Effect of functionalized graphene oxide with hyperbranched POSS polymer on mechanical and dielectric properties of cyanate ester composites

Abstract

In this paper, a hyperbranched silsesquioxane polymer grafted graphene oxide (HPP-GO) was fabricated successfully. Hexamethylenediamine was firstly grafted onto the graphene oxide sheets. Then chloropropyl silsesquioxane and hexamethylenediamine were repeatedly grafted onto the hexamethylenediamine coated graphene oxide, which was confirmed by X-ray photoelectron spectroscopy, Fourier-transform infrared spectra, X-ray diffraction and transmission electron spectroscopy. Subsequently, the HPP-GO was incorporated into dicyclopentadiene bisphenol dicyanate ester (DCPDCE) to prepare composites. The effects of HPP-GO on the mechanical, dynamic mechanical, dielectric and thermal properties of DCPDCE resin were investigated systematically. Results show that the appropriate content of HPP-GO can enhance the mechanical properties including impact and flexural strengths of DCPDCE resin. When HPP-GO content is 0.6 wt%, the corresponding composite exhibits higher storage modulus, glass transition temperature and onset degradation temperature values than pure DCPDCE resin. In addition, the dielectric constant and loss of HPP-GO/DCPDCE composites are decreased with the increasing content of HPP-GO. These attractive properties of HPP-GO/DCPDCE composites suggest that the method proposed herein is effective to prepare high-performance graphene oxide/polymer composites.

---

1. Introduction

Cyanate ester resins are important engineering thermosetting resins for high-temperature adhesives, electronic devices and composite matrices based on their good thermal resistance, very low dielectric constant and loss, low moisture absorption and small thermal expansion coefficient, etc.^1–3 However, the cured cyanate ester resins contain high cross-linked network of triazine groups, resulting in their brittleness. This has limited their use in many cutting-edge applications. Therefore, many attempts have been made to improve the toughness of cyanate ester resins, such as polymer blending,^4,5 fiber reinforcing^6,7 and nanofiller filling.^8,9 Among the various modifiers, nanofillers have attracted increasing attention because the properties of composites can be improved dramatically at low nanofillers loadings.^10–14 For example, Lin et al.¹⁵ prepared phenyl isocyanate functionalized GO to reinforce bisphenol A dicyanate ester resin, which resulted in a significant increase of 30% in the hardness of the composites with the addition of 1.25 wt% functionalized GO. John¹⁶ found the introduction of 4 vol% nanoclay can achieve an improvement of tensile strength by 94%. Dominguez et al.¹⁷ showed that only 1.0 wt% of multi-walled carbon nanotubes can lead to a 600% increase in storage modulus of cyanate ester composites at 200 °C.

Recently, graphene oxide (GO) has attracted increasing interest as a nanofiller for improving the properties of polymer composites because of its unique graphitized planar structure, good mechanical properties, unique thermal properties and low manufacturing costs, etc.^18–22 GO consists of a two-dimensional sheet of covalently bonded carbon atoms, bearing hydroxyl, carboxyl, and epoxide groups on its basal planes and edges, which provide versatile sites for further functionalization.^23,24 GO/polymer nanocomposites have been intensively studied in the recent years with the aim to obtain high-performance materials.^25–29 However, The reinforcing efficiency between GO and polymer is much lower than the anticipated value because of the poor dispersion of GO sheets in matrix and the insufficient polymer–GO interactions. Therefore, in order to prepare high performance GO/cyanate ester composites, it is of great value to design functionalized GO with a specific chemical structure that can form covalent bonds with cyanate ester matrix to achieve uniform dispersion of GO in cyanate ester matrix and improve the interfacial interaction between GO and the matrix.

Polyhedral oligomeric silsesquioxanes (POSS) are a family of nanoscale chemical structures that contain a silicon–oxygen core based on (SiO_1.5)_n and have each apex (silicon atom) connected to some organic groups.^30,31 It is this combination of an inorganic core covered with an organic shell at the molecular level that has led POSS structures to being labeled as organic–inorganic hybrid materials. POSS are attracting increased attention because they combine the advantages of both the organic moiety such as flexibility, solubility and high reactivity and inorganic moiety as rigidity, strength, and thermal stability.^32–35 On one hand, POSS can be incorporated into the polymer chains to modify the polymeric materials which eventually acquires enhanced mechanical, thermal and other physical properties compared to those of pure polymer systems.^36–38 On the other hand, POSS can also be used as surface modifiers to improve the dispersibility of nanoparticles. For example, Yadav et al.³⁹ grafted POSS onto carbon nanotubes using a click chemistry reaction. The resultant MWNT–POSS hybrid shows much better dispersion stability than pristine MWNTs in tetrahydrofuran. Gomathi et al.⁴⁰ adopted POSS to functionalize nanostructures of TiO₂, ZnO and Fe₂O₃, and the obtained POSS-functionalized nanostructures are highly soluble in many organic solvents, which enables the functionalized nanostructures to be used for preparing polymer-nanostructure composites. In our previous studies, we prepared a hybrid filler based on polyhedral oligomeric silsesquioxane and silica (POSS–SiO₂) to reinforce cyanate ester resin, and found that the dispersibility of SiO₂ in the matrix is dramatically improved with the introduction of POSS. Also, the mechanical, thermal and tribological properties of POSS–SiO₂/DCPDCE composites are markedly enhanced with appropriate content of POSS–SiO₂.⁴¹

Hyperbranched polymers are special kinds of polymers with highly branched structure. Comparing with their linear analogues, they possess considerably lower viscosity, better solubility, and high density of functional groups, so they can be utilized advantageously as surface modifiers.^42–47 In our previous studies, we have reported that hyperbranched polymer grafted GO is an effective approach to improve the dispersibility of GO in cyanate ester matrix.⁴⁸ This study further designed hyperbranched POSS polymer grafted GO to achieve uniform dispersion of functionalized GO in the cyanate ester matrix and construct flexible and covalently bonded interphase structures between functionalized GO and cyanate ester matrix. The hyperbranched POSS polymer grafted onto GO can not only prohibit the restacking of the GO sheets but also act as a “bridge connection” between GO and cyanate ester resin. Results show that the introduction of HPP-GO can significantly improve mechanical, dielectric and thermal properties of cyanate ester resin. The attractive features suggest that HPP-GO/cyanate ester composites have great potential in the fabrication of high-performance materials such as electronic packaging, printed circuit board, coating, etc., for cutting-edge industries.

2. Experimental

2.1. Reagents and materials

The GO nanosheets were produced from natural graphite flakes by the modified Hummer's method.⁴⁹ Chloropropyl POSS was synthesized according to the literature.⁵⁰ Dicyclopentadiene bisphenol dicyanate ester (DCPDCE) was purchased from Jiangdu Wuqiao Resin Plant (Jiangsu, China). Ethanol was purchased from Tianjin Tianli Chemical Reagents Co., Ltd. N′,N-Dimethylformamide was purchased form Tianjin Fuyu Fine Chemical Co. Ltd (Hebei, China). Hexamethylenediamine was purchased from Sinopharm Chemical Reagents Co., Ltd. 2-(7-Aza-1H-benzotriazole-1-yl)-1,1,3,3-tetramethyluronium hexafluorophosphate (HATU) was purchased from Shanghai Dibai Chemical Reagents Co., Ltd (Shanghai, China). Triethylamine was purchased from Guangdong Guanghua Sci-Tech Co., Ltd (Guangdong, China). Other reagents were all commercial products with analysis grades. All the chemicals were used as received.

2.2. Preparation of hexamethylenediamine grafted GO (N-GO)

GO (0.25 g) and DMF (60.0 ml) were added into a glass breaker, followed by ultrasonication with a power of 300 W for 30 min. Then the mixture was transferred to a 250 ml three-mouth flask holding a mechanical stirrer, reflux-condenser, constant-pressure funnel and thermometer. After that, HATU (20 mg) and hexamethylenediamine (8 g) were slowly added into the flask. The reaction mass was heated to 60 °C and maintained at that temperature for 6 h. The synthetic route was shown in Fig. 1. After the reaction, the mixture was filtrated and washed with copious amounts of ethanol. Finally, the resulting product was collected and dried in a vacuum oven at 60 °C for 12 h. The resulting product was abbreviated N-GO.

Fig. 1 The synthesis route of N-GO.

2.3. Preparation of hyperbranched grafted GO (HPP-GO)

Grafting of hyperbranched POSS polymer from the GO surfaces was achieved by repeating two reactions: (1) the reaction of chloropropyl POSS with N-GO and (2) the reaction of terminal groups of POSS with hexamethylenediamine.

The first step of the reaction was carried out as follows: In a 250 ml three-mouth round-bottom flask holding a mechanical stirrer, reflux-condenser, and thermometer, N-GO (0.8 g), chloropropyl POSS (0.45 g), DMF (60.0 ml) and triethylamine (5 ml) were added. Then the reaction mass was heated to 90 °C at which the reaction continued for 6 h. After the reaction, the mixture was filtered and washed with copious amounts of ethanol.

The second reaction step was carried out as follows: In a 250 ml three-mouth round-bottom flask holding a mechanical stirrer, reflux-condenser, and thermometer, GO obtained from the first step, hexamethylenediamine (0.6 g), DMF (60.0 ml) and triethylamine (5 ml) were added, and the mixture was stirred at 90 °C for 5 h. After the reaction, the mixture was filtered and washed with copious amounts of ethanol. Both the first and second reaction steps were repeated three times to grow the POSS polymer from the N-GO surfaces. The resulting product was abbreviated HPP-GO. The synthetic route of HPP-GO was shown in Fig. 2.

Fig. 2 The synthesis route of HPP-GO.

2.4. Preparation of HPP-GO/DCPDCE nanocomposites

The DCPDCE was heated to 100 °C in a glass beaker and kept at this constant temperature until melting. The appropriate amount of HPP-GO was then mixed with the melted DCPDCE using a mechanical high shear dispersion process. The mixture, consisting of prepolymer and HPP-GO, was heated to 100 °C in an oil bath and kept at this temperature for 15–20 min with stirring. Then the mixture was put into a preheated mold with release agent followed by degassing at 130 °C for about 1 h in a vacuum oven. After that, the mixture was cured and post-cured via the procedures of 160 °C/1 h + 180 °C/1 h + 200 °C/2 h + 220 °C/2 h and 240 °C/2 h, respectively. Finally, the mold was cooled to room temperature and demolded to get the samples of HPP-GO/DCPDCE systems.

The samples of pure DCPDCE resin were prepared in the same manner as above.

2.5. Characterization

Fourier transform infrared (FTIR) spectrum was recorded between 400 and 4000 cm⁻¹ with a resolution of 2 cm⁻¹ on a Nicolet FTIR 5700 spectrometer (USA).

X-ray photo-electron spectroscopy (XPS, Thermal Scientific K-Alpha XPS spectrometer) was used to investigate the surface elemental composition of GO, N-GO and HPP-GO. The analysis was performed under 1027 Torr vacuum with an Al Kα X-ray source using a power of 200 W.

X-ray diffraction (XRD) investigation was carried out using a X'Pert Pro MPD diffractometer (PANalytical, Holland) with Cu Kα radiation (λ = 0.154178 nm). The tube voltage was 36 kV, and the current was 20 mA. Scans were taken over the 2θ range of 5 to 85° with the scanning rate of 0.02° s⁻¹.

Transmission electron microscopic (TEM) images were obtained by a HITACHIS-600 (Japan) transmission electron microscope operating at 200 kV.

Impact strength was determined according to GB/T 2571-1995. The sample dimension was (80 ± 0.02) × (10 ± 0.02) × (4 ± 0.02) mm³.

Flexural strength was measured according to GB/T2570-1995. The sample dimension was (80 ± 0.02) × (15 ± 0.02) × (4 ± 0.02) mm³.

Scanning electron micrographs (SEM) were performed on a HITACHIS-570 instrument. For SEM samples preparation, the fracture surface of the specimens was sputtered with a thin layer (about 10 nm) of gold by vapor deposition on a stainless steel stub using a vacuum sputter coater.

Dynamic Mechanical Analysis (DMA) scans were performed using DMA/SDTA861e apparatus. DMA tests were carried out from 25 °C to 350 °C with a heating rate of 3 °C min⁻¹ at 1 Hz. The sample dimension was (45 ± 0.02) × (6 ± 0.02) × (3 ± 0.02) mm³.

Thermal gravimetrical analysis (TGA) tests were performed by using Perkin Elmer TGA-7 (USA) at a heating rate of 10 °C min⁻¹ in a nitrogen atmosphere.

The dielectric constant and loss factor were measured by a high frequency QBG-3 Gauger and a S914 dielectric loss test set (China) at the frequency range from 10 MHz to 60 MHz. The sample dimension was (25 ± 0.02) × (25 ± 0.02) × (3 ± 0.02) mm³.

3. Results and discussions

3.1. Characteristics of HPP-GO

3.1.1 XPS analysis. XPS was used to detect the surface composition and functional groups of GO, N-GO and HPP-GO based on the chemical shift observations. The XPS survey spectra of GO, N-GO and HPP-GO are presented in Fig. 3. GO exhibits C 1s and O 1s peaks. For the sample of N-GO, new slight reflections are found at 405 eV besides the signals of C 1s and O 1s, corresponding to N 1s. As shown in Fig. 3c, HPP-GO shows significant increased amount of N 1s comparing to that of N-GO, which is attributed to the grafted hexamethylenediamine. Also, HPP-GO exhibits additional Si 2s, Si 2p and Cl 1s peaks compared with N-GO, which are attributed to the composition of the grafted hyperbranched POSS polymer on GO, confirming the success of the modification.

Fig. 3 XPS survey spectra of GO (a), N-GO (b) and HPP-GO (c).

To provide further evidence on the grafting of hyperbranched POSS polymer on the GO surface, high resolution C 1s spectra of GO, N-GO and HPP-GO spectra are shown in Fig. 4. The C 1s XPS spectrum of GO shows a considerable degree of oxidation with four peaks appearing at 284.6, 286.5, 287.6 and 288.6 eV, which could be ascribed to the non-oxygenated ring C atoms, the C atoms in hydroxyl groups, the C atoms in epoxy groups, and C atoms in carboxyl groups, respectively. Compared with GO, the C 1s XPS spectrum of N-GO exhibits a new peak at 285.3 eV, originating from the C–N bonds. After the grafting of the hyperbranched POSS polymer on the surface of N-GO, an additional peak is found at 283.9 eV, which is ascribed to C–Si bonds. These results indicate that the hyperbranched POSS polymer attaches to the surfaces of GO through chemical reaction.

Fig. 4 High resolution C 1s spectra of GO (a), N-GO (b) and HPP-GO (c).

3.1.2 FTIR analysis. Further evidence on the successful grafting of the hyperbranched POSS polymer onto the GO can be demonstrated by FTIR spectra, as shown in Fig. 5. Fig. 5a shows the FTIR spectrum of the GO. The absorption peaks of GO at 1626 cm⁻¹, 1055 cm⁻¹ and 1737.79 cm⁻¹ are ascribed to C=C, C–O and C=O vibrations, respectively. In the case of hexamethylenediamine grafted GO, as seen in Fig. 5b, the peaks at around 3434 cm⁻¹ and 1632 cm⁻¹, which are assigned to –OH and C=C, respectively, continue to be observed. Additional new peaks at 2923 cm⁻¹ and 2857 cm⁻¹ are generated by the stretching vibrations of methyl groups and methylene groups, respectively. After the grafting of the hyperbranched POSS polymer on the surface of N-GO, a obvious change can be observed in the FTIR spectrum as shown in Fig. 5c, a strong absorption peak is found at 1119 cm⁻¹, which is ascribed to Si–O–Si groups in POSS, and the peak at 1035 cm⁻¹ is corresponding to the Si–O–C vibrations. These results also provide an evidence of the successful covalent grafting of the hyperbranched POSS polymer onto the GO through the chemical reaction.

Fig. 5 FTIR spectra of GO (a), N-GO (b) and HPP-GO (c).

3.1.3 XRD analysis. The GO, N-GO and HPP-GO structures were further investigated by X-ray diffraction (XRD), and spectra are shown in Fig. 6. The XRD spectrum of GO exhibits a sharp diffraction peak at 2θ = 11.36°, corresponding to a d-spacing of 0.78 nm. Following modification with hexamethylenediamine, the diffraction peak of N-GO significantly shifted to a smaller angle 2θ = 8.9° and weakened intensity compared to that of GO, revealing that the hexamethylenediamine was inserted in the GO sheets to further increase the interlayer distances and structural heterogeneity. It is interesting to note that the sharp diffraction peak disappears in the XRD profile of HPP-GO. This result may be attributed to the fact that the presence of long molecular chains on the GO surface can restrain the restacking of the planar GO nanosheets, inducing ordered graphitic stacking to well-disorder structures. In summary, analyzing the results from XPS, FT-IR, and XRD show that the graphene oxide sheets were successfully functionalized by hyperbranched POSS polymer.

Fig. 6 XRD spectra of GO (a), N-GO (b) and HPP-GO (c).

3.1.4 TEM analysis. To characterize the morphologies of GO and HPP-GO, TEM was used to observe the morphology of the structures (as shown in Fig. 7). The GO sheets (as shown in Fig. 7a) exhibited a smooth carpet like structure with some wrinkles. After surface functionalization (as shown in Fig. 7b), some dark regions on the surface of HPP-GO can be observed, which can be attributed to the hyperbranched POSS polymer layer attached onto the GO surface from both sides.

Fig. 7 TEM images of GO (a) and HPP-GO (b).

3.2. Mechanical properties of HPP-GO/DCPDCE systems

The impact and flexural strengths of HPP-GO/DCPDCE systems are shown in Fig. 8 and 9, respectively. The impact and flexural strength values of HPP-GO/DCPDCE systems are shown in Table S1 (see ESI†). Compared to the pure DCPDCE, the HPP-GO/DCPDCE systems show increased impact and flexural strengths. The impact and flexural strength are increase by 77% and 56% for the system with 0.6 wt% HPP-GO, which suggests that HPP-GO has very high reinforcing efficiency. The increases in mechanical properties can be understood from the following reasons. Firstly, the grafted hyperbranched POSS polymer on the surface of GO nanosheets can bring a large interlayer spacing, preventing the restacking of GO nanosheets, which results in the efficient load transfer from DCPDCE matrix to GO sheets. Secondly, the rigid POSS in the interface layer zone can efficiently distribute the stresses concentrated on the tip of the cracks. Thirdly, the flexible hyperbranched polymer chains can also provides a high ability for transferring the stresses from the DCPDCE matrix to HPP-GO sheets, preventing the crack propagation. Fourthly, the HPP-GO surface can provide abundant amine-functional groups for forming enhanced interfacial bonding through the reaction between amine groups and –NCO in HPP-GO. Therefore, the mechanical properties of HPP-GO/DCPDCE system are increased as the contents of HPP-GO from 0.0 wt% to 0.6 wt%. However, when the fillers content is high enough (>0.6 wt%), the flexural strength of the composites decrease with the increasing concentration of fillers. This occurs because further loading causes the GO sheets to stack together, reducing the improvement return in flexural strength.

Fig. 8 Relationship of impact strength on HPP-GO content for HPP-GO/DCPDCE systems.

Fig. 9 Relationship of flexural strength on HPP-GO content for HPP-GO/DCPDCE systems.

In order to further confirm the effect of HPP-GO on the toughness of DCPDCE resin, Fig. 10 gives SEM images of the fracture surfaces of HPP-GO/DCPDCE composites. The pure DCPDCE resin exhibits a brittle fracture surface as shown in Fig. 10a, while with the incorporation of HPP-GO to DCPDCE resin, the HPP-GO/DCPDCE systems exhibit relatively rougher fracture surfaces, which indicates that the addition of HPP-GO can toughen DCPDCE resin. For the 0.2 wt% HPP-GO/DCPDCE system as shown in Fig. 10b, the surface is coarse and some ductile sunken regions can be observed, which can absorb the energy of fracture and limit the propagation of the cracks into big cracks.⁵¹ When the HPP-GO content is to 0.6 wt%, as shown in Fig. 10c, the fracture surface of the composite is much rougher than those of pure DCPDCE and 0.2 wt% HPP-GO/DCPDCE system, and there exist large amount of ductile sunken areas, exhibiting a typical rough feature. When the concentration of HPP-GO is up to 0.8 wt%, as shown in Fig. 10d, the fracture surface is very rough and large amount of ductile sunken areas can also be observed. However, some agglomerates appear, which are considered as flaws and crack initiation sites. The presence of flaws in composite may decrease the mechanical properties of composite,⁵² while the sensitivity of impact strength and flexural strength to the flaws may be different, so in this study, the optimal HPP-GO contents for reaching the maximum flexural strength and impact strength is different. This phenomenon also has been reported by other ref. 53 and 54.

Fig. 10 SEM images of the fracture surfaces of the HPP-GO/DCPDCE systems ((a) 0.0 wt%, (b) 0.2 wt%, (c) 0.6 wt%, (d) 0.8 wt%).

3.3. Dynamic mechanical properties of HPP-GO/DCPDCE systems

Fig. 11 shows the overlay plots of storage modulus as a function of temperature for pure DCPDCE resin and 0.6 wt% HPP-GO/DCPDCE system. As can be observed, the HPP-GO/DCPDCE system have obvious higher storage modulus values than pure DCPDCE resin from 50 to 260 °C. This phenomenon can be explained by the following reasons. Firstly, the addition of HPP-GO in matrix introduces rigid POSS structure, and the massive units of the relatively rigid POSS tend to retard and restrict that segments' motions in DCPDCE. Secondly, GO grafted with hyperbranched POSS polymer can achieve good dispersion of the HPP-GO sheets in DCPDCE matrix, further forming a 3D network. The HPP-GO sheets can act as interlock points in the 3D network, which can effectively inhibit polymeric chains' entanglements and slips under heating conditions. Thirdly, the strong interfacial adhesion between HPP-GO and DCPDCE matric can also restrict the segments' motions, increasing the storage modulus of HPP-GO/DCPDCE composite.

Fig. 11 Overlay curves of storage modulus as a function of temperature for DCPDCE resin and 0.6 wt% HPP-GO/DCPDCE system.

Fig. 12 shows the overlay plots of loss tangent (Tan Delta) as a function of temperature for pure DCPDCE resin and 0.6 wt% HPP-GO/DCPDCE system. Since the tan δ peak occurs in a temperature range over which the polymer changes from a glassy state to an elastic state, hence the peak temperature of which is often taken as the glass transition temperature (T_g) of a polymer.⁵⁵ In this paper, the temperature corresponding to tan δ peak is considered as the T_g value of the materials. It is noted that the whole peak of HPP-GO/DCPDCE system shifts toward higher temperature by about 20 °C compared with pure DCPDCE resin. In fact, there is a negative influence coming from the flexible isoureas structure which is formed through the reaction between –NH₂ groups in HPP-GO and –OCN groups, but it is not the dominant role. The well dispersed HPP-GO nanosheets possess high steric hindrance, so the segmental motion of polymers in hybrids is restricted efficiently, leading to increased Tg values. In addition, the 0.6 wt% HPP-GO/DCPDCE system exhibits a single peak as that of DCPDCE resin, reflecting that 0.6 wt% HPP-GO is distributed uniformly in the matrix.

Fig. 12 Overlay curves of loss tangent as a function of temperature for DCPDCE resin and 0.6 wt% HPP-GO/DCPDCE system.

3.4. Dielectric properties of HPP-GO/DCPDCE systems

Fig. 13 and 14 shows overlay plots of dependence of dielectric constant and loss on frequency for pure DCPDCE resin and modified systems. It can be seen that the dielectric constants and losses of HPP-GO/DCPDCE systems are lower than those of DCPDCE resin, and the higher the HPP-GO loading is, the lower the dielectric constant. The decrease in dielectric constant and loss may be explained by the following reasons.

Fig. 13 Dielectric constants of DCPDCE resin and HPP-GO/DCPDCE systems versus frequency.

Fig. 14 Dielectric loss factors of DCPDCE resin and HPP-GO/DCPDCE systems versus frequency.

Firstly, The GO surface possesses a high density of oxygen-containing functional groups. The carbon–oxygen bonds were formed on the GO surface, and the sp²-hybridized carbon atoms in graphite were transferred to the sp³-hybridized carbon atoms in GO. This process reduces conjugation and confines π-electrons, so GO exhibits very low dielectric constant values.⁵⁶ The HPP-GO sheets may be more insulating than GO because the grafted hyperbranched POSS polymer possesses many unoccupied spaces, so the addition of HPP-GO will reduce the dielectric constant of DCPDCE matrix. Secondly, the POSS structure in HPP-GO is nanoporous, so the incorporation of POSS structure into DCPDCE resin is like introducing “air bubbles” into the matrix, while the dielectric constant of air is very low (about 1). This is also one reason that HPP-GO/DCPDCE system has lower dielectric constant and loss than DCPDCE resin. Thirdly, a good adhesion between HPP-GO and DCPDCE matrix will reduce the interfacial polarization and restrict the mobility of chain segments, leading to reduced dielectric constant and loss.

3.5. Thermal properties of HPP-GO/DCPDCE systems

TG-DTG curves for pure DCPDCE and 0.6 wt% HPP-GO/DCPDCE system as a function of temperature at the heating rate of 10 °C min⁻¹ are displayed in Fig. 15. The onset degradation temperature (T_di) is defined as the temperature at which the weight loss is 5%. T_di is usually used to evaluate the thermal degradation and thermal stability of materials. Compared to that of pure DCPDCE resin, the T_di value of HPP-GO/DCPDCE system is increased by about 24 °C, suggesting that 0.6 wt% HPP-GO/DCPDCE composite has better thermal stability than pure DCPDCE resin. There are three reasons proposed for the improvement to the thermal properties of the DCPDCE composite. First, this significant improvement in thermal stability could be partial attributed to the so-called “tortuous path” effect of graphene sheets, which retards the permeation of heat and the escape of volatile degradation products.^57,58 Second, the HPP-GO nanosheets can act as physical interlock points in the cured resin, which can provide a sterically hindered environment and restrain the mobility of polymer chains. Thirdly, the amine groups on the surface of HPP-GO can react with –NCO groups, constructing a covalently bonded interphase structure between HPP-GO and DCPDCE resin, which is beneficial in restricting the movements of the polymer segments and thus increasing the energy consumption of polymer chains degradation.

Fig. 15 Overlay TGA and DTG curves of DCPDCE and 0.6 wt% HPP-GO/DCPDCE system.

4. Conclusions

In this study, a hyperbranched POSS polymer was grafted onto the GO surface. Then the functionalized GO was incorporated into DCPDCE resin to prepare HPP-GO/DCPDCE composite materials. On the basis of the results of mechanical properties, the HPP-GO/DCPDCE composite with 0.6 wt% HPP-GO loading exhibits an approximate 77% increase in impact strength and a 56% improvement in flexural strength compared to pure DCPDCE resin. Meanwhile, the addition of HPP-GO tends to increase the storage modulus and T_g value of the composite. Compared with pure DCPDCE resin, the HPP-GO/DCPDCE composites exhibit decreased dielectric constants and losses. TGA results show that 0.6% HPP-GO/DCPDCE composite also possesses better thermal stability than pure DCPDCE resin. Functionalized GO with hyperbranched POSS polymer presented herein will provide an effective strategy to improve the mechanical properties and dielectric properties in nonpolar polymers.

Acknowledgements

This work was financially supported by Research Fund for the Doctoral Program of Higher Education (20136102110049).

References

1. I. Hamerton and J. N. Hay, Polym. Int., 1998, 47, 465–473.
2. A. Ganesan and S. Muthusamy, Polym. Compos., 2009, 30, 782–790.
3. J. T. Reams, A. J. Guenthner, K. R. Lamison, V. Vij, L. M. Lubin and J. M. Mabry, ACS Appl. Mater. Interfaces, 2012, 4, 527–535.
4. C. Marieta, M. D. Rio, I. Harismendy and I. Mondragon, Eur. Polym. J., 2000, 36, 1445–1454.
5. P. Ren, G. Liang and Z. Zhang, Polym. Compos., 2006, 27, 402–409.
6. P. Badrinarayanan, M. K. Rogalski and M. R. Kessler, ACS Appl. Mater. Interfaces, 2011, 4, 510–517.
7. J. Li, Z. Wu, C. Huang and L. Li, Compos. Sci. Technol., 2014, 104, 81–88.
8. K. Liang, H. Toghiani and C. U. Pittman, J. Inorg. Organomet. Polym. Mater., 2011, 21, 128–142.
9. X. Zhang, G. Liang, J. Chang, A. Gu, L. Yuan and W. Zhang, Carbon, 2012, 50, 4995–5007.
10. Y. Liu, C. Zhang, H. Tan, X. Yan, H. Wei and Z. Guo, Macromol. Mater. Eng., 2014, 300, 64–69.
11. Q. He, T. Yuan, X. Yan, D. Ding, Q. Wang, Z. Luo, T. D. Shen, S. Wei, D. Cao and Z. Guo, Macromol. Chem. Phys., 2014, 215, 327–340.
12. J. Zhu, RSC Adv., 2013, 3, 22790–22824.
13. J. Zhu, J. Mater. Chem., 2010, 20, 4937–4948.
14. Z. Guo, T. Pereira, O. Choi, Y. Wang and H. T. Hahn, J. Mater. Chem., 2006, 16, 2800–2808.
15. Q. Lin, L. Qu, Q. Lü and C. Fang, Polym. Test., 2013, 32, 330–337.
16. B. John, C. P. R. Nair and K. N. Ninan, Mater. Sci. Eng., A, 2010, 527, 5435–5443.
17. D. D. Dominguez, L. Matthew and T. M. Keller, Macromol. Chem. Phys., 2009, 210, 1709–1716.
18. Y. Shao, J. Wang, M. Engelhard, C. Wang and Y. Lin, J. Mater. Chem., 2010, 20, 743–748.
19. J. Shen, B. Yan, T. Li, Y. Long, N. Li and M. Ye, Composites, Part A, 2012, 43, 1476–1481.
20. H. M. Kim, J. K. Lee and H. S. Lee, Thin Solid Films, 2011, 519, 7766–7771.
21. M. Han, J. Yun, H. I. Kim and Y. S. Lee, J. Ind. Eng. Chem., 2012, 18, 752–756.
22. J. N. Gavgani, H. Adelnia and M. M. Gudarzi, J. Mater. Sci., 2014, 49, 243–254.
23. S. Stankovich, R. D. Piner, S. B. T. Nguyen and R. S. Ruoff, Carbon, 2006, 44, 3342–3347.
24. Y. W. Zhu, S. Murali, W. W. Cai, X. S. Li, J. W. Suk, J. R. Potts and R. S. Ruoff, Adv. Mater., 2010, 22, 3906–3924.
25. H. Yang, C. Shan, F. Li, Q. Zhang, D. Han and N. Li, J. Mater. Chem., 2009, 19, 8856–8860.
26. S. N. Tripathi, P. Saini and D. Gupta, J. Mater. Sci., 2013, 48, 6223–6232.
27. W. Li, B. Zhou, M. Wang, Z. Li and R. Ren, J. Mater. Sci., 2015, 50, 5402–5410.
28. M. Zeng, J. Wang, R. Li, J. Liu, W. Chen, Q. Xu and Y. Gu, Polymer, 2013, 54, 3107–3116.
29. X. Wang, J. Jin and M. Song, Eur. Polym. J., 2012, 48, 1034–1041.
30. J. J. Deng, J. T. Polidan, J. R. Hottle, C. E. Farmer-Creely, B. D. Viers and A. R. Esker, J. Am. Chem. Soc., 2002, 124, 15194–15195.
31. D. Gnanasekaran, K. Madhavan and B. S. R. Reddy, J. Sci. Ind. Res., 2009, 68, 437–464.
32. O. Monticelli, A. Fina, A. Ullah and P. Waghmare, Macromolecules, 2009, 42, 6614–6623.
33. D. Jiang, L. Xing, L. Liu, X. Yan, J. Guo, X. Zhang, Q. Zhang, Z. Wu, F. Zhao, Y. Huang, S. Wei and Z. Guo, J. Mater. Chem. A, 2014, 2, 18293–18303.
34. D. Jiang, L. Liu, J. Long, L. Xing, Y. Huang, Z. Wu, X. Yan and Z. Guo, Compos. Sci. Technol., 2014, 100, 158–165.
35. D. Jiang, L. Xing, L. Liu, S. Sun, Q. Zhang, Z. Wu, X. Yan, J. Guo, Y. Huang and Z. Guo, Compos. Sci. Technol., 2015, 117, 168–175.
36. N. Yong, S. Zheng, K. Nie, N. Yong and S. Zheng, Polymer, 2004, 45, 5557–5568.
37. K. Liang, H. Toghiani, G. Li and C. U. Pittman, J. Polym. Sci., Part A: Polym. Chem., 2005, 43, 3887–3898.
38. E. T. Kopesky, G. H. Mckinley and R. E. Cohen, Polymer, 2006, 47, 299–309.
39. S. K. Yadav, S. S. Mahapatra, H. J. Yoo and J. W. Cho, Nanoscale Res. Lett., 2011, 6, 122–128.
40. A. Gomathi, K. Gopalakrishnan and C. N. R. Rao, Mater. Res. Bull., 2010, 45, 1894–1898.
41. M. M. Zhang, H. X. Yan, C. Liu and J. P. Zhang, Polym. Compos., 2015, 36, 1840–1848.
42. C. Gao and D. Yan, Prog. Polym. Sci., 2004, 29, 183–275.
43. M. E. Mackay, G. Carmezini, B. B. Sauer and W. Kampert, Langmuir, 2001, 17, 1708–1712.
44. C. Schüll and H. Frey, Polymer, 2013, 54, 5443–5455.
45. J. K. Paulasaari and W. P. Weber, Macromolecules, 2000, 33, 2005–2010.
46. V. Geiser, Y. Leterrier and J. A. E. Manson, Macromolecules, 2010, 43, 7705–7712.
47. D. Foix, X. Ramis, F. Ferrando and A. Serra, Polym. Int., 2012, 61, 727–734.
48. M. M. Zhang, H. X. Yan, X. W. Yang and C. Liu, RSC Adv., 2014, 4, 45930–45938.
49. W. S. Hummers and R. E. Offeman, J. Am. Chem. Soc., 1958, 80, 1339.
50. U. Dittmar, B. J. Hendan, U. Flörke and H. C. Marsmann, J. Organomet. Chem., 1995, 489, 185–194.
51. X. Zhang, O. Alloul, Q. He, J. Zhu, M. J. Verde, Y. Li, S. Wei and Z. Guo, Polymer, 2013, 54, 3594–3604.
52. X. Zhang, X. Yan, J. Guo, Z. Liu, D. Jiang, Q. He, H. Wei, H. Gu, H. A. Colorado, X. Zhang, S. Wei and Z. Guo, J. Mater. Chem. C, 2015, 3, 162–176.
53. G. Khanbabaei, J. Aalaie, A. Rahmatpour, A. Khoshniyat and M. A. Gharabadian, J. Macromol. Sci., Part B: Phys., 2007, 46, 975–986.
54. C. Kaynak and C. C. Tasan, Eur. Polym. J., 2006, 42, 1908–1921.
55. W. Ling, A. Gu, G. Liang, L. Yuan and J. Liu, Polym. Adv. Technol., 2009, 21, 365–370.
56. J. Y. Wang, S. Y. Yang, Y. L. Huang, H. W. Tien, W. K. Chin and C. C. M. Ma, J. Mater. Chem., 2011, 21, 13569–13575.
57. W. H. Liao, S. Y. Yang, S. T. Hsiao, Y. S. Wang, S. M. Li, C. C. M. Ma, H. W. Tien and S. J. Zeng, ACS Appl. Mater. Interfaces, 2014, 6, 15802–15812.
58. X. Wang, L. Song, H. Yang, W. Xing, B. Kandola and Y. Hu, J. Mater. Chem., 2012, 22, 22037–22043.

---

Footnote

† Electronic supplementary information (ESI) available. See DOI: 10.1039/c6ra01053a

---