Extreme mechanical reinforcement in graphene oxide based thin-film nanocomposites via covalently tailored nanofiller matrix compatibilization

Accepted Manuscript Extreme mechanical reinforcement in graphene oxide based thin-film nanocomposites via covalently tailored nanofiller matrix compatibilization Toby Sainsbury, Sam Gnaniah, Steve J. Spencer, Sandro Mignuzzi, Natalie A. Belsey, Keith R. Paton, Amro Satti PII:

S0008-6223(16)31035-1

DOI:

10.1016/j.carbon.2016.11.061

Reference:

CARBON 11501

To appear in:

Carbon

Received Date: 1 September 2016 Revised Date:

21 November 2016

Accepted Date: 22 November 2016

Please cite this article as: T. Sainsbury, S. Gnaniah, S.J. Spencer, S. Mignuzzi, N.A. Belsey, K.R. Paton, A. Satti, Extreme mechanical reinforcement in graphene oxide based thin-film nanocomposites via covalently tailored nanofiller matrix compatibilization, Carbon (2016), doi: 10.1016/j.carbon.2016.11.061. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

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Table of Contents Graphic

Thin film polymer-functionalized graphene oxide nanocomposites demonstrating extreme mechanical

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enhancement due to the chemically compatibilized nanofiller-polymer matrix interaction.

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ACCEPTED MANUSCRIPT Extreme Mechanical Reinforcement in Graphene Oxide Based Thin-Film Nanocomposites via Covalently Tailored Nanofiller Matrix Compatibilization Toby Sainsbury,a,* Sam Gnaniah,a Steve J. Spencer,a Sandro Mignuzzi,a Natalie A. Belsey,a Keith R. Paton,a Amro Sattib National Physical Laboratory, Teddington, London, TW11 0LW, United Kingdom

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a

b

Leitat Technological Center, Carrer de la Innovació, 2, 08225 Terrassa, Barcelona, Spain

* Corresponding author

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E-mail address: [email protected] Abstract Oxide

Poly(Bisphenol

(GO) nanosheets

have

been covalently

functionalized with

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Graphene

A-co-epichlorohydrin)(PBE)

a

polymer

structurally

analogous

to

polycarbonate and derived from Bisphenol-A. PBE-functionalized-GO (PBE-GO) was utilized to form PBE-GO:PBE thin-film nanocomposites by dispersion of the PBE-GO within a solution of the PBE polymer followed by solvent evaporation to form thin-films. Extremely high levels of reinforcement were observed by mechanical analysis of the PBE-

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GO:PBE nanocomposites films relative to the pure PBE polymer and pristine GO:PBE control specimens. The impact of the covalent functionalization resulting in the direct chemical compatibilization between the nanofiller and the polymer matrix was characterised using tensile mechanical analysis and was evaluated by analysis of the

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mechanical enhancements observed. These include maximum enhancements for the PBEGO:PBE nanocomposites above the PBE control in Young’s modulus (Y) of 26 % increase,

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>63 % increase in ultimate tensile strength (UTS), as well as more than ∼20 fold increases in the strain-to-failure (εB), and >30 fold enhancement of the toughness (T). The implications

of

tailored

nanofiller-matrix

compatibilization

through

covalent

functionalization are considered in the context of the formation of ultra-strong thin films and the rational design of nanocomposites over a range of alternative polymer systems. Keywords: graphene oxide, functionalization, nanocomposite, Poly(Bisphenol A-co-epichlorohydrin) Introduction In recent years graphene and its chemically derived form; graphene oxide, have seen unprecedented levels of research attention on account of an expansive array of technological applications.1,2

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ACCEPTED MANUSCRIPT Motivated by the array of impressive intrinsic physical properties associated with graphene and graphene oxide (GO) such as mechanical, barrier, electronic and thermal, one of the critical focus areas for 2-Dimensional (2-D) carbon nanomaterials research has been the integration of these materials within solvents, polymer matrices and on substrates.3,4 Graphene based nanocomposites in which the interface between filler and matrix have been tailored have been shown to exhibit enhanced performance in applications including flexible packaging, engineering nanocomposites,

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printed electronics, memory devices, and super capacitors.5-9 It is clear that without efficient integration at the interface between nano and bulk materials, mechanical stress transfer, phonon conduction and electron transport will be diminished, and thus the effective utilisation of the nanomaterial compromised.2-4 In response to this, chemical functionalization and the surface

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modification of graphene and GO has received considerable attention.2-4,10,11

Specifically,

the

covalent functionalization of the basal plane in graphene as well as the derivatization of GO

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functional groups has facilitated precise tailoring of the surface chemistry using a diverse array of molecular, bio-molecular and condensed phase materials.10-13 This has been achieved via a number of chemical strategies which include primarily the direct chemical functionalization of graphene using diazonium and related radical intermediates,3,10,14 as well as the derivatization of graphene oxide functional groups via amide, ester and carbamate formation.11,13 Most notably, for many of the envisaged applications of graphene and graphene oxide as a physical barrier,15,16 electrical17 and

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thermal conductor18, or as a mechanical reinforcing agent within polymer matrices,19 the ability to disperse and efficiently control the interface between the materials has critical implications for the performance of the nanocomposite systems produced.20,21 This has prompted a large amount of research interest in creating chemical bonds between polymers and graphene/GO substrates by the

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covalent attachment of polymerization initiators, and by covalently grafting polymers to graphene/GO substrates.10-13,22-24

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Guided primarily by a choice of industrially relevant polymers, investigation of the ‘grafting from’ technique are currently of extremely high interest, whereby polymer chains are covalently bound to GO substrates and the modified material dispersed within a parent polymer matrix.22-25 Bisphenol-A based polymers which are conventionally used for bulk mechanical applications represent one of the highest volume chemical classes produced globally.26 Including polycarbonate and epoxy resins, bisphenol-A based systems have attracted enormous research directed towards the enhancement of mechanical properties such as modulus and fracture toughness due to the commercial market size and potential economic benefits of bisphenol-A based nanocomposites.27,28 Despite several notable examples of mechanical property enhancement of graphene/GO-epoxy and graphene-polycarbonate nanocomposites,29-32 widespread examples of mechanical reinforcement where stiffness, strength 2

ACCEPTED MANUSCRIPT and toughness parameters are simultaneously enhanced are relatively rare.11,23,27,28 A large number of reports exist in the literature concerning the formation of polymer nanocomposites using graphene and GO which describe increases in stiffness (Young’s modulus, Y) and strength (ultimate tensile strength (UTS)) yet decreasing values in strain to failure (εB) and toughness (T).33-35 Realistic applications of nanocomposites within commercial markets will require the strain-to-failure of the

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material and toughness to be enhanced over and above the values of the base polymer matrix in order to add significant value.26,27,36 Chemical strategies to address the issues of compatibilization and integration of graphene and GO within polymer systems are therefore of high importance.10,12 Although exfoliated graphene represents a more structurally perfect material, GO, which exists typically as monolayer sheets and contains intrinsic chemical functionality, offers a great deal of

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potential in the context of readily tailoring the surface chemistry to that of the matrix and being able to utilize intrinsic mechanical properties which persist despite the highly oxidative preparative

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treatment.11,13

In view of the above factors, we have developed a GO-based nanocomposite system with an analogue of polycarbonate, poly(Bisphenol A-co-epichlorohydrin)(PBE) in this work. The covalent functionalization of GO with PBE polymer chains was investigated in order to compatibilize the nanofiller material within a PBE matrix by creating identical chemical interfaces between the

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materials. This was achieved by covalently coupling PBE polymer chains to carboxylic acid groups on GO via ester formation. Through the designed chemical compatibilization between nanofiller and matrix the impact of chemically tailored GO surfaces on the mechanical performance of the nanofiller with PBE matrix relative to unmodified GO was investigated. The wider implication of this

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also considered.

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methodology for the controlled interaction between 2-D nanomaterials and polymer matrices was

Experimental

Materials: Chemicals and reagents were supplied by Sigma Aldrich and were used as received. Processing: Sonication was carried out using a Branson 2510EMT sonic bath. Centrifugation was performed using an Eppendorf, Model 5430 R centrifuge. Characterisation: Transmission Electron Microscopy (HR-TEM) images were acquired using a JEOL2010 operated at 200 kV. Samples for TEM were prepared by evaporating a drop of a dilute suspension of graphene onto a lacy-carbon copper TEM grid. Scanning Electron Microscopy (SEM)

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ACCEPTED MANUSCRIPT images were acquired using a ZEISS Supra at an accelerating voltage of 5 kV and a nominal working distance of 2.5 mm. Atomic Force Microscopy (AFM) measurements were made using an Asylum Research Cypher system in tapping mode using PPP-NCHR silicon cantilevers at a resonant frequency of ∼290 kHz. Fourier Transform Infrared spectroscopy (FTIR) measurements were recorded using a Nicolet 6700, with a diamond ATR accessory. Raman spectra were recorded using a Renishaw inVia

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Raman Microscope at λ = 514.5 nm and a total laser power incident on the sample of 2 mW laser excitation. X-ray Photoelectron Spectroscopy (XPS) was performed using a Kratos Axis Ultra DLD system using an Al monochromated X-ray source operated at 15 kV, 5 mA emission. Analysis conditions used were 160 eV pass energy, 1 eV steps, 0.2 sec dwell per step and 2 sweeps. Samples for XPS were prepared by evaporation of graphene from solution onto Si-wafer substrates. X-ray

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diffraction measurements were conducted using a Siemens D5000 diffractometer in conjunction with a Cu-Kα X-ray tube (40 kV, 40 mA) filtered using a Ni filter and anti-scatter and divergences slits

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of 1mm under standard θ-2θ conditions. Time-of-Flight Secondary Ion Mass Spectrometry (ToFSIMS) analysis was performed in the negative ion mode using a ToF-SIMS IV mass spectrometer (IONTOF) with a 25 keV Bi3+ primary ion source, a target current of 0.125 pA, and with a final ion dose of 8.2 x 1010 ions per cm2. Optical Absorption measurements were recorded using a Perkin Elmer Lambda 850 using a 2 mm path-length quartz cuvette. Thermogravimetric Analysis (TGA) was performed using a Perkin-Elmer Pyris-1 TGA system in air. The temperature was scanned from 30 to

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1100 oC at 10 oC/min. Mechanical characterisation of nanocomposite thin-films was performed using a TA Instruments GSA-G2 DMA system in tensile analysis configuration. Preparation of Graphene Oxide (GO)

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GO was prepared according to the Hummers method.37 Powdered graphite, (<20 μm particle size, synthetic) (0.50 g) and NaNO3 (0.25 g, 2.94 x 10-3 mol) were combined in H2SO4 (95.0-98.0%, 16.5 mL). After 10 min of stirring, KMnO4 (1.5 g, (.5 x 10-3 mol) was slowly added over 30 min to the

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stirring solution which was maintained at 0 oC in an ice bath. The mixture was then allowed to stir for 30 min at 35 oC in an oil bath. Water (distilled, deionized, 18.2 M.Ω.cm)(23 mL) was added slowly to the solution which increased in temperature to 98 oC, and was then followed by the addition of a further aliquot of water (distilled, deionized, 18.2 M.Ω.cm)(70 mL). H2O2 (30 %, 1.75 mL) was then added to the stirring solution. The resulting dark solution was then divided into equal aliquots in plastic centrifuge tubes and was centrifuged (10 min, 3000 rpm) in order to precipitate the oxidised graphite material and to allow the removal of any residual fulvic acid graphitic-type fragments. Graphite oxide (GO) deposited on the bottom of the centrifuge tubes was then carefully resuspended by sonication in order to allow the removal of any fragment materials which may be present at the surface of the GO nanosheets. The GO dispersion was then washed 3 times with HCl 4

ACCEPTED MANUSCRIPT solution (10 %) followed by ethanol. Each washing step was followed by centrifugation and removal of the supernatant fraction using a Pasteur pipette. Residual GO powder was allowed to dry overnight (30 oC, oven) and was recovered as a dark brown solid. XPS analysis of GO prepared by this technique was used to determine a nominal O/C ratio of 0.21, and were found to have an average flake size of ∼ 1 µm determined by SEM and AFM.

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Covalent Functionalization of GO with Poly(Bisphenol A-co-epichlorohydrin)(PBE)

GO was covalently coupled to PBE chains via N,N’-Dicyclohexylcarbodiimide (DCC) mediated esterification of GO-carboxylic acid groups and PBE-hydroxyl groups. GO (5.0 x 10-2 g) was dispersed by ultrasonication in dimethylformamide (DMF)(25 mL) for 15 mins, followed by the addition of DCC

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(1.5 x 10-1 g, 7.3 x 10-4 mol) and 4-(dimethylamino)pyridine (DMAP)(4.7 x 10-3 g, 3.9 x 10-5 mol) under magnetic stirring. Poly(Bisphenol A-co-epichlorohydrin)(PBE)(2.5 x 10-1 g, Mw ∼ 40k) was then gradually added over 60 mins and allowed to dissolve in the stirring suspension. The suspension was

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stirred for 48 h. The reaction mixture was then carefully filtered using membrane filtration apparatus (Millipore, 47 mm) in conjunction with an Anodisc filter membrane (Whatman, 0.02 μm) followed by washing through with DMF (200 mL) and then with 2-propanol (300 mL). Excess solvent was deliberately used at this step to create a large dilution and ensure that as well as molecular byproducts, that any fulvic acid graphitic-type fragments, which may be present, were removed. The

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solid PBE-functionalized GO was then sonicated off the filter membrane (5 mins) into DMF (25 mL) and the suspension was filtered once again using DMF (200 mL) and 2-propanol (300 mL). The washing and filtration was repeated once more in order to ensure all traces of the DCC, DMAP, and any graphitic fragments were removed. GO was then sonicated off the final membrane filter into

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DMF (40 mL) and centrifuged for 25 min at 4000 rpm in order to separate any trace of unreacted PBE material from PBE-functionalized GO. The supernatant portion was discarded and the PBEfunctionalized GO precipitate retained. The precipitate was once again dispersed in DMF (40 mL) by

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sonication and centrifuged once again to remove any traces of unreacted materials or soluble graphitic fragments. PBE-functionalized GO was then isolated from residual DMF by membrane filtration (Anodisc, 0.02 μm) and washed using 2-propanol (200 mL) and methanol (100 mL). PBE-GO solid was dried of any residual solvent using rotary evaporation followed by storage prior to characterization using Schlenk vacuum apparatus. Preparation of PBE-GO Thin-Film Nanocomposites A series of thin film nanocomposite samples were prepared consisting of PBE-functionalized GO within PBE (PBE-GO:PBE) as well as a series of control specimens of unmodified GO in PBE (GO:PBE). Specimens across a range of mass fraction loadings (0.1 wt. %, 0.25 wt. %, 0.5 wt. %, 0.75 wt. %, 1.0

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ACCEPTED MANUSCRIPT wt. %) were prepared by the dispersion of PBE-GO and GO (1 mg, 2.5 mg, 5.0 mg, 7.5 mg, 10 mg) within DMF (50 mL) by mild sonication (15 mins). PBE (1.0 g) was added gradually over 60 mins to the respective sonicating GO suspensions. Following sonication each suspension was allowed to stir for a further 4 h. Aliquots of the nanocomposite dispersions (10 mL) were deposited onto glass microscope slides and were dried under vacuum at 35 °C for 48h. The nanocomposite thin-films

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were removed from the glass substrates and were cut into specimens for mechanical testing using a mechanical punch apparatus. Prior to mechanical analysis, mechanical specimens were stored in vacuum at ambient temperature using a vacuum oven. Thin-film nanocomposites were measured to have an average thickness of 0.12 ± 0.02 mm and were cut to specimen dimensions corresponding

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to ASTM testing standard D-638 Type 1. Results and Discussion

Preparation and Characterisation of GO: In this work GO acts as the fundamental nanofiller material

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for the development of PBE-functionalized GO substrates. It is therefore critical to characterize the GO in terms of surface chemistry and structure which will enable both chemical derivitization as well as ultimate mechanical performance respectively. It is vital that the chemical functional groups which enable the coupling chemistry are present in sufficient numbers to provide an efficient functionalization density yet not in such a high number whereby oxidation has degraded the GO

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lattice structure. GO was prepared according to a modified Hummer’s preparative technique.37 This involved using a bulk graphite starter material a nominal flake size with <10 μm lateral dimensions. It is well known that the oxidative preparative procedure for GO introduces minor dislocations and holes where strain regions and defects are preferentially oxidized.38, Indeed, the production of GO

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under highly oxidative conditions is known to produce graphitic-type fulvic acid fragments which may bind to the surface of the GO nanosheets and may indeed hinder subsequent surface chemical processes.39 Therefore, we have taken particular care during the preparation of GO to ensure that

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the preparative procedure does not excessively oxidise the GO nanosheets. We have also taken great care to remove any such debris through washing and re-suspension steps adopted in our preparative procedure. Following the preparation of GO, structural characterisation was used to verify that the GO retains its 2-D structure and is viable for mechanical application. Figure 1a shows a representative AFM image of a GO flake indicating lateral dimensions of the order of 1 μm and a height profile demonstrating a flake height of ∼1.2 nm. TEM of a GO flake indicates an overall integrity to the material without major dislocations which would affect the 2-D structure, Figure 1b. SEM analysis of GO flakes on a silicon wafer substrate also verifies the structure as well as the dispersed nature of individual nanosheets, Figure 1c.

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Figure 1. (a) AFM micrograph of a typical GO nanosheet showing height profile inset. (b)

TEM image of GO nanosheet protruding from a hole region of a lacey carbon TEM grid.

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(c) SEM image of dispersed GO nanosheets on silicon wafer substrate.

Characterization of the prepared GO was carried out using XRD, XPS, Raman spectroscopy and FTIR

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and confirms the presence of the principle functional groups: hydroxyl, carboxyl and ketone. Figure 2a shows the XRD pattern for starting graphite material and GO. The XRD pattern for graphite shows the principle (002) peak present at ∼26.70° which corresponds to an interlayer spacing of ∼ 0.34 nm, as well as (100), (101) and (004) peaks present at 42.60°, 44.77° and 54.85° respectively.40 The XRD pattern for GO displays a broadened (002) peak of significantly lower intensity at 10.5 ° which corresponds to an interlayer spacing of ∼ 0.83 nm in addition to a small peak at 42.5 ° believed to be

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a combination of (100) and (101) reflections. The broadening of the (002) peak and the increased interlayer spacing confirms that the chemical modification of the sp2 hybridized structure has distorted the planar lattice structure by the introduction of oxygen containing functional groups.

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This is consistent with the literature and is indicative of the formation of GO.40,41 XPS was used to characterize the covalent bonding in GO. Analysis of the C1s signal for GO centred at 284.6 eV indicated multiple shoulder signals that were fitted using Gaussian curves, Figure 2b.13,42

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Signals at 285.0 eV, 286.6 eV, 287.5 eV and 289.0 eV were assigned to sp3 carbon-carbon bonding, hydroxyl, ketone and carboxylic acid groups respectively.42 The carbon to oxygen ratio (O/C) was determined to be 0.21. Survey spectra for GO and graphite as well as the C1s core level spectrum for graphite are contained in the supporting information (SI, Figure S5-S7). The introduction of the oxygen containing functional groups was additionally characterized using Raman spectroscopy to identify the introduction of the characteristic D (disorder) band in the spectrum of GO which corresponds to sp3 hybridized carbon bonding as part of the GO structure.42,43 Figure 2c shows the Raman spectrum for the graphite starter material and includes the G-band at 1580 cm-1 due to the in-plane E2g mode and the 2D peak at 2725 cm-1 due to second-order zone boundary phonons.43 The spectrum of GO indicates the G band which has shifted to 1585 cm-1 as well as the presence of a 7

ACCEPTED MANUSCRIPT large D band at 1350 cm-1 assigned here to lattice disorder due to the presence of oxygen containing functional groups at defect sites, dislocations and strain regions.43 FTIR spectroscopy was also used to chemically characterize GO by identifying the vibrational bands associated with the functional groups, Figure 2d. The spectrum of GO indicates the presence of a number of intense vibrational bands which are assigned to oxygen containing functional groups as well as the C=C bonding.45 A strong band centred at 3366 cm-1 is assigned to the hydroxyl (-OH) groups of alcohol and carboxylic

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acid groups. A band at 1715 cm-1 is assigned to the carboxylic acid carbonyl group, however it is noted that at slightly higher wavenumbers (1740 cm-1) a slight shoulder band is also observed. This is assigned to ketone groups and is masked beneath the strong vibration of carboxylic acid carbonyl. The C-OH bending vibration is present at 1411 cm-1, while C-O stretching vibrations are observed at

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1222 cm-1 and 1040 cm-1. A strong vibration band at 1614 cm-1 is assigned to C=C stretching of the GO lattice. It is clear from the combined chemical characterization that the GO has substantial

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chemical functionality as a result of the strongly oxidizing preparative procedure. The presence of these functional groups, particularly the carboxylic acid groups underpins this work and it is therefore gratifying that the characterization indicates that a substantial proportion of these groups

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exist as part of the GO structure.

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ACCEPTED MANUSCRIPT Figure 2. (a) XRD pattern for graphite and GO, indicating substantial shifting of the (002) peak corresponding to an increase in the inter-plane spacing. (b) XPS C1s signal for GO indicating covalent C-O bonding for hydroxyl, ketone, and carboxylic acid groups. (c) Raman spectroscopy for graphite and GO showing substantially enhanced D band indicative of disorder induced by oxidative functionalization. (d) FTIR spectrum of GO showing vibrational bands assigned to covalent bonding of hydroxyl, carboxylic acid and C=C bonding.

Preparation of Covalently Functionalized PBE-GO: Following the preparation of GO, PBE chains were

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covalently bound to the surface of the nanosheets by the formation of ester linkages. Specifically, ester formation was mediated using a carbodiimide activating agent (N,N’-dicyclohexylcarbodiimide (DCC)) to couple GO carboxylic acid groups and hydroxyl groups which are present on each repeat

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unit of the PBE chains, Figure 3.

Figure 3. Reaction scheme for the DCC mediated esterification between GO and PBE.

Following the functionalization of GO with PBE chains, PBE-GO material was subjected to a rigorous washing procedure to ensure that all traces of any unreacted PBE material as well as any trace of

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fulvic acid graphitic type fragments were removed. The product material (PBE-GO) was characterized using FTIR to determine the presence of polymer chains attached to the GO nanosheets and the

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covalent ester bonding between the GO and PBE, Figure 4a.

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Figure 4. (a) FTIR spectra of GO, PBE polymer, and PBE-functionalized GO. (b) UV-Vis spectra of GO, PBE and PBE-functionalized GO. (c) TGA traces for GO, PBE and PBE-functionalized GO indicating mass loss and the derivative mass loss traces as a function of temperature.

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ACCEPTED MANUSCRIPT The FTIR spectrum of GO in Figure 4a indicates the position of the carboxylic acid C=O centred at ∼1715 cm-1. The spectrum of PBE indicates multiple bands which correspond to hydroxyl groups, aromatic C=C bonding, aromatic C-H (aC-H), alkyl C-H, as well as ether C-O vibrational bands. Importantly in this work, clear assignment of hydroxyl groups is possible due to the strong stretching band centred at 3324 cm-1 and the corresponding C-OH bending vibration at 1231 cm-1 and C-O

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stretching band at 1037 cm-1 (Figure 4a (PBE)). Following the coupling of GO carboxylic acid groups and PBE hydroxyl groups the formation of the ester is confirmed by a shift in the position of the carbonyl C=O absorption (Figure 4a (PBE-GO)). The position of the free acid C=O at 1715 cm-1 has shifted to 1735 cm-1 indicative of ester formation. The hybrid spectrum of the PBE-GO is dominated by the molecular signature of the PBE chains as expected. However, broadening in the region of the

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–OH stretching band is observed, with the central location of the band being shifted approximately ∼20 cm-1 to ∼3385 cm-1. The position of the C-O stretching band, which is most intense in GO at 1040

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cm-1, is also significantly reduced in intensity and shifted downfield by ∼6 cm-1 to 1034 cm-1 in the coupled PBE-GO. While it is complex to assign shifting behaviour accurately to the result of the ester coupling reaction, it is speculated that the observed behaviour may be assigned to reduced vibrational freedom and stiffening of the GO structure as a result of surface and side group bonding interactions.

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PBE-GO was characterized using UV-Vis in order to identify PBE bound to the surface of the GO nanosheets, Figure 4b. The UV-Vis spectrum of GO shows a peak at 235 nm corresponding to the ππ* transition of aromatic C-C bonds.46 The spectrum of the unbound PBE polymer exhibited peaks as a shoulder at 228 nm and peaks at 278 nm and 450 nm. Following PBE-functionalization of GO, the

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spectrum of PBE-GO exhibited peaks corresponding directly to the PBE polymer chains as a shoulder at 230 nm and as a peak at 280 nm. The peak at 450 nm which had been observed in the spectrum of the free PBE polymer was not observed in the spectrum of the PBE-GO due to the strong

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scattering contribution of GO in this region. The analysis of PBE-GO using UV-Vis spectroscopy confirms the presence of PBE chains grafted from the surface of the GO nanosheets. In order to quantify PBE chains on the surface of the GO nanosheets, TGA was performed on the samples, Figure 4c. The TGA trace of weight (%) as a function of temperature for GO exhibited significant decomposition at ∼175 °C which is also observed as a defined peak in the derivative weight (%) as a function of temperature. This reasonably low degradation temperature for GO in comparison with graphene is attributed to the oxidized nature of the nanosheets which undoubtedly contain oxidized holes and oxidized defects in the structure which act as preferential sites for structural decomposition as the temperature is increased.41 The TGA trace for PCE exhibits a well-defined decomposition profile with a maximum loss in the region of 400-500 °C which is verified by the single 11

ACCEPTED MANUSCRIPT peak in the derivative weight (%) trace at 420 °C. The TGA trace for PBE-GO exhibits a combined signature of the decomposition of both GO and PBE. The derivative weight (%) shows peaks centred at 175 °C and 420 °C which are identical to the decomposition of the individual GO and PBE components. The TGA trace indicates that ∼ 40% of the total mass of PBE-GO is due to PBE chains and confirms the coupled hybrid structure of the GO nanosheets grafted with PBE polymer chains.

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Preparation and Characterisation of PBE-GO Thin-Film Nanocomposites: Utilization of PBE chains which are covalently grafted to the surface of GO nanosheets was investigated by forming PBE thinfilm nanocomposites across a range of mass fraction loadings from 0.1 wt. % to 1.0 wt. %. A critical feature for the formation of nanocomposites is the efficient dispersion of the nanofiller within the

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polymer matrix. In cases where aggregation of the nanofiller occurs, these aggregates may act as defect centres which may facilitate premature failure.28,33-35 Therefore a combination of an optimum dispersion of the nanofiller as well as optimized interfacial interaction between nanofiller and matrix

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is required.20,21 In this work, the functionalization of GO with PBE chains provides direct chemical compatibility and optimises the interfacial interactions between the PBE host matrix and the nanofiller. This is achieved by creating dispersions of PBE-GO across a range of mass fraction loadings (0.1, 0.25, 0.50, 0.75, 1.00 wt. %) in a relatively large volume of solvent and gradually introducing the polymer which is dissolved and dispersed amongst the solubilized PBE-GO

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nanosheets, Figure 5a. In this manner, by the addition of more PBE to a sonicating suspension of PBE-GO, the polymer chains act as mechanical buffers which prevent the sheets from aggregating and reach an optimum dispersion of the PBE-GO within the PBE solvent solution. Subsequent removal of the solvent by evaporation allows thin films of the nanocomposites to be cast (Figure

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5b,c) and punched to test specimen dimensions which correspond to ASTM-D638 Type-1 international standard specification for tensile testing (Figure 4d).47 Thin-film nanocomposites with PBE-GO in PBE (PBE-GO:PBE) and the control GO without any functionalization in PBE (GO:PBE) were

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tested across the series of mass-fraction loadings from 0.1 wt. % to 1.0 wt. % (minimum of six specimens per loading fractions) using a micro-mechanical tensile testing apparatus (Figure 5e). Representative tensile mechanical properties of the thin-film nanocomposites PBE-GO:PBE and GO:PBE are shown in Figure 5f and 5g respectively. It is noted that the PBE control material which has no nanofiller content was found to exhibit stiff and brittle behaviour with a Young’s modulus of 1.20 GPa, an ultimate tensile strength (UTS) of 40.7 MPa and a low strain-to-failure of ∼8.7 %. PBE-functionalized-GO:PBE nanocomposite thin films were found to exhibit substantially improved tensile mechanical properties relative to the control PBE, Figure 5f,h. Young’s modulus values for PBE-GO:PBE nanocomposites were all increased across the range of mass fraction loadings with a

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ACCEPTED MANUSCRIPT maximum value of 1.59 GPa at 0.25 wt. %, an increase of ∼30 % over that of the PBE control (1.20 GPa, Figure 5h). While such an increase in Young’s modulus is not surprising on addition of GO nanofiller material, what is important is the ability of the nanofiller to continue to allow the transfer of the tensile stress in order to reach higher strength values. The inset in Figure 5f shows the tensile stress-strain curves for the PBE-GO:PBE nanocomposites as well as the PBE control between 0-10 %

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strain. This clearly shows that at the initial onset of mechanical strain the nanocomposites are substantially stiffened relative to the PBE control and display mechanical properties which deviate from those of the control PBE material. Ultimate tensile strength (UTS) values for PBE-GO:PBE nanocomposites were significantly higher across the series than that of the PBE control reaching a maximum of 67 MPa, an increase of ∼60 % over that of the PBE control (41 MPa, Figure 5h). This is

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indicative of efficient stress transfer to the PBE-GO substrates dispersed within the PBE matrix. Failure of the nanocomposite at the point of reaching a high UTS value would simply be indicative of

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a strengthened but embrittled polymer. Critical to the prevention of failure at high UTS is the optimised interface between matrix and filler which upon the increase in tensile stress allows the material to yield, harden and start to neck as the material becomes drawn but importantly does not fail. The PBE-GO:PBE series exhibits exceptional increases in strain-to-failure (εB) values indicating an extremely high degree of interaction between the PBE-GO and the PBE matrix, Figure 5h. Maximum enhancement of the % strain-to-failure (εB) of 183 % was achieved for 1.0 wt. % PBE-GO, an increase

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of more than 20 fold compared to that of the PBE control (9%). Correspondingly, the toughness, taken as the integrated area under the stress-strain curve, indicate that the maximum toughness achieved for 1.0 wt. % PBE-GO is representative of more than a 35 fold increase in toughness, 9819 MPa over 269 MPa for PBE, Figure 5h. The extremely high degree of reinforcement achieved over

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the series of PBE-GO:PBE nanocomposites is attributed to an optimized nanofiller-matrix interaction on account of the chemical compatibilization between identical PBE polymer chains. To test this

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hypothesis, identical control experiments were conducted by preparing GO-PBE nanocomposites which used GO which did not have surface grafted PBE polymer chains.

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Figure 5. (a) Photograph of PBE-GO:PBE dispersions in DMF. (b) PBE-GO:PBE solutions deposited on glass microscope slides. (c) Representative PBE-GO:PBE thin-film at 0.50 wt. % loading following drying and removal from glass substrate. (d) PBE-GO:PBE thin-film specimen for mechanical characterization. (e) Mechanical characterization of PBE-GO:PBE specimen using DMA micro-tensile apparatus. (f) Representative stress-strain curves for PBE-GO:PBE nanocomposites with inset showing initial strain

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ACCEPTED MANUSCRIPT (0-10 %). (g) Representative stress-strain curves for GO:PBE nanocomposites with inset showing initial strain (0-10 %). (h) Mechanical characterization data showing Young’s Modulus (Y), Ultimate Tensile Strength (UTS), strain-to-failure (εB), and Toughness values measured for PBE and PBE-GO:PBE, GO:PBE nanocomposites as a function of mass fraction loading.

Tensile mechanical properties of the control nanocomposites of GO:PBE were measured across the mass fraction loading range 0.1 wt. %, 0.25 wt. %, 0.5 wt. %, 0.75 wt. % 1.0 wt. % prepared in an

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identical fashion to the PBE-GO:PBE range. Young’s modulus values for the GO:PBE range averaging ∼0.94 GPa were found to be markedly lower than that of the PBE control (1.20 GPa, Figure 5h). This is particularly significant as the addition of GO to the PBE matrix may be expected to give modulus values comparable to that of the PBE-GO:PBE series upon initial tensile stress simply on account of

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the inclusion of GO.28 The inset in Figure 5g shows that with the initial onset of strain the GO:PBE nanocomposites exhibit Young’s modulus values which are generally similar to that of the PBE

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control. This indicates that the addition of GO to the PBE matrix is not reinforcing the matrix to any great extent and may be indicative of poor dispersion within the matrix. UTS values for the GO:PBE series were found to be essentially unchanged from the PBE control (41 MPa) or slightly higher in the case of 0.1 wt. % (57 MPa) and 1.0 wt. % (48 MPa) loadings, Figure 5h. In stark contrast to the extremely high strain-to-failure (εB) values for the PBE-GO:PBE series, the control GO:PBE series exhibited comparatively low values of ∼6 %, all lower than the PBE control, 9 %. The combination of

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lower strain-to-failure and marginal increase in UTS values illustrates that addition of GO to the PBE matrix did not efficiently reinforce the polymer but had a detrimental effect by embrittling the matrix and causing failure at lower strain. Toughness of the GO:PBE nanocomposites is correspondingly low. The integrated area under the curves indicates that in the cases where higher

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UTS was achieved (0.1 wt. %, and 1.0 wt. %), toughness values are moderately higher than the rest of the series at 255 MPa and 207 MPa respectively, however across the GO:PBE series, the

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toughness values are lower than that of the PBE control, 269 MPa, Figure 5h.

The addition of a series of mass fraction loadings of PBE-GO to PBE was found to enhance the Young’s modulus of the polymer, with extremely high values of UTS, strain-to-failure and toughness were achieved, indicating a true enhancement of the mechanical properties of the PBE polymer. The exceptional enhancements in mechanical properties exhibited by PBE-GO:PBE nanocomposites relative to the control GO:PBE nanocomposites and PBE are assigned to the impact of PBE polymer chains covalently grafted to the GO nanosheets. It is clear that the tailored chemistry has a 15

ACCEPTED MANUSCRIPT substantial impact on the failure mechanism of the PBE-GO:PBE nanocomposite since the GO-PBE control nanocomposites fail at relatively low degrees of strain (<10 %). The mechanism of interaction which supports these properties is assigned to the chemical compatibility and mechanical interaction between PBE chains attached to the GO sheets and the PBE matrix. The attractive forces between PBE chains in the unmodified polymer are due to a combination of inter- and intra-chain

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attractive forces consisting of aromatic π-π interaction as well as hydrogen bonding interactions between ether and hydroxyl groups, and methylene and hydroxyl hydrogen atoms. We would assume of course that a great deal of additional mechanical entanglement of interlocking polymer chains provide integrity to the PBE matrix.

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Following characterization of PBE-GO nanosheets using FTIR, UV-Vis spectroscopy and TGA, the attractive forces between the PBE-GO and the PBE matrix are dictated by the aforementioned PBEPBE attractive forces. Mechanical characterization of the series of control samples using unmodified

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GO resulted in only moderate increases in the UTS values as well as reduced strain-to-failure values (εB). Presumably GO will have hydrogen bonding interactions and possibly even some slight degree of π-π interaction, as well a mechanical wrapping or entanglement with the PBE matrix. However, despite a small degree of presumably favourable attractive forces, ultimately the mismatch in surface energy between the GO and PBE matrix are likely to cause aggregation of the nanofiller

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during processing and the inclusion of defect centres within the PBE matrix. It is also assumed that attractive forces between GO functional groups and the PBE matrix would also exist for the PBE-GO material but it is believed that the covalently attached PBE chains dictate firstly the dispersion of the nanosheets within the forming PBE matrix and secondly the interlocking entanglement by direct

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chemical and mechanical interaction as the nanocomposite thin films are formed. The result of compatibilized surface chemistry of PBE-GO is that the PBE-GO:PBE nanocomposites exhibit effective reinforcement across the principle mechanical parameters characterized: Young’s modulus, UTS,

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strain-to-failure (εB), and toughness (T). The significant enhancement in the mechanical properties of the PBE-GO:PBE thin-film nanocomposites has been attributed to the tailoring of the interfacial chemistry between GO and the PBE matrix. The compatibilization of GO nanosheets produced via the covalent grafting of PBE chains at GO hydroxyl groups is a versatile approach which may be applied to a range of alternative polymer systems. Previous reports in the literature have demonstrated closely related methodology by forming poly(L-lactic acid) (PLLA) functionalized GO and MWNTs to yield impressive enhancement in mechanical properties.22-24,35,48,49 Unlike approaches which demonstrate the formation of polymer chains covalently bound to GO via initiator mediated polymerisation, the ‘grafting from’ approach

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ACCEPTED MANUSCRIPT described in this work offers a number of distinct advantages. The intrinsic lack of chemical compatibilization between initiator-modified nanofillers within solvent or monomer systems is overcome in this work by achieving direct chemical compatibilization between polymer chains (PBE) and the polymer-functionalized GO nanosheets (PBE-GO). An additional advantage of the ‘grafting from’ approach is that it allows nanomaterials with tailored surface chemistry to be prepared

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separately and controllably added to monomer/polymer systems in the same manner as conventional anti-oxidant, pigment or toughening additives. A further distinct advantage is the commercial availability of end-terminated and functionalized oligomer and polymer species which may be bound to surface-modified nanomaterials through appropriate coupling chemistry such as amide, ester, ether and urethane formation. The formation of separate nanomaterial additives also

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gives the opportunity to assess the quality of a functionalized nanomaterial additive prior to use in order to assess the efficiency of the modification process. This is impossible in the ‘grafting to’

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scenario where nanomaterials may be co-dispersed within solvent or polymer systems and the efficiency of the grafting surface chemistry may be undetermined prior to mechanical or thermal analysis. Conclusion

In order to tailor the chemical compatibilization between GO and PBE, GO has been covalently

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functionalized using PBE polymer chains. Characterization of GO has been performed using a range of microscopy and spectroscopic characterization techniques including TEM, AFM, SEM, XRD, XPS, FTIR, Raman spectroscopy, UV-Vis and TGA to determine the structural integrity and the chemical characteristics of the GO starting material. Covalent coupling of the PBE chains to GO via ester

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formation has been confirmed by characterization using FTIR. Thin-film nanocomposites of PBE, PBE with GO and PBE-GO across a series of mass fraction loadings have been formed (0.10, 0.25, 0.50, 0.75, 1.00 wt. %). The addition of PBE-GO to PBE was found to significantly enhance the mechanical

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properties above that of PBE while the addition of unmodified-GO was found to generally decrease the mechanical properties. Maximum enhancements in Young’s modulus (∼30%, 1.59 GPa), UTS (∼60%, 67 MPa), a more than 20 fold increase in the strain-to-failure (εB) (183 %) and a more than 35 fold increase in toughness was achieved (9819 MPa) for PBE-GO:PBE over that of pristine PBE. These enhancements in mechanical properties are assigned to the contribution of PBE chains grafted to GO and thus the direct compatibilization of the nanofiller to the PBE matrix. The advantages of the ‘grafting from’ approach for GO have been discussed in the context of this work and also in the context of the application of similar processing methodologies to alternative polymer and monomer systems. The approach described here is particularly significant due to the fact that the nanomaterial additives have been demonstrated to have significant enhancement of the mechanical properties of 17

ACCEPTED MANUSCRIPT polymers at low percentage loadings (≤ 1%). The relevance of this work in terms of industrial application is considered in the context of the enhancement of PBE, which typically exhibits brittle and low strength properties. GO is rapidly gaining research attention for its ability to transform the mechanical properties of polymers such as PVA and epoxy resin by enhancing strength and fracture toughness and to create commercial products.50-53 In view of this, the enhancement of PBE following

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the addition of GO illustrates the potential to transform the mechanical performance of polymers at very low mass fraction loadings of nanofiller materials. This may indicate the potential for typically low performing polymers to be applied in new functional embodiments following their modification and may create new markets such as packaging materials for electronics, barrier layers or protective coatings and paints for materials which require high mechanical strength and toughness within

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aerospace or automotive applications.

Supporting Information

AFM analysis of GO nanosheets (Figure S3,S4). XPS survey spectra for graphite and GO (Figure S5,S6), C1s core level spectrum for graphite starting material (Figure S7). Corresponding Author

Notes

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* Telephone: +44 20 8943 6434, E-mail: [email protected]

The authors declare no competing financial interest.

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Acknowledgements

The authors acknowledge funding through the UK National Measurement System (NMS) strategic

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capability programme.

References

1. C. Ferrari, et al., Science and technology roadmap for graphene, related two-dimensional crystals, and hybrid systems, Nanoscale, 2015, 7(11) 4598-4810.

2. D. Chen, H. Feng and J. Li, Graphene Oxide: Preparation, Functionalization, and Electrochemical Applications, Chem. Rev., 2012, 112(11), 6027-6053.

18

ACCEPTED MANUSCRIPT 3. V. Georgakilas, M. Otyepka, A. B. Bourlinos, V. Chandra, N. Kim, K. C. Kemp, P. Hobza, R. Zboril and K. S. Kim, Functionalization of Graphene: Covalent and Non-Covalent Approaches, Derivatives and Applications, Chem. Rev., 2012, 112, 6156-6214.

4. Y. Zhu, S. Murali, W. Cai, X. Li, J.W. Suk, J.R. Potts and R.S. Ruoff, Graphene and graphene

RI PT

oxide: synthesis, properties, and applications. Adv. Mater., 2010, 22(35), 3906-3924.

5. Stevens, E. Dessiatova, D. A. Hagen, A. D. Todd, C. W. Bielawski and J. C. Grunlan, Low-

Temperature Thermal Reduction of Graphene Oxide Nanobrick Walls: Unique Combination

ACS Appl. Mater. Interfaces, 2014 6(13), 9942-9945.

SC

of High Gas Barrier and Low Resistivity in Fully Organic Polyelectrolyte Multilayer Thin Films,

M AN U

6. Y. Yang, , W. Rigdon, X. Huang and X. Li, Enhancing graphene reinforcing potential in composites by hydrogen passivation induced dispersion, Sci. Rep., 2013, 4.

7. F. Torrisi, T. Hasan, W. Wu, Z. Sun, A. Lombardo, T. S. Kulmala, G.-W. Hsieh, S. Jung, F. Bonaccorso, P. J. Paul, D. Chu and Andrea C. Ferrari, Inkjet-Printed Graphene Electronics,

TE D

ACS Nano, 2012 6(4), 2992-3006.

8. M. S. Choi, G.-H. Lee, Y.-J. Yu, D.-Y. Lee, S. H. Lee, P. Kim, J. Hone and W. J. Yoo, Controlled charge trapping by molybdenum disulphide and graphene in ultrathin heterostructured

EP

memory devices, Nat. Commun., 2013, 4, 1624.

9. W. K. Chee, H. N. Lim, Z. Zainal, N. M. Huang, I. Harrison and Y. Andou, Flexible Graphene-

AC C

Based Supercapacitors: A Review, J. Phys. Chem. C, 2016 120(8), 4153-4172.

10. A. Hirsch, J. M. Englert and F. Hauke, Wet Chemical Functionalization of Graphene, Acc. Chem. Res., 2013, 46, 87-96.

11. D.R. Dreyer, A.D. Todd and C.W. Bielawski, Harnessing the chemistry of graphene oxide, Chem. Soc. Rev., 2014, 43(15), 5288-5301.

19

ACCEPTED MANUSCRIPT 12. J. M. Englert, C. Dotzer, G. Yang, M. Schmid, C. Papp, J. M. Gottfried, H.-P. Steinrück, E. Spiecker, F. Hauke and A. Hirsch, Covalent Bulk Functionalization of Graphene, Nat. Chem., 2011, 3, 279-286.

13. D.R. Dreyer, S. Park, C.W. Bielawski and R.S. Ruoff, The chemistry of graphene oxide, Chem.

RI PT

Soc. Rev., 2010, 39(1), 228-240.

14. J. Park, M. Yan, Covalent Functionalization of Graphene with Reactive Intermediates, Acc. Chem. Res., 2013, 46, 181-189.

SC

15. K. Choi, S. Nam, Y. Lee, M. Lee, J. Jang, S. J. Kim, Y. J. Jeong, H. Kim, S. Bae, J.-B. Yoo, S. M. Cho, J.-B. Choi, H. K. Chung, J.-H. Ahn, C. E. Park and B. H. Hong, Reduced Water Vapor

ACS Nano, 2015, 9(6), 5818-5824.

M AN U

Transmission Rate of Graphene Gas Barrier Films for Flexible Organic Field-Effect Transistors,

16. P. Tzeng, B. Stevens, I. Devlaming and J. C. Grunlan, Polymer–Graphene Oxide Quadlayer Thin-Film Assemblies with Improved Gas Barrier, Langmuir, 2015, 31(21), 5919-5927.

TE D

17. J. Xiong, F. Jiang, H. Shi, J. Xu, C. Liu, W. Zhou, Q. Jiang, Z. Zhu and Y. Hu, Liquid Exfoliated Graphene as Dopant for Improving the Thermoelectric Power Factor of Conductive PEDOT:PSS Nanofilm with Hydrazine Treatment, ACS Appl. Mater. Interfaces, 2015, 7(27),

EP

14917-14925.

18. K. M. F. Shahil and A. A. Balandin, Graphene–Multilayer Graphene Nanocomposites as Highly

AC C

Efficient Thermal Interface Materials, Nano Lett., 2012, 12(2), 861-867.

19. Xiang, P. J. Cox, A. Kukovecz, B. Genorio, D. P. Hashim, Z. Yan, Z. Peng, C.-C. Hwang, G. Ruan, E. L. G. Samuel, P. M. Sudeep, Z. Konya, R. Vajtai, P. M. Ajayan and J. M. Tour, Functionalized Low Defect Graphene Nanoribbons and Polyurethane Composite Film for Improved Gas Barrier and Mechanical Performances, ACS Nano, 2013, 7(11), 10380-10386.

20. Tkalya, M. Ghislandi, R. Otten, M. Lotya, A. Alekseev, P. van der Schoot, J. Coleman, G. de With and C. Koning, Experimental and Theoretical Study of the Influence of the State of

20

ACCEPTED MANUSCRIPT Dispersion of Graphene on the Percolation Threshold of Conductive Graphene/Polystyrene Nanocomposites, ACS Appl. Mater. Interfaces, 2014, 6(17), 15113-15121.

21. K. D. Papadimitriou, E. N. Skountzos, S. S. Gkermpoura, I. Polyzos, V. G. Mavrantzas, C. Galiotis and C. Tsitsilianis, Molecular Modelling Combined with Advanced Chemistry for the

RI PT

Rational Design of Efficient Graphene Dispersing Agents, ACS Macro Letters, 2016, 5(1), 2429.

22. Voylov, T.Saito, B. Lokitz, D. Uhrig, Y. Wang, A. Agapov, A. Holt, V. Bocharova, A. Kisliuk and A. P. Sokolov, Graphene Oxide as a Radical Initiator: Free Radical and Controlled Radical

SC

Polymerization of Sodium 4-Vinylbenzenesulfonate with Graphene Oxide, ACS Macro Lett.,

M AN U

2016, 5(2), 199-202.

23. L. Yan, Y. B. Zheng, F. Zhao, S. Li, X. Gao, B. Xu, P. S. Weiss and Y. Zhao, Chemistry and physics of a single atomic layer: strategies and challenges for functionalization of graphene and graphene-based materials, Chem. Soc. Rev., 2012, 41(24), 97-114.

24. H. K. F. Cheng, N. G. Sahoo, Y. P. Tan, Y. Pan, H. Bao, L. Li, S. H. Chan and J. Zhao, Poly(vinyl

TE D

alcohol) Nanocomposites Filled with Poly(vinyl alcohol)-Grafted Graphene Oxide, ACS Appl. Mater. Interfaces, 2012, 4(5), 2387-2394.

25. M. Zhang, Y. Li, Z. Su and G. Wei, Recent advances in the synthesis and applications of

Bisphenol A (BPA): 2016 World Market Outlook and Forecast up to 2020 -

AC C

26.

EP

graphene–polymer nanocomposites, Polym. Chem., 2015, 6(34), 6107-6124.

https://mcgroup.co.uk/researches/bisphenol-a-bpa

27. A. J. Kinloch, Toughening Epoxy Adhesives to Meet Today’s Challenge, Mater. Res. Soc., 2003, 28(06), 445–448.

28. N. Domun, H. Hadavinia, T. Zhang, T. Sainsbury, G. H. Liaghat and S. Vahid, Improving the fracture toughness and the strength of epoxy using nanomaterials–a review of the current status, Nanoscale, 2015, 7(23), 10294-10329.

21

ACCEPTED MANUSCRIPT 29. X. Tang, Y. Zhou and M. Peng, Green Preparation of Epoxy/Graphene Oxide Nanocomposites Using a Glycidylamine Epoxy Resin as the Surface Modifier and Phase Transfer Agent of Graphene Oxide, ACS Appl. Mater. Interfaces, 2016, 8(3), 1854-1866.

30. J. R. Potts, S. Murali, Y. Zhu, X. Zhao and R. S. Ruoff, Microwave-Exfoliated Graphite

RI PT

Oxide/Polycarbonate Composites, Macromolecules, 2011, 44(16), 6488-6495.

31. D. R. Bortz, E. G. Heras and I. Martin-Gullon, Impressive Fatigue Life and Fracture Toughness Improvements in Graphene Oxide/Epoxy Composites, Macromolecules, 2012, 45(1), 238-

SC

245.

32. M. Naebe, J. Wang, A. Amini, H. Khayyam, N. Hameed, L. H. Li, Y. Chen and B. Fox,

M AN U

Mechanical property and structure of covalent functionalised graphene/epoxy nanocomposites, Sci. Rep., 2014, 4.

33. J. Wei, T. Vo and F. Inam, Epoxy/graphene nanocomposites–processing and properties: a review, RSC Adv., 2015, 5(90), 73510-73524.

TE D

34. Z. Li, R. J. Young and I. A. Kinloch, Interfacial Stress Transfer in Graphene Oxide Nanocomposites, ACS Appl. Mater. Interfaces, 2013, 5(2), 456-463.

35. X. Zhao, Q. Zhang and D. Chen, Enhanced mechanical properties of graphene-based

EP

poly(vinyl alcohol) composites, Macromolecules, 2010, 43, 2357–2363.

AC C

36. T. Ramanathan, A.A. Abdala, S. Stankovich, D.A. Dikin, M. Herrera-Alonso, R.D. Piner, D.H. Adamson, H.C. Schniepp, X. Chen, R.S. Ruoff and S.T. Nguyen, Functionalized graphene sheets for polymer nanocomposites, Nat. Nanotechnol., 2008, 3(6), 327-331.

37. W. S. Hummers and R. E. Offeman, Preparation of Graphitic Oxide, J. Am. Chem. Soc. 1958, 80, 1339-1339.

22

ACCEPTED MANUSCRIPT 38. K. Erickson, R. Erni, Z. Lee, N. Alem, W. Gannett and A. Zettl, Determination of the local chemical structure of graphene oxide and reduced graphene oxide, Adv. Mater., 2010, 22, 40, 4467-4472.

39. J. P. Rourke, P. A. Pandey, J. J. Moore, M. Bates, I. A. Kinloch, R. J. Young and N. R. Wilson,

Sheets, Angew. Chem. Int. Ed., 2011, 50(14), 3173-3177.

RI PT

The Real Graphene Oxide Revealed: Stripping the Oxidative Debris from the Graphene-like

40. G. Sun, X. Li, Y. Qu, Preparation and characterization of graphite nanosheets from

SC

detonation technique, Mater. Lett., 2008, 62, 703-706.

41. D. A. Dikin, S. Stankovich, E. J. Zimney, R. D. Piner, G. H. B. Dommett, G. Evmenenko, S.-B. T.

2007, 448, 7152, 457-460.

M AN U

Nguyen and R. S. Ruoff, Preparation and characterization of graphene oxide paper, Nature,

42. D. Yang, , A. Velamakanni, G. Bozoklu, S. Park, M. Stoller, R. D. Piner, S. Stankovich et al., Chemical analysis of graphene oxide films after heat and chemical treatments by X-ray

TE D

photoelectron and Micro-Raman spectroscopy, Carbon, 2009, 47(01), 145-152.

43. K. N. Kudin, B. Ozbas, H. C. Schniepp, R. K. Prud'Homme, I. A. Aksay and R. Car, Raman spectra of graphite oxide and functionalized graphene sheets, Nano Lett., 2008, 8(01), 36-41.

EP

44. A. C. Ferrari, J. C. Meyer, V. Scardaci, C. Casiraghi, M. Lazzeri, F. Mauri, S. Piscanec, D. Jiang, K. S. Novoselov, S. Roth and A. K. Geim, Raman Spectra of Graphene and Graphene Layers,

AC C

Phys. Rev. Lett., 2006, 97, 187401-1−187401-4.

45. S. Stankovich, R. D. Piner, S.-B. T. Nguyen and R. S. Ruoff, Synthesis and exfoliation of isocyanate-treated graphene oxide nanoplatelets, Carbon, 2006, 44(15), 3342-3347.

46. D. Li, M. B. Mueller, S. Gilje, R. B. Kaner and G. G. Wallace, Processable aqueous dispersions of graphene nanosheets, Nat. Nanotechnol., 2008, 3(02), 101-105.

47. ASTM, Standard Test Method for Tensile Properties of Plastics. ASTM Standard D638-14. American Society for Testing and Materials; 2014.

23

ACCEPTED MANUSCRIPT 48. W. Li, Z. Xu, L. Chen, M. Shan, X. Tian, C. Yang, H. Lv and X. Qian, A facile method to produce graphene oxide-g-poly (L-lactic acid) as an promising reinforcement for PLLA nanocomposites, Chem. Eng. J., 2014, 237, 291-299.

RI PT

49. Q. Jing, W. Liu, Y. Pan, V. V. Silberschmidt, L. Li and Z. Dong, Chemical functionalization of graphene oxide for improving mechanical and thermal properties of polyurethane composites, Mater. Des., 2015, 85, 808-814.

50. X. Zhao, Q. Zhang, D. Chen, P. Lu. Enhanced mechanical properties of graphene-based

M AN U

SC

poly(vinyl alcohol) composites, Macromolecules, 2010, 43, 2357-2363.

51. H. Ning, J. Li, N. Hu, C. Yan, Y. Yiiu, L. Wu, F. Liu, J. Zhang, Interlaminar mechanical properties of carbon fiber reinforced plastic laminates modified with graphene oxide interleaf, Carbon, 2015, 91, 224-233.

K. Krishnamoorthy, K. Jeyasubramanian, M. Premanathan, G. Subbiah, H. S. Shin, S. J. Kim,

TE D

52.

Graphene oxide nanopaint, Carbon, 2014, 72, 328-337.

EP

53. A. J. Shakir, D. C. Culita, J. Calderon-Moreno A. Musuc, O. Carp, G. Ionita, P. Ionita, Covalently grafted TEMPO on graphene oxide: A composite material for selective oxidations

AC C

of alcohols. Carbon, 2016, 105, 607-614.

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Table of Contents Graphic

Thin film polymer-functionalized graphene oxide nanocomposites demonstrating extreme

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mechanical enhancement due to the chemically compatibilized nanofiller-polymer matrix

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interaction.

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