Electroactive phases of poly(vinylidene fluoride)

Progress in Polymer Science 39 (2014) 683–706

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Progress in Polymer Science journal homepage: www.elsevier.com/locate/ppolysci

Electroactive phases of poly(vinylidene fluoride): Determination, processing and applications P. Martins a,b,1 , A.C. Lopes a,1 , S. Lanceros-Mendez a,b,∗ a b

Centro/Departamento de Física, Universidade do Minho, 4710-057 Braga, Portugal INL – International Iberian Nanotechnology Laboratory, 4715-330 Braga, Portugal

a r t i c l e

i n f o

Article history: Received 27 January 2013 Received in revised form 4 July 2013 Accepted 11 July 2013 Available online 20 July 2013

Keywords: Poly(vinylidene fluoride) Copolymers Piezoelectric Electroactive phases

a b s t r a c t Poly(vinylidene fluoride), PVDF, and its copolymers are the family of polymers with the highest dielectric constant and electroactive response, including piezoelectric, pyroelectric and ferroelectric effects. The electroactive properties are increasingly important in a wide range of applications such as in biomedicine, energy generation and storage, monitoring and control, and include the development of sensors and actuators, separator and filtration membranes and smart scaffolds, among others. For many of these applications the polymer should be in one of its electroactive phases. This review presents the developments and summarizes the main characteristics of the electroactive phases of PVDF and copolymers, indicates the different processing strategies as well as the way in which the phase content is identified and quantified. Additionally, recent advances in the development of electroactive composites allowing novel effects, such as magnetoelectric responses, and opening new applications areas are presented. Finally, some of the more interesting potential applications and processing challenges are discussed. © 2013 Elsevier Ltd. All rights reserved.

Contents 1. 2.

3.

Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Identification and quantification of the electroactive ˇ and !-phases of poly(vinylidene fluoride) . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.1. Fourier transformed infrared spectroscopy . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.2. X-ray diffraction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.3. Differential scanning calorimetry . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Processing of the ˇ-phase of poly(vinylidene fluoride) . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.1. From the ˛-phase of PVDF . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.2. From the melt . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.3. From solvent casting . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

684 684 685 686 686 688 688 689 690

Abbreviations: DMAc, N,N-dimethylacetamide; DSC, differential scanning calorimetry; FE, ferroelectric; FTIR, Fourier transformed infrared spectroscopy; ME, magnetoelectric; PAN, poly(acrylonitrile); pDDA, amphiphilic poly(N-dodecylacrylamide); PE, paraelectric; PLLA, polylactic acid; PLGA, poly(lactic-glycolic acid); PMMA, poly(methyl methacrylate); PVDF, poly(vinylidene fluoride); P(VDF-CTFE), poly(vinylidene fluoride-chlorotrifluoroethylene); P(VDF-HFP), poly(vinylidene fluoride-co-hexafluoropropene); P(VDF-TrFE), poly(vinylidene fluoride-trifluoroethylene); P(VDF-TrFE-CTFE), poly(vinylidenefluoride/trifluoroethylene/chlorotrifluoroethylene); PZT, lead zirconate titanate; PVC, poly(vinyl chloride); PEMFC, proton exchange membrane fuel cells; Tc , Curie temperature; Tm , melting temperature; XRD, X-ray diffraction. ∗ Corresponding author at: Centro/Departamento de Física, Universidade do Minho, 4710-057 Braga, Portugal. Tel.: +351 253 604073; fax: +351 253 604061. E-mail address: [email protected] (S. Lanceros-Mendez). 1 Equal contribution. 0079-6700/$ – see front matter © 2013 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.progpolymsci.2013.07.006

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P. Martins et al. / Progress in Polymer Science 39 (2014) 683–706

3.4. PVDF copolymers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.5. Composites: inclusion of fillers within the polymer matrix . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Nucleation of the !-phase of poly(vinylidene fluoride) . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.1. From the melt . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.2. From the ˛-phase . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.3. Composites: inclusion of fillers within the polymer matrix . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Applications and additional effects of fillers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5.1. Magnetoelectric composites . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5.2. Improvement of the piezoelectric response for sensor and actuator applications . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5.3. Development of polymer based membranes . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5.4. Electrolytes for battery applications . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5.5. Hydrogen production . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5.6. Biomedical applications . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Conclusions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Acknowledgements . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

1. Introduction Electroactive polymers are among the most interesting classes of polymers used as smart materials in various and numerous applications, such as sensors, actuators, energy harvesting and as biomaterials in the biomedical field, among others [1]. Amid the few polymers presenting piezo, pyro or ferroelectricity, such as Nylon-11 [2], polylactic acid (PLLA) [3] and poly(lactic-co-glycolic acid) (PLGA) [4], poly(vinylidene fluoride) (PVDF) and its copolymers have the best all-around electroactive properties, being therefore the polymer of choice for the increasing number of possible applications [5–8]. This semi-crystalline polymer shows a complex structure and can present five distinct crystalline phases related to different chain conformations designed as all trans (TTT) planar zigzag for the ˇ-phase, TGTG′ (trans-gauche–trans-gauche) for the ˛ and ı phases and T3 GT3 G′ for ! and ε phases [7,9,10]. Fig. 1 shows the most investigated and used for applications PVDF phases: ˛, ˇ and !-phases. Many of the interesting properties of PVDF, in particular those related with its use as sensor or actuator, are related to the strong electrical dipole moment of the PVDF monomer unit (5–8 × 10−30 C m) which is due to the electronegativity of fluorine atoms as compared to those of hydrogen and carbon atoms [7,11]. In this way, each chain possesses a dipole moment perpendicular to the polymer chain. The monomer units and therefore the dipolar

691 693 694 694 694 694 696 696 697 698 699 700 700 701 701 701

moments are then packed in a morphology which can show an overall dipolar contribution per unit cell as in the polar ˇ, ! and ı phases. Here, the ˇ-phase is the one with the highest dipolar moment per unit cell (8 × 10−30 C m) when compared to the other two phases [12]. The ˛ and ε phases are non-polar due to antiparallel packing of the dipoles within the unit cell [7,13,14]. Since the ˇ and ! phases are the most electrically active phases, their promotion within the material is an on-going pursuit [15] due to the strong interest in application areas such as sensors, actuators, batteries, filters, chemical warfare protection, magnetoelectric, and, more recently, in the biomedical field [16–19]. Different strategies have been therefore developed to obtain the electroactive phases of PVDF, mainly focusing on the development of specific processing procedures and the inclusion of specific fillers. Another important issue is that some of the reported results regarding the identification and quantification of both ! and ˇ-phases are contradictory: due to the similarity of the ˇ and !phase specific conformations, their characteristic Fourier transformed infrared spectroscopy, FTIR, bands and X-ray diffraction peaks typically used for the identification of the phases either coincide or are very close to each other, making difficult to distinguish among both phases [20–22]. Due to the increasing interest and large potential of this electroactive polymer, this review is focused on the determination of each electroactive phase, the strategies used to obtain them, and the additional effects of fillers, together with the use of some of them as PVDF nucleating agents. Finally, some of the most interesting and challenging applications will be outlined. 2. Identification and quantification of the electroactive ˇ and !-phases of poly(vinylidene fluoride)

Fig. 1. Schematic representation of the chain conformation for the ˛, ˇ and ! phases of PVDF.

The two phases of PVDF firstly discovered were the ˇ and ˛-phase, which are clearly identified by, for example, FTIR and X-ray diffraction [23]. However, the third discovered phase !-phase has caused some confusion in its identification and has been mistakenly reported as the ˇ-phase [22]. These mistakes persisted for a long time


685

and still happen in the present due the considerable less attention directed to the properties of the !-phase since, until recently, it was difficult to prepare single-polymorph high !-phase content films [24]. A careful interpretation of the results provided by FTIR, XRD and DSC are enough to identify the correct phase of PVDF [25–33]. However, the use of more than one technique is sometimes needed due to the superposition of the peaks of ˇ, ! and ˛-phases on each of the different techniques. The combination of the different techniques allows therefore the correct identification of the phases.

FTIR results are also commonly used to quantify the electroactive phase content of PVDF, however, this quantification is not performed in a single and uniform way. In [34] it is assumed that FTIR absorption follows the LambertBeer law and calculated the absorption coefficients, K˛ and Kˇ, at the respective wavenumber of 766 and 840 cm−1 . In this way, according to Gregorio et al. [34], the relative fraction of the ˇ-phase in a sample containing just ˛ and ˇ-PVDF is:

2.1. Fourier transformed infrared spectroscopy

where F(ˇ), represents the ˇ-phase content; A˛ and Aˇ the absorbance at 766 and 840 cm−1 ; K˛ and Kˇ are the absorption coefficients at the respective wavenumber, which values are 6.1 × 104 and 7.7 × 104 cm2 mol−1 , respectively. This expression has been widely used in diverse works, such as [35,36]. The quantification of the !-phase [33] has been more difficult, as exclusive FTIR bands corresponding to the !phase appear as shoulders. This makes them very useful for qualitative identification of the corresponding phase, but hinders quantitative classification, as the baseline correction would always include the other crystalline phases [37,38]. The small absorption coefficient of the band at 512 cm−1 makes it less sensitive and so, assuming that the ˇ crystalline phase is not present, the 835 cm−1 band has been used for !-phase quantification, despite its coincidence with ˇ-phase. For the calculation of the absorption coefficient of the ˛ and ˇ-phases, the non-saturated bands 762 and 1275 cm−1 were used, respectively. Further, an absorption coefficient for the amorphous phase was also used. The calculation of the phase content of the different phases was then based on the following equations:

The Fourier transformed infrared spectroscopy, FTIR, spectra of PVDF provides valuable information about its structure allowing to distinguish between the different crystalline forms. However, some bands are simultaneously common to the ˇ and !-phases [25] and some others are associated to the amorphous phase of the polymer [27]. Further, depending on the preparation conditions, the same film can contain one or more than one crystalline phase structures, fact that usually creates some qualitative and quantitative problems in the identification and quantification of the phases [26]. ˛-Phase of PVDF is most easily detected by FTIR absorption as it presents a large number of characteristic bands which are exclusive of this phase, such as the absorption bands at 489, 614, 766, 795, 855 and 976 cm−1 [26] (Table 1). As previously mentioned, the ! and ˇ-phases, due to the similar polymer chain conformation, show a similar number of bands and most of them appear at similar wavenumbers [26]. This is, for instance, the case of the bands at 512 cm−1 for the !-phase, that is very close to the band at 510 cm−1 for the ˇ-phase [25]. On the same way, some authors classify the strong peak at 840 cm−1 as a characteristic of the ˇ-phase [25,31] while others consider it common to both phases [32]. It has been recently accepted that the band at 840 cm−1 is common to both polymorphs but it is a strong band just for the ˇ-phase, whereas for the !-phase it appears as a shoulder of the 833 cm−1 band, characteristic of !-PVDF (Fig. 2b) [26,27]. However there are some distinguishable FTIR absorption bands. The bands at 431, 776, 812, 833 and 1233 cm−1 are exclusively of the !-phase, while the ones at 445 and 1279 cm−1 are exclusively ˇ-phase, the band at 1279 cm−1 appearing just as a shoulder. The characteristic bands of each crystalline phase are summarized in Table 1. Table 1 Absorption FTIR bands characteristics of ˛, ˇ and !-PVDF.

Wavenumber (cm

−1

)

˛

ˇ

!

408 532 614 766 795 855 976

510 840 1279

431 512 776 812 833 840 1234

F(ˇ) =

Aˇ (Kˇ /K˛ )A˛ + Aˇ

(1)

A762 = K˛762 · X˛ · t

(2)

A1275 = Kˇ1275 · Xˇ · t

(3)

835 A835 = (Kˇ835 · Xˇ + K!835 · X! + Kam (1 − Xtotal ))t

(4)

where Aj is the baseline-corrected absorbance at j cm−1 , j

Ki is the absorption coefficient at j cm−1 for the i phase, Xi is the mole fraction of the i phase, Xtotal is the total crystallinity, and t is the thickness in micrometers. The total crystallinity of each sample was calculated by DSC and the thickness of the film (2–9 !m) was calculated from the IR absorption band at 1070 cm−1 , which absorption coefficient is not dependent from the crystalline phase of polymer, and corresponds to: A1070 = 0.095t + 0.07. The values obtained for the absorption coefficients were 0.0259, 835 , K 762 , K 835 , 0.365, 0.150, 0.140 and 0.132 !m−1 , for Kam ˛ ! Kˇ1275 and Kˇ835 , respectively [33].

The two aforementioned methods were combined in order to calculate the crystalline phase content of !- and ˛phase PVDF when just these two phases are present in the material [20,39]. The resulting equation used is similar to Eq. (1) but the absorption coefficients are the ones referred in [33]. So, assuming that the crystalline phase of the

686


α γ γ +β

α

b)

γ

β

Transmittance (a.u.)


a)

γ α

β -1

840 cm

γ

833 cm-1

800

1000

1200

1400

-1

Wavenumber (cm )

810

840 cm-1

840

870

-1

Wavenumber (cm

)

Fig. 2. FTIR-ATR spectra of ˛, ! and ˇ-PVDF with identification of the ˛, ˇ and !-phase characteristics bands (a) and a detail of the ˇ and !-characteristic region commonly used in the literature for the identification of the phases (b).

polymer is either in the ˛ or !-phases, the !-phase content is given by: F(!) =

A! (K! /K˛ )A˛ + A!

(5)

where Aj is the baseline-corrected absorbance at j cm−1 , and K˛ and K! are the absorption coefficients at the respective wavenumbers, which values are 0.365 and 0.150 !m−1 , respectively. Further theoretical studies based on density functional theory [40,41] allowed the calculation of the vibrational spectra of ˛ and ˇ-PVDF chains with 5–15 monomer units. Additional characteristic vibrational modes were found that can be used to the identification of the ˛ and ˇ-phases. In particular, peaks at 1347, 930, 463 and 449 cm−1 are potential references for further identification of ˇ-PVDF. Despite all these developments, several problems may be encountered in specific samples in analyzing the FTIR spectrum since the preparation of an ideal single phase of PVDF is generally difficult due to its semicrystalline nature and the FTIR spectrum is dependent on molecular mass distribution, head-to-head and tail-to-tail defects, crystalline nature and orientation and thickness of samples [40]. Therefore, other experimental techniques, such as Xray diffraction and differential scanning calorimetry, have been explored for phase identification. 2.2. X-ray diffraction X-ray diffraction (XRD) results have been often also used to determine phases of PVDF [25,42,43]. Despite the different crystal structures presenting similar peaks, some of them are exclusive of each phase, allowing therefore their identification. Nevertheless, once again, the !-phase is the one that raises some identification problems as until recently, no samples containing exclusively the !phase had been processed, and therefore, no characteristic diffractogram of this phase had been obtained [25,29]. However, the comparison of the diffractograms of exclusive ˛-phase samples and ˛ + !-phase samples allows the

identification of diffraction peaks belonging exclusively to the !-phase [25]. More recently, samples of PVDF have been completely crystallized in the !-phase by the incorporation of clays [20] and the XRD diffraction peaks could be confirmed. ˚ is taken, When X-ray diffraction (K˛1 , # = 1.5405600 A) all ˛, ˇ and !-phases have an intense peak around 20◦ , but only ˛ and ! show other peaks close to 18◦ what makes them easily distinguishable from the ˇ-phase [29]. In this way, the ˇ-phase presents just a well define peak at 2$ = 20.26◦ relative to the sum of the diffraction at (1 1 0) and (2 0 0) planes [25,29]. In the same region, ˛phase presents more characteristic peaks at 2$ = 17.66◦ and 18.30◦ relative to diffractions in planes (1 0 0), (0 2 0) and (1 1 0), respectively. In addition, ˛-phase also presents a peak at 2$ = 26.56◦ corresponding to the (0 2 1) diffraction plane [25,29,44]. Finally, !-phase presents a superposition of peaks at 18.5◦ and 19.2◦ associated to planes (0 2 0) and (0 0 2), respectively and a more intense peak can be detected at 2$ = 20.04◦ corresponding the (1 1 0) crystalline plane. Similarly to ˛-phase, !-phase present a weaker peak at the region of 26.8◦ attributed to the (0 2 2) plane. Concluding, if, on the one hand, FTIR results has led to some confusion between ˇ and !-phases, and, on the other hand, XRD results has led to some confusion between ˛ and !-phase, the combination of these two techniques allows a perfect distinction among all the main phases of PVDF. The diffraction crystal planes and diffraction angles for each of the phases of PVDF are summarized in Table 2 and Fig. 3. 2.3. Differential scanning calorimetry Differential scanning calorimetry (DSC) is a thermoanalytical technique that has been complementary to the other identification techniques and also used for the identification of the crystalline phase of PVDF. Depending on the crystalline phase of PVDF, different melting peaks appear on the DSC thermogram. As the characteristics of the DSC peaks not only depend on the crystalline phase, but also on

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P. Martins et al. / Progress in Polymer Science 39 (2014) 683–706 Table 2 Diffraction crystal plane and angle of the different phases of PVDF.

ˇ-PVDF !-PVDF

2$

Crystal plane

17.66◦ 18.30◦ 19.90◦ 26.56◦ 20.26◦ 18.5◦ 19.2◦ 20.04◦

(1 0 0) (0 2 0) (1 1 0) (0 2 1) (1 1 0) (2 0 0) (0 2 0) (0 0 2) (1 1 0)

Fig. 4. Characteristics DSC melting temperatures of the ˛, ˇ and !-phases.

and therefore, DSC is not used to distinguish this two phases, but just to calculate the crystalline percentage of the film [20,45]. However, PVDF films crystallized in the !-phase show different endotherms when compared to ones present in the ˛ and ˇ-phases which also depend on the origin of the !-phase. In this way, when the !-phase is obtained by crystallization from the melt, the melting temperature is about 8 ◦ C higher than that of the ˛-phase, i.e., around 179–180 ◦ C [25]. On the other hand, when the !-phase is obtained from the ˛ → ! transformation, the melting temperature is about 18 ◦ C higher than the ˛-phase melting temperature, i.e., around 189–190 ◦ C [25,27], and is represented in this article by ! ′ . Some other small peaks or shoulder can appear when the sample is subjected to some specific thermal histories. For example, the quenching of the sample after isothermal crystallization at 166 ◦ C leads to the formation of very small ˛-phase spherulites (∼8 !m) which fusion occurs at 160 ◦ C [25] (Fig. 4). In this way, despite being an indicator of which phase is present in the polymer film, DSC should be just a complementary technique to FTIR or XRD as its features are not only dependent on the crystalline phase, but are also affected by crystalline defects which are particularly enhanced by the presence of fillers in polymer composites. On the other hand, this technique is particularly interesting for the assessment of the existence of ! phase PVDF. As a conclusion, the abovementioned experimental techniques allow the identification and quantification of the most relevant phases of PVDF, being necessary, in some cases, the combination of at least two of the techniques to clearly identify a specific phase (Fig. 5). FTIR allows to clearly distinguish ˛-phase, however, ˇ and !-phases bands are very similar. On the other hand, XRD show coincident diffraction peaks for the ˛ and !phases, allowing, on the other hand, a clear identification of the ˇ-phase. The combination of these two techniques

(110) (200)

β −PVDF

(002)

Intensity (a.u.)

(110) (020)

γ−PVDF (110) (020) (100)

(021)

α−PVDF

5

10

15

20

25

30

2θ Fig. 3. XRD patterns of ˛, ! and ˇ-PVDF with the identification of the corresponding diffraction crystal planes for each phase.

the characteristics of the morphology such as defects and crystalline size, among others, the existing literature does not allow to define a melting temperature for the different phases, but a temperature range. According to Prest and Luca [34], the expected melting temperature, represented by the endothermal peak, of the ˛-phase of PVDF is 172 ◦ C while Gregorio and Cestarini [34] refers that it occurs at 167 ◦ C. On the other hand, ˇ-crystallites present a melting temperature similar to the one for ˛-PVDF [25,27,34],

b

20.26

Intensity (a.u.)

β −PVDF

19.20 20.04

18.50 γ−PVDF

18.30 17.66

19.90 26.56

α−PVDF

5

10


a

β -PVDF

20

2θ

25

30

γ-PVDF

812 1234

833; 840

766

15

1279

840

855

800

α-PVDF

976

1000

1200

Wavenumber (cm -1 )

1400

c Heat Flow (Endo Up)

˛-PVDF

167-172 C

179-180 C

β-PVDF

γ-PVDF

α-PVDF

100

120

140

160

o

180

Temperature ( C)

200

Fig. 5. Synopsis of main experimental features of the different experimental techniques for correct identification of the PVDF phase: (a) XRD (K˛1 , ˚ (b) FTIR and (c) DSC. # = 1.5405600 A);

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allows the correct identification of the phases. DSC, can be used to confirm the results, in particular with respect to the existence of the !-phase. Along this review electroactive PVDF samples, will be referred as: (i) “from ˛-phase”, when obtained from ˛-PVDF; (ii) “from melting”, when obtained by subjecting PVDF to temperatures above 180 ◦ C; (iii) “from solvent casting”, when PVDF is dissolved and in its processing is not exceeded the temperature of 180 ◦ C. 3. Processing of the ˇ-phase of poly(vinylidene fluoride) The ˇ-phase is the one that presents the best piezoelectric, pyroelectric and ferroelectric properties [46] and is most commonly obtained by mechanical stretching of the ˛-phase [7,36,46], from melt under specific conditions such as high pressure [15,47,48], external electric field [18,49–52] and ultra-fast cooling [53,54]; from solution crystallization at temperatures below 70 ◦ C [19,34,55] or by the addition of nucleating fillers such as BaTiO3 [56], clay [35,57,58], hydrated ionic salt [59], PMMA [60], TiO2 [61] or nanoparticles such as ferrite [61], palladium [62] or gold [63] (Fig. 6). Also the development of PVDF copolymers such as poly(vinylidene fluoride-trifluoroethylene) (PVDFTrFE) has allowed to obtain the material in the electro active phase [64,65]. 3.1. From the ˛-phase of PVDF Taking advantage of a stretching mechanism, several authors have obtained ˇ-PVDF from initial ˛-phase PVDF

samples [13,30,36,46,66]. In this kind of ˇ-phase formation, the applied stress to the film results in alignment of polymer chains into the crystals so that an all-trans planar zigzag (TTT) conformation was inducted. Such mechanism allows the dipoles of the polymer chains to align normal to the direction of the applied stress [7]. It was reported that the maximum ˇ-phase content was achieved at a temperature of 80 ◦ C and a stretch ratio, R, of 5, but such samples still showed 20% of the original ˛-phase. Accompanying the phase transformation, an orientation of the polymer chains was observed. The stretching process also influenced the degree of crystallinity of the polymer. Additional poling of the samples also improved the ˛–ˇ-phase transformation [46]. Similar results were obtained in [7] in which the maximum ˇ-phase content was obtained at 90 ◦ C and at a stretch ratio, R, of about ∼4.5–5. It was found that R affected the ˇ-phase content more than stretching temperature. The most likely reason for this observation is the chain ordering and disordering of crystallites during film stretching and the effect of chain mobility in disordered regions to induce ˇ-phase into the crystals. The ferroelectric switching behavior and piezoelectric response of PVDF films prepared by drawing at R from 1 to 5 and temperatures from 80 to 140 ◦ C was reported in [67]. It was shown that R and temperature deeply influence the ˛–ˇ-phase transformation. Coercive electric field, remnant polarization and saturation polarization increased with increasing ferroelectric ˇ-phase content in the sample. In a similar way, samples with higher ˇ-phase content have shown higher d33 piezoelectric coefficients, in particular the maximum d33 value, 34 pC N−1 , was observed for the sample with higher ferroelectric ˇ-phase content (R = 5 at temperature 80 ◦ C). In the R = 5 samples at higher

Fig. 6. Methods for obtaining the ˇ-phase conformation of PVDF from melt and from ˛-phase through solvent casting and with the addition of fillers.


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temperatures, the ˇ-phase content decreases and, accordingly, the piezoelectric coefficient also drops. Since higher ˇ-phase content implies larger availability of dipoles able to be oriented, such variation of the piezoelectric coefficients is explained by the different number of oriented dipoles. 3.2. From the melt Even when widely used for applications, the piezoelectric activity of mechanically drawn PVDF film is not satisfactorily strong and not thermally stable, therefore, more active and thermally stable PVDF films are desired for practical use, for example, at higher temperatures [47]. By increasing temperature mechanically drawn ˇform of PVDF undergoes depolarization at temperatures far below the melting point and the Curie temperature. This effects is ascribed to the fact that the ˇ-form crystals in a mechanically stretched film are very tiny in size (about 10 nm thick) and include many conformational defects [68]. Further, the mechanical stretching is a processing step difficult/impossible to achieve when the material is directly deposited into devices or in micro applications. In this way, alternative methods have been reported; one of them being high pressure crystallization of PVDF. It has been proven that pressure quenching at high temperatures (280 ◦ C and 500 MPa) yielded PVDF of a new crystalline form having a melting point at 185 ◦ C, 27 ◦ C higher than that of the ˛-PVDF melt crystallized at ordinary pressures [69]. Later it was reported that the new crystalline form is a mixture of ˛ and ˇ-form, and that pressures higher than 500 MPa yielded only the PVDF ˇ form [48]. PVDF was also crystallized under high pressures exploring the pressure-temperature diagram phase [47]. Crystallization in a metastable hexagonal phase comprising extended chain lamellar crystals of a mixture of ˇ and ! forms was observed. Such mixture was later transformed into ˇ-form by poling at an elevated temperature (120 ◦ C). High pressures above 4000 kg cm−2 were also used to promote the ˇ-PVDF crystallization from melt [70]. Confirming the positive influence of high pressure in the ˇ-phase formation reported, Scheinbeim et al. [71] verified that increasing quenching pressures from 200 MPa to 700 MPa, lead to increases in the ˇ-phase content of the samples from 0% to 85%. It appeared that a final quenching pressure of ∼800 MPa would be enough to produce a sample containing 100% ˇ-phase using an initial pressure of 100 MPa. An initial pressure larger than 145 MPa was necessary for producing ˇ-phase if the final pressure was ∼630 MPa. The most important result reported in [71] is that the polymorphic form obtained by pressure quenching depends on which region the crystallization versus pressure curve is crossed by the melting temperature versus pressure curve (Fig. 7). That is, if a path crossed the ˛-phase crystallization curve (a–a′ ), the ˛-phase of PVDF was obtained, whereas if a path crossed the ˇphase crystallization curve (c–c′ ), it was obtained ˇ-phase. Finally, whenever a processing path crossed the intermediate region (b–b′ ), a mixture of both phases is obtained. Melt-crystallization of PVDF was also investigated in [54] by non-isothermal crystallization at ultra-high

Fig. 7. Plot of the melting and crystallization temperatures of PVDF, with the initial (a, b and c) and final states (a′ , b′ and c′ ) of the pressurequenched samples connected by gray straight line paths.

cooling rates ranging between 30 to 3000 K s−1 as well as at constant temperatures after quenching at 6000 K s−1 . An increase of the cooling rate above 150 K s−1 lead to the formation of the ˇ phase, observed by a low temperature shoulder in the crystallization exotherm in addition to the ˛ modification. At cooling rates above 2000 K s−1 just the low temperature exothermic peak was observed that is attributed to the crystallization of pure ˇ-phase. Through a systematic study of the effect of the quenching temperature on the polymorphism of PVDF, it was proposed that the appearance of the characteristic ˇphase vibrational bands, and therefore of the ˇ-phase of PVDF, strongly depend on the quenching temperature [53]. Quenching is a process in which PVDF homopolymer is strengthened and hardened. This is performed by heating PVDF films at a certain temperature, where the polymerization cannot occur and then rapidly cooling the polymer. During the annealing process, polymerization can occur only in certain local microscopic areas allowing PVDF molecules aggregate together to form the distinct phases [72,73]. In [53], films were either cast from cyclohexanone solutions on glass substrates or prepared by compression molded. The resulting two types of films were the deposited on a thin sheet of mica, melted at 210 ◦ C for 10 min and then quickly quenched in a bath in the temperature range between −50 and 70 ◦ C. The quenched samples were then annealed in an oven at 120 ◦ C for 50 h. There was no visible difference between the molded and solution-cast films. It was observed that when the PVDF quenching temperature is below 30 ◦ C, the samples exhibit essentially the ˇ-phase and the bands which correspond to the ˛-phase completely disappear [53]. When the quenching temperature is above 40 ◦ C, the bands assigned to the ˛-phase increased. For a quenching temperature of 70 ◦ C, the samples exhibited mainly the ˛-phase. A mechanism for the ˇ-PVDF phase formation in highrate quenching was proposed [74]: assuming that the ˇ-phase nucleation rate has a maximum at lower temperatures than that of the ˛-phase, the nucleation of the ˇ-phase was predominant at low temperature and that of the ˛-phase at high temperatures. For high-rate quenching, high-temperature nucleation was suppressed and nucleation proceeded mostly at low temperatures resulting in the formation of the ˇ-phase.

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The effects of quenching on crystallization of PVDF from the melt have been also studied in [75,76], on which the formation of the ˛-phase is the only result reported. A likely explanation for that observation is the occurrence of a skin-core morphology in the samples quenched when the thickness of the films increases, as the films in the present study have a thickness of ∼200 !m, larger that the previously reported (∼30 !m) [72,75,76]. In this way, it can be expected that a large amount of ˛-phase was formed at the core of the samples, while only a small amount of the ˇ-phase is produced at the surface [53]. Thin films in the ˇ-phase can be also generated directly on a substrate by evaporative deposition methods, which is interesting for microelectronic applications, however issues related to the thermal decomposition of PVDF molecular chains on vacuum heating (∼350 ◦ C) which cause a severe reduction in molecular weight, crystallinity, and the extent of ferroelectric response prevailed [77]. Thus, less destructive methods such as the electrospray technique [78], ultrasonic atomization [79] and the use of VDF oligomers [80] instead of PVDF to overcome the problems due to pyrolytic decomposition have been used [77]. Despite all these developments, the process of evaporative deposition remains still complicated and requires high vacuum and stringent conditions to control the morphology of the thin films [81]. 3.3. From solvent casting An alternative route for the formation of the ˇ-PVDF is the temporary application of electric fields. The application of high electric fields to PVDF that resulted in a contribution to the residual polarization which lied within the crystalline regions of the polymer was reported in [82] by Davis et al. In addition to the alignment of dipoles within the all trans polar ˇ-form, the polarization of the non-polar ˛form apparently occurred in two stages. At intermediate fields (near 1 MV cm−1 ), the large dipole normal to the chain axis apparently interacted with the applied field so as to align the dipoles parallel to neighboring molecular chains and create a stable polar phase with a little (if any) change in molecular conformation. At higher fields (near 5 MV cm−1 ), a crystal phase transformation to ˇform involving a change in molecular conformation was observed. A more recent approach, but which also stems from the use of high electric fields, is the use of the electrospinning method [18,49]. A schematic representation of this method is shown in Fig. 8 which presents the three main components of the electrospinning; a high voltage supplier (HV), a steel needle of small diameter and a metal collecting screen where the fibers are collected [18,83]. This method is appropriate and efficient when the objective is the preparation of submicron to nano-scale fibers and relies in an electrostatically driven jet of polymer solution or melt. In the case of PVDF fibers, their morphologies and the crystal forms were controlled simply through the adjustment of electrospinning parameters such as the applied voltage, flow rate, needle diameter and the distance from

Fig. 8. Schematic representation of the electrospinning method.

the needle to the collector [18]. It was reported by Ribeiro et al. [18] that the gradual increase in the applied voltage between 15 and 30 kV had as a result a decrease in the ˇ-phase fraction from 86% to 81%. The same study also revealed that the rotation speed of the rotation collector used for the fibers orientation had considerable effects on the ˇ-phase fraction. A slight increase in the rotation speed from 500 rpm to 750 rpm caused the increase from 45% to 85% in the ˇ-phase fraction. Consecutive increases in speed between 750 rpm and 2000 rpm had no significant effects on the ˇ-phase fraction. In this way, PVDF materials with submicron to nanoscale fibers with ˛ or ˇ or !-phase domination could be fabricated directly by electrospinning without the need of any post-treatment. It was also found by Zheng et al. [49] that adding a solvent with a low boiling point, such as acetone (56 ◦ C), decreasing the electrospinning temperature (from 60 ◦ C to 20 ◦ C), decreasing the feeding rate (from 75 !L min−1 to 5 !L min−1 ) and distance between the tip and the collector (from 20 cm to 10 cm) has as a consequence an increase in ˇ-phase content. The possible mechanism for the formation of ˇ-phase on PVDF produced by electrospinning is explained by the high voltage applied to the electrospinning solution and high stretching ratio, R, of the jets [82]. Remarkably, the high ratio of stretching during the electrospinning process is in some way similar to uniaxial mechanical stretching, which can cause the ˛ to ˇ-phase transition [7,36,46]. Analogously, the low crystallization temperature that arises from the low environmental temperature during electrospinning and the rapid evaporation of the solvent could all lead to the predominance of ˇ-phase of PVDF. The spin-coating, in addition to be an alternative to the evaporative deposition, is also interesting as readily results in the formation of the ˇ-phase of PVDF. Benz et al. [59] prepared PVDF thin films (2 !m thick) with a high ˇ-phase content on highly-polished single-crystal silicon wafers by spin coating PVDF solutions, with acetone/DMF as a solvent. The spin speed and the humidity conditions were


reported to be the main factors determining the ˇ-phase formation and the surface roughness of the thin films [59]. Nanoscale thin films of PVDF containing the ˇcrystalline phase were also directly prepared through heat-controlled spin coating without the influence of external stimuli either in the form of additives or posttreatments [77]. Sample preparation was carried out at different temperatures and it was found that at elevated temperatures (40–70 ◦ C), PVDF was crystallized into the ˇphase, while at near-ambient conditions (20–30 ◦ C) it was crystallized into the ˛-phase. The stability of spin-coatinginduced molecular stretching and the rate of solvent evaporation were the main factors that determined the crystalline structure and morphology of the films. The heat supplied from an external source facilitated the formation of flat and homogenous films with higher ˇ-phase content. In the same way, Cardoso et al. [84] reported a larger control in the fabrication of ˇ-PVDF thin films by spincoating with controlled thickness and morphology. Smooth and flat PVDF films with ˇ-phase content from 0 to 75% can be obtained just by varying spin-coating conditions and the post-thermal annealing temperature of the films. PVDF ultra-thin films were prepared by Jiang et al. [85] via Langmuir–Blodgett solution deposition technique. Results have shown that oriented PVDF ultrathin films can be prepared by Langmuir–Blodgett technique and contain poled ˇ-phase without poling treatment [85]. There are several advantages which make ferroelectric Langmuir–Blodgett ultrathin films particularly attractive candidates for molecular devices. The most important of these is that the sequential deposition of single monolayers enables the symmetry of the film to be precisely defined; in particular, layers of different materials can be built up to produce a highly polar structure. Secondly, the polarization of a Langmuir–Blodgett film is in situ “frozen in” during deposition, which avoids the poling procedure for device applications [85,86]. It was found [87] that such Langmuir–Blodgett (LB) deposition process leads to the direct formation of ferroelectric ˇ-phase in PVDF homopolymer ultra-thin films, in which the molecular chains were parallel to the substrates and the dipoles were aligned perpendicular to the substrates. Theoretical analysis and experimental results showed that the mechanism of forming the ˇ-phase and the dipole orientation were attributed to the hydrogen bonds between the PVDF molecules and water formed through the LB deposition. Taking advantage of the same LB deposition and assisted by amphiphilic poly(N-dodecylacrylamide) (pDDA) nanosheets a facile and novel method for preparing highly dense LB nanofilms of PVDF with precisely adjustable film thickness from several to hundreds of nanometers was developed [88]. It was shown that the polymer nanofilms without any post-processing comprised dominant ferroelectric ˇ phase of ∼95% and negligible paraelectric ˛-phase. Furthermore, through control of the surface pressure, controllable PVDF crystal morphologies were achieved with the ˇ-phase of PVDF dominating in all cases. The self-orienting behavior in the LB nanofilms indicated that the hydrogen-bonding interaction between PVDF and water molecules at the

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air/water interface played a key role in obtaining the ˇ-phase of the polymer [87]. The higher surface pressure for the deposition of the mixed LB nanofilm is thus more effective for obtaining higher contents of ˇ-PVDF crystals. Exclusively ˇ-phase films were also obtained from crystallization of PVDF from solution with N,N-dimethyl formamide or dimethyl acetamide at temperatures below 70 ◦ C [19,30]. Those films showed a high degree of porosity, which makes them opaque (milky) and fragile. The pore microstructure caused a degradation of the electrical properties of the material (lower dielectric constant) not allowing its poling, which is essential for the applications involving the piezo-, pyro- and ferroelectric effects. Further, the mechanical properties are also affected by the high porosity and the films cannot be oriented by stretching due to the large fragility [19]. In a later study, the porosity problem was solved through an original approach: after applying a pressure of 7.5 × 106 Pa, perpendicular to the surface of the films at elevated temperatures (140–160 ◦ C), the pores in the original sample were eliminated [19]. With respect to the previous porous films obtained by crystallization from solution, these PVDF films have shown better mechanical properties, are transparent and possess excellent flexibility. Such ferroelectric porous PVDF membranes with tailored interconnected porosity also have large potential for biomedical and energy applications as will be noted ahead in the review [89,90]. crystalline-state chain conformations of The isothermally solution-crystallized PVDF in N,Ndimethylacetamide (DMAc), and cyclohexanone was studied in [91]. DMAc with higher dipole moment produced PVDF films with larger trans states, ˇ and !-phases, while PVDF films with lower trans state were obtained using cyclohexanone, which has a lower dipole moment. According to the proposed mechanism, the proton exchange rate in DMAc was slower at low temperatures; hence, it provided larger intermolecular interactions between PVDF and DMAc. In cyclohexanone there were no variations in polarity, thus, no variation in PVDF chain conformation was observed. The largest content of the ˇ-phase was obtained at temperatures below 90 ◦ C. In fact, at lower temperatures the polymer coil dimensions expand and all-trans planar conformation in the PVDF chain is induced.

3.4. PVDF copolymers With the aim to improve the PVDF properties and to adapt it to the increasing technological demands, different copolymers of PVDF have been developed. Poly(vinylidene fluoride-Trifluoroethylene), P(VDFTrFE), is one of the most studied copolymer (Fig. 9). Contrary to PVDF, and in specific molar rations, it presents always the ferroelectric ˇ crystalline phase, once the addition of the third fluoride in the TrFE monomer unit with a large steric hindrance, favors the all-trans conformation and induces therefore the ferroelectric ˇ-phase independently of the used processing method: melt or solution casting. This situation occurs when the VDF

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Fig. 9. Schematic representation of the P(VDF-TrFE) repeat units.

Fig. 11. Schematic representation of the P(VDF-CTFE) repeat units.

content is between 50 and 80% [6], corresponding also to the ferroelectric behavior of the copolymer. Moreover, the high degree of crystallinity and the preferred orientation of well grown crystallites explain the higher remnant polarization present (∼110 mC m−2 ) when compared with PVDF, which, in turn, gives rise to a larger electromechanical coupling factor, k, that is translated into a higher efficiency in mechanical to electrical transformation [92]. Furthermore, this copolymer also shows the Curie Temperature (Tc) below the melting temperature (Tm), contrary to what happen in PVDF, and, in this way, allows the study of ferroelectric (FE) to paraelectric (PE) phase transition [93,94]. The Curie temperature (Tc) can vary from 55 to 128 ◦ C for an amount between 55 and 88 mol% of VDF. The correspondent enthalpy changes, %Hc has the value of 16 cal g−1 what can be explored in cooling device applications [92]. In order to enhance this effect, and make P(VDF-TrFE) applicable as a terpolymer, highenergy electron irradiation at elevated doses was used to destroy the stability of polymer ferroelectric domains, by the creation of defects. This will cause the transition of the ferroelectric phase to the relaxor state and paraelectric phase. Those effects are related to the increase in the amorphous phase fraction of the copolymer and the development of polar nano-regions, resulting in the increase of the electro-mechanical response of it [95–97]. Beyond the production of films, P(VDF-TrFE) copolymer can also be produced in the form of thin-films by spin coating, allowing a suitable control of sample thickness, ideal for the production of microstructures [98,99]. In the other hand, P(VDF-TrFE) can also be produced in the form of a membrane with a controlled microporosity that can be used, for instance, for lithium-ion battery applications [100,101]. Poly(vinylidene fluoride-co-hexafluoropropene), P(VDF-HFP), consist in the incorporation of the amorphous phase of hexafluropropylene on the PVDF homopolymer, as represented in Fig. 10. This copolymer has been mainly studied for applications on the area of polymer electrolytes of rechargeable lithium batteries and for the production of membranes for organophilic pervaporation. This interest

can be explained by the fact that P(VDF-HFP) is chemically inert and present a lower crystallinity when compared with PVDF, which is explained by the presence of the bulky CF3 groups [102]. The presence of ferroelectric properties in this copolymer is strongly dependent on the preparation method of the film. While slow cooled films do not present displacement versus electric field (D–E) hysteresis and so, no ferroelectricity, the opposite occurs for solvent-cast and quenched samples that present the typical ferroelectric behavior [102]. The highest values of remnant polarization, Pr , were obtained for solvent cast samples that reaches values up to 80 mC m−2 for 5% of HFP, value that decreases as the HFP content increases. Finally, the higher piezoelectric coefficient (d31 ) of 30 pC N−1 presented by this copolymer when compared to PVDF, makes it a promising material in some piezo and ferroelectric application areas, such as the development of magnetoelectric sensors and actuators [102–104]. PVDF has also been modified by the introduction of chloride trifluoride ethylene (CTFE) on the polymer chain, producing the P(VDF-CTFE) copolymer (Fig. 11), which final properties are dependent of CTFE content. In this way, the Tg can vary from −40 (Tg of PVDF) to 45 ◦ C (Tg of PCTFE) and the semicrystalline state is only obtained for CTFE content lower than 16 mol%, while an amorphous state is present for higher CTFE concentrations [105]. This copolymer also presents piezoelectric properties, higher electrostrictive strain response and higher dielectric constant when compared with PVDF polymer. The introduction of bulky CTFE makes the structure loose, which result in an easier orientation of dipoles under external electric field. The piezoelectric constant, d33 for this copolymer, for example, reaches the value of 140 pC N−1 [106]. Furthermore, the introduction of CTFE in PVDF-TrFE with different percentage of CTFE, P(VDF-TrFE-CTFE) (Fig. 12) allows the creation of materials with even lower Curie temperature than the one presented by PVDF-TrFE, and with similar level of electrocalorific effect [105,107]. Similarly to high-energy electron irradiation, the addition of chemical monomers like CTFE also creates defects in the P(VDF-TrFE) ferroelectric structure and, also in this case,

Fig. 10. Schematic representation of the P(VDF-HFP) repeat units.

Fig. 12. Schematic representation of the P(VDF-TrFE-CTFE) repeat units.

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P. Martins et al. / Progress in Polymer Science 39 (2014) 683–706 Table 3 Summary of some relevant electromechanical and dielectric properties of PVDF polymer its copolymers.

d31 (pC N−1 ) or longitudinal electrostrictive strain (*) d33 (pC N−1 ) k33 ε′

PVDF

P(VDF-TrFE)

P(VDF-HFP)

P(VDF-CTFE)

8–22 [111] −24 to −34 [114] 0.20 [92] [117] 6–12 [16]

12 [112] −38 [115] 0.29 [92] 18 [118]

30 [103] −24 [116] 0.36 [102] [119] 11 [119]

5.5% (*) [113] −140 [106] 0.39 [113] 13 [105]

the polymers are converted from a normal ferroelectric to a relaxor ferroelectric [108,109], what make them multifunctional due to their unique properties in comparison with other ferroelectric polymers, such as high electrostrictive strain and high dielectric constant [110]. A summary of PVDF and PVDF copolymer properties is presented in Table 3. 3.5. Composites: inclusion of fillers within the polymer matrix Some of the aforementioned methods used to obtain the ˇ-PVDF usually induce undesired structural deformations or microstructural limitations which may hinder specific applications such as electro-optical sensors and non-volatile memories [120,121]. Hence, alternative ways for obtaining the electroactive phases of PVDF have been developed. In recent investigations, the ˇ-phase of PVDF has been obtained by doping the polymer with fillers such as BaTiO3 [56], clays [35,57,58], hydrated ionic salts [59], PMMA [60], TiO2 [61], nanoparticles such as ferrite [61] palladium [62], gold [63] and carbon nanotubes [122]. Even when the nucleation mechanism cannot be universal, some important advances have been performed in order to unveil the main interactions driving the nucleation of the electroactive phases of the polymer by specific fillers. In this way, by designing experimental conditions in order to change the surface charge of the fillers [123,124], it was found that the piezoelectric ˇform of the polymer increased (until ∼90%) when ferrite nanoparticles with larger negative electrostatic charge were added, being this effect dependent on the amount of nanoparticles introduced into the polymer solution. The nucleation of the electroactive ˇ-phase in such nanocomposites obtained by melt processing was attributed to the interaction between the negatively charged particles and the polymer CH2 groups, having a positive charge density [124]. The obtained piezoelectric coefficient in such nanocomposites (∼33 pC N−1 )[17] are in agreement with the values reported for pure polymer with the same ˇ-phase content [67]. Starting also from nanocomposites obtained by melt processing, it was reported by Mendes et al. [56] that the electroactive ˇ-PVDF is nucleated by the presence of the BaTiO3 ceramic filler with a maximum ˇ-phase of ∼80%, being this effect strongly dependent on the filler size and almost independent of the filler content. Therefore, the nucleation of the ferroelectric phase should be strongly influenced both by geometrical factors due to the nanosize of the fillers and, in particular, by interactions at the interface between the local electric field around the nanoparticles and the PVDF dipoles. These

local field–dipole interactions can be related to the ones reported in [123]. After some studies regarding the processing of electroactive ˇ-phase PVDF with the addition of clays [57,58], PVDF–clay nanocomposites were prepared via one meltmixing process [35]. Formation of ˇ-PVDF has been observed in the nanocomposites with organically modified clays, irrespective of their percentage. Cast extrusion PVDF was used to produce nanocomposites films by Sadeghi et al. [125]. The obtained clay/PVDF composite was prepared by melt extrusion using a twin screw extruder equipped with a slit die. XRD results showed the formation of the electroactive ˇ-phase. Nevertheless, from all clay/PVDF composites reported so far, the highest ˇ-phase fractions (∼99%) were obtained for the phosphonium surfactant modified clay [35]. Changing the focus to ˇ-PVDF prepared by solvent casting with the addition of the fillers, similar mechanisms to the one in Fig. 13 [123] were also reported in montmorillonite/PVDF and graphite/PVDF nanosheets composites [126,127]. Similar ˇ-phase nucleation results have been recently reported in [62,63] in palladium-doped PVDF films and in gold–nanoparticle and gold–nanoshell doped PVDF respectively. The reason behind the ˇ-phase promotion in the palladium/PVDF composites was also described on the basis of the dipole–surface charge interaction model. In the gold/PVDF composites the ˇ-polymorph was more prominent with higher concentration of nanoparticles (0.5 wt.%), increasing also for the case of the gold nanoshells. Ongoing efforts in obtaining ˇ-PVDF include the crystallization of the polar ˇ phase when a hydrated salt was added to the polymeric solution [59]. However, simply

Fig. 13. Schematic representation of the ˇ-phase formation mechanism proposed in [14].

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adding water to a polymer solution does not induce the ˇphase presumably due to the evaporation of water. Thus, the proposed essential characteristic for depositing PVDF directly in the ˇ-phase from solution is to force water to remain in the drying film for a sufficiently long period of time. It was revealed that hydrogen bonding between the water and the polar C-F bonds was likely responsible for this phase control. The preparation of nano-TiO2 /PVDF composite films in acetone/DMF solution allowed to obtain the formation of the ˇ-phase and micro-fibrillar structures on PVDF. Such composites resulted in a considerable improvement of the mechanical properties, room temperature dielectric constant and electric breakdown field relative to those of pure PVDF. Due to these contributions, these kind of composites can 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 [128]. Finally, nucleation of ˇ-PVDF with carbon nanotubes has been also achieved with solution sonication and mechanical mixture approaches [122]. It was found that ˛phase coexists with the ˇ-phase in the composite prepared by sonicating the carbon nanotube/PVDF mixture solution, while no ˇ-phase can be observed in the composite prepared from the mechanical mixture route. With the help of density functional theory calculations, it was explained that a large amount of energy is required for the formation of the ˇ-phase on the carbon nanotube surface. Such required energy was obtained by the sonication process. The different forms for obtaining the electroactive ˇphase of PVDF are summarized in Table 4. 4. Nucleation of the !-phase of poly(vinylidene fluoride) The !-phase is also piezoelectric but this effect is weaker when compared to ˇ-phase PVDF due the presence of a gauche bond every fourth repeat C C units [20]. Similarly to what happens with the ˇ-phase, the !phase of PVDF is difficult to obtain excepting through melt crystallization at extremely high temperatures and at very slow cooling rates, with an ˛ or ˇ to ! solid state phase transformation by annealing at temperatures close to melt temperature or with the addition of fillers [19,20] (Fig. 14 and Table 5). 4.1. From the melt Kim et al. studied the crystalline properties of PVDF and its copolymer films prepared at different temperatures with subsequent slow cooling to room temperature [129]. !-Phase crystals were confirmed to be formed if the evaporation was performed at Tcrystallization < Tevaporation < Tmelting . Two types of !-crystallites were produced during isothermal crystallization of PVDF [130]. One of these was originated from the nucleation and growth of both low growth rate ˛-crystallites coexisting with high growth rate !-crystallites. !-Crystallites were also produced form a solid–solid phase transition from the already grown ˛ crystallites to the ! form (! ′ -crystallites) [24,130,131].

PVDF films with a large !-phase content were obtained after isothermal crystallization at ∼170 ◦ C for at least 20 h [24,132]. It was also found that after isothermal crystallization at 158 ◦ C for 15 h samples consists of 40.5% ˛-crystallites and 7.2% ! ones, but after isothermal crystallization at the same temperature for 95 h, the sample contains 11.2% ˛ crystallites and 34.8% ! ones [24,132]. All these works suggested that at a high crystallization temperature the growth rate of the !-crystallites is very low compared to that of the ˛ ones, and that the phase transition of the ˛ crystallites to the ! ′ -crystallites proceeds very slowly[130]. Therefore, it is difficult to obtain a !-phase rich film by isothermal crystallization. 4.2. From the ˛-phase The ! phase can be then formed from the skeleton of the ˛-phase morphology by a thermally induced solid-state ˛–! transformation [24,133]. When the ˛phase is annealed in the vicinity of the melting point, the vibrational energy of the crystal allowed conformational fluctuations to occur within the lattice and that those conformations associated with the higher-melting !-form were preferentially stabilized relative to the lower-melting trans-gauche–trans-gauche′ structure of the ˛-phase [24]. Lovinger et al. [134] reported that parts of PVDF ˛ spherulites undergo a transformation to the highermelting "-form when crystallized at high temperatures. Usually, this transformation is originated at the periphery of the ˛ spherulites where the lamellae are in contact with, and oppositely directed to, lamellae of !-spherulites. Those !-spherulites are formed only at high temperatures and at a slow linear rate, which increases with temperature. At very high temperatures ≥160 ◦ C, this transformation is also initiated at some ˛ nuclei; a few of the ˛-spherulites also exhibit areas of the high-melting phase irregularly dispersed within their interiors. In samples crystallized below ∼155 ◦ C, the transformation is of very limited extent, even after prolonged annealing at higher temperatures [24,134]. 4.3. Composites: inclusion of fillers within the polymer matrix In the same paper in which Prest et al. demonstrated the !-phase can be formed from the ˛-phase in the way discussed above [24], the authors also reported that certain surfactants modify the melting behavior of the ˛ and ˇ-phases in a manner that accelerates the rapid re-crystallization of the !-form. It was proposed that the organo-modified silicone copolymer L-520 surfactant delayed the diffusion of molecules in the melting process and thus increased the probability of the stabilization of conformational fluctuations that resulted in the !-phase formation. The greatest yield of !-phase crystals were obtained when surfactant-impregnated ˛ or ˇ-phase samples were heated slowly through their melting regions. This behavior was indicative of a melting and re-crystallization process. It was also reported that a !-form-rich PVDF film is isothermally crystallized by adding a KBr powder as a nucleating agent to the polymer–solvent system produced


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Table 4 Methods and features for obtaining the ˇ-phase conformation of PVDF. ˇ

Sample preparation

Observations

From the melt

Pressure quenching at high temperatures (280 ◦ C and pressures higher than 500 MPa) Pressure quenching at high pressure, 700 MPa, with an initial pressure of 100 MPa Non-isothermal crystallization at ultra-high cooling rates above 150 K s−1 Non-isothermal crystallization at ultra-high cooling rates above 2000 K s−1 Stretching mechanism (maximum ˇ-phase content achieved at 80 ◦ C and a R of 5) Additional poling of the samples also improved the ˛ to ˇ-phase transformation Application of high electric field

Pure ˇ-phase

From ˛-phase

From solvent casting

Spin-coating

Langmuir Blodgett deposition process

Copolymers

From solution with N,N-dimethyl formamide or dimethyl acetamide at temperatures below 70 ◦ C P(VDF-TrFE), by melt or solution casting

P(VDF-HFP), when prepared by solvent casting and quenched P(VDF-CTFE), stretched and well annealed Fillers

Ferrite nanoparticles BaTiO3 ceramic Organically modified clays Montmorillonite clay Graphite nanosheets Palladium nanoparticles Gold nanoparticles and nanoshells

Hydrated salts added to the polymeric solution Nano TiO2 Carbon nanotubes

Formation of partial ˇ-phase Formation of pure ˇ-phase

The samples still showed 20% of original ˛-phase; Many conformational defects.

Electrospinning relies on the use of high electric fields: it allows the production of sub-micro to nano-scale fibers with a ˇ-phase fraction up to of 86% without any post treatment. Spin speed and humidity conditions allows to control the ˇ-phase content (from 0 to 75%), the thickness and the morphology Direct formation of ultra-thin ˇ-phase films without any post-processing High degree of porosity and fragility. Degradation of the electrical properties. Formation of ˇ-phase independently of the processing method; Can be converted to relaxor if subjected to a high-energy electron irradiation Chemically inert; Lower crystallinity High electrostrictive strain response and piezoelectric response ˇ-Phase content ∼90% d33 ∼ −33 pC N−1 ˇ-Phase content ∼90% (dependent on filler size) Prepared by melt extrusion ˇ-phase content higher than 90% Prepared by solvent casting Prepared solvent casting Prepared by heat-controlled spin coating Prepared through solution mixing of Au-NPs or Au-NSs with PVDF; Interesting novel optical properties ˇ-Phase control by solvent water content; water induces surface roughness in the deposited film. Prepared in acetone/DMF solution Prepared by solution sonication and mechanical treatment

Table 5 Methods and features for obtaining the !-phase conformation of PVDF. !

Sample preparation

Observations

From the melt From ˛-phase

Melt crystallization at extremely high temperatures (∼170 C) ˛-Phase crystals are annealed in the vicinity of the PVDF melting point

Fillers

Organo-modified silicone copolymer L-20

◦

KBr powder Clay

Zeolite

Only 34.8% of !-phase In samples crystallized below ∼155 ◦ C, the !-phase content is very small, even after prolonged annealing at higher temperatures. Melting and slow recrystallization through polymer melting region; Produce at a relatively high temperature (165 ◦ C) for a short time (45 min) Melting at temperatures below 200 ◦ C and crystallization at room temperature Total crystallization on !-phase; Very high optical transmittance in the visible Melting at temperatures below 200 ◦ C and crystallization at room temperature Total crystallization of the !-phase; Increase on sample electrical conductivity.

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Fig. 14. Transitions to the !-conformation of PVDF from melt, from ˛-phase and with the addition of fillers.

at a relatively high temperature (165 ◦ C) for a short time (45 min) [130]. Although it was reported that the ˇ-crystallites preferentially grow on the KBr surface, this behavior must be attributed to the epitaxial growth of the !-form PVDF crystallites on KBr surfaces, since the work in [131] was performed under slightly different conditions and with a different PVDF grade. Most recently, the nucleation from melt of the electroactive !-PVDF at crystallization temperatures lower than 200 ◦ C was reported with the addition of two distinct fillers, Montmorillonite clay and NaY zeolite, respectively [20,39]. This nucleation effect was attributed to interaction between the negatively charged delaminated clays and the dipolar moments of the PVDF. Since clay has a higher contact area, it needed less quantity (0.25 wt.%) than zeolite (16 wt.%) to achieve the full crystallization of the electroactive !-phase. The piezoelectric d33 coefficient for a PVDF samples with 91 wt.% !-phase content was −7 pC N−1 . The different forms for obtaining the electroactive !phase of PVDF are summarized in Table 5. 5. Applications and additional effects of fillers In addition to the nucleation of the electroactive phases of PVDF, most of the fillers used for nucleation purposes also induce additional effects that bring added value to the use of PVDF nanocomposites for technological applications. Those effects will be discussed in the next sub-sections together with a brief presentation of the main challenges in these specific research fields. Although their piezoelectric and pyroelectric coefficients are inferior to those of ferroelectric ceramics, PVDF and PVDF based composites have advantages of

low permittivity, low thermal conductivity, softness and flexibility, good impedance matching to air and water, and relatively low cost. They have many major applications as actuators, in vibration control, ultrasonic transducers, batteries, filters, chemical warfare protection, magnetoelectric, and, more recently, in the biological field [16–19]. Their lesser-known applications include tactile sensors, ferroelectric devices, energy conversion devices, shock sensors, thermal and optical property measurement devices, pyroelectric infrared arrays, and dust sensors in interplanetary studies [135]. 5.1. Magnetoelectric composites Magnetoelectric (ME) materials have been attracting an increased interest due to its(their) innovative potential applications in data storage, switching, modulation, polarization, filters, waveguides, magnetic sensors (Fig. 15), transducers, spin wave generation, resonators and energy harvesting, just citing the most common [136]. In such ME materials the electrical polarization can be varied in the presence of an applied magnetic field or the induced magnetization can be modified with a presence of an applied

Fig. 15. Magnetic sensor with ME PVDF based nanocomposite [137].


electric field, allowing the applications mentioned above [137]. As the ME effect in single-phase ME materials is of difficult technological application due to the requirement of low temperatures and the weak ME coupling [138], the interest in such ME materials decreased until the emerging of the magnetoelectric multiferroic composite materials concept [139]. Most of the reported ME composites are based on piezoelectric phase ceramics such as barium titanate (BaTiO3 ) or lead zirconate titanate (PZT) due to their large piezoelectric coefficients. Such high piezoelectricity allows ME coefficients three orders of magnitude higher than in single phase materials [140], but, on the other hand, ceramic composites are fragile, limited by reactions at the interface regions, expensive, dense, brittle and can fail during operation [141,142] limiting in this way the incorporation into technological applications [61]. As regards the magnetostrictive phase components, Terfenol-D is the most used material due to its large magnetostriction, but its cost and brittle behavior are nevertheless the larger disadvantages of this alloy [140,143] fostering therefore investigation in novel magnetostrictive materials for ME applications. In this way, ferrite/PVDF based ME materials appeared as a reasonable solution to overcome the aforementioned problems [144]. In the literature it is only possible to find the use (P(VDF-TrFE)), as piezoelectric constituent of ME nanocomposites since the copolymers of PVDF, containing VDF contents between 55 and 82 mol percentage, contrary to what happens to the PVDF homopolymer, such copolymers crystallize from the melt in the ferroelectric phase which is an essential factor for the preparations of PVDF based ME nanocomposites [145,146]. Two main investigations can be found in the literature regarding ferrite/P(VDF-TrFE) ME nanocomposites. Martins et al. introduced CoFe2 O4 [147] and Ni0.5 Zn0.5 Fe2 O4 [61] ferrite nanoparticles into the polymer matrix and the resulting ferrite/P(VDF-TrFE) nanocomposites exhibit a ME effect that is dependent on the ferrite loading. The resultant ME films showed a maximum ME coefficient value of 41.3 mV cm−1 Oe−1 and 1.35 mV cm−1 Oe−1 in the case of the P(VDF-TrFE)/CoFe2 O4 and P(VDFTrFE)/Ni0.5 Zn0.5 Fe2 O4 , respectively. A significant ME voltage coefficient value around 40 mV cm−1 Oe−1 was also obtained in a similar CoFe2 O4 /P(VDF-TrFE) nanocomposites in [148]. With the reports [61,147] regarding the nucleation of the electroactive ˇ-phase of PVDF in ferrite/PVDF nanocomposites, it is possible to obtain the ME effect in the PVDF homopolymer. This finding will allow to reduce the production costs of the polymeric based ME nanocomposites since PVDF is substantially cheaper than P(VDF-TrFE) [149]. Such low-cost may allow significant and attractive applications such as didactics, toys and disposable devices [144]. Despite all these innovative developments in combining properties of fillers and polymers in the optimization of ME materials, there are still some aspects that require further attention and an intensive study. One important issue is the effect of nanofiller shape anisotropy in the ME response of the nanocomposite. Most of the existing studies report

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on nanocomposites with almost spherical filler nanoparticles. The use of anisotropic nanofillers may promote the emergence of new effects and the fabrication of anisotropic sensors and actuators. Other underexplored field lies in the creation of ME materials that can be used in applications that require simultaneously low and high magnetic fields either by the incorporation of distinct magnetostrictive nanoparticles in the same piezoelectric polymer matrix or by the use of multilayer systems with different magnetostrictive alloys. 5.2. Improvement of the piezoelectric response for sensor and actuator applications Since the discovery of piezoelectricity in the late 19th century, a great deal of interest has been shown in this electromechanical property [150]. Many materials including crystals and ceramics are under investigation and several are being used for piezoelectric devices for actual applications [151]. Piezoelectric ceramics such as piezoelectric lead zirconate titanate (PZT) and barium titanate (BaTiO3 ) and polymers such as PVDF have been widely used for sensor and actuator applications [152]. However, the current performance challenges are related to the use of such piezoelectric ceramics/piezoelectric polymers as flexible, simple fabricated and easy shaped smart materials [152,153]. In order to obtain those optimal material properties, some works have been previously dedicated to the study of BaTiO3 /PVDF composites and other similar filler/polymer systems. One of the main effects of BaTiO3 inclusion is the increase of the dielectric constant with increasing filler content [154], whereas the characteristics of the polymer matrix is just slightly affected as demonstrated by the small variation of the melting temperature of PVDF (160–158 ◦ C) with BaTiO3 inclusion (0 to 90 wt.%). Similar findings were also reported in [56], once it was verified that the dielectric constant of the composite increases significantly when nanosize BaTiO3 particles were used as a nucleating agents. This increase was particularly important for the smaller filler sizes and also shows strong concentration dependence. Mendes et al. [56] also stated, as previously mentioned in this review, that BaTiO3 can additionally stimulate the nucleation of the electroactive phase of PVDF (ˇ-phase fraction of 82%) and with that promotes the possibility of tailoring the piezoelectric response by proper tuning of the response of both polymer and filler to meet the requirements necessary to their use as specific sensor and actuators. Taking advantage of the features of BaTiO3 /PVDF composites, a recent study from Kakimoto et al. [155] presented a unique piezoelectric vibration harvesting structure composed of fibrous BaTiO3 ceramic fillers inserted into PVDF sheets. Such composite material was designed to produce large electrical response when an stretching movement was applied in the longitudinal direction, along fiber orientation. The measured piezoelectric response (d31 ) was enhanced to 14.8 pm V−1 in the PVDF sheet having a high orientation ratio (∼83%) of fibrous BaTiO3 filler, which was more substantial than the cases of the reference specimens

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using spherical BaTiO3 powders and fillers with low BaTiO3 orientation ratios (≈38%), 11.4 and 10.8 pm V−1 respectively. Furthermore, a highest energy output of 0.63 J m−3 per unit cycle, that was ∼26% higher value than pure PVDF sheet, was achieved at the frequency of 100 Hz for the sheet having highly oriented BaTiO3 . By using the same materials as in [155] but with different morphologies, Nunes-Pereira et al. [156] reported on the energy harvesting efficiency of electrospun PVDF, P(VDF-TrFE) and composites of P(VDF-TrFE) with BaTiO3 fillers deposited on interdigitated electrodes. It was found that the best energy harvesting performance was obtained for pure PVDF fibers, with power outputs up to 0.02 !W and 25 !W under low and high mechanical deformation respectively. The BaTiO3 /P(VDF-TrFE) composites show reduced power output mainly due to increased mechanical stiffness, the power output of the composites being higher for the non-piezoelectic smaller fillers (10 nm and 100 nm). The obtained values of power output, among the highest found in the literature, the easy processing and the low cost and robustness of the system, demonstrate its applicability for low power electronics. The piezoelectric response of PVDF was also increased with the addition of PZT and carbon nanotubes [157]. It was discovered that mixing a moderate amount of carbon nanotubes into the PVDF matrix increases the dielectric constant of the matrix (∼55) and induces the transformation of PVDF from the ˛-phase to the ˇ-phase. The piezoelectric response of the PVDF/PZT composite was improved up to ∼40 pC N−1 by dispersing a few carbon nanotubes (volume fraction of approximately 0.9%). The dielectric constant obtained in this tri-component composite (PVDF + PZT+ carbon nanotubes) being ∼140. There is an increasing need to correspond to the present demands of the piezoelectric sensor and actuator industry, such as the manufacture of miniaturized piezoelectric devices with low cost and power consumption and the fabrication of high-quality micrometer-size patterns onto a variety of hard and flexible substrates. These demands can be achieved by ink-jet printing [158,159]. Thus it is necessary the development of piezoelectric inks and printing optimization, in which fillers may have a predominant role, similar to that played before in the development of piezoelectric composites. It is also important to understand the nonlocal piezoelectricity of these systems. The characterization of surfaces and interfaces of piezoelectric nanocomposites is still controversial, so further conclusive studies are required [160], that will allow further improving material characteristics. 5.3. Development of polymer based membranes Membrane technology has gained a large importance in the last 30 years, competing with long established technologies for water desalination, food processing and emerging as an unique solution in medical applications [161–163]. Among several novel membrane materials, the emergence of the first developments in porous PVDF based membranes has attracted particular increasing interest due to their potential applications as filters [164]; as polymer electrolyte for applications in rechargeable

Fig. 16. Batteries application of Zeolite/PVDF nanocomposite [168].

batteries (Fig. 16) [165] (see next sub-chapter) and in biomedical applications [18,166]. The hydrophobicity (non-wettability) of the membranes, its porous structure and mechanical strength plays an essential role in the development of such applications [167]. Commercial microporous membranes such as PVDF are often prepared by the so-called immersion precipitation process [169]. In such process, the polymer solution is cast on a substrate and then immersed into a non-solvent coagulation bath to induce a series of liquid–solid and/or liquid–liquid phase separations [170]. After a drying procedure, the liquid phase is removed and a porous membrane is formed. It has been well established that the properties of the membrane such as the size and number of pores can be modified by changing the conditions of the casting dope or coagulation bath [169]. A number of strategies such as changing the casting temperature or dope dissolution temperature, the addition of non-solvent in the dope and the use of additives have been used to manipulate those properties [170]. Among these methods, incorporation of a filler in the casting dope as a blend is known to be very effective and uncomplicated when compared to the previous strategies [169,171,172]. In this way, besides nucleating the electroactive ˇphase, PMMA fillers can have an important role in the development of PVDF based membranes such as the one reported in [173] by Nunes et al. where the water permeability was increased and retention decreased with increasing PMMA content due to the rising porosity. PMMA has also been employed as a coupling agent for enhancing the interfacial compatibility between PVDF and polycarbonate in membranes formed from these two polymers [174–176]. Owing to their high compatibility, phase separation in the blends of PVDF and PMMA can be controlled to be within nanoscale, and in some cases, bicontinuous structures with nanophase domains can be obtained [176]. In the same way, results reported in [177] demonstrated that the addition of nano-TiO2 into the PVDF matrix strongly affects the properties of the membranes. Among all the prepared membranes, TiO2 /PVDF membranes containing 10 vol.% TiO2 exhibited hydrophilicity


(contact angle ∼59.7◦ ) higher than that of unmodified PVDF membranes (contact angle ∼82.8◦ ). The antifouling characteristics were thus improved in the hybrid membranes as proven by the smooth change in normalized permeate flux of bovine serum albumin when compared to the dramatically drop observed on the PVDF membranes. This fact is to be ascribed to the improved hydrophilic characteristics of the membranes. With the objective to prepare mixed matrix membranes, zeolites such as the zeolite 4A and Fe–zeolites have been also incorporated into the PVDF matrix [178,179]. In the case of zeolite 4A, the contact angle of composite membranes was reduced from 90◦ to 70◦ and the single gas permeability of the composite membranes increased with increasing inorganic filler content [178] reaching the maximum value of 9.65 Barrer in the PVDF composite with 32 wt.% of 4A zeolite to the He permeability. Regarding the use of Fe–zeolites, the contact angle and sorption measurements evidenced the higher affinity of the reacting solution (acetonitrile containing phenol) for the zeolite/PVDF membrane [179]. It was also reported that the use of different catalyst forms (free or entrapped within the polymeric matrix) in an open batch reactor evidenced the best performance in terms of selectivity for the entrapped catalyst (97% vs. 69%). It can be tempting to conclude that polymer based membrane technology has already reached the maturity, nevertheless many opportunities still exist. Efforts should continue on the development of high flux/low pressure composite membrane with enhanced salt rejection as well as improved resistance to fouling, chlorine, solvent, etc. Continuous improvements in composite membrane performances with respect to permeability, selectivity and stability will widen the applications of membranes to new areas [180]. One of the areas in which polymer based membranes have a well-defined potential are for the development of separator membranes (electrolytes) in Li-ion battery applications. Due to its specificity, they are presented in the next chapter. 5.4. Electrolytes for battery applications In recent years, there has been an increasing interest in the preparation of polymer electrolytes with high ionic conductivity, good mechanical strength and thermal stability, as these polymer electrolytes play a major role not only in lithium lithium/polymer ion batteries but also in other electrochemical devices such as super capacitors and electrochromic devices [181,182]. Such polymer electrolytes are expected to produce safe, high energy density and flexible lithium polymer batteries [182]. Although, some polymer electrolytes have been extensively studied, many of them present some problems regarding their practical applications. Poly(ethylene oxide) (PEO)-based electrolytes showed conductivity which ranges from 10−8 to 10−4 S cm−1 at temperatures between 40 and 100 ◦ C [183], poly(acrylonitrile) (PAN)-based electrolytes undergo severe passivation when in contact with lithium metal anode [184], PMMA polymer electrolytes lose their mechanical strength when they are plasticized

699

[182] and poly(vinyl chloride) (PVC)-based electrolytes offered low ionic conductivity at ambient temperatures [185]. Very recently, PVDF as a host has drawn the attention of many researchers due to its appealing properties. It has high anodic stability due to strong electron withdrawing functional group and has a high dielectric constant of ε = 8.4 that helps for larger ionization of lithium salt [182,186,187]. In this way, ionic salts, besides nucleating the ˇ-phase of PVDF can be incorporated into the polymer matrix in order to obtain high functional polymer electrolytes. Stephan et al. reported that among the lithium salts, LiBF4 , LiCF3 SO3 and LiClO4 the PVDF composite films containing LiBF4 offer highest conductivity (from 10−7 to 10−4 S cm−1 at 0 ◦ C), having as detrimental effect the severe passivation with lithium metal. Such membrane was found to be stable thermally up to 70 ◦ C [182]. The results from gas permeation data obtained in [172] revealed an increase in the mean pore size of PVDF membrane (from 20 to 50 nm), coupled with a more uniform pore size distribution as the amount of LiClO4 added was increased from 1 to 3 wt.%. An increase in coagulation temperature (from 25 to 50 ◦ C) was found to be advantageous in the production of a network membrane pore structure with higher permeation performance with aid of a sufficient amount of additive. However, over excessive additive would have an adverse effect, reducing its performance. The interaction between Li+ ions and fluorine in the PVDF matrix was studied in [188] and it was observed that the addition of such filler effectively disrupts the crystallinity of PVDF. Of significance result is the continuous increase of the ion conductivity (from 8.46 × 10−8 to 2.70 × 10−4 S cm−1 ) when the salt concentration is increased from 5 wt.% to 20 wt.%. For the use in portable electronics or electric vehicles, the demand of batteries with high charge/discharge rate, high power density, high ionic conductivity, long cycle life, and safe operation will increase. The successful design of new assembly technologies, the discovery of new polymer and polymer based composite materials will promote the development of the next generation of batteries, that allied to the flexibility of the polymer membranes will allow power sources for versatile-shaped electronic devices. Namely, replacing the polyethylene separator membrane of conventional Li-ion batteries with piezoelectric PVDF films or composites, the piezoelectric potential from the created by mechanical straining will act as a charge pump to drive Li ions to migrate from the cathode to the anode accompanied by charging reactions at the electrodes [189]. This novel approach can be applied to fabricate a self-charging power cell for sustainable driving of micro/nano-systems and personal electronics such as rollup displays, wearable electronic gadgets, and biomedical devices [190,191]. At last, for the use of polymer composite materials in battery applications some work is still needed. In particular, proper surface functionalization has to be achieve to improve compatibility between fillers and/or electrodes with the polymer matrix and therefore to improve the performance and stability of the membrane. Nevertheless,

700


polymer composites are one of the most promising ways to improve PVDF based separator membranes. The innovation and progress in this field is intimately related to the fabrication of multilayers or hierarchical pore structures of the polymer matrix in order to enhance the thermal, electrical, mechanical and electrochemical properties of the battery separators and to improve compatibility with electrodes [187]. 5.5. Hydrogen production In recent years, proton exchange membrane fuel cells (PEMFCs) have attracted more attention as an environment-friendly energy generation device [192]. Hydrolysis of chemical hydrides is becoming increasingly important as a potential in situ hydrogen supply method for PEMFCs [193]. Among chemical hydrides, sodium borohydride (NaBH4 ) is the most extensively investigated due to its high hydrogen storage density (10.8 wt.%), pure hydrogen production and safety in alkaline solution, where hydrogen can be generated from the catalytic hydrolysis of NaBH4 via a suitable catalyst, as described in the following chemical reaction [194]: NaBH4 + 2H2 O → 4H2 + NaBO2 + ∼300 kJ In order to accelerate the hydrolysis of alkaline NaBH4 solution, one catalyst should be utilized [195]. Among other catalysts, it was proved that PVDF nanofiber composite (CoCl2 /PVDF) showed excellent catalytic activity (the volume of hydrogen generated increasing from 100 ml to 325 ml for a 5000 s reaction time) and recycling stability (for 5 cycles) [196]. In a subsequent work [193] it was also proved that doping CoCl2 /PVDF composites with the Y-zeolite would effectively adjust the surface wetting properties of PVDF nanofiber and enhance its surface hydrophilicity since the contact angle decreases from 135◦ to 100◦ when the concentration of Y-zeolite is increased from 0 to 8 wt.%. This new composite nanofiber catalyst showed good catalytic activity in hydrogen production from NaBH4 hydrolysis, which exhibited higher hydrogen generation rate (1.63 ml min−1 ) than PVDF-based nanofibers without zeolite modification (1.38 ml min−1 ). Especially, this catalyst can be easily recovered and continuously used in applications when compared to conventional powderliked catalyst. Future improvements in the activity and stability of the catalyst are expected with the impregnation of CoCl2 into the pores of the Y-zeolite and adopting nano-Y-zeolite particles with smaller particle size [193]. 5.6. Biomedical applications Novel and interesting materials are continuously being developed with the objective of being used in biomedical applications. Among these, polymer based materials have confirmed to be an innovative and valuable choice as biomaterials in tissue engineering applications, smart prostheses, and sensors, among others [197].

A large number of polymers are widely used in various applications. This is mainly due to their availability in a wide variety of compositions, properties and forms (solids, fibers, fabrics, films, and gels), and the possibility of being fabricated readily into complex shapes and structures [198]. In the last years, the potential of electro-active polymers has been recognized for biomedical applications due to its ability to convert mechanical, thermal, or magnetic signals into electrical ones. It this sense, these materials can be used as smart scaffolds to stimulate cell growth and compatibility, biosensors, mechanical sensors, and actuators, among others [197]. From the short choice of electro-active polymers, including poly(lactic acid) and poly-(hydroxybutyrate), PVDF and its co-polymers are still the ones with the best electroactive performance, showing the highest piezo, pyro, and ferroelectric responses [5,10]. The possibility of tailoring PVDF properties and microstructure, allows new and challenging applications in the biomedical area, not only in device applications but also induce targeted cell responses [199]. Previous studies investigated the biocompatibility of PVDF films and showed that PVDF is a very promising material for biomedical applications [197,199]. In this way, the effect of the phase of PVDF and its polarization state on the fibronectin adsorption and cell response has been previously studied in [199]. It was shown that polarization of PVDF modified the conformation of adsorbed fibronectin at the material surface and therefore cell adhesion on the fibronectin-coated substrates. The in vitro attachment and metabolic activity of L929 cells on PVDF films shows that ˛-phase PVDF supports higher cell metabolic activity and cell spreading as compared to ˇ-phase PVDF [200], the nucleation of the distinct phases allowing therefore to tailor cell activity. The influence of the polarization of PVDF biological response of cells cultivated under static and dynamic conditions was recently reported by Ribeiro et al. [201]. It was observed that the MC3T3-E1 osteoblast cell culture exhibited different responses in the presence of charged PVDF films. The positively charged ˇ-PVDF films promote higher osteoblast adhesion and proliferation (∼500 cell mm−2 ), which is higher under dynamic conditions on poled samples, showing that the surface charge under mechanical stimulation improves the osteoblast growth. This interesting study showed that electroactive membranes and scaffolds can provide the necessary electrical stimuli for the growth and proliferation of specific cells. The interest in PVDF as a biomaterial was also extended to its composites. The addition of silver particles potentiated antimicrobial properties [202], zeolites opened the way to controlled drug release [39] and ferrite nanoparticles allowed the production of ME and multiferroic composites for sensors and cell stimulation [61,147]. Cell viability and proliferation in some of the aforementioned composites were performed in vitro both with Mesenchymal Stem Cells differentiated to osteoblasts and Human Foreskin Fibroblast 1 [197]. Taken together both in vitro and in vivo findings, it was concluded that for biological applications in direct contact with the cells zeolite


and clay composites are biocompatible materials promoting cell response and not showing in vivo pro-inflammatory effects, being therefore suitable for biological applications and tissue engineering. Other innovative uses of the electroactivity of PVDF and its nanocomposites in the biomedical field include highly sensitive flexible pressure sensors with microstructured rubber dielectric layers [203], hybrid nanogenerator for concurrently harvesting biomechanical and biochemical energy [204], tactile sensors for robotic applications [205], bio-photoacoustic detection [206], bioactuators [207] and biosensors [208,209]. Despite all these studies and innovative potential applications, this is an emerging field with important research needs for the increasing use of piezoelectric polymers and polymer nanocomposites as biomedical materials. In particular, it is important the identification of specific targets and respective ligands for these polymers since they are critical to obviate off-target toxicity, to establish the risk/benefit analysis in relation to long term safety polymeric nanocomposites, to get better insight into its in vivo behavior in humans and to develop proper theoretical animal models [210,211]. In particular, it is interesting to evaluate which specific tissues are more suitable to react to electromechanical stimulation. In the specific research field of the drug-delivery materials, there is the need to meet manufacturing (scale-up) and quality guidelines to produce consistently performing quality polymer based composite biomaterials [210]. Important topics of basic research in the near future also include sterilization and cleaning methods compatible with polymers and multifunctional materials and methods to reliably predict the performance of polymer-based biomaterials under specific conditions such as the ones found in living organisms. These concepts will enable and support innovations in the field of implantable piezoelectric polymer based bio-systems, among others. 6. Conclusions Poly(vinylidene fluoride), PVDF, and its copolymers are among the most challenging and interesting polymers for the development of advanced applications due to its largest dielectric constant, piezoelectric, pyroelectric and ferroelectric effects. Applications in the areas of biomedicine, energy generation and storage, filtration, sensor and actuators area being used and developed based on the fine control of the processing conditions to achieve the electroactive phase of the material. The main characteristics of the electroactive phases of PVDF and copolymers have been reviewed, the experimental techniques to identify them in a proper way, as well as the different processing strategies to achieve the desired materials characteristics. Recent advances related to the the development of electroactive composites have been presented and discussed together with some of the latest application possibilities in areas such as magnetoelectric materials, sensors and actuators, filtration membranes, Li-ion battery separation membranes, hydrogen production and biomedical applications, among others. All together, it can be concluded that PVDF, its copolymers and composites, mainly due to their

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