Fabrication of Stretchable Nanocomposites with High Energy Density

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Fabrication of Stretchable Nanocomposites with High Energy Density and Low Loss from Crosslinked PVDF Filled with Poly(dopamine) Encapsulated BaTiO3 Yunchuan Xie, Yangyang Yu, Yefeng Feng, Wanrong Jiang, and Zhicheng Zhang ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.6b14166 • Publication Date (Web): 09 Jan 2017 Downloaded from http://pubs.acs.org on January 10, 2017

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ACS Applied Materials & Interfaces

Fabrication of Stretchable Nanocomposites with High Energy Density and Low Loss from Crosslinked PVDF Filled with Poly(dopamine) Encapsulated BaTiO3 Yunchuan Xie†, Yangyang Yu†, Yefeng Feng††, Wanrong Jiang†, Zhicheng Zhang††* †

Department of Materials Chemistry, School of Science, Xi’an Jiaotong University, Xi’an, P. R. China, 710049. Department of Applied Chemistry, MOE Key Laboratory for Nonequilibrium Synthesis and Modulation of Condensed Matter, School of Science, Xi’an Jiaotong University, Xi’an, P. R. China, 710049. ††

ABSTRACT: In this report, a simple solution-cast method was employed to prepare poly(dopamine) (PDA) encapsulated BaTiO3 (BT) nanoparticles (PDA@BT) filled composites using PVDF matrix crosslinked by the free radical initiator. The effects of both the particle encapsulation and matrix crosslinking on the mechanical and dielectric properties of the composites were carefully investigated. The results suggested that the introduction of BT particles improved permittivity of the composites to ~30 at 100 Hz when particle contents of only 7 wt% were utilized. This was attributed to the enhanced polarization, which was induced by high permittivity ceramic particles. Compared to bare BT, PDA@BT particles could be dispersed more homogeneously in the matrix, and the catechol groups of PDA layer might form chelation with free ions present in the matrix. The latter might depress the ion conduction loss in the composites. Other results revealed that the formation of hydrogen-bonding between the PDA layer and the polymer, especially the chemical crosslinking across the matrix, resulted in increased Young’ modulus by ~25%, improved breakdown strength by ~40%, and declined conductivity by nearly one order of magnitude when compared to BT filled composites. The composite films filled with PDA@BTs indicated greater energy storage capacities by nearly 190% when compared to the pristine matrix. More importantly, the excellent mechanical performance allowed the composite films to adopt uni- or bi-axially stretching, a crucial feature required for the realization of high breakdown strength. This work provided a facile strategy for fabrication of flexible and stretchable dielectric composites with depressed dielectric loss and enhanced energy storage capacity at low filler loadings (< 10 wt%). KEYWORDS: stretchable; poly(dopamine); surface modification; chemical crosslinking; Dielectric loss; breakdown strength; energy storage capacity;

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1. INTRODUCTION With the rapid development of advanced power and electronic systems, energy storage capacitors with low cost, compact size and high performance are increasingly investigated. The dielectric materials, which can undergo fast charge and discharge of electrostatic energy with high efficiency, have attracted considerable attention for their valuable use in high energy storage fields and power density devices, such as a high-power laser, microwave, rail-gun, and electromagnetic ejection techniques

1-5

. For simple linear response dielectric materials, the maximum energy density could

2 simply be defined as U e = 1 2 ε r ε 0 (E b ) (where ε r is the relative dielectric permittivity, ε 0 is the

vacuum permittivity with the value of 8.85×10-12 F/m, and Eb is the dielectric breakdown strength). Apparently, to obtain high Ue, both the εr and Eb of dielectric materials should simultaneously be as high as possible. During the past two decades, more attention has been focused on fabrication of dielectric materials with high εr using compositing techniques to overcome the shortcoming of low εr in polymeric materials 6-11

. One way is based on the usage of organic or inorganic conductive particles, including nickel

powder

12

, polyaniline

13

and graphene

14-15

, as fillers to construct polymer based composite dielectrics

according to percolative and interfacial polarization theories 16-19. Although this kind of composites may lead to higher εr reaching up to 102 orders of magnitude, their inherent percolative characteristics resulted in larger leakage loss and significantly lower Eb near the percolation threshold. This, in turn, limited their high-field electrical applications. Introducing ferroelectric ceramic particles with superior dielectric constants (εr > 103) into polymer matrices, especially poly(vinylidene fluoride) (PVDF) based polymers with εr of 10, is another well-investigated strategy to produce high dielectric composites. However, the realization of high εr in such composites requires larger volume content of ceramic (> 30 vol%) to be added into the polymer matrix according to either series or parallel models 20,21. Moreover, the uneven electrical field distribution and deteriorated flexibility of the composites often lead to


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significantly lowered Eb. Consequently, methods that would lead to simultaneous improvement of εr and Eb have emerged as one of the major challenges of composite dielectrics, in which tremendous efforts have been devoted particularly in polymer/ceramic composites. For instance, in order to prepare composites with improved compatibility between two phases, Atom transfer radical polymerization (ATRP) and Reversible addition-fragmentation chain transfer (RAFT) polymerization technologies have been performed through “graft from” route on ceramic surfaces to form “core-shells” structured ceramic particles, such as PS@BaTiO3 22, PMMA@BaTiO3

23

, and PGMA@BaTiO3 particles

24

. Thanks for

their enhanced interfacial compatibilities, particles could well be dispersed and incorporated into polymer matrices and composite dielectrics with high dielectric constants and low dielectric losses 25-27. Besides, to reduce the filling content, nanofillers with high aspect ratio, have been employed instead of particles with irregular structures to yield effective improvement in εr of the composites under relatively low volume fraction when compared to their spherical counterparts. This includes 1D BaTiO3 fibers 29

, barium strontium titanate (Ba0.7Sr0.3TiO3) nanowires

30

, and 2D montmorillonite clay

31

28-

, BN

nanosheets 2. Nevertheless, both surface-initiated polymerization which are employed to graft polymer brushes from inorganic particles and production of non-spherical nanoparticles as high-εr fillers could partially solve the problems of composites at low-cost and high-efficiency in the large-scale production. Previously published studies have mostly focused on improving the dielectric properties of composites but, the influence of mechanical and film processing properties on their dielectric properties have barely been mentioned. In fact, many studies suggested the close relationship between the electrical and mechanical properties of the materials. The well-known Stark-Garton equation ( Eb = 0.6 (Y ε 0ε r ) ) clearly states that increased mechanical modulus often favors elevated dielectric 0.5

breakdown strength of dielectric materials. Meanwhile, the state-of-the-art dielectric films like polypropylene have to be biaxially stretched in order to induce high Eb, and thus improve Ue and stability under high electric fields. The introduction of inorganic particles usually decrease elongation and reduce the flexibility of target composites. This means that high-εr composites would be difficult to ACS Paragon Plus Environment

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fabricate in large scale, which may significantly reduce their breakdown strength and limit their application to high electric fields. In this study and by taking into consideration the above-mentioned problems, a series of PVDF/BT flexible composites with εr of up to 30 and Eb of 300 MV/m were fabricated using a facilitated method. The surface modification of BT was performed using a low-cost poly(dopamine) (PDA) encapsulation process. Dopamine (DA), a bio-inspired molecule with catechol and amine groups, was utilized due to its powerful ability to deposit thin polymer layers (PDA) on nearly any bulk material surface. The dopamine based method has already been demonstrated its high-efficiency and easy-handling in improving compatibility of inorganic ceramics with polymer matrices

32-34

. PVDF bearing internal

double bonds named as P(VDF-CTFE-DB) was synthesized from partially dehydrochlorinated poly(vinylidene fluoride-chlorotrifluoroethylene) (P(VDF-CTFE)) precursor and was employed for crosslinking purpose. The composites were fabricated through a cast of suspensions containing modified BT, P(VDF-CTFE-DB) and benzoperoxide (BPO) as crosslinking initiator onto glass slides, followed by removal all volatiles and curing at elevated temperature. The data suggested that both the surface modification of BT and the crosslinking of polymer matrix were responsible for the depressed dielectric loss, significantly improved breakdown strength, and elevated energy storage density of the resulting dielectric composites. Most importantly, the fine flexibility and high elongation characteristics were well retained in the resulting composites. These features were important for large scale use through uni- or bi-axially stretching.

2. EXPERIMENTAL SECTION 2.1 Fabrication of “0-3” type nanocomposites. A typical solution cast process35-36 was used to fabricate “0-3” type nanocomposite films as follows and illustrated in Scheme 1. BT or PDA@BT particles were dispersed into DMF by ultrasonication and vigorously stirring for 30 minutes before transferred into P(VDF-CTFE-DB) solution in DMF with a concentration of 5 wt%. The composite


4

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films were fabricated by casting the suspensions bearing fillers of varied weight contents and polymer matrix onto glass slides, followed by removing all the volatiles at 90 oC for 4 h. The increased weight contents of BT fillers, 0.0, 2.0, 3.8, 5.0 and 7.0 wt%, were designed in resultant samples. For PDA@BT filled composites, BPO (5 wt% of polymer matrix) was used to cure the polymer matrix at 150 oC after the films were molded at 90 oC. “0-3” type composite films with thickness of ~20 µm were peeled off from the slides and sputtered with Au on both surfaces as electrodes for subsequent electrical measurements. As illustrated in Scheme 1, PDA@BT particles exhibit much better dispersion in DMF than neat BTs, which benefits the following fabrication process of PDA@BT filled composites.

Scheme 1. Schematic illustration for the surface modification of BT particles and preparation of crosslinked polymer nanocomposite films. Insert digital photograph: BT particles (left) and PDA@BT particles (right) in DMF, settled for 2 hours after ultrasonic treatment for 10 minutes. Table 1. Abbreviations for the materials and composites. Full names

Abbreviations

Uncrosslinked P(VDF-CTFE-DB) Crosslinked P(VDF-CTFE-DB)

PVDF c-PVDF

Neat BaTiO3

BT

PDA modified BaTiO3

PDA@BT

Uncrosslinked P(VDF-CTFE-DB)/neat BaTiO3

PVDF/BT

Crosslinked P(VDF-CTFE-DB)/neat BaTiO3

c-PVDF/BT

Crosslinked P(VDF-CTFE-DB)/modified BaTiO3

c-PVDF/PDA@BT


5


For expression convenience, the materials and composites in present work were indicated by abbreviations as listed in Table 1.

3. RESULTS AND DISCUSSION 3.1. Chemical structure and morphology of PDA@BT. The chemical structure and morphology of PDA modified BT were characterized by FT-IR, TEM, TGA and XRD. 100

BT Weight percent (%)

BT

Wavenumber

1500 1000 (cm-1)

500

0

70 100

100

200

200

300

300

400

500

400

600

700

500

800

600

700

800

Temperature (°C)

d)

(111)

(100)

c)

(222)

2000

0

b)

(311)

2500

83.5%

80

20

(310)

3000

PDA@BT

(300)

3500

90

40

(220)

4000

97.9%

(211)

a)

BT

100

60

(210)

-N-H 1625

-C-N1327 -C-C1500

PDA@BT

(200)

PDA@BT

80

(110)

Transmittance (a.u.)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

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PDA@BT BT 20

30

40

50

60

70

80

90

2θ (°) Figure 1. a) FT-IR spectra, b) TGA curves of BT and PDA@BT nanoparticles, c) TEM image of PDA@BT particles, and d) XRD spectra of BT and PDA@BT particles. The FT-IR results of BT and PDA@BT particles were compared, and the results are shown in ACS Paragon Plus Environment

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Figure 1a. The new absorption bands appeared at 1625 cm-1 and 1327 cm-1 in PDA@BT were mainly attributed to the bending vibration of N-H bond and C-N bond stretching vibration in the aromatic amine, respectively. The stretching vibration of aromatic C-C bond could be confirmed by the absorption peak at 1500 cm-1. The latter suggested the successful coating of PDA molecules onto BT particles surfaces 37-38. The TGA results of two samples are presented in Figure 1b. The profiles exhibited residual weight of PDA@BT by about 83.5 wt% at 600 oC, which was nearly 14.4% lower than that of neat BT samples (97.9 wt%). This could be assigned to the thermal degradation of encapsulated PDA layer, indicating that PDA content in PDA@BT particles was around 15 wt%. The coating structure of PDA around BT could well be characterized with TEM scanning. Figure 1c depicted a distinct contrast between BT core and PDA shell. The average thickness of the PDA layer was estimated to about 12 nm according to the average counting calculations shown in Figure 1c. The XRD patterns of neat BT and PDA@BT particles are illustrated in Figure 1d. Apparently, all the characteristic diffraction peaks can well be assigned to the crystal form of BaTiO3. The latter was well consistent with the standard cards (JCPDS 31-0174) 39. Both the samples were confirmed to have cubic perovskite crystal form, where the PDA coating process did not alter the crystal form of BT particles.

3.2. Microstructure and morphology of composites. The chemical composition and morphology of the composites were characterized by FT-IR, XRD, and SEM. The FT-IR spectra of two series of composites, including PVDF/BTs and c-PVDF/PDA@BTs, are shown in Figure 2a and 2b. A comparison with the matrix revealed no particular new absorption peaks in PVDF/BT composites profiles, meaning that introduction of BT fillers did not alter the chain conformation or crystal form of the PVDF matrix. However, several obvious changes were found after PDA@BT were introduced into the crosslinked PVDF matrix. The absorption at ~1722 cm-1 assigned to the stretching vibration of C=C


7


double bonds on main chain of the matrix, as well as bands at ~1396 cm-1 assigned to the stretching vibration of C-H bonds were significantly shifted to ~1700 cm-1 and ~1372 cm-1, respectively. The latter might be induced by the strong coupling interaction between the -C-F bond of PVDF and -NH- or -OH groups on PDA@BT particles (-C-F⋅⋅⋅⋅H-N- or -C-F⋅⋅⋅⋅H-O-)

40

, resulting in stronger interactions

between the two phases (matrix and filler). The increased absorption intensity peak at ~950 cm-1 could be attributed to the deposition of PDA layer on BT 38. 0 wt%


0 wt%

2.0 wt% 3.8 wt% 5.0 wt% 7.0 wt%

PVDF/BT

a)

3500

3000

1722 1396

2500

2000

1500 1000 -1 Wavenumber (cm )

500

1722

2.0 wt%

1396

3.8 wt% 5.0 wt% 7.0 wt%

b)

3500

c-PVDF/PDA@BT 1700 3000

1372

950 1500 1000 500 -1 Wavenumber (cm ) 2500

2000

Figure 2. FT-IR spectra of polymer nanocomposites with different a) BT and b) PDA@BT particles contents.

PVDF PVDF/BT c-PVDF/PDA@BT

Intensity (a.u.)


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

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20

30

40

50

2θ (°)

60


70

80

8

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Figure 3. XRD spectra of polymer nanocomposites with 2.0 wt% BT and PDA@BT particles. Figure 3 presents the XRD spectra of the polymer matrix and nanocomposites filled with BT and PDA@BT. It will be noted that the polymer matrix mostly presented an amorphous phase. Where the diffraction signals in the composites finely coincided with those from the polymer matrix and neat BT ceramic. This indicated that, complexation between the PVDF matrix and either BT or PDA@BT caused no significant change in crystalline performance of the polymer matrix or BT fillers (Figure 2b).

Figure 4. Cross-section images of polymer nanocomposite with a) 5.0 wt%, c) 7.0 wt% BT particles, and b) 5.0 wt%, d) 7.0 wt% PDA@BT particles.

The cross-section morphology of two series of composites was observed through SEM, as presented in Figure 4. All samples showed rough surface morphologies with ravine-like structures, suggesting typical ductile fractures. The comparison with BT filled composites revealed that the surface coated particles presented much more evenly and separately dispersion states in PDA@BT filled


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composites. Meanwhile, unlike the distinct interface of BT particles with the matrix, the interface between PDA@BT particles and the matrix appeared indistinguishable, suggesting enhanced interactions between the two phases through the PDA coating (see Figure S2). The reason for that had to do with coating PDA onto BT, which did not only improve the dispersal uniformity of ceramics but also enhanced the reciprocal interaction between the two phases. The latter could explain the enhanced mechanical properties of c-PVDF/PDA@BT composites (more details in next section).

3.3. Mechanical properties of composites. In order to distinguish between the effect of surface modification of fillers and chemical crosslinking of polymer matrix on mechanical properties of the composites, three series of composites, including PVDF/BTs, c-PVDF/BTs and c-PVDF/PDA@BTs were prepared and studied. The stress-strain test results are illustrated in Figure 5. The increase in filler loading content raised Young’s modulus of all three composites. This was due to the contribution of ceramic particles on the modulus of composites. For example, Young’s modulus of PVDF/BT, cPVDF/BT and c-PVDF/PDA@BT with 7 wt% filler loadings increased respectively by 190%, 170% and 220% when compared with the pristine P(VDF-CTFE-DB) matrix. Furthermore, PVDF/BT samples showed that increased loading contents of ceramic particles continuously reduced their elongation at break (Figure 5a). This phenomenon had widely been observed in many inorganic particles filled polymer composite systems. However, simply crosslinking PVDF matrix should result in significantly increased elongation at break, as shown in Figure 5b. This could mostly be ascribed to the three-dimension network formed during crosslinking since crosslinked PVDF had larger elongation (of over 300%) when compared with uncured PVDF (less than 250%). Most notably, the elongation of c-PVDF/PDA@BTs remarkably improved if compared to those of PVDF/BTs and c-PVDF/BTs bearing consistent content of fillers (Figure 5c). Interestingly, the crosslinking of PVDF matrix led to continuously boosted elongation in the composites bearing modified BT particles, where surface modification of BT particles favored the


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improvements in elongation as well. The elongation of c-PVDF/PDA@BTs with 7 wt% PDA@BTs was even slightly larger than that of c-PVDF, which might be attributed to both the crosslinking of PVDF matrix and the strong intermolecular forces exerted between the matrix and PDA cell on BT. As a result, induced stresses at the break point of c-PVDF/PDA@BTs were the largest among three sets of composites. 100

a)

PVDF/BT 0 wt% 2.0 wt% 3.8 wt% 5.0 wt% 7.0 wt%

80 60 40 20 0

0

50

100

150

200

250

Tensile stress (MPa)

100

Tensile stress (MPa)

b) 80 60 40 20 0

300

c-PVDF/BT 0 wt% 2.0 wt% 3.8 wt% 5.0 wt% 7.0 wt%

0

50

100

150

200

250

300

Strain (%)

Strain (%) 100

c) Tensile stress (MPa)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


80 60

c-PVDF/PDA@BT 0 wt% 2.0 wt% 3.8 wt% 5.0 wt% 7.0 wt%

40 20 0

0

50

100

150

200

250

300

Strain (%) Figure 5. Stress-strain curves of a) PVDF/BTs, b) c-PVDF/BTs and c) c-PVDF/PDA@BTs with different particle loadings under room temperature. In general, the addition of inorganic particles may elevate the modulus and harm the flexibility of polymer matrix due to interfacial defects like voids and air bubble formed by foreign phases. However, thanks to synergistic effects caused by improved interfacial interactions between two phases and


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crosslinked 3D network, the modulus and elongation at break for PDA@BT filled composites could simultaneously be enhanced. This may offer a strategy to prevent inevitable elongation losses during the fabrication of inorganic particles filled polymer composites. Figure 6a presents the resistance of the composite films to a strong dissolving solvent of PVDF matrix, where PVDF/BT and c-PVDF/PDA@BT composites with 7 wt% fillers were submerged in acetone and stirred for 2 h. It can be seen that the PVDF/BT sample was completely dissolved, while the original shape of c-PVDF/PDA@BT film was well maintained during the process thanks to chemically crosslinked PVDF matrix. As indicated in Figure 6b, higher elongation allowed uniaxial stretching of cPVDF/PDA@BT in larger elongation, which was rather important for achieving high breakdown strength in dielectric composites.

Figure 6. Digital figures of polymer nanocomposite a) BT particles (left) and PDA@BT particles (right) in acetone, and b) PDA@BT filled composite film before (left) and after (right) tension. c) Schematic diagram of the composite.

3.4. Dielectric properties of composites. The frequency dependence of dielectric constant of


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two series of nanocomposites measured on a HP4284A LCR instrument at room temperature are gathered in Figure 7a and 7c. The dielectric constant of both composites gradually reduced as a function of frequency, which was rather similar to the changing tendency of the neat polymer matrix. The introduction of high-εr BT particles was the reason for the elevated permittivity at consistent frequency, which has been well investigated and finely predicted using a series or parallel models 21. If compared with BT filled composites, the composites filled with PDA@BT showed slightly lower dielectric constants at moderate frequency. For example, the permittivity at 100 Hz of composites containing 7.0 wt% PDA@BT filler reduced to 28 with respect to 30 obtained with BT filled composites. The latter might be caused by the strong interfacial interaction exerted between the two phases and chemical crosslinking across the polymer matrix, which, in turn, might depress the contribution from the interfacial polarization on the dielectric performance of the composites 41. 40

30 25

0.30 0.25

20 15 10 5


0.35

Tanδ

Dielectric constant

0.40


35

0.20 0.15 0.10 0.05

a)

0 2 10

3

10

10

4

5

10

b)

0.00 2 10

6

10

3

10

40 35

10

5

10

6

0.30 0.25

Tanδ

25

c-PVDF/PDA@BT 0 wt% 2.0 wt% 3.8 wt% 5.0 wt% 7.0 wt%

0.35

0 wt% 2.0 wt% 3.8 wt% 5.0 wt% 7.0 wt%

30

20 15 10

0 2 10

5

10

0.40 c-PVDF/PDA@BT

5

4

10

Frequency (Hz)

Frequency (Hz)

Dielectric Constant

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


0.20 0.15 0.10

c)

0.05 3

10

4

10

5

10

6

10

d)

0.00 2 10

Frequency (Hz)

3

10

4

10

10

6

Frequency (Hz) ACS Paragon Plus Environment

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Figure 7. Frequency dependence of dielectric constant [a), c)] and dielectric loss [b), d)] of polymer nanocomposite with different BT and PDA@BT particles contents under room temperature. It is well accepted that dielectric losses in polymer materials at a lower frequency (

Fabrication of Stretchable Nanocomposites with High Energy Density

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