Interfacial piezoelectric polarization locking in printable Ti3C2Tx MXene-fluoropolymer composites

Piezoelectric fluoropolymers convert mechanical energy to electricity and are ideal for sustainably providing power to electronic devices. To convert mechanical energy, a net polarization must be induced in the fluoropolymer, which is currently achieved via an energy-intensive electrical poling process. Eliminating this process will enable the low-energy production of efficient energy harvesters. Here, by combining molecular dynamics simulations, piezoresponse force microscopy, and electrodynamic measurements, we reveal a hitherto unseen polarization locking phenomena of poly(vinylidene fluoride–co–trifluoroethylene) (PVDF-TrFE) perpendicular to the basal plane of two-dimensional (2D) Ti3C2Tx MXene nanosheets. This polarization locking, driven by strong electrostatic interactions enabled exceptional energy harvesting performance, with a measured piezoelectric charge coefficient, d33, of −52.0 picocoulombs per newton, significantly higher than electrically poled PVDF-TrFE (approximately −38 picocoulombs per newton). This study provides a new fundamental and low-energy input mechanism of poling fluoropolymers, which enables new levels of performance in electromechanical technologies.


Ti3C2Tx MXene nanosheets and properties
The X-ray powder diffraction (XRD) patterns of the Ti3AlC2 MAX phase and the Ti3C2Tx MXene nanosheets were obtained using a powder diffractometer (X'Pert Powder, PANalytical) equipped with a Cu Kα radiation (40 kV, 30 mA) with an X-ray wavelength (λ) of 1. 54  The Ti3C2Tx MXene nanosheets were synthesized from the Ti3AlC2 parent ternary carbide precursor (MAX phase) by selective etching of the aluminium layer (A-group element) using a mixture of lithium fluoride (LiF) and hydrochloric acid (HCl) at room temperature for 24 h. [1][2][3] The subsequent intercalation of water (H2O) molecules and Li + ions within the negatively charged surface resulted in a volume increase during washing with ultra-pure water, indicating the self-delamination of multi-layered Ti3C2Tx to few/single layers. 4 The delamination and the removal of the aluminium was confirmed by the downshifting of the (002) peak and a disappearance of the aluminium peak at 2θ of 39° in the XRD spectra ( Supplementary Fig.   1a). 5 The TEM image ( Supplementary Fig. 1b) and DLS data ( Supplementary Fig. 1c) showed that the Ti3C2Tx nanosheets exhibited an average lateral size of approximately 310 nm. The AFM image showed that the Ti3C2Tx nanosheets exhibited a clean surface at the edge ( Supplementary Fig. 1d). The thickness profile of the Ti3C2Tx nanosheets ( Supplementary Fig.   1d, inset) showed an average height of 1.6 nm, corresponding to single-layer Ti3C2Tx nanosheets. The C 1s spectral region ( Supplementary Fig. 1f), shows four singlet peaks, C-Ti, C-Ti-O, C-C, and C-O. The C-Ti peak corresponds to internal bridging C atoms. The C-Ti-O peak, occurring at slightly higher binding energies, arises from the long-range influence of oxygenbased surface termination on the electronic state of the internal C atoms. The C-C signal, while anomalous given the crystal structure of Ti3C2Tx, is always observed in literature 7 and is understood to arise from residual hydrocarbons 8 in the XPS chamber. The C-O peak occurs as Ti is an extremely mobile metal, known to leave vacancies and thus slightly altered stoichiometry. 9

Preparation of PVDF-TrFE and Ti3C2Tx/PVDF-TrFE inks
Recently, we described the dissolution and recycling of SEA extrusion printed PVDF-TrFE co-polymer films using acetone as the only solvent. 10 Here, N,N-dimethylformamide (DMF) was completely eliminated as a solvent for extrusion printing entirely and replaced by acetone.
Acetone has inherent advantages over DMF and other solvents commonly used to dissolve fluoropolymers, with faster evaporation rates that enable rapid crystallization and drying of SEA extrusion printed polymer films. 11 In particular, DMF exhibits a high boiling point (>150 °C at 101.3 kPa), 12 low vapor pressure (<0.5 kPa at 21 °C) 13 and high toxicity. 12 Comparatively, acetone exhibits a low boiling point (56 °C at 101.3 kPa), 14 high vapor pressure (26 kPa at 21 °C) 13 and reduced toxicity, reported as one of the least toxic industrial solvents, 14 and is thus better suited for SEA extrusion printing.
Initially, pristine PVDF-TrFE inks were prepared in acetone, which were used to optimize the SEA extrusion printing parameters. These inks were prepared by a simple mixing method, whereby PVDF-TrFE powder was slowly added to acetone under mechanical stirring. The pristine PVDF-TrFE co-polymer inks were prepared at PVDF-TrFE co-polymer concentrations of 35 wt%, 40 wt% and 45 wt%, based on the concentrations of inks prepared in the previously reported DMF:acetone solvent mixture. 15 The prepared inks were viscous ( Supplementary Fig.   2), moving slower when tilted to a 45° angle as the concentration increased. The rheological optimization of the inks for printing is shown further in this document. Similarly, Ti3C2Tx/PVDF-TrFE inks were prepared at Ti3C2Tx concentrations at 0.02 wt%, 0.10 wt%, 0.20 wt% and 0.50 wt%. Here, a small aliquot of the Ti3C2Tx stock dispersion in DMF (4.4 mg mL -1 ) was added to acetone to form dispersions at 0.00 mg mL -1 , 0.10 mg mL -1 , 0.52 mg mL -1 , 1.05 mg mL -1 and 2.61 mg mL -1 in acetone. Subsequently, the PVDF-TrFE powder was added slowly to the Ti3C2Tx dispersions in acetone at 23 °C while stirring, at 40 wt% relative to the mass of the dispersion, to form the Ti3C2Tx/PVDF-TrFE inks. The inks were stirred until homogeneous, then sealed with parafilm and stored at -5 °C to minimize solvent evaporation.
Throughout the experimental procedure, the stability of the Ti3C2Tx nanosheet dispersion was monitored in the Ti3C2Tx/PVDF-TrFE ink, for up to five months ( Supplementary Fig. 3, middle). The Ti3C2Tx/PVDF-TrFE ink was compared to a single-walled carbon nanotube (SWCNT)/PVDF-TrFE ink, which we have recently reported ( Supplementary Fig. 3, right). 10 Notably, after five months of storage, all three inks exhibited similar flow properties to recently prepared inks. The SWCNTs were found to aggregate in the SWCNT/PVDF-TrFE ink, causing occasional blocking of the nozzle during printing. Conversely, the Ti3C2Tx nanosheets showed minimal aggregation in the Ti3C2Tx/PVDF-TrFE ink due to exceptional electrostatic interactions between the Ti3C2Tx nanosheets and the PVDF-TrFE co-polymer (Supplementary

Ti3C2Tx nanosheets and the PVDF-TrFE co-polymer
The density of the PVDF-TrFE co-polymer melt was investigated as a function of the monomer units (alternatively the molecular weight) to validate the interatomic potential used for the simulations (Supplementary Fig. 4). The density was found to increase with increasing number of monomer units, reaching an asymptotic plateau corresponding to 1.42 g cm -3 . The value obtained using MD simulations was in excellent agreement with the value of 1.49 g cm -3 provided by the manufacturer of the PVDF-TrFE co-polymer (Solvay), validating the interatomic potential used in the MD simulations. The distribution of the local density of the PVDF-TrFE co-polymer film (70 chains) was investigated as a function of distance from the substrate, for a graphene substrate ( Supplementary Fig. 5, black line) and a Ti3C2Tx nanosheet substrate ( Supplementary Fig. 5, red line). The local density distribution was calculated from the mass within each separation distance interval, normalized to the volume within said interval and averaged over the duration of the simulation (1.8 ns timespan). The shaded regions correspond to the minimum and maximum values of the local density at each separation distance interval relative to the graphene or Ti3C2Tx nanosheet substrate. The layer adjacent to the substrate was found to adsorb to both the graphene and the Ti3C2Tx nanosheet substrates, exhibiting a local density of approximately 2.3 g cm -3 and 1.6 g cm -3 , respectively. The larger local density of the PVDF-TrFE co-polymer film adjacent to the graphene substrate indicates that the PVDF-TrFE copolymer chains are more packed than those adjacent to the Ti3C2Tx nanosheet substrate, as the latter possesses an increased surface roughness due to the OH termination, thus inducing steric effects in the PVDF-TrFE co-polymer chains. The local density of the layers further away from the substrate is practically identical in both the Ti3C2Tx/PVDF-TrFE and the graphene/PVDF-TrFE systems. Notably, the PVDF-TrFE co-polymer film adsorbs closer to the Ti3C2Tx nanosheet relative to graphene, as the first local density peak appears at a lower separation.
This decreased separation indicates a stronger attractive interaction between the Ti3C2Tx nanosheet and the PVDF-TrFE co-polymer in comparison to the graphene substrate. This enhanced adhesion phenomenon was investigated by adhesion strength studies. Here, a PVDF-TrFE co-polymer chain was placed in close proximity to either the Ti3C2Tx nanosheet or graphene sheet substrate. A force was applied to the PVDF-TrFE co-polymer perpendicular to the basal plane of the substrate to measure the desorption force. The force was increased from 0.00 pN to 6.95 pN with a step size of 0.695 pN, monitoring the position of the PVDF-TrFE co-polymer chain for the desorption from the substrate. It was found that the PVDF-TrFE co-polymer chain desorbed form the graphene substrate at approximately 2.78 pN, whereas the required desorption force increased on the Ti3C2Tx nanosheet to approximately 4.17 pN, indicating a greater adhesion strength at the interface between the Ti3C2Tx nanosheet and the PVDF-TrFE co-polymer.
The distribution of the H and F atoms in the PVDF-TrFE co-polymer film was further investigated as a function of the separation from the Ti3C2Tx nanosheet substrate to investigate whether preferential orientation of these dipolar atoms in the PVDF-TrFE co-polymer were giving rise to the polarization locking mechanism ( Supplementary Fig. 6). The datapoints represent average values for 14 PVDF-TrFE co-polymer chains and the shaded areas represent the minimum and maximum number of H and F atoms over the entire simulation. The H and F probability distributions were observed to be approximately equal at all separations from the Ti3C2Tx nanosheet substrate, indicating the PVDF-TrFE did not preferentially orient on the substrate. Small deviations at a low separation (up to 2 Å) were observed, whereby the H atoms were found closer to the Ti3C2Tx nanosheet substrate relative to the F atoms. This was attributed to the shortest non-covalent hydrogen bonds between the H atoms of the PVDF-TrFE copolymer and the hydroxyl terminations (Tx) of the Ti3C2Tx nanosheet substrate. To understand the phase distributions of the PVDF-TrFE co-polymer film adjacent to the Ti3C2Tx nanosheet substrate, the probability distributions of the dihedral angles were monitored for a 14-chain PVDF-TrFE co-polymer film ( Supplementary Fig. 7). The PVDF-TrFE copolymer consists of three commonly found phases, namely the α phase (non-polar), γ phase (semi-polar) and the β phase (highly polar). 16 These phases correspond to spatial conformation of the bonds, either trans (T) or gauche (G). The α phase is thermodynamically favored in fluoropolymers, due to its trans-gauche (TGTG'TGTG') conformation, which consists of 50% trans bonds and 50% gauche bonds. 17 Conversely, the β phase is an all-trans (TTTTTTTT) conformation, which spatially separates the H moieties on one C atom from the F atoms on the adjacent C atom, giving rise to a strong H-F dipole moment. 18 The γ phase is a stable intermediate state between the α phase and the β phase, as evidenced by the 75% trans and 25% gauche fraction (TTTGTTTG'), giving rise to dipole moments which result in a lower maximum polarization relative to the β phase. 17 Hence, the distribution of the dihedral angles and subsequently the phase fractions can provide insight on the changes in local electroactivity of the PVDF-TrFE co-polymer film adjacent to the Ti3C2Tx nanosheet substrate. prevalence of the β phase (26%), or a near-even distribution of the α phase (48%) and gamma phase (52%), or a combination of the two. Importantly, while the PVDF-TrFE generally crystallizes into the β phase due to the third F atom in the TrFE monomer, these values suggest a large presence of the α phase. 10 Indeed, at the local level, the experimental data observed the presence of the α and γ phases adjacent to the Ti3C2Tx nanosheet (Fig. 4d, e); however, the presented further in this document suggest the β phase as the primary conformation in the bulk of the Ti3C2Tx/PVDF-TrFE composites.
Similarly, the temporal evolution of the dihedral angles was repeated for a 70-chain PVDF-TrFE co-polymer film on the Ti3C2Tx nanosheet or graphene substrate ( Supplementary Fig. 8).
Similar to the 14-chain PVDF-TrFE co-polymer films ( Supplementary Fig. 7b), the larger films on a Ti3C2Tx nanosheet substrate exhibited a larger trans fraction (approximately 65%) relative to the gauche fraction (approximately 35%). Interestingly, when simulated adjacent to a graphene substrate, the same 70-chain PVDF-TrFE copolymer film exhibited a lower fraction of trans bonds (approximately 57%) and subsequently a higher fraction of gauche bonds (approximately 43%).

Rheological printing optimization of pristine PVDF-TrFE in acetone
To optimize the ink system (PVDF-TrFE in acetone) for SEA 3D printing, the rheological properties of the inks were first studied for PDVF-TrFE (35 wt%, 40 wt% and 45 wt%) loadings. PVDF-TrFE powder (75 mol% VDF, 25 mol% TrFE, Mw = 420 kDa) was slowly added into acetone and stirred at 23 °C until the powder completely dissolved, forming viscous inks ( Supplementary Fig. 2). The rheology of these inks was assessed using an MCR 702 rheometer (Anton Paar GmbH) in a cone-plate geometry, with a cone diameter of 25 mm, a cone angle of 2° and a gap at 102 μm (CP25-2, Anton Paar GmbH). The temperature in all measurements was held at 5°C.

Steady state rheology
Initial steady-state logarithmic shear rate ramps were used to compare the viscosity (η) of the PVDF-TrFE ink in acetone to that of the commonly reported solvent mixture of DMF and acetone (40:60 vol%), with polymer concentration at 35 wt% ( Supplementary Fig. 10a). 10,15 Both inks showed non-Newtonian (shear thinning) behavior, which is required for extrusion printing. 19 At a shear rate of 0.01 s -1 , which corresponded to the resting state (prior to and post printing), the η of the ink in the DMF:acetone solvent system was measured at 670 Pa s, drastically lower than that of the ink in acetone as the solvent, measured at 430,000 Pa s. The extreme increase in the viscosity at low shear represents a three order of magnitude increase in shape retention capability of the ink directly upon printing, further aided by the faster evaporation rate of acetone relative to DMF. Interestingly, the viscosity of the ink with acetone as the only solvent exhibited a lower η (7 Pa s) at high shear rate (1,000 s -1 , corresponding to conditions during printing) relative to the ink with DMF:acetone as the solvent system (12 Pa s). This signifies a lower pressure is required to extrude the same volume of ink, following the Pa s (2.7%) and 113 Pa s (0.5%) in the 45 wt% and 35 wt% inks, respectively. This suggests the higher concentration of PVDF-TrFE assists in stabilizing the entanglement in the polymer chains; however, the application of shear nonetheless reduces the entanglement between the polymer chains, correlating to a pseudo-1D material. 22 In region IV, the 35 wt% ink was found to drop in η to the lower value of that in region II, suggesting that the disentanglement is irreversible, whereas the η of the 45 wt% ink was found to be consistent throughout the region, with the same values as region II. Surprisingly, upon decrease in shear rate in region V, the η of the 35 wt% ink decreased to similar values as the high shear rate region II, unable to reliably recover to the values of region III, confirming the disentanglement effects and therefore proving unsuitable for a printing system where the printed ink must retain its shape.

Oscillatory rheology
Oscillatory rheology was employed at 5 °C to further probe the hypothesis of gel formation and determine the flow parameters in the PVDF-TrFE/acetone inks with PVDF-TrFE concentration at 35 wt%, 40 wt% and 45 wt% ( Supplementary Fig. 11). 19 Oscillatory frequency (ω) sweeps were performed (Supplementary Fig. 11a-c), which can give insight into the timedependent flow properties of the inks. 23 The tests were performed with fixed shear strain (γs) at 1%. All measured samples exhibited a similar trend in the storage (G') and loss (G'') moduli as a function of the ω, confirming the increased η ( Supplementary Fig. 10a) relative to PVDF-TrFE/(DMF:acetone) inks was due to enhanced swelling of the fluoropolymer, which has lower dependence on the fluoropolymer concentration. 24 Furthermore, minimal deviation in the slope of G' and G'' over the entire measured ω range strongly suggested the formation of a strongly bound gel, which was solid-like (G' > G'') for all measured frequencies. Similarly, the oscillatory γs sweeps ( Supplementary Fig. 11d-f) at low ω (1 Hz) exhibited similar characteristics between all three PVDF-TrFE concentrations in acetone. All three tested samples showed solid-like behavior (G' > G'') at low γs, followed by a liquid-like region (G' < G'') at high γs (>100%). 19 Interestingly, the flow point (γs at cross-over of G' and G'') was found to decrease with increasing concentration, which, while counterintuitive, suggests the strong binding between the PVDF-TrFE and acetone. 25 As the concentration increases, the number of polymer-solvent contact points decreases (increasing polymer-polymer binding points), therefore the gel becomes weakened and is able to flow with a lower γs. In translating this theory to extrusion printing, all of the three tested inks were suitable for printing; however, a lower flow point would decrease the required pressure input to extrude the sample, meaning the inks with higher PVDF-TrFE concentration are preferred for the printing. 20 Finally, oscillatory shear stress (σs) cycling was undertaken to probe the recovery parameters of G' and G'' within the inks and determine the optimal PVDF-TrFE concentration (35 wt%, 40 wt% or 45 wt%) in acetone for extrusion printing ( Supplementary Fig. 12). 19 Here, the σs was cycled at constant ω (1 Hz) between 1 Pa and 5 kPa, representing the induced σs at rest and during printing, respectively, held constant for at least 70 s ( Supplementary Fig. 12a). At PVDF-TrFE (35 wt%), the ink was unable to maintain σs at 6 kPa, whereas the inks containing 40 wt% and 45 wt% PVDF-TrFE exhibited consistent response to the input σs. The value for the tan(δ), or the ratio of G'' and G' was <1 (marked by grey horizontal line) for all samples at 1 Pa σs and increased to > 1 upon the application of 5 kPa σs for 70 s ( Supplementary Fig. 12b).
For the inks containing 40 wt% and 45 wt% PVDF-TrFE, the tan(δ) remained constant throughout the 5 kPa σs region and completely recovered for all concentrations after 70 s at 1 Pa σs. During the second cycle, the 35 wt% ink was found to flow with the lowest resistance, represented by a tan(δ) value of 20,000 (instrument limit), whereas the 40 wt% and 45 wt% inks retained similar values to the first cycle. Supplementary Fig. 12c shows the complex viscosity (η*) of the PVDF-TrFE inks. As expected, the starting η* was found to increase with increasing PVDF-TrFE concentration. During the first high σs cycle, the η* was found to decrease significantly for the 35 wt% ink as a function of time and unable to recover to initial values in the subsequent low stress cycle. Conversely, the 40 wt% and 45 wt% PVDF-TrFE inks exhibited full recovery after two high stress cycles.  Fig. 12d), the G' was observed to decrease rapidly as a function of time at σs = 5 kPa, with an average decrease over the timespan of greater than 1,000-fold. In the subsequent low-σs period, the slope of G' was higher than that of G"; however, the G' was unable to recover to the initial value of 9,600 Pa, reaching a maximum of 1,736 Pa. Throughout the second σs = 5 kPa cycle, the G' for the 35 wt% ink decreased significantly to below 1 mPa and subsequently exhibited a significantly lower slope during recovery. Conversely, the 40 wt% ( Supplementary Fig. S12e) and 45 wt% ( Supplementary Fig. 12f) PVDF-TrFE inks were stable under high-σs for at least one cycle and showed considerably higher G' recovery relative to the PVDF-TrFE (35 wt%) ink, from initial values of 11,800 Pa and 28,300 Pa, to final maxima of 7,000 Pa and 27,500 Pa, respectively. While the PVDF-TrFE (40 wt%) ink exhibited similar characteristics throughout the second σs = 5 kPa cycle ( Supplementary Fig.   12e), the PVDF-TrFE (45 wt%) ink was unable to consistently recover to initial values ( Supplementary Fig. 12f). Therefore, PVDF-TrFE (40 wt%) ink was selected for further experiments involving the incorporation of Ti3C2Tx nanosheets.

Mechanical properties
The tensile mechanical properties of the SEA extrusion printed Ti3C2Tx/PVDF-TrFE films were measured by a dynamic mechanical tester (ElectroForce 5500, TA Instruments). Samples, with a length of 27 mm and a width of 5 mm ( Supplementary Fig. 13a), were secured in grips by friction adhesive, with the distance between grips set at 5 mm ( Supplementary Fig. 13b).
The width (w) and thickness (t) of each sample is given in Supplementary Table 1. The films were extended parallel to the printing axis at a rate of 0.01 mm s -1 . Notably, the instrument displacement limit was approximately 11 mm, significantly below the breaking strain of the sample.  The tensile strain (γt) was calculated from the data obtained during tests following Equation S1 , where L is the displacement and L0 is the distance between grips at the beginning of the test (5 mm): The tensile stress (σt) was calculated from the data obtained during tests, using the crosssectional area of the sample (Acs), following Equation S2: Here, F is the measured force, t is the thickness of the sample, and w is the width of the sample (cut to approximately 5 mm).
Due to the low displacement limit of the instrument, the strain at break was approximated via empirical measurements ( Supplementary Fig. 13c), namely extending by hand. While these tests were merely representative, the samples were found to stretch to at least 65 mm prior to breaking, corresponding to 1,300% of the L0 (5 mm). Additionally, the final transparency in the extended regions was observed to be higher when the extension rate was slower ( Supplementary Fig. 13d).

Optical properties
The As expected, the transmittance was found to decrease as the Ti3C2Tx nanosheet concentration increased, from 94% for pristine PVDF-TrFE co-polymer to 20% for the Ti3C2Tx/PVDF-TrFE (0.50 wt%) film ( Supplementary Fig. 14a,c,d). Surprisingly, the addition of Ti3C2Tx nanosheets did not significantly increase the scattered light intensity and haze ( Supplementary Fig. 14b,c,

Raman analysis
The evolution of the Raman spectra when the Ti3C2Tx nanosheets are added to the PVDF-TrFE co-polymer reveals a clear suppression of out-of-plane vibrational modes occurring in Ti3C2Tx/PVDF-TrFE films. These modes occurring at 700 -720 cm -1 and 200 cm -1 correspond to the out-of-plane A1g vibrational modes for oxygen functional groups bound to the Ti3C2Tx lattice, whereas the peaks between 250 cm -1 and 700 cm -1 all correspond to in-plane Eg vibrational modes. 27 Notably, the A1g modes at 700 -720 cm -1 disappear almost completely, even in the Ti3C2Tx/PVDF-TrFE (0.50 wt%) films, with no difference between solvent-cast and extrusion printed films (Fig. 4a). In contrast, the intensity of the higher energy A1g mode at 200 cm -1 appears unchanged or even have an increased intensity relative to the main E2g modes (Fig. 4a). While this contrast in intensity change appears anomalous, it supports the data for well exfoliated flakes in literature. 27 Here, it should be noted that the Raman spectrum of the Ti3C2Tx nanosheets was attained by drop-casting Ti3C2Tx nanosheets in DMF on a silicon wafer, likely resulting in restacking and stronger A1g modes. The absence, or weak intensity, of these A1g modes in the Ti3C2Tx/PVDF-TrFE films therefore implies two key points, (1) that the PVDF-TrFE co-polymer is an excellent stabilizing agent for the Ti3C2Tx nanosheets as there is no evidence of restacking; and (2) there is a strong binding between the PVDF-TrFE copolymer and Ti3C2Tx nanosheets (and subsequent polymer densification) such that the A1g modes are even further weakened and shifted. 28 These results confirm the strong electrostatic binding as predicted by MD simulations (Fig. 2a).
Raman mapping of the surface of the Ti3C2Tx/PVDF-TrFE films showed a significantly variable response in the Iβ/Iγ ratio ( Supplementary Fig. 15). This variation was most noticeable for the SEA extrusion printed Ti3C2Tx/PVDF-TrFE (0.02 wt%) film and decreased with an increased Ti3C2Tx nanosheet loading up to 0.50 wt%, where the sample presented a homogenous ratio. This improvement in sample homogeneity at higher Ti3C2Tx nanosheet loadings is hypothesized to be due to the discrepancy in the state of the PVDF-TrFE co-polymer when it is bound to the Ti3C2Tx nanosheet basal plane. At higher Ti3C2Tx nanosheet loadings, we propose a high proportion of the PVDF-TrFE co-polymer is within the electrostatic sphere of influence (between 1 nm and 10 nm) of the Ti3C2Tx, thus presenting a homogenous Iβ/Iγ. 29 The data from these maps was averaged and used for describing the average sample spectra and Iβ/Iγ (Fig. 4a,b). The solvent cast Ti3C2Tx/PVDF-TrFE film shows a higher variation in Iβ/Iγ ( Supplementary Fig. 14f) which is attributed to the lack of homogenization of 2D materials by the extrusion printing process.

Attenuated total reflection Fourier transform infrared (ATR-FTIR) spectroscopy
ATR-FTIR spectroscopy was performed on the samples using an ALPHA II spectrometer (Bruker). Absorbance spectra were collected by taking an average of 128 individual scans at a resolution of 1 cm -1 , between 600 cm -1 and 4000 cm -1 .
ATR-FTIR spectroscopy was used to estimate the fraction of phases in the SEA extrusion printed Ti3C2Tx/PVDF-TrFE films, for Ti3C2Tx nanosheet concentrations at 0.00 wt%, 0.02 wt%, 0.10 wt%, 0.20 wt% and 0.50 wt% (Supplementary Fig. 16a). The peak commonly attributed to the α phase (766 cm -1 ) was not distinctly visible in all the measured spectra, suggesting the low fraction of the α phase in the bulk of the samples. 17 Notably, as was determined by Raman microscopy (Fig. 4d,e), the α phase was present in close proximity to the Ti3C2Tx nanosheet surface; however, the ATR-FTIR spectra suggested a low fraction of the α phase in the bulk of the Ti3C2Tx/PVDF-TrFE films. The peak at 840 cm -1 , indicative of the total electroactive phase (consisting of β and γ phases, denoted as β+γ), was present in all measured samples. 16 The separate peaks γ phase (1235 cm -1 ) and β phase (1290 cm -1 ) were both observed, confirming the presence of both electroactive phases (β+γ); however, the separate peaks could not be deconvoluted as the γ phase peak was present as a shoulder. The total electroactive phase fraction (Fea) was calculated using ATR-FTIR data by Equation

S5
: Here, Iea is the intensity of β+γ peak, Iα is the intensity of α peak, Kα and Kea are the absorption coefficients for the peaks at 766 cm -1 and 840 cm -1 , with values of 6.1 x 10 4 cm 2 mol -1 and 7.7 x 10 4 cm 2 mol -1 , respectively. 16 The Fea of the pristine PVDF-TrFE co-polymer film was 87.0% ( Supplementary Fig. 16b), significantly higher relative to the pristine PVDF-TrFE films SEA extrusion printed from a solvent mixture of DMF and acetone (40:60 vol%). 10 The highest Fea value was observed at 87.5% for the Ti3C2Tx/PVDF-TrFE (0.02 wt%) film, although this value exhibited little deviation from that of the pristine PVDF-TrFE film. Notably, at Ti3C2Tx/PVDF-TrFE (0.50 wt%), the Fea was found to decrease to 82.0%, consistent with the local α phase formation in the PVDF-TrFE on the surface of the Ti3C2Tx nanosheets (Fig. 4e).
The phase distribution in the SEA extrusion printed films was analyzed with XRD ( Supplementary Fig. 17). Two main peaks were visible in the spectra. The broad peak at 18. Spectra offset for clarity.
The primary fingerprint region ( Supplementary Fig. 17b) region was further deconvoluted to investigate the distribution of phase fractions ( Supplementary Fig. 18). The region required four peaks to ensure the correct fit. The strongest peak was assigned to the β phase (blue), the broad second peak was assigned to the γ phase (purple), the third peak was attributed to the α phase (yellow) and the final peak corresponded to the amorphous regions of the polymer (gray). The deconvoluted XRD spectra ( Supplementary Fig. S18a-e) were used to calculate the phase fractions ( Supplementary Fig. 18f) within the PVDF-TrFE co-polymer in the SEA extrusion printed Ti3C2Tx/PVDF-TrFE (0.00 wt%, 0.02 wt%, 0.10 wt%, 0.20 wt%, 0.50 wt%) films from the intensities for the respective peaks following Equations S6a-c: Here, Iα, Iβ and Iγ correspond to the intensities for the peaks found at 19.4°, 20.2° and 18.1°, respectively. The phase distributions were found to correlate well with the data obtained from Raman spectroscopy (Fig. 4b) and FTIR spectroscopy (Supplementary Fig. 16b). In particular, the Fea calculated from the FTIR spectra for the Ti3C2Tx/PVDF-TrFE films at Ti3C2Tx nanosheet concentrations below 0.50 wt% (87%) matched closely with the sum of Fβ and Fγ calculated from the deconvoluted XRD (between 85% and 90%). Furthermore, the Raman spectroscopy has suggested the primary phases in the bulk are the β and γ phases (Fig. 4a), with the β phase as the primary component, which is in close agreement with the XRD data.

Differential scanning calorimetry (DSC)
The crystallinity of the PVDF-TrFE co-polymer in the SEA extrusion printed Ti3C2Tx/PVDF- The DSC thermograms ( Supplementary Fig. 19) showed two endothermic peaks for all analyzed samples, centered at approximately 105 °C and 142 °C. 10 The peak at 105 °C corresponded to the ferroelectric to paraelectric transition (Curie temperature, Tc), whereby the samples exhibit piezoelectric properties below the Tc and lose polarization above the Tc. 33 The The PFM of the Ti3C2Tx/PVDF-TrFE films was carried out in lithography mode, 10 whereby a bias was applied to individual regions, monitoring the piezoelectric response through the converse piezoelectric effect (γ3 = d33E3, whereby γ3 is the out-of-plane strain, d33 is the piezoelectric coefficient and E3 is the out of plane electric field). 16,36,37 The lithography mode was chosen to obtain data below the poling field, where typical ferroelectric hysteresis loops cannot be formed, such that the poling state of the material would be minimally altered. 10 The applied voltage was between -20 V and +20 V, in increments of 2 V (Supplementary Fig. 21 using purpose-built Matlab code, which undertook pixel-by-pixel operations ( Supplementary   Fig. 24) to form the data obtained in Fig. 5b by multiplying the A signal (Supplementary Fig.   24b) by the cosine of the φ signal ( Supplementary Fig. 24a,c) and dividing by the Q factor of the cantilever (Qf) for each applied bias (Supplementary Fig. 24d). The Qf was measured in the tuning stage directly prior to the measurement (Supplementary Table 2). To calculate the effective d33, the Acos(φ)/Qf data was separated by the applied bias and averaged, obtaining a plot for Acos(φ)/Qf as a function of the applied bias for the samples (Supplementary Fig. 24d).  In the converse piezoelectric effect, the d33 is given as shown in Equation S8a: 37

Supplementary
Here, the superscript σ denotes constant stress. In order to minimize the stress on the sample, the cantilever is required to possess a sufficiently low spring constant, in this case approximately 0.3 N m -1 . It should be noted, however, that the cantilever will nonetheless apply stress to the sample, therefore restricting the expansion in the measured material and providing an underestimate to the calculated d33. The out-of-plane strain is given as γ3 = L/L0, whereby L is the magnitude of the expansion or contraction, and L0 is the material thickness. In PFM, L corresponds to the normalized amplitude, shown in Equation S8b: Therefore, the out-of-plane strain then takes on the form shown in Equation S8c : Moreover, the E3 is given as the V applied per unit distance. In the case of PFM, as the material expands upon applied V, the expansion should include the distance of expansion, shown in

Equation S8d
: Hence, substituting Equation 8b for L, the expression becomes as shown in Equation S8e : Indeed, the data shown in Supplementary Fig. 24d was found to be linear over the measured range, taking into account the deviation over the measured scan area. The data was fit with a linear trendline for each sample, whereby the slope of the trendline, accounting for the error in each sample, was the effective d33 value.
The accuracy of the nanoscale polarization measurements via PFM has been widely debated in recent literature, demonstrating the values can underrepresent or overrepresent the macroscale d33 both due to an empirical calculation methodology and the localized measurement approach. 36,38 Nonetheless, these PFM experiments are able to show trends of the d33 in samples with similar composition, as has been demonstrated here. Notably, the most accurate methodology is to utilize a single cantilever, as has been undertaken in these experiments. The utilization of multiple cantilevers has the potential to vary in the spring constant and therefore dampen the amplitude signal, subsequently changing the observed trends.

Macroscale displacement field measurement under compressive stress
The macroscale energy harvesting experiments were undertaken by the application of cyclic compression force and monitoring of the generated surface charge, configured to replicate the quasi-static Berlincourt method. 39 The cyclic compression was applied by a dynamic mechanical tester (ElectroForce 5500, TA Instruments) following a sinusoidal force pattern ( Supplementary Fig. 29a, b). : Here, D3 is the electric displacement field, d33 is the piezoelectric charge coefficient, σ3 is the applied stress, ε σ 33 is the dielectric permittivity at constant stress, E3 is the electric field and the subscripts correspond to the directionality, in this instance all parallel to the thickness axis.
Notably, in short circuit conditions where the input impedance of the load (in this instance the charge amplifier) is significantly lower than the output impedance of the PEG, the charge is transferred with no resistance. 41 In this instance, minimal voltage is generated and therefore E3 ≈ 0 V m -1 . Hence, the ε σ 33 can be ignored and the expression takes on the form shown in Equation S10 : The d33 can then be directly calculated from the input stress and the resultant electric displacement field, as shown in Equation S11 : Here, the superscript E denotes a constant electric field (∂E3 = 0). The stress is calculated from Equation S2 and is shown as a function of time in Supplementary Fig. 26c, d. The load cell of the mechanical tester, used to apply the stress, was cylindrical with a radius (r) of 12.5 mm.

27) was measured (Supplementary
for a circular segment: 42 C − ( − ℎ) U 2 ℎ − ℎ 1 C (S12)  Similarly, the electric displacement field is a value normalized to the active area (AD), shown in Equation S12 : In this instance, the AD corresponded to the area with sputter coated electrodes on both sides, which is under impact ( Supplementary Fig. 27, purple line, Supplementary Fig. 28, dashed purple line). The load cell was placed on the PEG such that the load cell did not make contact with the Cu foil adhesive ( Supplementary Fig. 27, Supplementary Fig. 28). The AD was measured as 2.25 cm 2 (2.25 x 10 -4 m 2 ), based on the active electrode dimensions at 15 mm length and 16 mm width, whereby the major part of the electrode was compressed ( Supplementary Fig. 27).
In order to ensure the generated charge arose only from the piezoelectric effect, the dependence of D3 on σ3 was investigated. In piezoelectric materials, as demonstrated in Equation S11, the slope must be linear, corresponding to the d33. In the instances where the slope is not linear, either the constant E requirement is not satisfied, or contributions from contact electrification and/or flexoelectricity are present. [43][44][45][46] A representative SEA extrusion printed Ti3C2Tx/PVDF-TrFE (0.50 wt%) PEG was tested for this dependence. The minimum force was set at 5 N, to minimize the effects from contact electrification. The resultant data is shown in Supplementary   Fig. 29. The slope was found to be linear, confirming the sole contribution of the direct piezoelectric effect to the measured surface charge.
Supplementary Fig. 29: The generated electric displacement field (∂D3) as a function of the input stress (∂σ3) for the SEA extrusion printed Ti3C2Tx/PVDF-TrFE (0.50 wt%) PEG. The solid line represents a linear fit to the data and the overlay represents a 95% confidence interval.
Upon the analysis of the generated surface charge (Q) in the SEA extrusion printed pristine PVDF-TrFE co-polymer PEG and the SEA extrusion printed Ti3C2Tx /PVDF-TrFE (0.50 wt%) PEG (Fig. 5d), it was evident that a significant enhancement in the energy harvesting was observed upon the incorporation of the Ti3C2Tx nanosheets into the PVDF-TrFE co-polymer.
In this instance, the D3 is related to the polarization (P3) following equation S13: 37 Here, we were able to neglect the previously reported enhancements in the dielectric permittivity (ε) in Ti3C2Tx/fluoropolymer composites 44 due to the absence of external electric fields (E ≈ 0 V m -1 ), therefore the enhancements in Q (and subsequently D3) were attributed directly to enhancements in the P3. The polarization is given as the sum of the individual dipole moment vectors (μ3) within a given volume (V), as shown in Equation S14, supporting the enhanced dipole moment alignment within the materials, as the dipole moment magnitude and the volume were constant.
The d33 of the Ti3C2Tx/PVDF-TrFE (0.50 wt%) PEG (at -52.0 pC N -1 ) was found to be higher than that of completely poled PVDF-TrFE in literature (at approximately -38 pC N -1 ), 16,48 suggesting that the electrical poling technique commonly utilized in literature does not completely polarize the pristine PVDF-TrFE co-polymer. 18,49 The presence of dielectric breakdown at a high poling electric field strength is hypothesized as the limiting factor in achieving completely polarized PVDF-TrFE for their utilization as PEGs. 50 Overcoming the limitation posed by the dielectric breakdown has profound opportunities in a multitude of fields where piezoelectric materials are used. The dipole locking mechanism from a nanomaterial template, as described in this study, has tremendous potential to unlock new applications for flexible piezoelectric materials, where the cost and energy input during manufacture is currently limiting commercial adoption.

Measurement of dielectric properties
The investigation of the dielectric properties of the SEA extrusion printed Ti3C2Tx/PVDF-TrFE was undertaken on the fabricated PEGs. An LCR meter (4284A, Hewlett Packard) was swept between 20 Hz and 1 MHz frequency at 0.5 V, with the probes connected directly to the wires of the PEG. Three individual PEG samples were measured at each Ti3C2Tx nanosheet concentration. The measured capacitance (C) was normalized via Equation S15 to the thickness (t) and the overlapping electrode area (AD) of the Ti3C2Tx/PVDF-TrFE film and the permittivity of free space (ε0 ≈ 8.854 x 10 -12 F m -1 ) in order to obtain the dielectric constant (εr): 50 The resultant dependence of the εr on the frequency is shown in Supplementary Fig. 30a for the pristine PVDF-TrFE and the 0.50 wt% Ti3C2Tx/PVDF-TrFE films.
Supplementary Fig. 30: The dielectric constant for the SEA extrusion printed Ti3C2Tx/PVDF-TrFE composites, shown as a function of a frequency and b the Ti3C2Tx concentration (100 Hz frequency). The error bars were obtained by testing three separate films at each Ti3C2Tx concentration and represent the mean ± SD.
The εr was found to increase slightly between the pristine PVDF-TrFE film and the Ti3C2Tx/PVDF-TrFE film, exhibiting a similar response to an increasing frequency. Notably, the error between the samples was observed to overlap at all frequencies, therefore it was concluded that the increase was not significant. This is in accordance with the data recently presented by Tu et al. 47 , demonstrating a 76% increase in the εr at 3.5 wt% of Ti3C2Tx nanosheets (nanosheet size between 1 μm and 2 μm) in poly(vinylidene fluoridetrifluoroethylene-chlorofluoroethylene) (PVDF-TrFE-CFE), a ter-polymer of the PVDF-TrFE co-polymer. Notably, the nanosheet dimensions in this study are significantly smaller in size (approximately 300 nm) and the maximum Ti3C2Tx nanosheet concentration is significantly lower (0.5 wt%), thus the lower increase in the εr of the composites presented in this manuscript is expected. Furthermore, the trend of the εr with increasing Ti3C2Tx nanosheet concentration ( Supplementary Fig. 30b) exhibits little correlation between the two parameters, and the εr of the 0.20 wt% Ti3C2Tx/PVDF-TrFE film is equal to that of the pristine PVDF-TrFE film. This data confirms that the increase in the d33 of the Ti3C2Tx/PVDF-TrFE composites does not arise from an increased εr, as discussion in the previous section.

Piezoelectric voltage coefficient and piezoelectric figure of merit
The measurement of the d33 by the macroscale (Berlincourt) method and the determination of the εr enables the subsequent calculation of the piezoelectric voltage coefficient (g33) and consequently the piezoelectric figure of merit (FOM). 16,39 The g33 is calculated following Equation S16, corresponding to the generated Q data presented in Fig. 5d.

"
(S16) The average g33 of the SEA extrusion printed pristine PVDF-TrFE PEG, measured from 60 individual compression cycles, was found to be 341 mV m N -1 . As expected, the partial polarization from the shear stress at the nozzle wall during the SEA extrusion printing process results in a g33 value lower than that of literature reports for electrically poled PVDF-TrFE (approximately 380 mV m N -1 ). 16 More importantly, the SEA extrusion printed 0.50 wt% The average FOM for the SEA extrusion printed pristine PVDF-TrFE PEG was calculated as 9.7 x 10 -12 Pa -1 . As expected, the lower d33 and g33 of the pristine PVDF-TrFE PEG relative to literature values for electrically poled PVDF-TrFE (14.4 x 10 -12 Pa -1 ) results in a 33% lower FOM. Conversely, the FOM of the 0.50 wt% Ti3C2Tx/PVDF-TrFE PEG at 20.9 x 10 -12 Pa -1 is significantly higher (45%) than electrically poled PVDF-TrFE, which is largely attributed to the strong electrostatic interactions between the PVDF-TrFE co-polymer and the Ti3C2Tx nanosheets. Further, the FOM of the 0.50 wt% Ti3C2Tx/PVDF-TrFE PEG is 115% higher than the SEA extrusion printed pristine PVDF-TrFE PEG. Unlike the pristine PVDF-TrFE PEG, which is restricted to the SEA extrusion printing process in order to exhibit partial polarization from induced shear stresses during deposition, the polarization-locked Ti3C2Tx/PVDF-TrFE solution possesses the flexibility for processing via conventional polymer film deposition techniques (e.g., solvent casting) while retaining the high piezoelectric properties (as demonstrated by PFM in Fig. 5c). Importantly, the FOM of the Ti3C2Tx/PVDF-TrFE PEG is 45% higher than electrically poled PVDF-TrFE, demonstrating a viable low energy deposition technique to produce effective flexible piezoelectric energy harvesting devices on a massproduced scale.
The comparison of the polarization-locked Ti3C2Tx/PVDF-TrFE PEGs presented in this work to piezoelectric materials reported in literature (Supplementary Table 4) demonstrates that the polarization-locked PEG possesses the highest g33 and FOM values reported to date, taking into account the analysis of Deutz et al. 39 and van den Ende et al. 51 , which investigate perovskite structure materials and polymer-perovskite composites, respectively.  Ref.