A Rational Design of Hyperbranched Aromatic Polyamide

Article pubs.acs.org/JPCB

Achieving High Performance Electric Field Induced Strain: A Rational Design of Hyperbranched Aromatic Polyamide Functionalized Graphene−Polyurethane Dielectric Elastomer Composites Tian Chen,*,†,‡ Jinhao Qiu,*,†,§ Kongjun Zhu,†,§ Jinhuan Li,‡ Jingwen Wang,‡ Shuqin Li,‡ and Xiaoliang Wang∥,⊥ †

State Key Laboratory of Mechanics and Control of Mechanical Structures, Nanjing University of Aeronautics and Astronautics, Nanjing, 210016, China ‡ College of Materials Science and Technology, Nanjing University of Aeronautics and Astronautics, Nanjing, 210016, China § College of Aerospace Engineering, Nanjing University of Aeronautics and Astronautics, Nanjing, 210016, China ∥ Key Laboratory of High-Performance Polymers Materials and Technology of Ministry of Education, Nanjing University, Nanjing, 210093, China ⊥ School of Chemistry and Chemical Engineering, Nanjing University, Nanjing, 210093, China ABSTRACT: Dielectric elastomers have great potentials as flexible actuators in micro-electromechanical systems (MEMS) due to their large deformation, light weight, mechanical compliancy, and low cost. The low dielectric constant of these elastomers requires a rather high voltage electric field, which has greatly limited their applications. In this work, a diaphragm-type flexible microactuator comprising a hyperbranched aromatic polyamide functionalized graphene (HAPFG) filler embedded into the polyurethane (PU) dielectric elastomer matrix is described. The rational designed HAPFG sheets exhibits uniform dispersion in PU matrix and strong adhesion with the matrix by hydrogen-bond coupling. Consequently, the HAPFG−PU composites possess high dielectric performance and low loss modulus. The effect of hyperbranched aromatic polyamide functionalized graphene on high voltage electric field induced strain was experimentally investigated using the Fotonic sensor. The high electric field response of the composite was discussed by applying different kinds of alternating-current field. In addition, a comparison of the breakdown strength between the HAPFG−PU composite and the pure PU was carried out.

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Polyurethane was first reported to show an electric field induced contraction by Z. Ma et al. in 1994.25 Z. Ma et al. showed that polyurethane could produce thickness strains of 3% under 20 MV/m electric field by using metal electrodes of 20 nm thick gold. These values are greater than those obtained from piezoelectric ceramic and piezoelectric polymer. However, just like all existing DEs, polyurethanes have rather low dielectric constants. A high dielectric constant is desirable to obtain large electric field induced strain outputs and to reduce the driving voltage. The reported methods to increase the dielectric constant of DEs may be divided into three kinds: (1) introduction of organic dipole groups to the polymer matrix;21−24 (2) incorporation of high dielectric constant ceramic fillers;12,14−16,35 and (3) incorporation of conductive nanofillers.29,36−41,43 The main drawback of method 1 is the polar group significantly reduces the polymers’ dielectric

INTRODUCTION The emergence in the early 1990s of “electroactive polymers” (EAPs) that react to electrical stimulation with large changes in shape has inspired the creativity of many scientists and engineers and has widened the already broad spectrum of applications for polymer materials.1−4 Because of the low cost, the light weight of materials, and the inherent compliance ability of polymers to be tailored to particular applications, EAPs can be utilized in microactuators, sensors, microrobotics, micro air vehicles, and artificial muscle prosthetic device applications. In general, EAPs are classified as “ionic” or “electronic” with different modes of action.3−5 Ionic EAPs consist of materials such as polymer gels, conducting polymers, and ionic polymer−metal composites. Electronic EAPs consist of materials such as electrostrictive elastomers, ferroelectric polymers, and dielectric elastomers (DEs). Among them, DEs offer excellent performance in an unusually broad range of applications covering actuators and sensors.4−6 Commercial acrylic adhesives,7−11 silicones,12−24 and polyurethanes25−42 are the most studied DEs. © 2015 American Chemical Society

Received: September 2, 2014 Revised: February 20, 2015 Published: March 5, 2015 4521

DOI: 10.1021/jp508864b J. Phys. Chem. B 2015, 119, 4521−4530

The Journal of Physical Chemistry B

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EXPERIMENTAL SECTION Graphene Sheets Functionalized with Hyperbranched Aromatic Polyamide (HAPFG). Graphene oxide (GO) was prepared from natural graphite via the modified Hummers method.58 The synthetic strategy of HAPFG is depicted in Scheme 1. The preparation procedure of HAPFG was based on

strength and increases sensitivity to moisture. Method 2 requires the use of a large amount of ceramics. The agglomeration of ceramics at high loading contents is inevitable, and reduces the uniformity of the films and enhances the brittleness. Method 3 can achieve a high dielectric constant at a relatively low filler loading fraction. Q. Zhang et al. discovered that the embedment of carbon nanotubes into an electroactive polymer could increase the dielectric constant and electric field induced strain of the polymer.43 B. Guiffard et al. adopted amorphous carbon black nanopowder to enhance the electromechanical performance of polyurethane.36 The highest strain amplitude value observed was obtained for 1 wt % carbon−polyurethane composite (Sd = −7.4% at E = 17.8 MV/m). By incorporation of carbon-coated SiC nanowires into PU matrix, they revealed that a loading of 0.5 wt % SiC@C presented the maximal observed thickness strain of about 18% at 18 MV/m (20 μm thick film).37 K. Yuse et al. showed that 24% thickness strain at 15 MV/m could be achieved by using 1.7 wt % nanocarbon black ink in micellar form.38 The conductive carbon nanofillers are considered to be an effective way to generate large-strain and low-powered electroactive polymer. Recently, graphene has attracted increasingly paramount interest in the field of polymer nanocomposites.44−50 Owing to high electrical conductivity, high surface-to-volume ratio, and exceptional in-plane functional and mechanical properties, graphene can induce a dramatical improvement in electrical, thermal, optoelectronic, and mechanical properties of the resulting graphene−polymer composites at very low filler fractions.51−63 In our previous study, we demonstrated that the thickness strain of polyurethane dielectric elastomer could be enhanced by incorporating PMMA functionalized graphene nanoplates into the polymer matrix.64 A fine control of dispersion of the graphene remains the major problem for the effective reinforcement of the mechanical properties and adding functional properties to the graphene-based composites. Surface functionalization of graphene has been established as an efficient means to achieve superior dispersion and enhance the quality of the interface between the graphene and polymer host.48,49 Organic groups such as SO3H,65 alkyl amine,66 and organic isocyanate groups67 were applied to covalently couple to the surface of graphene. Even if great progress has been achieved in increasing the physical and chemical properties of host polymers by functionalization of graphene fillers, the dielectric properties of functionalized graphene−polymer composites still fall below the expected values because of a lack of sufficient functionalized groups. Hyperbranched polymers not only have a highly branched and nonentangled architecture but also have lots of terminal groups. The large number of terminal groups can reinforce the interactivity with the host polymer matrix by hydrogen bonding or covalent linkage.68 Furthermore, hyperbranched polymers show lower solution viscosities and higher solubility in comparison with their linear analogues.69 For these reasons, hyperbranched aromatic polyamide functionalized graphene was designed and adopted to increase the dielectric property and electric field induced strain of polyurethane dielectric elastomer in this work. To our best knowledge, it is the first time that the electromechanical actuation performances of hyperbranched aromatic polyamide functionalized graphene−polyurethane dielectric elastomer intelligent materials have been investigated in detail.

Scheme 1. Scheme for the Synthetic Strategy of HAPFG

the previous report with some modifications.59 In a typical experiment, GO (100, 150, and 200 mg) was first dispersed in DMF (50 mL) containing ethylenedimine (EDA, 1, 2, and 4 mL), dicyclohexylcarbodiimide (DCC, 0.5, 1.0, and 2.0 g), and 4-dimethylaminopyridine (DMAP, 0.1, 0.4, and 0.5 g); afterward, the mixture was heated to 60 °C under stirring. After reaction for 8 h, the EDA-modified GO was obtained. Subsequently, the EDA-modified GO was dispersed in Nmethyl pyrrolidone (NMP, 20, 30, and 50 mL) solution containing DMAP (0.1, 0.3, and 0.5 g), pyridine (1 mL), and triphenyl phosphite (TPP, 2 mL). LiCl (0.1 g) and 3,5diaminobenzoic acid (3,5-DABA, 0.5, 1.0, and 3.0 g) were then added to the dispersion, and the mixture was heated to 120 °C. After reaction for 10 h, the mixture was heated to 180 °C and kept for 15 h for reduction of GO. After that, the mixture was put into 200 mL of vigorously stirred methanol. The HAPFG-x (x = 1, 2, 3) was obtained by filtration, washing with DMF, and centrifugal separation. The sample code HAPFG-x (x = 1, 2, 3) means three different preparation processes that are performed under the same reacting conditions by using different reactant dosages. Reduced graphene oxide (RGO) was prepared for comparison by reduction of GO with hydrazine hydrate. Functionalized Graphene−Polyurethane Dielectric Elastomer Composites (HAPFG−PU). HAPFG−PU films were prepared using the solution cast method as described elsewhere.64 In a typical experiment, PU (A polyether-type thermoplastic polyurethane elastomer TPU58887, Estane) granules were completely dissolved in DMF at around 75 °C. Subsequently, HAPFG was added to the solution under 4522

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The Journal of Physical Chemistry B magnetic stirring and then ultrasonically stirred for at least 4 h. The weight percentage of functionalized graphene was limited to below 3.0%. The mixed solution was then poured onto a Petri dish and dried at 60 °C in air for 12 h and then in a vacuum at 60 °C for 12 h to remove DMF. The films were removed from the Petri dish using a little bit of ethanol and were dried again at 130 °C for 4 h under air. Samples without any additives, denoted pure PU, were also prepared for comparison. The film thicknesses varied between 80 and 120 μm. Characterization. The X-ray diffraction (XRD) patterns are collected by a Bruker D8 Advance diffractometer. The infrared spectroscopy (IR) was recorded with the sample/KBr pressed pellets using a Bruker Vector-22 FT-IR spectrometer. The ultraviolet and visible spectrum (UV) was characterized by a Lambda 20 PerkinElmer UV−vis instrument. X-ray photoelectron spectroscopy (XPS) analysis was performed on a Perkin eElmer PHI550 spectrometer. Scanning electron microscopy (SEM) was taken on a JSM-5610LV instrument. Transmission electron microscopy (TEM) was obtained by a FEI Tecnai G2 instrument. The atomic force microscope (AFM) images were taken on a Shimadzu SPM-9500J3 atomic force microscope with the noncontact mode. Dielectric properties were measured by a HP4294A impedance analyzer. The electric breakdown strength was tested by a Beijing Electromechanical Research Institute dielectric withstand voltage test. The thermogravimetric analysis (TG) curves were recorded by a Diamond TG/DTA in N2 at a heating rate of 10 °C/min. The tensile properties of films were measured according to ASTM: D882 using a CM5105 universal testing machine at room temperature using a tensile rate of 10 mm/ min. The dynamic mechanical analysis (DMA) curves were measured by using a TA Model Q-800 dynamic mechanical analyzer. Samples are heated from −150 to 200 °C at a heating rate of 3 °C/min and at a constant frequency of 1 Hz with a static force of 0.02 N. The electric field induced strain measurement was measured by using the MTI Fotonic Sensor (MTI-2100) as described elsewhere.70 The “deflection strain” is defined as S = d/t, where d is the deflection displacement measured at the center of the diaphragm and t is the sample thickness.

Figure 1. FT-IR spectra of GO and EDA-GO-x (a); HAPFG-x (b).

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RESULTS AND DISCUSSION As shown in Scheme 1, GO was first modified by surface linkage of ethylenedimine through an amidation reaction. Subsequently, the EDA modified GO was grafted with hyperbranched aromatic polyamide by using 3,5-DABA as a monomer.71 Finally, the hyperbranched aromatic polyamide grafted GO was chemically reduced to obtain HAPFG. The hyperbranched aromatic polyamide functionalization was verified by FT-IR, UV−vis, XPS, TEM, and SEM. As shown in Figure 1a, the peak at 1740 cm−1 can be identified to represent the carbonyl moieties (CO) on its surface for GO. For EDA-GO, the peaks at 2930 and 2854 cm−1 representing the asymmetric and symmetric vibrations of CH2, the peak at 1640 cm−1 representing secondary amide (CO) stretching, and the peak at 1568 cm−1 representing secondary amide NH bending and CN stretching demonstrate that EDA has been successfully decorated onto the graphene.72 The absorbance peaks are significantly enhanced when the EDA grafting increases. HAPFG has similar absorbance peaks to EDA-GO. Besides, the characteristic absorbance peaks of metasubstituted phenyl are obviously much higher for HAPFG-3 in

Figure 2. UV spectra of EDA-GO-3 and HAPFG-3.

Figure 1b. As displayed in Figure 2, when compared with GO, the red-shift of the typical UV peak of EDA-GO (corresponding to π−π* transitions of aromatic CC) from 220 to 260 nm indicates that electronic conjugation of GO was restored, in agreement with preceding reports.73 As for HAPFG, a characteristic UV absorption peak is shown at 250−260 nm, indicating that GO was reduced. XPS was used here to evaluate the chemical composition of GO and HAPFG-3. The typical XPS spectra of the GO and HAPFG-3 are shown in Figure 3. The N 1s peak is not detected for GO, whereas the strong N 1s peak appears in the HAPFG-3. The atomic percentage of doped nitrogen is about 12.25 wt % according to the XPS, as shown in Table 1. The morphology of EDA-GO and HAPFG was investigated by TEM and SEM. As displayed in Figure 4a, EDA-GO shows a similar morphology to GO. The surface of the EDA-GO sample is relatively smooth. On the contrary, the surface of HAPFG appears to be rough and clearly covered by the addends (Figure 4b), which are attributed to the grafting of hyperbranched 4523

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Figure 4. TEM images of EDA-GO-3 (a) and HAPFG-3 (b) and SEM images of HAPFG-3 (c, d).

Figure 3. (a) XPS survey spectra of GO and HAPFG-3 and (b) typical high-resolution XPS of N 1s for HAPFG-3.

aromatic polyamide. As shown in Figure 4c and d, the SEM morphology of HAPFG shows an isotropic structure and very desultory stacking. There are a lot of wrinkles upon the surface of HAPFG. The morphology indicated that hyperbranched aromatic polyamide chains were homogeneously grafted onto the surfaces of graphene, which is consistent with the TEM results. Figure 5a shows a SEM top view picture of 3.0 wt % HAPFG3−PU. With HAPFG, no obvious aggregations can be observed. The SEM picture of the HAPFG−PU sample’s cross section confirms a homogeneous dispersion of nanosheets in the PU elastomer (Figure 5b). In order to reveal the roughness of the HAPFG−PU surface, we examined at least five AFM images of the HAPFG−PU surface at a scale of 2 × 2 μm2 in noncontact mode. The smooth surface provides evidence of good compatibility and interactions between HAPFG and PU (Figure 5c,d). The XRD results of PU and HAPFG−PU, shown in Figure 6, reveal that there is no detectable crystalline phase in PU and its composite. Furthermore, the XRD patterns of HAPFG−PU do not show any obvious characteristic peaks for graphene, indicating the uniform dispersion of HAPFG in PU matrix. The strong force between PU and HAPFG through the NHCOO group of PU and the NHCO group of hyperbranched aromatic polyamide leads to a uniform distribution of HAPFG in PU.

Figure 5. SEM patterns of HAPFG3−PU composites (a, top view; b, cross-section view) and AFM patterns of HAPFG3−PU (c, d).

TG curves for PU and HAPFG−PU are shown in Figure 7, and the shift in the onset of the thermal degradation temperature with weight percent loading of the HAPFG is observable. The PU sample possesses a single stage in the thermal degradation around 300−400 °C. All the HAPFG−PU samples exhibit similarly significant improvements in the thermal stability as the pure PU. The results confirm that the HAPFG can enhance the thermal stability of the PU and slow down the weight loss during thermal degradation. The TG curve of the 3% HAPFG−PU hybrid shows the amount of

Table 1. XPS Results of the Tested HAPFG-3 peak

binding energy (eV)

fwhm (eV)

raw area (cps eV)

RSF

atomic mass

atomic conc %

mass conc %

O 1s N 1s C 1s

531.524 399.824 284.824

3.177 2.055 1.815

9042.5 4696.3 18436.4

0.780 0.477 0.278

15.999 14.007 12.011

12.23 11.12 76.65

15.38 12.25 72.37

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Figure 6. XRD patterns of pure PU and HAPFG3−PU hybrids with different weight percentages of HAPFG-3.

Figure 9. DMA curves of pure PU, 1% RGO−PU, and different x% HAPFG−PU: (a) storage modulus; (b) loss modulus.

Figure 7. TG curves of pure PU and different HAPFG3−PU hybrids.

Figure 8. Stress−strain curves of pure PU, 1% RGO−PU, and different x% HAPFG−PU.

weight residue increased by 5.5% in comparison with the pure PU at 600 °C. Representative stress−strain curves of the HAPFG−PU hybrids are shown in Figure 8. The addition of HAPFG−PU significantly increases the initial modulus of the composites but reduces the ductility. As the mass percentages of HAPFG increase, the stress at a given strain tends to increase progressively and the strain at a break tends to decrease slowly. In order to get an in-depth understanding of the dynamic mechanical properties, the storage modulus and loss modulus values of the HAPFG−PU samples are measured (Figure 9). In the storage modulus vs temperature plot (Figure 9a), there is a break at about −45 °C, indicating a glass transition in the system. In the loss modulus plot (Figure 9b), the glass transition temperature may be computed from the peak. As shown in Figure 9, all the storage moduli and loss moduli of the

Figure 10. Relative permittivity (a) and dielectric loss (b) of HAPFxRGO-y%-PU with different weight fractions of HAPFG (1.0−3.0 wt %) measured at room temperature as a function of frequency. 4525

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Figure 12. Electric induced displacement of 3% HAPFG3−PU at super high sinusoidal electric field above 3000 V (a). Maximum thickness strain as a function of the applied electric field for pure PU and PU hybrids with different weight fractions of fillers (b).

Figure 13. Variations of the breakdown strength of HAPFG−PU and RGO−PU composites for different weight percentages of nanosheets.

RGO−PU significantly increases with the addition of rigid graphene nanosheets, which indicates the significant enhanced stiff properties of RGO/PU nanocomposites. The enhanced stiff property of RGO/PU may be due to the reinforcement by the high intensity graphene nanoplates. Moreover, the existence of large aggregates of graphene nanosheets increases the stiff property of RGO/PU. On the contrary, the flexible coating of hyperbranched aromatic polyamide chains onto graphene reinforces the interaction between the PU matrix and graphene sheets; thus, it is of benefit to homogeneous dispersion of fillers into the polymer matrix. This greatly reduces the loss modulus

Figure 11. High voltage strength−time plots (a, sinusoidal wave; c, square wave; e, sharp wave) and electric induced displacement−time plots (b, d, f) of 1% HAPFG3−PU with different kinds of electric fields.

1.0 wt % RGO−PU composites are much higher than those of the pure PU and HAPFG−PU at the temperature range studied here. The storage modulus and loss modulus of the 1.0 wt % 4526

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The Journal of Physical Chemistry B Table 2. Comparison of Several Conductor−Insulator Polyurethane DEs material

dielectric constant

dielectric loss

max actuation strain

HAPFG3−PU in this work (3 wt %) CNTs-g-PH10/PU (0.5 wt %)40 14PANI/15PolyCuPc/85PU29 carbon black/PU (1.0 wt %)36 SiC@C−PU (0.5 wt %)37 CB-ink/PU (1.7 vol %)38 MWCNT-COOH/PU (1.0 vol %)39 PMMA-RGO/PU (1.5 wt %)64 Fe3C−PU (2.5 wt %)41 CuPc−PU (8.78 vol %)77

217a 8.6b 800c 5.4a 25d 7.1e 330f 28.21a 6.6a 391g

0.14a 30b 0.6c 0.05a 0.8d 4.8e 10f 0.68a 0.2a 0.19g

92% thickness (98 μm, 38 MV/m) 1.1% transverse (30 × 10 × 1 mm3, 10% prestrain, 5 MV/m) 9.3% thickness (50 μm, 20 MV/m) 7.4% thickness (72 μm, 20 MV/m) 18% thickness (20 μm, 18 MV/m) 24% thickness (80 μm, 15 MV/m) 18% thickness (50 μm, 10 MV/m) 32.8% thickness (3D average strain, 110 μm, 9.1 MV/m) 130% thickness (80 μm, 20 MV/m) 17.7% in area (80 μm, 10 MV/m)

a

Measured on an impedance analyzer at 1000 Hz. bMeasured on a vector network analyzer at 50 Hz. cMeasured on an impedance analyzer at 1 Hz. Measured on a lock-in amplifier at 0.1 Hz. eMeasured on an LCR meter at 1000 Hz. fMeasured on an impedance meter at 0.1 Hz. gMeasured on an impedance analyzer at 100 Hz. d

near 100 μm, which is almost the thickness of the film itself. Figure 12b summarizes the results of the maximum electric induced strain versus the maximum electric field of several types of membrane actuators: 1.0, 2.0, and 3.0 wt % HAPFG3− PU, 1 wt % RGO−PU, and pure PU. It is obvious that the HAPFG−PU presented higher electric induced strain levels than the PU and RGO−PU counterparts. The strain curves consist of three distinguishable regions. The first region, where strain increases slightly with electric field, is a region of low strain (below 7 MV/m). The second region, where strain increases gradually with electric field, is a region of moderate strain (7−20 MV/m). Two different kinds of curve shapes were found in the high-strain region (above 20 MV/m). Strain curves of HAPFG3−PU show a steep slope, whereas curves of 1% RGO−PU and PU are relatively smooth. Excitingly, 92% electric induced strain of 3% HAPFG3−PU was achieved at 38 MV/m. This is much higher than the thickness strain obtained on the other polyurethane elastomer hybrids.25−29,36−39 We further investigated the high field electric breakdown strength of RGO−PU and HAPFG−PU composites at room temperature. As shown in Figure 13, the breakdown strength of pure PU is 83.3 MV/m. The breakdown strength of RGO−PU falls below 35 MV/m at 3 wt % RGO. Differently, with the addition of HAPFG nanosheets, the breakdown strength of HAPFG−PU only decreases to 74.8 MV/m at 1 wt % HAPFG and remains above 60 MV/m at higher HAPFG loadings up to 3 wt %. On the one hand, the hyperbranched aromatic polyamide functionalized layer can restrict the tunneling current of the HAPFG−PU composites at high electric field and thus increase the breakdown strength of the HAPFG− PU.75 On the other hand, the hyperbranched polymer layer formed on the surface of RGO nanosheets can be considered as a buffer layer and help to lower the local electric field concentrated inside the polymer matrix and thus improve the breakdown strength.76 For comparison, various conductor−insulator polyurethane composite materials are summarized in Table 2. High dielectric constants have been obtained in most of the composites; however, the reported electric induced strain is generally lower than HAPFG3−PU in this work except for Fe3C−PU. The primary advantage here is the novel chosen of a hyperbranched aromatic polyamide functionalized graphene−polyurethane dielectric elastomer composite for the microactuator. HAPFG−PU composites have a relatively high dielectric constant, low loss modulus, and high field strength. Its overall

of the corresponding PU hybrids. This is of great benefit to the electric induced strain properties. Figure 10 shows dielectric properties of PU and HAPFG− PU with different grafting contents of hyperbranched aromatic polyamide and different weight percentages of HAPFG. The dielectric constant of HAPFG−PU increased with the increase of grafting and filler fraction. With low grafting and filler fractions, the dielectric loss of HAPFG−PU decreases with the increase in frequency, giving a carrier-dominated response. With high grafting and filler fractions, the dielectric loss increases, giving a dipole-dominated response. This dielectric behavior indicates that interfacial polarization dominates for the capacitance of these microcapacitors in the HAPFG−PU and results in a high dielectric constant for the HAPFG−PU.74 Obviously, a great improvement of dielectric constant for HAPFG3−PU hybrids is observed. The dielectric constant of 3% HAPFG3−PU at 1 kHz is 217, 55 times higher than that of PU. Because the hyperbranched functionalization largely improves the dispersion of graphene in polymer matrix and the hyperbranched polymer can serve as dielectric layers to prevent graphene backbones from direct contact in the polymer matrix, more microcapacitors are constructed in HAPFG−PU, leading to a higher dielectric constant. Figure 11 shows the typical results of electric induced strain behaviors versus time with four to five cycles of a sinusoidalshaped, square-shaped, and sharp-shaped electric field. As shown in Figure 11, when the maximum electric field is raised, the maximum electric induced strain increases. This is obviously evidence of the electromechanical behavior of the PU EAPs being electrostrictive in nature. It is worth noticing that the displacement between the polymer film surface and sensor contact increases when the applied electric field is increasing, which indicates a compression of the film. Interestingly, the displacement signals have different shapes under different shapes of the applied electric field. When the shape of an electric field is sinusoidal, the displacement has a sinusoidal shape. When the shape of an electric field is square, the displacement has a sharp shape. The displacement shows increasing oscillation when the applied electric field is sharpshaped. We focused on the sinusoidal shape electric field and found an extra high electric induced strain at high sinusoidal electric field above 3000 V. Figure 12a shows the different thickness displacements of 3% HAPFG3−PU under high voltage (3000− 4000 V). The thickness displacements increase gradually with the rise of voltage. The maximum thickness displacement is 4527

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electromechanical properties are much higher than that of the other reported EAPs.

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CONCLUSIONS Hyperbranched aromatic polyamide functionalized graphene− polyurethane dielectric elastomer composites have been synthesized as promising candidate materials for flexible microactuators. Hyperbranched aromatic polyamide functionalized chains reinforce the interaction between the polyurethane matrix and graphene nanosheets and prevent graphene nanosheets from aggregating and connecting, which are favorable to the construction of microcapacitors in the matrix and the inhibition of leakage current. The dielectric analysis verified that the interface effect of this hyperbranched polymer layer significantly improved the dielectric performance. The dynamic mechanical behavior of the HAPFG−PU composites confirmed the strong interface bonding strength weakened the rigidity of the graphene−polyurethane composites. High field electric induced strain and electric breakdown strength electromechanical performance test proved the unique structure of EAPs filled with hyperbranched aromatic polyamide functionalized graphene had an advantage over the single component EAP. Generally, the HAPFG−PU composites have a high dielectric constant, low dielectric loss, low loss modulus, high electric induced strain, and high breakdown strength, all being essential to the application in MEMS.

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AUTHOR INFORMATION

Corresponding Authors

*Phone/Fax: +862552112626. E-mail: [email protected]. *Phone/Fax: +862584891123. E-mail: [email protected]. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was supported by Project Funded by the Priority Academic Program Development of Jiangsu Higher Education Institutions (PAPD), the National High Technology Research and Development Program of China (863 Program: 2013AA041105), the National Natural Science Foundation of China (11372133), the Aeronautical Science Foundation of China (No. 2014ZF52070), the Funding for Outstanding Doctoral Dissertation in NUAA (BCXJ13-11), the Funding of Jiangsu Innovation Program for Graduate Education (KYLX_0264), and the Fundamental Research Funds for the Central Universities (NS2015062).

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