Multiscale-structuring of polyvinylidene ﬂ uoride for energy harvesting: the impact of molecular-, micro- and macro-structure of Materials Chemistry A REVIEW

Energy harvesting exploits ambient sources of energy such as mechanical loads, vibrations, human motion, waste heat, light or chemical sources and converts them into useful electrical energy. The applications for energy harvesting include low power electronics or wireless sensing at relatively lower power levels (nW to mW) with an aim to reduce a reliance on batteries or electrical power via cables and realise fully autonomous and self-powered systems. This review focuses on ﬂ exible energy harvesting system based on polyvinylidene ﬂ uoride based polymers, with an emphasis on manipulating and optimising the properties and performance of the polymeric materials and related nanocomposites through structuring the material at multiple scales. Ferroelectric properties are described and the potential of using the polarisation of the materials for vibration and thermal harvesting using piezo- and pyro-electric e ﬀ ects are explained. Approaches to tailor the ferroelectric, piezoelectric and pyroelectric properties of polymer materials are explored in detail; these include the in ﬂ uence of polymer processing conditions, heat treatment, nanocon ﬁ nement, blending, forming nanocomposites and electrospinning. Finally, examples of ﬂ exible harvesting devices that utilise the optimised ferroelectric polymer or nanocomposite systems are described and potential applications and future directions of research explored.


Introduction
Energy harvesting is a topic of intense interest where ambient sources of energy are harvested and converted into useful electrical energy. The concept of harvesting energy covers applications that produce relatively small levels of power (nW to mW) from mechanical loads, vibrations, human motion, waste heat, light or chemical sources. The power generated by the harvesting material or device is used for applications such as low power electronics, lighting or wireless sensor systems. The ability to use energy harvesting to produce autonomous selfpowered systems is of interest to reduce reliance on power cables or batteries that require regular replacement or recharging; there are also environmental benets from reduced battery usage.
Piezoelectric and pyroelectric materials are particularly attractive for a variety of energy harvesting applications. This includes the potential to convert mechanical vibrations into electrical energy via the direct piezoelectric effect and the conversion of thermal uctuations into electrical energy via the pyroelectric effect. Since both properties are oen present simultaneously, it provides a potential route to harvest energy from multiple sources so that the design of hybrid systems is possible.
A number of excellent reviews have been published in the area of piezoelectric and pyroelectric energy harvesting, which concentrate on nano-scale materials and devices, including 'nanogenerators', [1][2][3][4][5][6][7][8][9][10] and surveys of the various potential devices. [11][12][13][14][15][16] Li et al. recently presented a review on ferroelectric polymers for energy, including storage applications 17 and Fan et al. reviewed exible nanogenerators for harvesting and selfpowered electronics. 18 Prateek et al. 19 have overviewed recent progress on ferroelectric polymer based nanocomposite for energy storage and capacitor applications. The emphasis of this review is to describe mechanically exible systems based on polyvinylidene uoride (PVDF) polymers for energy harvesting, with an emphasis on manipulating the properties and energy harvesting performance of ferroelectric polymers and polymer nanocomposites through structuring the materials at multiple scales. This includes effort to tailor the ferroelectric, piezoelectric and pyroelectric properties of polymer materials by manipulating the polymer processing conditions, heat treatment, nano-connement, blending, forming nanocomposites, electrospinning and creating porous structures. PVDF is chosen in this review as it the most commonly used piezoelectric polymer due to its high piezoelectric coefficient compared with other bulk polymers. A comparison with other piezoelectric materials will be undertaken later in Table 1 and a detailed comparison of the range of piezoelectric polymers has been undertaken by Ramadan et al. 20 Before discussing the properties of ferroelectric PVDF-based polymers it is rst of interest to introduce the mechanisms of harvesting via the piezoelectric and pyroelectric effects so that the important properties for energy harvesting can be identied.

Piezoelectric and ferroelectric materials
Piezoelectric materials undergo a change in electrical polarisation when a mechanical stress is applied and this is termed the direct piezoelectric effect. This effect can create an electrical current in an external circuit and can therefore be exploited as an electromechanical generator to harvest energy. Piezoelectric materials also undergo a mechanical strain when subjected to an electric eld, this is termed the converse piezoelectric effect and is more relevant to actuator, acoustic emitter and vibration damping applications and will not discussed here.
The phenomenon of piezoelectricity is best introduced by considering a crystalline solid and its distribution of ions within an individual unit cell. The most common case is when the centres of the positive and negative charges are noncentrosymmetric within the unit cell planes when no pressure is applied; this is the case in ceramics such as barium titanate (BaTiO 3 ), lead zirconate titanate (PZT), aluminium nitride (AlN) and zinc oxide (ZnO). Of the 32 crystallographic classes that depend on the geometry and symmetry of the unit cell, 21 are non-centrosymmetric and all but one (due to other symmetry elements) exhibit piezoelectricity. It is the lack of symmetry in the distributions of ions in these crystalline materials that lead to the presence of an electrical dipole that leads to the piezoelectric response. Due to the lack of symmetry, piezoelectric materials for energy harvesting possess a well-dened polar axis and the application of stress relative to this polar axis is of importance for energy harvesting applications; the polar axis is oen termed the 3-direction or z-direction.
Ferroelectrics are a subgroup of piezoelectric materials, and are characterised by a spontaneous polarisation in the unstrained state and the capability to re-orientate the polarisation direction via an applied electric eld. The BaTiO 3 and the PZT family of ceramic materials are commonly used ferroelectric ceramics due to their good piezoelectric properties. The subject of this review is ferroelectric polymers, such as polyvinylidene uoride (PVDF) and its copolymers. Example copolymers include polyvinylidene-diuoride triuoro-ethane, P(VDF-co-TrFE), and we will see later that these are crystalline polymers with aligned molecular chains where the polarity is due to the alignment of polarised covalent bonds. The temperature limits of a ferroelectric are determined by the Curie temperature (T c ) and when BaTiO 3 is heated above T c its crystal structure is cubic and since the unit cell is symmetrical there is no polarisation and the material in no longer ferroelectric. Below the T c the unit cell is tetragonal and non-centrosymmetric, thereby leading to ferroelectric properties. A similar behaviour is observed in PVDF based materials where the ferroelectric b-phase converts to the paraelectric a-phase when the material is heated above T c .
While a ferroelectric contains many individual unit cells with a corresponding electrical dipole, there are regions where the unit cells have equal polarisation directions, these are known as domains. When a ferroelectric is cooled below the T c there is no net polarisation as the domains are effectively randomly distributed through the volume of the material. For the material to exhibit piezoelectricity the randomly orientated electrical dipoles are aligned in a common direction to achieve a net polarisation in a process called 'poling'. The process of electrical alignment, or 'poling', is essential in converting an inactive ferroelectric polymer into a useful electromechanically active material. Poling involves the application of an electric  57) eld, oen at elevated temperatures, to orientate the polar axis of the domains to those directions allowed by the crystal symmetry, which are nearest to that of the applied electric eld.
To further facilitate the poling process ferroelectric polymers oen require mechanical stretching to enhance the alignment of the molecular chains and these factors will be discussed later in the review. Fig. 1 shows a simple schematic of the poling process and the material initially consists of randomly orientated domains (Fig. 1a). The material is then heated to an elevated temperature below T c , and an electric eld is applied to orientate the domains (Fig. 1b). The material is then cooled to ambient temperature while the electric eld remains applied. On removal of the electric eld at ambient temperature, although some domains may change conguration (Fig. 1c), the material retains a net polarisation producing a poled ferroelectric with piezoelectric behaviour. Once poled, the material has a polar axis (the z-or 3-direction), which for the case in Fig. 1a-c is through the thickness of the material. The material is now piezoelectric and will exhibit the direct and converse piezoelectric effects.
A polarisation-electric eld (P-E f ) loop of PVDF is shown in Fig. 1d. At point 1 the material has no net polarisation, as the dipoles are randomly orientated (as in Fig. 1a). However, as the electric eld is applied the dipoles align with the applied electric eld to produce a polarisation at point 2 (as in Fig. 1b). Once the electric eld is removed, the PVDF has a remanent polarisation (P r ), where the material is similar to Fig. 1c and the material is piezoelectric. Clearly a high P r is of interest to achieve a high piezoelectric activity for harvesting applications. If an electric eld is applied in the opposite direction the polarisation can be returned to zero (point 4) and the polarisation direction can be reversed; which is dened as the coercive eld (E c ).

Direct piezoelectric effect for energy harvesting
The direct piezoelectric effect, where a force leads to the production of an electrical charge, can be used for energy harvesting. The presence of a polarisation with aligned dipoles and domains leads to the presence of a charge at each surface of the material and free charges, such as ions or electrons, are attracted to the charged surfaces of the material. This is shown schematically in Fig. 2a where the polarisation of the material is observed along with its bound surface charge. No current will ow under this stress-free condition if an electrical load it connected across the surfaces. The origin of the direct piezoelectric effect stems from the behaviour of the surface charge as the material is subjected to a stress that changes the polarisation level. If the application of a compressive stress leads to a decrease in the polarisation of the material, as in Fig. 2b, the free charges at the surface generate a current across an electrical load applied to the material. In Fig. 2c the application of a tensile load in the opposite direction increases the polarisation level and current ows in the opposite direction to balance the surface charge. If an alternating stress in applied to the piezoelectric the piezoelectric will generate an AC current through the load impedance.
The charge (Q) generated across the opposite faces of a piezoelectric material of area (A) when a mechanical stress (Ds) is applied between the faces is given by eqn (1). 22 where d 33 is the piezoelectric coefficient (C N À1 ); the 33-mode corresponds to the application of stress in the polarisation direction, which is the case in Fig. 2. Other modes can be utilised to harvest energy, for example in the 31-mode the applied stress is normal to the polarisation direction and d 15 is a shear mode. It is possible to determine typical voltages and current generated under an applied stress by considering two extremes of load impedance in Fig. 2. Under open circuit conditions, where the load impedance is innite, we can use the relationship Q ¼ CV, where C (¼A Â 3 T 33 /h) is the material capacitance to determine the voltage (V) from eqn (2).
where h is the thickness and 3 T 33 is the permittivity at constant stress in the polarisation direction. Since the energy stored in a capacitor in 1/2CV 2 , the energy (E) as a result of an applied stress is given by eqn (3). Therefore, for a given area and thickness the energy from an applied stress can be maximised by selecting materials with a high d 33 ; this is termed a harvesting gure of merit (FoM) to assess the performance of a material for vibration harvesting off-resonance. Under short circuit conditions, where the load impedance is zero, a current (I) will be generated and since (1) we arrive at eqn (4).
Measurements at both open circuit and closed circuit are commonly made in energy harvesting applications, although it should be noted at these conditions there is effectively no power since at open circuit conditions there is no current and at closed circuit conditions the potential difference is zero.

Pyroelectric effect for energy harvesting
While the piezoelectric effect is the generation of charge with applied stress, the pyroelectric effect is the generation of charge with a temperature change. As with piezoelectric materials, pyroelectrics are polar materials and exhibit a spontaneous polarization in the absence of an applied electric eld 23 and in ferroelectric polymers the polarisation is a consequence of the alignment of polarised covalent bonds. 24 As described in Section 1.2, the polarisation leads to the presence of bound charge on each surface of the material, as in Fig. 3a, and the origin of pyroelectric behaviour stems from the change in polarisation level with material temperature. 23 If a pyroelectric is heated (dT/dt > 0), as in Fig. 3b, there is a decrease in its polarisation as dipoles within the material lose their orientation due to thermal vibrations. This fall in the polarisation leads to a decrease in the number of free charges bound to the material surface; 23 see Fig. 3c. If the material is under an open circuit condition the free charges remain at the electrode surface and an electric potential is generated across the material. 25 If the material is under short circuit conditions an electric current ows between the two polar surfaces of the material. If the pyroelectric is then cooled (dT/dt < 0), as in Fig. 3b, the dipoles regain their orientation leading to an increase in the level of spontaneous polarization, thus reversing the electric current ow under short circuit conditions as free charges are now attracted to the polar surfaces; see Fig. 3d.
Eqn (5) denes the relationship between the developed pyroelectric charge (Q) due to a change in temperature (DT) for a material of surface area (A) with a pyroelectric coefficient (p). 26 Since current is dQ/dt, eqn (6) provides the short circuit pyroelectric current (i p ) as a function of the rate of temperature change (dT/dt) 27 with electrodes orientated normal to the polar direction.
The pyroelectric coefficient of an unclamped material, under a constant stress and electric eld, is dened by eqn (7),  where P s is spontaneous polarisation 28 and subscripts s and E correspond to conditions of constant stress and electric eld respectively. We have seen for vibration harvesting there is a need to dene the direction of applied load relative to the poling direction (e.g. 33-and 31-mode) and while the pyroelectric coefficient is a vector quantity, the electrodes that collect the charges are oen normal to the polar direction and so the measured quantity is oen treated as a scalar. 29 Under opencircuit conditions the charge created due to a temperature change, leads to a potential difference across the material, given by eqn (8).
Under open circuit conditions the material behaves as a capacitor and the energy stored in the pyroelectric is 1/2CV 2 , such that the energy (E) generated from a change in temperature is given by eqn (9).
Therefore, for a given area and thickness the energy for a temperature change can be maximised by maximising selecting a materials with a high p 2 3 T

33
; which is a gure of merit to asses the performance of a material for pyroelectric harvesting. This gure of merit has similarities with the vibration gure of merit d 33 2 3 T 33 which is not surprising given the similarity in the charge generation for both mechanisms result from a change in polarisation of the material as a result of stress, d ij ¼ dP s /ds (C N À1 or C m À2 /N m À2 ), and temperature, p ¼ dP s /dT (C m À2 K À1 ). Since there is a requirement for a pyroelectric material to be polar and exhibit a level of polarisation all pyroelectric materials are piezoelectric (pyroelectrics are a sub-class of piezoelectric materials so that all pyroelectrics are also piezoelectric). In ferroelectric pyroelectric materials, such as PVDF, the orientation, and sign, of the spontaneous polarisation can be switched by reversing the direction of the applied electric eld. These ferroelectrics are a sub-class of pyroelectric materials so that all ferroelectrics are both pyroelectric and piezoelectric.
Clearly for energy harvesting applications the gures of merit indicate that there is a need for high piezoelectric activity to maximise the piezoelectric coefficients (such as d 33 and d 31 ) and high pyroelectric coefficients (p); this also relates to a need for a high remnant polarisation (P r ) aer the poling process. Since the origins of the piezoelectric and pyroelectric effect originate from an electrical dipole from the polymer molecule there is a need to maximise the fraction of the non-symmetrical phase of the polymer and ensure it can be easily poled. The gures or merit also indicate a low relative permittivity is benecial for harvesting, although a large electrical dipole or ease of polarisation oen leads to a high permittivity so that the optimum choice of material based on a combination of d 33 , d 31 , p and 3 T 33 can be complex. Other factors include the Curie temperature (T c ) since it indicates the temperature at which the polymer transforms from a ferroelectric to paraelectric phase and indicates the temperature at which piezoelectric/pyroelectric properties are lost. The ferroelectric phase of polymers such as PVDF and its copolymers will now be discussed in more detail.

Piezo-and pyro-electric polymers
We have seen that ferroelectric and piezoelectric materials can be ceramic or polymeric. Electroactive polymers have obvious advantages over ceramics in specic applications in terms of their ease of processing at low-temperatures, low density, low stiffness, exibility and mechanical robustness, such as toughness and high strains to failure. There are also benets in terms of biocompatibility for implantable harvesters and sensors. 30 Polymer based piezoelectrics have found diverse applications in smart and multifunctional systems which include transducers, sensors, actuators, energy harvesting and storage devices, in the form of bres, foams, thin lms, textiles and coatings. Ferroelectric polymers that contain net molecular dipole moments in their macromolecular structure are of interest for energy harvesting and storage due to their high levels of polarisation. The molecular dipoles in these materials can be spontaneously polarised and oriented by the application of an electric eld, temperature variation, mechanical stretching, and via interactions with nanoparticles; 31 thereby leading to piezo-and pyro-electric response of the polymers. Polarised and electroactive polymers of interest include vinylidene uoride (VDF)-based uoropolymers, odd-numbered nylon, poly-L-lactide, polyurethane, and liquid crystal elastomers. The electroactive performance is highly dependent on their macromolecular structure which includes aspects of their crystalline structure, chain conformation, dipole orientation in crystalline regions towards the applied poling eld, 31,32 processing conditions and post-treatment methods. Compared to ferroelectric ceramics, electroactive polymers tend to have a lower permittivity and lower piezoelectric coef-cients. Table 1 shows typical properties of electroactive polymers along with common piezoelectric ceramics. Due to their relatively low piezoelectric activity, signicant effort has been undertaken to enhance the piezoelectric coefficients of polymers and tailor the permittivity, and the approaches used to improve the electrical properties include modifying the molecular structure, forming composites, controlling processing conditions and post-treatment methods such as mechanical stretching and electrical poling. In this review we will discuss the latest research progress and strategies employed to enhance the piezo-and pyro-electric properties of electroactive polymers and composites, in particular PVDF and its copolymers due to their high piezoelectric coefficient and ease of fabrication, with a particular emphasis on energy harvesting applications. structure and variety of crystalline forms. Table 1 indicates PVDF is a polymer with high piezoelectric coefficients compared to other polymers and the origins of the negative piezoelectric coefficients of PVDF in Table 1 has recently been reported by Katsouras et al. 58 The PVDF homopolymer contains 59.4 wt% uorine and 3 wt% hydrogen. 31 The presence of uorine atoms with a large van der Waals radius (1.35Å, versus hydrogen 1.2Å) and electronegativity in the polymer chain [-CH 2 -CF 2 -] induces a dipole moment perpendicular to the chain in each monomer unit. 32,59 PVDF has approximately 50% crystallinity, and exhibits ve different polymorphs: a (phase II), b (phase I), g (phase III), d and 3, which are related to the molecular chain conformations. Fig. 4a shows the a, b and g phases. 42 The a-crystal phase is hexagonal, with aligned polymer chains anti-parallel to each other, in the conformation of trans-gauche-trans-gauche 0 (TGTG 0 ). 42 The polar b-crystal is orthorhombic and has an alltrans planar zigzag conformation with their dipoles parallel to the b-axis, and contribute to the highest dipolar moment per unit cell (8 Â 10 À30 C m). 60 The phase transition from the ferroelectric b-phase to the paraelectric a-phase is dened as the Curie transition, which is a process that is highly dependent on the polymer chain structure, processing condition and posttreatment. Achieving a higher b-phase fraction in PVDF leads to higher piezo-, pyro-and ferroelectric properties. 31,32,34 The polar gand d-phase have the conformation of TTTTGTTTG 0 and TGTG 0 , respectively, which are also responsible for the piezo-and pyro-electric properties of PVDF, together with the b-phase. 31 2.1.1 Characterisation of ferroelectric polymers. The polymorphs of ferroelectric polymers are generally identied and quantied by using combined Fourier transform infrared spectroscopy (FTIR), X-ray diffraction (XRD) and differential scanning calorimetry (DSC) techniques. FTIR is an important approach to evaluate the polar b-phase content, a to b transition, as well as the dipolar orientation in the material. With regards to the FTIR spectra of PVDF, the vibrational bands at 530 cm À1 (CF 2 bending), 615 cm À1 and 765 cm À1 (CF 2 bending and skeletal bending) and 795 cm À1 (CH 2 rocking) refer to the aphase; the vibrational bands at 510 cm À1 (CF 2 bending) and 840 cm À1 (CH 2 rocking) correspond to the ferroelectric b-phase. The fraction of b-phase, F (b) , in the complete crystalline phase can be determined with eqn (10) 62 in which X b and X a are the crystalline mass fractions of aand bphase and A a and A b correspond to the absorbance intensities at 764 and 841 cm À1 , respectively. The K 841 (7.7 Â 10 4 cm 2 mol À1 ) and K 764 (6.1 Â 10 4 cm 2 mol À1 ) parameters are the absorption coefficients at the respective wavenumbers. 32 The polar g-phase has a similar chain conformation as the b-phase, and share similar characteristic peaks in the FTIR spectrum. The polymorph can be also identied by XRD. The characteristic peaks at 2q ¼ 17.6 , 18.3 , 19.9 , and 26.5 correspond to the (100), (020), (110), and (021) reections of the a-phase of PVDF. The (200) and (110) reections of the b-phase and (021) reection of the g-phase overlap at 2q ¼ 20.4 . 32,59 Differential scanning calorimetry (DSC) is oen used to study the Curie transition and crystallisation behaviour of polymers during a change from the ferroelectric b-phase to the paraelectric a-phase. Measurements are generally conducted in the temperature range of 0 to 220 C for PVDF and its copolymers, to understand the relationship of heat ow and temperature during non-isothermal and isothermal crystallisation processes. A value of 104.6 J g À1 was used as the heat of fusion of the perfect crystalline a-phase of PVDF. 63 The efficiency of nucleation agents can be evaluated using eqn (11); 64 where T c-nucl , T c1 , T c2 are the crystallization temperatures of the nucleated, non-nucleated and self-nucleated polymer, respectively. The melting temperature and crystalline temperature of polymers are dependent on the molecular structure, molecular weight and any annealing process. The surface morphology, chain alignment and dipole orientation, ferroelectric properties have been characterised by FTIR-grazing incident reection absorption spectroscopy, grazing incident wide angle X-ray diffraction, 65 atomic force microscopy, small-angle neutron scattering, 66 dynamic contact electrostatic force microscopy, and polarization-electric eld hysteresis measurements. 67 Measurements of the ferroelectric and piezoelectric properties can be undertaken by polarisation-electric eld measurements, piezo-response force microscopy and second harmonic generation microscopy.
2.1.2 Formation of ferroelectric b-phase. A variety of methods have been developed to increase the fraction of the ferroelectric b-polymorph in PVDF, including co-polymerisation with a second monomer, blending with other polymers, forming composites with nucleating agents or nanollers, tailoring the processing conditions, mechanical stretching and electrical poling. The formation of the non-polar a-phase of PVDF is kinetically more favourable which can readily form by crystallization from the melt at moderate or high levels of supercooling 68 (<160 C), or from solution-crystallisation of a xylene/ acetone mixture, monochlorobenzene, or dimethylformamide (DMF) solutions. 62 While the polar b-phase is the most thermodynamically polymorph, it can only be formed under special conditions. These include (i) melt crystallization at high pressure 65 or very high cooling rates, 68,69 (ii) solution-casting from highly polar solutions, such as hexamethyl phosphoramide, 70 (iii) vapour deposition of oligomeric PVDF 71,72 or (iv) by blending with surface charged nanoparticles or nucleating agents. The aphase can transform to the b-phase by mechanical-stretching and electrically poling at electric elds as high as 100 kV mm À1 under elevated temperature of 80-165 C. 31 From rst-principle simulations and a generalised solid-state nudged elastic band method, 73 it was suggested that the electric eld transforms the a-phase to the polar d-phase, and then to the b-phase. This is shown in Fig. 4b where the a-phase forms from the melt and progressively transforms to the b-phase at high electric elds. The inuence of cold drawing and annealing at high pressure to form the b-phase is also shown in Fig. 4b. 61 The crystalline structure of PVDF has been studied during a micro-injection process and in the presence of polar additives. 74 In contrast to crystallisation in the static state, the application of high stress during micro-injection promotes the formation of b-phase. The presence of polar additives also aids in the stabilization of the metastable TT conformation (as shown in Fig. 4b) via an interaction between the polar groups and the pCF 2 groups of PVDF, which benets the formation of b-PVDF. As a result, PVDF micro-components with rich b-phases have been injection-moulded under the combined effects of an externally applied force and polar group interaction.
The a to b transformation is highly dependent on the stretching rate and temperature. By stretching PVDF blown lms biaxially or uniaxially, 75 a high 86.5% content of highly oriented b-phase was achieved during drawing at 87 C, with a drawing rate of 50 mm min À1 and a stretch ratio of 6.5. At this stretch ratio, a high d 33 piezoelectric coefficient of 33 pC N À1 was obtained. Other studies have shown that the b-phase is more favourably formed in the temperature range of 70-100 C and at a stretch ratio of 3-5. 76 Higher stretching temperatures reduce the efficiency of the phase conversion to b-phase and the conversion into b-phase only takes place for stretch ratios above 5. It has also been shown that annealing of the PVDF thin lms at 90 C transforms the g-phase to the b-phase and this is re-ected in a decrease of the relative permittivity. 77 Recently, nanofabrication technologies such as electrospinning, nanoimprint lithography 78 and the horizontal Langmuir-Schaefer technique 79 have attracted attention to produce non-volatile memory and exible piezoelectric sensors. In particular, electrospinning can produce a high b-phase content in PVDF nanobers via a one-step process that avoids additional processing stages such as mechanical stretching or electrical poling; this process will be described later in Section 2.2.8. In the following sections, the latest progress in understanding the effects of co-polymerisation, processing condition, post-treatment, addition of nanoparticles and electrospinning on the piezo-and pyro-electric properties of PVDF and its copolymers are discussed. Harvesting devices exploiting these optimised properties are then described.
2.2 Approaches to enhancing b-phase formation 2.2.1 Copolymerisation and phase transition. Copolymerisation is an effective method to tune the polymorph structure and phase transition behavior of PVDF. The comonomer type and composition ratio determine the crystalline structure of the polymer and degree of crystallinity which directly affect the ferro-, piezo-and pyro-electric properties. A variety of VDF-copolymers and terpolymers have been synthesized by incorporation of triuoroethylene (TrFE), chlorotri-uoroethylene (CTFE) or hexauoropropylene (HEP) comonomer. The molecular structures of these commonly studied monomers and the copolymers are shown in Fig. 5 and typical properties have been shown in Table 1 for comparison.
The introduction of a bulky co-monomer to PVDF chains generally facilitates the formation of ferroelectric b-phase due to a steric hindrance effect. 80 For example, for P(VDF-co-TrFE) in Fig. 5a the introduction of a third uoride atom in the TrFE [-CHF-CF 2 -] co-monomer unit forces the polymer chains to align in an extended planar zigzag all-trans conformation below T c when the TrFE content is over 11 mol%. 81 When the TrFE is above 20 mol%, the formation of b-phase is independent of the processing conditions and electric poling. 82 For P(VDF-co-HFP), shown in Fig. 5b, the presence of bulky -CF 3 groups in the PVDF chains provides more space to allow dipoles to re-orient. With a HFP content of 5 mol%, a high value of remnant polarization (P r $ 80 mC m À2 ) was achieved for solvent cast lms, resulting in a high d 31 piezoelectric coefficient of 30 pC N À1 . 83 In the case of the P(VDF-co-CTFE) copolymer, Fig. 5c, a CTFE content lower than 16 mol% (ref. 42) led to a d 33 piezoelectric coefficient as high as 140 pC N À1 . 43 Therefore, the advantages of PVDF copolymers over PVDF is that the b-phase is always present regardless of processing methods, which can be attributed to the steric hindrance of the co-monomer that expands the interchain distance and reduces the activation energy for the a to b phase transition. 32,84 Additional annealing, mechanical stretching or electrical poling processes can be used to increase the degree of crystallinity and further align the CF 2 dipoles, which leads to higher piezo-and pyro-electric properties than PVDF homopolymers. By introducing a chloro-containing third monomer termonomer, such as chlorotriuoroethylene (CTFE), to P(VDF-co-TrFE) copolymers, as in Fig. 5d, some of the ferroelectric b-phase transforms to the g-phase when the termonomer units reach 7 mol%. 85 As a result, the T c of the mixed ferroelectric phase polymer was reduced to almost ambient temperature and exhibits a high relative permittivity (3 0 > 70), a slim polarization hysteresis and ferroelectric relaxor behaviour.
For ter-polymers, such as poly(VDF-TrFE-TFP), the termonomer 3,3,3-triuoropropene (TFP) with a bulky side group leads to -CF 3 behaving as crystalline defects in the polymer chains and this results in small crystals and a small degree of crystallinity. This is reected by a T c that reduces from 72 C for a conventional P(VDF-co-TrFE) copolymer to 65 C for the terpolymer. 86 The TFP is mainly an amorphous phase and the bulky -CF 3 groups hinders the mobility of the chain segments, leading to an increase of the glass transition temperature (T g ) from approximately À20 C for a standard copolymer to 0 C for a terpolymer with 9 mol% of TFP. The effect of the ter-monomer on the polarization behaviour of copolymer thin lms (20 mm) were compared for P(VDF-co-TrFE) (65/35), poly(VDF-TrFE-CTFE) (61/26/13) and poly(VDF-TrFE-TFP) (62/32/6). As shown in Fig. 6, the poly(VDF-TrFE-CTFE) exhibited a typical relaxor ferroelectric behaviour with a slim hysteresis loop and low coercive eld (E c ¼ 13 MV m À1 ), remnant polarization (P r ¼ 10 mC m À2 ) and saturation polarization (P sat ¼ 58 mC m À2 ). Both P(VDF-co-TrFE) and poly(VDF-TrFE-TFP) exhibited wide ferroelectric hysteresis loops with a high coercive eld, E c ¼ 63 MV m À1 , indicating that both polymers have a similar polarisation response. The poly(VDF-TrFE-TFP) has a lower P r ¼ 33 mC m À2 and saturation polarisation, P sat ¼ 49 mC m À2 , as compared to P(VDF-co-TrFE) copolymer, which may be due to the lower crystallinity of the terpolymer and the lower polarizability of the -CF 3 group of the TFP compared to the CFE and CTFE units. 86 2.2.2 Processing conditions and inuence on Curie temperature. As shown in Fig. 4b, the paraelectric a-phase is generally formed in PVDF on cooling from the melt or from solution-cast lms. A mixture of a-, b-, or g-phases is generally observed in PVDF copolymers when crystallised from the melt or from solutions under a range of conditions. The conversion of the ferroelectric b-phase to the paraelectric a-phase takes place thermally when the material is heated above the T c , which denes the upper use temperature for piezoelectric and pyroelectric applications such as energy harvesting. 31 The T c of b-PVDF is typically $170 C (ref. 87) and the introduction of a co-monomer to the polymer chains can restrict the degree of crystal growth and form nanosized crystalline domains, 61,85 which increases the inter-chain distance and dipolar mobility, and decreases the T c . For example, a low T c of 105 C was measured for a copolymer of 70/30 composition ratio. 88 A previous study 89 demonstrated that T c is 70-80 C for a 65/ 35 mol% P(VDF-co-TrFE) copolymer, which was observed as a decrease in the intensity of the XRD diffraction peak at 2q ¼ 19.5 (b-phase), and the emergence of a second peak at 2q ¼ 18.2 (a-phase) as the temperature increased above 80 C. The all-trans peak vanished at 100 C (lower than the melting temperature, T m $ 153 C). In the temperature between T c and T m , the polymer chains move more freely to reorganise into higher degree of crystallinity. In comparison, a copolymer of 73/ 27 mol% had a T c of 100-110 C and the all-trans ferroelectric peak disappeared on heating to 140 C, while T m was 150 C. For a 78/22 mol% copolymer, the T c was increased further to 130-145 C and T m was estimated to be 149 C. 89,90 For a 81/19 mol% copolymer, the T c was $120 C, and the T m was near 150 C. 81  During the manufacturing process, polymers suffer from tension, shearing and compression during melt-extrusion, injection moulding and compression moulding; and tension and electrical poling during the electrospinning process. In addition to mechanical stretching or electrical poling, the polar b-phase of PVDF can also be induced by application of a shearing history (shear rate, shear strain and shear temperature). 91 A recent study found an increase in b-phase content from 38 to 84% for PVDF 120-150 mm thick lms when sheared at 220 C with an applied shear rate of 10 s À1 for 10 s, followed by isothermal crystallization at 155 C. The samples sheared at a higher temperature generally had a higher nuclei density, smaller spherulites and a higher b-phase fraction, as shown in Fig. 7. A large deviation between the shearing temperature and crystallization temperature was shown to facilitate the formation of b-phase.
2.2.3 Inuence of annealing conditions on phase transition. The applied annealing treatment affects two important phase transition temperatures, namely the Curie temperature and melting temperature. The T c of PVDF copolymers relies on the chemical composition of the copolymers, and is also determined by the polymer annealing conditions. Copolymers with 50-80% VDF show a T c in the range of 70-140 C, 92 which is absent in pure PVDF.
Generally, a higher annealing or crystallisation temperature and/or longer annealing time above T c yields lamellar thickening and improved crystalline packing, lower defect density and higher crystallinity of the paraelectric a-phase. These factors result in an increase in T m and decrease T c . 67,92 When annealing the polymer below T c in the b-phase state, an increase of the T c is expected. 93 Solution-cast P(VDF-co-TrFE) (72/28) thin lms annealed at 120 C for 3 h (above T c , but below T m ) had a preferential chain orientation that was aligned parallel to the substrate surface with a high degree of crystallinity, and high dipole alignment upon electrical poling. These factors lead to a large polarization in 100 nm thin lms, P r ¼ 7.8 mC cm À2 and an E c ¼ 0.75 MV cm À1 . 67 Therefore, to achieve high piezoelectric coefficients, it is essential to control the annealing procedures through the T c . 92 As an example of the inuence of such processing conditions, for a solution-cast P(VDF-co-TrFE) (81/19) copolymer that was annealed at 120 C (close to the T c ) for 16 h the T c increased to 128 C but no change of T m was observed. In contrast, the T c increased to 125 C and T m increased to 151 C aer annealing at a higher temperature of 140 C for 16 h. 81 The changes of T c were associated with changes in the Gibbs free energies in the orthorhombic b-phase and hexagonal a-phase during annealing. Annealing of the material below its T c favors the removal of gauche defects in the ferroelectric phase which increases the degree of crystallinity and b-phase concentration. Annealing above T c facilitates the transformation of the paraelectric aphase into the b-phase. The b-phase concentration in the as-cast lm increased from 75% to 93% aer annealing at 130 C for 2 h (Fig. 8). The increase in b-phase concentration led to higher piezoelectric activity and increased the d 33 piezoelectric coefficient from an original value of 13 pC N À1 to 18 pC N À1 aer 4 h of annealing with a three-fold increase in elastic modulus and 10-fold reduction of oxygen permeability. 81 For energy harvesting applications high piezoelectric coefficients, such as d 33 , are benecial, see eqn (2)-(4).
In the case of P(VDF-co-TrFE) copolymers with a VDF content in the range of 65-82 mol%, the T c shows a large thermal hysteresis, it is lower during the cooling process (T Y c ) compared to the heating process (T [ c ). For the P(VDF-co-TrFE) 75/25 mol%, the T Y c ¼ 57-79 C was much lower than that of T [ c ¼ 124 C. 35 Due to the large difference in the free energy between the hexagonal a-phase phase and orthorhombic b-phase phase at T Y c , the phase transition proceeds rapidly even when the lm is cooled slowly. Therefore, TGTG 0 sequences in the hexagonal aphase phase are quenched partly in the orthogonal direction as conformational defects. When annealed above T c , P(VDF-co-TrFE) copolymers undergo a thickening of crystallites, resulting in a strong increase in crystallinity, up to 90% or more for a VDF content in the range 70-80%. 35,93,94,[94][95][96] The T c decreases as the TrFE content increases, and vanishes for P(VDF-TrFE) 50/50, due to cooperative movements in the ferroelectric b-phase near the transition temperature. 87 A maximum b-phase fraction of 66.3% was obtained from the sequential treatment of "annealing and cooling-pressing-electrical poling" while control materials only exhibited a b-phase fraction of 45.3%. 97 For a copolymer crystallising from the melt, the T c was affected by the cooling rate, suggesting that a higher cooling rate leads to VDF-rich crystalline regions due to the TrFE units having difficulty entering into the crystal lattice. In contrast, quenching the material from the melt leads to a broader endotherm, similar to dimethylformamide (DMF) crystallized samples annealed at 120 C. 93 Melt processing, in particular extension ows, also affect the crystallization and chain orientation of polymers. For example, in P(VDF-co-TrFE) copolymers, the extensional ow at a normal stress level of 6.35 Â 10 5 Pa produced an all-trans crystal structure in the 66/34 mol% copolymer. A 7 C higher transition temperature and lower melting temperature was observed in the two-dimensional ow rate range of 0.002-0.03 cm 2 s À1 for the 75/25 mol% copolymer. 98 In both PVDF and P(VDF-co-TrFE) copolymers, the a-phase tends to form when cooling from high temperature and at high cooling rates. The presence of increasing amounts of TrFE increases the formation of b-phase. A lower fabrication temperature favours the formation of polar band g-phases. 99 When an as-cast P(VDF-co-TrFE) (80/20) lm was heated to 200 C for 4 h, then quenched at 100 C, a 43% degree of crystallinity was obtained and the a-phase dominated; when quenched at À20 C, a 56% degree of crystallinity was obtained, and almost 100% b-phase was formed. Therefore, the annealing process affects the local distribution of the TrFE comonomer units in the crystalline lattices, which is reected by the varied T c and crystallinity. It is largely accepted that b-phase melting occurs in the range 165-172 C; a-phase crystals in the range 172-175 C with the g-phase melting between 175 and 180 C. 100 The ferroelectric behaviour of PVDF based materials have been investigated as a function of temperature and poling frequency, as shown in Fig. 9 and 10. 101 For a uniaxially  stretched P(VDF-co-TrFE) (50/50) lm, paraelectric behavior was obtained due to the nucleation of electric eld-induced ferroelectric nanodomains inside the paraelectric matrix when it was poled at a high poling frequency (1 kHz) at 100 C and above the T c of 64 C. These ferroelectric nanodomains were highly reversible and could be rapidly depolarized upon removal of the poling eld, as shown in Fig. 10C. At an intermediate poling frequency (10 Hz) at 100 C, an 'antiferroelectric-like' behaviour was observed, which could be attributed to the competition between depolarization and polarization elds upon reversed poling, this corresponds to the behaviour in Fig. 10B. Finally, at a low poling frequency (1 Hz) at 100 C, conventional ferroelectric behaviour with a rectangular hysteresis loop was observed since the small, reversible ferroelectric domains have sufficient time to grow into large irreversible ones. For energy harvesting applications the ability to pole the material and achieve a high remnant polarisation (P r ) which is modulated with stress or temperature is of interest for piezoelectric or pyroelectric harvesting, as in Fig. 10A. For energy storage applications, the antiferroelectric-like ( Fig. 10B) or paraelectric ( Fig. 10C) behaviour is more desirable due to the ability to store and recover energy, as given by the larger areas in the D-E f loops which are indicated by the purple areas in Fig. 10.
A b-phase content of 74% and an overall crystallinity of 42.6% was achieved for melt-quenched stretched PVDF lms. 102 A higher b-phase content of 82% and a d 33 of 21 pC N À1 was obtained by stretching PVDF lms that were solution-cast from N,N-dimethylacetamide (DMAc) solutions. 103 It is reported that an overall crystallinity of 52-60% with a b-phase fraction of 53% was present in PVDF lms prepared by spin coating from an acetone/DMF solution and then stretching uniaxially. 104 The pyroelectricity of a ferroelectric VDF-oligomer [CF 3 (CH 2 CF 2 ) 17 I]evaporated lm was investigated utilizing a low-frequency sinusoidal heat source. 72 The pyroelectric coefficient (p) of the 600 nm-thick VDF oligomer lm was À68 mC m À2 K À1 at 37 C, which is larger than those reported for ferroelectric polymers. The VDF oligomer is therefore a promising material for pyroelectric thin-lm infrared detectors and pyroelectric harvesting.
2.2.4 Nanoconnement. Nanoconnement has recently been used to facilitate the formation of polar b-phase. PVDF nanowires can be predominantly crystallized into the b-phase by connement in a nanoporous structure (e.g. in 200 nm channels), leading to the formation of well aligned polymer chains and crystallites arranged perpendicularly to the channel walls. 105 PVDF nanowires exhibited a higher remnant polarization (P r ¼ 19 mC cm À2 ) than the saturation polarisation (P s ¼ 9 mC cm À2 ), which was attributed to the combination of the polarization of the ferroelectric polymer with the charge derived from the superposition of leakage current to the displacement current. In contrast, thin lms of PVDF without any poling stage did not show any ferroelectric behaviour. P(VDF-co-TrFE) (70/30) nanowires formed by templating exhibited a P r ¼ 7.4 mC cm À2 and P s ¼ 9.6 mC cm À2 , that are similar to as-prepared P(VDF-co-TrFE) thin lms. Maximum d 33 values of À8.2 and À6.5 pm V À1 were obtained for the P(VDF-co-TrFE) and PVDF nanowires. Non-poled P(VDF-co-TrFE) thin lms had a d 33 of À15 pC N À1 , while in comparison, poled lms and bulk materials of both P(VDF-co-TrFE) and PVDF had a d 33 in the range of À20 to À30 pC N À1 (or pm V À1 ).
A signicant change in the crystallisation behaviour of P(VDF-co-TrFE) was observed when conned in a nanoporous template with pores less than 40 nm in diameter, 38 as compared to the bulk material. The copolymer crystallised into an oriented ferroelectric crystal lamellae perpendicular to the pore axis. Both melting and crystallization temperatures decreased with a decrease in pore diameter, and the T c was weakly affected. The presence of ferroelectric phase was preserved down to pore sizes as small as 15 nm. The results imply that nanoconnement enhances the formation and orientation of the ferroelectric b-phase and can enhance ferroelectricity and piezoelectricity in nanoscale P(VDF-co-TrFE) materials.
For a P(VDF-co-TrFE) copolymer with a composition ratio of 72.2/27.8 mol%, 88 the polymer chain length affects the conformations. It was found that a short polymer chain length leads to a higher crystal size and promotes a higher b-phase content; the d 33 piezoelectric coefficient was enhanced by decreasing the molecular weight (M w ) of the copolymer. A maximum d 33 value of À50 pC N À1 was achieved for the composition 72.2/27.8 mol% with a molecular weight of 470 kg mol À1 . Interestingly, the pyroelectric properties were enhanced for the lowest polymer crystalline grain size studied. A pyroelectric coefficient of 37.8 mC m À2 K À1 was obtained with the composition 71/29 mol% with a molecular weight of 505 kg mol À1 . These differences may be related to how the polarisation changes with either stress (thereby inuencing the piezoelectric coefficient) or temperature (thereby inuencing the pyroelectric coefficient).
When a PVDF based material is crystallised and conned in two-dimensional (2D) cylindrical anodic aluminium oxide (AAO), the crystal c-axis becomes perpendicular to the porous cylinder axes, which favours dipole switching. The polarity of the nanopore surface and the surface hydroxyl groups on the template nanopores promote the formation of the polar phase, with a b-phase content of 40.5%. As shown in Fig. 11a, the use of an oxygen plasma treatment increases the number of negatively charged OH À and O 2À groups on the surface of the Al 2 O 3 nanopore template, thereby increasing its hydrophilicity compared to the pristine template; see Fig. 11a and b.
Therefore, the presence of highly polar surface increases the template interaction with the PVDF dipoles, and forces the polymer chains to arrange in a TT conrmation, leading to a polar b-phase fraction of 40.2%. The use of a 3-aminopropyltrimethoxysilane (APMS) treatment increased the hydrophobicity of the nanopore surface, and enhanced the interfacial interactions with the polymer chains, and enhanced the polar phase formation to 71%, as shown in Fig. 11c. 106 However, no polar phases have been obtained from 2D connement down to 35 nm pores aer melt recrystallization.
When recrystallized from the melt in conned 3D polystyrene nanospheres 180 nm in diameter, PVDF exhibited both polar band g-phases, rather than the kinetically favoured aphase, 107 see Fig. 12. These results indicate that PVDF can nucleate homogeneously at a high crystallization rate in the 3D nanoconned space, which does not rely on the polarity of the  substrate. The use of 3D nanoconnement is thought to be more effective than 2D nanoconnement in producing polar crystalline phases in PVDF.
By varying the quenching temperature, the crystal structure and properties of PVDF can be tailored for different applications, as shown in Fig. 13. 108 For solution-cast lms, when quenched below 0 C, more b-phase was formed which was aligned to create a self-polarised material with a high d 33 of $49.6 pm V À1 (see Fig. 13c). This is much higher than typical values of À20 to À35 pm V À1 produced by mechanical stretching and poling. These values are comparable to PVDF nano-bres which show a d 33 coefficient of À54 pm V À1 , which originates from the high g-phase content with $75% crystallinity. 109 The high d 33 of the lm was attributed to the high ($100%) b-phase content and high crystallinity of 56%, as shown in Fig. 13a and b. Fast quenching at a low temperature induced a strong thermal eld gradient, which can cause crystals to align along the thermal eld direction. The interaction with polar water can also be a possible reason for the alignment of the b-crystals. 109,110 The self-polarised b-phase and high piezoelectric coefficient of the PVDF lms make them suitable for energy harvesting applications, such a vibration harvesting (Fig. 13d) and examples of devices are described at the end of the review (Section 3). Further details of ferroelectric polymer nanostructures under connement has been described 111 and recently nanostructured arrays of PVDF based polymers have been formed by nanoimprint lithography. 112 2.2.5 Surface charge and nucleation. The nucleation behaviour of a polymer depends on the properties of the nucleating agents, such as surface charge, surface area, lattice matching, concentration and dispersion, as well as the interfacial interactions between the nucleating agents with PVDF chains. 64 Nucleating agents reduce the nucleation energy barrier, increase crystallization kinetics, crystallisation temperature, melting temperature, and the degree of crystallinity. To evaluate the effects of surface charge on the nucleation efficiency, nucleating agents with a positive charge (phosphonium, pyridinium, pyrrolidinium, ammonium, and sulfonium salts), negative charge (sulfate and phosphate salts) and neutral agents (avanthone) have been melt-compounded with PVDF. 64 The addition of 0.5 wt% of nucleating agent altered the crystallization behaviour of PVDF, especially when the melting temperature of the nucleating agent was close to the melting temperature of PVDF. Only the a-phase was formed in the absence of nucleation agents with negative or neutral nucleation agents, while band g-phases were formed in PVDF with positively charged nucleation agents. Fig. 14 shows examples of the chemical structures of positive, negative and neutral nucleating agents.
Nanoparticles can also act as nucleating agent for polymers by increasing the crystallisation kinetics and degree of crystallinity. For PVDF composites lled with Na(M)Y zeolites, the negatively charged zeolite framework induces the formation of the polar g-phase in PVDF when prepared by solution-cast and melt compression. 114 The use of Na(Y) in the composite induced 100% of g-phase in PVDF due to the negatively charged iondipole interactions between the zeolite framework and the PVDF. The exchangeable (M) ion size in the zeolite also affects the PVDF phase structure. The presence of small sized lithium (Li) ions increased the crystallisation of polar phases, as compared to Na(Y), while the presence of potassium (K) and cesium (Cs) has the opposite effect. This behaviour depends on the interaction competition between the alkali ion-zeolite structure and alkali ion-molecular dipoles. Using larger sized ions and stronger interactions with zeolite structures reduced the interactions with polar polymer chains, leading to a lower fraction of polar phase. The relative permittivity ranged from 47 to 3 for PVDF/Na(Li)Y and PVDF/Na(Cs)Y, respectively, while the electrical conductivity of PVDF/Na(Li)Y was three orders of magnitude higher than that of PVDF/Na(Cs)Y. An increased relative permittivity can oen be related to increased conductivity in the material.
With the addition of graphene oxide (GO) nanoplatelets to PVDF, 115 only the g-phase was induced to form in PVDF when crystallized from the solution, and only a-phase forms from melt crystallization. When GO was combined with a positively charged surfactant (cetyltrimethylammonium bromide, CTAB), more g-phase crystals were formed in PVDF during isothermal melt crystallization at 160 C for 80 min. It is suggested that there are two distinct stages during the melt crystallization of PVDF/GO composites in the presence of CTAB, i.e., a simultaneous growth of gand a-phases, which is followed by an ato gphase transition. The T m was increased by 10 C for the g-phase PVDF/GO composites. The addition of CTAB was thought to strengthen the ion-dipole interactions between the GO and PVDF, as compared to weak p-dipole interaction between GO and PVDF.
Similarly, the effects of the surface charge on the crystal structure of PVDF were studied using nanoparticles of CoFe 2 O 4 whose surface was treated with three types of surfactants, i.e., anionic (SDS), nonanionic (Triton X-100), and cationic (CTAB). 116 A higher fraction of polar b-phase was obtained when CoFe 2 O 4 nanoparticles with a negative electrostatic charge were added. The b-phase content reached 30% and 90% for CoFe 2 O 4 / PVDF and CoFe 2 O 4 -SDS/PVDF, respectively, with a resultant d 33 of 23 and 33 pC N À1 respectively. This behavior was attributed to the interaction between the negatively charged magnetic CoFe 2 O 4 particles and the positive polymer CH 2 groups.
The ion-dipole interactions are described as (i) negatively charged surfaces with the CH 2 dipoles in PVDF, or (ii) the positively charged surfaces with the CF 2 dipoles in PVDF. By comparing the effects of positively-and negatively-charged organic nucleation agents, 113 as in Fig. 15, the positively charged nucleation agents interact more strongly with the partially negative CF 2 dipoles of PVDF, resulting in a lower nucleating energy barrier and increased crystallization kinetics and crystallization temperature for PVDF; this induces the formation of polar (b, g) phase formation. The addition of positively charged organic molecules in solid PVDF induced a 100% of g-phase formation, while the b-phase PVDF becomes dominant when in the molten state. Almost 100% of polar phases were formed by the addition of 3 wt% of positively charged small molecules (1butyl-3-methylimidazolium hexauorophosphate). The band g-phases can give rise to ferro-, pyro-, and piezoelectricity for harvesting, and the g-phase can lead to increased transparency 117 and exibility 113 of PVDF.
In addition to tailoring the processing conditions and heattreatment conditions, recent approaches to tailoring the piezoand pyro-electric properties are to use polymer blending and form nanocomposites, which will be now discussed.  121,122 For example, the miscibility of PVDF with amorphous PMMA is driven by hydrogen bonding between the carbonyl group of PMMA and CH 2 of PVDF and dipole-dipole interactions between CH 2 of PMMA and CF 2 of PVDF. 119 The blending of amorphous PMMA with PVDF decreased the crystallinity and crystal size of PVDF which were mostly in the polar b-phase when the PMMA content was above 25 wt%. 118 With the assistance of rigid PMMA segments, the b-phase was easier to pole and was reversible with electric eld changes. A higher PMMA content ($40 wt%) and quenching process favoured the ato bphase transition. Therefore, PVDF/PMMA blends with a high bphase content exhibited a high relative permittivity, and a low dielectric and energy loss due to the presence of PMMA. 118 Semicrystalline PBS is miscible with PVDF, and can promote the formation of b-phase in PVDF, depending on the PBS content and quenching temperature. 120 The b-phase fraction increased with increasing PBS content, up to 50 wt%, while the addition of greater amounts of PBS can restrict the nucleation and growth of the b-phase, thus resulting in more a-phase. Lowering the quenching temperature benets the formation of b-phase in PVDF.
Semicrystalline PA11 has comparable coercive elds and remnant polarization with PVDF, but a low piezoelectric response (<4 pC N À1 , see Table 1) at room temperature because of its higher glass transition temperature (ca. 50 C) compared to PVDF ($À30 C). 123,124 When blending ferroelectric PA11 and PVDF, the glass transition temperature and melting point of PA11 decreases with increasing PVDF concentration, due to the dipolar intermolecular interactions between the polar amide groups (-NH-CO-) in PA11 and the polar -CF 2 groups in PVDF. 121,125 It has also been found that the hydrogen-bonded structure of PA11 becomes more disordered as the PVDF concentration is increased. The PVDF developed a larger proportion of polar band g-phase in blends with a high PA11 concentration compared to pure PVDF under similar melt quench conditions. 121 During the uniaxial drawing process, the phase transformation of PVDF from the nonpolar a-phase to the polar b-phase is more complete, with more ordered b-crystals than in pure PVDF. These structural changes and the dipolar interactions between PA11 and PVDF lead to a 30% increase in piezoelectric properties and improved high-temperature stability of the blends at temperatures up to 160 C. 121 These properties enables this new polymeric blend material to be used in electroactive applications, such as sensing and harvesting in more hostile environments.
The formation of unique orientation textures due to conned crystal growth of PVDF in PVDF/PA6 blends of drawn lms have been reported. 122 Uniaxial-drawn PVDF/PA6 blends were heat-treated at 180 C for 3 minutes to melt the PVDF component, followed by non-isothermal crystallization of PVDF at a cooling rate of 0.5 C min À1 . For PVDF/PA6 as a 50/50 blend, stretched domains of PVDF with a diameter of 0.2-0.5 mm were dispersed in a PA6 phase. Spatial connement of the crystal growth resulted in the alignment of the crystal b-axis along the long axis of the domains, since the PVDF crystallized as thin cylindrical domains. The orientation behaviour is different from the oriented crystallization of PVDF/PA11 125 in which trans-crystallization from the interface causes an a-axis orientation in the drawing direction. It is thought that the domain size inuenced the mechanism of oriented crystallization.
A recent study showed that a P(VDF-co-HFP)/poly(ether imide) blend with a poly(ether imide) content larger than 80% created more free volume in the polymer matrix, thereby providing more space for dipole orientation polarization and providing an enhancement of the relative permittivity of the polymer blends. 126 2.2.7 Formation of nanocomposites. Polymer-based nanocomposites have attracted interest for high energy density capacitors with high energy density (U e ), 19 since these heterogeneous materials combine the high relative permittivity and dielectric properties of nanollers, such as ferroelectrics, with the high breakdown strength, low dielectric loss, and lightweight nature of polymeric matrices.
As discussed, PVDF and its copolymers contain mainly of aphase ($40%) 127 or a trace amount of b-phase when cooled from the melt; Fig. 4. A signicant body of research has demonstrated that nanoparticles with a negatively charged surface can interact with the positively charged -CH 2 groups, and induce the formation of the polar b-phase, with fractions up to 100%, 64 140 ), and conducting polymers. The effect of nanoparticle additions on the structure and piezoelectric properties of PVDF polymers generally depend on their nucleation efficiency, supercooling effect and interfacial interactions. The nucleation effect of the nanoparticle additions strongly depends on the particle size, 129,135 shape, 141 surface chemistry, 135 concentration, 130 dispersion, interfacial interaction and processing conditions.
2.2.7.1 Ceramic particles in nanocomposites. As described in Section 1, the generation of a piezoelectric potential under strain is due to the displacement of the positive and negative charges in the b-phase of PVDF. With the incorporation of ferroelectric inorganic particles, such as BaTiO 3 , the piezoelectric potential can be further enhanced. 129 BaTiO 3 hollow nanospheres 141 with particle sizes of z20 nm and surface area of 297 m 2 g À1 were surface-treated to enhance their compatibility with PVDF. Both the large specic surface area and increased surface functionality of the hollow spheres increased the b-phase content and degree of crystallinity at 16 wt% of ller. The changes in crystallinity led to a high relative permittivity of 109 and high energy density (U e z 21.7 J cm À3 ) with a high retained breakdown strength (E b ¼ 3.81 Â 10 3 kV cm À1 ) compared to pristine PVDF (3 0 z 11.6 and U e z 2.16 J cm À3 at 3.98 Â 10 3 kV cm À1 ). Surface functionalised Ba 0.6 Sr 0.4 TiO 3 nanobers 142 showed good dispersion and strong interfacial adhesion with PVDF matrix, which resulted in high relative permittivity (3 0 $ 12) and high breakdown strength (3900 kV cm À1 ) for the composite containing 2.5 vol% of nanobers. The maximum energy storage density reached 7.5 J cm À3 for a PVDF nanocomposite under the electric eld of 3900 kV cm À1 , which is three times that of pure PVDF (2.8 J cm À3 at 4000 kV cm À1 ). A recent study on BaTi (1Àx) Zr x O 3 (BTZO) nanocube lled PVDF composites for harvesting reported that 143 the exible PVDF/ BTZO composite lms exhibit a high electrical output up to $11.9 V and $1.35 mA compared to PVDF/BaTiO 3 lms which had an output of 7.99 V and 1.01 mA when subjected to a cyclic stress at 21 Hz with a constant load (11 N). The doping of various amount of Zr 4+ (x ¼ 0, 0.05, 0.1, 0.15, and 0.2) into the Ti 4+ site of BaTiO 3 can further enhance the piezoelectric performance of the particles. The piezoelectric potential of the PVDF/BTZO composites was adjusted by varying the composition ratio of BTZO and PVDF, electrical poling and mechanical loading. The composite generators were used to measure different water velocities at an outlet pipe and the peak power of the piezoelectric nanogenerator varied from 0.2 to 15.8 nW for water velocities ranging from 31.43 to 125.7 m s À1 .
2.2.7.2 Metal and metal salts in nanocomposites. For metal and metal salt lled PVDF composites, it was shown that 5 wt% of copper induced the formation of up to 90% of b-phase due to the high interfacial area between the well-dispersed nanoparticle surface and the polymer. 135 Superior ferro-and piezoelectret properties in a self-poled, porous hybrid ferroelectretic polymer nanocomposite lm was achieved by in situ generation of platinum nanoparticles embedded in a P(VDF-co-HFP) matrix. 134 The cooperative functionality between the self-polarized b-phase and the micropores as charge trapping sites was achieved by using a simple solvent evaporation method. The resulting porous hybrid lm exhibited a square-shaped hysteresis loop with large remnant polarization (P r z 61.7 mC cm À2 ), high piezoelectric charge coefficient (d 33 $ À686 pC N À1 ), and high dielectric properties (relative permittivity 3 0 ¼ 2678 and tan d ¼ 0.79 at 1 kHz); this high permittivity may be associated with the high loss (tan d). An open-circuit output voltage of 18 V with a 17.7 mA short-circuit current were generated under a 4 MPa stress. The high piezoelectric energy conversion efficiency (h piezo z 0.2%) of the nanogenerator demonstrated its potential for piezoelectric-based energy harvesters.
The addition of metal salts, such as cerium(III)/yttrium(III) nitrate hexahydrate (1-30 wt%), 144 or Fe 2 O 3 -Co 3 O 4 nanoparticles 138 to PVDF induced the formation of b-phase due to a strong ion-dipole interaction through the formation of hydrogen bonds between the water molecules of the salts and the -CF 2 dipoles of the polymer chains; this led to a large permittivity. A 15 wt% SiO 2 -loaded PVDF lm exhibited a high permittivity, due to the homogeneous dispersion and interfacial interaction of SiO 2 nanoparticles in the PVDF matrix. 131 For a P(VDF-co-HFP) matrix lled with mesoporous SiO 2 nanorods, the anisotropic shape of the SiO 2 nanorods and ordered mesopores doubled the amount of the -OH groups present. This increased the intermolecular interactions and enhanced the bphase content and ferroelectric properties of P(VDF-co-HFP). 145 Using this approach, the nucleation of the ferroelectric b-phase in PVDF can be inuenced by the geometry of the llers through the interface interactions between the local electric eld of the ller and PVDF dipoles, such as ion-dipole and dipole-dipole interactions.
Ferroelectric P(VDF-co-CTFE) terminated with phosphonic acid groups were synthesised and subsequently coupled with ZrO 2 llers. 136 The functional chain end-groups can form covalent coupling with the ZrO 2 surface, thereby providing the nanocomposites with stability and a uniform ller dispersion. The presence of 9.1 wt% ZrO 2 increased the crystallization temperature of the copolymer from 82 to 91 C, and increased the crystallinity from 18.3 to 22.8%. As a result of the intimate coupling between the ller and matrix, the interfacial interaction regions between the polymer and ZrO 2 increased the energy density at high electric elds. The energy density reached a maximum of 11.2 J cm À3 at 270 MV m À1 with 9.1 wt% of ZrO 2 , this was a 60% increase in comparison to the copolymer matrix. The improvement in the energy storage capability of the nanocomposite was ascribed to changes in polymer microstructure and an increase in crystallinity, the interfacial region and the rise of the electric displacement induced by the incorporation of the nanollers.
The incorporation of a cerium(III)-N,N-dimethylformamidebisulfate [Ce(DMF)(HSO 4 ) 3 ] complex 140 into PVDF provided a higher yield (99%) of the polar band g-phases. An enhancement of the output voltage ($32 V) of a nanogenerator based on a non-electrically poled cerium(III) complex containing PVDF composite lm was achieved by repeated human mechanical loading, whereas pure PVDF did not show any response. This high electrical output was due to the electrostatic interactions between the uoride of PVDF and the positive charge cloud of the cerium complex via H-bonding and/or a bipolar interaction between the opposite poles of the cerium complex and PVDF, as shown in Fig. 16a-c. The capability of the composite lm to charge a capacitor shows its capability as a piezoelectric-based energy harvester. The cerium(III) complex doped PVDF composite lm exhibited an intense photoluminescence in the UV region, which may be due to participation of an electron cloud from the negative pole of the bipolarized PVDF. This fact may also be of interest for the development of exible solid-state UV light emitters, as shown in Fig. 16d. 2.2.7.3 Nanoclay additions. Nanoclay, in particular, layered montmorillonite has been shown to provide excellent reinforcement and barrier functions to polymers. A variety of organoclays modied with cationic surfactants have been investigated for modifying PVDF-based polymers. Commercial organoclays, such as Cloisite®Na, Cloisite®6A, Cloisite®15A, Cloisite®20A, Cloisite®25A, Cloisite®30B, Lucentite STN, Nanocor Inc I.34TCN have been studied. The main difference between the organoclays above are the different surface treatment and polarity. It was found that the organoclay promotes the formation of b-phase of PVDF, while unmodied clay shows little or no inuence. Solution-casting 146 has been shown to facilitate dispersion and strengthening of the interaction between the organoclay and PVDF, as compared to meltprocessing.
The effects of nanoclay inclusions on the crystallisation behaviour of PVDF depend on the dispersion, exfoliation and interfacial interactions of nanoclay in the PVDF matrix. Three types of organoclay modied by different surfactants were studied: ammonium (C18), pyridinium, and phosphonium. It was found that the three types of organoclay could be readily dispersed in PVDF melts. The ammonium-clay showed the best dispersion and phosphonium clay was more efficient in forming the polar b-phase. 127 High b-phase fractions of $99% were obtained at 5 wt% of octadecyltriphenylphosphonium bromide modied clay with a 28.2 wt% organic content. In comparison, the addition of 5 wt% of unmodied clay leads to only 23% of bphase in the composites. The presence of nanoclay also acted as a nucleating agent and enhanced the melting and crystallization temperature of PVDF by 10 and 13 C, respectively. For the composite containing 3 wt% of nanoclay, the relative permittivity at 1 kHz was approximately 9.6 with a dielectric loss less than 0.05, the P r was 6.3 mC cm À2 , as compared to 5.0 mC cm À2 for pristine PVDF. A higher loading of nanoclay, over 5 wt%, led to a higher dielectric loss, increased leakage current and early breakdown due to the aggregation of the clay in the PVDF matrix. By tuning the surface charge and dispersion of the nanoclay in PVDF, the energy density of the PVDF/nanoclay composites was increased from 5.34 to 5.91 J cm À3 at 1 wt% clay content and the highest energy density could reach 10.2 J cm À3 . 147 In addition to promoting b-phase formation, organoclay platelets can also stabilize the b-phase from depolarisation. [148][149][150] The organoclay may play three roles in inuencing the crystal structure and properties of PVDF: 68,149 (i) a supercooling effect, (ii) a heterogeneous nucleating effect and (iii) intermolecular interactions between the polymer chains and silicate sheets. To investigate the role of organoclay, solutioncast PVDF/organoclay (Lucentite STN) composite lms with different thermal treatments were studied. This included annealing at 160 C/24 h, melt-quenching followed by annealing treatment 160 C/5 h and melt-slow cooling (10 C min À1 ). The as-cast samples exhibited nearly pure g-phase irrespective of the organoclay content, the melt-slow cooled samples showed a-phase at lower clay concentration, and a mixture of aand gphases at a higher clay content. The samples prepared by meltquenching followed by annealing had a mixture of a-, band gphases across a range of clay contents. Upon heating, the bphase initially transformed to the thermodynamically stable gphase before melting. The increased melt-crystallization temperature of PVDF with the addition of organoclay indicated that the clay did not show any super-cooling effect. Ultra-thin lms of PVDF/organoclay (Lucentite STN) nanocomposites have been prepared by heat-controlled spin coating. 71 The Lucentite STN favours the formation of b-phase and orientation in nanoscale thin lms, irrespective of preparation temperature, which leads to a high remanent polarization. In comparison, thick PVDF/organoclay nanocomposites lms only contained aphase aer quenching and slow-cooling from the melt, while the nanoscale thin nanocomposite lms showed a mixture of band g-crystalline phases without any a-crystalline phase. Therefore, the sample thickness also affects the crystal structures of the PVDF composites.
Organoclay platelets can retard the relaxation of the PVDF chains and stabilize conformation due to the presence of iondipole interaction between the exfoliated nanoclay layers and the PVDF. An increase of organoclay concentration should therefore increase the level of interaction. For organically modied montmorillonite (OMMT) modied PVDF, the a-phase dominates when the OMMT concentration is below 0.025 wt%, while the b-phase fraction exceeds the a-phase fraction when the OMMT concentration is between 0.025 and 0.5 wt%. 151 It is found that an increase of organoclay above 0.5 wt% introduces gauche defects in the all-trans conformation and results in a mixture of band g-crystals in the polymer. 151 In this case, the interactions between the clay and PVDF is responsible for crystal modication rather than the heterogeneous nucleating actions and super-cooling effect. 63,148,152 The presence of a second polymer which is compatible with PVDF can also benet the formation of the polar ferroelectric phase. 153 In addition to the well-studied PVDF/PMMA blends, as described in Section 2.2.6, the addition of a polymeric compatibilizer such as acrylic rubber (ACM) can facilitate the intercalation and dispersion of nanoclay in the PVDF matrix, and promote polar band g-phase formation. 153 A relative permittivity of 16 (at 10 Hz) was obtained for PVDF/ACM/clay (90/10/5 wt%), which was 40% higher than that of a PVDF/clay (100/5 wt%) nanocomposite without ACM.
Therefore, organically modied clays generally have good dispersion and strong interfacial interaction with PVDF. Exfoliated organoclay nanosheets promote and stabilise the epitaxial growth of the b-phase PVDF 149,154 mainly by the stronger interfacial interactions rather than heterogeneous nucleating effects or super-cooling effect. 149 The PVDF/clay nanocomposites have mainly b-phase, and also have g-phases coexisting at low crystalline temperature (T c < 155 C), the coexistence of gand b-phases was found at a high crystalline temperature range (T c > 155 C). The b-PVDF nucleation is inuenced by the geometrical factors, interactions at the interface between the nanoparticles and the PVDF dipoles and particle concentration. 31,32, 34 2.2.7.4 Carbon nanotubes additions. Single-walled (SWCNT) and multiwalled carbon nanotubes (MWCNT) have been used to modify PVDF and its copolymers. Utilizing a solution-mixing method using dimethylacetamide (DMAc) as the solvent, 155 pristine MWCNT (diameter 10-50 nm, length 4-10 mm) induced the formation of both aand b-phases in PVDF under sonication, while no b-phase was formed in mechanically mixed mixtures. According to density functional theory, more energy is required to form the TT b-phase, and more TGTG 0 a-phase can transform to the TT conformation under sonication, and the TT molecular chains can bind the CNT surface more tightly than the TGTG 0 polymer chains.
For SWCNT lled P(VDF-co-TrFE) composites prepared by solution-casting followed by annealing at 70 C for 38 h, 156 the pyroelectric coefficients of the composites were higher than that of pure P(VDF-co-TrFE) copolymer. The increase in the pyroelectric coefficients of the composites as a function of temperature is more gradual compared the copolymer, as shown in Fig. 17a. This may be due to the interfacial interactions between the SWCNTs and the polymer chains which hinder the mobility of the dipoles at higher temperatures, thereby limiting the overall contribution to pyroelectricity. The measured d 31 coef-cient was 25 pC N À1 , which is higher than the 20 pC N À1 for the pure PVDF-TrFE lm with no additions, see Fig. 17b.
By directly melt-compounding PVDF with pristine MWCNTs of diameter of 20-40 nm and length 5-15 mm through a twinscrew extruder, Yuan et al. 157 reported the relative permittivity of PVDF/MWCNTs composites to be as high as 3800, which is three orders of magnitude higher than pristine PVDF, while maintaining a low conductivity level at 10 À5 S m À1 . The enhancement was explained by a reinforced Maxwell-Wagner-Sillars effect due to the wrapping of PVDF chains on the MWCNTs surface and increased interfacial interactions, but the effect of the pristine MWCNTs on the piezoelectric active phase of PVDF was not reported and is of interest for further study.
Without additional post-treatment, only a trace amount of bphase PVDF could be formed in PVDF/CNT nanocomposites when processed by melt-cooling 158,159 and melt-spinning. 160 One of the benets of the inclusion of nanoparticles in PVDF is that the crystalline structures of PVDF can be tuned using surface modied particles, without additional annealing or electrical poling treatment. The nucleating efficiency of three types of surface functionalised MWCNTs in b-phase PVDF were compared. 159 The contents of the functional groups on the MWCNTs surface were carboxyl (-COOH $ 3 at%), amino (-NH 2 $ 0.5 at%) and hydroxyl (-OH $ 1.2 at%), respectively. It was demonstrated that the NH 2 -MWCNTs induced the highest percentage of b-phase (17.4%) in PVDF, 159 followed by OH-MWCNT (11.6%) and unmodied MWCNTs (9.4%). The nanocomposites containing COOH-MWCNTs had the lowest amount of b-phase (4.7%). It is believed that the combined effects of the dispersion of MWCNTs and the interfacial interactions account for the formation of b-phase in PVDF, as shown in Fig. 18. The NH 2 -MWCNTs and COOH-MWCNTs were meltblended with PVDF/PMMA blends, 161 and only the NH 2 -MWCNTs favoured the formation of b-phase, where both the pure blends and COOH-MWCNTs lled composites exhibited only b-phase aer melt-mixing. This was ascribed to the   Ionic liquids are a good dispersant for CNTs. It was found that the addition of 2 wt% of pristine MWCNTs only induced 6.1% b-phase in PVDF 162 aer cooling from the melt. However, when the MWCNTs were modied with an ionic liquid (1-butyl-3-methylimidazolium hexauorophosphate [BMIM] + [PF6] À ), almost 100% b-phase was achieved in PVDF. This was thought to be due to the interactions between the planar imidazolium cations of the ionic liquid that were wrapped on the MWCNTs surface with the -CF 2 bonds of the PVDF, thereby improving the dispersion of MWCNTs in the PVDF melt. In comparison, the incorporation of only a ionic liquid depressed the PVDF melt crystallization due to the miscibility of ionic liquid with PVDF and led to a higher content of polar band g-phase compared to using pristine MWCNTs. This indicates the synergetic effects of MWCNTs and ionic liquids on the b-phase formation in PVDF during melt cooling. MWCNT covalently modied with 3-aminoethyl imidazolium bromide 162 led to the formation of 100% bphase in PVDF at only 1 wt% concentration by both solventcasting and melt-blending processes. 163 It was concluded that covalently linked ionic liquids with MWNTs is an efficient method to induce polar phase formation in PVDF compared to physically modied ionic liquids and MWNTs.
In addition to ionic liquids, MWCNTs modied by ester groups (e.g. -COOC 2 H 5 ) 164 and PMMA 165 have been used to fabricate PVDF/CNT nanocomposites through solution-mixing, in which the fraction of b-phase was highly dependent on the modication of MWCNTs and their dispersion in the matrix. The ester (-COOC 2 H 5 )-functionalized MWCNTs (FMWCNTs) 164 were well-dispersed in the PVDF matrix and the FMWCNTs facilitated the transformation of ato the b-phase by the interaction of the pC]O group in the FMWCNTs and the pCF 2 group of PVDF. However, only a maximum of 50% b-polymorph PVDF was achieved, even at a high loading of the FMWCNT in the melt-cooled samples. PMMA-graed MWNTs via nitrene chemistry could induce the formation of an almost fully b-phase PVDF at a 5 wt% loading 165 from solution-cast PVDF nanocomposites.
For a PVDF/PMMA (50/50) blend, the introduction of a small amount of acidized CNTs (0.2 and 0.5 wt%) accelerated PVDF crystallisation, and two types of spherulites were observed in the PVDF/PMMA blends, namely ring-banded spherulites and compact spherulites. The acceleration of the crystallization of the PVDF phase was attributed to the joint effects of CNT nucleation and inducing phase uctuation between PVDF and PMMA during the crystallization process. 166 As mentioned in Section 2.2.6, the fraction of PMMA in the blends governs the origin of polymorphism in the PVDF, 118 while the effect of steady shear was found more pronounced in the blends rich in a-phase crystals. 167 Composites based on PVDF lled with NH 2 -MWNTs showed two distinct structural relaxations in dielectric loss owing to mobility connement of PVDF chains and smaller cooperative lengths. The aspect ratio of CNTs also affect the nature of the dispersion, nucleation and reinforcement of polymers. Two MWNTs with a similar diameter of 10-20 nm, and different lengths 5-15 mm (L-MWNT) and 1-2 mm (S-MWNT) were mixed with PVDF in DMF solutions. It was found that b-phase structures were formed for L-MWNT at 2 wt%, and a mixture of aand b-phases were detected at below 2 wt% of L-MWNT. For the short S-WMNT, a mixture of aand b-phases coexisted at a concentration at or below 2 wt%, which is believed to be due to the high-aspect-ratio CNT surface having more zigzag carbon atoms, which match with the all-trans conformation of the bphase. 168 Post-treatment can also induce a greater ato b-phase transition and dipole orientation in PVDF nanocomposites. By mechanical stretching, 169 a high phase transformation from ato b-phase ($96%) was achieved for nanocomposites with greater than 1.0 wt% carbon nanobers (CNF). The AC conductivity of CNF/PVDF composites decreased signicantly when the percolation threshold was raised from 1.0 to 4.2 wt% CNFs aer stretching. This was attributed to the reduced crystallinity induced by the phase transformation from ato b-phase as well as the CNF re-orientation.
The thermal conductivity of PVDF is approximately 0.22 W mK À1 , and was increased to 0.47 W mK À1 aer mixing the polymer with 10 wt% of CNT. With the addition of only 1 wt% GO, the thermal conductivity of PVDF/CNT/GO nanocomposites reached 0.95 W mK À1 . In comparison, the thermal conductivity of PVDF/GO with 1 wt% of GO was only $0.27 W mK À1 . Therefore, the presence of low intrinsic thermal conductive GO (0.14-2.87 W mK À1 (ref. 170)) may be of benet for high heat transfer rates in pyroelectric applications; for example eqn (6) indicates a high rate of temperature change are needed to produce high pyroelectric currents. Although the crystallinity of the matrix in the PVDF/CNT/GO composites is decreased in comparison with PVDF/CNT composites, a large number of polar g-phase crystallites were induced. The presence of GO facilitated the dispersion of CNTs and the formation of a denser CNT/GO network structure in the PVDF matrix is thought to be the reason for the enhanced thermal conductivity. 171 PVDF composites containing TiO 2 coated MWCNTs nanoparticles have been prepared through a solution-cast method, 172 followed with mechanical rolling. A highly oriented structure with both PVDF lamella and TiO 2 -MWCNT core-shell structures (TiO 2 @MWCNTs) were formed, and such an aligned structure led to enhanced breakdown strength and a piezoelectric coefficient d 33 of $41 pC N À1 when the particle loading was at 0.3 wt%; this is almost double that of the pure PVDF (see Table 1). MWCNTs were coated with a continuous layer of TiO 2 nanoparticles (TiO 2 @MWCNTs) by a simple hydrothermal process and TiO 2 @MWCNTs/PVDF composites were prepared by solution casting. Compared to the pristine MWCNTs/PVDF composites, the TiO 2 @MWCNTs/PVDF composites had enhanced permittivity and lower dielectric loss. In addition, the breakdown strength of the TiO 2 @MWCNTs/PVDF composites was also improved, which is favourable for enhanced ferroelectric properties and storage. 173 Therefore, the effects of CNT on the nucleation of b-phase of PVDF is mainly dependent on the surface charge-dipole interactions (surface chemistry), the particle size and concentration, as well as processing conditions (solvent, temperature, annealing, shearing, post-treatment).
2.2.7.5 Graphene additions. Graphene and its derivatives are effective nanoparticles for the modication of PVDF and its copolymers. The oxygen functional groups on the surface of GO or reduced graphene oxide (RGO) can interact with the -CF 2 in PVDF via electrostatic interaction and/or hydrogen bonding, thereby beneting the nucleation of ferroelectric gor bphase. 174,175 With a solvent casting route, approximately 100% of the polar ferroelectric b-phase was formed in PVDF/GO lms with a 0.1 wt% GO content, 175 and 80% of b-phase formed in PVDF/ RGO lms, 176 which led to enhanced mechanical properties, relative permittivity and electric polarization. However, PVDF/ graphene nanocomposites oen suffer from higher dielectric loss and lower breakdown strength compared to composites employing insulating piezoelectric ceramics, this is due to the presence of high electric elds in the composite due the presence of electrically conductive ller. 177 The volume fraction of the llers should therefore be as low as possible in order to maintain high breakdown strength and energy density.
Surface modication of graphene or GO provides a route for enhancing the dispersion and interfacial interactions with polymers, thereby leading to higher composite performance at lower ller loadings. To evaluate the effects of surface functionalisation of graphene on the formation of b-phase in PVDF, 178 a thermally-reduced exfoliated graphene (EG) with various oxygen-containing functional groups including uorination-EG, ozone-EG, and PMMA-g-EG were added to PVDF through a solution-mixing method with a 0.5 wt% ller content. The b-phase fraction of PVDF for the variety of systems was PMMA-g-EG > ozone-EG > uorination-EG > EG and was related to its specic interaction between the C]O group of PMMA and the CF 2 group of PVDF. In the frequency range of 10 2 to 10 7 Hz, the relative permittivity of the composites showed a linear relationship with the content of the carbonyl groups of the llers, in the order of PMMA-g-EG > ozone-EG > uorination-EG > EG. Among all the composites studied, the PVDF/PMMA-g-EG composite had the highest relative permittivity (3 0 ¼ 18) at 100 Hz.
When using an ionic liquid (IL) to modify graphene, for example, 1-hexadecyl-3-methylimidazolium bromide, 179 both graphene and the IL play a positive role in the crystallisation of the b-phase in PVDF. This was ascribed to the existence of graphene-cation interaction between the imidazolium cation and the aromatic carbon ring structure, and the electrostatic interaction between the -CF 2 group of the polymer backbone and imidazolium cation. The addition of 3 wt% of IL functionalized GO to PVDF promoted the formation of b-phase, and increased the melting temperature and glass transition temperature. 180 The maximum increase in the storage modulus (73%), Young's modulus (333%) and tensile strength (628%) was at 3% ller concentration. A sharp increase of DCconductivity of the composite to $10 À2 S cm À1 occurred at the percolation threshold of 0.1 wt%. The relative permittivity increased from 3 0 $ 7.5 for pure PVDF to 3 0 $ 13.5 for 3% ller with a percolation threshold at 0.1 wt%. PVDF/NH 2 -treated graphene nanodots (GNDs)/RGO nanocomposites were developed to utilise the synergetic effects of the RGO sheets and NH 2 -treated GNDs. 181 The RGO sheets promoted the formation of b-phase by disrupting the nucleation of a-phase in the PVDF. The -NH 2 groups attached on the surfaces of GNDs to effectively interact with the PVDF, thereby stabilising the polar bor g-phases. The resulting PVDF/NH 2 -treated GND/RGO nanocomposites exhibited a higher relative permittivity (3 0 z 61) and larger energy density (U e z 14.1 J cm À3 ) compared to pristine PVDF (3 0 z 11.6 and U e z 1.8 J cm À3 ).
The addition of GO, magnetic iron oxide (Fe 3 O 4 ) nanoparticles of size 10-12 nm, or the combination of both GO and Fe 3 O 4 led to the transformation of PVDF from the a-phase to the b-phase, which may be due to the heterogeneous nucleating effects of the nanoparticles and interfacial interactions. 182 With the addition of 5 wt% GO and 5 wt% Fe 3 O 4 , the maximum saturation polarisation was 0.065 mC cm À2 , this is higher than PVDF (mainly a-phase, 0.038 mC cm À2 ) and PVDF/Fe 3 O 4 (0.058 mC cm À2 ), but lower than that of PVDF/GO (0.1 mC cm À2 ), which may be due to the conducting nature of Fe 3 O 4 . A higher permittivity was observed for the composite lms (3 0 ¼ 12-19) than the pure PVDF lm (3 0 $ 6), with a dielectric loss of approximately 0.6. The enhanced magnetic, ferroelectric, dielectric, magneto-dielectric coupling and structural properties of the PVDF/Fe 3 O 4 -GO nanocomposite lms were attributed to the homogeneous dispersion and good alignment of Fe 3 O 4 nanoparticles and GO in the PVDF matrix, along with a conversion of nonpolar a-phase PVDF to the polar b-phase.
For RGO-ZnO lled PVDF lms, 183 a ato b-phase transformation was observed and a maximum content of 83% of bphase was determined by FTIR. A decrease in the size of the spherulitic crystal structure of PVDF/RGO-ZnO nanocomposites was observed and a combination of RGO, ZnO and Fe 3 O 4 led to approximately a 50% decrease of b-phase content in the PVDF lms. 176 However, with a Fe-doped RGO, a 99% of polar g-phase was formed in PVDF when the Fe-RGO content was 2 wt%, 184 which was thought to be due to electrostatic interactions among the -CH 2 and -CF 2 dipoles of PVDF and the delocalized p-electrons and oxygen functionalities of Fe-RGO via ion-dipole and/or hydrogen bonding interactions, as shown in Fig. 19a and b. The nanocomposite lms generated an open circuit output voltage and short circuit current up to 5.1 V and 0.254 mA respectively on loading with a human nger as a piezoelectric energy harvester. In addition, the nanocomposite showed a higher electrical energy density of z0.85 J cm À3 at an electric eld of 537 kV cm À1 , higher than that of pristine PVDF of 0.27 J cm À3 , as shown in Fig. 19c and d. This is of interest for combined harvesting and storage systems.
2.2.8 Electrospinning and nanocomposites. Electrospinning has become a promising technique for producing exible piezoelectric polymers as it can induce the formation of more polar crystal phases and align the molecular dipoles with a single processing stage, which is ascribed to the simultaneous mechanical stretching and electrical poling. 185 In addition, other characteristics such as ease of setup, scalability, generation of continuous nanobers and membranes with high specic surface area and high porosity are also benecial for producing energy harvesting devices. A variety of devices, such as high frequency transducers, implanted biosensors, vibration absorbers and composite pressure sensors have been reported. 186 A typical electrospinning apparatus consists of a syringe with a metal needle connected to a high voltage dc power supply. When the polymer solution is ejected through the metal tip at high voltage (7-30 kV), the pendant droplet at the syringe tip becomes electried, and is distorted into a conical shape, known as the 'Taylor cone'. This distortion is caused by the electrostatic repulsion between the surface charges and the coulombic force exerted by the external electric eld. 187 When the electrostatic repulsion between surface charges overcomes the surface tension of the solution, an electried jet of polymer solution is ejected from the syringe tip and deposits onto a grounded collector, while mechanical stretching and electrical poling occur simultaneously during the bre solidication process. Therefore, electrospinning can be a single process to induce preferential orientation and formation of ferroelectric dipoles in PVDF and its copolymers. 188,189 The electrospinning parameters, such as applied voltage and electrode-to-collector distance (from less than 1 cm, up to 20 cm), solvent polarity, polymer molecular weight and concentration, and additives (BaTiO 3 ,190 CNTs,191,192 Ag-CNT, 188 clay, 193 graphene, SiO 2 , and ionic liquid 194 ) have considerable effects on the bre morphology and properties. Subsequent treatments such as annealing and cooling, electrical poling, and compression can be applied to further enhance the crystallinity of the b-phase of the PVDF and its copolymers. 97,[191][192][193]195 2.2.8.1 b-phase content during electrospinning. Electrospinning favours the formation of b-phase formation or induces an ato b-phase transformation due to its unique in situ electrical poling and mechanical stretching effects. The b-phase content in polymeric bres is dependent on the solvent type, additives and electrospinning parameters. For example, when electrospinning PVDF from DMF solution, a b-phase fraction of 75% and a crystallinity in the range of 49-58% can be produced. It is evident that solvents with a higher dipole moment produce PVDF lms with g-phase crystallinity, while solvents with lower dipole moment result in the nonpolar a-phase. The formation of the ferroelectric b-phase can be controlled by the composition of the solvent, notably the water content. 196 Using a hydrated salt in the casting solvent formed lms with a high bphase content. Table 2 summarises some typical results for effects of electrospinning parameters on the formation of the polar b-phase.
By adjusting the electrospinning conditions to narrow the bre diameter or increasing the electrospinning voltage, more b-phase was obtained, which is ascribed to the increased elongation and a larger electric eld that acts as an increased poling eld on the polymer jet. 197 When electrospinning PVDF from a N-methyl-2-pyrrolidinone/acetone (5/5 v/v) mixed solution at 16 wt% concentration, PVDF nanobers with 95% of b-phase were collected from a rotating drum at 800 rpm. 198 When electrospinning 20 wt% PVDF from DMF/acetone (6/4 v/v) solution 199 in the presence of 3 wt% of tetrabutylammonium chloride (TBAC), almost pure b-phase bres were produced. The TBAC salt is believed to facilitate the degree of hydrogen bonding between water molecules and the uorine atoms of the PVDF and hence induce more trans-conformation. In contrast, only the aand g-phase was detected in the spin-coated samples from the same solutions. It is believed that the TBAC additive induced local conformational changes and electrospinning promotes inter-chain interactions.
A higher fraction of polar b-phase can generally lead to higher d 33 piezoelectric coefficients, 34 and therefore a higher voltage output in energy harvesting applications, however the piezoelectric performance of PVDF bers is also affected by the total dipole moment, i.e., the orientation of dipoles that is closely related to the poling effect. 2.2.8.2 Orientation of molecular dipoles during electrospinning. Given the high applied electrostatic elds and polymer jet characteristics of the electrospinning process, it is believed that electrospinning can not only facilitate more ferroelectric phase formation, but also induce dipole orientation in the polymer bres. A number of studies have reported exible, high-output piezoelectric nanogenerators based on PVDF bers using near-eld electrospinning (NFES) or conventional far-eld electrospinning (FFES) processes. 200,201,[201][202][203] However, the effects of the electrospinning process on the dipolar orientation of the polymer bres are still not fully understood and con-icting results have been reported. 202,203 During electrospinning, a polymer solution experiences a number of forces. The rst is a shear force when it ows through a capillary needle at a high rate. The second is a coulombic force when the jet is elongated and accelerated by the applied electric eld. When a rotation disk collector is used to collect the bers, a mechanical force may also be applied. These three forces can cause polymer chains to be aligned and/ or stretched in the spinning direction. Using a rotating drum or disk as a bre collector can exert an additional mechanical stretching, and promotes the formation and alignment of the caxis of the b-phase crystallites along the ber axis. 199 One study 204 reported that rotation of the collecting drum has a pivotal role in the enhancement of b-phase formation and degree of the orientation of the bers in the electrospun mats. However, another study found that the degree of orientation and the polymorphism behaviour of the bres did not vary signicantly with either the rotating disk speed or the size of the spinneret used. 199 This implies that the formation of the bphase is likely to be caused by the coulombic force imposed by the electric eld rather than the mechanical and shear force exerted by the rotation disk collector and spinnerets.
The coulombic force may cause conformational changes to the straighter TT conformation, and hence promote the formation of b-phase. To seek evidence of dipole orientation during the electrospinning process, Sun et al. 198 compared electrospun PVDF nanobers with bres made by a force-spinning process (mechanical stretching without electrostatic force), in which the nanobers were produced by an electrostatic force and a centrifugal force. It was found that both brous mats formed showed aligned PVDF bres with a high bphase content of 95%. However, the force-spun PVDF bers did not show piezoelectricity, although it had a high b-phase content, since the mechanical force may only induce a phase transformation, but fails to align the ferroelectric dipoles. In comparison, the FFES process can lead to an electric poling eld ($1 kV cm À1 ) on the polymer jet, thus inducing polarization of the crystallites with a preferential orientation along the bres; see Fig. 20.
A study on PVDF/polar polyacrylonitrile and PVDF/nonpolar polysulfone nanobers 205 concluded that mechanical stretching is more effective than electric poling to induce the formation of ferroelectric phases. The b-phase was more stable in the polar polyacrylonitrile blends than in the nonpolar polysulfone/PVDF blends, which indicates that apart from the mechanical stretching, the local electric eld-dipole interactions determine the nucleation and growth of polar PVDF phases. This reects that electrical poling is more effective than mechanical stretching in enhancing the piezoelectric properties of PVDF nanobers. 205 2.2.8.3 Synergistic effects of nanoparticles and electrospinning. While the benets of nanocomposites have been discussed, the addition of nanoparticles has also been used in the electrospinning process. The possible mechanism of the formation of b-phase PVDF with the aid of templates with a large surface area is proposed in Fig. 21.
Both OMMT 149 or modied CNTs, 162,206 with their intrinsically large aspect ratio, are capable of interacting with the positive charged -CH 2 or negative -CF 2 groups of PVDF, resulting in the large scale conformation of polar phases, in particular the b-phase. Clay can increase and stabilise the polar phase in PVDF bres, 207 and act as a processing agent to facilitate electrospinning. 208 Particles with a smaller aspect ratio, such as Ag, metal salts or ionic liquid, also induce the formation of polar b-phases, but are not as effective as the particles with a high aspect ratio. Electrospun PVDF/room-temperature ionic liquid (RTIL) composite nanobers 206 based on 1-butyl-3-methylimidazolium hexauorophosphate [BMIM][PF6], exhibited almost 100% of b-phase, which differs from the dominant gphases observed for melt-blended PVDF/RTIL blends. 209 Two different nanoscale particles: carboxyl-MWCNTs and Ag-doped MWCNTs were used to modify the piezoelectric coefficient of PVDF electrospun bers. 188 The high voltage electrostatic eld generated extensional forces on the polymer chains that aligns dipoles. The Ag-CNTs lled PVDF electrospun bers exhibited the highest piezoelectric coefficient (d 33 ¼ 54 pm V À1 ) in contrast to PVDF/CNT bers (35 pm V À1 ) and pure PVDF (30 pm V À1 ). In another study, silver nanowires (AgNWs) doped PVDF bers were electrospun in a mixed DMF/acetone solution. 210 The b-phase content in the PVDF was increased by the presence of AgNWs, and the piezoelectric coefficient d 33 was 29.8 pC N À1 for the nanobers webs containing 1.5 wt% AgNWs, which is close to that of P(VDF-co-TrFE) (77/23).
Additives can aid the electrospinning process, and also improve the piezoelectric performance of the composites bers. The presence of inorganic salt LiCl of 0.00133 wt% in PVDF DMF/acetone (4/6) solution in 16 wt% concentration, 211 decreased the ber diameter from 340 nm to 65 nm due to the increased charge carried by the jet, and increased the b-phase formation to 93.6%. The voltage response of PVDF was increased from 1.89 V to 8 V for composite bers with 0.00133 wt% of LiCl, which is of interest for harvesting. Alternatively, increasing the collector drum speed up to 8.4 ms À1 lead to the alignment of b-crystallites along the ber axis without a significant effect on the formation of b-phase and output voltage. The addition of acetic acid signicantly reduced the density and size of the beads, while the addition of TBAC effectively removed all traces of beads from the electrospun bers. The b-phase enhancement induced by TBAC is likely to be caused by the hygroscopic nature of the salt, which retains water in the bers and leads to hydrogen bonding between the water molecules and the uorine atoms of PVDF. 196 Non-woven ber mats of PVDF/BaTiO 3 nanocomposites 212 with randomly oriented ber diameters ranging between 200 nm and 400 nm have been fabricated. The PVDF electrospun mats exhibited a mixed crystalline phase consisting of both aand b-phases. However, the addition of BaTiO 3 was found to result in an increase in the b-phase content in electrospun PVDF/BaTiO 3 mats. By increasing the bre alignment in the PVDF brous membranes, the piezoelectric response was enhanced. 213 For randomly orientated membranes with a BaTiO 3 concentration of 20 wt%, the output voltage was 0.1 V and corresponded to a 7 mm displacement at 1 Hz. 185 By uniaxial-alignment of the composite bres, the piezoelectric output voltage was increased with increasing BaTiO 3 concentration, and reached 0.48 V at 6 mm deection, which is 1.7 times higher than PVDF bers of 0.27 V when subjected to the same deformation, see Fig. 22. 214 The increased piezoelectric response of uniaxial-aligned BaTiO 3 -PVDF nanobers in comparison with randomly oriented PVDF bers and thin lms suggests their possible uses in energy harvesting and as power sources in miniaturized electronic devices such as wearable smart textiles and implantable biosensors.
2.2.8.4 Electrospinning technology. A variety of electrospinning methods have been developed in order to produce aligned or doped nanobers for energy applications. For example, PVDF nanobers were produced by a bubble electrospinning technique, where the production rate of the obtained PVDF nanobers was higher (13.93 mL h À1 ) than that of the reported traditional electrospinning (1-5 mL h À1 ). 222 The crystallinity of the PVDF nanobers was higher than in traditionally synthesized PVDF powders. PVDF bers were also investigated via non-uniform eld electrospinning where the air pressure and PET substrate thickness were varied to align the PVDF bers, and increase the level of piezoelectricity. 223 PVDF/Fe 3 O 4 nanocomposite bers were prepared under magnetic eld assisted electrospinning. 224 Two Helmholtz coils were mounted on the electrospinning apparatus to create a uniform magnetic eld. The PVDF in DMF/acetone (3/1) solution with a concentration of 18 wt% were mixed with Fe 3 O 4 nanoparticles in the size of 20-30 nm at PVDF ratios of 1 : 5, 1 : 10 and 1 : 15. The application of an electromagnetic eld during ber deposition resulted in improved orientation of the polymer ow towards the grounded electrode and led to smoother bers with diameters in the range of hundreds of nanometers. A magnetic eld response of the nanobers with higher magnetic elds was observed.
During the conventional electrospinning process, the fast evaporation of the solvent and the Coulomb forces generated among the induced charges inside the electried jet causes the polymer jet to follow a curved shape. 225 These cause the elec-tried jets to spin in a looping path towards the collector, and thus the bending instability results in randomly oriented nanobers. 226 Well-aligned nanobers are oen required for energy, sensor and biomedical applications, 221,227 but their production by electrospinning remains technically challenging. Different techniques have been developed to produce unidirectional aligned bers, such as modication of electrostatic eld, 228-230 mechanically wiring, and using a specially designed collector, including a fast rotating mandrel collector 221,231 or a parallel-electrode collector. 232 Centrifugal electrospinning 221 applies a centrifugal force and disperses a PVDF solution through a capillary, which causes elongation and thinning of the solution jet, and the ber is produced with no applied voltage. In a recent study, a hybrid centrifugal electrospinning process has been developed by integration of the concepts of the parallel-electrode electrospinning with centrifugal dispersion, that can produce highly aligned bers at a large scale, 221 see Fig. 23. Aligned PVDF bers were prepared across three-inch electrode gaps at an applied voltage of 15 kV and a spinneret rotation speed of 200 rpm from a 20 wt% PVDF solution with 3 wt% TBAC. Randomly oriented bers as well as aligned bers produced by the parallel  electrodes were also prepared. All ber samples were 25 mm long with a cross-sectional area of 0.76 mm 2 . The samples were embedded in PDMS, connected via electrodes to a voltageoutput analyzer and clamped onto a nanomechanical tester to produce controlled strain rates, as shown in Fig. 23a. The aligned bers shown in Fig. 23b produced an output voltage of $3.04 mV at a strain of 0.10 compared to the only 0.059 mV for the randomly oriented bers. Although the individual PVDF bers in the randomly oriented specimen were likely to be poled, the randomly oriented poling directions in the membrane's bulk macrostructure resulted in a negligible response. A schematic of the centrifugal system is shown in Fig. 23f.
In a rotating collector conguration, the polymer bers are deposited and wrapped around a rotating mandrel. The degree of ber alignment mainly depends on the mandrel rotational speed. By simply mounting two polyimide lms onto a rotating cylindrical collector with a 1 mm gap from the collector surface, the air ow generated by the rotation of the collector can be redirected, and the electric eld at the gap between the lms can be altered; this is termed the 'end-point' control assembly method, 233 as shown in Fig. 24. The deposition of nanobers can therefore be controlled by the relative position of the gap and the collector. With this manufacturing setup, PVDF ber bundles were electrospun when operated at the tip-to-collector distance of 6 cm, external applied voltage of 8 kV, and the two polyimide lms separated by a distance of 5 mm. It was found that both the alignment and positioning controllability of the electrospun PVDF bers were improved as compared with conventional electrospinning with a rotating cylindrical collector.
The electrospinning process can be classied according to the spinning distance, for example, near-eld electrospinning (NFES, spinning distance < 1 cm), 'short-distance' electrospinning (spinning distance of 1-8 cm) 195,[234][235][236][237][238][239] and conventional electrospinning (spinning distance > 8 cm). The main characteristic of NFES is that the jet does not have a 'whipping' movement within the short spinning distance, and the macromolecular chains can align along the ber length. In the conventional electrospinning process, the polymer jet undergoes whipping and solidication before it deposits on the substrate. The electrospinning of PVDF at a short spinning distance range (1-8 cm) can form highly interconnected bers throughout the web. The mechanical properties are higher than bers made by conventional electrospinning due to the inter-ber connections stabilizing the brous structure. 195 Both NFES and short-distance electrospun PVDF nanobers could be more suitable for the development of robust energy harvesters.
In order to grow b-phase extended-chain crystallites in PVDF bers, a high electric eld (1.6 Â 10 7 V m À1 ) elongates the polymer jet and a strong extensional force is applied when the glass tube collector rotated at 700-2100 rpm. A further introduction of CNTs enhances the b-phase formation, and the benets of CNT additions have been discussed in Section 2.2.7.4. The PVDF bers exhibited a large deection under high electric eld. Bending resulted in a mechanical strain (0.05-0.1%) that was distributed along the PVDF bers, and the maximum voltage and current output of the piezoelectric energy harvesters were 43.6 mV pp and 240 nI pp at a 15 Hz impact frequency, respectively. A exible energy harvester was made by patterning of ordered PVDF/poly(g-methyl L-glutamate) (PMLG) composite bers produced by NFES onto a PET lm. The NFES process improved the piezoelectric properties of the PMLG/ PVDF composites, resulting in better orientation of their dipoles, a high ultimate stress (27.5 MPa), and a high Young's modulus (2.77 GPa). The energy harvester could capture ambient energy with a maximum peak voltage of 0.08 V, a power of 637 pW, and the energy conversion efficiency is 3.3%. The electro-mechanical energy conversion efficiency of this PVDF/ PMLG energy harvester was up to three times higher than those of pure PVDF and PMLG based energy harvesters. 237 The NFES and hollow cylindrical near-eld electrospinning (HCNFES) process was used to fabricate piezoelectric PVDF/MWCNT nanobers. 240 When comparing NFES and HCNFES, the HCNFES process applies a high electric eld with an in situ mechanical stretching for alignment of dipoles along the longitudinal direction of PVDF nanober. Therefore, the PVDF nanobers fabricated using HCNFES can be of smaller diameter and higher b-phase content with higher piezoelectric activity for harvesting.

Examples of structuring ferroelectric polymers in devices
This section describes some of the potential harvesting devices that exploit the properties of ferroelectric properties that have been optimised using the methods described above. The emphasis of this section is not on the design of the devices but to provide examples on how the optimisation of the ferroelectric properties of the polymers using the processes described in this review can lead to improved performance. Exploitation of the ferroelectric materials will be achieved by a combination of both harvester device design and optimisation of the structure of the piezoelectric material at a range of scales (molecular-, micro-and macro-structure). Since all pyroelectrics are piezoelectric, researchers have attempted to combine both pyroelectric and piezoelectric harvesting approaches. Table 3 compares the relevant equations for a pyroelectric subjected to a temperature change (DT) and piezoelectric subjected to a stress (Ds) with similarities in the relationships between current, voltage and stored energy. Due to their similarities there is interest in hybrid piezoelectric-pyroelectric harvesting systems 241-244 whereby a combination of temperature change and stress is applied. In such systems care must be taken to ensure the changes in polarisation are constructive and enhance the power generation of the harvesting device. 241 As discussed by Sebald et al., the frequencies of temperature and vibration can be different and there is a need to optimise the electronics for such a hybrid system. 245 Flexible and stretchable piezo-and pyro-electric generators have attracted considerable attention in energy harvesting systems, especially in wearable and portable electronics, electronic skins, touch screens and electronic textiles. The challenges of exible and stretchable nanogenerators include the limited output power generation, low efficiency at low frequency, mechanical weakness, low stability and durability. The exibility and stretchability of the devices rely on the intrinsic mechanical properties of polymers and elastomers, or they can be realised by intelligent design of the device structures. 246 For instance, through the fabrication of foam-like structures, micro-patterning, 243 arch-shape assembly, electrospun membranes and textiles. The use of structuring technologies expands the range of potential materials from elastomer to semi-crystalline polymers, and hybrid nanocomposites. Nanostructuring can also enhance output power density by integration of piezo-and pyro-electric generators with other harvesting mechanisms, such as the triboelectric effect, and energy storage devices.

Curved and multi-layered structured devices
A stretchable hybrid piezo-and pyro-electric energy harvester was fabricated by using micro-patterned PDMS/CNTs composite as an electrode, which also makes the nanogenerator stretchable and exible. Graphene was used as a top electrode for improved heat transfer, and a P(VDF-co-TrFE) copolymer layer provided piezo-and pyro-electric performance, as described in Section 2.2.1. Under the effects of a thermal gradient and mechanical strain from 0 to 30%, the pyroelectric output voltage was stable at 0.4 V. 243 The potential of the material to harvest both mechanical loads (s) and temperature changes (DT) was examined. The total change in polarisation is expressed as: A multilayered exible PVDF curved generator was studied by investigating the effects of using a curved structure and tailoring the substrate material. 246 By stretching the device to a 1 cm displacement in the x-direction, the PVDF lm, without a polyimide substrate, generated a low output voltage of only 0.1 V, while with a substrate it generated over 100 V. In addition, PVDF with a curved structure generated 86 V, in comparison to a at PVDF lm that only produced 15.7 V. As a result, the attachment to a substrate to a curved structure allowed the PVDF to be subjected to a higher external force, and induces a higher stress and produces a higher output power. The multilayered curved nanogenerator produced a peak output voltage of $200 V and a peak output current of $2.7 mA. The output power density of the generator reached $17 mW cm À2 , which could illuminate over 950 LED bulbs. This concept was reported to provide the groundwork to enhancing the output power of conventional piezoelectric generators, thereby enabling novel approaches to realizing self-powered systems.
A hybrid piezoelectric/triboelectric generator 247 was fabricated by vertically-stacking two functional layers, as shown in Fig. 25. An arch-shaped piezoelectric generator with a Au/PVDF/ Au structure on the top, and a touch-and-release type triboelectric generator with a PTFE/Al structure on the bottom. In the Table 3 Comparison of relevant equations for pyroelectric p ¼ dP s /dT (C m À2 K À1 ) and piezoelectric systems d ij ¼ dP s /ds (C N À1 or C m À2 /N m À2 ) and 3 T 33 is permittivity at constant stress 16

Parameters
Pyroelectric Piezoelectric

33
Â h Â Ds triboelectric generator, electrical charges are generated due to the charge transfer between two thin organic/inorganic lms that exhibit a different surface electron affinity. The arch-shaped piezoelectric generator produced electricity by the d 31 mode. The device had a common shared electrode as a bottom electrode of the piezoelectric generator and as a driving electrode of the triboelectric generator. With two full-wave bridge diodes connected to the hybrid generator, the piezoelectric and triboelectric outputs were effectively combined. The hybrid open-circuit output voltage and current density were approximately 370 V and 12 mA cm À2 , respectively, and the output power density was $4.44 mW cm À2 . The hybrid nanogenerator produced a high output power even at a mechanical force of as small as 0.2 N, enough to light 600 LEDs.

Nanoconnement based devices
P(VDF-co-TrFE) nanowires with a diameter of approximately 196 nm were prepared by using an AAO template-wetting method, 248 with the nanoconnement promoting high crystallinity and preferential molecular dipole orientation in the polymer nanowires; as discussed in Section 2.2.4. The as-prepared selfpoled polymer nanowire exhibited a peak electrical output voltage of 3 V and peak output current of 5.5 nA when subjected to a strain rate of 0.1% s À1 , which is comparable to those reported highest efficiency values estimated from electrospun nanowire generators. 234 The energy conversion efficiency was approximately 11%.

Foam-like structured devices
Meso-and micro-porous structures are important for tailoring the mechanical energy harvesting performance of polymer lms. 249 Porous PVDF nanogenerators have been fabricated by templating or electrospinning methods (Section 4) to further enhance the piezoelectric properties while providing the materials with stretchability and exibility. In addition to tailoring the crystalline structures of PVDF-based materials, the pore size and porosity of the materials can also affect the properties. By using a templating method, mesoporous PVDF thin lms were prepared by casting PVDF/ZnO nanoparticles solutions, following by etching of the ZnO nanoparticles, as shown in Fig. 26A. The benets of nanocomposites and electrospinning nanocomposites have been discussed in Section 2.2.7 and 2.2.8, respectively. The pore size and porosity of the lms were controlled by the size and volume fraction of the ZnO particles. 249 With a ZnO particle size of 35-45 nm, the output voltage increased from 3.5 to 11 V as the porosity increased from 6.5 to 32.6% when the ZnO mass fraction increased from 10 to 50%, but the voltage decreased to 8.3 V as the porosity further increased to 45.5% at a ZnO mass fraction of 70 wt%. The mixing of ZnO with PVDF helped to from a higher fraction of bphase in the PVDF, and reached a maximum fraction of b-phase at a ZnO content of 50 wt%, where the output voltage and current of the porous PVDF lm nanogenerator 28 mm thick was 11 V and 9.8 mA at a frequency of 40 Hz, respectively. Porous PVDF lms were fabricated with a ZnO nanowire array template-assisted preparation method, as shown in Fig. 26B. 250 With a sonic power of 100 dB at 100 Hz, the porous PVDF with a pore size of 2 mm showed an enhancement in the peak-to-peak piezoelectric potential (open circuit voltage)/ piezoelectric current (short circuit current) with an output of 2.6 V/0.6 mA, which is over 5.2 times (piezoelectric potential) and six times (piezoelectric current) higher compared to that of the bulk lm (0.5 V/0.1 mA). The device could produce a rectied power density of 0.17 mW cm À3 . For the porous PVDF lms with pores of 700-900 nm and 2.7 mm in thickness, an output voltage  2.84 V and an electric eld of 1.05 Â 10 6 V m À1 were measured, which is ascribed to the geometrical strain connement effect, as shown in Fig. 26C. 251 The pore structure can reduce the strain on the orthogonal direction, to allow the development of strain that is conned to the axial direction. The porous structured nanogenerators showed a higher performance compared to ber-like generators, such as a single PVDF nanober (30 mV/3 nA, 6.5 mm diameter/600 mm length), 234 PVDF nanobers (20 mV/0.3 nA, sample thickness and area were not indicated), 202 PVDF mats (z1 V, 100 mm thickness/1 cm 2 area, output current not indicated), 252 and PVDF nanober membranes (z2.2 V/z4.5 mA, 140 mm thickness/2 cm 2 area). 203 Porous P(VDF-co-TrFE) copolymer lms have also been prepared via an immersion precipitation and phase inversion method, 253 where the porosity can be varied by changing the type of solvent and nonsolvent. By using 15 wt% of P(VDF-co-TrFE) (60/40) copolymer in MEK with ethanol as a nonsolvent, a porous lm with an average pore size of 2.2 mm and 33% porosity was produced. As shown in Fig. 26D, the porous lm had a coercive eld, E c of 30.5 MV m À1 , similar to those of commercial copolymer thin lm, but the relative permittivity 3 0 was 39.9 which is higher than commercially available thin lm material (3 0 $ 28.8). The pyroelectric coefficient of commercial thin lm and porous lms were 4.5 Â 10 À5 and 4.3 Â 10 À5 C m À2 K À1 at room temperature, respectively, and both exhibited peak values of 1.4 Â 10 À4 and 1.2 Â 10 À4 C m À2 K À1 at 50 C, respectively, which is the transition temperature from the ferroelectric b-phase to the paraelectric a-phase phase. In this case, the porous structure improved the thermal response of the lms, but reduced the pyroelectric coefficient. Therefore, porous structure and porosity are a key features to tune the properties of the materials for harvesting applications.

Fibrous structured devices
As mentioned for electrospinning (Section 2.2.8), the NFES technique can produce in situ poled and well-aligned nanobers in a continuous manner due to the strong electric elds (>10 7 V m À1 ) 254 and stretching forces during the electrospinning process. It can realize direct-write and integrated PVDF nanogenerators with high energy conversion efficiency. 234,255 The piezoelectric coefficient d 33 is À63.25 pm V À1 for a single NFES bre, 238 or À57.6 pm V À1 for PVDF brous mats, 256 which is twice as large than that reported for PVDF thin-lms (about À15 pm V À1 ). The typical electrical outputs of the direct-write PVDF nanogenerators are 5-30 mV and 0.5-3 nA. 215 In comparision, randomly distributed PVDF nanobers prepared by conventional electrospinning process did not show a measurable electrical output. A number of serially integrated layers of the fabricated nanober were encapsulated in a device consisting of Fig. 26 Images of (A) mesoporous PVDF thin film templated with ZnO nanoparticles, and the voltage output of a PVDF thin film NG (fabricated from a 50% ZnO mass fraction mixture) generated during one cycle of surface oscillation. The blue and red curves were collected under forward and reverse connections, respectively; 249 (B) cross section SEM images. Scale bars are 5 mm. A porous PVDF nanostructure using a ZnO nanowire template, and the piezoelectric potential (blue) and piezoelectric currents (red) from the porous PVDF and bulk structure under the same force; adapted with permission from ref. 250 a PDMS polymer on a PVC substrate, and an increase of output voltage was obtained due to the increasing number of layers via a summing up the external piezoelectric potential. A nanogenerator that comprised 20 000 rows of well-aligned PVDF NFs was able to create a peak output voltage of $4 V and a current of 75 nA. A three-layer integrated nanogenerator could reach a maximum output voltage and current up to 20 V and 390 nA, respectively. By attaching the nanober-based device on the human nger under a folding-releasing cycle at $45 , the output voltage and current could reach 0.8 V and 30 nA, respectively, making the nanober based devices promising for exible and wearable electronics. 238 Fig . 27a shows a hollow cylindrical near-eld electrospinning setup. 215 Large arrays of PVDF bers were produced with a high concentration of b-phase. The bers were deposited on a PET substrate and xed with copper foil electrodes at both ends to form a exible PVDF energy harvester, as shown in Fig. 27b. At a strain of 0.05%, a maximum peak voltage and current was measured to be 76 mV and 39 mA, respectively, as shown in Fig. 27c. 215 The co-doping of Ce 3+ and graphene to PVDF electrospun nanobers helped to align and achieve in situ poling of the polar ferroelectric crystal phases. 220 The sensitivity and energy harvesting performance of the composite bres were tested for different applications. As an ultrasensitive pressure sensor, it could detect a pressure as low as 2 Pa. As an acoustic nanogenerator, it generated an output voltage of 11 V and 6.8 mW of maximum power by the application of a 6.6 kPa of pressure amplitude, which could light up 10 blue light emitting diodes (LEDs). For sound-driven power generation, when a sound intensity of $88 dB was applied to the nanogenerator, an AC output voltage of approximately 3 V was obtained but no current was reported; however the device could power three blue LEDs. These properties are ascribed to the synergistic effects of graphene, Ce 3+ doping and electrospinning. The electrospun nanogenerator was reported to be promising for application as a pressure sensor, mechanical energy harvester, and effective power source for portable electronic and wearable devices.
Single PVDF nanober prepared by NFES only show an electrical output in tens of millivolts under bending. 234 With a metallic coaxial needle injector, hollow PVDF ber tubes were produced by NFES. Near-eld electrospinning was shown to enhance the energy harvesting performance of hollow PVDF piezoelectric bers. 235 The elongation of the bre tubes was greater than that of solid PVDF bers, with a tensile strength of 32.5 MPa. The output voltage of the ber tubes was 71.7 mV with an external load resistance of 6 MU; the output power was also signicantly greater (856 pW) than the solid PVDF ber which has output voltage of 45.7 mV and the maximum output power of 347 pW. As a result, the power generation of the PVDF ber tubes, while small in magnitude, was 2.46 times higher than that of the solid bers. With a rotating glass tube collector, highly aligned PVDF nanobers were produced by the NFES process. 236 All-bre piezoelectric fabrics nanogenerators 257 were fabricated by assembly of knitted high b-phase ($80%) piezoelectric PVDF monolaments as the spacer yarn interconnected between silver coated polyamide multilament yarn layers acting as the top and bottom electrodes. With an effective area of 15 cm Â 5.3 cm, the 3D spacer piezoelectric fabrics exhibited an output power density in the range of 1.10-5.10 mW cm À2 at applied impact pressures of 0.02-0.10 MPa. This was reported to be higher than 2D woven and nonwoven piezoelectric structures. The piezoelectric fabrics can also be coupled with the knitted and screen printed supercapacitors for textile based energy harvesting storage in wearable electronics. 258 Clearly, these harvesting devices demonstrate that the optimisation of PVDF at multiple scales can improve the performance of energy harvesting systems.

Copolymer and nanocomposite based devices
Copolymers and nanocomposite materials have been used to create energy harvesting devices. Bhavanasi et al. reported enhanced piezoelectric energy harvesting performance in bilayer lms of poled PVDF-TrFE and graphene oxide (GO). Bilayer lms of poled P(VDF-co-TrFE) and GO 259 exhibited a voltage output of 4 V and power output of 4.41 mW cm À2 for energy harvesting as compared to poled P(VDF-co-TrFE) lms alone that had a voltage output of 1.9 V and power output of 1.77 mW cm À2 . The enhanced voltage and power output in the presence of the GO lm was thought to be due to the combined effect of electrostatic contribution from the GO, residual tensile stress, enhanced Young's modulus of the bilayer lms, and the presence of space charge at the interface of the P(VDF-co-TrFE) and GO lms. The higher Young's modulus and relative permittivity of GO was thought to facilitate efficient transfer of mechanical and electrical energy. Zhao et al. 260 produced piezoelectric nanocomposites composed of barium titanate nanoparticles with a PVDF matrix. By a simple and solvent evaporation process, the BaTiO 3 nanoparticles and PVDF composites where formed with an oriented structure and the highest open-circuit voltage reached 150 V. Siddiqui et al. 261 produced lead free and exible piezoelectric nanogenerators based on a piezoelectric nanocomposite thin lm of barium titanate nanoparticles embedded in a highly crystalline polyvinylidene P(VDF-TrFE) polymer for charge storage, Fig. 28. The nanocomposite with up to 40 wt% of BaTiO 3 in the nanocomposite produced an output voltage of 9.8 V and output power density of 13.5 mW cm À2 under cyclic bending, comparable to the output of conventional inorganic polycrystalline lead zirconate titanate devices. The high performance of nanocomposite system was thought to be attributed to the high effective piezoelectricity of the crystalline P(VDF-TrFE) and BaTiO 3 .

Summary and perspective
This review has provided a detailed examination of the use of exible PVDF based ferroelectrics for energy harvesting applications. The potential mechanisms by which a ferroelectric can harvest vibrations or thermal uctuations has been explored. In terms of improving the ferroelectric properties for harvesting applications, a wide variety of properties of interest can be tailored. This can include the piezoelectric coefficients (such as d 33 and d 31 ), relative permittivity, pyroelectric coefficient, Curie temperature (T c ), the electric eld required to activate and align dipoles in ferroelectric domains (coercive eld, E c ) and the remnant polarization (P r ) which is related to the number of dipoles "locked" in aer the poling process. The main ferroelectric phase of PVDF is the b-phase and approaches to maximise the fraction of the ferroelectric phase to optimise the ferroelectric, piezoelectric and pyroelectric properties of polymer materials are described in detail; these include the inuence of polymer processing conditions, co-polymerisation, heat treatment, nanoconnement, blending, porosity, ionic liquids, forming nanocomposites that contain dielectric or conductive llers and electrospinning. A wide variety of processing conditions and methods are available to optimise the properties above for the desired application; a combination of these methods can also be exploited. While a number of these approaches have been used to design and tailor exible material for harvesting applications, such as forming nanocomposites and electrospinning, some areas have yet to be explored, such as the use of nanoclays and ionic liquids. The majority of work concentrates on processing and experimental characterisation and opportunities exist for modelling these complex materials are a range of scales to inform the experimental efforts. The range of interactions and potential benets are highlighted in Fig. 29. We have seen that concept of harvesting typically covers systems at relatively lower power levels (nW to mW) for applications such as low power electronics or for wireless sensors and reduce reliance on batteries, electric cables and provide autonomous sensors or devices. Less effort has considered the use of Olsen types cycle for higher power/larger thermal harvesting applications and are worthy of investigation.
Applications include exible and stretchable piezo-and pyroelectric generators have attracted considerable attention in energy harvesting systems, especially in wearable and portable electronics, electronic skin, medical devices touch screen and electronic textile with stretchable properties. Since ferroelectric polymers are both piezoelectric and pyroelectric properties, there is the prospect to harvest energy from multiple sources including vibration and thermal uctuations so that the design of hybrid systems is possible. This can also be combined with additional 'non-ferroelectric' based harvesting mechanisms in polymers, such as tribo-electric effects. Ferroelectric polymers, such as PVDF and its copolymers can also be used as energy storage materials due to their high dielectric strength, permittivity and relaxor properties; therefore there is also scope for combining harvesting and storage applications, such as multilayer structures and devices. While ceramic based materials with high Curie temperature and high ferroelectric activity are available the ability of polymers to be processed at lowtemperatures, their low density, bio-compatibility, low stiffness, exibility and mechanical robustness, such as toughness and high strains to failure will ensure there remains a number of applications where these materials are of continued interest for harvesting applications.