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nano-TiO2 composite film
Applied Surface Science 257 (2011) 3831–3835
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Applied Surface Science journal homepage: www.elsevier.com/locate/apsusc
Review
Preparation and electroactive properties of a PVDF/nano-TiO2 composite film Ningli An a , Hongzhong Liu a,∗ , Yucheng Ding a , Min Zhang b , Yiping Tang a a b
State Key Laboratory for Manufacturing Systems Engineering, Xi’an Jiaotong University, Xi’an 710049, China College of Chemistry & Chemical Engineering, Shaanxi University of Science & Technology, Xi’an 710021, China
a r t i c l e
i n f o
Article history: Received 15 July 2010 Received in revised form 7 December 2010 Accepted 14 December 2010 Available online 21 December 2010 PACS: 77.55.+f 82.35.Np Keywords: PVDF TiO2 Composite materials Solution Thin films Electroactive
a b s t r a c t In the present study, poly(vinylidene fluoride) (PVDF)/nano-TiO2 electroactive film was prepared by coating a substrate with an acetone/DMF solution, which was evaporated at a high temperature (110 ◦ C). The crystallisation behaviour, dynamic mechanical properties and electroactive properties of this PVDF/nanoTiO2 electroactive film were investigated. The cross-section and surface of the film were observed with a scanning electron microscope (SEM). X-ray diffraction (XRD) results showed that the film containing the PVDF  phase, the desired ferroelectric phase, was obtained by crystallising the mixed solution of nano-TiO2 and PVDF at 110 ◦ C, while the film containing the ␣ phase was obtained from the crystallisation of the pure PVDF solution at the same temperature. It was found that the storage modulus, the room-temperature dielectric constant and the electric breakdown strength of the composite films were much higher than those of a pure PVDF film. TiO2 improved the mechanical properties and electroactive properties of the film. The results indicate that PVDF/nano-TiO2 composite films can be applied to the fabrication of self-sensing actuator devices. © 2010 Elsevier B.V. All rights reserved.
Contents 1. 2. 3.
4.
Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3831 Experiments . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3832 Results and discussion . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3832 3.1. Crystalline structure and microstructure of the films . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3832 3.2. Mechanical properties of the films . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3833 3.3. Electroactive performance of the film . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3833 Conclusion . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3834 Acknowledgements . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3835 References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3835
1. Introduction PVDF as a piezoelectric electroactive polymer (EAP) has enormous potential for use in biological self-sensing actuators, according to previous studies [1,2]. However, its low strain and low elastic energy density limit its applications. Hence, one of the major challenges is to obtain a high strain in the film. In recent
∗ Corresponding author at: State Key Laboratory for Manufacturing Systems Engineering, Xi’an Jiaotong University, Department of Mechanical Engineering, No. 28, Xianning West Road, Xi’an, Shaanxi 710049, China. Tel.: +86 29 83395021; fax: +86 29 83399508. E-mail address: [email protected] (H. Liu). 0169-4332/$ – see front matter © 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.apsusc.2010.12.076
years, Zhang et al. explored the high electrostrictive performance of copolymers and terpolymers in the PVDF family [3,4]. However, the polymerisation processes of these high-performance polymers were especially complex. Dang et al.’s experiment of blending fillers with PVDF resulted in an increase of the dielectric constant of the electroactive polymer [5–7]. According to the energy conversion law, when the applied field E is constant, a higher electric energy density can be obtained by increasing dielectric constant [3]. Among the five known crystalline forms of PVDF, namely ␣, , ␥, ␦ and , the  phase results in good piezoelectric and ferroelectric properties. Traditionally, the  phase is obtained by stretching films prepared by melting. Obviously, such a mechanical stretching process is not suitable for the preparation of thin films directly on
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substrates [8]. Instead, such a film could be prepared by evaporating an appropriate solution below 70 ◦ C, which facilitates coating the film directly on a substrate to obtain the  phase. Some previous studies [8–10] indicate that the films made at low temperatures were porous and fragile, while an evaporation temperature above 110 ◦ C is desirable. However, the ␣ phase is the predominant crystal structure when the evaporation temperature is above 110 ◦ C rather than the  phase [9]. Some methods have been proposed for coating films containing the  phase at evaporation temperatures above 110 ◦ C, such as the use of blending clay [11–13], hydrated ionic salt [14] or PMMA [15] with PVDF. To conveniently fabricate patterns during the preparation of the thin film, an integrated method was used. Nano-sized TiO2 particles were dispersed into the PVDF polymer matrix, and the electroactive film was evaporated from solution at 110 ◦ C. TiO2 , as a photocatalyst [16,17], could promote the induced degradation of PVDF via ultraviolet irradiation. The photocatalytic lithography method could be used to fabricate the patterns on the film in large batches to lower the cost. Although many attractive enhancements have been obtained in the conductivity performances and mechanical properties of PVDF/TiO2 nanoparticle composite films [18,19], studies on the crystal formation from blend solutions at high evaporation temperatures are scarcely found in the literature. In the present study, a PVDF/nano-TiO2 composite film was prepared, and its crystallisation behaviour, dynamic mechanical properties, dielectric constant and electroactive performance were investigated.
Fig. 1. XRD patterns of the PVDF and PVDF/nano-TiO2 films: (a1 ) PVDF film evaporated at 110 ◦ C, (a2 ) PVDF film evaporated at 70 ◦ C, (b1 ) PVDF/nano-TiO2 film evaporated at 110 ◦ C and (b2 ) PVDF/nano-TiO2 film evaporated at 70 ◦ C.
placement data of the film. An AC electric sine wave field at 0.5 Hz was used to measure the strain response.
2. Experiments
3. Results and discussion
In the experiment, PVDF powder with a molecular weight of 534,000 was obtained from the Sigma–Aldrich Corporation. TiO2 powder (P25) with an average particle size of 21 nm was obtained from the Degussa Corporation. The polymer was dissolved in a solvent composed of 70% acetone and 30% dimethylformamide (DMF). The polymer concentration was 50 g/L. The PVDF powder and milled TiO2 were added to the mixed solution of acetone and DMF in sequence. The resulting solution was placed in an ultrasonic water bath at 30 ◦ C for an hour. The thin films were prepared with a spin-coating process on a silicon substrate (1 0 0) or poured on a glass plate at 110 ◦ C and then cooled to room temperature. Samples a1 and a2 were prepared from the pure PVDF solution, while samples b1 and b2 were prepared from the PVDF solution with 5 wt% TiO2 . Samples a1 and b1 were prepared at 110 ◦ C, while samples a2 and b2 were prepared at 70 ◦ C. X-ray diffraction (XRD, RIGAKU) analysis was conducted to reveal the crystalline structure of the film by using the Cu K␣1 line (40 keV/100 mA). The scanning covered the 2 range of 10–60◦ with steps of 0.02◦ . The thickness of the film was measured with a stylus profilometer (Ambios XP, American). The morphology of the film was observed with a scanning electron microscope. A dynamic mechanical analysis was carried out in a tension mode on film samples by using a dynamic mechanical analyser. The tests were performed in isochronal conditions (frequency 20 Hz). The samples were heated at a rate of 3◦ /min. An HP8720ES network analyser was used to measure the relative dielectric constant and dielectric loss at low frequencies from 10−2 Hz to 107 Hz. For this measurement, the samples were cut into discs with diameters of 20 mm, and layers of Pt (0.2 m) were deposited on the upper and lower sides with a magnetron sputtering system (ACS-4000-C4, ULVAC) at room temperature. An electrostriction test on the film with a thickness of 30 m was performed on a circular frame (radius R = 10 mm). To test the electroactive performance, a 100-nm-thick carbon layer (radius R = 10 mm) was imprinted on both sides of the polymer film. A laser scanning vibrometer (LK-G80, Keyence) was used to collect the dis-
3.1. Crystalline structure and microstructure of the films Fig. 1 shows the XRD patterns of the PVDF films and the PVDF/nano-TiO2 composite films prepared at different temperatures. The results show that the PVDF films predominantly contained the ␣ phase when the acetone/DMF solution was evaporated at 110 ◦ C, while  phase was obtained at 70 ◦ C, which agrees well with previous studies for PVDF films [9]. The crystallisation behaviour of the PVDF/nano-TiO2 composite film prepared at 110 ◦ C was obviously different from that of the pure PVDF films. The PVDF/nano-TiO2 composite film had strong peaks (1 1 0) and (2 0 0) corresponding to the typical  phase and obviously weakened peaks (0 2 0) and (1 2 0) corresponding to the ␣ phase. Nano-TiO2 had a strong effect on the crystallinity of the PVDF molecules, which has been proven by experiments [20]. Due to the addition of TiO2 , a large fraction of the  phase instead of the ␣ phase was formed in the composite film from the evaporation of the solution at 110 ◦ C. On the other hand, nano-TiO2 retards the crystallisation of PVDF and preserves the amorphous structure in the PVDF/nano-TiO2 composite film [18], which sacrifices the degree of crystallinity. Fig. 2(a) and (b) shows the cross-sections of the pure PVDF film and of the PVDF/nano-composite film, respectively. The two samples were obtained from the crystallisation of solutions at 110 ◦ C. It was found that the pure PVDF film’s surface formed globules with diameters between 20 and 30 m. This micrograph suggests that the sample predominantly contains the ␣ phase [21]. The PVDF/nano-composite film was observed to possess an interpenetrating network of long fibre chains. These chains directly link adjacent crystalline lamellae. Fig. 2(c) shows that the TiO2 nanoparticles were dispersed in the polymer matrix and micro-sized aggregates exist. These results indicate that the transition from the metastable  phase to the stable ␣ phase was prevented by the homogeneous dispersion of TiO2 nano-particles in the PVDF matrix owing to the match of crystal lattice of TiO2 with the  phase of PVDF, which is similar to the effect of adding clay to obtain the  phase of PVDF [11]. Hydrogen bonds were formed between the
N. An et al. / Applied Surface Science 257 (2011) 3831–3835
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Fig. 3. Dynamic mechanical analysis of PVDF and PVDF/nano-TiO2 films: (a) storage modulus vs. temperature and (b) loss modulus vs. temperature.
Fig. 2. The SEM images taken at a 7◦ tilt: (a) pure PVDF (the insert at the bottom left shows close-up of the cross-section) and (b) PVDF/nano-TiO2 (The insert at the top right shows close-up of the surface). The insert at the bottom left shows close-up cross-section). (c) SEM backscattered electron image of the PVDF/nano-TiO2 film’s surface.
the loss modulus of the PVDF/nano-TiO2 composite film increases below room temperature, which indicates that the PVDF/nanocomposite film has an improved mechanical strength at lower temperatures. Apparently, TiO2 exerts a reinforcement effect on the PVDF/nano-composite film [26]. The micro-fibrillar structure and lamellae of the PVDF/nano-TiO2 composite film leads to its enhanced hardness and elastic properties [27]. The high surface area of the nano-TiO2 can result in frictional energy dissipation during the interfacial sliding of nano-particles within the composite film [11,28]. 3.3. Electroactive performance of the film
hydroxyl group on the surface of the TiO2 particles [22,23] and the polar C–F bonds of the PVDF [24]. The movement and arrangement of the polymer chains were restricted. The polymer chains were oriented to the surface of TiO2 to form an orderly arrangement. This mechanism induces formation of the  phase of PVDF. Therefore, the composite film retained the  crystal phase at 110 ◦ C. The variation in the phase content is accompanied by changes in the film microstructure [25]. 3.2. Mechanical properties of the films As shown in Fig. 3(a), the storage modulus of the PVDF/nanoTiO2 composite film is higher than that of pure PVDF throughout the temperature range from −50 ◦ C to 150 ◦ C. Fig. 3(b) shows how
In Fig. 4(a), the relative dielectric constant of the PVDF/nanoTiO2 composite film increased considerably to about 20 units at low frequencies due to the addition of TiO2 , which has a higher dielectric constant by nature and causes interfacial polarisation of the Maxwell–Wagner type [29]. Fig. 4(b) shows that the dielectric loss of the sample with TiO2 increased in the lower frequency range below 1 Hz but decreased in the higher frequency range above 104 Hz. The decrease in the dielectric loss at low frequencies could be attributed to the polarisation of the -phase PVDF induced by the TiO2 . The low frequency dielectric loss is sensitive to the existence of polar radicals with dipole moments [29]. The relationships between the displacement of the film centre and the voltage were acquired by positioning a laser scanning
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Fig. 5. (a) Amplitude of response of the film to an electric field with an applied field frequency of 0.5 Hz. (b) Amplitude of response of film under the electric field with the different applied field amplitudes and an applied field frequency of 0.5 Hz. Fig. 4. (a) PVDF and PVDF/TiO2 dielectric constant. (b) PVDF and PVDF/TiO2 dielectric loss.
vibrometer detector over the centre of the upper electrodes of the films. Fig. 5(a) shows the amplitude of responses of the composite films under an electric field of 90 V/m and a pure film under an electric field of 50 V/m. The displacement response frequency is 0.5 Hz. The vibration system of the PVDF/nano-TiO2 composite film achieved a relatively late dynamic stability, which can be attributed to the improvement in the storage modulus accompanied by a decrease in the vibration damping [30]. Nevertheless, the strain amplitudes induced by the electric field are enhanced in PVDF/nano-TiO2 composite film, as shown in Fig. 5(b). The composite film was broken when the electric field strength reached about 150 V/m, while the pure PVDF film was broken when electric field strength was only 50 V/m. The increased amorphous structure in the PVDF/nano-TiO2 composite and enhanced mechanical properties resulted in the increase of the electric field breakdown strength. As for the PVDF/nano-TiO2 film, 45 V/m could be taken as a threshold. At this point, it is interesting to notice that the strain of the composite film is smaller than that of the pure PVDF film below this threshold and rapidly increases above this threshold. This is because stress relaxation is weakened above the threshold when electric energy density is sufficient to overcome the mobility and movement of the polymer chains. This study clearly proves that the electroactive properties of the nano-TiO2 composite film are significantly improved by dispersing nano-TiO2 particles in the PVDF. The enhancement of strain in composite film is assumed to be controlled by two factors. On the one hand, an important role is played by the amorphous phase tie chains bordering the crystals. The amorphous phase makes a certain con-
tribution to the electroactive behaviour of the material [25,31]. On the other hand, the -phase affects the electroactive properties of polymer [25]. The present study reveals that the -phase PVDF induced by the presence of nano-TiO2 in the composite film is randomly oriented, while a preferred orientation of the dipoles can be induced by certain electric fields thus increasing the electroactive response of composite film. Therefore, the strain of the composite film rapidly increased when the electric field reached a certain strength. Furthermore, both the mobility of TiO2 and its dispersion in the soft cross-links of PVDF resulted in local strains in the composite films. The variation in the dielectric constant is accompanied by changes in the strain under applied electric fields [1]. The increased toughness resulting from the ability of nanoparticles to dissipate energy [11] might further enhance the strain.
4. Conclusion This study shows that -phase poly(vinylidene fluoride) and the micro-fibrillar structures were formed in the composite film, which resulted in the considerable improvement of the mechanical properties, room temperature dielectric constant and electric breakdown field relative to those of pure PVDF. Due to these contributions, the PVDF/nano-TiO2 composite film has a distinctive electroactive response. Furthermore, the preparation of the PVDF/nano-TiO2 composite film in the acetone/DMF solution allowed the electrostrictive film to be easily deposited on a substrate by spin-coating without the conventional stretching process, which can provide a possible approach for fabricating self-sensing actuator devices.
N. An et al. / Applied Surface Science 257 (2011) 3831–3835
Acknowledgements This work is supported by National Natural Science Foundation of China (No. 90923040, 50975226), National Basic Research Program of China (Nos. 2009CB724202), New Century Excellent Talents (NCET-08-0447), and the Fundamental Research Funds for the Central Universities. References [1] R. Hayakawa, Y. Wada, Adv. Polym. Sci. 11 (1973) 1–55. [2] M. Oshiki, E. Fukada, J. Mater. Sci. 10 (1975) 1–6. [3] Q.M. Zhang, H. Li, M. Poh, F. Xia, Z.Y. Cheng, H. Xu, C. Huang, Nature 419 (2002) 284–287. [4] K. Ren, S. Liu, M. Lin, Y. Wang, Q.M. Zhang, Sens. Actuators A 143 (2008) 335–342. [5] L. Wang, Z.M. Dang, Appl. Phys. Lett. 87 (2005) 042903. [6] Z.M. Dang, H.Y. Wang, H.P. Xu, Appl. Phys. Lett. 89 (2006) 112902. [7] Z.M. Dang, H.P. Xu, D. Xie, L. Li, Mater. Lett. 61 (2007) 511–515. [8] X. He, K. Yao, B.K. Gan, Sens. Actuators A 139 (2007) 158–161. [9] R. Gregorio, E.M. Ueno, J. Mater. Sci. 34 (1999) 4489–4500. [10] W. Ma, J. Zhang, X. Wang, Appl. Surf. Sci. 254 (2008) 2947–2954. [11] D. Shah, P. Maiti, E. Gunn, D.F. Schmidt, D.D. Jiang, C.A. Batt, E.P. Giannelis, Adv. Mater. 16 (2004) 1173–1176.
[12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27]
[28] [29] [30] [31]
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L. Priya, J.P. Jog, J. Polym. Sci. B: Polym. Phys. 40 (2002) 1682–1689. T.U. Patro, M.V. Mhalgi, D.V. Khakhar, A. Misra, Polymer 49 (2008) 3486–3499. M. Benz, W.B. Euler, O.J. Gregory, Macromolecules 35 (2002) 2682–2688. W. Ma, J. Zhang, S. Chen, X. Wang, Appl. Surf. Sci. 254 (2008) 5635–5642. L. Zan, L. Tian, Z. liu, Z. Peng, Appl. Catal. A 264 (2004) 237–242. X. Zhao, Z. Li, Y. Chen, L. Shi, Y. Zhu, Appl. Surf. Sci. 245 (2008) 1825–1829. Y.J. Wang, D. Kim, Electrochim. Acta 52 (2007) 3181–3189. L.Y. Yu, H.M. Shen, Z.L. Xu, J. Appl. Polym. Sci. 113 (2009) 1763–1772. X. Cao, J. Ma, X. Shi, Z. Ren, Appl. Surf. Sci. 253 (2006) 2003–2010. R. Gregorio Jr., J. Appl. Polym. Sci. 100 (2006) 3272–3279. K.M. Kim, N.G. Park, K.S. Ryu, S.H. Chang, Polymer 43 (2002) 3951–3957. C. Arrouvel, M. Digne, M. Breysse, H. Toulhoat, P. Raybaud, J. Catal. 222 (2004) 152–166. T.M. Becker, J.A.K. Bauer, J.E.D. Bene, M. Orchin, J. Organomet. Chem. 629 (2001) 165–170. J. Gomes, J.S. Nunes, V. Sencadas, S. Lanceros-Mendez, Smart Mater. Struct. 19 (2010) 065010. J.L. Yang, Z. Zhang, H. Zhang, Compos. Sci. Technol. 65 (2005) 2374–2379. P. Podsiadlo, A.K. Kaushik, E.M. Arruda, A.M. Waas, B.S. Shim, J. Xu, H. Nandivada, B.G. Pumplin, J. Lahann, A. Rammamoorthy, N.A. Kotov, Science 318 (2007) 80–83. D. Devaprakasam, P.V. Hatton, G. Möbus, B.J. Inkson, Tribol. Lett. 34 (2009) 11–19. T. Tanaka, IEEE Trans. Dielectr. Electr. Insul. 12 (2005) 914–928. L. Meirovitch, Principles and Techniques of Vibrations, Prentice Hall, Englewood Cliffs, 1997. V.V. Kochervinskii, Russ. Chem. Rev. 68 (1999) 821–857.
Contents lists available at ScienceDirect
Applied Surface Science journal homepage: www.elsevier.com/locate/apsusc
Review
Preparation and electroactive properties of a PVDF/nano-TiO2 composite film Ningli An a , Hongzhong Liu a,∗ , Yucheng Ding a , Min Zhang b , Yiping Tang a a b
State Key Laboratory for Manufacturing Systems Engineering, Xi’an Jiaotong University, Xi’an 710049, China College of Chemistry & Chemical Engineering, Shaanxi University of Science & Technology, Xi’an 710021, China
a r t i c l e
i n f o
Article history: Received 15 July 2010 Received in revised form 7 December 2010 Accepted 14 December 2010 Available online 21 December 2010 PACS: 77.55.+f 82.35.Np Keywords: PVDF TiO2 Composite materials Solution Thin films Electroactive
a b s t r a c t In the present study, poly(vinylidene fluoride) (PVDF)/nano-TiO2 electroactive film was prepared by coating a substrate with an acetone/DMF solution, which was evaporated at a high temperature (110 ◦ C). The crystallisation behaviour, dynamic mechanical properties and electroactive properties of this PVDF/nanoTiO2 electroactive film were investigated. The cross-section and surface of the film were observed with a scanning electron microscope (SEM). X-ray diffraction (XRD) results showed that the film containing the PVDF  phase, the desired ferroelectric phase, was obtained by crystallising the mixed solution of nano-TiO2 and PVDF at 110 ◦ C, while the film containing the ␣ phase was obtained from the crystallisation of the pure PVDF solution at the same temperature. It was found that the storage modulus, the room-temperature dielectric constant and the electric breakdown strength of the composite films were much higher than those of a pure PVDF film. TiO2 improved the mechanical properties and electroactive properties of the film. The results indicate that PVDF/nano-TiO2 composite films can be applied to the fabrication of self-sensing actuator devices. © 2010 Elsevier B.V. All rights reserved.
Contents 1. 2. 3.
4.
Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3831 Experiments . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3832 Results and discussion . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3832 3.1. Crystalline structure and microstructure of the films . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3832 3.2. Mechanical properties of the films . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3833 3.3. Electroactive performance of the film . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3833 Conclusion . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3834 Acknowledgements . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3835 References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3835
1. Introduction PVDF as a piezoelectric electroactive polymer (EAP) has enormous potential for use in biological self-sensing actuators, according to previous studies [1,2]. However, its low strain and low elastic energy density limit its applications. Hence, one of the major challenges is to obtain a high strain in the film. In recent
∗ Corresponding author at: State Key Laboratory for Manufacturing Systems Engineering, Xi’an Jiaotong University, Department of Mechanical Engineering, No. 28, Xianning West Road, Xi’an, Shaanxi 710049, China. Tel.: +86 29 83395021; fax: +86 29 83399508. E-mail address: [email protected] (H. Liu). 0169-4332/$ – see front matter © 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.apsusc.2010.12.076
years, Zhang et al. explored the high electrostrictive performance of copolymers and terpolymers in the PVDF family [3,4]. However, the polymerisation processes of these high-performance polymers were especially complex. Dang et al.’s experiment of blending fillers with PVDF resulted in an increase of the dielectric constant of the electroactive polymer [5–7]. According to the energy conversion law, when the applied field E is constant, a higher electric energy density can be obtained by increasing dielectric constant [3]. Among the five known crystalline forms of PVDF, namely ␣, , ␥, ␦ and , the  phase results in good piezoelectric and ferroelectric properties. Traditionally, the  phase is obtained by stretching films prepared by melting. Obviously, such a mechanical stretching process is not suitable for the preparation of thin films directly on
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substrates [8]. Instead, such a film could be prepared by evaporating an appropriate solution below 70 ◦ C, which facilitates coating the film directly on a substrate to obtain the  phase. Some previous studies [8–10] indicate that the films made at low temperatures were porous and fragile, while an evaporation temperature above 110 ◦ C is desirable. However, the ␣ phase is the predominant crystal structure when the evaporation temperature is above 110 ◦ C rather than the  phase [9]. Some methods have been proposed for coating films containing the  phase at evaporation temperatures above 110 ◦ C, such as the use of blending clay [11–13], hydrated ionic salt [14] or PMMA [15] with PVDF. To conveniently fabricate patterns during the preparation of the thin film, an integrated method was used. Nano-sized TiO2 particles were dispersed into the PVDF polymer matrix, and the electroactive film was evaporated from solution at 110 ◦ C. TiO2 , as a photocatalyst [16,17], could promote the induced degradation of PVDF via ultraviolet irradiation. The photocatalytic lithography method could be used to fabricate the patterns on the film in large batches to lower the cost. Although many attractive enhancements have been obtained in the conductivity performances and mechanical properties of PVDF/TiO2 nanoparticle composite films [18,19], studies on the crystal formation from blend solutions at high evaporation temperatures are scarcely found in the literature. In the present study, a PVDF/nano-TiO2 composite film was prepared, and its crystallisation behaviour, dynamic mechanical properties, dielectric constant and electroactive performance were investigated.
Fig. 1. XRD patterns of the PVDF and PVDF/nano-TiO2 films: (a1 ) PVDF film evaporated at 110 ◦ C, (a2 ) PVDF film evaporated at 70 ◦ C, (b1 ) PVDF/nano-TiO2 film evaporated at 110 ◦ C and (b2 ) PVDF/nano-TiO2 film evaporated at 70 ◦ C.
placement data of the film. An AC electric sine wave field at 0.5 Hz was used to measure the strain response.
2. Experiments
3. Results and discussion
In the experiment, PVDF powder with a molecular weight of 534,000 was obtained from the Sigma–Aldrich Corporation. TiO2 powder (P25) with an average particle size of 21 nm was obtained from the Degussa Corporation. The polymer was dissolved in a solvent composed of 70% acetone and 30% dimethylformamide (DMF). The polymer concentration was 50 g/L. The PVDF powder and milled TiO2 were added to the mixed solution of acetone and DMF in sequence. The resulting solution was placed in an ultrasonic water bath at 30 ◦ C for an hour. The thin films were prepared with a spin-coating process on a silicon substrate (1 0 0) or poured on a glass plate at 110 ◦ C and then cooled to room temperature. Samples a1 and a2 were prepared from the pure PVDF solution, while samples b1 and b2 were prepared from the PVDF solution with 5 wt% TiO2 . Samples a1 and b1 were prepared at 110 ◦ C, while samples a2 and b2 were prepared at 70 ◦ C. X-ray diffraction (XRD, RIGAKU) analysis was conducted to reveal the crystalline structure of the film by using the Cu K␣1 line (40 keV/100 mA). The scanning covered the 2 range of 10–60◦ with steps of 0.02◦ . The thickness of the film was measured with a stylus profilometer (Ambios XP, American). The morphology of the film was observed with a scanning electron microscope. A dynamic mechanical analysis was carried out in a tension mode on film samples by using a dynamic mechanical analyser. The tests were performed in isochronal conditions (frequency 20 Hz). The samples were heated at a rate of 3◦ /min. An HP8720ES network analyser was used to measure the relative dielectric constant and dielectric loss at low frequencies from 10−2 Hz to 107 Hz. For this measurement, the samples were cut into discs with diameters of 20 mm, and layers of Pt (0.2 m) were deposited on the upper and lower sides with a magnetron sputtering system (ACS-4000-C4, ULVAC) at room temperature. An electrostriction test on the film with a thickness of 30 m was performed on a circular frame (radius R = 10 mm). To test the electroactive performance, a 100-nm-thick carbon layer (radius R = 10 mm) was imprinted on both sides of the polymer film. A laser scanning vibrometer (LK-G80, Keyence) was used to collect the dis-
3.1. Crystalline structure and microstructure of the films Fig. 1 shows the XRD patterns of the PVDF films and the PVDF/nano-TiO2 composite films prepared at different temperatures. The results show that the PVDF films predominantly contained the ␣ phase when the acetone/DMF solution was evaporated at 110 ◦ C, while  phase was obtained at 70 ◦ C, which agrees well with previous studies for PVDF films [9]. The crystallisation behaviour of the PVDF/nano-TiO2 composite film prepared at 110 ◦ C was obviously different from that of the pure PVDF films. The PVDF/nano-TiO2 composite film had strong peaks (1 1 0) and (2 0 0) corresponding to the typical  phase and obviously weakened peaks (0 2 0) and (1 2 0) corresponding to the ␣ phase. Nano-TiO2 had a strong effect on the crystallinity of the PVDF molecules, which has been proven by experiments [20]. Due to the addition of TiO2 , a large fraction of the  phase instead of the ␣ phase was formed in the composite film from the evaporation of the solution at 110 ◦ C. On the other hand, nano-TiO2 retards the crystallisation of PVDF and preserves the amorphous structure in the PVDF/nano-TiO2 composite film [18], which sacrifices the degree of crystallinity. Fig. 2(a) and (b) shows the cross-sections of the pure PVDF film and of the PVDF/nano-composite film, respectively. The two samples were obtained from the crystallisation of solutions at 110 ◦ C. It was found that the pure PVDF film’s surface formed globules with diameters between 20 and 30 m. This micrograph suggests that the sample predominantly contains the ␣ phase [21]. The PVDF/nano-composite film was observed to possess an interpenetrating network of long fibre chains. These chains directly link adjacent crystalline lamellae. Fig. 2(c) shows that the TiO2 nanoparticles were dispersed in the polymer matrix and micro-sized aggregates exist. These results indicate that the transition from the metastable  phase to the stable ␣ phase was prevented by the homogeneous dispersion of TiO2 nano-particles in the PVDF matrix owing to the match of crystal lattice of TiO2 with the  phase of PVDF, which is similar to the effect of adding clay to obtain the  phase of PVDF [11]. Hydrogen bonds were formed between the
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Fig. 3. Dynamic mechanical analysis of PVDF and PVDF/nano-TiO2 films: (a) storage modulus vs. temperature and (b) loss modulus vs. temperature.
Fig. 2. The SEM images taken at a 7◦ tilt: (a) pure PVDF (the insert at the bottom left shows close-up of the cross-section) and (b) PVDF/nano-TiO2 (The insert at the top right shows close-up of the surface). The insert at the bottom left shows close-up cross-section). (c) SEM backscattered electron image of the PVDF/nano-TiO2 film’s surface.
the loss modulus of the PVDF/nano-TiO2 composite film increases below room temperature, which indicates that the PVDF/nanocomposite film has an improved mechanical strength at lower temperatures. Apparently, TiO2 exerts a reinforcement effect on the PVDF/nano-composite film [26]. The micro-fibrillar structure and lamellae of the PVDF/nano-TiO2 composite film leads to its enhanced hardness and elastic properties [27]. The high surface area of the nano-TiO2 can result in frictional energy dissipation during the interfacial sliding of nano-particles within the composite film [11,28]. 3.3. Electroactive performance of the film
hydroxyl group on the surface of the TiO2 particles [22,23] and the polar C–F bonds of the PVDF [24]. The movement and arrangement of the polymer chains were restricted. The polymer chains were oriented to the surface of TiO2 to form an orderly arrangement. This mechanism induces formation of the  phase of PVDF. Therefore, the composite film retained the  crystal phase at 110 ◦ C. The variation in the phase content is accompanied by changes in the film microstructure [25]. 3.2. Mechanical properties of the films As shown in Fig. 3(a), the storage modulus of the PVDF/nanoTiO2 composite film is higher than that of pure PVDF throughout the temperature range from −50 ◦ C to 150 ◦ C. Fig. 3(b) shows how
In Fig. 4(a), the relative dielectric constant of the PVDF/nanoTiO2 composite film increased considerably to about 20 units at low frequencies due to the addition of TiO2 , which has a higher dielectric constant by nature and causes interfacial polarisation of the Maxwell–Wagner type [29]. Fig. 4(b) shows that the dielectric loss of the sample with TiO2 increased in the lower frequency range below 1 Hz but decreased in the higher frequency range above 104 Hz. The decrease in the dielectric loss at low frequencies could be attributed to the polarisation of the -phase PVDF induced by the TiO2 . The low frequency dielectric loss is sensitive to the existence of polar radicals with dipole moments [29]. The relationships between the displacement of the film centre and the voltage were acquired by positioning a laser scanning
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Fig. 5. (a) Amplitude of response of the film to an electric field with an applied field frequency of 0.5 Hz. (b) Amplitude of response of film under the electric field with the different applied field amplitudes and an applied field frequency of 0.5 Hz. Fig. 4. (a) PVDF and PVDF/TiO2 dielectric constant. (b) PVDF and PVDF/TiO2 dielectric loss.
vibrometer detector over the centre of the upper electrodes of the films. Fig. 5(a) shows the amplitude of responses of the composite films under an electric field of 90 V/m and a pure film under an electric field of 50 V/m. The displacement response frequency is 0.5 Hz. The vibration system of the PVDF/nano-TiO2 composite film achieved a relatively late dynamic stability, which can be attributed to the improvement in the storage modulus accompanied by a decrease in the vibration damping [30]. Nevertheless, the strain amplitudes induced by the electric field are enhanced in PVDF/nano-TiO2 composite film, as shown in Fig. 5(b). The composite film was broken when the electric field strength reached about 150 V/m, while the pure PVDF film was broken when electric field strength was only 50 V/m. The increased amorphous structure in the PVDF/nano-TiO2 composite and enhanced mechanical properties resulted in the increase of the electric field breakdown strength. As for the PVDF/nano-TiO2 film, 45 V/m could be taken as a threshold. At this point, it is interesting to notice that the strain of the composite film is smaller than that of the pure PVDF film below this threshold and rapidly increases above this threshold. This is because stress relaxation is weakened above the threshold when electric energy density is sufficient to overcome the mobility and movement of the polymer chains. This study clearly proves that the electroactive properties of the nano-TiO2 composite film are significantly improved by dispersing nano-TiO2 particles in the PVDF. The enhancement of strain in composite film is assumed to be controlled by two factors. On the one hand, an important role is played by the amorphous phase tie chains bordering the crystals. The amorphous phase makes a certain con-
tribution to the electroactive behaviour of the material [25,31]. On the other hand, the -phase affects the electroactive properties of polymer [25]. The present study reveals that the -phase PVDF induced by the presence of nano-TiO2 in the composite film is randomly oriented, while a preferred orientation of the dipoles can be induced by certain electric fields thus increasing the electroactive response of composite film. Therefore, the strain of the composite film rapidly increased when the electric field reached a certain strength. Furthermore, both the mobility of TiO2 and its dispersion in the soft cross-links of PVDF resulted in local strains in the composite films. The variation in the dielectric constant is accompanied by changes in the strain under applied electric fields [1]. The increased toughness resulting from the ability of nanoparticles to dissipate energy [11] might further enhance the strain.
4. Conclusion This study shows that -phase poly(vinylidene fluoride) and the micro-fibrillar structures were formed in the composite film, which resulted in the considerable improvement of the mechanical properties, room temperature dielectric constant and electric breakdown field relative to those of pure PVDF. Due to these contributions, the PVDF/nano-TiO2 composite film has a distinctive electroactive response. Furthermore, the preparation of the PVDF/nano-TiO2 composite film in the acetone/DMF solution allowed the electrostrictive film to be easily deposited on a substrate by spin-coating without the conventional stretching process, which can provide a possible approach for fabricating self-sensing actuator devices.
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Acknowledgements This work is supported by National Natural Science Foundation of China (No. 90923040, 50975226), National Basic Research Program of China (Nos. 2009CB724202), New Century Excellent Talents (NCET-08-0447), and the Fundamental Research Funds for the Central Universities. References [1] R. Hayakawa, Y. Wada, Adv. Polym. Sci. 11 (1973) 1–55. [2] M. Oshiki, E. Fukada, J. Mater. Sci. 10 (1975) 1–6. [3] Q.M. Zhang, H. Li, M. Poh, F. Xia, Z.Y. Cheng, H. Xu, C. Huang, Nature 419 (2002) 284–287. [4] K. Ren, S. Liu, M. Lin, Y. Wang, Q.M. Zhang, Sens. Actuators A 143 (2008) 335–342. [5] L. Wang, Z.M. Dang, Appl. Phys. Lett. 87 (2005) 042903. [6] Z.M. Dang, H.Y. Wang, H.P. Xu, Appl. Phys. Lett. 89 (2006) 112902. [7] Z.M. Dang, H.P. Xu, D. Xie, L. Li, Mater. Lett. 61 (2007) 511–515. [8] X. He, K. Yao, B.K. Gan, Sens. Actuators A 139 (2007) 158–161. [9] R. Gregorio, E.M. Ueno, J. Mater. Sci. 34 (1999) 4489–4500. [10] W. Ma, J. Zhang, X. Wang, Appl. Surf. Sci. 254 (2008) 2947–2954. [11] D. Shah, P. Maiti, E. Gunn, D.F. Schmidt, D.D. Jiang, C.A. Batt, E.P. Giannelis, Adv. Mater. 16 (2004) 1173–1176.
[12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27]
[28] [29] [30] [31]
3835
L. Priya, J.P. Jog, J. Polym. Sci. B: Polym. Phys. 40 (2002) 1682–1689. T.U. Patro, M.V. Mhalgi, D.V. Khakhar, A. Misra, Polymer 49 (2008) 3486–3499. M. Benz, W.B. Euler, O.J. Gregory, Macromolecules 35 (2002) 2682–2688. W. Ma, J. Zhang, S. Chen, X. Wang, Appl. Surf. Sci. 254 (2008) 5635–5642. L. Zan, L. Tian, Z. liu, Z. Peng, Appl. Catal. A 264 (2004) 237–242. X. Zhao, Z. Li, Y. Chen, L. Shi, Y. Zhu, Appl. Surf. Sci. 245 (2008) 1825–1829. Y.J. Wang, D. Kim, Electrochim. Acta 52 (2007) 3181–3189. L.Y. Yu, H.M. Shen, Z.L. Xu, J. Appl. Polym. Sci. 113 (2009) 1763–1772. X. Cao, J. Ma, X. Shi, Z. Ren, Appl. Surf. Sci. 253 (2006) 2003–2010. R. Gregorio Jr., J. Appl. Polym. Sci. 100 (2006) 3272–3279. K.M. Kim, N.G. Park, K.S. Ryu, S.H. Chang, Polymer 43 (2002) 3951–3957. C. Arrouvel, M. Digne, M. Breysse, H. Toulhoat, P. Raybaud, J. Catal. 222 (2004) 152–166. T.M. Becker, J.A.K. Bauer, J.E.D. Bene, M. Orchin, J. Organomet. Chem. 629 (2001) 165–170. J. Gomes, J.S. Nunes, V. Sencadas, S. Lanceros-Mendez, Smart Mater. Struct. 19 (2010) 065010. J.L. Yang, Z. Zhang, H. Zhang, Compos. Sci. Technol. 65 (2005) 2374–2379. P. Podsiadlo, A.K. Kaushik, E.M. Arruda, A.M. Waas, B.S. Shim, J. Xu, H. Nandivada, B.G. Pumplin, J. Lahann, A. Rammamoorthy, N.A. Kotov, Science 318 (2007) 80–83. D. Devaprakasam, P.V. Hatton, G. Möbus, B.J. Inkson, Tribol. Lett. 34 (2009) 11–19. T. Tanaka, IEEE Trans. Dielectr. Electr. Insul. 12 (2005) 914–928. L. Meirovitch, Principles and Techniques of Vibrations, Prentice Hall, Englewood Cliffs, 1997. V.V. Kochervinskii, Russ. Chem. Rev. 68 (1999) 821–857.