Remarkably variable dielectric and magnetic properties of poly(vinylidene fluoride) nanocomposite films with triple-layer structure

Composites Science and Technology 107 (2015) 107–112

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Remarkably variable dielectric and magnetic properties of poly(vinylidene fluoride) nanocomposite films with triple-layer structure Li Ren a, Jun Zhao a, Si-Jiao Wang a, Jun-Wei Zha a, Guo-Hua Hu b, Zhi-Min Dang a,⇑ a Laboratory of Dielectric Polymer Materials and Devices, Department of Polymer Science and Engineering, School of Chemistry and Biological Engineering, University of Science and Technology Beijing, Beijing 100083, PR China b Université de Lorraine – CNRS, Laboratoire Réactions et Génie des Procédés, UMR 7274, ENSIC, 1 rue Grandville, BP 20451, Nancy F-54000, France

a r t i c l e

i n f o

Article history: Received 21 August 2014 Received in revised form 24 November 2014 Accepted 6 December 2014 Available online 12 December 2014 Keywords: B. Electrical properties B. Magnetic properties A. Polymer-matrix composites (PMCs) A. Layered structures

a b s t r a c t The variable dielectric and magnetic properties of poly(vinylidene fluoride) (PVDF) nanocomposite films with triple-layer structure and with electrically conductive multi-walled carbon nanotubes (MWCNTs) and magnetic iron oxide (Fe3O4) nanoparticles as fillers, and PVDF as polymer matrix are systematically investigated. The single layer of MWCNTs/PVDF, Fe3O4/PVDF, and pure PVDF is denoted as A, B and P, respectively. The multilayered films with different arrangements are prepared by a simple two-step method. Scanning electron microscopy (SEM) shows good adhesion between the layers after stacking. The experimental results indicate that triple layer AAA films have higher dielectric constant (e) and lower dielectric loss (tan d) than the single layer A film with the same thickness in the frequency range of 102–104 Hz due to the interfacial effect. The nanocomposite films with structure of ABA and APA keep low tan d when the filler content is high due to the intermediate B and P layer as an inter-barrier to prevent the formation of conductive network in these samples. To replace the intermediate B layer of BBB films by A layer produces much higher saturation magnetization (Ms) for the filler content above 2.0 wt% due to the synergistic effect. But the replacement of the B layer by P layer causes nearly no significant change of Ms. This work demonstrates a new method to tune the dielectric and magnetic properties of nanocomposites. Ó 2014 Elsevier Ltd. All rights reserved.

1. Introduction In recent years, polymer-based nanocomposites have attracted tremendous attention in academia and industry because of their unique properties and potential applications [1,2]. Most of the polymer nanocomposites possess both the advantages of polymer matrix such as light weight, easy processing, and low cost and those of inorganic nanofillers such as special functionalities and high thermal stability. Therefore, the physiochemical properties of nanocomposites can be adjusted by changing the type of polymers, the kind and size of fillers, the distribution of fillers in the system, and the interface features between polymers and fillers [2–11]. In most cases the specific needs of materials could be realized by tuning these parameters of the system [12–14]. So far, great success has been achieved to optimize the properties of the polymer-matrix nanocomposites. To obtain better ⇑ Corresponding author. E-mail address: [email protected] (Z.-M. Dang). http://dx.doi.org/10.1016/j.compscitech.2014.12.008 0266-3538/Ó 2014 Elsevier Ltd. All rights reserved.

comprehensive properties including multiple functionalities, hybrid fillers are often used [15,16]. For example, both high dielectric constant and low dielectric loss could be obtained by using cofillers of surface-functionalized graphene nanosheets and barium titanate (BaTiO3) nanoparticles in poly(vinylidene fluoride) (PVDF) [15]. Besides, both high thermal conductivity and high electrical resistivity could be achieved by controlling the spatial distribution of hybrid fillers of multi-walled carbon nanotubes (MWCNTs) and silicon carbide (SiC) nanoparticles in immiscible PVDF/polystyrene (PS) 30/70 (vol%) blends [16]. However, it is usually hard to precisely control the spatial distribution of different fillers in the hybrid system to form a required micro-structure with high stability. In addition, stacking of double and even multiple polymer nanocomposite films with different compositions and functions has emerged as a useful strategy to better control the dielectric properties of polymer nanocomposites in the past decade [17–19]. For example, compared with the single layer multi-walled carbon nanotube (MWCNT)/cyanate ester (CE) nanocomposites, higher

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dielectric constant and much lower dielectric loss were found in the double-layer materials with a layer of polyethylene (PE) and a layer of MWCNT/CE resin composite [17]. Besides, triple layer materials with carbon nanofiber (CNF)/PVDF nanocomposites intercalated by a pure PVDF layer have enhanced the dielectric constant and the low dielectric loss [18]. A versatile processing technique called coextrusion was used to fabricate PVDF/polycarbonate (PC) multilayer films with even thousands of nanolayers [19]. Such materials have a high energy density and low dielectric loss. In this work, first MWCNTs with high electric conductivity and iron (II, III) oxide (Fe3O4) nanoparticles with excellent magnetic properties are added into PVDF matrix to form single layer functional materials, respectively. PVDF is used as the polymer matrix because of its unique chemical, mechanical, dielectric, and piezoelectric properties [20]. Then the single layers are stacked together in various ways to investigate the variation of the dielectric and magnetic properties. This work could shed some light on the better optimization of such polymer nanocomposites. 2. Experimental

dissolved in 100 mL deionized water in a 250 mL three necked round bottom flask in a water bath at 80.0 °C with mechanical stirring under N2 protection. And then 30 mL 25 wt% NH3H2O was quickly added to the solution. After a few seconds the solution turned black and it was further stirred for 1 h. Then a mixture of 6 mL ethanol and 6 mL oleic acid was added to the suspension and the system was further stirred for 2 h. The obtained precipitate was separated from the solution by using a magnet and washed with deionized water for three times and with ethanol twice. The cleaned precipitate was dried in vacuum oven at 70.0 °C for 24 h. The diameter of the obtained surface modified Fe3O4 nanoparticle was measured to be ca. 25–35 nm. 2.3. Surface modification of MWCNTs The purchased MWCNTs were surface-modified according to the literature [24,25]. First ca. 3.0 g MWCNTs were soaked into a large amount of nitric acid in a round bottom flask for refluxing 6 h, and the mixture was separated by filtering and washed with deionized water for several times. Then the treated MWCNTs were dried in vacuum oven at 80.0 °C for 12 h.

2.1. Materials 2.4. Preparation of single layer and multilayer nanocomposite films PVDF (FR904) powders with a density of 1.78 g cm3 and a melt flow index of 26.0 g (10 min1) were supplied by Shanghai 3F New Materials Co. (China). MWCNT (95% purity) with density of 2.10 g cm3, inner diameter of 5–10 nm, outer diameter of 10–20 nm, and length of 10–30 lm were purchased from Chengdu Organic Chemicals Co. Ltd. (China). FeCl24H2O (AR), FeCl36H2O (AR), Oleic acid (AR), nitric acid (AR) and NH3H2O (AR) were purchased from Xilong Chemical Co. Ltd. (China) and used without further purification. N,N-dimethylformamide (DMF) (AR) with density of 0.95 g cm3 at 25.0 °C was supplied by Beijing Chemical Works (China). 2.2. Preparation and surface modification of Fe3O4 nanoparticles Fe3O4 nanoparticles were prepared by a conventional coprecipitation method and surface-modified according to the literature [21–23]. Firstly, 7.83 g FeCl36H2O and 2.87 g FeCl24H2O were

PVDF powders were dried completely in a vacuum oven at 80 °C overnight before use. First, MWCNTs or Fe3O4 nanoparticles with various weight ratios were added into 20 mL DMF, and the resultant solution was sonicated at room temperature for 2 h to form the suspension. Then ca. 3.2 g PVDF powders were also added into DMF solvent and the mixture was mechanically stirred in water bath at 70.0 °C for 3 h to ensure the complete dissolution of PVDF. The suspension was then cast into thin films on a clean glass plate by spin-coating at room temperature. The obtained films were kept in a vacuum oven at 70.0 °C for 12 h to completely remove any remnant solvent. The MWCNT/PVDF and Fe3O4/PVDF single layer nanocomposite films with different filler concentrations of ca. 1.0–5.0 wt% possessed a thickness of ca. 8 lm and ca. 24 lm and are denoted as A8, B8, A24, and B24, respectively. The pure PVDF films prepared with a similar procedure to the nanocomposites had a thickness of ca. 8 lm and are denoted as P8. The prepared

Fig. 1. Schematic presentation of different stacking to form multilayer structure.

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A8, B8, and P8 films in rectangular shape with area of 25  76 mm2 were stacked together in different ways such as A8A8A8, B8B8B8, A8B8A8, B8A8B8, A8P8A8, and B8P8B8. The stacked triple-layer films were then put on a hot presser and compressed at 175.0 °C under a pressure of 10 MPa for 10 min. Very little change of thickness was found for all layers after processing, which could be explained as follows. Generally the processing temperature for PVDF matrix was 200.0 °C or above as reported in reference [18,26]. The temperature could ensure the complete flow of PVDF matrix. In this work, the processing temperature was 175.0 °C, the pressure was 10 MPa and kept 10 min. The temperature, pressure, and time were relatively low. So the slight diffusion of the PVDF molecules happened at the interface between two layers to achieve the self-adhesion. Because the diffusion was limited by the low temperature, pressure and time, the thickness of the nanocomposite films had no significant change. Fig. 1 shows the schematic presentation of different stacking to form multilayer structure. The various sandwiched structures including ABA, APA, BAB, and BPB were built up for the comparison. Triple layers of the same samples such as AAA and BBB were also prepared to see the effect of the interfaces. 2.5. Scanning electron microscopy (SEM) observation The samples were fractured in liquid nitrogen and then the fractured surface was sputtered with gold. Morphology observation on cross-section surface of the samples was performed on scanning electron microscopy (SEM S4700, Hitachi, Japan) with an accelerating voltage of 20 kV. 2.6. Measurements of electrical properties and magnetic properties For the electrical property measurements, samples in square shape with area of ca. 1.0 cm2 were cut from the prepared films

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and both sides of the samples were coated with silver as electrodes. Their dielectric properties including of alternating current (AC) conductivity (r) of the samples were measured using an impedance analyzer (Agilent 4294A) in the frequency range from 102 to 106 Hz at room temperature. For the magnetic properties measurements, samples with a mass of ca. 20–35 mg were cut from the prepared films and their magnetic properties were measured using a Vibration Sample Magnetometer (BKT-4500Z, Beijing Zetian Technology Co. Ltd., China) at room temperature.

3. Results and discussion 3.1. The phase morphologies of the nanocomposite films Fig. 2 displays the SEM images of MWCNT/PVDF and Fe3O4/ PVDF nanocomposite films with filler content of 2.0–5.0 wt%, sandwiched samples of ABA, APA, BAB, and BPB, in which the contents of fillers in these sandwiched samples are all 2.0 wt%. By comparing Fig. 2(a1) and (b1) with Fig. 2(a2)–(a4) and (b2)–(b4), it can be seen that the fillers, MWCNT and Fe3O4 are well dispersed in the PVDF matrix for the filler content of 2.0 wt%. However, some aggregations of fillers can be seen in both nanocomposites when the filler concentration reaches 3.0 wt%. In this work, the surface modification of MWCNTs and Fe3O4 nanoparticles and the bath-type sonication are used in order to achieve a good dispersion of fillers. It obtains a good effect when the filler content is low. However, more serious aggregation is observed in both nanocomposites with higher filler contents (as indicated by the red circles in Fig. 2(a2)– (a4) and (b2)–(b4)). Besides, some voids and defects can be seen in the Fe3O4/PVDF nanocomposites with increasing fillers content (as indicated by the green circles in Fig. 2(b2)–(b4)). These results can be attributed to the dispersive state of fillers. When the filler content is high, the bath-type sonication has no significant effect for

Fig. 2. SEM images of the MWCNT/PVDF and the Fe3O4/PVDF nanocomposite films with 2.0 wt% (a1, b1), 3.0 wt% (a2, b2), 4.0 wt% (a3, b3), and 5.0 wt% (a4, b4) filler content, sandwiched samples of ABA (c1), APA (c2), BAB (d1), and BPB (d2), (the content of fillers in c1, c2, d1 and d2 are all 2.0 wt%). The red circles in (a2–a4) and (b2–b4) indicate the aggregation and the green circles in (b2–b4) indicate the voids. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)

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the dispersion of fillers. Maybe a high power sonication treatment could effectively avoid the presence of aggregates, voids and defects in the nanocomposite films. Such aggregations and voids could have some effect on the dielectric and magnetic properties of the nanocomposite films. By comparing among Fig. 2(c1), (c2), (d1) and (d2), it can be seen that there is no visible gap between the two components for ABA, APA, BAB, and BPB structures, which indicates good adhesion between the components. Similar results have been reported in the system of a sandwich structure composed of CNF/PVDF composite layers and a pure PVDF interlayer, in which the bonding zones were perfectly integrated on both sides of the PVDF middle layer [26]. Besides, the surface of A and B layers is much rougher than that of the P component, which is due to the brittle-to-ductile transition induced by the fillers [27]. The thickness of each layer in the multilayered structure is almost the same as that of the as-prepared single layer, which means that the hot compression with a low temperature does not cause any significant compression of the samples. 3.2. The electrical properties of the nanocomposite films Frequency dependence of dielectric constant (e) and dielectric loss (tan d) of the pure PVDF and PVDF nanocomposite films with 2 wt% fillers stacked together in different ways are shown in Fig. 3. It can be seen that the addition of fillers increases both the e and the tan d of the pure PVDF. For the A24 and A8A8A8 nanocomposite films, their dielectric properties are influenced by the frequency and multilayered structure. It shows the e of A8A8A8 is higher and the tan d is lower than A24 with the same thickness in the frequency range of 102–104 Hz. The higher e for the A8A8A8 than that for the A24 is because of the significant interfacial polarization in the former and the lower tan d for the A8A8A8, which may be due to the weak leak loss at interface of the multilayered structure in an applied electric field [18]. Therefore, a combined effect of both the increasing interfacial polarization and the decreasing leak loss

at interface result in the improvement of dielectric properties for this sample and it is more significant in low frequency range. Besides, it can be seen that the replacement of the intermediate A8 layer by B8 layer causes decreasing e of the nanocomposite film and the intermediate A8 layer by P8 layer causes further decreasing e of the nanocomposite film (as shown in Fig. 3(a)). The results can be explained as follows: for the same filler content and frequency, the e of A8, B8, and P8 follows a sequence: A8 > B8 > P8, therefore the e of A8A8A8, A8B8A8, and A8P8A8 follows a sequence: A8A8A8 > A8B8A8 > A8P8A8. Besides, the insulating layer of B8 and P8 used in A8B8A8 and A8P8A8 structure, respectively, plays an inter-barrier role between the two MWCNT/PVDF nanocomposite films, which would retrain the polarization of MWCNTs in the nanocomposite films with triple-layer structure result in the low e and tan d. Meanwhile, the insulating layers as barriers to avoid the conductive network formed in the triple-layer structure, which would further decrease the tan d (as shown in Fig. 3(b)) [28,29]. Fig. 4 shows the filler content dependence of e and tan d (a), and r (b) at 100 Hz for the nanocomposite films of A24, A8A8A8, A8B8A8, and A8P8A8. It can be seen from Fig. 4(a) that the e and tan d for all the samples increase with increasing filler content, when the filler content reaches 3.0 wt%, the tan d for all the nanocomposite films increases as higher as 0.1. From Fig. 4(b), we can find that as the content reaches 3.0 wt%, the A24 and A8A8A8 nanocomposite films have already become conductors rather than insulators. This result can be ascribed to the conductive network formed in these nanocomposite films. For the A8B8A8 and A8P8A8 samples, the e increases with increasing filler content up to 3.0 wt%. However, it is not improved further with increasing filler content from 3.0 wt% to 5.0 wt%. This is because of that the effect of dielectric properties in nanocomposites is not only a function of the filler content, but also that of the dispersion of the fillers. When the filler content is 4.0 wt% and 5.0 wt%, the worse dispersion state of the fillers and the formation of some voids or defects in the samples are present as indicated by the red circles and green circles in Fig. 2(a3)–(a4)

Fig. 3. Frequency dependence of e (a) and tan d (b) of pure PVDF and the PVDF nanocomposite samples with 2.0 wt% fillers stacked together in different ways.

Fig. 4. Filler content dependence of e and tan d (a), and r (b) at 100 Hz of the pure PVDF and PVDF nanocomposite films stacked together in different combinations.

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and (b3)–(b4). According to the Maxwell–Wagner–Sillars (MWS) polarization mechanism, the e of the nanocomposites is strongly dependent on the interfacial effect [2,30]. The presence of aggregates reduces the interface areas between the fillers and polymer matrix thus suppressing the attribution of the interfacial polarization effect. This is responsible for the e of samples with 4.0 and 5.0 wt% filler content being no higher than that of the 3.0 wt% samples. Besides, the more serious aggregates and voids in A8B8A8 and A8P8A8 nanocomposite films is one reason for the slight decrease of the tan d. For the change of r of A8B8A8 and A8P8A8 nanocomposite films, it can be understood by the change in tan d of these nanocomposite films. The relationship between the r, tan d, and imaginary part (e00 ) of the relative complex permittivity is demonstrated as below:

e00 ¼ r=2pf e0

ð1Þ

e00 / tan d

ð2Þ

In Eq. (1), the f and e0 represent frequency and vacuum dielectric constant, respectively. Eqs. (1) and (2) clearly show that the change of r is similar to the change of tan d. So the r increases with increasing filler content up to 3.0 wt%, and then the r decreases for the filler content of 4.0 wt% and 5.0 wt%. It is worth noting that the A8B8A8 and A8P8A8 nanocomposite films are still insulators even for the filler content as high as 5.0 wt% due to the inter-barrier present in those samples [28,29]. These results indicate that the sandwich/ multi-layered structure is an effective method to tune the dielectric properties of nanocomposite films.

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with the filler content of 2.0 wt% shows that the replacement of the intermediate B8 layer by A8 layer produces much higher M but the replacement by P8 layer causes nearly no change of M. The former might be due to the synergistic effect between conductive nanoparticles and magnetic nanoparticles, while the latter does not have such effect. Besides, the replacement of the intermediate B8 layer by both A8 and P8 layers significantly increases the coercivity. Our explanation is that the intermediate layer of non-magnetic film is seen as the doping in the magnetic samples and results in the decrease of dipolar interaction between the magnetic particles [32]. Meanwhile, such a doping can cause the asymmetry of coercivity. It can be seen from Fig. 5(b) that the Ms of B8A8B8 is lower than that of B24 and is almost the same as that of B8B8B8 and B8P8B8 when the filler content is 1.0 wt%. However, the Ms of B8A8B8 is sharply increases when the filler content reaches 2.0 wt%, and then the increase trend of Ms slows down as the filler content increases from 3.0 wt% to 5.0 wt%. These results can be explained by the dipolar interaction and the synergistic effect comes from the coupling interactions [33]. The dipolar interaction between magnetic particles is dominant for the low content of MWCNTs e.g. 1.0 wt%, while the synergistic effect is dominant when the content reaches 2.0 wt%, which may be due to the MWCNTs in the contact to each other in the intermediate layer. However, the voids and defects in Fe3O4/PVDF films increase with increasing filler content, which results in the reduction of dipolar interaction and synergistic effect in these nanocomposite films. So the increase trend of Ms slows down.

4. Conclusion 3.3. The magnetic properties of the nanocomposite films Fig. 5 shows the magnetic hysteresis loops of the PVDF nanocomposite films with filler content of 2.0 wt% stacked together in different ways and the filler content dependence of saturation magnetization (Ms) of such PVDF nanocomposite films. It can be seen that the magnetic hysteresis loops of all the samples are S-like curves (as shown in Fig. 5(a)) and the Ms of all the samples increases with increasing filler content (as shown in Fig. 5(b)). According to Fig. 5(a), for B24 and B8B8B8 nanocomposite films with 2.0 wt% Fe3O4 nanoparticles, the latter has a little lower M than the former for the same magnetic field (H), which might be due to the weak interaction among the Fe3O4 nanoparticles for the multilayered structure in an applied magnetic field [31]. However, it is interesting to note that the coercivity of samples B24 and B8B8B8 nanocomposite films with 2.0 wt% Fe3O4 nanoparticles remain almost unchanged, which indicates that the multilayer structure itself has little effect on the movement of the domain walls. The comparison among B8B8B8, B8A8B8, and B8P8B8 nanocomposite films

In summary, the variable dielectric and magnetic properties of polymer nanocomposite triple-layer films had been systematically investigated. The SEM showed good adhesion between the layers after stacking. It was found that triple layer AAA films had higher e and lower tan d than the single layer film with the same thickness in the frequency range of 102–104 Hz due to the interfacial effect. The replacement of the intermediate A layer of AAA film by B and P layer caused the change in both of the e and tan d. The e and tan d of ABA and APA nanocomposite films increased with increasing filler content up to 3.0 wt%, and then decreased when the filler content reached 4.0 wt% and 5.0 wt%. To replace the intermediate B layer of BBB films by A layer produced much higher Ms for the filler content above 2.0 wt% due to the synergistic effect. But the replacement of the B layer by P layer caused nearly no significant change of Ms. Besides, the replacement of the intermediate B layer by both A and P layers significantly increased the coercivity. This work could shed some light on the optimization of such polymer nanocomposites.

Fig. 5. Magnetic hysteresis loops of the PVDF nanocomposite films with filler content of 2.0 wt% stacked together in different ways (a) and the filler content dependence of Ms for the PVDF nanocomposite films stacked together in different combinations (b).

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Acknowledgments This work was financially supported by NSF of China (Grant Nos. 51425201, 51207009 and 51377010), Ministry of Education of China through Doctor Project (Grant No. 20130006130002), the National Basic Research Program of China (973 Program) (Grant No. 2014CB239503), and Development fund for Graduate Student Education in University of Science and Technology Beijing.

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