Impressive Fatigue Life and Fracture Toughness Improvements in

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Impressive Fatigue Life and Fracture Toughness Improvements in Graphene Oxide/Epoxy Composites Daniel R. Bortz,*,† Erika Garcia Heras,‡ and Ignacio Martin-Gullon† † ‡

Department of Chemical Engineering, University of Alicante, Alicante, Spain Grupo Antolín Ingeniería, Crta Irun 244, Burgos, Spain ABSTRACT: Epoxy systems have proven popular having important applications in aerospace and wind energy, but fracture and fatigue resistance of this polymer remain less than desired. Graphene oxide, a form of atomically thin carbon, possessing impressive multifunctional properties and an ideal interface for interacting with polymer matrices, has emerged as a viable reinforcement candidate. In this work, we report enhancements of 28
111% in mode I fracture toughness and up to 1580% in uniaxial tensile fatigue life through the addition of small amounts (e1 wt %) of graphene oxide to an epoxy system. Graphene oxide was uniquely synthesized by unraveling and splaying open helical-ribbon carbon nanofibers. The resulting oxygenated basal planes and edges of the graphene oxide sheets were observed to promote onset of the cross-linking reaction and led to an increase in total heat of reaction effecting slightly higher glass transition temperatures of the cured composites. Measured improvements were also detected in quasi-static tensile and flexural stiffness and strength. The addition of only 0.1 wt % graphene oxide yielded a ∼12% increase in tensile modulus. At 1 wt %, flexural stiffness and strength were 12 and 23% greater than the unmodified epoxy. Sheets were observed to be well-dispersed and at various orientations within the matrix, enabling their large, 2D, and zero bulk dimensions to pin incipient matrix cracks, a toughening mechanism not typically detected in nanocomposites.

he 2004 report1 on the isolation of the most basic graphitic building block, atomically thin carbon known as graphene, truly excited science. Since then, enormous electron mobility2 and thermal conductivity3 values of these materials have been reported. Mechanically, single-layer graphene, the thinnest known material, has also been shown to possess the highest measured modulus and break strength of any substance to date.4 The range of potential applications for this material is indeed exciting, and although more discoveries and intrinsic measurements are to be made, translating its properties to the macro-scale has now become a prime scientific focus. The introduction of graphene sheets into polymeric matrices for instance has been proposed as an alternative5 or supplementation6 to more traditional carbon nanotube (CNT) reinforcement. Several studies have recently shown improved quasi-static,7 fatigue,8 and electrical properties5 for graphene-based polymer composites. Numerous reviews on the subject have also become available.9
11 Some debate has arisen over the appropriate graphene structure for use in polymer composites, e.g., graphene oxide (G
O) or a reduced form containing fewer surface and edge oxygen groups. Although promising results have been reported by several groups working with reduced forms of graphene in epoxy matrices, e.g., Rafiee et al.,8,12,13 it appears that from the most recent literature14,15 the oxide form may be more advantageous for composite synthesis and future scale-up operations. The “wrinkled” surface of reduced graphene is often credited as a favorable medium for creating a robust interface able to locally interlock with the matrix. At the same time, the atomically

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r 2011 American Chemical Society

smooth surface has been shown to effect low interfacial strength in strain-dependent Raman spectroscopy measurements.16 Furthermore, restacking of the sheets following chemical or thermal reduction due to their instability in popular solvents used in polymer processing can hinder performance, similar to the consequences of agglomeration in CNT composites. The latter concern however has been partially alleviated through the use of surfactants17 or polymer blending prior to reduction.5 The functional groups (epoxide, hydroxyl, carboxyl, and carbonyl) present on the basal planes and edges of G
O facilitate solubility in water or protic solvents. Moreover, the groups may further improve matrix affinity and allow for additional surface chemistry tailoring if desired. Unfortunately, these compounds also result in G
O being electrically insulating, thus limiting the multifunctionality observed in other graphite-based nanocomposites. Literature with respect to graphene synthesis is predominantly populated with methods relating to the intercalation of graphite by oxidation and subsequent exfoliation of the planar graphite oxide layers. More recently, fabricating graphene nanoribbons by methods pertaining to the “unzipping” of CNTs have emerged as alternatives with hopes for improving quality and yield. Kosynkin et al.18 first highlighted the common oxidative method by applying a Hummers and Offeman19 based treatment to lengthwise unzip and roll open both single and multiwall carbon Received: July 8, 2011 Revised: November 5, 2011 Published: December 02, 2011 238

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Figure 1. Series of high-resolution TEM images documenting (a) the unraveling of the graphene layers of a helical-ribbon carbon nanofiber (reprinted with permission from ref 21; copyright 2007 Elsevier), (b) a few-layer sheet having ∼8 μm lateral dimension, and (c) what appears to be a monolayer graphene oxide sheet with ∼2 μm lateral dimension.

nanotubes (SWCNTs and MWCNTs). Herein, we expand on our teams work previously presented by Varela-Rizo et al.20 in which the technique of Kosynkin and colleagues was modified to successfully unravel and splay open the graphene layers of helicalribbon carbon nanofibers (CNFs).21 The result is mostly monoand few-layer G
O sheets (Figure 1) that are highly soluble in water and polar organic solvents, a key feature for their use in composite systems. Conversion is high but not 100%, meaning some oxidized fiber fragments remain unexfoliated. The unique mechanism responsible for formation of the sheets has yet to be fully explained though is probably similar to the chemical unzipping of CNTs described by Kosynkin et al. The G
O sheets were subsequently introduced in small amounts (e1% by weight) to a commercially available thermosetting epoxy system. We report the effect on the epoxide cure kinetics by monitoring the cross-linking reaction via differential scanning calorimetry (DSC). Glass transition temperature (Tg) as well as thermo- and quasi-static mechanical properties were subsequently measured. We further demonstrate large improvements in fracture toughness, and up to 1580%, over the full range of applied stress amplitudes in uniaxial tensile fatigue life of the graphene oxide nanocomposites.

A typical process involved suspending a commercially available helicalribbon CNF (GANF, Grupo Antolín Ingeniería, Spain) in concentrated sulfuric acid (H2SO4, 95
97%, VRW) followed by treatment with potassium permanganate (KMnO4, VRW) at elevated temperature. After oxidation was complete the reaction was terminated by pouring over an ice/hydrogen peroxide (H2O2, 30%, VRW) mixture. Multiple water washings and filtrations followed by a final ethanol (VRW) washing and filtration concluded the process. The material collected from the final filtering was dried under vacuum. The oxidized solid was subsequently suspended in acetone at a concentration of 2.5 mg/mL and exfoliated by high-energy sonication to yield graphene oxide. Highresolution transmission electron microscopy confirmed unraveling and isolation of the fiber’s graphene planes (Figure 1). Lateral dimensions typically measured 0.5
10 μm. Graphene Oxide Dispersion in the Epoxy Matrix. A bisphenol A/F diglycidyl ether blend (Resoltech 1800/1805, Eguilles, France) was used as the polymer matrix. The resin was added to the graphene oxide/acetone suspension and heated to 60 °C for 12 h to allow for slow solvent evaporation. An additional 12 h of heating under vacuum conditions ensured complete acetone removal. The resulting dispersion was passed through a three-roll calender mill with gap settings of 5 μm and roller speeds of 30, 60, and 180 rpm. Dilution with neat resin produced final graphene nanocomposites at concentrations of 0.1, 0.25, 0.5, and 1 wt %. The diamine hardener, 1,2-diaminocyclohexane, was added to the dispersions at a weight ratio of 100:17 and mixed with a laboratory mixer at 5000 rpm for 5 min. The resulting mixtures were poured into a silicone mold and degassed at
1 bar for 30 min. The individually cast specimens, having dimensions of 250  25  2.5 mm

’ EXPERIMENTAL DETAILS Graphene Oxide Preparation. A detailed description and characterization of graphene oxide synthesis was previously reported.20 239

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Table 1. Cure Reaction and Thermomechanical Properties of the Various Neat and Reinforced Epoxy Systemsa Ton (°C)

Tp (°C)

H (J/g)

E00 p (°C)

Tg (°C)

neat matrix

60.50 (0.68)

81.99 (0.47)

209.10 (6.65)

100.73 (1.11)

126.30 (0.10)

0.1 wt % graphene oxide

58.96 (0.50)

81.92 (0.86)

207.15 (3.32)

106.18 (2.58)

126.46 (0.47)

0.25 wt % graphene oxide

58.37 (2.51)

82.78 (0.58)

232.15 (11.81)

108.39 (1.85)

126.89 (0.15)

material

0.5 wt % graphene oxide

57.43 (2.19)

82.78 (0.01)

229.25 (13.79)

117.64 (0.79)

129.33 (0.09)

1 wt % graphene oxide

54.80 (0.10)

83.53 (0.01)

228.35 (5.30)

114.90 (1.52)

128.05 (0.88)

a Onset temperature of curing (Ton), peak exothermic temperature (Tp), and total heat of cross-link reaction (H) as measured by DSC for neat and nanocomposite materials. Loss modulus peak (E00 p) and glass transition temperature (Tg) as measured by DMA for neat and nanocomposite materials. Standard deviations are in parentheses.

for static tensile and tension
tension fatigue testing, 63.5  12.7  4 mm for mode I fracture toughness testing, and 60  12.7  3.2 mm for flexural and dynamic mechanical analysis (DMA), were released from the mold after 12 h at 40 °C. Recommended by the manufacturer, initial curing for 15 h at 60 °C and a postcure of 6 h at 110 °C were carried out in a programmable laboratory furnace. Finally, in order to remove stress concentrating surface defects, specimen surfaces were smoothed with 600-grit silicone carbide paper prior to characterization. Characterization. X-ray photoelectron spectroscopy (XPS) experiments were carried out on a VG-Microtech Multilab 3000 system. The X-ray microprobe pass energy was fixed at 26.0 eV. DSC scans (TA Instruments Q100) to study the effect of graphene oxide on the cure kinetics of the cross-link reaction were performed by preparing a small dispersion and placing a 4
5 mg sample in an open aluminum pan. Scans were in flowing air (50 mL/min) from 25 to 160 °C at a heating rate of 2 °C/min. Samples were quickly cooled, and secondary scans were then executed to ensure a complete cure was achieved. Glass transition temperatures were taken from the tan δ peak on fully cured samples by single cantilever mode DMA (TA Instruments 2980) from 30 to 180 °C at 2 °C/min and 1 Hz. Quasi-static tension and tensile fatigue testing were performed on a servo-hydraulically controlled Instron 8516. Displacement rates during static testing of the nanocomposites were the ASTM D 3039 recommended 2 mm/min. Modulus was measured with a clip on type extensometer over a gauge length of 25 mm between 0.001 and 0.003 absolute strains. Fatigue testing was performed at R = 0.1 (R is the ratio of minimum to maximum stress in a fatigue cycle) and 5 Hz. Two specimens were tested at each of five stress levels. Flexural properties were measure according to ASTM D 790 at a span-to-thickness ratio of 16 and at a · crosshead rate of 1.4 mm/min (ε = 0.01 min
1). Fracture toughness of the nanocomposites was measured following ASTM D 5045 using single edge notched bend (SENB) specimens. A 1 mm wide sharpened notch was machined at the midpoint of each specimen. A precrack was then initiated by tapping a fresh razor blade in the notch. Specimens were loaded to failure at 10 mm/min using an Instron 3344 equipped with a 2 kN load cell. Load
deformation values were recorded, and precrack lengths were measured post mortem by optical microscopy. The critical-stress-intensity parameter (KIc) was calculated according to the standard; critical strain energy release rate (GIc) was tabulated by integration of each load
deformation curve. Sheet resistance was measured on prepared films by painting two rectangular electrodes (25  2 mm) parallel to one another and separated by 25 mm. The resistance between the two electrodes was then quantified by using a two-point measurement device operating at 10 V, fulfilling the requirements of ASTM D257. Scanning electron microscopy (SEM) of gold-coated fracture surfaces was performed using a JEOL JSM-840 SEM at 15 kV. Transmission electron microscopy (TEM) of ultramicrotomed G
O nanocomposites (∼80 nm thickness) positioned on a standard copper grid was performed at 200 kV using a JEOL 1020 TEM.

’ RESULTS AND DISCUSSION Cross-Link Reaction Monitoring, XPS, and Thermomechanical Properties. After the intercalation reaction and exfoliation

in acetone, graphene oxide/epoxide resin dispersions were prepared by the controlled and thorough evaporation of acetone followed by three-roll calender milling. Cross-linking was achieved through the stoichiometric addition of a diamine hardener (1,2diaminocyclohexane). Results from DSC scans on unreacted dispersions and DMA scans on fully cured specimens are given in Table 1. Onset of the cure reaction (Ton) was observed to decrease as a function of G
O concentration, up to a maximum of about 6 °C at 1 wt %. Peak heat flow (Tp), i.e., the temperature of maximum exothermic heat flow during the cure reaction, remained largely unaffected. Integrating the area of each exothermic peak also revealed an increase in the total heat of reaction, H, in the nanocomposites. This effect though appeared saturated at 0.25 wt % G
O. The acceleration effect in the early stage of the cure reaction could be explained by an increase in the oxygencontaining functionalities on the surfaces and edges of the sheets.18 Such an increase in oxygen groups could produce a catalytic effect on epoxide ring opening, resulting in higher initial reaction rates. Comparison of the XPS spectra of the parent CNFs and G
O corroborates the existence of a greater quantity of surface oxygen functionalities in the graphene oxide materials (Figure 2). The O 1s/C 1s peak area ratio increased from 0.05 for the CNFs to 0.68 in the case of G
O. This led to an increase in the relative mass of oxygen from 5% to 40%. Subsequent DSC scans did not display any further exothermic behavior, confirming a complete cure had been achieved. The second scans however did not allow for an accurate measure of Tg. Thus, glass transition temperatures were taken from the tan δ peak of DMA scans and exhibited increases of 2
3 °C at 0.5 and 1 wt % G
O loadings as compared to the neat sample. A more substantial shift to higher temperatures, up to ∼17 °C, was detected in the peak of the loss modulus curve. This behavior suggests a change in molecular mobility relative to the neat epoxy matrix.22 Horizontal shifting by the loss modulus curve indicates the entire polymer matrix was effectively altered by interactions with the graphene oxide surface. If a region of nonbulk polymer near the surface of the filler (interphase) had been created, the loss modulus curve would have instead exhibited broadening toward higher temperatures; this was not observed. Combined, these results indicate G
O inclusion could impart slightly higher cross-link densities to the network architecture, though the modest Tg increases suggest only minor overall affect. The 2D, zero bulk structure of the G
O planes did not however appear to stimulate the formation of an interphase zone as is common in materials having lower cross-link densities.23 240

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strength decreased. At 1 wt % G
O, tensile modulus was nearly equal to the neat matrix and strength was slightly less than the control. Flexural modulus and strength similarly saw the largest incremental increase at the lowest G
O concentration, 9 and 18%, respectively, at 0.1 wt % G
O. Unlike the tensile results though, gains continued as a function of graphene oxide content across the full range of concentrations. At 1 wt % G
O, flexural modulus and strength were 12 and 23% greater than the control group. More significant improvements were observed in fracture toughness. Graphene oxide addition led to enhancements of the critical stress intensity factor (KIc) of approximately 28
63% with respect to the control. Similarly, improvements of 29
111% were detected in the critical strain energy release rate (GIc). The rate of improvement was quite impressive up to 0.5 wt % G
O, whereas doubling the graphene oxide concentration to 1 wt % G
O revealed a saturation of the toughening effect, only increasing KIc by a further ∼1% over the 0.5 wt % G
O loaded composite. The characteristic brittleness and low fracture toughness of cured epoxy systems are a product of their high cross-link densities, which result in poor absorption of energy during fracture. These factors frequently lead to mirror-like fracture surfaces, e.g., Figure 4a. Evidence of toughening caused by fiber pullout, rupture, crack bridging, and the nucleation of voids surrounding debonded fibers are often visible on CNT reinforced epoxy composite fracture surfaces.24
26 These mechanisms appeared absent in the current G
O composites. Initial SEM evidence indicated a departure from the smooth fracture surface by the composite materials (Figure 4b). The coarse, multiplane features on the composite fracture surface suggest that the graphene oxide sheets have induced the deflection of propagating crack fronts. This process introduces off-plane loading that generates new fracture surfaces, thus

Nanocomposite Mechanical and Electrical Properties. Monotonic material performance of the graphene oxide based nanocomposites is summarized in Figure 3. Tensile modulus was enhanced by 12% at 0.1 wt % G
O. Higher loadings also displayed increased tensile stiffness when compared to the control but more modest than the 0.1 wt % G
O group. Ultimate tensile strength showed a maximum improvement of about 13% in samples with 0.5 wt % G
O compared to the control group. Above this concentration both tensile stiffness and

Figure 2. High-resolution XPS C 1s spectra and fitting curves: (a) parent carbon nanofibers and (b) graphene oxide and O 1s spectra and fitting curves; (c) parent carbon nanofibers and (d) graphene oxide.

Figure 3. Quasi-static mechanical and electrical properties of both neat epoxy and epoxy composites containing various weight fractions of graphene oxide: (a) Tensile modulus and ultimate tensile strength, (b) flexural modulus and ultimate flexural strength, (c) critical stress intensity factor and critical strain energy release rate, and (d) surface resistivity. 241

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Figure 4. Fracture surfaces indicating the machined notch, precrack, and propagation zone for the neat polymer system (a) and 1 wt % graphene oxide composite (b); crack propagation is from bottom to top in both micrographs. Note the rougher propagation zone of the reinforced specimen. A highresolution view (c) of the river marks leading back to the crack initiation site in a neat specimen. Evidence of crack pinning (d) was observed by the “bowlines” (white circles and arrow) left behind when the crack front bowed out between graphene oxide sheets. Crack propagation is from right to left.

Table 2. Theoretical Values of CTOD (δ) Assuming a Measured Matrix Yield Stress (σy) of 65.1 MPa

increasing the required strain energy for the continuation of fracture.27 Under higher magnification, evidence of crack pinning was detected (Figure 4b). Crack pinning occurs when a propagating crack encounters a series of impenetrable obstacles, the crack front bows out between particles but remains pinned at the particles. The process is identifiable by the bow-shaped lines left behind on fracture surfaces. Typically, nanometer-sized particles and fibers cannot account for the pinning of cracks because their relative size is much smaller than the crack tip opening displacement (CTOD). The micrometer 2D dimensions of the current G
O sheets though are understood to be large enough to explain the observation of pinning. The CTOD, δ, can be estimated according to28 δ¼

KIc 2 G ð1
v2 Þ ¼ Ic Eσy σy

GIc (J/m2)

δ (μm)

neat matrix

257

3.9

0.1 wt % graphene oxide

332

5.1

0.25 wt % graphene oxide

439

6.7

0.5 wt % graphene oxide 1 wt % graphene oxide

528 542

8.1 8.3

material

High conductivities and low percolation thresholds have been reported by chemically reduced graphene-based composite materials.5 The oxygen functionalities blanketing the surfaces and edges maintain G
O as chiefly insulative, thus limiting conductivity even after network formation is achieved. Nonetheless, surface resistivity was shown to decrease by 4 orders of magnitude by G
O addition. Even moderate reductions in sheet resistance can be significant for applications such as electrostatic painting. Results from tension dominated (R = 0.1) uniaxial fatigue tests conducted at a single graphene oxide concentration of 0.5 wt %

ð1Þ

where E, σy, and v2 are the Young’s modulus, yield stress, and Poisson’s ratio of the unmodified matrix. Table 2 gives calculated CTODs on the order of 5
8 μm for the various loadings, similar to the 2D length scale of the G
O sheets observed in Figure 1, corroborating the SEM evidence. 242

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Figure 5. Fatigue performance of both neat epoxy and epoxy composites containing 0.5 wt % graphene oxide. The solid lines represent leastsquares fits taking into account the abscissa contains the dependent variable. Neat matrix data reprinted with permission from Elsevier.30

are given in Figure 5. Significant improvements in mean fatigue life (taken as the linear fit29) were realized over the full range of applied stress amplitudes. Mean life at the highest comparable stress level, 40 MPa, was 420% greater than that of the control. Enhancements at lower stresses were more impressive, 1580% at 25 MPa. Prior results30 along with other studies31 have demonstrated improvements in the fatigue life of epoxy systems via CNT/CNF reinforcement. Typically, these gains are observed at low applied stress/strain regimes, where imbedded filaments can more readily interact with the surrounding matrix. Improvements are seldom reported at loads approaching 50% of the ultimate tensile strength. Here, fatigue life is often on par with or even inferior to the virgin material due to what can be described as competition between well-dispersed and nested regions of filaments. That is, high stress intensities due to loading in conjunction with the stress concentrating effects of the agglomerated areas ultimately overcome the beneficial effects of isolated filaments. The data shown in Figure 5 do not follow this trend. Instead, a shift to longer lives and slight “flattening” of the S
N curve (smaller slope) is observed. Complementary results were recently reported by Rafiee et al.8 and Yavari et al.32 In the former communication, the authors reported a significant decrease in the fatigue crack propagation rate of a graphene/epoxy nanocomposite through the addition of 0.125 wt % of what was referred to as partially oxygenated graphene sheets. Although they did not report on fatigue life, the authors described a 25-fold reduction in crack growth rate at ΔK = 0.5 MPa 3 m1/2. In the other letter, a graphene/epoxy composite was used to fabricate a glass fiber laminate to form a three-phase hierarchical composite. These structures were then subjected to fatigue loading in both bending and tensile configurations. Increases in fatigue life were more substantial in bending, up to 2 orders of magnitude greater than the traditional glass fiber composite. In tension, fatigue life was improved by 200
400% over the neat matrix composite. In both reports the authors claimed that the wrinkled texture of the thermally reduced graphene sheets were able to interlock with the epoxy matrix and led to improved interfacial adhesion. Accordingly, the mechanism(s) responsible for the added longevity in the current graphene oxide composites were further examined. The presence of well-dispersed G
O sheets at various

Figure 6. High-resolution TEM observation of two graphene oxide sheets embedded within the epoxy matrix. A large sheet (a) with ∼11 μm lateral distance with basal planes orientated perpendicular to the viewing angle and a sheet (b) with basal planes oriented parallel to the viewing angle containing ∼7 graphene layers tapering to a single dangling sheet (white arrow).

orientations within the matrix (Figure 6) suggests that some of these sheets would have their basal planes oriented to intercept incipient cracking. The large surface area of the sheets, combined with their anticipated strong interfacial bonding (due to their extremely high specific surface area) and enormous tensile properties, could then act to arrest crack growth. Such an explanation could further account for the earlier observation of crack pinning on fractured surfaces. It is unclear if a subsequent frictional pullout or bridging mechanism, often observed in fibrous nanocomposites, could further explain the improvements in fatigue and toughness. To this regard, Raman spectroscopy based interfacial stress transfer studies of graphene nanocomposites have generated some confusion indicating both low interfacial shear stresses16 and large debonding strains.33 Surface chemistry differences in such atomically thin materials may lead to significant inconsistencies in these measurements, particularly since the likely presence of surface groups in the oxide form can lead to a degree of covalent interaction at the graphene/matrix interface. 243

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G
O be described? Forthcoming research will answer these and other fundamental questions and provide more accurate descriptions of the atomic level interactional behavior of these materials.

Figure 7. Comparison between measured and modeled (Halpin
Tsai) normalized stiffness’s of graphene oxide nanocomposities. Adequate agreement is observed up to 0.25 wt % before divergence from modeled behavior.

Therefore, to gain further insight into the effectiveness of G
O/matrix stress transfer, a modified version13 of the classic Halpin
Tsai model34 was evaluated to compare predicted and measured composite stiffness values. The modified method takes the G
O sheets as effective rectangular solid fibers with appropriate length (l), width (w), and thickness (t) values and estimates the ratio of elastic modulus of the nanocomposite (Ec) to the neat matrix (Em) as   ðEG
O =Em Þ
1 1 þ ððw þ lÞ=tÞ VG
O Ec 3 ðEG
O =Em Þ þ ððw þ lÞ=tÞ   ¼ ðEG
O =Em Þ
1 Em 8 1
VG
O ðEG
O =Em Þ þ ððw þ lÞ=tÞ   ðEG
O =Em Þ
1 1 þ 2 VG
O 5 =E Þ þ 2 ðE  G
O m  þ ð2Þ ðEG
O =Em Þ
1 8 1
VG
O ðEG
O =Em Þ þ 2

where the G
O elastic modulus (EG
O) is assumed 1 TPa,4 Em is 2.99 GPa, and G
O weight fractions were converted to volume fractions based on the standard density of graphite (F = 2.25 g/cm3) and the manufacturer's listed matrix density (F = 1.12 g/cm3). Since the sheets have no genuine length or width and were of various graphene layers, approximations based on observed dimensions from our TEM characterization had to be introduced. Thus, average l, w, and t were assumed equal to 5 μm, 5 μm, and 4 nm, respectively. Predicted and measured normalized composite tensile moduli are compared in Figure 7. Reasonable correlation between the values is observed up to 0.25 wt % G
O. Higher loadings however trend significantly below the model. This disagreement could be explained by several circumstances. First, as is the case with fibrous nanocomposites, nonuniform dispersion could account for some level of discrepancy between predicted and measured performance. Since G
O was fabricated in a top-down process, unconverted CNF fragments could also contribute to the formation of agglomerated regions within the matrix. Moreover, the unconverted fragments indicate that G
O do not constitute the total reinforcement weight fraction, thereby introducing an additional gap between the model and reality. Finally, more basic questions regarding the adequacy of the theoretical prediction can also be asked. For example, how does the matrix interact with atomic-sized thicknesses, and how should stress-based interlayer interactions in few-layer

’ CONCLUSIONS In this paper, we have demonstrated significant toughness and fatigue life improvements through the addition of uniquely synthesized and scalable graphene oxide sheets to a thermosetting epoxy system. The fatigue life improvements were observed over the full range of experimental stress levels offering potential performance improvements to numerous applications. We have further shown the effects on the cross-linking reaction and monotonic tensile and flexural properties disseminated by the sheets. The oxygen-containing functionalities present on the basal plane surfaces and edges of the graphene oxide sheets triggered onset of the cure reaction at lower temperatures and resulted in slightly larger exotherms in calorimetric scans. Consequently, moderately higher glass transition temperatures and shifting of the loss modulus were measured in the nanocomposites. Decreases in sheet resistance of 4 orders of magnitude occurred. A larger decrease was not observed due to the insulating nature of graphene oxide. Improvements in stiffness and strength were modest but significant, particularly the increases at low graphene oxide weight fractions. The notable toughening and longevity enhancements were products of the deflection and pinning of small-scale propagating cracks. The latter is not typically observed in nanofilled polymer composites; however, the 2D dimensions of the graphene oxide sheets were similar to the calculated crack tip opening displacements based on measured fracture energies. Thus, the occurrence and consequent microscopic evidence of pinning are considered justified. ’ AUTHOR INFORMATION Corresponding Author

*E-mail: [email protected].

’ ACKNOWLEDGMENT This work was supported in part by the Investigacion en Nuevos Materiales para su Aplicacion en la Industria Aeronautica (NACAR) project through Grupo Antolín Ingeniería, Burgos, Spain. ’ REFERENCES (1) Novoselov, K.; Geim, A.; Morozov, S.; Jiang, D.; Zhang, Y.; Dubonos, S.; Grigorieva, I.; Firsov, A. Electric field effect in atomically thin carbon films. Science 2004, 306 (5696), 666–669. (2) Heersche, H. B.; Jarillo-Herrero, P.; Oostinga, J. B.; Vandersypen, L. M. K.; Morpurgo, A. F. Bipolar supercurrent in graphene. Nature 2007, 446 (7131), 56–59. (3) Balandin, A. A.; Ghosh, S.; Bao, W.; Calizo, I.; Teweldebrhan, D.; Miao, F.; Lau, C. N. Superior thermal conductivity of single-layer graphene. Nano Lett. 2008, 8 (3), 902–907. (4) Lee, C.; Wei, X.; Kysar, J. W.; Hone, J. Measurement of the elastic properties and intrinsic strength of monolayer graphene. Science 2008, 321 (5887), 385–388. (5) Stankovich, S.; Dikin, D. A.; Dommett, G. H. B.; Kohlhaas, K. M.; Zimney, E. J.; Stach, E. A.; Piner, R. D.; Nguyen, S. B. T.; Ruoff, R. S. Graphene-based composite materials. Nature 2006, 442 (7100), 282–286. 244

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dx.doi.org/10.1021/ma201563k |Macromolecules 2012, 45, 238–245