results in a higher mode I and mode II fracture toughness than adding 0.5 wt.% short COOH-CNTs in the same composite. Hence, 1 wt.% was selected as the loading for short COOHCNTs in the present investigation. To achieve the desired loading of nanofillers in the cured epoxy containing epoxide and curing agent, 1.4 times the desired weight percent of CNT or CNF in the cured epoxy was added to 15–100 g of epoxide since the desired epoxide and curing agent mixture ratio is 100:40. For example, if the desired CNT loadings in the cured epoxy were 0.5 wt.% and 0.25 wt.%, the loadings of CNT in the epoxide were 0.7 wt.% and 0.35 wt.%, respectively. The mixture was magnetically stirred at 160–240 rpm for 15 min. The nano-filled epoxide was sonicated in an ultrasonic bath operating at 45 W and 38.5 kHz for 2–8 h as specified in Table 4. At one hour intervals during sonication, the resin mixture was removed from the bath and magnetically stirred on a hot plate for 5 min to homogenize the mixture. After the nanofillers were dispersed into the epoxide at 0.7 wt.% or 1.4 wt.% using a bath sonication method, a small drop of CNT/epoxide mixture was deposited on a glass slide and a cover glass slide was placed on top of the mixture. Photographs of thin films of nanofilled epoxide mixtures considered for the current investigation are shown in Fig. 2. After dispersing 0.7 wt.% long GPS-CNT in the epoxide, the GPS-CNT/epoxide mixture was highly viscous and dispersion was not uniform (Fig. 2d). Hence, of all the formulations originally considered, only the 0.5 wt.% long GPSCNT/epoxy was considered to be unsuitable for fabricating glass fiber composites using the wet filament winding method. The loading of long GPS-CNT in epoxy (epoxide and curing agent) was therefore adjusted to 0.25 wt.% for making filament wound glass fiber composites (Fig. 2c). The compositions of nanofilled resin used in this investigation are summarized in Table 4 along with bath sonication time and temperature. After sonication, the nanofilled epoxide mixture was magnetically stirred at 120–200 rpm for 15 min and calculated amounts of curing agent and air-release agent were then stirred into the mixture. The baseline and nanofilled resins were degassed for 30 min before being used to make fiber reinforced composites. 2.2. Fabrication of toughness specimens composite interlaminar fracture Unidirectionally reinforced composite sheets were manufactured by wet-winding 10 layers of impregnated S2-glass fiber tow onto a flat mandrel as shown in Fig. 3a. While still on the mandrel, the 356 by 305 mm impregnated sheets were consolidated in a press at room temperature. Two 10-layer sheets of impregnated material were removed from the mandrel and stacked in an aluminum mold with their fibers parallel to each other, with a 0.0127-mm-thick PTFE film (DuPont, Wilmington, DE) placed over a portion of the plate at the midplane to serve as a starter crack. The composite was cured in a hot press held under a pressure of 240 kPa at 80 C for 2 h and then 125 C for 3 h. The fiber volume fraction of the cured composite was determined to be 62 ± 2% according to procedure G in ASTM standard D3171 [30]. Schematics of the composite layup arrangement are shown in Fig. 3b. The baseline composite was made using the baseline epoxy resin in all 20 layers. The multi-scale reinforced composites were made by using one or both of the following two approaches. In the first approach, the modified Fig. 2 – Photographs of nanofiller/epoxide thin films collected on a glass slide after ultrasonic dispersion with cover glass slide on top of the liquid mixture (images and weight fractions pertain to the mixtures before adding curing agent): (a) 1.4 wt.% short COOH-CNT in epoxide, (b) 0.7 wt.% short GPS-CNT in epoxide, (c) 0.35 wt.% long GPS-CNT in epoxide, (d) 0.7 wt.% long GPS-CNT in epoxide, (e) 0.7 wt.% UF-CNF in epoxide and (f) 0.7 wt.% GPS-CNF in epoxide. 1320 CARBON 5 0 ( 2 0 1 2 ) 1 3 1 6 –1 3 3 1 Fig. 3 – Schematic showing (a) filament winding process, (b) fiber reinforced composite layup and (c) interlaminar fracture toughness specimen configuration. was uniformly spread onto the wet prepreg using a Nylon roller. A water-cooled diamond abrasive saw was used to cut interlaminar fracture coupons from the cured plates. A typical interlaminar fracture specimen configuration is shown in Fig. 3c. Fig. 4 – DCB test setup. resin was used in the resin bath to impregnate the middle six layers symmetrically disposed about the midplane of the laminate. In the second approach, a calculated amount of nanofilled resin, equivalent to the weight of resin in two plies of cured composite plate (14 g), was added to the surfaces of both sheets facing the midplane to serve as a nanofilled interlayer region on the fracture plane. The nanofilled interlayer 2.3. Experiments 2.3.1. Carbon nanotube characterization GPS functionalization of CNTs was characterized using thermogravimetric analysis (TGA) and X-ray photoelectron spectroscopy (XPS). Thermogravimetric analyses of VGCNTs, short OH-CNTs, and short and long GPS-CNTs were performed using a TA Instruments thermogravimetric analyzer Q500. Typical sample mass ranged from 5 to 10 mg. Samples were analyzed in platinum pans at a heating rate of 15 C/min to 800 C in an atmosphere of air flowing at 40 ml/min. X-ray photoelectron spectra of as received VGCNTs, oxidized VGCNTs, and GPS functionalized VGVNTs were taken using a Kratos Ultra X-ray photoelectron spectrometer with monochromatic Al ka radiation under high vacuum (<1 · 107 Torr). The survey CARBON 5 0 ( 20 1 2 ) 1 3 1 6–13 3 1 1321 Fig. 5 – ENF test: (a) specimen geometry and (b) test setup. Fig. 6 – TGA weight loss results for as received VGCNTs, short OH-CNTs, short GPS-CNTs (starting material is short OH-CNT), and long GPS-CNTs (starting material is VGCNT): (a) weight loss curves and (b) first derivative of weight loss curves. scan of each sample was conducted in the binding energy range of 0–1350 eV using a spectrometer pass energy of 80 eV, step size of 0.5 eV and dwell time of 150 ms. High resolution scans were conducted on each sample in the binding energy ranges of 524–542 eV for O 1s, 275–300 eV for C 1s,and 95– 113 eV for Si 2p, using a spectrometer pass energy of 20 eV, a step size of 0.1 eV, and a dwell time of 3000 ms. All binding energies were referenced to carbon 1s at 284.8 eV. 2.3.2. Interlaminar fracture toughness (IFT) testing Mode I IFT tests were conducted using double-cantilever beam (DCB) specimens prepared according to ASTM standard D5528-01 [31]. Three repetitions of the DCB test were conducted for each type of composite material (except for the 0.25 wt.% long GPS CNT modified composite, for which two repetitions were conducted). The DCB test setup is shown in Fig. 4. The length, width, and thickness of the DCB specimens were approximately 150 · 25 · 3.5 mm. The initial crack length of the DCB specimens was approximately 45 mm. The DCB specimens were loaded through piano hinges using a servo-hydraulic load frame in stroke control (1 mm/min). Applied force was measured using a 110 N load cell and displacement was measured on the actuator. Crack propagation length was measured using an instrumented long distance microscope. The modified compliance calibration (MCC) method as specified in [31] was used for compliance calibration of the DCB specimen and for calculating the mode I IFT. Mode II IFT tests were conducted according to JIS K7086 [32] using end notched flexure (ENF) specimens loaded in 3-point flexure as shown in Fig. 5. Three repetitions of the ENF test were conducted for each type of material. The length, width, and thickness of the ENF specimens were approximately 150 · 25 · 3.5 mm and the initial crack length, a0 was approximately 30 mm. The specimen was loaded in stroke control using a servo-hydraulic test frame and loading point displacement was measured on the actuator. Load was measured by a 14.7 kN load cell using a 2.2 kN calibrated load range. Compliance calibration tests were conducted to obtain the compliance vs. crack length relationship for a given specimen similar to the methods described in [33]. For this purpose, six initial crack lengths (approximately a0 + 9, a0 + 6, a0 + 3, a0, a0 3 and a0 6 mm) were created by longitudinally offsetting the specimen in the bending fixture. Crack length in the compliance calibration tests was measured by an instrumented long-distance telescope. Care was taken to prevent crack onset during the compliance calibration procedures. Crack propagation length (a1) was calculated using the compliance calibration relation in the form of Eq. (1). " #1=3 3 Cð8bh Þ A a1 ¼ B where b and 2h are the specimen width and thickness, respectively (Fig. 5a). The compliance of the specimen, C, 1322 CARBON 5 0 ( 2 0 1 2 ) 1 3 1 6 –1 3 3 1 Fig. 7 – XPS spectrum analysis of VGCNT, oxidized CNT, and GPS functionalized VGCNT samples: (a) survey scans and (b) high resolution scans. Table 5 – Concentration of elements in VGCNT, oxidized VGCNT, and GPS functionalized VGCNT samples. VGCNT type Unfunctionalized Oxidized VGCNT GPS functionalized VGCNT Atomic concentration (%) C O Si 99 89 89 1.2 8.3 9.2 0.21 2.3 2.2 was determined by the measured displacement to load ratio (d/P) at the loading point. Parameters A and B were determined experimentally for each ENF specimen by the intercept and slope of the straight line fitted to the data points in a C(8bh3) vs. a3 plot by linear least squares. The compliance calibration (CC) method as specified in [33] was used for calculating the mode II IFT by Eq. (2). GII ¼ 3P2 B 2 a 2b 8bh3 O/C ratio 0.012 0.093 0.10 O/Si ratio 5.5 3.6 4.2 The critical load and displacement for crack onset was determined to be the load and displacement where the specimen compliance had increased by 5% compared to the initial compliance. 2.3.3. Microscopy Fracture surfaces of interlaminar fracture toughness specimens were examined within 1 cm distance from initial crack front using a Philips XL 30 (FEI, Hillsboro, OR) scanning CARBON Matrix: 1 wt% short COOH-CNT epoxy Baseline epoxy 1323 5 0 ( 20 1 2 ) 1 3 1 6–13 3 1 0.25 wt% long GPS-CNT epoxy 0.5 wt% short GPS-CNT epoxy 700 670 616 2 Mode I fracture resistance, J/m 643 600 603 500 400 300 200 Specimen 1 Specimen 2 Specimen 3 Nonlinear fit 127 100 133 Specimen 1 Specimen 2 Specimen 3 Nonlinear fit 136 Specimen 1 Specimen 2 Specimen 3 Nonlinear fit 132 Specimen 1 Specimen 2 Nonlinear fit 0 0 10 20 30 0 10 20 0 30 10 20 30 0 10 20 30 Crack propagation length, mm (a) Mode I Matrix: 1 wt% short COOH-CNT epoxy Baseline epoxy 1000 0.5 wt% short GPS-CNT epoxy 0.25 wt% long GPS-CNT epoxy Mode II fracture resistance, J/m 2 900 779 800 713 681 700 565 600 668 606 554 505 500 400 300 Specimen 1 Specimen 2 Specimen 3 Specimen 1 Specimen 2 Specimen 3 200 Specimen 1 Specimen 2 Specimen 3 Specimen 1 Specimen 2 Specimen 3 100 0 0 3 6 9 12 15 0 3 6 9 12 15 0 3 6 9 12 15 0 3 6 9 12 15 Crack propagation length, mm (b) Mode II Fig. 8 – Fracture resistance curves for composites made with CNTs added in the matrix of the middle six plies and in the baseline composite: (a) mode I and (b) mode II. See Table 4 for the specimen formulations. See Figs. 4 and 5 for the mode I and mode II test configurations, respectively. electron microscope. A gold layer of approximately 50 Å was coated over the fracture surface of the SEM specimen to prevent charging. 3. Results and discussion 3.1. CNT functionalization characterization Weight loss curves obtained from TGA analysis of VGCNTs, short OH-CNTs, short GPS-CNTs, and long GPS-CNTs are shown in Fig. 6a. For the pristine VGCNT, only one main peak at around 634 C appears in the differential weight loss curve (Fig. 6b). With oxidization, the main peak in the differential weight loss curve was shifted downward to around 558 C, which corresponds well with reported findings that defects and functional groups on the CNT wall decrease the thermal stability of CNTs [34]. For the differential weight loss curves of GPS functionalized CNTs, besides the main peak within 660–700 C, another peak appears within 320–360 C and the residual weights are 6% and 8% of the original weights for the 1324 CARBON 5 0 ( 2 0 1 2 ) 1 3 1 6 –1 3 3 1 Table 6 – Mode I and mode II onset and propagation fracture toughness for the baseline composite and composites made with CNTs added in the matrix of the middle six plies. Reported toughness values are averages and values in parentheses are the coefficients of variation expressed in percent. Matrix material Mode I toughness Onset J/m Baseline epoxy (no nanofiller) 1 wt.% short COOH-CNT epoxy 0.5 wt.% short GPS-CNT epoxy 0.25 wt.% long GPS-CNT epoxy 127 133 132 136 2 (4.2) (11) (3.6) (12) Compared to baseline – +5% +4% +7% J/m 616 643 603 670 short GPS-CNT and long GPS-CNT, respectively. The introduction of new peaks into the differential weight loss curve and the increase in residual weight indicate the successful attachment of new functional groups to VGCNTs and short OH-CNTs using the GPS treatment method described in Section 2.1.2. The elemental content of the added functional groups was characterized by XPS spectrum analysis. Wide scans were conducted on as-received VGCNT, oxidized VGCNT and GPS functionalized VGCNT samples to detect the elements present (Fig. 7a) [35]. The typical survey spectrum shows distinct carbon 1s and oxygen 1s peaks for all samples, indicating that carbon and oxygen are the major elements. Low intensity silicon peaks appear in the binding energy range of 100–160 eV for the GPS functionalized sample. Element concentrations in functionalized and unfunctionalized VGCNT samples were calculated based on the high resolution scan spectra as shown in Fig. 7b (see Table 5). The presence of the Si 2p peak at 102.7 eV in the spectrum of the GPS-VGCNT sample was attributed to the attachment of glycidyloxypropyl-trimethoxysilane group on the CNT surface. However, the Si 2p peak at 103.3 eV in the as-received VGCNT sample spectrum and at 106.0 eV in the oxidized VGCNT sample spectrum were considered to be caused by the presence of silicon contamination in a different chemical state (e.g., silica). The Si 2p and O 1s spectra from the oxidized VGCNT sample show evidence of differential charging. The peak widths and binding energies are anomalously high. These artifacts do not change the qualitative interpretation of the spectra, which principally reveal that the oxidation treatment has increased the O concentration, without removing the low levels of Si oxide(s) present in the as-received VGCNTs. Assuming the GPS is attached to the CNT surface in the same way as described in [23], the chemical formula for functional groups attached to the CNT can be written as –(SiO4C6H11)n. Based on this assumed chemical formula of the functional group and assuming all the silicon detected in the XPS spectrum is from the GPS functional group, the weight fraction of the functional groups on GPS-CNTs is approximately 23%. Hence, it is expected that when the weight percent filler loading for GPS-CNF/epoxy and UF-CNF/epoxy are the same, the volume occupied by CNFs in the GPS-CNF/epoxy is about 23% less than that in the UF-CNF/epoxy. 3.2. Interlaminar fracture toughness 3.2.1. Effect of adding CNTs in the matrix Mode II toughness Propagation The mode I and mode II interlaminar fracture resistance curves for S2-glass/epoxy composites with 1 wt.% short 2 Onset Compared to baseline (1.9) (1.9) (1.4) (1.1) – +4% 2% +9% J/m 565 505 606 668 2 (1.3) (3.9) (4.4) (5.0) Compared to baseline – 11% +7% +18% Propagation J/m2 681 554 713 779 (3.5) (6.9) (4.4) (4.0) Compared to baseline – 19% +5% +14% COOH-CNTs, 0.5 wt.% short GPS-CNTs, 0.25 wt.% long GPSCNTs, and with no nanofiller are shown in Fig. 8. Typically, the mode I propagation fracture toughness increases greatly within the initial 8 mm of crack propagation and reaches a plateau beyond about 20–25 mm. Therefore, a nonlinear curve in the form of GIR ¼ GIP P1 expðP2 DaÞ was fit to the mode I fracture resistance (GIR) vs. crack growth (Da) experimental data for each type of material, where GIP, P1, and P2 are curve fitting parameters. The value of parameter GIP is considered to be the mode I propagation fracture toughness for discussion purposes. Relative to the mode I case, the increase in mode II propagation fracture toughness with crack length is small and cannot to be fitted with a single type of functional relationship with the increase of crack length. Hence, mode II propagation fracture toughness was defined as the average mode II fracture toughness at a crack propagation length of 6 mm beyond the initial crack front. The mode I and mode II onset and propagation fracture toughness values for composites made with the three types of CNTs and the baseline case are labeled in Fig. 8 and summarized in Table 6. The results show that adding short COOH-CNTs at 1 wt.% in ply matrix had no significant effect on the mode I fracture toughness and decreased the mode II fracture toughness. Adding short GPS-CNTs at 0.5 wt.% into the matrix had no significant effect on either the mode I or mode II fracture toughness, as well. However, adding longer GPS-CNTs at 0.25 wt.% in ply matrix increased the mode I and mode II fracture toughness values by 7–18%. While the increase in mode I toughness from crack onset to propagation for all laminates made with and without CNTs ranges from 356% to 393%, the increase in mode II toughness from onset to propagation is only 10–20%. It is noteworthy that, in a previous related investigation of an S2-glass fiber composite made with the same matrix system as used presently [36], the addition of 0.5 wt.% unfunctionalized CNTs into the matrix material led to a slight reduction in mode I propagation and mode II onset interlaminar fracture toughness. 3.2.2. Effect of adding CNF- or CNT-filled resin interlayer The second method used for evaluating the effects of CNTs or CNFs on fracture toughness employs CNF- or CNT-filled resin at the fracture plane of IFT specimens made without any nanofillers in the prepregged matrix material. In Fig. 9, the mode I and mode II fracture resistance curves obtained for laminates made with an un-functionalized CNF epoxy interlayer and a GPS-functionalized CNF epoxy interlayer at the fracture plane are compared to the fracture resistance curves of the baseline material with no nanofillers whatsoever. A significant improvement in both mode I and mode II onset and propagation toughness (30–57% relative to the baseline) was obtained by adding either the UF-CNF or GPS-CNF resin interlayer. However, GPS functionalization of CNFs showed relatively little improvement in all toughness measures in comparison to material made with the unfunctionalized CNF interlayer. Numerical toughness values of these materials are labeled in Fig. 9 and listed in the first two lines of Table 7. The effects of interlayer nanofiller length on toughness were evaluated using 0.5 wt.% short GPS-CNTs and 0.25 wt.% long GPS-CNTs. The FRP composites in this com- Interlayer: None 1325 5 0 ( 20 1 2 ) 1 3 1 6–13 3 1 CARBON parison had 0.5 wt.% short GPS-CNTs in the middle six plies. Recalling Section 3.2.1, it has been shown that the addition of 0.5 wt.% short GPS-CNTs in the middle six plies without a nanofilled interlayer had no significant effect on toughness. The mode I and mode II fracture resistance curves for the composites made with and without the GPS-CNT interlayers are shown in Fig. 10. Numerical toughness results are shown as well in Fig. 10 and in the last two rows of Table 7. Compared to the composite with no filler, adding a short GPS-CNT interlayer increased the fracture toughness values by 52–86%. An even higher increase was obtained by using long GPS-CNTs in the interlayer (79–109%). It is apparent that 0.5 wt% UF-CNF epoxy 0.5 wt% GPS-CNF epoxy Mode I fracture resistance, J/m 2 1200 1000 968 903 800 616 600 400 Specimen 1 Specimen 2 Specimen 3 Nonlinear fit 127 200 Specimen 1 Specimen 2 Specimen 3 Nonlinear fit 165 Specimen 1 Specimen 2 Nonlinear fit 176 0 0 10 20 30 0 10 20 0 30 10 20 30 Crack propagation length, mm (a) Mode I Interlayer: 0.5 wt% UF-CNF epoxy 0.5 wt% GPS-CNF epoxy None Mode II fracture resistance, J/m 2 1200 963 996 1000 800 600 843 783 681 565 400 Specimen 1 Specimen 2 Specimen 3 200 Specimen 1 Specimen 2 Specimen 3 Specimen 1 Specimen 2 Specimen 3 0 0 3 6 9 12 15 0 3 6 9 12 15 0 3 6 9 12 15 Crack propagation length, mm (b) Mode II Fig. 9 – Fracture resistance curves of composites made with baseline resin in all plies, with and without a UF-CNF or GPS-CNF interlayer: (a) mode I and (b) mode II. See Table 4 for the specimen formulations. See Figs. 4 and 5 for the mode I and mode II test configurations, respectively. 1326 CARBON 5 0 ( 2 0 1 2 ) 1 3 1 6 –1 3 3 1 Table 7 – Mode I and mode II onset and propagation fracture toughness of composites with and without a nanofilled interlayer at the fracture plane. Reported toughness values are averages and values in parentheses are the coefficients of variation expressed in percent. Interlayer material Mode I toughness Onset J/m Baseline composite (no nanofiller) 0.5 wt.% UF-CNF epoxy 0.5 wt.% GPS-CNF epoxy 0.5 wt.% short GPS-CNT epoxya 0.25 wt.% long GPS-CNT epoxy a a 127 165 176 193 228 2 Compared to baseline (4.2) (14) (8.4) (8.6) (8.4) Mode II toughness Propagation – +30% +39% +52% +79% J/m 616 903 968 1025 1201 2 (1.9) (1.5) (4.1) (1.3) (1.5) Onset Compared to baseline J/m – +47% +57% +66% +95% 2 Propagation J/m2 Compared to baseline 565 (1.3) 783 (3.7) 843 (11) 981 (4.4) 1078 (6.1) – +39% +49% +74% +91% 681 996 963 1266 1423 Compared to baseline (3.5) (6.7) (2.3) (1.9) (6.3) – +46% +41% +86% +109% The FRP composite has 0.5 wt.% short GPS-CNTs in the matrix of the middle six plies. Interlayer: 0.25 wt% long GPS-CNF epoxy 0.5 wt% short GPS-CNT epoxy None 1201 Mode I fracture resistance, J/m 2 1200 1025 1000 603 800 600 Specimen 1 Specimen 2 Specimen 3 Nonlinear fit 400 132 200 0 0 10 20 Specimen 1 Specimen 2 Specimen 3 Nonlinear fit 193 0 30 10 20 Specimen 1 Specimen 2 Specimen 3 Nonlinear fit 228 0 30 10 20 30 Crack propagation length, mm (a) Mode I Interlayer: 0.5 wt% short GPS-CNT epoxy None 0.25 wt% long GPS-CNF epoxy 1600 Mode II fracture resistance, J/m 2 1423 1400 1266 1200 1000 1078 981 713 800 600 606 Specimen 1 Specimen 2 Specimen 3 400 200 Specimen 1 Specimen 2 Specimen 3 Specimen 1 Specimen 2 Specimen 3 0 0 3 6 9 12 15 0 3 6 9 12 15 0 3 6 9 12 15 Crack propagation length, mm (b) Mode II Fig. 10 – Fracture resistance curves of composites made with 0.5 wt.% short GPS-CNT modified resin in middle six plies, with and without a nanofilled interlayer at the fracture plane: (a) mode I and (b) mode II. See Table 4 for the specimen formulations. See Figs. 4 and 5 for the Mode I and Mode II test configurations, respectively. CARBON 5 0 ( 20 1 2 ) 1 3 1 6–13 3 1 1327 Fig. 11 – Representative optical microscope images of cross section of composite specimens: (a) baseline composite with no fillers anywhere (black spots with lateral dimension greater than 30 lm seen on the center top and right lower corner of the image are voids); (b) composite with no nanofiller in the ply matrix and an UF-CNF interlayer; (c) composite with short GPSCNTs in the plies and a short GPS-CNT interlayer. (Microscales were used to locate the midplane on the cross section. The white bars on the left of images show the markings of the microscale, which are separated by 100 lm. The dashed line indicates the midplane position.) Fig. 12 – SEM images of mode I fracture surfaces: (a) baseline S2-glass/epoxy composite (no filler); (b) composite with a GPS-CNF epoxy interlayer and no CNTs in the ply matrix; (c) composite with a UF-CNF epoxy interlayer and no CNT in the ply matrix; (d) composite with a long GPS-CNT epoxy interlayer and short GPS-CNT in the ply matrix. Highlighted areas show evidence of matrix toughening. 1328 CARBON 5 0 ( 2 0 1 2 ) 1 3 1 6 –1 3 3 1 Fig. 13 – SEM images of mode I fracture surfaces of S2-glass/epoxy composites made with (a) short COOH-CNTs in ply matrix, (b) short GPS-CNTs in ply matrix and interlayer and (c) short GPS-CNTs in ply matrix and long GPS-CNTs in interlayer. Highlighted areas show evidence of matrix toughening. Fig. 14 – SEM images of fracture surfaces highlighting dense CNFs wrapped around glass fibers: (a) and (b) mode I fracture surface of composite with an UF-CNF epoxy interlayer and no filler in ply matrix; (c) mode II fracture surface of composite with a GPS-CNF interlayer and GPS-CNTs in the ply matrix. Fig. 15 – SEM images of fracture surfaces of composites made with short GPS-CNTs in the ply matrix and long GPS-CNTs in the interlayer: (a) DCB specimen (cavities in circles were possibly created by CNT pull-out) and (b) ENF specimen showing GPSCNTs bridging the local matrix crack. using high aspect ratio, small diameter nanofillers in the interlayer produced the highest fracture toughness of the S2-glass/epoxy composite. Considering all the materials made with CNF and CNT fillers inside the plies or in the interlayer region, the highest improvement in fracture toughness is obtained by adding a long GPS-CNT epoxy interlayer. It should be noted that adding a resin-rich interlayer at a ply interface may increase the interlaminar fracture toughness. The effect of adding a neat resin interlayer on mode I and mode CARBON 5 0 ( 20 1 2 ) 1 3 1 6–13 3 1 1329 [37,38]. To determine if the present method of placing nanofiller in the interlayer region resulted in a resin-rich interface, cross sections of untested, polished composite specimens were inspected in an optical microscope. The sections were cut perpendicular to the fibers, away from the tip of the PTFE film. It was found that the glass fibers were uniformly packed throughout the entire thickness for all composite specimens made with or without a nanofilled interlayer (Fig. 11). Due to the low viscosity of the epoxy resin system used to manufacture the composites, it is plausible that the resin in the interlayer region bled out during the consolidation process. 3.3.2. Fig. 16 – SEM image showing dense UF-CNFs wrapped around a glass fiber, from the mode I fracture surface of a composite made with an UF-CNF epoxy interlayer. II fracture toughness of the same S2-glass/epoxy composite as used presently was investigated by the authors in [29]. The results show that adding a neat resin interlayer to a composite with short COOH-CNTs in the adjacent plies increased the mode I fracture toughness by 8–11% and increased the mode II fracture toughness by 16–18% compared to the COOH-CNT composite without neat resin interlayer. The results from the current investigation show much higher improvement in mode I and mode II fracture toughness by adding a nanofiller modified resin interlayer (52–95% in mode I and 79–109% in mode II fracture toughness). Hence, the significant improvement in fracture toughness is considered to be due to nanofillers remaining in the interlayer region following laminate consolidation rather than a significantly resin-rich layer of material. This topic is discussed further in Section 3.3.1 with reference to photomicrographs of the interlayer region. 3.3. Microscopy of fracture specimens 3.3.1. Region of fracture plane It is known that a discrete layer of polymer resin (typically 20 lm thick) placed in the interlayer region of a fiber composite can improve the interlaminar fracture toughness Nanofiller toughening mechanisms Based on observations of numerous SEM images of mode I and II fracture surfaces, the effects of CNTs or CNFs on the fracture behavior of the multi-scale reinforced composites are twofold. Firstly, nanofillers roughen the matrix fracture surface and increase matrix fracture surface area, as shown by representative SEM images in Figs. 12 and 13. Secondly, CNFs appear to pack tightly around the glass fibers and obstruct pullout of the glass fibers from the matrix during crack propagation. Such behavior was not observed for shorter nanofillers. Representative SEM images of fracture surfaces of composites made with functionalized CNTs are shown in Fig. 13 and SEM images of fracture surface of composites made with CNFs in the interlayer are shown in Fig. 14. Even though the long GPS-CNTs are not long enough to serve as obstacles for glass fiber pull-out, the SEM images show many long GPS-CNTs pulled from the matrix and bridging matrix cracks (see Fig. 15). Based on these and other similar observations in the SEM, more nanofiller particles were found on the fracture surface when they were added in the interlayer compared to the composite made with fillers added only in the ply matrix. It is plausible that, during the material consolidation process, the epoxy in the interlayer was mostly bled out and the fillers were entrapped by glass fibers, leaving concentrated fillers in close proximity to the glass fibers at the fracture plane (see Fig. 16). The CNFs deposited around the glass fibers can serve as effective obstacles for fiber/matrix debonding (Fig. 14) and long CNTs can bridge matrix cracks in the narrow gaps between glass fibers (Fig. 15). Fig. 17 – SEM images a CNT pulled out from matrix bonded to a glass fiber: (a) short COOH-CNT pulled out from matrix with epoxy partially bonded to CNT surface and (b) unfunctionalized CNT pulled out of the matrix in a related investigation (as received CNT outer diameter is 10–20 nm). 1330 CARBON 5 0 ( 2 0 1 2 ) 1 3 1 6 –1 3 3 1 Fig. 18 – SEM images of CNFs pulled out from the matrix in an S2-glass/epoxy composite: (a) and (b) GPS-CNFs with epoxy layers bonded to the CNF surface and (c) UF-CNFs. The as-received unfunctionalized CNF outer diameter is 60–150 nm. 3.3.3. Effect of CNT/CNF functionalization After CNT/CNF functionalization by either simple oxidization or GPS treatment, an epoxy layer bonded to CNTs or CNFs was observed as shown in Figs. 17a and 18a and b. However, this behavior is not typically observed for unfunctionalized fillers (see Figs. 17b and 18c). These observations indicate that better adhesion between the fillers and the epoxy matrix was obtained with functionalization. 4. Conclusions Carbon nanotubes (CNT) and carbon nanofibers (CNFs) with different aspect ratios and surface functional groups were placed by two different methods in S2-glass/epoxy composite laminates made by filament winding and hot pressing. CNTs and CNFs functionalized with silane coupling agent (GPSCNTs and GPS-CNFs) showed improved bonding to the epoxy matrix. When CNTs were dispersed in the matrix of the multi-scale composite, it was observed that adding short COOHCNTs at 1 wt.% and short GPS-CNTs at 0.5 wt.% of the matrix had no significant effect on the mode I and mode II interlaminar fracture toughness was slightly increased by using dispersed long GPS-CNTs at 0.25 wt.%. However, significant improvements in toughness were seen when placing a CNTor CNF-rich interlayer at the fracture plane of the composite. Specifically, adding a CNF interlayer resulted in a 30–57% improvement in mode I fracture toughness and 39–49% improvement in mode II fracture toughness compared to the baseline material without nanofiller. Adding a GPS-CNT resin interlayer at the fracture interface resulted in a 52– 95% improvement in mode I fracture toughness and 74– 109% improvement in mode II fracture toughness compared to the baseline material. The fracture test results indicate that using higher aspect ratio CNTs placed in an interlayer between fiber reinforced plies is the most effective means of improving fracture toughness of the investigated laminated composite material. Microscopic evaluation of the fracture surfaces of composites made with an interlayer showed that dense concentrations of nanofillers were trapped in microsized channels between glass fibers. Based on the increased toughness of the composites with an interlayer, it appears that the placement of dense concentrations of CNTs and CNFs in between the glass fibers is the best method of inhibiting delamination propagation. Acknowledgements This research project was supported by Penn State Vertical Lift Research Center of Excellence, the US Government under Agreement No. W911W6-06-2-0008, the US Army Research Office under grant W911NF-10-1-0267, the Pennsylvania State University Materials Research Institute, Materials Characterization Lab, the National Science Foundation under Cooperative Agreement No. ECS-0335765, and the Department of Engineering Science and Mechanics at Penn State. The U.S. Government is authorized to reproduce and distribute reprints notwithstanding any copyright notation thereon. The views and conclusions contained in this document are those of the authors and should not be interpreted as representing the official policies, either expressed or implied, of the U.S. Government. Dr. Nicole Brown is thanked for providing the TGA equipment and giving valuable advice on TGA characterization. Dr. David Fecko of AGY is thanked for providing the S2-glass fiber. R E F E R E N C E S [1] Compston P, Jar P-YB, Davies P. Matrix effect on the static and dynamic interlaminar fracture toughness of glass-fibre marine composites. Composites Part B 1998;29(4):505–16. [2] Hojo M, Ochiai S, Gustafson C-G, Tanaka K. Effect of matrix resin on delamination fatigue crack growth in CFRP laminates. Eng Fract Mech 1994;49(1):35–47. [3] Sela N, Ishai O. Interlaminar fracture toughness and toughening of laminated composite materials: a review. Composites 1989;20(5):423–35. [4] Brunner AJ, Murphy N, Pinter G. Development of a standardized procedure for the characterization of interlaminar delamination propagation in advanced composites under fatigue mode I loading conditions. Eng Fract Mech 2009;76(18):2678–89. [5] Kawaguchi T, Pearson RA. The effect of particle-matrix adhesion on the mechanical behavior of glass filled epoxies. Part 2. A study on fracture toughness. Polymer 2003;44(15):4239–47. [6] Singh RP, Zhang M, Chan D. Toughening of a brittle thermosetting polymer: effects of reinforcement particle size and volume fraction. J Mater Sci 2002;37:781–8. [7] Yee AF, Pearson RA. Toughening mechanisms in elastomermodified epoxies. J Mater Sci 1986;21(7):2462–74. CARBON 5 0 ( 20 1 2 ) 1 3 1 6–13 3 1 [8] Scott JM, Phillips DC. Carbon fiber composites with rubber toughened matrices. J Mater Sci 1975;10(4):551–62. [9] Johnsen BB, Kinloch AJ, Mohammed RD, Taylor AC, Sprenger S. Toughening mechanisms of nanoparticle-modified epoxy polymers. Polymer 2007;48(2):530–41. [10] Yoo M, Sharma A, Bakis CE. Comparison of interlaminar fracture toughening of filament wound glass/epoxy composites by using MWCNTs or flexible resin. In: SAMPE symposium and exposition, Society for the Advancement of Materials and Process Engineering, Covina, CA. Paper No. B116; 2009 (CD ROM). [11] Spindler-Ranta S, Bakis CE. Carbon nanotube reinforcement of a filament winding resin. In: Proceedings 47th international SAMPE symposium and exhibition, Society for the Advancement of Materials and Process Engineering, Covina, CA, 2002. p. 1775–87. [12] Zhu R, Pan E, Roy AK. Molecular dynamics study of the stressstrain behavior of carbon-nanotube reinforced Epon 862 composites. Mater Sci Eng A 2007;447:51–7. [13] Coleman JN, Khan U, Blau WJ, Gun’ko YK. Small but strong: a review of the mechanical properties of carbon nanotubepolymer composites. Carbon 2006;44(9):1624–52. [14] Frankland SJV, Harik VM, Odegard GM, Brenner DW, Gates TS. The stress–strain behavior of polymer-nanotube composites from molecular dynamics simulation. Compos Sci Technol 2003;63(11):1655–61. [15] Gojny FH, Wichmann MHG, Fiedler B, Schulte K. Influence of different carbon nanotubes on the mechanical properties of epoxy matrix composites – A comparative study. Compos Sci Technol 2005;65:2300–13. [16] Gojny FH, Wichmann MHG, Köpke U, Fiedler B, Schulte K. Carbon nanotube-reinforced epoxy-composites: enhanced stiffness and fracture toughness at low nanotube content. Compos Sci Technol 2004;64(15):2363–71. [17] Arai M, Noro Y, Sugimoto K-i, Endo M. Mode I and mode II interlaminar fracture toughness of CFRP laminates toughened by carbon nanofiber interlayer. Compos Sci Technol 2008;68(2):516–25. [18] Fan Z, Santare MH, Advani SG. Interlaminar shear strength of glass fiber reinforced epoxy composites enhanced with multi-walled carbon nanotubes. Composites Part A 2008;39(3):540–54. [19] Seyhan AT, Tanoglu M, Schulte K. Mode I and mode II fracture toughness of E-glass non-crimp fabric/carbon nanotube (CNT) modified polymer based composites. Eng Fract Mech 2008;75(18):5151–62. [20] Zhou Y, Pervin F, Lewis L, Jeelani S. Fabrication and characterization of carbon/epoxy composites mixed with multi-walled carbon nanotubes. Mater Sci Eng A 2008;475:157–65. [21] Yang K, Gu M, Guo Y, Pan X, Mu G. Effects of carbon nanotube functionalization on the mechanical and thermal properties of epoxy composites. Carbon 2009;47(7):1723–37. [22] Kim JA, Seong DG, Kang TJ, Youn JR. Effects of surface modification on rheological and mechanical properties of CNT/epoxy composites. Carbon 2006;44(10):1898–905. [23] Ma PC, Kim J-K, Tang BZ. Effects of silane functionalization on the properties of carbon nanotube/epoxy nanocomposites. Compos Sci Technol 2007;67(14):2965–72. 1331 [24] Qiu J, Zhang C, Wang B, Liang R. Carbon nanotube integrated multifunctional multiscale composites. Nanotechnology 2007;18(27):275708. [25] Davis DC, Wilkerson JW, Zhu J, Hadjiev VG. A strategy for improving mechanical properties of a fiber reinforced epoxy composite using functionalized carbon nanotubes. Compos Sci Technol 2011;71(8):1089–97. [26] Reed CW. The chemistry and physics of the interface region and functionalization. In: Nilson JK, editor. Dielectric polymer nanocomposites. New York: Springer Science; 2010. p. 111–3. [27] Pluedemann EP. Silane coupling agents. New York: Plenum Press; 1982, p. 96. [28] ASTM Standard D638. Standard test method for tensile properties of plastics. ASTM International, American Society for Testing and Materials: West Conshohocken, PA; 2008. doi:10.1520/D0638-10. [29] Zhu Y, Bakis CE. Effects of functionalized carbon nanotubes on mode I and mode II interlaminar fracture toughness of a hybrid glass fiber/MWCNT/epoxy composite. In: Proceedings of the American Society for Composites—25th Technical Conference, September 2010, Dayton, OH (CD-ROM): DEStech, Lancaster, Pennyslvania, USA. [30] ASTM Standard D3171. Standard test method for constituent content of composite materials. ASTM International, American Society for Testing and Materials: West Conshohocken, PA; 2009. doi:10.1520/D3171-09. [31] ASTM Standard D5528-01. Standard test method for mode I interlaminar fracture toughness of unidirectional fiberreinforced polymer matrix composite. In: Annual Book of ASTM Standards, American Society for Testing and Materials: West Conshohocken, PA; 2002. p. 254–63. [32] JIS K 7086. Testing methods for interlaminar fracture toughness of carbon fibre reinforced plastics. Japanese Standards Association: Tokyo, Japan; 1993. p. 651–55. [33] Davidson BD, Altonen CS, Polaha JJ. Effect of stacking sequence on delamination toughness and delamination growth behavior in composite end-notched flexure specimens. In: Deo RB, Saff CR, editors. Composite materials: testing and design (twelfth volume), ASTM STP 1274. Philadelphia: American Society for Testing and Materials; 1996. p. 393–413. [34] Yu MF, Files BS, Arepalli S, Ruoff RS. Tensile loading of ropes of single wall carbon nanotubes and their mechanical properties. Phys Rev Lett 2000;84(24):5552–5. [35] Moulder JF, Stickle WF, Sobol PE, Bomben KD. Handbook of X-ray photoelectron spectroscopy. Eden Prairie, MN: Perkin-Elmer Corporation Physical Electronics Division; 1992. [36] Zhu Y, Bakis CE. Quasi-static and cyclic interlaminar cracking behavior of glass fiber/MWCNT/epoxy hybrid composites. In: Proceedings 66th Forum. American Helicopter Society: Phoenix, AZ, USA; 2010. [37] Singh S, Partridge IK. Mixed-mode fracture in an interleaved carbon-fibre/epoxy composite. Compos Sci Technol 1995;55(4):319–27. [38] Ozdil F, Carlsson LA. Mode I interlaminar fracture of interleaved graphite/epoxy. J Compos Mater 1992;26(3):432–59.