Enhanced Dielectric Constant, Ultralow Dielectric Loss, and High

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C: Energy Conversion and Storage; Energy and Charge Transport

Enhanced Dielectric Constant, Ultralow Dielectric Loss and High Strength Imide-Functionalized Graphene Oxide/Hyperbranched Polyimide Nanocomposites Asma Iqbal, Seok Hwan Lee, Humaira Masood Siddiqi, O Ok Park, and Toheed Akhter J. Phys. Chem. C, Just Accepted Manuscript • DOI: 10.1021/acs.jpcc.8b00493 • Publication Date (Web): 05 Mar 2018 Downloaded from http://pubs.acs.org on March 7, 2018

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The Journal of Physical Chemistry

Enhanced

Dielectric

Constant,

Ultralow

Dielectric Loss and High Strength ImideFunctionalized Graphene Oxide/Hyperbranched Polyimide Nanocomposites Asma Iqbala,b, Seok Hwan Leeb, Humaira Masood Siddiqia*, O OkParkb*, Toheed Akhterb,c* a

Department of Chemistry,Quaid-i-Azam University, Islamabad, Pakistan, 45320.

b

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. c

Department of Chemistry, School of Science, University of Management and Technology, C-II,

Johar Town, Lahore, Pakistan. ABSTRACT We achieved for the first time the formation of a charge transfer complex 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 aminefunctionalized GO via a step-by-step condensation and thermal imidization approach. This FGO exhibited excellent compatibility with hyperbranched PI due to formation of a charge transfer complex between two domains. In viscoelastic measurements the dynamic storage modulus and

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glass transition temperature of flexible PI-FGO nanocomposites increased linearly with increasing FGO contents. The synthesized nanocomposites revealed high mechanical properties with tensile strength as high as 1.122 GPa. Thermogravimetric analysis demonstrates that these nanocomposites films exhibit high thermal stability up to 550 °C. Remarkably, the dielectric constant increases up to 42.47 at 8 wt% FGO loading with dielectric loss as low as 0.0018 while maintaining the breakdown strength as high as 147.3 ± 4.5MV/m. INTRODUCTION Energy storage technologies are gaining importance rapidly with the development in solar and wind power production sector1-3. Among energy storage devices, capacitors can demonstrate high power density due to their fast charge and discharge ability4-6. Efficiency of capacitors directly depends on dielectric materials with high permittivity7. Previously, commonly used dielectrics are based on ceramics, however, their frangibility, high density, and poor plasticity hinders them to meet the requirements of many applications8. On the bases of flexibility, low dielectric loss, appropriate breakdown strength, polymers were found more suitable dielectric materials than inorganic ceramics8-9. However, development of polymer dielectrics that can function well at high temperature is a key issue. For instance, dielectrics based on biaxially oriented polypropylene, poly vinylidene fluoride etc. can work at a maximum temperature of 105 °C, after that these lose their utility10-11. 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 resistance12-16.


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However, PIs possess inherently low dielectric constant in the range of 2.8-5 while their dielectric loss varies in the range of 0.001-0.03; this low dielectric loss is highly desirable for practical applications8,

17-19

. Mostly two approaches are adapted for increasing the dielectric

constant of PIs. First one is adding high-k ceramic reinforcement like Al2O320-21, nano-TiO216, CaCu3-Ti4O1222-23, and BaTiO324-26 into the PI matrix. However, to improve dielectric constant significantly, a large amount of ceramic filler is needed which results in brittleness of final PI matrix and hindering the development of polyimide based dielectrics. The second approach is introducing conductive filler including graphene, CNTs, and carbon black etc. In our previous study we reported γ-ray irradiated MWCNTs-PI nanocomposites with a higher dielectric constant of 36.25, however, their dielectric loss also remained higher due to leakage of current because of direct connection between conductive CNTs at higher filler loading13. Recently graphene is gaining importance as conductive filler due to its remarkable mechanical, thermal and electrical properties27-30. It has been incorporated in various polymer matrices including polyaniline31-32, carboxylated nitrile rubber33-35, poly vinylidene fluoride36-39 for various applications. Similarly, various research groups have studied the dispersion of graphene oxide (GO) in PI matrix to improve dielectric constant and mechanical properties40-44. For the efficient reinforcement of polymer matrices with graphene two key factors must be considered; (i) uniform dispersion of graphene into the polymer matrix (ii) and force of attraction between polymer and graphene. Modification of GO by attaching various functional groups to its surface i.e. carboxyl groups, amino groups etc has been the most exploited approach to circumvent these two challenges1,

8, 45-47

. However, in all the previously published reports improvement in

dielectric constant was not significant, also, dielectric loss increased tremendously with increase


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in dielectric constant which can restrict the practical application of these PI-GO nanocomposites as dielectric materials. In all these studies, PI matrix consists mostly of linear PI chains and functional groups attached to modified GO were mostly lacking structural similarity with PI leading to weak interfacial interaction and difficulty in efficient dispersion of GO in PI matrix at a higher loading of filler1, 44. Consequently, GO agglomerated developing direct connection with each other which can cause increase in dielectric loss. It is well established that the high mechanical and thermal properties of PIs are due to formation of a charge transfer complex (CTC) between adjacent PI chains leading to strong forces of attraction between them48-49. In the present study, we first time report the enhancement of compatibility of GO with hyperbranched PI via 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 was dispersed in 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 all previously reported PI-GO nanocomposites. Our aim of synthesis of hyperbranched PI instead of linear chain PI was to increase density of the CTC. This hyperbranched PI ensured the formation of 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 dielectric loss to go to higher value. To best of our knowledge, this is the first report of hyperbranched PI-GO nanocomposites with improved interfacial interaction by CTC formation between matrix and filler which afforded high dielectric constant, low dielectric loss, outstanding temperature resistance and mechanical properties.


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EXPERIMENTAL Materials and Measurements (3-Aminopropyl) trimethoxysilane and graphene oxide was purchased from Sigma Aldrich. 4,4oxidianiline (ODA) and pyromelliticdianhydride (PMDA) were supplied by Sinopharm Chemical Reagent Co., Ltd. N,N-dimethylformamide (DMF) 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. Instrumentation The melting points were determined using Gallenkamp (Sanyo) model MPDBM 3.5 with digital thermometer. FTIR spectra were recorded on Bruker α-alpha-P model using ATR method. 1HNMR and

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C-NMR spectral analysis were done on bruker (300 MHZ,) spectrophotometers

respectively in DMSO-d6 using TMS as an internal standard reference. The elemental analysis was carried out using CHNS analyzer, thermoscientific (Flash 2000series). The morphology of surface of nanocomposite was observed by field emission scanning electron microscopy (FESEM, S4800 Hitachi Japan) with an accelerating voltage of 10 kV. All samples were sputtered with platinum prior to 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 oC/min and a scanning range from room temperature to 800 oC. X-ray photoelectron spectroscopy (XPS) was conducted on a 60 (GENESIS EDAX, US) with Al Kα radiation (hm = 1486.4 eV). Viscoelastic behavior of nanocomposite films was investigated using TA Q800 DMA device in tension mode in


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temperature range of 30-500 °C. Mechanical tests were conducted on 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 hours before analysis. The dielectric properties measurements were carried out on CONCEBT 40 broadband dielectric spectroscope (Novocontrol Technologies, KG, Germany). The breakdown strength was measured using a dc dielectric strength tester with 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. Monomer Synthesis Synthesis of 4,4'-oxybis(N,N-bis(4-nitrophenyl)aniline) (1)The precursor of tetra-amine, i.e tetra-nitro compound, was synthesized by the cesium fluoride mediated N-arylation of 4-fluoro1-nitrobenzene by 4,4'-oxydianiline as shown in Figure 1. For this purpose, 4-fluoro-1nitrobenzene (42 mmol), cesium fluoride (46 mmol), 4,4'-oxydianiline (20 mmol), and DMSO (50 mL) were taken in a two necked round bottom flask equipped with a condenser and magnetic stirring bar. The reaction mixture was heated under inert atmosphere of nitrogen at 120 °C for 24 hours. The consumption of reactants was analyzed by thin layer chromatograph (n-Hexane: EtOAc, 3:1). After the completion of reaction, mixture was allowed to cool to room temperature. Then 150 mL of distilled water was added resulting in the precipitation of tetra-nitro compound which was filtered off, washed several times with distilled water, and purified by column chromatography. Yield: 90%. m.p: 182 oC.Rf= 0.54 (n-Hexane: EtOAc, 3:1). FT-IR (ῡ/cm-1): 1309 (C-N of ter amine stretch), 1570, 1337 (-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


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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 for C36 H24N6O(%): Calcd C= 63.16, H= 3.53, N= 12.28, Found C= 62.07, H= 3.65, N=11.78 Synthesis of N1,N1'-(4,4'-oxybis(4,1-phenylene))bis(N1-(4-aminophenyl)benzene-1,4-diamine) (2) For reduction of tetra-nitro compound into tetra-amine monomer, a 250 mL three neck round bottom flask, equipped with a stirring bar and reflux condenser, was charged with (1) (2.92 mmol), 10% Pd/C (0.2 g), and ethanol (70mL) and heated at reflux temperature. Hydrazine hydrate (9 mL) was added drop-wise into this boiling solution/suspension and the reaction mixture was refluxed for 24 hours. The progress of reaction was monitored by thin layer chromatography (n-Hexane: EtOAc, 1:1). After complete conversion of tetra-nitro compound into tetra-amino compound, Pd/C was filtered off. The solvent was evaporated from filtrate under reduced pressure, and the obtained solid product was recrystallized from ethanol to afford light pink solid. Yield: 65%. Mp. 145 oC, Rf = 0.31 (n-hexane: EtOAc, 2:1), FT-IR (ῡ/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 oC) (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= 9Hz, 4H, Ph-H,7,7’), 4.9 (s, NH2, 8H).

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C-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 for C36 H32 N6 O (%).: Calcd C= 76.57, H= 5.71, N= 14.88 , Found C= 76.07 , H= 5.85, N=14.38


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

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


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Figure 3

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C-NMR of N1-(4-aminophenyl)-N1-(4-(4-(bis(4-aminophenyl)amino)phenoxy)

phenyl)benzene 1,4-diamine (2) Surface Modification of Graphene Oxide Imide based functionalities were incorporated on the surface of graphene oxide aiming for formation of CTC with PI chains leading to uniform dispersion of FGO into PI matrix, as shown in Figure 4. For this purpose, a homogeneous dispersion of graphene oxide in anhydrous THF (0.1 mg/mL) was prepared by sonication for 30 mins at room temperature. To this dispersion (3aminopropyl)trimethoxysilane (10:1, APTMS:GO) was added and mixture was refluxed for 15 hours. Afterwards, the solution was filtered to recover APTMS-modified GO which was washed repeatedly with THF and dried in vacuum oven. The dried APTMS-modified GO was dispersed in anhydrous 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 hours. 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 addition of 2 (10:1),


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followed by addition of PMDA, and then by 2 to the surface of FGO resulting in preparation of FGO containing an imide surface modifier and amine end groups.

Figure 4 Surface functionalization of GO by attachment of imide based oligomers Synthesis of polyimide For synthesis of high molecular weight hyper branched polyimide (Figure 5), tetra-amine (0.5 mmol) was dissolved in anhydrous DMAc (3 mL) in a sealed glove box under inert environment. When tetra-amine was completely dissolved in 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. In order to compensate for any forfeited anhydride functional


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groups and to ensure anhydride as chain-end groups, a slight excess of PMDA (1 %) was added to the reaction mixture after 6 hours of stirring. Reaction mixture was stirred at about 0-5 °C for another 12 hours. Afterwards, poly(amic acid) (PAA) solution was transferred to teflon petri dish and film was casted by heating at 70 °C for 6 hours. Thermal imidization of PAA into PI was carried out by heating the PAA film at 100 °C, 200 °C, and 300 °C each for one hour. Preparation of PI/GO nanocomposites (PGC) For comparative study and evaluation of property benefits following two types of polyimide graphene nanocomposites (PGCs) were synthesized. a) PGC-FGO A suspension of surface functionalized graphene oxide (FGO) in DMAc was prepared by sonication for three hours. PAA solution in DMAc was added to this FGO suspension. The mixture was sonicated for another two hours. Afterwards, PAA/FGO suspension was transferred to teflon petri dish and heated at 70 °C for 12 hours to cast PAA/FGO films. Subsequently, PAA/FGO films were converted to PI/FGO nanocomposite films by thermal imidization. b) PGC-GO Here, the polyimide grapheme nanocomposites were prepared by the same method as discussed above except for incorporation of unmodified graphene oxide (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.


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Table 1: Nomenclature and composition of different types of PGCs PGCs

Sample codes

PI

FGO

GO

(wt %)

(wt %)

(wt %)

Pristine PI

PI

100

00

00

PGC-FGO

2-FGO

99

2

00

4-FGO

96

4

00

8-FGO

92

8

00

2-GO

99

00

2

4-GO

96

00

4

8-GO

92

00

8

PGC-GO

Figure 5 Synthesis of Polyimide


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Results and Discussions Monomer Synthesis Tetra-amine was synthesized in two step. The first step was cesium fluoride mediated Narylation of 4,4'-oxydianilinewith 4-fluoronitrobenzene, giving tetra-nitro compound. Second step was Pd/C catalyzed reduction of tetra-nitro compound to tetra-amine as shown in Fig 1. The structures of synthesized monomer and its precursor were evaluated by spectroscopic methods and elemental analysis. Firstly, 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 tetra-nitro compound was further confirmed by NMR spectroscopy. In the 1H NMR spectrum, the resonance peak for protons of amino group of 4,4'-oxydianiline was not observed. Also, a doublet at 8.19 ppm, which was assigned to eight protons located ortho to nitro groups, confirmed the successful synthesis of tetra-nitro compound. Similarly, in the FTIR spectrum of tetra-amine the characteristic absorption bands for -NH stretching of amino group appeared at 3249and 3190 cm-1. Furthermore, the absorption bands corresponding to nitro groups disappeared in the FTIR spectrum of tetra-amine monomer. Comprehensive evidence for complete conversion of tetra-nitro compound into tetra-amine was furnished by NMR spectroscopic analysis (Figure 2 & 3). In 1H NMR spectrum a singlet at 4.84 ppm was assigned to eight protons of amino group. The disappearance of resonance peak at 8.19 ppm (for eight protons at ortho position of tetra-nitro compound) also verified the completion of reduction of nitro into amine. Likewise, protons ortho to the amino groups resonated at 6.54ppm as a doublet in the 1H NMR spectrum of tetra-amine monomer.


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Synthesis of PI High molecular weight PI was synthesized by mixing stoichiometric amount of tetra-amine monomer with pyromelliticdianhydride. Firstly, the reaction between these two monomers resulted in the synthesis of poly(amic acid) (PAA), which was thermally imidized to convert PAA in to PI. Firstly, poly(amic acid) (PAA) was synthesized by the reaction between these two monomers, which was thermally imidized to convert PAA in to 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 to PI was evaluated by the disappearance of absorption band at 1650 cm-1 that was assigned to amide carbonyl group (Figure 6b). Furthermore, 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.


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Figure 6 FTIR of (a) [GO (A), GO modified with (3-Aminopropyl)trimethoxysilane (B), FGO (C), (b) [PAA and PI] Surface Functionalization of GO FTIR spectroscopy 0f FGO

To improve dispersion of GO in PI matrix and develop charge

transfer complex between GO and PI matrix, surface functionalized GO 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 tetra-amine 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. FTIR spectrum of (3-aminopropyl)trimethoxysilane 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 (3-


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aminopropyl)trimethoxysilane modified GO with PMDA produces amide carbonyl and carboxylic acid groups which were thermally imidized to imide functionalities. The terminal anhydride groups in this modified GO were reacted with tetra-amine. Thus, imide oligomers on the surface of GO were confirmed from 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. X-ray photoelectron spectroscopy (XPS) 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 the oxygen increased significantly as compared to unmodified graphene GO. This increment in the percentage of oxygen also provides a comprehensive evidence of presence of imide oligomers on the surface of FGO.


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Figure 7XPS spectra of (a) GO and (b) FGO, inset graph shows the N1s scan Morphological studies FE-SEM micrographs of the fractured surfaces of neat PI, PGC-FGO, and PGC-GO containing 4 wt% FGO and GO, respectively, are presented in Figure 8 (a-d). As it is obvious from Figure 8(a) that neat PI exhibits uniform and even surface which is commonly observed in conducting polymers50. On the other hand, Figure 8(b) reflects that the PGC-FGO film has a rough surface with some folds and wrinkles which are due densely packed crumpled FGO sheets creating wrinkles and bent up in the surface of film. These wrinkles and folds are even more prominent at high resolution as depicted in Figure 8 (c). FGO sheets in PGC-FGO films are well wrapped and laminated with thick polyimide layer that can be attributed to enhanced force of attraction via formation of charge transfer complex between FGO and polyimide matrix. This charge transfer complex with polyimide matrix prevented aggregation of FGO by isolating them from each other. Whereas in case of PGC-GO composite film, where unmodified GO does not form charge transfer complex with matrix, GO sheets are not well dispersed and are aggregated in PI matrix


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as obvious from Figure 8(d). Thus, it can be concluded that formation of charge transfer complex between PI matrix and FGO resulted in well dispersed filler state.

Figure 8SEM images (a, b, c) PGC-FGO and (d) PGC-GO Thermogravimetric analysis of PGCs Thermogravimetric analysis (TGA) was employed to investigate thermal stability of synthesized polyimide-GO nanocomposites (PGC- FGO, PGC-GO). The TGA thermograms are shown in Figure 9a whereas the thermal parameters are presented in Table S1. Thermal stability was evaluated in term of temperature at 5 % weight loss (Td), 10% weight loss (T10), and residual mass at 1000 °C (R1000). As 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. Thermal degradation temperature Td for neat PI is 483 °C, whereas those for PGCFGO and PGC-GO are 48 °C and 30 °C higher in comparison to neat PI, respectively.


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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 formation of charge transfer complex between PI matrix and FGO resulted in improved interfacial interaction between the two components imparting greater thermal stability to PGC-FGO system as compared to PGCGO.

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


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The delayed decomposition process of polymer matrix shows that formation of a charge transfer complex between two phases impeded diffusion of oxygen into PI matrix and reduced the degradation at elevated temperature. In case of PGC-FGO, FGO acted as an effective heat and mass transport barriers due to formation of charge transfer complex, and hence much suitable for preserving the PI matrix from thermal degradation. Viscoelastic properties Figure 9(b, c) presents the viscoelastic properties of synthesized nanocomposites reinforced with (2-8 wt %) FGO and GO. A charge transfer complex was created between polyimide 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 9 (b), whereas, Table S2 displays Tg and E’ values at room temperature and at Tg also. As it is obvious from Figure 9(b) that addition of FGO improved the values of E’ for PGC-FGO as well as for PGC-GO as compared to 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 same temperature, whereas, this value was 4.243 GPa for PGC-GO at the same temperature and wt% loading of un-functionalized GO. The effect of formation of charge transfer complex between FGO and PI matrix is more prominent from the E’ values at Tg. In rubbery plateau i.e. the region above the 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 while the 8-GO has E’ of 0.819 GPa. Thus, it can be inferred that significantly enhanced forces of attraction between FGO and PI matrix via charge transfer complex restricted the segmental motion of polymer chains, resultantly, improving the E’profile both in rubbery as well as in


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glassy state. Similarly, the effect of enhanced forces of attraction between FGO and PI matrix through formation of charge transfer complex 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 9 (c). It can be observed from Figure 9 (c) that in case of PGC-FGO the curves became more broadened with the addition of FGO and their maxima shift to higher temperature consistently. However, in case of PGC-GO, although Tan δ curves shift to higher temperature as compared to pristine PI, however, they remain almost unchanged with increasing content of GO. In case of PGC-FGO, as FGO contents increase from 0-8 wt%, Tg move 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 presence of charge transfer complex between FGO and PI matrix which hindered the chain segmental motions. Due to 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 the Tg to higher value. Mechanical properties Typical stress-strain curves of neat PI and PGC-FGO nanocomposites are presented in Figure 9(d). The interfacial interactions between filler and matrix, and hence, control of aggregation of filler is a key factor for improving the mechanical properties of nanocomposites51-52. As clear from Figure 9(d) that with increase in FGO contents, tensile strength linearly improved from 91.88 MPa for pure PI to 122.47 MPa for PGC-FGO8. The consistent increase in mechanical properties of nanocomposites on addition of FGO can be attributed to strong adhesion of FGO with PI matrix. As the strain moves to higher value, coiled polymer chains straighten to accommodate the deformation. This uncoiling of PI chains can create more room for charge


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transfer complex between filler and matrix, consequently external load is transferred to discontinuous high modulus FGO phase from continuous low modulus matrix phase via a shear stress at PI-FGO interface making nanocomposites more stiffer with higher load carrying ability in axial direction53-55. Dielectric properties The dielectric constant and dielectric loss of synthesized nanocomposites were investigated as a function of FGO/GO contents. The variations in dielectric constant along with variations in graphene oxide contents are shown in Figure 10a. The permittivity enhanced remarkably as the FGO contents reached and exceeded the percolation threshold as narrated by percolation theory1. The dielectric constant achieved maximum value of 42.47 at 8 wt% FGO loading. This value is higher by a factor of 5.5 than pristine PI matrix with dielectric constant of 7.72, confirming that FGO is homogeneously dispersed in PI matrix56. 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 polyimide graphene oxide nanocomposites19, 57. The dramatic increase of dielectric constant with increasing FGO contents is ascribed to the micro-capacitor network model13, 58. The contiguous FGO domains, which are completely coated with polymer matrix, are acting as electrodes and polymer bed present between them is acting as dielectric layer. In this way various microcapacitors are generated inside nanocomposites. At a higher filler contents like 8 wt%, microcapacitors increased and 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


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with PI matrix, which in turn can be attributed to remarkable compatibility of FGO with PI matrix due to formation of charge transfer complex between FGO and PI. As obvious from Figure 10a, in case of PGC-GO where pristine GO was dispersed in PI matrix, dielectric constant remained much lower as compared to PGC-FGO. Moreover, value of dielectric constant decreased at 8 wt% GO loading as compared to 4 wt%, indicating poor dispersion of GO in PI matrix, which is also indicated by increased dielectric loss at 8wt% GO content.

Figure 10 (a) Dielectric constant and dielectric loss of PGC-FGO and PGC-GO as a function of filler contents (b) breakdown strength of PGC-FGO as function of FGO contents Along with high dielectric constant, high breakdown strength is also an important factor for governing large energy storage of dielectric materials8. 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 PI matrix and PGC-FGO nanocomposites at room temperature as a function of FGO contents. As obvious from Figure 10b that pure PI has very high breakdown strength of 301.8 ± 5.1 MV m-1. Breakdown strength decreased to 172.4 ± 4.3 MV m-1 by addition of 2 wt% FGO, which is due to different electrical


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properties of FGO and PI matrix. With further increase in FGO contents, breakdown strength decreased further. However, at 8wt% loading of FGO the breakdown strength is 147.3 ± 4.5 MV m-1 which is fairly higher as compared to previously reported PI-GO nanocomposite1, 44. Conclusions In conclusion, we have successfully developed charge transfer complex 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 formation of charge transfer complex between PI chains and imide moieties on surface of FGO. PGC-FGO nanocomposites exhibit dielectric constant far higher than previously reported PI based nanocomposites. This goal was achieved by efficient dispersion of FGO and its strong binding with hyperbranched PI matrix. The extended branches in 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 PGCFGO 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. Supporting Information Available The Supporting Information contains Tables of data of thermogravimetric analysis

and

viscoelastic properties. Conflicts of interests


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The authors declare no conflicts of interests. AUTHOR INFORMATION Corresponding Author *Humaira Masood Siddiqi ([email protected], [email protected]) *O. Ok Park ([email protected]) *ToheedAkhter ([email protected]) Author Contributions The manuscript was written through contributions of all authors. All authors have given approval to the final version of the manuscript. 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 (20-3821/NRPU/R&D/HEC/14). REFERENCES 1.

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Enhanced Dielectric Constant, Ultralow Dielectric Loss, and High

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