Trimellitic Anhydride Modified Graphene

Subscriber access provided by University of Winnipeg Library

Article

Polyethylene Terephthalate/Trimellitic Anhydride Modified Graphene Nanocomposites Shigeru Aoyama, Issam Ismail, Yong Tae Park, Yuki Yoshida, Christopher W. Macosko, and Toshiaki Ougizawa ACS Appl. Nano Mater., Just Accepted Manuscript • DOI: 10.1021/acsanm.8b01525 • Publication Date (Web): 22 Oct 2018 Downloaded from http://pubs.acs.org on October 22, 2018

Just Accepted “Just Accepted” manuscripts have been peer-reviewed and accepted for publication. They are posted online prior to technical editing, formatting for publication and author proofing. The American Chemical Society provides “Just Accepted” as a service to the research community to expedite the dissemination of scientific material as soon as possible after acceptance. “Just Accepted” manuscripts appear in full in PDF format accompanied by an HTML abstract. “Just Accepted” manuscripts have been fully peer reviewed, but should not be considered the official version of record. They are citable by the Digital Object Identifier (DOI®). “Just Accepted” is an optional service offered to authors. Therefore, the “Just Accepted” Web site may not include all articles that will be published in the journal. After a manuscript is technically edited and formatted, it will be removed from the “Just Accepted” Web site and published as an ASAP article. Note that technical editing may introduce minor changes to the manuscript text and/or graphics which could affect content, and all legal disclaimers and ethical guidelines that apply to the journal pertain. ACS cannot be held responsible for errors or consequences arising from the use of information contained in these “Just Accepted” manuscripts.

is published by the American Chemical Society. 1155 Sixteenth Street N.W., Washington, DC 20036 Published by American Chemical Society. Copyright © American Chemical Society. However, no copyright claim is made to original U.S. Government works, or works produced by employees of any Commonwealth realm Crown government in the course of their duties.

Page 1 of 44 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

ACS Applied Nano Materials

Polyethylene Terephthalate/Trimellitic Anhydride Modified Graphene Nanocomposites Shigeru Aoyama,†§Issam Ismail,†å* Yong Tae Park,†¶Yuki Yoshida,ë Christopher W. Macosko,†

and Toshiaki Ougizawa‡* †Department

of Chemical Engineering and Materials Science, University of Minnesota, Minneapolis, MN 55455, USA,

‡Department

of Materials Science and Engineering, Tokyo Institute of Technology, 2-12-1-S8-33, Ookayama, Meguro-ku, Tokyo 152-8552, Japan, Japan,

§Films

& Film Products Research Laboratories, Toray Industries, Inc., 1-1-1 Sonoyama, Otsu, Shiga 520-8558, Japan,

å

ë

Department of Chemical Engineering, The Petroleum Institute, Abu Dhabi, 2533UAE,

Materials Science Laboratories, Toray Research Center, Inc., 3-3-7 Sonoyama, Otsu City, Shiga, 5208567, Japan, Department of Mechanical Engineering, Myongji University, 1st Engineering Bldg., 116

¶

Myongji-ro, Cheoin-gu, Yongin, Gyeonggi-do, 449-728, Korea,

1 ACS Paragon Plus Environment

ACS Applied Nano Materials 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 2 of 44

[email protected], [email protected], [email protected], [email protected], [email protected] and [email protected],

*To whom correspondence should be addressed, Tel: +81 3 5734 2423. Fax: +81 3 5374 2423. E-mail: [email protected]

(T.O).

Tel:

+971-2-607-5835.

Fax:

+971-2-607-5200.

E-mail:

[email protected] (I.I). ABSTRACT Graphene was modified with trimellitic anhydride groups and its Polyethylene terephthalate (PET)-based nanocomposites were prepared by melt-mixing. Percolation thresholds observed from changes to the electrical conductivity and storage moduli of nanocomposite melts suggest that the dispersion levels of unmodified graphene and those of modified graphene in PET matrix were the same. An enhancement of G’ and unexpectedly higher [] for modified graphene at low concentration suggest that that PET chains were grafted on the graphene surface creating a coupled network via covalent bonding. The bulk mechanical properties of amorphous nanocomposites were evaluated by tensile-testing. The nanocomposites with modified graphene also displayed an enhanced Young’s modulus as well as higher elongation compared to nanocomposites prepared with unmodified graphene. Differential scanning calorimetry, Fourier transform infrared spectroscopy and Raman spectroscopy results obtained on stretched nanocomposites suggest that both strain-induced orientation and strain-induced crystallization were suppressed by the modified graphene.


Page 3 of 44 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


KEYWORDS graphene, poly(ethylene terephthalate), trimellitic anhydride group, polymer-graphene nanocomposite, mechanical properties

1. Introduction Polymer nanocomposites are engineered by incorporating nanoscale fillers to modify properties relative to a pure polymer matrix. Two-dimensional platelet nano-fillers outperform spherical nanofillers in the enhancement of mechanical and gas-barrier properties as well as dimensional stability.1,2 Modified clay3-10 and graphene11-21 are widely used as platelet nano-fillers. Given its high modulus (25022,23-1000 GPa24) and electrical conductivity, graphene in particular has attracted considerable attention.14, 16,18,21,25-27. Poly(ethylene terephthalate) (PET) is a thermoplastic and semi-crystalline polymer with high performance characteristics such as: high glass transition temperature (Tg), good mechanical properties, high chemical resistance and ease of formation. Due to its high performance, PET is used for many industrial applications such as fibers, films, and bottles28. Biaxially-orientated PET films in particular are currently the most widely used polymeric substrate material in the fields of display technologies, electronic devices, storage media, electronic insulators, automobile, molding, package, printing and architecture29. PET resin is usually extruded into sheets followed by additional processing such as tentering and annealing above Tg 29-34. Additionally, products are often processed into the desired shapes via cutting, coating, laminating and molding by the end user, and most of these processes are conducted by automatic processing systems.　



Page 4 of 44

Several studies have reported on the enhancement of the properties of PET/graphene nanocomposites,

16-20.

Zhang et al.

16

melt-mixed PET with graphene and showed a greater

increase in electrical conductivity than with graphite. Feng et al. 17 reported that PET/graphene nanocomposites prepared by in-situ polymerization showed an increase of four orders of magnitude in the electrical conductivity compared to melt-blending. Bandla et al 18 and Li et al 19 enhanced the mechanical properties of PET by the addition of graphene. Most studies only focus on the enhancement of properties with incorporation of large amount of nano-fillers into PET. However, in most cases, at least one of the important mechanical property of PET such as elongation at break was decreased drastically with an increase in nano-filler content. Even that these disadvantages restrict the scope of industrial applications of such modified PET-based nanocomposites and cause many production and quality control issues such as cracking and dust contamination, most of previous studies for PET-based nanocomposites don’t show the elongation results. Shim et al

35

reported the

grafting of alkyl chains on graphene oxide (GO) and subsequent improvement of Young’s modulus, tensile strength and elongation at break of PET-based nanocomposites. However, these nanocomposites were formed by solution processing because of the lower thermal stability of GO36, rendering such systems impractical for melt processing applications. The purpose of this work is to improve the properties of melt-mixed PET/graphene nanocomposites and suppress their brittleness. In this work, we synthesized surface-modified graphene with trimellitic anhydride group and melt-mixed into PET. Transmission electron microscopy (TEM) and melt rheology were used to investigate the internal structure (i.e. dispersion level and interface 4 ACS Paragon Plus Environment

Page 5 of 44 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


structure) of the nanocomposite samples. Higher-order structures of stretched nanocomposites were investigated to clarify the mechanism of improvement of the mechanical properties by differential scanning calorimetry (DSC), Fourier-transform infrared spectroscopy (FTIR) and Raman spectroscopy.

2. Experimental Section 2.1. Materials PET pellets provided from Toray Plastics America (North Kingstown, RI) were ground into powder (30 mesh particle size) by Polyvision (Manchester, PA). The intrinsic viscosity of PET powder after grinding was 0.61 dl/g (in ortho-chlorophenol). Multi-layered graphene was used as nano-filler. Graphene G1 (XG Science, xGnP-C750, thickness 2 nm, diameter < 2 µm) was used as received. PET powder and graphene were dried in a vacuum oven at 120 °C for more than 12 hrs before meltblending. Anhydrous N,N-dimethyl formamide (DMF, EMD Chemicals Inc.) and anhydrous pyridine (Sigma Aldrich), 1,2,4-benzenetricarboxylic anhydride acid chloride (Trimellitic anhydride acid chloride, AK Scientific, Inc.) were used as received.

2.2. Synthesize of surface modified graphene The surface-modified graphene (modi-G1 from G1) was synthesized by the following procedures. Dried graphene (0.5 g) in a 2000mL flask was purged with dried nitrogen gas, then 500 mL of anhydrous DMF was added to the flask. Graphene was dispersed using a bath sonicator (Branson 2510, 100W) at room temperature for 1 hr and then stirred for 24 hrs (dispersion 1). 1,2,4Benzenetricarboxylic anhydride acid chloride (2.7 g) was taken in a 100 mL flask and purged with 5 ACS Paragon Plus Environment


Page 6 of 44

dried nitrogen gas, then 2.1g of anhydrous pyridine (used as catalyst) and 30 ml of anhydrous DMF were added and the mixture stirred until the solution became clear (solution 1). Solution 1 was transferred to dispersion 1 and the mixture was then stirred at 60 °C for 10 hrs. After reaction, the graphene dispersions were filtered through nylon membrane (0.20 μm; GNWP, Merk Millipore Ltd.) to collect the dispersed graphene. Filtered graphene was collected, re-dispersed in 500mL of DMF and then refiltered to remove unreacted material and byproducts. This step was repeated three times, after which the graphene cake was dried overnight in a vacuum oven at 90 °C.

2.3. Preparation of PET based nanocomposites The nanocomposites of PET with graphene were prepared by melt-blending. Two sets of mixtures were prepared: one of PET with unmodified graphene, while the other was PET with modified graphene. Each set of mixtures consisted of different loadings of graphene (from 0 to 12 wt% for G1 and modi-G1) and each of the resulting mixtures was fed in turn into a recirculating twin-screw extruder (Microcompounder, DACA Instruments) at 280 °C with N2 purge. Mixing was conducted at 360 rpm for 8 min and the composites were then extruded into an ice/water bath for cooling. The obtained nanocomposites were then dried in a vacuum oven at 120 °C for more than 12 hrs. Thin samples were prepared for TEM and mechanical property-testing by pressing the obtained nanocomposites to 120–180 μm thickness between fiber-reinforced Teflon sheets at 270 °C at 1–1.5 MPa for 2 min followed with quenching in an ice-water mixture to minimize crystallization.

2.4. Characterization Fourier-transform infrared (FT-IR) spectra of graphene and modified graphene were collected using a Magna IR-750 spectrometer (Nicolet Instrument, Madison WI) with the ATR method 38-42 6 ACS Paragon Plus Environment

Page 7 of 44 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


X-ray photoelectron spectra (XPS) of graphene and modified graphene were collected using a Surface Science model 555 spectrometer with MgKα radiation for investigation of physical and chemical states of graphitic materials38-42. After degassing in an introduction chamber (around 1 hr), samples were placed into a hemispherical analyzer where the pressure is maintained below 1.5 × 10-7 torr. An accelerating voltage of 12 kV at 250 W power was used for data acquisition of survey scans and high resolution scans. The peak analysis of the C1s and O1s in high resolution spectrum was conducted by using the Valdés’s method 40. Thermogravimetric Analysis (TGA) of the graphenes and modified graphenes was carried out using Perkin Elemer Pyris Diamond TG/DTA 6300. A sample (2-5 mg) was heated from a room temperature to 600 °C with a heating rate of 10 °C/min under N2 gas flow. Raman spectra of the graphene powder was measured by an Alpha 300R confocal Raman microscope equipped with a UHTS200 spectrometer and a DV401 CCD detector from WITec (Ulm, Germany). An air-cooled argon-ion laser operating at a wavelength of 514.5 nm was used as an excitation source with the output power varying between 2–20 mW. The Raman spectra were recorded with an integration time of 30–60 s at room temperature using a Nikon 100× air objective. TEM images of the polymer-graphene composites were obtained on a FEI Tecnai T12 microscope using an accelerating voltage of 120 kV. Thin sections (70nm) of rigid films were obtained using a microtome (Leica Ultracut) to obtain thin sections of the composites at room temperature with a diamond knife and then transferring the sections onto 400-mesh Ni grids. Rheological measurements were carried out using a strain-controlled rotational rheometer (ARES, TA Instruments) at 270 °C under N2 atmosphere. 0.6–0.7 g of nanocomposites samples were dried at 120 °C for at least 12 hrs then loaded between 25 mm parallel plates. The samples were squeezed into around1 mm thick disks by slowly lowering the upper plate. Using a dynamic strain sweep at 1 rad/s, the critical strain, γcrit, was recorded for which the storage modulus, G’, dropped to 90% of its limiting low strain value. Then, the dynamic moduli were measured at γ < γcrit from 100 rad/s to 0.1 rad/s. 7 ACS Paragon Plus Environment


Page 8 of 44

DSC measurements were performed using a TA Instruments Q1000 (TA Instruments, New Castle, DE). All measurements were carried out in an N2 atmosphere. The amorphous sheet samples were heated first to 280 °C at 10 °C/min and held at 280 °C for 1 min (first heating scan). For investigating the high order structure, heat of cold crystallization, ΔHcc, heat of fusion, ΔHm, were obtained from the first heating scan. The crystallinity, Xc, was calculated by equation (1)

Xc= (ΔHm - ΔHc)/ ΔHm0 × 1/(1-w) × 100

(1)

where ΔHm0 is the heat of fusion for 100% crystalline PET (140 J/g)43 and w is the weight fraction of the nano-fillers in the nanocomposites Measurements of mechanical properties were conducted on a Rheometrics Solid Analyzer, RSA-G2 (TA Instruments, New Castle, DE). The samples were cut to 2 mm wide dumbbells with 6 mm of effective length and loaded on the film fixtures. Strain-stress curve was collected at a stretching rate of 0.5 mm/s. Fourier transform infrared spectroscopy (FT-IR) was used to investigate the internal structure of stretched nanocomposites sheets. FT-IR spectra were measured for the stretched nanocomposite sheets using a Bruker ALPHA Spectrometer with the ATR method. XRD patterns were acquired using a Bruker-AXS (Madison, Wisconsin) microdiffractometer with Cu Ka radiation generated at 45 kV and 40 mA, a 0.8mmbeam collimator, and a 2D CCD detector. The samples were cut to small pieces and fixed on the stage and the scattering pattern was measured by transmittance method via both through -direction.


Page 9 of 44 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Raman spectra from stretched nanocomposites was measured by a T-64000 triple Raman spectrosmeter from Jobin Yvon (Bensheim, Germany) with 600 grooves/mm gratings and a slit width of 100 μm. A polarized air-cooled argon-ion laser operating at a wavelength of 514.5 nm was used as an excitation source with 5 mW of the output power. The Raman spectra by polarized incident light in the parallel direction to stretched direction and perpendicular direction are recorded at the same point at room temperature using a 100× air objective and rotating stage, respectively44.

3. Results and Discussion

3.1. Surface modification of graphene Graphene is hydrophobic in nature and adheres poorly to polyester 45. To improve interface adhesion between PET and graphene, surface modification of graphene was conducted as shown in scheme 1. It is generally believed that graphene contains various oxygen-containing surface groups such as C=O, C−OH, C-O-C, COOR, anhydride and lactone, COOH, chemisorbed water and oxygen

38-42.

Among these different types of oxygen-containing groups, three groups in particular provide viable reaction sites for chemical modification, namely: carboxylic acid, hydroxyl, and epoxy. Acid chlorides can react with the hydroxyl group and carboxyl group46,47. Therefore, trimellitic anhydride groups were grafted on the graphene surface by this reaction using pyridine as catalyst.



Page 10 of 44

Scheme 1. Synthetic procedure for functionalization of graphene

The FT-IR spectra of unmodified and modified graphenes are shown in Figure 1. Because of the lower content of chemically-bonded oxygen on the graphene, some specific peaks are hidden by the strong absorption of the graphite structure and carbonyl groups of graphene48. Nevertheless, modified graphene showed clear aromatic C-C stretching peak in the benzenoid region (wavenumber around 1470 cm-1 in Figure 1)49,50. Additionally, the peak around 1200 cm-1 was attributed to C-OH vibration stretching and the peak around 1580 cm-1 were attributed to the stretching of the C=C stretching in quinoid ring49. A salient qualitative observation is that the peak intensity for the C-OH vibration stretching is decreased compared to C=C stretching after the proposed reaction with acid chloride. As the quinoid ring does not react with acid chloride, these peak changes


Page 11 of 44 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


demonstrate that a functionalization reaction has taken place on the surface resulting in the reaction of OH groups and their subsequent replacement with aromatic benzene rings tethered to the graphene surface by esterification reaction.

Figure 1. FT-IR spectra of G1and modi-G1,. Raman spectra of unmodified and modified graphenes are shown in Figure S-1. Both samples show three prominent peaks: the D, G and 2D bands. Some specific peaks are hidden by the strong peak of the graphite structure and carbonyl groups of graphene due to the lower content of chemically-bonded oxygen on the graphene. The 2D peak is related with the number of the graphene layer.51 The G/2D intensity ratio IG/I2D for them are 2.6 and 2.7 respectively and peak shape and position of 2D shows no significant difference for G1 and modi-G1 (Figure S-1(b)). These results 11 ACS Paragon Plus Environment


Page 12 of 44

suggest that layer number of graphene is similar after reaction. 51 Additionally, the G band of modi-G1 is downshifted slightly (~1 cm-1) compared with G1 (Figure S-1(c)). The change in the G band position reflects the electrical condition of the graphene and can be used to indicate that charge transfer is taking place between graphene and any surface functional groups and/or additives, with electron donor groups causing a downshift and electron accepter groups causing an upshift in said peak.52 This downshift in the G band suggests that electrical density of graphene is increased after reaction which is consistent with functionalization taking place by the electron-donor trimellitic anhydride groups. XPS was also used for corroborative functional group identification. Survey scans and high resolution scan of C1s and O1s for the unmodified and modified graphenes are shown in Figure 2 and Figure S-2, with the analysis results of the O1s peak shown in Table S-1.

Figure 2. XPS high resolution scans of O1s of (a) G1 and (b) modi-G1.

Both graphenes consist of only carbon and oxygen (Figure S-2(a)). Peaks of N1s and Cl2s are not visible in both G1 and modi-G1. This result indicates that the modified graphene does not include the chemical functional modifier (1, 2, 4-Benzenetricarboxylic anhydride acid chloride) and catalyst 12 ACS Paragon Plus Environment

Page 13 of 44 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


(pyridine). Although the difference in the high resolution scan of C1s between unmodified and modified graphene is not clear (Figure S-2(b), (c)), the peak shape of high resolution scans of O1s were distinctly different between the two graphenes (Figure 2). The O1s peak of the carbon-containing material was deconvoluted into five components: C=O (531.1 eV), C−OH, C-O-C (532.3eV), COOR, anhydride and lactone (533.3 eV), COOH (534.2 eV), chemisorbed water and oxygen (536.1eV)

35-37.

G1 has a main peak (C−OH, C-O-C) at 532.3eV (Figure

2(a)), whereas the main peak was shifted in favor of COOR, anhydride and lactone after reaction. The intensity of the peak (C-OH, C-O-C) was reduced through the reaction, as seen in Figure 2 and Table S-1. This change in a major peak suggests that the acid chloride groups on the chemical functional modifier selectively reacted by alkoxy-de-halogenation with the hydroxyl groups on the graphene and formed ester bonds46,47,53, resulting in grafting of trimellitic anhydride groups on the graphene surface. The TGA curves of graphene and modified graphene are shown in Figure S-3. Around the meltprocessing temperature for PET (at 270°C), the recorded weight loss for both G1 and modi-G1 was around 5 wt%, which is a much lower value compared to graphene oxide (around 35 wt%) and functionalized graphene oxide (around 15 wt%) as reported in Shim’s work32. Therefore, the higher heat stability of modi-G1 compared with that reported for the functionalized 13 ACS Paragon Plus Environment


Page 14 of 44

graphene oxide in Shim’s work32 means that our material is better suited for melt-processing. The surface content of the trimellitic anhydride groups on the modified graphene was determined by the difference of weight residue between the G1 and modi-G1 at 600°C. In Figure S-3, modi-G1 showed 4% larger weight loss compared to G1, indicating that the organic component on the graphene surface increases after reaction. According to Scheme 1, the concentration of trimellitic anhydride groups in modi-G1 was estimated to be around 1 group per 340 carbon atoms of the parent graphene (G1).

3.2. Dispersibility and structure of PET-based nanocomposites. PET/unmodified graphene and PET/modified graphene nanocomposites were prepared by melt-mixing, melt pressing and quenching. In order to study the crystallinity of nanocomposites, representative DSC profiles for the first heating scan are shown in Figure S-4 and S-5. During the first heating scan from room temperature, all samples showed stepwise endothermic changes for glass transition, exothermic peaks for cold crystallization and endothermic peaks for crystal melting. Peaks corresponding to cold crystallization for all nanocomposites have been shifted to lower temperatures than for the neat-PET, Furthermore the peak shift for PET/modi-G1 is 1°C smaller than PET/G1. This suggest that graphene works as a nucleation agent and the nucleation effect of G1 was stronger than that for modi-G1. Despite the difference in cold crystallization behavior across G1 and modi-G1 composites, all samples displayed less than 5% crystallinity (i.e. nearly amorphous state) as shown in Figure S-5. 14 ACS Paragon Plus Environment

Page 15 of 44 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


XRD characterization results – both spectral and two-dimensional - of PET and PET-based nanocomposites are shown in Figure S-6 and S-7. All samples analyzed using 2D XRD showed the isotropic scattering feature amorphous halo of PET with no visible peaks specific to crystals of PET; while any diffraction patterns from the graphene were hidden by the amorphous halo of PET. These results also indicate that all samples are in the amorphous state without preferred orientation54. TEM micrographs were collected to understand the dispersion of nano-fillers in PET-based nanocomposites. Electrical resistance measurements were also conducted to estimate the level of dispersion of the two graphenes in the PET matrix. The TEM images of PET/graphene nanocomposites and the surface resistances of PET/graphene are shown Figure 3 and Figure S-8. The dispersion levels of PET/G1 and PET/modi-G1 seem to be the same, and some aggregates were observed in both images (Figure 3(a) (b) and Figure S-8).



Page 16 of 44

Figure 3. TEM images of (a) PET/G1 and (b) PET/modi-G1 nanocomposites with 2 wt% of nanofillers, (c) Electrical resistance of PET/G1 and PET/modi-G1 nanocomposites.

Both G1 and modi-G1 reduced the surface resistance with graphene concentration as expected. The onset concentration for electrical percolation of both G1 and modi-G1 was around 4 wt% (Figure3(c)), indicating that there is no significant difference in the level of dispersion of G1 versus modi-G1. The percolation threshold

45,

Φper, for randomly-oriented disks is inversely proportional to Af, which is the

ratio of disc diameter D to disc thickness, h. The particle average aspect ratio, Af, was estimated by equation (2) 45

Af = D/ h = 3Φsphere/2Φper

(2)


Page 17 of 44 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


where Φspher = 0.29, which is the onset of percolation of interpenetrating and randomly-packed spheres 56. For the weight to volume fraction conversion, the densities were assumed to be 2.28 g/cm3 for graphene

14,15,20

and 1.335 g/cm3 for amorphous PET20,57. The average particle aspect ratio, Af,

estimated by equation (2) is 20 for both graphenes. The lack of change in dispersion quality is because the reaction under investigation is mild. For starters the number of OH and COOH groups on the surface of G1 are expected to be much lower than for their graphene oxide counterpart, meaning that the reaction with acid chloride groups is only possible at a select number of locations on the surface. Moreover, it is expected to be difficult for the bulky trimellitic anhydride molecules to insert themselves in between graphene layer in G1 , especially given that trimellitic anhydride molecules contain aromatic rings with high electron density which would be repelled by conjugated - domains on the surface of graphene layer and thus prevent the anhydride molecules from intercalating properly between stacks. This is supported by the Raman data, for which 2D band is not changed after reaction in Figure S-1. Thus, the reaction has limited power to exfoliate the G1 and there is not a noticeable improvement in the exfoliation meaning that the dispersion levels of graphene in the resulting composites is almost the same for modi-G1 versus G1.

3.3. Rheological properties of PET based nanocomposites. Melt rheology measurements were conducted to estimate internal structures such as dispersion15,58 and the molecular network51 of PET based nanocomposites. Dynamic strain sweeps of PET/graphene melts at 1 rad/s and dynamic frequency sweeps are shown in Figure S-9. All nanocomposites exhibit enhanced melt elasticity with increasing filler-loading. This behavior is similar to that of other polymer/graphene nanocomposites systems15,59,50. In particular, the storage modulus, G’, of PET/modified G1 composite showed a greater increase and at a lower concentration than for unmodified graphene (G1). Both samples yield at large strain, as characterized by a decrease in G’ with 17 ACS Paragon Plus Environment


Page 18 of 44

an increased strain from the critical strain, γcrit. γcrit decreases at a higher concentration of graphene (Figure S-9(a),(b)), which indicates formation of more fragile filler networks15. γcrit values were plotted against filler volume fraction,  the results of which are shown in Figure S-10. For the weight to volume fraction conversion, the densities were assumed to be 2.28 g/cm3 for graphene14,15,20 and 1.335 g/cm3 for amorphous PET20,57. The decrease in γcrit becomes more pronounced once the concentration of inclusions reaches a value of around 0.025 (4 wt%,) to 0.049 (8 wt%) for both G1 and modi-G1. This observation implies that percolation of a graphene network occurs at these two volume fractions for G1 and modi-G1. Both dispersion uniformity and fragility of graphene networks of G1 and modi-G1 are thus similar from the perspective of melt rheology. In terms of dynamic frequency sweeps of PET/graphene melts (Figure S-9(c)(d)), G’ of PET/G1 at low frequency increases with graphene concentration and becomes independent of the frequency above a critical concentration. This is a typical solid-like response of nanocomposite melts with layered fillers 15, 45, 58, 61-64

and it indicates the formation of a network of filler–filler contacts. On the other hand, the

increase in G’ of PET/modi-G1 at low frequency was slightly larger than for PET/G1. G’ at ω = 0.1 rad/s for PET/graphene are plotted against filler concentration and the percolation thresholds, Φper, are obtained by linear regression, the results are shown in Figure 4.


Page 19 of 44 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Figure 4. G’ at ω = 0.1 rad/s of PET/G1 and PET/modi-G1 melts at 270 °C.

The drastic increase in G’ at low frequency occurs between 0.025 (4 wt%) and 0.049 (8 wt%) for both G1 and modi-G1, indicating that the dispersion levels of PET/G1 and PET/modi-G1 are the same. From a previous study, the particle aspect ratios, Af, which is a ratio of diameter D to thickness h, in the nanocomposites was estimated to be 14 for PET/G1 from rheology of the molten dispersions23. Therefore, Af of PET/modi-G1 was also estimated to be 14. On the other hand, G’ of PET/modi-G1 were higher than those of PET/G1 under Φper, In order to investigate the difference between the network structures of the G1- and modi-G1 composites, the complex viscosity * was calculated from G’ and G’’ data for the two sets of samples (Figure S-11).



Page 20 of 44

The  data of nanocomposites with 0-4wt% of graphene (i.e. under Φper) was then used to calculate intrinsic viscosity [] for G1 and modi-G1. The results are shown in Figure 5.

Figure 5. [η] of PET/G1and PET/modi-G1 melts at 270 °C.

The aspect ratio for G1 can be estimated using the following Kuhn-Kuhn relation for intrinsic viscosity of oblate spheroids and circular discs65:

[η] = 32/15π

(3)

where, is the mean aspect ratio. Using equation (3), the aspect ratio for G1 is estimated to be 24 which is consistent with the estimate of 20 from electrical conductivity, as well as the estimate of 14 20 ACS Paragon Plus Environment

Page 21 of 44 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


obtained from the percolation threshold equation as mentioned above. It can be shown that the agreement between the intrinsic viscosity estimate of the aspect ratio and the other two estimates is even closer by iteratively using the following generalized Kuhn-Kuhn relation66:

[η] = (1 + 3 + 3) 32Af /15π (1 + )

(4)

where  is the normalized standard deviation of aspect ratio. Equation (4) was used to estimate the standard deviation  of the G1 aspect ratio with an initial guess of 20 (based on the electrical conductivity estimate), yielding    thus suggesting that for G1, Af ≈ 20 ± 6, thus further indicating that the three estimates of Af are indeed consistent. On the other hand, [η] of PET/modi-G1 is around one-order of magnitude higher than PET/G1 even though the aspect ratios of both nanocomposites are the same. The [] for the modi-G1 dispersion yields an aspect ratio of around 240. However this value is physically unrealistic, based on the TEM images as well as estimated aspect ratios from both the electrical and rheological percolation thresholds for modiG1. These percolation thresholds are very similar to those for G1 (20 from electrical and 14 from rheological respectively). Thus, it can be assumed that the aspect ratio did not change significantly for the modi-G1 sheets following the modification, and the high [] for modi-G1 can thus only be interpreted as being the result of entanglement coupling between modi-G1 and the surrounding PET polymer matrix post-blending. Modi-G1 has trimellitic anhydride groups on the surface which can react with the hydroxyl end groups and carboy groups of PET by ester formation during the melt processing (see Supporting Information, Scheme S1, reaction I and II)47. The elasticity of coupled PET melts shows higher G’ than for linear PET

61.

Therefore, the enhancement of G’ and unexpectedly higher [] for

modified graphene at lower concentration both suggest that PET chains were grafted on the graphene surface, creating a coupled network via covalent bonding (Scheme S-2). 21 ACS Paragon Plus Environment


Page 22 of 44

3.4. Mechanical properties of PET-based nanocomposites. Representative strain-stress curves for PET/G1 and PET/modi-G1 at room temperature are shown in Figure S-12, in which the nanocomposite samples appear to have undergone yielding followed by a plastic deformation region that ended with strain hardening and ultimately the failure of the samples. Although G1 and modi-G1 do not display any discernible differences in dispersion level or aspect ratio in the PET composites, each set exhibited different behavior; while the draw stresses in the plastic deformation zone increased with an increase in strain for both samples sets, those of the PET/modi-G1 composites increased as a function of increase in the loading concentration of graphene, whereas no such concentration dependence was observed for PET/G1. Additionally, the elongation of PET/modi-G1 was higher than for PET/G1 at any concentration of graphene loading. The Young’s modulus and elongation at break of these samples are shown in Figure S-13. All nanocomposites exhibited higher Young’s modulus than neat PET (Figure S-13 (a)). Improvements of Young’s modulus were 4% for G1, 12% for modi-G1 with 4 wt% of graphene. PET/modi-G1 graphene showed about 5% higher Young’s modulus than PET/G1. Although elongation at break of all nanocomposites decreased with an increase of graphene content, those of PET/modi-G1 showed higher values than that of PET/G1 with concentrations from 1wt% to 4.0wt% (Figure S-13(b)). The relationship between Young’s modulus and elongation at break are plotted and shown in Figure 6.


Page 23 of 44 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Figure 6. Relationship between Young’s modulus and elongation at break of PET/G1 and PET/modiG1. PET/G1 first exhibited a drastic decrease in elongation at break, followed by a slight enhancement of the Young’s modulus. On the other hand, modi-G1 was found to suppress the decrease in elongation at break in PET/modi-G1. Thus, it may be concluded that modi-G1 improves the Young modulus and suppresses the decrease of elongation. According to previous research, these enhancement of modulus and draw stress in PET/modi-G1 indicates that strong interfacial adhesion by covalent bonding leads to effective stress transfer from PET matrix to graphene.67, 68

3.5. Structure of stretched PET/graphene nanocomposites 23 ACS Paragon Plus Environment


Page 24 of 44

In order to clarify the specific effects of modified graphenes, higher-order structures of stretched samples were investigated. 2D XRD data and XRD pattern data of 200% stretched samples are shown in Figure S-14 and S-15. All stretched samples showed the a pair of broad and intense arcs on the equator (Figure S-9 and S-10) and additional very weak and broad arc at around 16° and distinct peak at 2θ = 26.1° on the meridian (Figure S-14 and S-15(b)). Similar XRD pattern change with respected to stretching also observed stretched samples for various materials such as PET54,55 and elastomers62. According to previous research, these diffraction pattern changes indicates that crystalline phase and mesophase (i.e.: semicrystalline phase) with a middle degree of packing order between the crystalline phase and the amorphous phase are created and orientated during deformation54,55,62. Because of the very weaker diffraction intensity by XRD, the crystallinity of 200% stretched nanocomposites was investigated by DSC. Representative DSC profiles for the first heating scan and crystallinity of stretched samples are shown in Figure S-16 and Figure 7.


Page 25 of 44 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Figure 7. Crystallinity of 200% stretched nanocomposites. Samples were stretched at room temperature.

All samples showed, very small and wider exothermic peaks for cold crystallization and endothermic peaks for crystal melting. Despite there being no significant differences in the cold crystallization peaks, PET/G1 nanocomposites showed more change in the endothermic peaks than PET/modi-G1, indicating crystallinity of PET in PET/G1 nanocomposites more increased than PET/modi-G1 The crystallinity of all nanocomposite samples before deformation was less than 5% (See Figure S-5 in Supporting Information), whereas it increased up to 22% after deformation. The crystallinity of stretched PET/G1 was more than 2-4% higher than stretched neat PET and increased with graphene loading. Similar phenomena have been observed in other nanocomposite systems such as elastomer/layered double hydroxides69 and elastomer/clay. The large interfacial area between polymer 25 ACS Paragon Plus Environment


Page 26 of 44

and graphene could have led to an early orientation of polymer chains, thus promoting nucleation under stretching, and this might have led to acceleration of strain-induced crystallization63. On the other hand, the crystallinity of stretched PET/modi-G1 with 1% of graphene was slightly decreased compared to the neat one. Moreover, it was increased with an increase of graphene concentration, although the absolute values were still lower than those for PET/G1. These observations suggest that modi-G1 suppresses strain-induced crystallization. Fourier transform infrared spectroscopy (FT-IR) was used for the investigation of the deformation behavior of PET

71-73.

The FT-IR spectra of unstretched and 200%

stretched nanocomposites are shown in Figure S-17. The ethylene glycol unit in PET has two rotational conformations, trans-conformations and gauche-conformations. Each conformation has specific bands at 1340 and 1370 cm-1 respectively, which are assigned to CH2 wagging modes of the glycol segment 7173.

In Figure S-17, unstretched nanocomposites have clear peaks of the trans- and gauche-conformations.

The peak intensity for the gauche-conformations is greater than that for the trans-conformations. After 200% stretching, the peak intensity of the gauche-conformations is reduced, whereas that of the transconformations is increased. The interchange of these conformations is influenced by both the crystallinity and amorphous structure. The remaining gauche-conformations indicate that there is still some room for stretching the chain molecules There are both trans-conformations and gaucheconformations in the amorphous phase, whereas only trans-conformations exist in the crystalline phase 71.

Thus, the content of trans-conformations in crystal X trans, a, can be considered to be the same as the

crystallinity, Xc

74.

The content of total conformations, Xt, trans-conformations in crystal X

trans, c,

in

amorphous, X trans, a, and gauche conformations in amorphous, X gauche, a, are calculated by equations (5) (7).

X trans, c= Xc

(5)

X t (%), =X trans, a + Xtrans, c = I1340 / (I1340 + I1370) × 100

(6) 26

ACS Paragon Plus Environment

Page 27 of 44 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


X gauche, a = I1370 / (I1340 + I1370)

× 100

(7)

where I1340 and I1370 are intensity of peak at 1340 cm-1 , and 1370 cm-1 by FT-IR, respectively. The higher order structures of 200% stretched nanocomposites are shown in Figure 8. Although the total number of trans-conformations of PET/G1 is similar to neat PET, the trans-consumer in the crystal phase was increased due to the increase in crystallinity by the addition of graphene (Figure 7). On the contrary, the total number of trans-conformations of PET/modi-G1 was decreased slightly at 1 wt% of modi-G1 and then became similar with neat PET and PET/G1 at 2 wt%. However, the content of transconformations in the crystal phase in PET/modi-G1was still lower than that for PET/G1.

Figure 8. High order structures of 200% stretched nanocomposites.



Page 28 of 44

Additionally, raman spectroscopy was used to search for observable orientation of the PET in the 200% stretched nanocomposites. Representative raman spectra of amorphous PET and PET/graphene nanocomposites before and after 200% stretching are shown in Figure S-18. Before stretching, all the samples showed identical spectra for both of the orthnormal incident light directions to the stetching axis of the stretched nanocomposites. On the other hand, after stretching, the peak intensity is greater when light is shone in the direction parallel to the stretching axis for the composites compared with the peak intensity for light shone in the perpendicular direction. This suggests that the PET chains in the modi-G1 nanocomposites are oriented as a result of stretching. For this experiment, the ring-form C-C stretch (wavenumber ~1615 cm-1) of PET were selected for assessment because its intensity in the raman spectra is particularly sensitive to the polarization direction of incident light

75.

The intensity ratio for the two light

orientations (I1615// / I1615 ⊥ ) is correlated to the orientation degree of PET, and is used as the orientation parameter

44,75,76.

An enlarged view of the representative raman spectra of the ring-

form C-C stretch and the orientation parameter I1615// / I1615 ⊥ of 200% stretched nanocomposites are shown in Figure S-19 and 9.


Page 29 of 44 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Figure 9. Orientation parameter I1615// / I1615

⊥

of 200% stretched nanocomposites. Samples were

stretched at room temperature.

Moreover, the orientation paramater of 200% strethed PET/G1is observed to decrease with an increase in graphene content. This suggests that graphenes cause a disruption in the orientationof the PET chains surrounding them in amorphous domanins77. On the other hand, stretched PET/modi-G1 composites exhibited a larger decrease in the orientation parameter than PET/G1 for the same graphene loadings. The same trend for PET/modi-G1 composites was observed as for PET/G1 composites in this regrd, i.e.: the orientation parameter decreased with an increase in the graphene concentration, although the absolute values were still lower than those for PET/G1. These



Page 30 of 44

results thus suggest that both the strain-induced orientation and strain-induced crystallization of PET were suppressed by modified graphene.

3.6. Mechanism of improvement of mechanical properties Based on these results, a mechanism of improvement of mechanical properties is suggested and described below. During melt processing, it is thought that PET chains were grafted on the graphene surface by the reaction between trimellitic anhydride groups on the modified graphene and hydroxyl end groups of PET chains, thus enhancing the wettability of graphene to PET. The entanglement of grafted PET chains with matrix PET is thought to have increased interfacial adhesion between PET and graphene, thus leading to the formation of a coupled network. Higher interfacial adhesion prevents destruction of interface and enhances the stress-transfer efficiency between matrix polymer and nanofillers

66,67,78,

and this is thought to be the reason behind the increase in Young’s modulus and draw

stress of PET/modified graphene compared to PET/unmodified graphene. Because coupled points are unable to flow along the applied stress direction, the degree of orientation and crystallization of branched PET was decreased

79-81.

Therefore, both strain-induced orientation and crystallization of

PET/modified graphene were suppressed. Breakage of PET material was caused in the amorphous region by the breakdown of entanglement network in the amorphous region according to molecular conformational changes from gauche-conformations to trans-conformations

82-83.

The restriction on

molecular conformation changes in PET/modi-G1 during deformation results in the improvement of elongation at break of the nanocomposites.

4. Conclusions


Page 31 of 44 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Graphene with trimellitic anhydride group was synthesized and its PET-based nanocomposites were prepared by melt-mixing. Internal structures such as dispersion and molecular networks were investigated by melt rheology. From percolation thresholds, the dispersion levels of unmodified graphene and those of modified graphene in PET matrix were observed to be the same. On the other hand, the increase in storage modulus of nanocomposite melts at lower concentrations, up to 0.049 (8 wt%) at lower frequency, suggests the formation of a coupled network between PET and graphene during melt-mixing. The formation of a coupled network between PET and graphene is thought to play an important role in the improvement of mechanical properties of PET/graphene nanocomposites. PET/modified graphene showed about 5% higher Young’s modulus than PET/unmodified graphene. Furthermore, the decrease in elongation at break was suppressed more by the modified graphene than by unmodified graphene. DSC, FT-IR and Raman spectroscopy results of stretched nanocomposites suggest that both straininduced orientation and strain-induced crystallization were suppressed by modified graphene. Therefore, stronger interface adhesion and restriction of strain-induced orientation and crystallization by modified graphene are thought to have beneficial effects on the enhancement of mechanical properties of PET/graphene nanocomposites. This is the first reported case of improving the brittleness of melt-mixed PET /graphene nanocomposites.

ACKNOWLEDGMENTS This research was supported by Toray Industries through their membership in the University of Minnesota Industrial Partnership for Research in Interfacial and Materials Engineering (IPRIME). Parts of this work were performed in the University of Minnesota Characterization Facility, which receives partial support from the NSF through the MRSEC program.

SUPPORTING INFORMATION 31 ACS Paragon Plus Environment


Page 32 of 44

Reaction scheme between trimellitic anhydride group on graphene and PET, raman spectrrum, XPS survey and high resolution scan of C1s, and thermogravimetric analysis of graphene, rheology of nanomomposites melt. TEM image, crystallinity and tensile properties of amorphous PET based nanocomposites sheets, XRD , DSC, FT-IR spectrum and raman spectrum of un-stretched and 200% stretched PET based nanocomposites. This material is available free of charge via the Internet at http://pubs.acs.org.

REFERENCES (1) Brune, D. A.; Bicerano, J. Micromechanics of Nanocomposites: Comparison of Tensile and Compressive Elastic Moduli, and Prediction of Effects of Incomplete Exfoliation and Imperfect Alignment on Modulus. Polymer 2002, 43, 369–387. (2) Choi W.; Kim H.; Yoon, K.; Kwon, O.; Hwang, C. Preparation and Barrier Property of Poly(ethylene terephthalate)/Clay Nanocomposite using Clay-supported Catalyst. J. Appl. Polym. Sci. 2006, 100, 4875–4879. (3) Hasegawa, N.; Okamoto, H.; Kato, M.; Usuki, A. Preparation and Mechanical Properties of Polypropylene–Clay Hybrids Based on Modified Polypropylene and Organophilic Clay. J. Appl. Polym. Sci. 2000, 78, 1918-1922. (4) Pettarin, V.; Frontini, P. M.; Rodriguez Pita, V. J. R.; Lopes Dias, M.; Valenzuela Diaz, F. Polyethylene/(Organo-montmorillonite)

Composites

Modified

with

Ethylene/methacrylic

Acid

Copolymer: Morphology and Mechanical Properties. Composites Part A 2008, 39, 1822–1828. (5) Tsai, T. Y.; Li, C. H.; Chang, C. H.; Cheng, W. H.; Hwang, C. L.; Wu, R. J. Preparation of 32 ACS Paragon Plus Environment

Page 33 of 44 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Exfoliated Polyester/Clay Nanocomposites. Adv. Mater. 2005, 17, 1769–1773. (6) Sharma, R.; Joshi, H.; Jain, P. Short Review on the Crystallization Behavior of PET/Clay Nanocomposites. J. Chem. Eng. and Mater. Sci. 2011, 2, 39-43. (7) Ou, C.; Ho, M.; Lin, J. Synthesis and Characterization of Poly(ethylene terephthalate) Nanocomposites with Organoclay. J. Appl. Polym. Sci. 2004, 91, 140–145. (8) Vermogen, A.; Masenelli-Varlot, K.; Séguéla, R. Evaluation of the Structure and Dispersion in Polymer-Layered Silicate Nanocomposites. Macromolecules 2005, 38, 9661–9669. (9) Pegoretti, A.; Kolarik, J.; Peroni, C.; Migliaresi, C. Recycled Poly (ethylene terephthalate)/Layered Silicate Nanocomposites: Morphology and Tensile Mechanical Properties. Polymer 2004, 45, 2751– 2759. (10) Gökkurt, T.; Durmus, A.; Sariboga, V.; Öksüzömer, M. A. F. Investigation of Thermal, Rheological, and Physical Properties of Amorphous Poly(ethylene terephthalate)/Organoclay Nanocomposite Films. J. Appl. Polym. Sci. 2013, 129, 2490-2501. (11) Geim, A. K.; Novoselov, K. S. The Rise of Graphene. Nature Mater. 2007, 6, 183–191. (12) Novoselov, K. S.; Geim, A. K.; Morozov, S. V.; Jiang, D.; Katsnelson, M. I.; Grigorieva, I. V.; Dubonos, S. V.; Firsov, A. A. Two-dimensional Gas of Massless Dirac Fermions in Graphene. Nature 2005, 438, 197–200. (13) Zhang, Y.; Tan, Y. W.; Stormer, H. L.; Kim, P. Experimental Observation of the Quantum Hall Effect and Berry's Phase in Graphene. Nature 2005, 438, 201–204. (14) Kim, H.; Abdala, A. A.; Macosko, C. W. Graphene/Polymer Nanocomposites. Macromolecules 2010, 43, 6515–6530. 33 ACS Paragon Plus Environment


Page 34 of 44

(15) Kim, H.; Macosko, C. W. Processing-property Relationships of Polycarbonate/Graphene Composites. Polymer 2009, 50, 3797–3809. (16) Zhang, H.; Zheng, W.; Yan, Q.; Yang, Y.; Wang, J.; Lu, Z.; Ji, G.; Yu, Z. Electrically Conductive Polyethylene Terephthalate/Graphene Nanocomposites Prepared by Melt Compounding. Polymer 2010, 51, 1191–1196. (17) Feng, R.; Guan, G.; Zhou, W.; Li, C.; Zhang, D.; Xiao, Y. In situ Synthesis of Poly(ethylene terephthalate)/Graphene Composites using a Catalyst Supported on Graphite Oxide. J. Mater. Chem. 2011, 21, 3931–3939. (18) Bandla, S.; Hanan, J. C. Microstructure and Elastic Tensile Behavior of Polyethylene terephthalateExfoliated Graphene Nanocomposites. J. Mater. Sci. 2012, 47, 876–882. (19) Li, M.; Jeong, Y. G. Poly (ethylene terephthalate)/Exfoliated Graphite Nanocomposites with Improved Thermal Stability, Mechanical and Electrical Properties. Composites: Part A, 2011; 42: 560– 566. (20) Aoyama, S.; Park, Y.T.; Ougizawa, T.; Macosko, C.W. Melt Crystallization of Poly

(ethylene terephthalate): Comparing Addition of Graphene vs. Carbon Nanotubes. Polymer 2014, 55, 2077-2085. (21) Li, M.; Jeong, Y. G. Influences of Exfoliated Graphite on Structures, Thermal Stability, Mechanical Modulus, and Electrical Resistivity of Poly(butylene terephthalate). J. Appl. Polym. Sci. 2012, 125, E532–E540. (22) Gómez-Navarro, C.; Burghard, M.; Kern, K. Elastic Properties of Chemically Derived Single Graphene Sheets. Nano Letters 2008, 8, 2045–2049.


Page 35 of 44 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


(23) Suk, J. W.; Piner, R. D.; An, J.; Ruoff, R. S. Mechanical Properties of Monolayer Graphene Oxide. ACS Nano 2010, 4, 6557–6564. (24) Lee, C.; Wei, X.; Kysar, J. W.; Hone, J. Measurement of the Elastic Properties and Intrinsic Strength of Monolayer Graphene. Science 2008, 321, 385–388. (25) Rahman, M. M.; Hussein, M.A.; Alamry, K. A,; Al-Shehry, F. M., and Asiri. A. M.; Polyaniline/Graphene/Carbon Nanotubes Nanocomposites for Sensing Environmentally Hazardous 4Aminophenol. Nano-Structures & Nano-Objects, 2018, 15 , 63-74. (26) William, C.; Groden, N.; Wilson. C. G. Graphene Nanoscrolls Fabricated by Ultrasonication of Electrochemically Exfoliated Graphene. Nano-Structures & Nano-Objects, 2017, 12 ,77-83. (27) Sangeeta, M.; Bajpai, A. K.; Bajpai. J. Poly (vinyl alcohol) Supported Thermally Reduced Graphene Oxide (TRGO) Nanosheets Exhibit Enhanced Electrical and Mechanical Behavior. NanoStructures & Nano-Objects, 2018, 14, 73-83. (28) Gupta, V. B.; Bashir, Z., Handbook of Thermoplastic Polyesters: Homopolymers, Copolymers, Blends, and Composites 2005,362-387. (29) Tsunekawa, T.; Higashioji, T; Hosokawa, H.; Kubota, A.; Ishizuka, I. Development and Industrialization of a Biaxially Stretched Nano-alloyed Film. Polymer Journal 2012, 44, 1170–1178. (30) Elenga, R.; Seguela, R.; Rietsch, F. Thermal and Mechanical Behaviour of Crystalline Poly (ethylene terephthalate): Effects of High Temperature Annealing and Tensile Drawing. Polymer 1991, 32, 1975-1982. (31) Salem, D. R. Orientation and Crystallization in Poly(ethylene terephthalate) during Drawing at High Temperatures and Strain Rates. Polym. Eng. Sci. 1999, 39, 2419-2430.



Page 36 of 44

(32) Lu, X. F.; Hay, J. Crystallization Orientation and Relaxation in Uniaxially Drawn Poly(ethylene terephthalate). Polymer 2001, 43, 8055–8067. (33) Mascia, L.; Fekkai, Z. Influence of Stress-induced Crystallization on the Dimensional Stability of Monoaxially Drawn Poly (ethylene terephthalate). Polymer 1993, 34, 1418-1422. (34) Salem, D. R. Crystallization During Hot-drawing of Poly (ethylene terephthalate) Film: Influence of the Deformation Mode. Polymer 1995, 36, 3605-3608. (35) Shim, S. H.; Kim, K. T.; Lee, J. U.; Jo, W. H. Facile Method to Functionalize Graphene Oxide and Its Application to Poly (ethylene terephthalate)/Graphene Composite. ACS Appl.

Mater. Interfaces 2012, 4, 4184−4191.

(36) Jeong, H. K.; Lee, Y. P.; Jin, M. H.; Kim, E. S.; Bae, J. J.; Lee, Y. H. Thermal Stability of Graphite oxide. Chem. Phys. Lett. 2009, 470, 255–258.

(37) Staudenmaier L. Method for the Preparation of Graphitic Acid. Ber. Dtsch. Chem. Ges. 1898, 31, 1481-1487.

(38) Zielke, U.; Hüttinger, K. J.; Hoffman, W. P. Surface-oxidized Carbon Fibers: I. Surface Structure and Chemistry. Carbon 1996, 34, 983-998. (39) Figueiredo, J.L.; Pereira, M.F.R.; Freitas, M.M.A.; Órfão. J.J.M. Modification of the Surface Chemistry of Activated Carbons. Carbon 1999, 37, 1379-1389.


Page 37 of 44 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


(40) Valdés H.; Sánchez-Polo, M.; Rivera-Utrilla, J.; Zaror, C. A. Effect of Ozone Treatment on Surface Properties of Activated Carbon. Langmuir 2002, 18, 2111-2116. (41) Stankovich, S.; Dikin, D. A.; Piner, R. D.; Kohlhaas, K. A.; Kleinhammes, A.; Jia, Y.; Wu, Y.; Nguyen, S. T.; Ruoff, R. S. Synthesis of Graphene-based Nanosheets via Chemical Reduction of Exfoliated Graphite Oxide. Carbon 2007, 45, 1558-1565. (42) Yang, D.; Velamakanni, A.; Bozoklu, G.; Park, S.; Stoller, M.; Piner, R. D.; Stankovich, S.; Jung, I.; Field, D. A.; Ventrice, C. A., Jr.; Ruoff, R. S. Chemical Analysis of Graphene Oxide Films after Heat and Chemical Treatments by X-ray Photoelectron and Micro-Raman Spectroscopy. Carbon 2009,47, 145-152. (43) Wunderlich B. Equilibrium melting of Flexible Linear Macromolecules. Polym. Eng. Sci.1978, 18, 431-436 (44) Litchfield D.W.; Baird D.G. The Role of Nanoclay in the Generation of Poly(ethylene terephthalate) Fibers with Improved Modulus and Tenacity. Polymer, 2008, 49 5027–5036 (45) Kim H.; Macosko C.W. Morphology and Properties of Polyester/Exfoliated Graphite Nanocomposites. Macromolecules 2008, 41, 3317–3327.

(46) Ignatov, V. N.; Kuznetsov, A. I.; Vasnev, V. A.; Gribova, I. A.; Vinogradova, S. V; Gureeva, G. I.; Lavrukhin, B. D.; Krasnov, A. I. Polycondensation in the Presence of a Filler as a New Way to Create Polymer Multicomponent Systems with Prespecified Properties. Macromolecules 1992, 25, 1652-1657. 37 ACS Paragon Plus Environment


(47) Sclavons, M.; Carlier, V.; Legras, R.

Page 38 of 44

Anhydride Chain Ends Functionalization of

Poly(ethylene terephthalate). Polym. Eng .and Sci. 1999, 39, 789-803.

(48) Park, S.; An, J.; Jung, I.; Piner, R. D.; An S. J. Li X.; Velamakanni, A.; Ruoff R.S. Colloidal Suspensions of Highly Reduced Graphene Oxide in a Wide Variety of Organic Solvents. Nano

Lett. 2009, 9, 1593-1597.

(49) Li, L.; Qin, Z.Y.; Liang, X.; Fan, Q.Q.; Lu, Y.Q.; Wu, W.H.; Zhu M.F. Facile Fabrication of Uniform Core−Shell Structured Carbon Nanotube−Polyaniline Nanocomposites. J. Phys.

Chem. C 2009, 113, 5502–5507.

(50) Tang, B.; Wu, P.,; Siesler, H. W. In Situ Study of Diffusion and Interaction of Water and Mono- or Divalent Anions in a Positively Charged Membrane Using Two-Dimensional Correlation FT-IR/Attenuated Total Reflection Spectroscopy. J. Phys. Chem. B 2008, 112, 28802887. (51) Graf, D.; Molitor F.; Ensslin K.; Stampfer, C.; Jungen, A.; Hierold, C.; Wirtz, L. Spatially Resolved Raman Spectroscopy of Single- and Few-Layer Graphene, Nano Lett. 2007, 7, 238-242. (52) Qi, X.; Pu, K.; Zhou, X.; Li H.; Liu, B.; Boey, F.; Huang. W.: Zhang, H. Small 2010, 6, 663–669.


Page 39 of 44 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


(53) Jia, Z.; Wang, Y. Covalently Crosslinked Graphene Oxide Membranes by Esterification Reactions for Ions Separation. J. Mater. Chem. A 2015, 3, 4405–4412.

(54) Ran, S.; Wang, Z.; Burger, C.; Chu, B.; Hsiao, B. S. Mesophase as the Precursor for StrainInduced Crystallization in Amorphous Poly(ethylene terephthalate) Film. Macromolecules 2002, 35, 10102-10107. (55) Matsuo, M.; Tamada, M.; Terada, T.; Sawatari, C.; Niwa, M. Deformation Mechanism of Poly(ethylene terephthalate) Film under Uniaxial Stretching, Macromolecules 1982, 15, 988-998. (56) Shante, V.K.S.; Kirkpatrick, S. An Introduction to Percolation Theory. Adv. Phys. 1971, 20, 325-357.

(57) Miller, R. G. J.; Willis, H. A. An Independent Measurement of the Amorphous Content of Polymers. J. Polym. Sci. 1956, 19, 485–494.

(58) Vermant, J.; Ceccia, S.; Dolgovskij, M.K.; Maffettone, P.L.; Macosko, C.W. Quantifying Dispersion of Layered Nanocomposites via Melt Rheology. J. Rheol. 2007, 51, 429-450. (59) Rosu, R. F,; Shanks, R. A; Bhattacharya, S. N. Dynamic Rheology of Branched Poly(ethylene terephthalate). Polym. Int. 2000, 49, 203-208. (60) Kim H.; Miura Y.; Macosko C.W. Graphene/Polyurethane Nanocomposites for Improved Gas Barrier and Electrical Conductivity. Chem. Mater. 2010, 22, 3441–3450. 39 ACS Paragon Plus Environment


Page 40 of 44

(61) Krishnamoorti R.; Giannelis E. P. Rheology of End-tethered Polymer Layered Silicate Nanocomposites. Macromolecules 1997, 30, 4097-4102. (62) Koener, H.; Price,.G.; Pearce, N A.; Alexander, M.; Vaia, R.A; Remotely Actuated Polymer Nanocomposites—Stress-Recovery of Carbon-Nanotube-Filled Thermoplastic Elastomers, Nature Materials 2004, 3, 115–120 (63) Solomon, M. J.; Almusallam, A. S.; Seefeldt, K F.; Somwangthanaroj A.; Varadan P. Polypropylene/Clay Hybrid Materials. Macromolecules 2001, 34, 1864-1872. (64) Abdel-Goad M.; Poetschke P.; Zhou D.; Mark J.E.; Heinrich G. Preparation and Rheological Characterization of Polymer Nanocomposites Based on Expanded Graphite. J. Macromol. Sci. Part A 2007, 44, 591-598. (65) Kuhn, W.;

Kuhn H. Dependence of Viscosity on Velocity Gradient in Very Dilute

Suspensions and Solutions. Helv. Chim. Acta 1945. 28, 97-127.

(66) Ismail, I.; Abdala, A. A.; Macosko, C. W. Modeling the Intrinsic Viscosity of Polydisperse Disks . J. Rheology 2017, 61, 997–1006. (67) Park, O.; Hwang, Y.; Goh, M.; Lee, J. H.; Ku, B.; You, N. Mechanically Strong and Multifunctional Polyimide Nanocomposites Using Amimophenyl Functionalized Graphene Nanosheets, Macromolecules, 2013, 46, 3505–3511.


Page 41 of 44 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


(68) Chen, D.; Zhu,H.; Liu, T. Effect of Chemical Modification of Graphene on Mechanical, Electrical, and Thermal Properties of Polyimide/Graphene Nanocomposites, ACS Appl. Mater. Interfaces, 2010, 2 , 3702–3708. (69) Pradhan, S.: Costa, F. R.; Wagenknecht, U.; Jehnichen, D.; Bhowmick, A.K.; Heinrich, G. Elastomer/LDH Nanocomposites: Synthesis and Studies on Nanoparticle Dispersion, Mechanical Properties and Interfacial Adhesion. Eur. Polym. J. 2008, 44, 3122–3132.

(70) Carretero-Gonzalez, J; Verdejo, R.; Toki, S; Hsiao, B. S,; Giannelis, E. P.; LópezManchado, M. A. Real-Time Crystallization of Organoclay Nanoparticle Filled Natural Rubber under Stretching. Macromolecules 2008, 41, 2295-2298.

(71) Guhrémont, J.; Ajji, A.; Cole, K. C.; Dumoulin, M. M. Orientation and Conformation in Poly(ethylene terephthalate) with Low Draw Ratios as Characterized by Specular Reflection Infra-red Spectroscopy. Polymer 1995, 36, 3385-3392.

(72) Ajji, A.; Cole, K. C.; Dumoulin M. M.; Brisson, J. Amorphous Orientation of Poly(ethylene terephthalate) by X-ray Diffraction in Combination with Fourier Transform Infra-red Spectroscopy. Polymer 1995, 37, 4023-4030.



Page 42 of 44

(73) Ajji, A.; Guhrémont, J.; Cole, K. C.; Dumoulin, M. M. Orientation and Structure of Drawn Poly(ethylene terephthalate). Polymer 1996, 37, 3707-3714. (74) Padibjo, S. R.; Ward, I. M. A Structural Study of the Tensile Drawing Behaviour of Poly(ethylene terephthalate). Polymer 1983, 24, 1103-1112. (75) Yang S.; Michielsen S.; Orientation Distribution Functions Obtained via Polarized Raman Spectroscopy of Poly(ethylene terephthalate) Fibers. Macromolecules, 2003, 36, 6484–6492 (76) Lesko C.C.C.; Rabolt J.F.; Ikeda R.M.; Chase B.; Kennedy A. Experimental Determination of the Fiber Orientation Parameters and the Raman Tensor of the 1614 cm-1 Band of Poly(ethylene terephthalate). J. Molecular Structure, 2000, 521,127–136 (77) J é ol S.; Fenouillot F.; Rousseau A.; Masenelli-Varlot,K.; Gauthier C.; Briois J, Drastic Modification of the Dispersion State of Submicron Silica during Biaxial Deformation of Poly(ethylene terephthalate) Macromolecules, 2007, 40, 3229-3237. (78) Schadler L.S.; Giannaris S.C.; Ajayan P.M. Load Transfer in Carbon Nanotube Epoxy Composites. Appl. Phys. Lett. 1998, 73, 3842-3844.

(79) Hess, C. ;Hirt, P.; Oppermann W. Influence of Branching on the Properties of Poly(ethylene terephthalate) Fibers. J. Appl. Polym. Sci.1999, 74, 728–734.


Page 43 of 44 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


(80) Bikiaris, D. N.; Karayannidis, G. P. Synthesis and Characterisation of Branched and Partially Crosslinked Poly (ethylene terephthalate). Polym Int 2003, 52, 1230–1239. (81) Rosu, R. F.; Shanks, R. A.; Bhattacharya S.N. Synthesis and Characterisation of Branched Poly(ethylene terephthalate). Polym Int 1997, 42, 267-275.

(82) Yeo, S.S.; Hsuan, Y.G. Evaluation of Creep Behavior of High Density Polyethylene and Polyethylene-terephthalate Geogrids. Geotextiles and Geomembranes 2010, 28, 409–421

(83) Van Den Heuvel, C.J.M.; Heuvel, H.M.; Faassen, W.A.; Veurink, J.; Lucas, L.J., Molecular Changes of PET Yarns during Stretching Measured with Rheo-optical Infrared Spectroscopy and Other Techniques. J. Appl. Polym. Sci. 1993, 49, 925–934.



Page 44 of 44

Graphical Abstract


Trimellitic Anhydride Modified Graphene

Recommend Documents