ELEMENTS THEREOF

This work was financially supported by a grant from the European Research Council through an ERC Starting Grant (Grant no. ERC-2014-STG-639526, NANOGEN).

BACKGROUND TO THE INVENTION

Field of the invention

The present invention relates to triboelectric nanogenerators, to elements for triboelectric nanogenerators and to methods for the manufacture of triboelectric nanogenerators and of elements for triboelectric nanogenerators. Related art

The rapidly growing demand for energy solutions for autonomous, wireless, portable and wearable electronic devices has prompted great interest in environment-friendly energy harvesting devices. In this respect, harvesting energy from ubiquitous mechanical vibrations can provide a viable power solution for these devices such as wireless sensors that may be implanted within the body for healthcare monitoring, or embedded within structures for early fault detection [Ref. 1 ]. Mechanical-to-electrical conversion can be achieved via electrostatic and electromagnetic generators, as well as through

piezoelectric materials [Refs. 1 , 2]. A more recent approach has been through the use of triboelectric generators based on contact electrification and electrostatic induction between materials having dissimilar electron affinities [Refs. 3, 4]. If two materials with different electron affinities make contact, the resulting charge transfer causes the material which gains charge to become negatively charged, and the other material which loses charge becomes positively charged; a phenomemon known as contact

electrification. When the two materials are separated, if there is no conducting path between the two surfaces, then these surfaces are able to maintain their induced charges as static electricity; this process is called electrostatic induction. A periodic potential difference can thus be generated across the materials as a result of a periodic relative motion between the two. Maximising the surface area over which the charge transfer takes place enhances the effect, and thus several nanostructuring routes have been adopted to achieve this, resulting in many different kinds of triboelectric

"nanogenerators" (TENGs) that have been reported over the past few years [Refs. 5-9].

For an efficient TENG device, the appropriate pairing of materials should be considered. Versions of empirical "triboelectric series" have enabled such selection through careful consideration of the positions of different materials with respect to one another in the series [Refs. 3, 10, 1 1]. Although the sequence of this triboelectric series can be affected by many variables, such as electron affinity, surface structure, and dielectric permittivity, etc., materials on the positive and negative series tend to consist of negatively charged and positively charged molecules, respectively [Refs. 12, 13]. As a result, TENG devices based on pairs of materials located on the extreme opposite ends of the triboelectric series are expected to show superior mechanical energy harvesting capability.

Interestingly, the majority of the research to date has almost exclusively focused on tribo- negative materials, in other words, electron-accepting materials, such as

polytetrafluoroethylene (PTFE) [Ref. 4]. These have been paired with aluminium or copper as the tribo-positive or electron-donating counterpart, even though these metals are not located on the extreme positive end of the triboelectric series. This is because the vast majority of materials on the positive side of the triboelectric series are biological or natural materials, such as human skin and cotton, and have relatively low mechanical stiffness and/or shape controllability [Ref. 14]. An indicative triboelectric series is illustrated in Fig. 1. SUMMARY OF THE INVENTION

Nylon, however, as shown in Fig. 1 , is an example of an exception among tribo-positive materials, being synthetic in nature with excellent mechanical properties that allow for easy control of its shape and subsequent integration into TENG devices. Therefore, the present disclosure is based on an investigation into Nylon as a potential tribo-positive candidate, seeking to enhance TENG performance and extend the range of TENG application. There have been a few reports where Nylon has been used as a tribo-positive material [Refs. 15-17], but these studies have focused on even-numbered Nylons, such as Nylon- 6 or Nylon-66, showing only slightly better output performance than other tribo-positive materials such as silk or aluminium [Ref. 18]. Odd-numbered Nylons, such as Nylon-1 1 , exhibit ferroelectric properties because of the dipole orientation resulting from the arrangement of polyamide molecules within adjacent chains, similar to more well-known ferroelectric polymers such as poly(vinylidene fluoride) (PVDF) and its copolymers [Refs. 2, 19-21]. Compared to these fluorinated polymers, Nylon-1 1 shows a similar degree of ferroelectric properties, such as dipole moment [Ref. 22] and piezoelectric coefficient [Ref. 23], with the added advantage of higher thermal stability [Refs. 24, 25]. The effect of remnant polarization in the performance of TENG devices based on ferroelectric polymers has been previously studied, whereby a high voltage was applied across the polymer to achieve dipolar alignment along the direction of the applied field. According to these studies, positively poled PVDF and its copolymer films showed more than double the power output when incorporated into TENGs as compared to the un-poled film [Refs. 17, 26]. This is due to an increase in surface charge density as a result of the remnant polarization [Ref. 27]. These studies indicate that poled Nylon-1 1 may be well-suited for TENG applications through inherent surface charge density modification. However, in order to maximize the required polarization property of this material, extreme processing conditions are required, including mechanical stretching and/or electrical poling process under high voltage (about 140 MV/m), which are major processing issues that need to be

addressed. In order to realise better polarization in Nylon-1 1 , the present inventors have realised that the material should preferably possess a pseudo-hexagonal structure with randomly oriented hydrogen bonds, referred to as the "δ'-phase". This is discussed in more detail below. This structure is considered to be beneficial for aligning the dipole moment [Refs. 28, 29]. This δ'-phase, however, is typically achieved through extremely fast

crystallization that is required to avoid the formation of large domain size [Refs. 30-33]. As a result, most of the studies regarding the δ'-phase Nylon-1 1 have been carried out on films grown via melt-quenching.

Although several alternative techniques, such as spin coating [Ref. 34], vapour deposition [Ref. 35], electrospinning [Ref. 36], thermal annealing [Ref. 37], and adding carbon nanostructures [Ref. 38], have been suggested, most of these reported approaches resulted in a lower crystallinity and/or still require harsh processing conditions, such as high temperature (290 °C) and high voltage (40 kV). Moreover, these methods involve further drawing and electrical poling process to enhance the remnant polarization of the Nylon-1 1 films [Ref. 23].

Based on the insight developed above, the present inventors have carried out research to reduce, ameliorate, avoid or overcome at least one of the above problems. The present inventors have realised that improved triboelectric generators are possible in which at least one of the triboelectric generator elements comprises a templated array of nanowires of a suitable triboelectric material. This constitutes a general aspect of the present invention. Although the insight developed above related to identification of Nylon-1 1 as a suitable triboelectric material, the applicability of the developments by the inventors is not necessarily limited to this material or related materials. Accordingly, in a first preferred aspect, the present invention provides a triboelectric generator having a first generator element and a second generator element, the first and second generator elements being arranged so that relative movement between them generates a potential difference between them due to a triboelectrification effect, wherein:

the first generator element comprises a first triboelectric material having a first electron affinity;

the second generator element comprises a second triboelectric material having a second electron affinity, different to the first electron affinity; and

the first generator element comprises a template structure having an array of channels extending in the template structure, the channels being substantially filled with the first material to define a templated array of nanowires of the first material.

In a second preferred aspect, the present invention provides a method for the manufacture of a triboelectric generator element comprising a templated array of nanowires of a first material, wherein a solution of the first material is allowed to fill an array of channels extending in a template structure by capillary wetting and the solvent is removed from the solution in the channels to solidify the first material into an array of self-poled nanowires.

In a third preferred aspect, the present invention provides a method for the manufacture of a triboelectric generator according to the first aspect, the method including

manufacturing the first generator element including the step:

a solution of the first material is allowed to fill an array of channels extending in a template structure by capillary wetting and the solvent is removed from the solution in the channels to solidify the first material into an array of self-poled nanowires,

the method further including the step of assembling the first generator element with the second generator element so that relative movement between them generates a potential difference between them due to a triboelectrification effect. In a fourth preferred aspect, the present invention provides a method of operating a triboelectric generator according to the first aspect, the method including causing relative movement between the first and second generator elements to generates a potential difference between them due to a triboelectrification effect.

The first, second, third and/or fourth aspect of the invention may have any one or, to the extent that they are compatible, any combination of the following optional features.

Preferably, the first material is a tribo-positive material.

The first material may, for example, be a polymeric material. Preferably, the polymeric material is capable of being solution-processed. In this regard, it is preferred that the polymeric material is capable of solidification from a solvent.

The first material may be a polymer with hydrogen bonding. For example, polymers containing carbonyl, carbonate and/or hydroxyl groups may be suitable. Therefore suitable example materials include Nylon, polyurethane, poly(methyl methacrylate) (PMMA). Poly(methacrylic acid) (PMAA), poly(acrylic acid) (PAA), polyvinyl alcohol) (PVA), poly(4-vinylphenol) (PVP).

The first material may, for example, be a Nylon material. Of particular interest here are odd-numbered nylons. In one preferred embodiment, the first material is Nylon-1 1 . In this case, preferably the material includes the pseudo-hexagonal polymer structure or the hexagonal/pseudo-hexagonal crystal structure. The material may include the δ'-phase.

Preferably, the degree of crystallinity of the first material is at least 30%, wherein the degree of crystallinity is determined using DSC according to the equation

AH

Crystallinity (%) =— ψ x 100 (%)

AH _m wherein AH _m is the equilibrium heat of fusion enthalpy of the first material and AH° _m is the equilibrium heat of fusion enthalpy of the perfect crystalline equivalent composition of the first material, and wherein AH _m is determined from the area under the DSC melting peak. Preferably, the degree of crystallinity of the first material is at least 35%, or at least 40%.

Preferably, the nanowires of the first material are self-poled.

Preferably, the template structure has a base face and a top face. The channels preferably open at the base face and the top face. In the method, the solution of the first material is allowed to fill the channels by contacting the base face of the template structure with the solution. The solution of the first material may therefore fill the channels by capillary wetting. Preferably, the removal of the solvent from the solution in the channels is assisted by a gas flow. Preferably, the gas flow is directed substantially parallel to the top face of the template structure. The gas flow speed may be at least 0.5 ms ^-1 . More preferably, the gas flow speed is at least 1 ms ^"1 , at least 1 .5 ms ^-1 , at least 2 ms ^-1 , at least 2.5 ms ^"1 , or at least 3 ms ^-1 . It is considered that operating with a gas flow in these ranges provides a suitable control over the removal of the solvent and therefore over the temperature distribution in the solution of the first material. In turn, this provides control over the crystallisation of the first material.

Preferably, crystals of the first material nucleate at a free surface of the solution exposed in the channel.

Additionally or alternatively, crystals of the first material nucleate as lamellae adjacent to the internal wall of the channel. In the resultant nanowires, preferably there are formed concentric lamellae, oriented substantially parallel to the internal wall of the channel. In this manner, it is found to be possible that the nanowires form with a high degree of crystallinity with a high proportion of a phase favourable for triboelectric properties.

Preferably, the width of the channels is at least 50 nm. This permits the solution to infiltrate the channels by capillary wetting. Preferably, the width of the channels is at most 500 nm. It is considered that having channels wider than this causes difficulties with the control of the crystallisation of the first material from the solution. The width of the channel may be expressed as a diameter, this to be understood as the diameter of a circle of equal area as the cross sectional area of the channel.

Further optional features of the invention are set out below.

BRIEF DESCRIPTION OF THE DRAWINGS Embodiments of the invention will now be described by way of example with reference to the accompanying drawings in which:

Fig. 1 shows a schematic triboelectric series of common materials, showing tribo-positive materials at the top and tribo-negative materials at the bottom.

Fig. 2 shows a schematic overview of the nanowire fabrication procedure used in preferred embodiments of the invention.

Fig. 3 shows a cross-section SEM image of a nanowire-filled AAO template. White threads indicate the Nylon-1 1 nanowires, which are stretched during the template breaking process.

Fig. 4 shows an SEM image of template-freed nanowires.

Fig. 5 shows SEM images of a single strand of Nylon nanowire, freed from the template. The right hand image shows an enlarged view of the region indicated in the left hand view.

Fig. 6 shows, on the left hand side, XRD patterns of nanowire-filled templates crystallized at various assisted gas-flow rates. Fig. 7 shows, based on the normalized XRD patterns, the average intensity of the peak at 21.6° and 22.6° plotted as a function of assisted-gas flow rate.

Fig. 8 shows, on the left hand side, XRD patterns of a nanowire-filled template (solid line, top), melt-quenched film (dotted line, middle), and silicon background (dashed line, bottom). The inset indicates the orientation of the nanowire-filled template with respect to the x-rays used in XRD. The right hand side of Fig. 8 shows a magnification of the δ'- phase range.

Fig. 9 shows an SEM image of an AAO template structure.

Fig. 10 shows Fig. 9 transformed into the black and white image, to allow image processing.

Fig. 1 1 shows the symmetry geometry used in simulations of nanowire formation.

Fig. 12 shows a perspective view of the numerical simulation results of turbulence flow generated by assisted gas-flow (3.1 m/s).

Fig. 13 shows a xz-plane view of the numerical simulation results of turbulence flow generated by assisted gas-flow (3.1 m/s).

Fig. 14 shows a plot of the relationship between the velocity of turbulent flow at different heights above the solution surface and assisted gas-flow rate.

Fig. 15 shows a perspective view of the numerical simulation result of heat transfer around the solution-filled nano-pore.

Fig. 16 shows a xz-plane view of the numerical simulation result of heat transfer around the solution-filled nano-pore. There is turbulent flow, and Fig. 16 shows two dominant cooling mechanisms: (i) evaporative cooling and (ii) thermal conductive cooling.

Fig. 17 shows the temperature gradient from the centre of the solution to the air (triangle points) and template wall (circular points). The initial temperature of the Nylon solution and assisted gas were taken to be 70°C and 20°C, respectively.

Fig. 18 shows a DSC thermogram of template-freed δ'-phase nanowire during the first heating with glass transition temperature Tg and melting temperature Tm.

Figs. 19A, 19B and 19C show schematic progressive cross sectional views of the polymer crystallisation process in the nano-dimensional pore of the GANT infiltration method. Fig. 20 shows DSC thermograms for the melt-quenched film (dotted line, top), additionally stretched film (dashed line, second from top), nanowires inside the template (solid line, second from bottom), and template freed nanowires (dot-dash chain line, bottom).

Fig. 21 shows FT-IR absorbance spectra for template-freed nanowires (dotted line) and additionally stretched film (solid line).

Fig. 22 shows a schematic view of the FT-IR sample for the stretched film. The IR spectra direction (red arrow) for the FT-IR measurement with the draw direction of the polymer materials (grey arrow).

Fig. 23 shows a schematic view of the FT-IR sample for the template-freed nanowires mat. The IR spectra direction (horizontal arrow) for the FT-IR measurement with the draw direction of the polymer materials.

Fig. 24 shows KPFM potential images of the melt-quenched Nylon-1 1 film and the top surface of the self-poled nanowires filled template device.

Fig. 25 shows a plot for the surface charge potential difference of the melt-quenched Nylon-1 1 film and the self-poled nanowires filled template device. The inset shows KPFM measured surface structure.

Fig. 26 shows open-circuit output voltages of TENGs with different combinations of materials. The output voltage increases from about 40 V in the device with aluminium to about 1 10 V in the device with δ'-phase Nylon nanowires.

Fig. 27 shows short-circuit output current densities of the same TENGs as in Fig. 26. Fig. 28 shows the power density of different TENG devices as a function of variable load resistance.

Fig. 29 shows a plot of charge accumulation of the nanowire-based TENG with respect to time. The left top and inset shows the circuit structure for the charge accumulation test. The right bottom inset shows a schematic view of the TENG.

Fig. 30 shows a schematic view of a TENG and various dimensions and parameters exemplified in Table 3.

Figs. 31 -34 show the results of simulations on the triboelectric potential difference of three different TENGs. Figs. 35-37 show the current density plotted against time for operation of three different TENGs.

Fig. 38-40 show the gradual increase and a decrease of voltage and current density across the various load resistors, respectively, for a TENG formed using: a nanowires filled template (Fig. 38); a melt-quenched film (Fig. 39); and aluminium (Fig. 40). In these plots, the power density is calculated by the multiplication of current density squared and load resistance.

Fig. 41 shows an SEM image of a template freed nanowire mat.

Fig. 42 illustrates the random direction of the remnant polarisation.

Fig. 43 shows an SEM image of region "b" of Fig. 41.

Fig. 44 shows an SEM image of region "c" of Fig. 43.

Fig. 45 shows the Voc performance of a TENG formed using a template freed nanowire mat.

Fig. 46 shows the Jsc performance of a TENG formed using a template freed nanowire mat.

Fig. 47 shows TENG performance under various input conditions. Short circuit current density of the nanowire device was measured under the application of a periodic impacting force at variable frequency between 2 Hz and 20 Hz with amplitude of 6 V. Fig. 48 also shows TENG performance under various input conditions. Short circuit current density of the nanowire device was measured under the application of a periodic impacting force at different amplitude between 3 V and 12 V with frequency of 5 Hz. Note that in the energy generator system, force amplitude of the magnetic shaker can be controlled by the applied voltage.

Fig. 49 shows the results of fatigue testing, carried out by recording the short circuit current density over time in response to continuous impacting at a frequency of 5 Hz and amplitude of 6 V on the same Nylon-1 1 nanowire based TENG device for 30 h (about 540,000 cycles impacting cycles in total). Data were recorded after 2 h (18 k cycles), 5 h (90 k cycles), 10 h (180 k cycles), 20 h (360 k cycles), and 30 h (540 k cycles).

Fig. 50 shows TENG device performance under various humidity conditions. The maximum (upper line) and minimum (lower line) peak of short circuit density were collected at certain humidity, the right hand image illustrating the collection of these maximum and minimum values.

DETAILED DESCRIPTION OF THE PREFERRED EMBODIMENTS, AND FURTHER OPTIONAL FEATURES OF THE INVENTION

Triboelectric nanogenerators (TENGs) have emerged as potential candidates for mechanical energy harvesting, relying on motion-generated surface charge transfer between materials with different electron affinities. In order to achieve suitable levels of energy harvesting performance, and considering the triboelectric series illustrated schematically in Fig. 1 , materials with electron-donating tendencies have not been studied sufficiently in the past. It is important to study these because they are far less common than electron-accepting counterparts. Nylons are notable synthetic organic materials with the electron-donating property, with odd-numbered Nylons such as Nylon- 1 1 , exhibiting electric polarization that can further enhance the surface charge density that is considered to be important to TENG performance.

The fabrication of Nylon-1 1 in the polarized δ'-phase typically requires extremely rapid crystallization, such as melt-quenching, as well as "poling" via mechanical stretching and/or large electric fields for dipolar alignment.

In the preferred embodiment of the present invention described below, there is disclosed a one-step fabrication process for forming a templated array of nanowires suitable for use as a TENG element. The preferred material provides enhanced surface charge density of highly crystalline "self-stretched" and "self-poled" δ'-phase Nylon-1 1 nanowires using a novel, facile gas-flow assisted nano-template (GANT) infiltration method. As will be explained in more detail below, when incorporated in a TENG, these Nylon-1 1 nanowires showed a ten-fold improvement in output power density as compared to a conventional aluminium based TENG device, under similar mechanical excitation. The embodiment uses a one-step, near room-temperature method. In this disclosure, this is termed a gas-flow assisted nano-template (GANT) infiltration method. This promotes the fabrication of highly crystalline "self-poled" δ'-phase Nylon-1 1 nanowires, the nanowires being formed within nanoporous anodized aluminium oxide (AAO) templates. The reference to "one-step" indicates that there is no need for subsequent processing such as stretching and/or poling. The gas-flow assisted method allows for controlled crystallization rate that manifests as a rapid solvent evaporation and a suitable temperature gradient within the nanopores of the template. This is predicted by finite- element simulations, resulting in the δ'-phase crystal structure. Furthermore, self- stretching and preferential crystal orientation originating from template-induced nano- confinement effects lead to self-polarization of the nanowires, increasing average crystallinity up to 40%. The self-poled δ'-phase Nylon-1 1 nanowires show enhanced surface charge distribution when compared to melt-quenched films as observed by Kelvin probe force microscopy (KPFM). Correspondingly, a TENG device based on GANT Nylon-1 1 nanowires showed a ten-fold increase in output power density compared to an aluminium based TENG device, when subjected to identical mechanical excitations.

Fabrication of Nylon-1 1 nanowires

Calculated wt% Nylon-1 1 (Sigma-Aldrich, [-NH(CH2)ioCO-] _n ) solutions were prepared in the solvent formic acid (Sigma-Aldrich, Reagent grade≥ 95%) at 70°C. When observing the nanowires using the GANT method, 25 mm diameter AAO templates (Anopore, Whatman) with a diameter of 200 nm and thickness of 60 μηη were placed on an 800 μΙ_ Nylon-1 1 solution droplet. To control the crystallization rate of the solution, assisted gas (air) flow was introduced upon the AAO template using a portable mini fan placed immediately next to the floating template. The rate of assisted gas was controlled by fan rotation speed and measured by an anemometer. The whole drying experiment was proceeded under room temperature. Detailed descriptions of the nanowire preparation set-up and treatment were reported previously [Ref. 25]. To obtain the Nylon-1 1 nanowire arrayed mats, we immersed the nanowire-filled AAO template in a 40 vol% phosphoric acid solution for 4 hr. Fabrication of Nylon-1 1 films

Nylon-1 1 films were produced by quenching the molten film (210 °C) into an ice bath. In the case of stretched films, quenched films were drawn to a draw ratio of 3 at room temperature.

Characterization

The Nylon-1 1 nanowires were visualized using field-emission scanning electron microscopy (FE-SEM, FEI Nova Nano SEM)). The crystal phase and crystalline quality of the Nylon-1 1 nanowires inside the AAO template and Nylon-1 1 films were measured using an X-ray diffraction (XRD) machine (Bruker D8) using Cu Ka radiation (λ = 1 .5418 Au) with a background silicon substrate. Numerical modeling of gas flow, heat transfer around nano-pores, and the triboelectric effect were achieved by COMSOL Multiphysics. Differential scanning calorimeter (DSC) was carried out using TA Instruments Q2000 DSC at a scanning rate of 5 °C/minute. Around 2 mg of samples was used and sealed into T zero aluminum DSC pans. Fourier transform infrared (FTIR) was performed using a Bruker Tensor 27 IR spectrometer in the reflection mode. Kelvin probe force microscopy (KPFM) measurements were carried out using Bruker Multimode 8 with Antimony (n) doped Si (tip radius < 35 nm, resonance frequency 150 kHz). AC voltages were applied from a lock-in amplifier. Film thickness was measured by stylus surface profilometer (Veeco Dektak 6M). To study the electrical output performance of the TENG device, a self-made energy harvesting measurement setup [Ref. 41 ] was utilized, recording the output open-circuit voltage (Voc) and short-circuit current (Isc) by a multimeter (Keithley 2002) and a picoammeter (Keithley 6487), respectively. In the present work, an impacting frequency of 5 Hz with an amplitude of about 1 mm was used. 24 μηη thick Aluminum film and 60 μηη thick melt-quenched Nylon-1 1 film with the same diameter (about 2 cm) were prepared to compare the TENG device performance with the nanowire one. To make the TENG device, one side of the Nylon-1 1 NW filled AAO template and the melt-quenched film was coated by sputter coating about 100 nm thick Au (using k550 Emitech) layer.

Results and discussion

Remnant polarization, without the need of subsequent stretching and/or high-voltage poling, is achieved via nanoconfinement during the template wetting process [Refs. 39- 41]. This method has been shown to enable self-poling of a ferroelectric polymer through grapho-epitaxial alignment in the lamellae [Refs. 20, 25, 41 , 42]. Recent work from our group has shown that the template-induced nanoconfinement led to self-poling in poly(vinylidene difluoride-trifluoroethylene) (P(VDF-TrFE)) nanowires [Ref. 20] as well as Nylon-1 1 nanowires [Ref. 25], resulting in a highly efficient piezoelectric nanogenerators [Ref. 41]. However, controlling the exact crystal structure through nanoconfinement methods remains challenging, in part owing to the crystal growth initiation site. For a typical template-wetting process where a polymer melt or solution is introduced into the top surface of the template via spin coating, nucleation is initiated from the surface of the bulk solution and propagated into the nano-sized pores [Ref. 43]. Such bulk-initiated nucleation and growth mechanism imposes restrictions on the control of polymer crystallinity, due to the difficulties in controlling the rate of crystallization. We mitigate this problem by using a reverse template infiltration technique with assisted gas flow.

Capillary forces allow the target solution to infiltrate from the bottom, directly exposing the solution to air. Furthermore, assisted gas flow enables precise control of the crystallization rate of the Nylon-1 1 polymer solution, resulting in the desired δ'-phase crystal structure. We note that this phase was not prominent in the XRD diffraction pattern reported in our previous work [Ref. 25] where a more typical template-wetting method was employed. We find that the strong presence of the δ'-phase in Nylon-1 1 nanowires leads to enhanced surface polarisation, and thus improved control of this crystalline phase through our GANT method is useful in enhancing the triboelectric energy harvesting performance. The GANT fabrication procedure is schematically illustrated in Fig. 2. To achieve δ'- phase Nylon-1 1 nanowires, the AAO template (diameter about 2 cm and thickness about 60 μηη) was placed on top of a 17. 5 wt% Nylon-1 1 solution in formic acid at 70 °C. At the same time, assisted gas flow (about 3 m/s) was introduced in a direction parallel to the template surface.

Stretched and entangled Nylon-1 1 nanowire strands (white thread-like regions) were observed in cross-sectional SEM images of a cleaved template - see Fig. 3. The observed deformation of the nanowires in Fig. 3 was a result of the template cleaving process. In contrast, well-aligned nanowire arrays were detected in samples after the template material was dissolved using mild acid. For example, Fig. 4 shows an SEM image of nanowires freed from the template. A single nanowire strand has uniform width and length of 200 nm and 60 μηη respectively, which are similar to the dimensions of the AAO template pore channels. Fig. 5 shows an SEM image of an individual nanowire, at low (left hand image) and high (right hand image) magnifications.

Gas-flow assisted crystal structure control

Detailed crystal structural characterization was subsequently carried out by X-ray diffraction (XRD, Bruker D8). At room temperature, a melt-quenched film with pseudo- hexagonal δ'-phase typically showed broader reflection peak at 2Θ of about 21 .6° corresponding to (hkO) plane, which is merged planes of (100), (010), and (1 10) [Ref. 44]. However, without assisted gas-flow, Nylon-1 1 nanowires displayed a relatively strong peak at 2Θ = 22.6° with a very weak δ'-phase peak (2Θ = 21 .6°). This is shown in Figs. 6 and 7. This diffraction pattern is consistent with our previous result regarding the Nylon-1 1 nanowire manufactured by a template wetting method [Ref. 25]. The GANT infiltration method allowed control of the crystal structure, wherein we were able to manipulate the rate of crystallization by adjusting the speed of gas flow. As shown in Fig. 6 (left panel), the relative peak intensity of the δ'-phase gradually increased with increasing the rate of assisted gas-flow up to a gas-flow rate of about 3 m/s, without any further increment thereafter. The relative changes in intensities between the peak at 21.6° and peak at 22.6° are depicted in Fig. 7, where the variation in average peak intensities is plotted as a function of gas flow rate. This result indicates that the crystal structure of nano-confined Nylon-1 1 nanowires could be well-controlled using different gas-flow rate, leading to the formation of pseudo-hexagonal δ'-phase Nylon-1 1 nanowires. It should be noted that there exists some confusion about the distinction between γ and δ-phase of Nylon-1 1 . Based on the recent work of Pepin et al., we identify the δ' phase in our Nylon-1 1 nanowires which has the required polar

characteristics [Ref. 44]. The crystalline phase corresponding to the 22.6° peak is less straightforward to identify as it has not been reported in the literature, though a possible explanation may be that this peak corresponds to the (210/010) plane of triclinic α'- phase.

In order to compare the crystallography of the nanowires, melt quenched Nylon-1 1 films with δ'-phase, were grown. Fig. 8 shows the XRD from both a melt-quenched film and nanowires in the AAO template. While the melt-quenched film showed a relatively broad δ'-phase peak at 20hk0 = 21.6°, the δ'-phase nanowires that were grown using the GANT method exhibited a δ'-phase peak with small full width at half-maximum. This means that the smectic phase has more organized (pseudo)-hexagonal lattice structure [Ref. 44]. It is to be noted that the volume of nanowires was only about 50% of the film with the same size and thickness. This was determined by imaging the bare AAO template by SEM (Fig. 9) and then transforming this into a black and white image (Fig. 10). The number of black pixels (pores) was counted and divided by the total number of pixels (whole surface area of the template). As a result, the average effective area was calculated to be about 48.25 %. Note that the XRD data for nanowires was taken while the nanowires were still embedded in the AAO template.

It can therefore be surmised that the actual amount of crystalline region in the nanowire per unit area is larger than that of the melt-quenched film. These results suggest that even through the conventional melt-quenching method, the presence of large amorphous regions cannot be prevented in the films due to the thickness of the film and the low thermal conductivity of Nylon [Ref. 33]. However, this problem is mitigated via the GANT method where the δ'-phase in Nylon-1 1 nanowires is realised to a larger extent.

The crystal structure of Nylon-1 1 have been extensively studied due to their extensive degree of polymorphism [Refs. S1-S6]. Nylon-1 1 displays several crystal structures depending on the processing condition as summarised below.

Table 1. Various crystal structures of Nylon-1 1

At room temperature, in Nylon-1 1 there have been demonstrated at least four crystalline forms (the triclinic a and α', the monoclinic β, and the pseudo-hexagonal γ phase) and one smectic-like pseudo-hexagonal phase [Ref. 44].

In terms of electrical property, although all Nylon-1 1 has a polar crystal structure due to its molecular configuration, the electric polarisation can be maximised from a specific type of crystalline phase. In the case of the most stable and well-organized a-phase, Nylon-1 1 does not display remnant polarisation due to the strong hydrogen bond, which is originated from the highly-packed crystal structure. In contrast, the highest electric polarisation could be observed from the meta-stable pseudo-hexagonal phase (γ, γ', δ, δ') [Refs. 28, 29, 33]. Disordered, short-range hydrogen bond and breakage of gauche bonding originated from rapid crystallisation are likely to enable dipole reversal [Ref. 33]. Despite disordered configuration in the perpendicular direction to the chain and randomly oriented hydrogen bonds, pseudo-hexagonal phase shows more ordered crystalline structure along the chain direction, which was referred to be a smectic-like phase, with aligned amide groups [Ref. S1 1].

The crystallization mechanism initiated by the GANT method was verified using finite element analysis using COMSOL Multiphysics.

Using COMSOL Multiphysics simulation, the simulation is demonstrated based on these three different effects: the turbulent flow of the assisted gas, heat transfer in all components, and the vaporisation of the solvent in the Nylon solution. For the turbulent flow, we assumed that the gas flow rate and pressure field are independent of the property of gas, such as moisture content level and temperature. Heat transfer in the model is considered to have two different aspects: conduction and convection. The heat transfer between the template wall and the solution is governed by conduction. In the gas flow, the heat transfer is originated from the convection and the effect of turbulent flow. The cooling effect during the solvent evaporation needs to be considered as well based on the heat of vaporisation H _va . To calculate the amount of vaporised solvent to the air, the material transport equation is used with the turbulent flow as a diffusion coefficient.

Fig. 1 1 illustrates the symmetry used in the COMSOL Multiphysics simulation. The geometry was built based on the experimental conditions. The inner diameter and thickness of the nanopore is 200 nm (measured from SEM images) and 100 nm, respectively. The height of the nanopore is assumed 5 μηη filled with 70°C formic acid with H _V a of 23.1 kJ/mol. The assisted gas is air with an initial temperature of 20°C and enters to the right side of the geometry. The inlet velocity of the turbulent flow was set by the gas flow rate. To further investigate the rapid crystallization process from the polymer solution, both the solvent evaporation and temperature of the solution have been considered. In terms of solvent evaporation, our simulations revealed that through the assisted gas-flow, the artificially generated dry and cool turbulent air flow effectively encourages the

evaporation of the solvent.

Figs. 12 and 13 show shows the induced turbulent flow by assisted gas within the nanopore channels. Fig. 14 shows that the speed of the induced turbulent flow increases with increasing assisted gas-flow velocity.

The velocity of the turbulence, evaluated 10 nm above the surface of the solution, was 36 μηΊ/s for an assisted-gas velocity of 3.1 m/s. This is a high flow rate considering the size of nanopores are about 200 nm in diameter. As a result, nucleation can be initiated at the surface of the exposed solution by rapid solvent evaporation.

In the computational model, we also considered the effect of the solution cooling within the pore channel: both turbulent flow induced evaporative cooling and thermal conductive cooling via the nano-template walls. The overall cooling mechanism is demonstrated in Figs. 15 and 16, with the turbulent flow around the exposed Nylon solution and heat transfer through the pore walls. As a result of cooling, a temperature gradient is generated both on the top surface and side of the Nylon solution. Fig. 17 shows the significant temperature gradient which is produced by evaporative cooling and thermal conduction as a function of the distance from the top and side surface of the solution, respectively. The crystal growth direction can be inferred from these heat transfer simulations, as the growth of the crystalline region is most likely to occur in the direction of the largest temperature gradient (see Figs. 19A-19C). In addition, considering the crystal growth direction of the GANT method, due to the size of the nano-pores, the crystal growth length (about 100 nm) can be limited as compared to the melt-quenched film (about 30 μηη), resulting in extremely rapid nucleation and growth. This is illustrated by the schematic illustrations of Figs. 19A, 19B and 19C showing the polymer crystallisation process in the nano-dimensional pore of the GANT infiltration method. Therefore, fast crystallization is achieved within the confines of the nanopores, even with mild external conditions, such as low gas flow rate (3 m/s) and near-room temperature conditions.

Self-poled δ'-phase nanowires through GANT method

Differential scanning calorimetry (DSC) was carried out to determine the thermal and structural properties of Nylon-1 1 nanowires as shown in Fig. 20, from which the melting temperature (Tm) and the melt crystallization temperature (Tc) were recorded. To investigate the drawing effect within the template pores, as-received melt-quenched films, stretched melt-quenched films and template-freed nanowire mats were also prepared. During the heating cycle, a single melting peak at 190 °C was observed in the melt-quenched film corresponding to the δ'-phase crystal formation [Refs. 34, 44]. In contrast, a significant change of melting points was obtained from the additionally stretched film as compared to the non-stretched film. The stretched film showed a dominant melting peak at 184.5 °C with a minor melting peak (190 °C), showing an increase in crystallinity from 29 % to 40 % [Ref. 32].

Using the average data of the four distinct DSC peaks of Nylon-1 1 , the crystallinity was calculated by means of the equation

AH

Crystallinity (%) =— x 100 (%)

AH _m

where AH _m and AH° _m are the equilibrium heat of fusion enthalpies of the semi-crystalline Nylon-1 1 samples and the perfect crystalline Nylon-1 1 (AH° _m is about 189 J/g) [Ref. 48], respectively. AH _m is achieved from the area under the DSC melting peak. The results of the calculations are given in Table 2 below. Table 2 - Enthalpy of fusion and degree of crystallinity of various morphologies of Nylon-

Film Stretched film Nanowire

w _m (Jig) 54 76 74

Crystallinity (%) 29 40 39

Typically, the degree of crystallinity of melt-quenched films is usually around 30% since the rate of crystallisation is very fast compared to other polymers [Refs. 33, 48]. The crystallinity of δ'-phase nanowires is similar to the crystallinity of stretched film due to the self-stretching effect.

Returning now to a consideration of Fig. 20, the formation of a new low-temperature melting peak and improved crystallinity may indicate the additionally aligned molecular chain structure developed by the mechanical stretching process [Ref. 45]. During the cooling cycle, similar to the melting point shift described above, we found that stretched films crystallize at higher temperatures (162°C) compared to the non-stretched film (158°C) due to a more ordered crystal structure.

The thermal properties of the nanowires can be explained on the basis of several factors nanowire size effect, nano-template effect and crystal ordering effect. Both nanowires within the template, as well as those freed from the tempate displayed double melting behaviour. The additional low temperature melting peak can be explained by nanoscale size effect. According to the Gibbs-Thomson equation [Refs. 46,47], the melting point depression AT_m is given by

m = T _m - r _m (d) = o _sl T _m /{dLH _f p _s )

where T _m is the normal (bulk) melting point, T _m (d) is the melting point of crystals of size d, Osi is the surface tension of the solid-liquid interface, d is particle size,AHf is the bulk enthalpy of fusion (per g of material), and p _s is the density of the solid [Ref. 46]. This shows that the melting point of the nanowires was depressed due to the nano-confined structure and surface tension of the Nylon-1 1 /template interface.

From nanowires in the template, the dominant melting peak was observed at a lower temperature (185 °C), as compared to the template-freed nanowires which displayed dominant melting peaks at a higher temperature (190 °C). This is because, in case of nanowires in the AAO template, both the very thin nanowire diameter as well as the surface tension between Nylon-1 1 and the nano-template interface affect the melting behaviour [Ref. 46]. In the cooling cycle, the nano-template effect was confirmed by the relatively broader crystallization temperature of the nanowire in the template. Thermal behaviour of template-freed nanowires, however, makes us infer an additional effect that influences the property of the nanowires. This is because cooling behaviour of template- freed nanowires and stretched film are substantially similar and have the same Tc (162 °C). In addition, the δ'-phase nanowires displayed much higher degree of crystallinity (39 %) than those of the melt-quenched film (29 %) [Refs. 33, 48]. The crystallinity of the δ'-phase nanowire is rather similar with those of stretched film (40 %), even though those nano-confined structure showed much lower latent heat of fusion AH _m [Ref. 47]. All of these phenomena suggests that, during the GANT process, δ'-phase nanowires were likely to have experienced crystalline ordering. Similar to the molecular alignment from the film-drawing process, grapho-epitaxial alignment originating from the template-pore walls resulted in highly ordered structure, giving rise to similarities in the thermal behaviour and high crystallinity observed in the stretched film. Our results indicate that the GANT method successfully enabled highly ordered δ'-phase nanowires to be formed without additional drawing process.

The molecular bond structure of δ'-phase nanowires was measured using room temperature Fourier transform infrared (FT-IR) spectroscopy and the results shown in Fig. 21 . A template freed δ'-phase nanowire mat [Ref. 25] and stretched film were prepared to confirm the preferential crystal orientation originating from nano-confinement. For more detailed analysis, the direction of the infrared spectra for drawing should be considered since the peak intensity can be changed depending on the chain alignment in the draw direction [Ref. 49]. Considering the drawing direction, parallel and

perpendicular direction infrared absorption spectra were measured from δ'-phase nanowires and stretched film, respectively (Figs. 22 and 23). According to the literature, FT-IR spectra of Nylon-1 1 have two important regions related to dipole alignment. The region 1500-1700 cm ^-1 contains the amide I and II mode and is assigned to hydrogen- bonded or free amide group. The band at 3300 cm ^-1 (amide A peak) is assigned to N-H stretching vibration and is sensitive to the hydrogen bond [Refs. 31 , 34, 50]. The region 1500-1700 cm ^"1 is shown in Fig. 21 for both stretched film and nanowires, containing the amide I and amide II bands. In this region, two materials show similar intensity except for 1635 cm ^-1 band because the conformation in the amorphous phase is expected to be the same [Ref. 34]. The 2920 and 2850 cm ^-1 bands of both materials are assigned to the antisymmetric and symmetric Ch stretching modes of the methylene groups, respectively. The significant differences in the absorption intensities of the N-H stretching modes at 3300 cm ^-1 and C=0 stretching modes at 1635 cm ^-1 show direct evidence for the self-polarization in the nanowires. The parallel direction absorbance peak in the template-freed nanowire mat was found to be higher for both bands compared to the perpendicular direction peak intensity of the stretched film. Considering the relatively higher peak intensity in the perpendicular direction spectra and peak intensity rise in the parallel spectra due to the electric poling process [Refs. 29, 49], the higher intensity of the nanowire mat suggests that nanowires obtained from the GANT method have preferential crystal orientation.

In order to investigate the direction of polarization and quantify the surface charge variations between Nylon-1 1 nanowires and melt-quenched films arising from the self- poled nature of the former, we employed Kelvin probe force microscopy (KPFM), where the surface charge potential difference between the material and an atomic force microscope (AFM) tip can be measured. The results are shown in Figs. 24 and 25.

Melt-quenched Nylon-1 1 film showed the surface potential difference of 82 mV. In contrast, the higher surface potential (426 mV) was observed from the self-poled δ'- phase nanowire-filled template. Furthermore, the enhanced surface charge potential difference of self-poled nanowires demonstrates the relationship between polarization and surface charge density. Our results indicate an "upward" direction of polarization, which corresponds to the expected charge donating property.

Nylon-1 1 Nanowire based TENG device

A TENG device was fabricated using self-poled nanowires embedded within AAO template. The bottom side of the AAO template was coated by an Au electrode (about 100 nm thickness), and an Au-coated Teflon film was prepared as a counterpart substrate. Aluminium films and melt-quenched Nylon-1 1 film were also prepared to compare the device performance with the nanowire-based TENG.

Figs. 26 and 27 show the open circuit voltage (Voc) and short circuit current density (Jsc) respectively, measured in response to the periodic impacting at a frequency (f) of 5 Hz and amplitude of 0.5 mm in an energy harvesting setup that has been previously described [Ref. 41]. The aluminium based TENG device showed a peak Voc of about 40 V and a peak JSC of about 13 mA-nr ² . Following the triboelectric series, higher TENG performance was observed from a melt-quenched Nylon-1 1 film based TENG with Voc of about 62 V and Jsc of about 21 mA-nr ² than from the Al based TENG. The self-poled Nylon-1 1 nanowire based TENG displayed further enhanced output performance with a peak Voc of about 1 10 V and a peak Jsc of about 38 mA-nr ² likely due to the self-poled nature of the nanowires.

We conducted analytical simulations using a finite element method (FEM) with COMSOL Multiphysics software to further confirm the effects of self-polarization on triboelectric potential. Fig. 30 shows the schematic structure of the TENG device and indicates the dimensions and parameters of the device used in the simulations. Table 3 shows values of the dimensions and parameters used in the simulations. Table 3 - dimensions and parameters used in simulation of TENG device

Symbol Meaning Value unit d max. separation distance 0.5 mm tl thickness of PTFE 100 μιη

thickness of Nylon film & AAO template 60 μιη

S surface area of the electrode 3.14 cm ²

VOC.AI Voc of aluminium based TENG 40 V

VoC.film Voc of melt-quenched film based TENG 60 V

VoC,NW Voc of Nylon NW filled template based TENG 1 10 V εο vacuum permittivity 8.85E-12 F / m

Sr. film relative permittivity of Nylon-1 1 film [Refs. 25, 45] 3.7

€r.AAO relative permittivity of AAO template 9

€r.NW relative permittivity of oriented Nylon [Refs.25, 45] 4.5

r.PTFE relative permittivity of PTFE 2.1

Pr remnant polarization of Nylon-1 1 [Ref. S15] 55 mC / m ²

To confirm the working mechanism of the self-poled nanowire based TENG, we carried out a theoretical analysis based on Gauss theorem [Refs. S13, 1 1]. In the dielectric-to- dielectric contact mode TENGs, the electric potential difference (AV) between two electrodes can be given by:

AV = E _x t _x + E ₂ t ₂ + E _air d

where E is electric field strength, t _x and t ₂ are the thickness of the two surfaces, d is the distance between two different layers, Q is the value of transferred charges, S is the area of the electrode, σ is the triboelectric charge density, ε ₀ is the vacuum permittivity, and e _rl and s _r2 are the relative permittivity (dielectric constant) of dielectric materials, respectively. In the conductor-to-dielectric structure, ti/e _r i can be ignored because the metal layer acts as triboelectric layer and electrode. Under open-circuit conditions, the value of transferred charges (Q) become zero since no charge is transferred between the two top and bottom electrodes. Thus, if we assume electric potential of the bottom electrode to be zero, the equation for the open-circuit voltage {Voc) can be calculated by

Vac = f (2)

Using experimentally determined V _oc of each TENGs, we can obtain a theoretical triboelectric charge density (o): When d is maximum (0.5 mm) and σ of aluminium, the melt-quenched Nylon film, and the nanowire-filled template are 0.78, 1 .06, and 1.95 μθ nr ² , respectively.

A COMSOL Multiphysics simulation demonstrated the triboelectric potential difference of three different structure TENGs with Teflon as the top part and: aluminium (Fig. 31 ), Nylon film (Fig. 32), and nanowires in the template (Fig. 33) as the bottom parts. Fig. 34 shows an enlarged view of the simulation results for the nanowire-air-alumina interface. In the simulation, we assumed that the length of the dielectric and electrodes are infinite because they are significantly larger than the thickness of the dielectric layer. The nanowire-filled alumina template structure (Fig. 33) was simplified in the simulation, considering the surface area of the nanowires (about 50 %). Because of the charge density (σ) difference, the nanowire sample shows the highest potential different between the top and bottom electrodes. In addition, the self-poled Nylon nanowire region exhibits relatively higher electric potential than the alumina template region. This indicates that further enhanced triboelectric performance of nanowire containing AAO template devices is mainly attributed to the self-polarization of the nanowires.

The ideal electric potential of self-poled Nylon-1 1 nanowire based TENGs can be calculated based on the experimental result of remnant polarisation [Ref. S15]. Under short-circuit conditions (V = 0), the transferred charges (Q _sc ) are given by Q Ξσάε _τ1 ε _τ2

The value of Q _sc can therefore be varied as a function of d. When the generator is fully released (d = 0.5 mm), the maximum transferred charges per unit area was theoretically calculated to be 0.65, 0.94, and 1 .75 μθ nr ² for aluminium, melt-quenched Nylon film, and the nanowire-filled template, respectively. The transferred charges (Q _sc ) are also calculated using the integration of short-circuit current (/ _sc ). The short circuit current is shown in Figs. 35, 36 and 37, corresponding to the devices of Figs. 31 , 32 and 33, respectively. Thus, in the work presented here, when compared with the aluminium and Nylon film based TENGs, a higher triboelectric potential is obtained from the self-poled nanowire based TENG, because the self-poled nanowires increase the charge density (o) on the electrified surfaces. The electrical power output of the TENG was measured across different resistors. Peak output power density of 1 .03 W-nr ² , 0.19 W-nr ² , and 0.099 W-nr ² were observed from the Nylon-1 1 nanowire, Nylon-1 1 (melt-quenched) film, and aluminium based device respectively under impedance-matched conditions at a load resistance of about 20 ΜΩ (Fig. 28 and Fig. 38 (nanowire-filled template), Fig. 39 (melt-quenched film) and Fig. 40 (aluminium)). The observed output power from the nanowire based TENG was about 6 times and about 10 times higher than those of a melt-quenched Nylon-1 1 film and aluminium based TENG, respectively. Such remarkable improvement in the output performance of δ'-phase nanowires can be rationalized as follows: the self-polarization of the nanowires can be expected to give rise to larger surface charge density, which can result in more transferred charges compared to the film surface. It should be noted that the surface area of the nanowires is only about 50 % as compared to the melt-quenched film. This indicates that polarization in the nanowire effectively further enhances the surface charge density of the device. The electrical output was found to increase with increasing impact frequencies and amplitudes (see Figs. 47 and 48). In addition, fatigue testing was carried out by continuously impacting the device for up to 30 hours at 5 Hz (about 540,000 cycles). Fig. 47 shows that the Nylon-1 1 nanowire-based TENG device exhibited negligible change in output current density over the entire period of continuous testing. Reliability tests under various humidity conditions were also carried out by impacting the Nylon-1 1 nanowire based TENG device within a humidity-controlled box. (See Figs. 49 and 50). Although Nylon is known to be prone to degradation in the presence of moisture, the Nylon-1 1 nanowire-based TENG showed reliable output performance up to high humidity condition (about 80 %), indicating that the AAO template serves to encapsulate and protect the nanowires from environmental factors.

To further confirm the effect of self-polarization on TENG devices, we also measured the output performance of the template-free nanowire mat where the nanowires are lying on the substrate with randomly oriented polarization directions. Fig. 41 shows an SEM image of a template freed nanowire mat. Fig. 42 shows a schematic view of the area indicated as "b" in Fig. 41 . The arrows in Fig. 42 indicate the local direction of the remnant polarisation, based on which it can be seen that there is no overall preferred direction for the remnant polarisation. Fig. 43 shows an enlarged SEM view of the area indicated as "b" in Fig. 41. Fig. 44 shows an enlarged SEM view of the area indicated as "c" in Fig. 43.

Figs. 45 and 46 show the output performance of the nanowire mat based TENG device. Interestingly, the template-freed nanowire mat-based TENG device generated

suppressed output performance (V ₀ c = 20V and Jsc = 4.0 mA-nr ² ), even though it was thinner and had a larger effective surface area. Due to the randomly oriented

polarization in the nanowire mat, the generation of triboelectric charges was likely inhibited during the contact and separation process. These results support the orientation directed polarization effects in the Nylon-1 1 nanowires while they are still aligned and embedded within the template. To confirm the energy generated by the δ'-phase nanowire, TENG is also feasible for energy storage. As illustrated in Fig. 29, a 470 F capacitor was connected to the device using a full-wave bridge rectifying circuit. The Nylon nanowire based TENG under mechanical pressure at 5 Hz for about 20 minutes successfully charged the capacitor with a charging speed of about 38 μθτηίη ^"1 . Notably, the accumulated charge increased with time as shown in Fig. 29, suggesting the Nylon nanowire-based TENG had excellent stability. In a demonstration, the electric power produced by the Nylon nanowire based TENG was used to directly turn on several commercial light-emitting diodes (LEDs). During contact and separation with 5 Hz frequency, 36 white LEDs were driven by the produced output voltage without the need for external energy storage devices.

Conclusions To summarize, we have demonstrated for the first time, a novel and facile gas-flow assisted nano-template (GANT) infiltration method for the fabrication of highly oriented, self-poled δ'-phase Nylon-1 1 nanowires, as a rarely synthesized tribo-positive material. Assisted gas-flow controlled the crystallization speed, resulting in the δ'-phase crystal structure. Preferential crystal orientation originated from the nano-confinement effect in the template, resulting in self-poling of the Nylon-1 1 nanowires with an increased average crystallinity of up to about 40%. When self-poled Nylon-1 1 nanowires were combined with counterpart tribo-negative surfaces, such as Teflon, the resulting output power were observed to be about 6 times and about 10 times higher than those of melt- quenched Nylon-1 1 film and aluminium based TENGs, respectively. The output power generated by the Nylon-1 1 nanowire-based TENG device was high enough to drive commercial electronic components such as LEDs and capacitors without external power sources. It is therefore now possible to provide enhanced surface charge density for TENG elements for use in high-performance TENG devices. While the invention has been described in conjunction with the exemplary embodiments described above, many equivalent modifications and variations will be apparent to those skilled in the art when given this disclosure. Accordingly, the exemplary embodiments of the invention set forth above are considered to be illustrative and not limiting. Various changes to the described embodiments may be made without departing from the spirit and scope of the invention.

All references referred to above are hereby incorporated by reference. List of non-patent references

1 K. A. Cook-Chennault, N. Thambi and A. M. Sastry, Smart Mater. Struct., 2008, 17, 43001.

2 S. Crossley, R. A. Whiter and S. Kar-Narayan, Mater. Sci. Technol., 2014, 30, 1613-1624. 3 Z. L. Wang, ACS Nano, 2013, 7, 9533-9557.

4 Z. L. Wang, J. Chen and L. Lin, Energy Environ. Sci., 2015, 8, 2250-2282.

5 F. R. Fan, Z. Q. Tian and Z. Lin Wang, Nano Energy, 2012, 1 , 328-334.

6 G. Zhu, Z. Lin, Q. Jing, P. Bai, C. Pan, Y. Yang and Y. Zhou, Nano Lett., 2013, 13, 847. 7 G. Cheng, Z. Lin, L. Lin, Z. Du and Z. L. Wang, ACS Nano, 2013, 7, 7383-7391.

8 C. K. Jeong, K. M. Baek, S. Niu, T. W. Nam, Y. H. Hur, D. Y. Park, G. Hwang, M. Byun, Z. L. Wang, Y. S. Jung and K. J. Lee, Nano Lett., 2014, 14, 7031-7038.

9 H. Kang, H. Kim, S. Kim, H. J. Shin, S. Cheon, J. H. Huh, D. Y. Lee, S. Lee, S. W. Kim and J. H. Cho, Adv. Funct. Mater., 2016, 26, 7717-7724. 10 J. Henniker, Nature, 1962, 196, 474.

1 1 G. Zhu, C. Pan, W. Guo, C. Y. Chen, Y. Zhou, R. Yu and Z. L. Wang, Nano Lett., 2012, 12, 4960-4965. 12 S. Shin, Y. H. Kwon, Y. Kim, J. Jung, M. H. Lee and J. Nah, ACS Nano, 2015, 9, 4621-4627.

13 S. Shin, Y. E. Bae, H. K. Moon, J. Kim, S. Choi and Y. Kim, ACS Nano, 2017, 1 1 , 6131-6138. 14 H. J. Kim, J. H. Kim, K. W. Jun, J. H. Kim, W. C. Seung, O. H. Kwon, J. Y. Park, S. W. Kim and I. K. Oh, Adv. Energy Mater., 2016, 6, 1-6.

15 T. Zhou, C. Zhang, C. B. Han, F. R. Fan, W. Tang and Z. L. Wang, ACS Appl. Mater. Interfaces, 2014, 6, 14695-14701.

16 Y. Zheng, L. Cheng, M. Yuan, Z. Wang, L. Zhang, Y. Qin and T. Jing, Nanoscale, 2014, 6, 7842-6.

17 J.-H. Lee, R. Hinchet, T. Y. Kim, H. Ryu, W. Seung, H.-J. Yoon and S.-W. Kim, Adv. Mater., 2015, 27, 5553-5558.

18 X. Pu, M. Liu, X. Chen, J. Sun, C. Du, Y. Zhang, J. Zhai, W. Hu and Z. L. Wang, Sci. Adv., 2017, 3, e1700015. 19 A. J. Lovinger, Science (80-. )., 1983, 220, 1 1 15-1 121 .

20 R. A. Whiter, Y. Calahorra, C. Ou and S. Kar-Narayan, Macromol. Mater. Eng.,

2016, 301 , 1016-1025.

21 Y. S. Choi, J. Sung, S. J. Kang, S. H. Cho, I. Hwang, S. K. Hwang, J. Huh, H. Kim, S. Bauer and C. Park, Adv. Funct. Mater., 2012, 1-9. 22 J. W. Lee, Y. Takase, B. A. Newman and J. I. Scheinbeim, J. Polym. Sci. Part B- Polymer Phys., 1991 , 29, 279-286.

23 B. A. Newman, P. Chen, K. D. Pae and J. I. Scheinbeim, J. Appl. Phys., 1980, 51 , 5161-5164.

24 Y. Takase, J. W. Lee, J. I. Scheinbeim and B. A. Newman, Macromolecules, 1991 , 24, 6644-6652.

25 A. Datta, Y. S. Choi, E. Chalmers, C. Ou, S. Kar-Narayan Adv. Funct. Mater.,

2017, 27, 1604262. 26 P. Bai, G. Zhu, Y. S. Zhou, S. Wang, J. Ma, G. Zhang and Z. L. Wang, Nano Res., 2014, 7, 990-997.

27 K. Y. Lee, S. K. Kim, J.-H. Lee, D. Seol, M. K. Gupta, Y. Kim and S.-W. Kim, Adv. Funct. Mater., 2016, 26, 3067-3073. 28 J. I. Scheinbeim, J. Appl. Phys., 1981 , 52, 5939-5942.

29 J. I. Scheinbeim, J. W. Lee and B. a Newman, Macromolecules, 1992, 25, 3729- 3732.

30 S. L. Wu, J. I. Scheinbeim and B. A. Newman, J. Polym. Sci. Part B-Polymer Phys., 1999, 37, 2737-2746. 31 H. Isoda and Y. Furukawa, J. Phys. Chem. B, 2015, 1 19, 14309-14314.

32 Q. Zhang, Z. Mo, H. Zhang, S. Liu and S. Z. D. Cheng, Polymer (Guildf)., 2001 , 42, 5543-5547.

33 Z. Zhang, M. H. Litt and L. Zhu, Macromolecules, 2016, 49, 3070-3082.

34 S. S. Nair, C. Ramesh and K. Tashiro, Macromolecules, 2006, 39, 2841-2848. 35 K. Meurisch, B. Gojdka, T. Strunskus, V. Zaporojtchenko and F. Faupel, J. Phys. D. Appl. Phys., 2012, 45, 55304.

36 M. Dhanalakshmi and J. P. Jog, Express Polym. Lett., 2008, 2, 540-545.

37 N. B. N. Wu S-L, Scheinbeim J. I., J. Polym. Sci. Part B Polym. Phys., 1996, 34, 3035-3053. 38 G. Mago, D. M. Kalyon and F. T. Fisher, J. Polym. Sci. Part B Polym. Phys., 201 1 , 49, 131 1-1321 .

39 Z. Hu, M. Tian, B. Nysten and A. M. Jonas, Nat. Mater., 2009, 8, 62-67.

40 M. C. Garcia-Gutierrez, A. Linares, J. J. Hernandez, D. R. Rueda, T. A. Ezquerra, P. Poza and R. J. Davies, Nano Lett., 2010, 10, 1472-1476. 41 R. A. Whiter, V. Narayan and S. Kar-Narayan, Adv. Energy Mater., 2014, 4, 1400519. 42 Y. Calahorra, R. A. Whiter, Q. Jing, V. Narayan and S. Kar-Narayan, APL Mater., 2016, 4.

43 M. Steinhart, P. Goring, H. Dernaika, M. Prabhukaran, U. Gosele, E. Hempel and T. Thurn-Albrecht, Phys. Rev. Lett., 2006, 97, 1-4. 44 J. Pepin, V. Miri and J. M. Lefebvre, Macromolecules, 2016, 49, 564-573.

45 P. FrQbing, A. Kremmer, R. Gerhard-Multhaupt, A. Spanoudaki and P. Pissis, J. Chem. Phys., 2006, 125, 214701-1-8.

46 C. L. Jackson and G. B. McKenna, J. Chem. Phys., 1990, 93, 9002.

47 L. H. A. S. L. Lai, J. Y. Guo, V. Petrova, G. Ramanath, Phys. Rev. Lett., 1996, 77, 99-102.

48 Q. Zhang, Z. Mo, S. Liu and H. Zhang, Macromolecules, 2000, 33, 5999-6005.

49 H. H. Yu and L. J. Fina, Macromolecules, 1994, 27, 6192-6200.

50 J. Jakes and S. Krimm, Spectrochim. Acta Part A Mol. Spectrosc, 1971 , 27, 19- 34. S1 A. Kawaguchi, I. Tokimitsu, Y. Fujiwara, M. Tabuchi and K. Monobe, J.

Macromol. Sci. Part B Phys., 1981 , 20, 1-20.

52 S. Rhee and J. L. White, J. Polym. Sci. Part B Polym. Phys., 2002, 40, 2624- 2640.

53 H. Search, C. Journals, A. Contact, M. lopscience and I. P. Address, Br. J. Appl. Phys., 1959, 10, 225-230.

54 G. H. E. Roguet, S. Tence-Girault, S. Castagnet, J.C. Grandidier, J. Polym. Sci. Part B Polym. Phys., 2007, 45, 3046-3059.

55 W.P. Slichter, J. Polym. Sci. Part A Polym. Chem., 1959, 36, 259-266.

56 R. Aelion, Ann. Chim. Appl, 1948, 3, 5-61 . S1 1 L. J. Mathias, D. G. Powell, J. P. Autran and R. S. Porter, Macromolecules, 1990, 23, 963-967. S13 S. Niu, S. Wang, L. Lin, Y. Liu, Y. S. Zhou, Y. Hu and Z. L. Wang, Energy Environ. Sci., 2013, 6, 3576.

S15 L. F. Brown, J. L. Mason, M. L. Klinkenborg, J. I. Scheinbeim and B. A. Newman, IEEE Trans. Ultrason. Ferroelectr. Freq. Control, 1997, 44, 1049-1059.