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Article Cite This: J. Phys. Chem. C 2018, 122, 6555−6565

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Enhanced Dielectric Constant, Ultralow Dielectric Loss, and HighStrength Imide-Functionalized Graphene Oxide/Hyperbranched Polyimide Nanocomposites Asma Iqbal,†,‡ Seok Hwan Lee,‡ Humaira Masood Siddiqi,*,† O. Ok Park,*,‡ and Toheed Akhter*,‡,§ †

Department of Chemistry, Quaid-i-Azam University, Islamabad 45320, Pakistan Department of Chemical and Biomolecular Engineering (BK21+ Graduate Program), Korea Advanced Institute of Science and Technology (KAIST), 291 Daehak-ro, Yuseong-gu, Daejeon 305-701, Republic of Korea § Department of Chemistry, School of Science, University of Management and Technology, C-II, Johar Town, Lahore, Pakistan

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‡

S Supporting Information *

ABSTRACT: We achieved for the first time the formation of a charge-transfer complex (CTC) between a novel hyperbranched polyimide (PI) and oligo-imide-functionalized graphene oxide (FGO), aiming for enhancing the dielectric properties of resulting PI−FGO nanocomposites. This novel hyperbranched PI was derived from a new diamine N1,N1′-(4,4′oxybis(4,1-phenylene))bis(N1-(4-aminophenyl)benzene-1,4-diamine). The imide moieties were integrated on amine-FGO via a step-by-step condensation and thermal imidization approach. This FGO exhibited excellent compatibility with hyperbranched PI because of the formation of a CTC between two domains. In viscoelastic measurements, the dynamic storage modulus and glass-transition temperature of flexible PI−FGO nanocomposites increased linearly with increasing FGO contents. The synthesized nanocomposites revealed high mechanical properties with a tensile strength as high as 1.122 GPa. Thermogravimetric analysis demonstrates that these nanocomposite films exhibit high thermal stability up to 550 °C. Remarkably, the dielectric constant increases up to 42.47 at 8 wt % FGO loading with a dielectric loss as low as 0.0018 while maintaining the breakdown strength as high as 147.3 ± 4.5 MV/m.

1. INTRODUCTION Energy storage technologies are gaining importance rapidly with the development in the solar and wind power production sector.1−3 Among energy storage devices, capacitors can demonstrate high power density because of their fast charge and discharge ability.4−6 Efficiency of capacitors directly depends on dielectric materials with high permittivity.7 Previously, commonly used dielectrics are based on ceramics; however, their frangibility, high density, and poor plasticity hinder them to meet the requirements of many applications.8 On the basis of flexibility, low dielectric loss, and appropriate breakdown strength, polymers were found to be more suitable dielectric materials than inorganic ceramics.8,9 However, the development of polymer dielectrics that can function well at high temperature is a key issue. For instance, dielectrics based on biaxially oriented polypropylene, polyvinylidene fluoride, and so forth can work at a maximum temperature of 105 °C, after that they lose their utility.10,11 Therefore, to meet the demand of high-temperature applications, polyimides (PIs) were employed as high-performance dielectrics because of their excellent thermal stability (decomposition temperature over 550 °C) coupled with mechanical strength and chemical resistance.12−16 However, PIs possess inherently low dielectric constant in the range of 2.8−5, whereas their dielectric loss varies in the © 2018 American Chemical Society

range of 0.001−0.03; this low dielectric loss is highly desirable for practical applications.8,17−19 Mostly two approaches are adapted for increasing the dielectric constant of PIs. First one is adding a high-k ceramic reinforcement such as Al2O3,20,21 nanoTiO2,16 CaCu3−Ti4O12,22,23 and BaTiO324−26 into the PI matrix. However, to improve the dielectric constant significantly, a large amount of ceramic filler is needed which results in brittleness of the final PI matrix and hinders the development of PI-based dielectrics. The second approach is introducing conductive fillers including graphene, carbon nanotubes (CNTs), carbon black, and so forth. In our previous study, we reported γ-ray-irradiated multiwall CNT−PI nanocomposites with a higher dielectric constant of 36.25; however, their dielectric loss also remained higher due to the leakage of current because of the direct connection between conductive CNTs at higher filler loading.13 Recently, graphene is gaining importance as a conductive filler because of its remarkable mechanical, thermal, and electrical properties.27−30 It has been incorporated in various polymer matrices including polyaniline,31,32 carboxylated nitrile Received: January 16, 2018 Revised: March 3, 2018 Published: March 5, 2018 6555

DOI: 10.1021/acs.jpcc.8b00493 J. Phys. Chem. C 2018, 122, 6555−6565

Article

The Journal of Physical Chemistry C

Figure 1. Synthesis of N,1N1′-(4,4′-oxybis(4,1-phenylene))bis(N1-(4-aminophenyl)benzene-1,4-diamine) (2).

rubber,33−35 and polyvinylidene fluoride36−39 for various applications. Similarly, various research groups have studied the dispersion of graphene oxide (GO) in the PI matrix to improve the dielectric constant and mechanical properties.40−44 For the efficient reinforcement of polymer matrices with graphene, two key factors must be considered: (i) uniform dispersion of graphene into the polymer matrix and (ii) force of attraction between the polymer and graphene. Modification of GO by attaching various functional groups to its surface, that is, carboxyl groups, amino groups, and so forth has been the most exploited approach to circumvent these two challenges.1,8,45−47 However, in all of the previously published reports, improvement in the dielectric constant was not significant; also, the dielectric loss increased tremendously with the increase in the dielectric constant which can restrict the practical application of these PI−GO nanocomposites as dielectric materials. In all of these studies, the PI matrix consists mostly of linear PI chains and functional groups attached to the modified GO was mostly lacking structural similarity with PI, leading to a weak interfacial interaction and difficulty in efficient dispersion of GO in the PI matrix at a higher loading of the filler.1,44 Consequently, GO agglomerated developing direct connection with each other which can cause increase in the dielectric loss. It is well-established that the high mechanical and thermal properties of PIs are due to the formation of a charge-transfer complex (CTC) between adjacent PI chains, leading to strong forces of attraction between them.48,49 In the present study, we first time report the enhancement of compatibility of GO with hyperbranched PI via the formation of a similar CTC between the GO filler and PI matrix. This formation of CTC was achieved by attaching imide oligomers on the surface of GO to introduce structural similarity with PI. When this functionalized GO (FGO) was dispersed in the PI matrix, a strong CTC was formed between imide functionalities on FGO and PI matrix, leading to efficient dispersion of FGO in the matrix affording the dielectric constant of nanocomposites higher than those of all previously reported PI−GO nanocomposites. Our aim of synthesizing hyperbranched PI instead of linear-chain PI was to increase the density of the CTC. This hyperbranched PI ensured the formation of the CTC with FGO to a higher extent on one hand and to ensure the complete wrapping of FGO with extended PI chains on other hand. This complete wrapping of FGO with hyperbranched PI prevented the direct connection of GO particles, impeding the dielectric loss to go to higher value. To the best of our knowledge, this is the first report of hyperbranched PI−GO nanocomposites with

improved interfacial interaction by the CTC formation between the matrix and filler which afforded high dielectric constant, low dielectric loss, outstanding temperature resistance, and mechanical properties.

2. EXPERIMENTAL SECTION 2.1. Materials and Measurements. (3-Aminopropyl)trimethoxysilane (APTMS) and GO were purchased from Sigma-Aldrich. 4,4-Oxydianiline (ODA) and pyromelliticdianhydride (PMDA) were supplied by Sinopharm Chemical Reagent Co., Ltd. N,N-Dimethylformamide and ethanol were provided by Alfa Aesar and were used as received. All other chemicals and reagents were of commercial grade and used as received, unless otherwise mentioned. The solvents were purified and dried by standard protocols. 2.2. Instrumentation. The melting points were determined using Gallenkamp (Sanyo) model MPDBM 3.5 with a digital thermometer. Fourier transform infrared spectroscopy (FTIR) spectra were recorded on a Bruker α-Alpha-P model using the attenuated total reflection method. 1H NMR and 13C NMR spectral analyses were done on Bruker (300 MHz) spectrophotometers in dimethyl sulfoxide (DMSO-d6) using tetramethylsilane (TMS) as an internal standard reference. The elemental analysis was carried out using a CHNS analyzer, Thermo Scientific (Flash 2000 series). The morphology of the surface of the nanocomposite was observed by field emission scanning electron microscopy (FE-SEM, S4800 Hitachi Japan) with an accelerating voltage of 10 kV. All samples were sputtered with platinum prior to the scanning electron microscopy (SEM) analysis. The thermal stability of the synthesized nanocomposites was measured by thermal gravimetric analysis (NETZSCH thermal analysis system TG209) under a nitrogen atmosphere with a heating rate 10 °C/min and a scanning range from room temperature to 800 °C. X-ray photoelectron spectroscopy (XPS) was conducted on a 60 (GENESIS EDAX, US) with Al Kα radiation (hν = 1486.4 eV). Viscoelastic behavior of nanocomposite films was investigated using a TA Q800 DMA device in the tension mode in the temperature range of 30−500 °C. Mechanical tests were conducted on a Shimadzu AGX-100NX universal testing machine at ambient temperature. The nanocomposite films were cut into strips and dried at 120 °C under vacuum for 24 h before analysis. The dielectric properties measurements were carried out on a Concept 40 broadband dielectric spectroscope (Novocontrol Technologies, KG, Germany). The breakdown strength was measured using a dc dielectric strength tester with 6556


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Figure 2. 1H NMR of N1-(4-aminophenyl)-N1-(4-(4-(bis(4-aminophenyl)amino)phenoxy)phenyl)benzene-1,4-diamine (2).

three-neck round-bottom flask, equipped with a stirring bar and a reflux condenser, was charged with (1) (2.92 mmol), 10% Pd/C (0.2 g), and ethanol (70 mL) and heated at reflux temperature. Hydrazine hydrate (9 mL) was added dropwise into this boiling solution/suspension, and the reaction mixture was refluxed for 24 h. The progress of reaction was monitored by thin-layer chromatography (n-hexane/EtOAc, 1:1). After complete conversion of the tetranitro compound into tetraamino compound, Pd/C was filtered off. The solvent was evaporated from the filtrate under reduced pressure, and the obtained solid product was recrystallized from ethanol to afford a light pink solid. Yield: 65%. mp 145 °C, Rf = 0.31 (n-hexane/EtOAc, 2:1), FT-IR (ν̅/cm−1): (3249, 3190 −NH2 stretch), (1310 C−N stretch for tertiary amine), (1630 N−H bend primary amine). 1 H NMR (DMSO-TMS 25 °C) (ppm): 6.51, (d, 8H, J = 8.7 Hz, Ph-H 2,2′), 6.7, (m, 12H, Ph-H, 3,3′,6,6′), 6.6 (t, J = 9 Hz, 4H, Ph-H,7,7′), 4.9 (s, NH2, 8H). 13C NMR: 150.11 (8), 145.6 (1), 145.4 (5), 137.1 (4), 126.9 (3,3′), 119.6 (7,7′), 119.2 (6,6′), 115.2 (2,2′). Anal. Calcd for C36 H32 N6 O (%): C, 76.57; H, 5.71; N, 14.88. Found C, 76.07; H, 5.85; N, 14.38. 2.4. Surface Modification of GO. Imide-based functionalities were incorporated on the surface of GO aiming for the formation of CTC with PI chains, leading to uniform dispersion of FGO into the PI matrix, as shown in Figure 2. For this purpose, a homogeneous dispersion of GO in anhydrous tetrahydrofuran (THF) (0.1 mg/mL) was prepared by sonication for 30 min at room temperature. To this dispersion, APTMS (10:1, APTMS/GO) was added and the mixture was refluxed for 15 h. Afterward, the solution was filtered to recover APTMS-modified GO which was washed repeatedly with THF

a sphere−sphere stainless electrode (DH, Shanghai Lanpotronics Co., China). The breakdown tests were carried out on five specimens of each film in an oil bath. 2.3. Monomer Synthesis. 2.3.1. Synthesis of 4,4′Oxybis(N,N-bis(4-nitrophenyl)aniline) (1). The precursor of tetraamine, that is, tetranitro compound, was synthesized by the cesium fluoride-mediated N-arylation of 4-fluoro-1-nitrobenzene by ODA, as shown in Figure 1. For this purpose, 4-fluoro1-nitrobenzene (42 mmol), cesium fluoride (46 mmol), ODA (20 mmol), and DMSO (50 mL) were taken in a two-necked round bottom flask equipped with a condenser and a magnetic stirring bar. The reaction mixture was heated under inert atmosphere of nitrogen at 120 °C for 24 h. The consumption of reactants was analyzed by a thin-layer chromatograph (nhexane/EtOAc, 3:1). After the completion of reaction, the mixture was allowed to cool to room temperature. Then, 150 mL of distilled water was added resulting in the precipitation of tetranitro compound which was filtered off, washed several times with distilled water, and purified by column chromatography. Yield: 90%. mp: 182 °C. Rf = 0.54 (n-hexane/EtOAc, 3:1). −1 ): 1309 (C−N of ter amine stretch), 1570, 1337 FT-IR (ν/cm ̅ (−NO2 stretch); 1H NMR (DMSO-d6), 8.19 (d, 8H of four phenyl rings ortho to nitro group), 7.33 (d, 8H of four phenyl groups meta to nitro), 7.23, (m, 8H of phenoxy rings). 13C NMR (DMSO-d6) 152.62, 146.98, 141.64, 138.03, 125.90, 126.26, 124.94, 124.0, 121.14. Anal. Calcd for C36H24N6O (%): C, 63.16; H, 3.53; N, 12.28. Found: C, 62.07; H, 3.65; N, 11.78. 2.3.2. Synthesis of N1,N1′-(4,4′-Oxybis(4,1-phenylene))bis1 (N -(4-aminophenyl)benzene-1,4-diamine) (2). For reduction of tetranitro compound into a tetraamine monomer, a 250 mL 6557


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Figure 3. 13C NMR of N1-(4-aminophenyl)-N1-(4-(4-(bis(4-aminophenyl)amino)phenoxy)phenyl)benzene 1,4-diamine (2).

transferred to a Teflon Petri dish and heated at 70 °C for 12 h to cast PAA/FGO films. Subsequently, PAA/FGO films were converted to PI/FGO nanocomposite films by thermal imidization. 2.6.2. PGC−GO. Here, the PGCs were prepared by the same method as discussed above except for the incorporation of unmodified GO. The amount of FGO/GO that was incorporated in the PI matrices was varied between 0 and 8 wt %. Table 1 shows the codes and the composition of different types of PGCs.

and dried in a vacuum oven. The dried APTMS-modified GO was dispersed in anhydrous N-methyl-2-pyrrolidone (NMP) (0.1 mg/mL) by sonication for 30 min. Subsequently, PMDA (10:1) was added to this dispersion followed by heating at reflux for 15 h. After cooling to room temperature, the GO with a terminal anhydride functional end group was filtered, washed with NMP several times, and dried under vacuum. The process was repeated with the addition of 2 (10:1), followed by the addition of PMDA, and then by 2 to the surface of FGO, resulting in the preparation of FGO containing an imide surface modifier and amine end groups. 2.5. Synthesis of PI. For the synthesis of high-molecularweight hyperbranched PI (Figure 3), tetraamine (0.5 mmol) was dissolved in anhydrous dimethylacetamide (DMAc) (3 mL) in a sealed glovebox under inert environment. When tetraamine was completely dissolved in the solvent, PMDA (1 mmol) was added to this solution. As the polymerization reaction started, the reaction mixture became very viscous which caused hindrance to the stirring of solution. Some extra amount of DMAc (02 mL) was poured for facile stirring of the solution. To compensate for any forfeited anhydride functional groups and to ensure anhydride as chain-end groups, a slight excess of PMDA (1%) was added to the reaction mixture after 6 h of stirring. The reaction mixture was stirred at about 0−5 °C for another 12 h. Afterward, the poly(amic acid) (PAA) solution was transferred to a Teflon Petri dish, and the film was casted by heating at 70 °C for 6 h. Thermal imidization of PAA into PI was carried out by heating the PAA film at 100, 200, and 300 °C each for 1 h. 2.6. Preparation of PI−GO Nanocomposites. For comparative study and evaluation of property benefits, following two types of PI graphene nanocomposites (PGCs) were synthesized. 2.6.1. PGC−FGO. A suspension of surface FGO in DMAc was prepared by sonication for 3 h. The PAA solution in DMAc was added to this FGO suspension. The mixture was sonicated for another 2 h. Afterward, the PAA/FGO suspension was

Table 1. Nomenclature and Composition of Different Types of PGCs PGCs

sample codes

PI (wt %)

FGO (wt %)

GO (wt %)

pristine PI PGC−FGO

PI 2-FGO 4-FGO 8-FGO 2-GO 4-GO 8-GO

100 99 96 92 99 96 92

00 2 4 8 00 00 00

00 00 00 00 2 4 8

PGC−GO

3. RESULTS AND DISCUSSION 3.1. Monomer Synthesis. Tetraamine was synthesized in two steps. The first step was cesium fluoride-mediated Narylation of ODA with 4-fluoronitrobenzene, giving a tetranitro compound. The second step was Pd/C-catalyzed reduction of the tetranitro compound to tetraamine, as shown in Figure 1. The structures of the synthesized monomer and its precursor were evaluated by spectroscopic methods and elemental analysis. First, in the FTIR spectrum, the presence of nitro groups in the product was analyzed by two absorption bands at 1596 and 1337 cm−1. The structure of this tetranitro compound was further confirmed by NMR spectroscopy. In the 1H NMR spectrum, the resonance peak for protons of the amino group 6558


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Figure 4. Surface functionalization of GO by the attachment of imide-based oligomers.

Figure 5. Synthesis of PI.

resonated at 6.54 ppm as a doublet in the 1H NMR spectrum of the tetraamine monomer. 3.2. Synthesis of PI. High-molecular-weight PI was synthesized by mixing stoichiometric amount of tetraamine monomer with PMDA. First, the reaction between these two monomers resulted in the synthesis of PAA, which was thermally imidized to convert PAA into PI. First, PAA was synthesized by the reaction between these two monomers, which was thermally imidized to convert PAA into PI. The synthesis of PAA and its conversion to PI was analyzed by FTIR spectroscopy. FTIR spectra of PAA and PI are given in Figure 6a. After thermal imidization, the conversion of PAA into PI was evaluated by the disappearance of the absorption band at 1650 cm−1 that was assigned to the amide carbonyl group (Figure 6b).

of ODA was not observed. Also, a doublet at 8.19 ppm, which was assigned to eight protons located in ortho to nitro groups, confirmed the successful synthesis of the tetranitro compound. Similarly, in the FTIR spectrum of tetraamine, the characteristic absorption bands for −NH stretching of amino group appeared at 3249 and 3190 cm−1. Furthermore, the absorption bands corresponding to nitro groups disappeared in the FTIR spectrum of the tetraamine monomer. Comprehensive evidence for complete conversion of tetranitro compound into tetraamine was furnished by NMR spectroscopic analysis (Figures 4 and 5). In the 1H NMR spectrum, a singlet at 4.84 ppm was assigned to eight protons of the amino group. The disappearance of the resonance peak at 8.19 ppm (for eight protons at the ortho position of the tetranitro compound) also verified the completion of reduction of nitro to amine. Likewise, protons ortho to amino groups 6559


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Figure 6. FTIR of (a) [GO (A), GO modified with APTMS (B), FGO (C)] and (b) [PAA and PI].

Figure 7. XPS spectra of (a) GO and (b) FGO; the inset graph shows the N 1s scan.

The terminal anhydride groups in this modified GO were reacted with tetraamine. Thus, imide oligomers on the surface of GO were confirmed from the presence of absorption bands for imide carbonyl groups at 1779 and 1719 cm−1, imide C− N−C bond at 1369 cm−1, and imide ring deformation at 721 cm−1. 3.3.2. X-ray Photoelectron Spectroscopy. The covalent bonding of imide oligomers to the surface of GO was investigated by XPS. The XPS spectrum is given in Figure 7. The atomic percentages of the constituent elements in the surface-functionalized graphene were found to be 9.96% for nitrogen (N 1s), 10.6% for oxygen (O 1s), 75.9% for carbon (C 1s), and 3.54% for silicon (Si 2p). The presence of nitrogen was attributed to imide oligomers, whereas silicon atoms came from 3-aminoisopropyltrimethoxysilane which acts like a bridge between imide oligomers and GO surface. Besides, the XPS spectrum revealed that the atomic percentage of oxygen increased significantly as compared to the unmodified GO. This increment in the percentage of oxygen also provides a

Furthermore, the appearance of new absorption bands at 1775 cm−1 (asymmetric stretching of CO), 1714 cm−1 (symmetric stretching of CO), 1365 cm−1 (symmetric stretching of C−N), and 720 cm−1 (imide ring deformation) also verified the complete thermal imidization of PAA into PI. 3.3. Surface Functionalization of GO. 3.3.1. FTIR Spectroscopy of FGO. To improve dispersion of GO in the PI matrix and develop CTC between GO and PI matrix, surface FGO was prepared. For this purpose, imide functionalities were mounted by reacting amino group on the surface of GO with PMDA, which, in turn, was reacted with the tetraamine monomer. The presence of imide moieties on the surface of GO was investigated by FTIR spectroscopy. FTIR spectra of different stages of modification of GO are shown in Figure 6b. The FTIR spectrum of APTMS-modified GO exhibits the −CH stretching vibrations at 2925 and 2870 cm−1, −NH bending vibrations at 1552 cm−1, and O−Si stretching vibrations at 1020 cm−1. Reaction of this APTMS-modified GO with PMDA produces amide carbonyl and carboxylic acid groups which were thermally imidized to imide functionalities. 6560


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resolution, as depicted in Figure 8c. FGO sheets in PGC−FGO films are well-wrapped and laminated with a thick PI layer that can be attributed to enhanced force of attraction via the formation of CTC between FGO and PI matrix. This CTC with a PI matrix prevented aggregation of FGO by isolating them from each other, whereas in the case of PGC−GO composite film, where unmodified GO does not form CTC with matrix, GO sheets are not well-dispersed and are aggregated in the PI matrix, as obvious from Figure 8d. Thus, it can be concluded that the formation of CTC between PI matrix and FGO resulted in a well-dispersed filler state. 3.5. Thermogravimetric Analysis of PGCs. Thermogravimetric analysis (TGA) was employed to investigate the thermal stability of synthesized PI−GO nanocomposites (PGC−FGO and PGC−GO). The TGA thermograms are shown in Figure 9a, whereas the thermal parameters are presented in Table S1. Thermal stability was evaluated in terms of temperature at 5% weight loss (Td), 10% weight loss (T10), and residual mass at 1000 °C (R1000). It is obvious from Figure 9a that neat PI did not show any significant weight loss before 550 °C. After this temperature, a fast degradation was examined. The thermal degradation temperature Td for neat PI is 483 °C, whereas those for PGC−FGO and PGC−GO are 48 and 30 °C higher in comparison to neat PI, respectively. Comparing thermal stabilities of PGC−FGO and PGC−GO, it is obvious that FGO exhibited a greater increase in Td, T10, and R1000 in contrast to GO. FGO is well-assimilated and strongly adhere to PI as compared to GO. Thus, it may be inferred that the formation of CTC between PI matrix and FGO resulted in improved interfacial interaction between the

comprehensive evidence of the presence of imide oligomers on the surface of FGO. 3.4. Morphological Studies. FE-SEM micrographs of the fractured surfaces of neat PI, PGC−FGO, and PGC−GO containing 4 wt % FGO and GO are presented in Figure 8a−d.

Figure 8. SEM images (a−c) PGC−FGO and (d) PGC−GO.

It is obvious from Figure 8a that neat PI exhibits uniform and even surface which is commonly observed in conducting polymers.50 On the other hand, Figure 8b reflects that the PGC−FGO film has a rough surface with some folds and wrinkles which are due to densely packed crumpled FGO sheets creating wrinkles and bent up in the surface of the film. These wrinkles and folds are even more prominent at high

Figure 9. (a) TGA of PGC−FGO and PGC−GO, (b) dynamic storage moduli of PI, PGC−FGO, and PGC−GO as a function of FGO/GO content, (c) tan δ of PI, PGC−FGO, and PGC−GO as a function of FGO/GO, and (d) stress−strain curves for PGC−FGO nanocomposites. 6561


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Figure 10. (a) Dielectric constant and dielectric loss of PGC−FGO and PGC−GO as a function of filler contents and (b) breakdown strength of PGC−FGO as a function of FGO contents.

two components imparting greater thermal stability to the PGC−FGO system as compared to PGC−GO. The delayed decomposition process of polymer matrix shows that the formation of a CTC between two phases impeded diffusion of oxygen into the PI matrix and reduced the degradation at elevated temperature. In the case of PGC−FGO, FGO acted as an effective heat and mass transport barriers because of the formation of CTC and hence much suitable for preserving the PI matrix from thermal degradation. 3.6. Viscoelastic Properties. Figure 9b,c presents the viscoelastic properties of synthesized nanocomposites reinforced with (2−8 wt %) FGO and GO. A CTC was created between PI oligomers on the surface of FGO and PI matrix, aiming for maximum load transfer between both components. The variations of dynamic storage modulus (E′) as a function of temperature for PGCs (PGC−FGO and PGC−GO) reinforced with 2−8 wt % FGO/GO are presented in Figure 9b, whereas Table S2 displays Tg and E′ values at room temperature and at Tg also. It is obvious from Figure 9b that the addition of FGO improved the values of E′ for PGC−FGO as well as for PGC−GO as compared to the pristine PI matrix. For example, the value at 30 °C was 2.588 GPa for pristine PI which increased to 4.586 GPa for PGC−FGO at 8% loading of FGO at the same temperature, whereas this value was 4.243 GPa for PGC−GO at the same temperature and wt % loading of unfunctionalized GO. The effect of the formation of CTC between FGO and PI matrix is more prominent from the E′ values at Tg. In rubbery plateau, that is, the region above Tg, all PGC−FGO composite films demonstrated remarkably higher values of E′ as compared to PGC−GO composite films, for instance, 8-FGO has a E′ of 1.885 GPa at Tg, whereas 8-GO has a E′ of 0.819 GPa. Thus, it can be inferred that significantly enhanced forces of attraction between FGO and PI matrix via CTC restricted the segmental motion of polymer chains, resultantly, improving the E′ profile both in rubbery as well as in glassy state. Similarly, the effect of enhanced forces of attraction between FGO and PI matrix through the formation of CTC on the cooperative movement of polymer chains and damping behavior during glass transition can be realized from variations in tan δ as a function of temperature, as shown in Figure 9c. It can be observed from Figure 9c that in the case of PGC−FGO, the curves became more broadened with the addition of FGO and their maxima shift to higher temperature consistently. However, in the case of PGC−GO, although tan δ curves shift to higher temperature as compared to pristine PI;

however, they remain almost unchanged with the increasing content of GO. In the case of PGC-FGO, as FGO contents increase from 0−8 wt %, Tg moves from 334 °C for pristine PI to 423 °C for 8-FGO and 386 °C for 8-GO. The shift of Tg to higher temperature can be attributed to the presence of CTC between FGO and PI matrix which hindered the chain segmental motions. Because of increased interactions between FGO and PI, the three-dimensional network requires a higher temperature for chain segmental motion as compared to pristine PI and PGC−GO composites, thus, shifting Tg to higher value. 3.7. Mechanical Properties. Typical stress−strain curves of neat PI and PGC−FGO nanocomposites are presented in Figure 9d. The interfacial interactions between the filler and matrix, and hence, the control of aggregation of the filler is a key factor for improving the mechanical properties of nanocomposites.51,52 As clear from Figure 9d that with the increase in FGO contents, the tensile strength linearly improved from 91.88 MPa for pure PI to 122.47 MPa for PGC−8-FGO. The consistent increase in mechanical properties of nanocomposites on addition of FGO can be attributed to strong adhesion of FGO with the PI matrix. As the strain moves to higher values, coiled polymer chains straighten to accommodate the deformation. This uncoiling of PI chains can create more room for the CTC between the filler and matrix, consequently external load is transferred to the discontinuous high-modulus FGO phase from the continuous low-modulus matrix phase via a shear stress at the PI−FGO interface, making nanocomposites more stiffer with higher load carrying ability in the axial direction.53−55 3.8. Dielectric Properties. The dielectric constant and dielectric loss of synthesized nanocomposites were investigated as a function of FGO/GO contents. The variations in the dielectric constant along with variations in GO contents are shown in Figure 10a. The permittivity enhanced remarkably as the FGO contents reached and exceeded the percolation threshold as narrated by percolation theory.1 The dielectric constant achieved a maximum value of 42.47 at 8 wt % FGO loading. This value is higher by a factor of 5.5 than the pristine PI matrix with a dielectric constant of 7.72, confirming that FGO is homogeneously dispersed in the PI matrix.56 Similarly, the dielectric loss for PGC−FGO at 8 wt % is 0.0018. To the best of our knowledge, this is the highest value of dielectric constant and lowest value of dielectric loss as compared to previously reported PI GO nanocomposites.19,57 6562



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The dramatic increase of dielectric constant with increasing FGO contents is ascribed to the microcapacitor network model. 13,58 The contiguous FGO domains, which are completely coated with the polymer matrix, are acting as electrodes, and the polymer bed present between them is acting as a dielectric layer. In this way, various microcapacitors are generated inside nanocomposites. At a higher filler contents like 8 wt %, microcapacitors increased and the dielectric bed between them decreased, resulting in increased capacitance and higher dielectric constant value. The key factor in achieving higher dielectric constant was uniform dispersion and complete wrapping of FGO with the PI matrix, which in turn can be attributed to remarkable compatibility of FGO with the PI matrix because of the formation of CTC between FGO and PI. As obvious from Figure 10a, in the case of PGC−GO where pristine GO was dispersed in the PI matrix, the dielectric constant remained much lower as compared to PGC−FGO. Moreover, the value of dielectric constant decreased at 8 wt % GO loading as compared to 4 wt %, indicating poor dispersion of GO in the PI matrix, which is also indicated by increased dielectric loss at 8 wt % GO content. Along with high dielectric constant, high breakdown strength is also an important factor for governing large energy storage of dielectric materials.8 Breakdown strength dictates the operating electric field as well as maximum energy storage density of a particular dielectric material. Figure 10b exhibits the breakdown strength of the PI matrix and PGC−FGO nanocomposites at room temperature as a function of FGO contents. It is obvious from Figure 10b that pure PI has a very high breakdown strength of 301.8 ± 5.1 MV m−1. The breakdown strength decreased to 172.4 ± 4.3 MV m−1 by the addition of 2 wt % FGO, which is due to different electrical properties of FGO and PI matrix. With further increase in FGO contents, the breakdown strength decreased further. However, at 8 wt % loading of FGO, the breakdown strength is 147.3 ± 4.5 MV m−1 which is fairly higher as compared to the previously reported PI−GO nanocomposite.1,44

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.jpcc.8b00493. Data of TGA and viscoelastic properties (PDF)

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AUTHOR INFORMATION

Corresponding Authors

*E-mail: [email protected], [email protected] (H.M.S.). *E-mail: [email protected] (O.O.P.). *E-mail: [email protected] (T.A.). ORCID

Humaira Masood Siddiqi: 0000-0002-9634-9383 Author Contributions

The manuscript was written through contributions of all authors. All authors have given approval to the final version of the manuscript. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The authors greatly appreciate the support of this research by the Higher Education Commission of Pakistan under the national research program for universities research project (203821/NRPU/R&D/HEC/14).

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REFERENCES

(1) Fang, X.; Liu, X.; Cui, Z.-K.; Qian, J.; Pan, J.; Li, X.; Zhuang, Q. Preparation and properties of thermostable well-functionalized graphene oxide/polyimide composite films with high dielectric constant, low dielectric loss and high strength via in situ polymerization. J. Mater. Chem. A 2015, 3, 10005−10012. (2) Gu, W.; Peters, N.; Yushin, G. Functionalized carbon onions, detonation nanodiamond and mesoporous carbon as cathodes in Liion electrochemical energy storage devices. Carbon 2013, 53, 292− 301. (3) Zhi, M.; Xiang, C.; Li, J.; Li, M.; Wu, N. Nanostructured carbonmetal oxide composite electrodes for supercapacitors: a review. Nanoscale 2013, 5, 72−88. (4) Evanko, B.; Yoo, S. J.; Chun, S.-E.; Wang, X.; Ji, X.; Boettcher, S. W.; Stucky, G. D. Efficient charge storage in dual-redox electrochemical capacitors through reversible counterion-induced solid complexation. J. Am. Chem. Soc. 2016, 138, 9373−9376. (5) Ghidiu, M.; Lukatskaya, M. R.; Zhao, M.-Q.; Gogotsi, Y.; Barsoum, M. W. Conductive two-dimensional titanium carbide clay with high volumetric capacitance. Nature 2014, 516, 78−81. (6) Lang, X.; Hirata, A.; Fujita, T.; Chen, M. Nanoporous metal/ oxide hybrid electrodes for electrochemical supercapacitors. Nat. Nanotechnol. 2011, 6, 232−236. (7) Wang, Q.; Zhu, L. Polymer nanocomposites for electrical energy storage. J. Polym. Sci., Part B: Polym. Phys. 2011, 49, 1421−1429. (8) Chen, Y.; Lin, B.; Zhang, X.; Wang, J.; Lai, C.; Sun, Y.; Liu, Y.; Yang, H. Enhanced dielectric properties of amino-modified-CNT/ polyimide composite films with a sandwich structure. J. Mater. Chem. A 2014, 2, 14118−14126. (9) Schroeder, R.; Majewski, L. A.; Grell, M. High-performance organic transistors using solution-processed nanoparticle-filled high-k polymer gate insulators. Adv. Mater. 2005, 17, 1535−1539. (10) Rabuffi, M.; Picci, G. Status quo and future prospects for metallized polypropylene energy storage capacitors. IEEE Trans. Plasma Sci. 2002, 30, 1939−1942. (11) Sun, L. L.; Li, B.; Zhao, Y.; Mitchell, G.; Zhong, W. H. Structure-induced high dielectric constant and low loss of CNF/PVDF

4. CONCLUSIONS In conclusion, we have successfully developed CTC between a novel hyperbranched PI and FGO which resulted in strong forces of attraction between PI matrix and FGO. GO was functionalized by integrating imide functionalities on its surface so that it can have structural similarity with PI, leading to the formation of CTC between PI chains and imide moieties on the surface of FGO. PGC−FGO nanocomposites exhibit dielectric constant far higher than previously reported PIbased nanocomposites. This goal was achieved by efficient dispersion of FGO and its strong binding with the hyperbranched PI matrix. The extended branches in the matrix completely wrapped FGO which effectively prevented the direct connection of GO particles with each other, leading to an ultralow dielectric loss. More importantly, these PGC−FGO nanocomposites have high breakdown strength and energy storage density, which is required by capacitors for energy storage. Along with high dielectric constant, prepared PGC− FGO nanocomposites have high thermal and mechanical properties which make them desirable for the multifunctional applications. We believe that unique properties of these PGC− FGO nanocomposites will make them potential dielectrics for applications in the field of capacitors. 6563


Article

The Journal of Physical Chemistry C composites with heterogeneous CNF distribution. Nanotechnology 2010, 21, 305702. (12) Iqbal, A.; Lee, S. H.; Park, O. O.; Siddiqi, H. M.; Akhter, T. Synthesis and characterization of blue light emitting redox-active polyimides bearing a noncoplanar fused carbazole-triphenylamine unit. New J. Chem. 2016, 40, 5285−5293. (13) Akhter, T.; Mun, S. C.; Saeed, S.; Park, O. O.; Siddiqi, H. M. Enhancing the dielectric properties of highly compatible new polyimide/[gamma]-ray irradiated MWCNT nanocomposites. RSC Adv. 2015, 5, 71183−71189. (14) Chou, Y.-H.; Tsai, C.-L.; Chen, W.-C.; Liou, G.-S. Nonvolatile transistor memory devices based on high-k electrets of polyimide/ TiO2 hybrids. Polym. Chem. 2014, 5, 6718−6727. (15) Bin, Y.; Oishi, K.; Koganemaru, A.; Zhu, D.; Matsuo, M. Catalytic effect of nickel under carbonization of polyimide films. Carbon 2005, 43, 1617−1627. (16) Liu, X.; Yin, J.; Kong, Y.; Chen, M.; Feng, Y.; Wu, Z.; Su, B.; Lei, Q. The property and microstructure study of polyimide/nano-TiO2 hybrid films with sandwich structures. Thin Solid Films 2013, 544, 54− 58. (17) Kim, J.; Kwon, J.; Kim, M.; Do, J.; Lee, D.; Han, H. Lowdielectric-constant polyimide aerogel composite films with low water uptake. Polym. J. 2016, 48, 829−834. (18) Volksen, W.; Miller, R. D.; Dubois, G. Low dielectric constant materials. Chem. Rev. 2010, 110, 56−110. (19) Chen, F.; Bera, D.; Banerjee, S.; Agarwal, S. Low dielectric constant polyimide nanomats by electrospinning. Polym. Adv. Technol. 2012, 23, 951−957. (20) Qadri, S. B.; Gorzkowski, E.; Rath, B. B.; Feng, J.; Qadri, S. N.; Kim, H.; Caldwell, J. D.; Imam, M. A. Nanoparticles and nanorods of silicon carbide from the residues of corn. J. Appl. Phys. 2015, 117, 044306. (21) Alias, A.; Ahmad, Z.; Ismail, A. B. Preparation of polyimide/ Al2O3 composite films as improved solid dielectrics. Mater. Sci. Eng., B 2011, 176, 799−804. (22) Chi, Q.; Sun, J.; Zhang, C.; Liu, G.; Lin, J.; Wang, Y.; Wang, X.; Lei, Q. Enhanced dielectric performance of amorphous calcium copper titanate/polyimide hybrid film. J. Mater. Chem. C 2014, 2, 172−177. (23) Dang, Z.-M.; Zhou, T.; Yao, S.-H.; Yuan, J.-K.; Zha, J.-W.; Song, H.-T.; Li, J.-Y.; Chen, Q.; Yang, W.-T.; Bai, J. Advanced calcium copper titanate/polyimide functional hybrid films with high dielectric permittivity. Adv. Mater. 2009, 21, 2077−2082. (24) Abe, K.; Nagao, D.; Watanabe, A.; Konno, M. Fabrication of highly refractive barium-titanate-incorporated polyimide nanocomposite films with high permittivity and thermal stability. Polym. Int. 2013, 62, 141−145. (25) Hamciuc, C.; Hamciuc, E.; Olariu, M.; Ciobanu, R. Thermal and electrical behaviour of some hybrid polyimide films containing barium and titanium oxides. Polym. Int. 2010, 59, 668−675. (26) Wang, Y.; Wu, X.; Feng, C.; Zeng, Q. Improved dielectric properties of surface modified BaTiO3/polyimide composite films. Microelectron. Eng. 2016, 154, 17−21. (27) Rodríguez-Pérez, L.; Herranz, M. Á .; Martín, N. The chemistry of pristine graphene. Chem. Commun. 2013, 49, 3721−3735. (28) Wu, J.; Pisula, W.; Müllen, K. Graphenes as potential material for electronics. Chem. Rev. 2007, 107, 718−747. (29) Ebrahimi, M.; Rosei, F. Materials science: Organic analogues of graphene. Nature 2017, 542, 423−424. (30) Naumov, A.; Grote, F.; Overgaard, M.; Roth, A.; Halbig, C. E.; Nørgaard, K.; Guldi, D. M.; Eigler, S. Graphene Oxide: a one-versus two-component material. J. Am. Chem. Soc. 2016, 138, 11445−11448. (31) Li, M.; Huang, X.; Wu, C.; Xu, H.; Jiang, P.; Tanaka, T. Fabrication of two-dimensional hybrid sheets by decorating insulating PANI on reduced graphene oxide for polymer nanocomposites with low dielectric loss and high dielectric constant. J. Mater. Chem. 2012, 22, 23477−23484. (32) Sun, J. Z.; Qin, A.; Tang, B. Z. Functional polyacetylenes: hybrids with carbon nanotubes. Polym. Chem. 2013, 4, 211−223.

(33) Tian, M.; Zhang, J.; Zhang, L.; Liu, S.; Zan, X.; Nishi, T.; Ning, N. Graphene encapsulated rubber latex composites with high dielectric constant, low dielectric loss and low percolation threshold. J. Colloid Interface Sci. 2014, 430, 249−256. (34) Wu, S.; Hui, K. S.; Hui, K. N.; Kim, K. H. Electrostatic-induced assembly of graphene-encapsulated carbon@nickel−aluminum layered double hydroxide core−shell spheres hybrid structure for high-energy and high-power-density asymmetric supercapacitor. ACS Appl. Mater. Interfaces 2017, 9, 1395−1406. (35) Kretinin, A. V.; Cao, Y.; Tu, J. S.; Yu, G. L.; Jalil, R.; Novoselov, K. S.; Haigh, S. J.; Gholinia, A.; Mishchenko, A.; Lozada, M.; et al. Electronic properties of graphene encapsulated with different twodimensional atomic crystals. Nano Lett. 2014, 14, 3270−3276. (36) Chu, L.; Xue, Q.; Sun, J.; Xia, F.; Xing, W.; Xia, D.; Dong, M. Porous graphene sandwich/poly(vinylidene fluoride) composites with high dielectric properties. Compos. Sci. Technol. 2013, 86, 70−75. (37) Zhang, X.-J.; Wang, G.-S.; Wei, Y.-Z.; Guo, L.; Cao, M.-S. Polymer-composite with high dielectric constant and enhanced absorption properties based on graphene-CuS nanocomposites and polyvinylidene fluoride. J. Mater. Chem. A 2013, 1, 12115−12122. (38) Shi, B.; Su, Y.; Zhang, L.; Huang, M.; Li, X.; Zhao, S. Facilely prepared Fe3O4/nitrogen-doped graphene quantum dot hybrids as a robust nonenzymatic catalyst for visual discrimination of phenylenediamine isomers. Nanoscale 2016, 8, 10814−10822. (39) Wang, W.-J.; Xia, J.-M.; Hai, X.; Chen, M.-L.; Wang, J.-H. A hybrid of carbon dots with 4-chloro-7-nitro-2, 1, 3-benzoxadiazole for selective detection of p-phenylenediamine. Environ. Sci.: Nano 2017, 4, 1037−1044. (40) Wang, J.-Y.; Yang, S.-Y.; Huang, Y.-L.; Tien, H.-W.; Chin, W.-K.; Ma, C.-C. M. Preparation and properties of graphene oxide/polyimide composite films with low dielectric constant and ultrahigh strength via in situ polymerization. J. Mater. Chem. 2011, 21, 13569−13575. (41) Alzate-Carvajal, N.; Basiuk, E. V.; Meza-Laguna, V.; Puente-Lee, I.; Farías, M. H.; Bogdanchikova, N.; Basiuk, V. A. Solvent-free onestep covalent functionalization of graphene oxide and nanodiamond with amines. RSC Adv. 2016, 6, 113596−113610. (42) Liao, W.-H.; Yang, S.-Y.; Hsiao, S.-T.; Wang, Y.-S.; Li, S.-M.; Tien, H.-W.; Ma, C.-C. M.; Zeng, S.-J. A novel approach to prepare graphene oxide/soluble polyimide composite films with a low dielectric constant and high mechanical properties. RSC Adv. 2014, 4, 51117−51125. (43) Luong, N. D.; Hippi, U.; Korhonen, J. T.; Soininen, A. J.; Ruokolainen, J.; Johansson, L.-S.; Nam, J.-D.; Sinh, L. H.; Seppälä, J. Enhanced mechanical and electrical properties of polyimide film by graphene sheets via in situ polymerization. Polymer 2011, 52, 5237− 5242. (44) Chen, M.; Yin, J.; Jin, R.; Yao, L.; Su, B.; Lei, Q. Dielectric and mechanical properties and thermal stability of polyimide−graphene oxide composite films. Thin Solid Films 2015, 584, 232−237. (45) Yoonessi, M.; Shi, Y.; Scheiman, D. A.; Lebron-Colon, M.; Tigelaar, D. M.; Weiss, R. A.; Meador, M. A. Graphene polyimide nanocomposites; thermal, mechanical, and high-temperature shape memory effects. ACS Nano 2012, 6, 7644−7655. (46) Zhi, X.; Mao, Y.; Yu, Z.; Wen, S.; Li, Y.; Zhang, L.; Chan, T. W.; Liu, L. γ-Aminopropyl triethoxysilane functionalized graphene oxide for composites with high dielectric constant and low dielectric loss. Composites, Part A 2015, 76, 194−202. (47) Qin, Y.; Peng, Q.; Ding, Y.; Lin, Z.; Wang, C.; Li, Y.; Xu, F.; Li, J.; Yuan, Y.; He, X.; Li, Y. Lightweight, superelastic, and mechanically flexible graphene/polyimide nanocomposite foam for strain sensor application. ACS Nano 2015, 9, 8933−8941. (48) Ni, H.-j.; Liu, J.-g.; Wang, Z.-h.; Yang, S.-y. A review on colorless and optically transparent polyimide films: Chemistry, process and engineering applications. Ind. Eng. Chem. Res. 2015, 28, 16−27. (49) Vanherck, K.; Koeckelberghs, G.; Vankelecom, I. F. J. Crosslinking polyimides for membrane applications: A review. Prog. Polym. Sci. 2013, 38, 874−896. (50) Lyu, Y.; Xie, C.; Chechetka, S. A.; Miyako, E.; Pu, K. Semiconducting polymer nanobioconjugates for targeted photo6564


Article

The Journal of Physical Chemistry C thermal activation of neurons. J. Am. Chem. Soc. 2016, 138, 9049− 9052. (51) Dalton, A. B.; Collins, S.; Muñoz, E.; Razal, J. M.; Ebron, V. H.; Ferraris, J. P.; Coleman, J. N.; Kim, B. G.; Baughman, R. H. Supertough carbon-nanotube fibres. Nature 2003, 423, 703. (52) Jung, Y. C.; Shimamoto, D.; Muramatsu, H.; Kim, Y. A.; Hayashi, T.; Terrones, M.; Endo, M. Robust Conducting and transparent polymer composites using surface-modified and individualized double-walled carbon nanotubes. Adv. Mater. 2008, 20, 4509− 4512. (53) Yoonessi, M.; Lebrón-Colón, M.; Scheiman, D.; Meador, M. A. Carbon nanotube epoxy nanocomposites: the effects of interfacial modifications on the dynamic mechanical properties of the nanocomposites. ACS Appl. Mater. Interfaces 2014, 6, 16621−16630. (54) Jia, X.; Zhang, Q.; Zhao, M.-Q.; Xu, G.-H.; Huang, J.-Q.; Qian, W.; Lu, Y.; Wei, F. Dramatic enhancements in toughness of polyimide nanocomposite via long-CNT-induced long-range creep. J. Mater. Chem. 2012, 22, 7050−7056. (55) Li, Y.; Shimizu, H. Toward a stretchable, elastic, and electrically conductive nanocomposite: morphology and properties of poly[styrene-b-(ethylene-co-butylene)-b-styrene]/multiwalled carbon nanotube composites fabricated by high-shear processing. Macromolecules 2009, 42, 2587−2593. (56) Cui, L.; Lu, X.; Chao, D.; Liu, H.; Li, Y.; Wang, C. Graphenebased composite materials with high dielectric permittivity via an in situ reduction method. Phys. Status Solidi A 2011, 208, 459−461. (57) Volksen, W.; Miller, R. D.; Dubois, G. Low dielectric constant materials. Chem. Rev. 2010, 110, 56−110. (58) Dang, Z.-M.; Wang, L.; Yin, Y.; Zhang, Q.; Lei, Q.-Q. Giant dielectric permittivities in functionalized carbon-nanotube/electroactive-polymer nanocomposites. Adv. Mater. 2007, 19, 852−857.

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