Introduction

The rapid development of energy harvesting technology is driving next-generation self-powered electronics to be more intelligent, highly integrated, and fully portable. Triboelectric nanogenerator (TENG), first demonstrated by Wang’s group in 2012, are a distinct class of energy harvesters that convert ambient mechanical energy into electricity through triboelectrification and electrostatic induction1. Alongside piezoelectric and pyroelectric systems, TENG represents one of the three major types of mechanical energy harvesting technologies for self-powered electronics2. Their emergence shows great potential to meet the growing demand for wearable microdevices, multifunctional sensors, and portable power solutions3. However, achieving industrial-scale adoption requires TENG manufacturing to meet demands for high-throughput precision and customizable production. Meanwhile, the material properties of TENG need to be optimized to meet the performance requirements of specialized application scenarios4.

Photocurable 3D printing has emerged as a promising solution for TENG fabrication, enabling the production of complex hierarchical architectures from multifunctional materials through layer-by-layer or point-by-point deposition5. This approach not only achieves cost-effective and efficient manufacturing but also allows precise control over electromechanical properties through tailored structural design6,7. For example, Yoon et al. developed a biomimetic-villus structure-based TENG for dust filtration, leveraging 3D printing to maximize surface area8. The device utilizes triboelectric charges generated between PTFE, ABS, and dust particles, creating electrostatic adhesion that captures ultrafine particulates (<2.5 μm) with ~40% efficiency over 75 min. Ideal TENG photocurable inks typically require balanced properties including good stability, low viscosity, optimal light penetration depth, minimal curing shrinkage, and rapid polymerization kinetics9. Furthermore, most available photocurable materials are petroleum-based resins, which pose sustainability challenges. The development of eco-friendly, 3D printable TENG inks remains an important research frontier. Palm oil (PO) is the world’s highest-yielding and most economical vegetable oil, make it a key driver for developing cost-competitive and sustainable polymers to replace petroleum-based alternatives10. In our prior work, we developed PO-derived methacrylate monomers exhibiting good flow characteristics with a low viscosity of 50 mPa·s and rapid UV-induced C = C conversion11. The developed sustainable photocurable inks offer potential for TENG applications, featuring both ecological advantages and efficient processing characteristics.

Another critical challenge in self-powered electronics is polymer flammability, which poses significant safety concerns for applications requiring thermal stability. Currently, most photocurable materials remain highly flammable, severely limiting their use in fire safety-critical scenarios12,13,14. The integration of flame retardancy into photocurable resins generally employs two strategies: (1) physical blending with additive flame retardants (e.g., ammonium polyphosphate, halogenated compounds)15,16, and (2) molecular design of reactive flame-retardant monomers containing phosphorus, nitrogen, or silicon17,18,19. Incorporating flame-retardant elements into the polymer molecular structure via covalent bonding can maintain or even enhance mechanical properties, representing the most common and optimal strategy for achieving durable flame retardancy. Phytic acid (PA), accounting for 65%‒85% of total plant phosphorus content, serves as the primary phosphorus reservoir of nature20. Its six phosphate groups have dual roles as an acid source and char-forming catalyst, enabling synergistic flame suppression through combined gas-phase radical quenching and condensed-phase char barrier formation21,22,23. However, the introduction of PA into 3D printable TENG inks faces challenges including poor compatibility with hydrophobic matrices, which impedes homogeneous dispersion, and an absence of polymerizable sites for integration into crosslinked networks24,25.

In this study, we developed a 3D printable TENG ink system comprising PO-derived methacrylate monomer (MPOEA) and a photocurable PA derivative (GPA). The MPOEA features dual polymerizable C = C bonds and flexible fatty acid chains11,26, while GPA was synthesized through glycidyl methacrylate (GMA) functionalization of PA via nucleophilic ring-opening reactions, enabling covalent integration into the photocrosslinked networks (Scheme 1)27. Employing LCD 3D-printing technology, a customized TENG was fabricated, exhibiting robust mechanical strength, superior shape fidelity, and compatibility. Its specially designed polymer substrate demonstrates outstanding self-extinguishing properties (UL94 V-0 rating) and thermal stability, delaying the degradation of material structure under extreme thermal environments, thereby extending the application potential of TENG in scenarios involving open flames or intense thermal radiation. This material, combining customizable design with active flame-retardant protection, provides new insights and technological solutions for developing next-generation self-powered electronics with high safety and reliability.

Scheme 1: Preparation and mechanism of flame-retardant bio-based polymers for customized TENGs.
scheme 1

a Synthesis of MPOEA from PO. b Synthesis of GPA from PA and GMA. c Formation of crosslinked network of MPOEA and GPA via 3D printing technique. d Construction of flame-retardant TENGs.

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Results and Discussion

Resin printability

Photocurable resins were formulated by uniformly blending GPA and MPOEA at various mass ratios, denoted as GPxMy, where x and y indicate the respective mass fractions of GPA and MPOEA. Real-time monitoring for UV curing of GP5M5, GP6M4, and GP7M3 resins was conducted using attenuated total reflectance-Fourier transform infrared spectroscopy (ATR-FTIR) technique (Supplementary Fig. 1). Monomer conversion was quantified by tracking the decreasing intensity of the C = C characteristic peak at 814 cm−1. The resins exhibit two distinct photopolymerization stages during the 30 s irradiation: an initial rapid conversion phase (0–15 s) followed by a plateau stage (> 30 s) (Fig. 1a). During the initial stage, the C = C conversions of GP5M5, GP6M4, and GP7M3 reach 83.08%, 57.95% and 40.10%, respectively. After 30 s, both GP5M5 and GP6M4 reach ~90% C = C conversion and plateau, whereas GP7M3 exhibits slower polymerization kinetics and remains in the active polymerization phase. The increased GPA content introduces greater steric hindrance effects, and its higher molecular weight reduces both the polymerization efficiency and limits final C = C conversion. In contrast, MPOEA serves as a reactive comonomer, preserving resin crosslinking efficiency. To further validate the photocuring kinetics and the formation of a robust solid network during printing, we performed photo-rheology measurements (Fig. 1b). For the representative GP6M4 resin, the storage modulus (G’) rapidly increased and plateaued within ~38 s of UV exposure, indicating the fast formation of a crosslinked network. This fast gelation is crucial for ensuring high shape fidelity and strong interlayer adhesion in LCD 3D printing. This finding aligns well with the C = C conversion kinetics observed via ATR-FTIR. The similarly rapid curing (~36 s) of both pure MPOEA and GPA confirms that their combination in GP6M4 preserves efficient photopolymerization. Viscosity, critical for resin processability and interlayer adhesion in 3D printing, is typically maintained below 1.3 Pa·s in commercial UV-curable formulations28. The viscosities of GP5M5, GP6M4, and GP7M3 resins are well below this threshold (Fig. 1c), satisfying the rheological requirements for 3D printing.

Fig. 1: LCD 3D printability of bio-based resins.
figure 1

a C = C conversion of GP5M5, GP6M4, and GP7M3 resins versus UV exposure time from ATR-FTIR spectra. b Evolution of storage modulus (G’) and loss modulus (G”) of GP6M4 resin under UV exposure. c Viscosities of GP5M5, GP6M4, and GP7M3. d Digital images of the formed resins. e LCD printing process. f Model for printing; g Digital images of the printed objects (I, II, and III). h SEM image of the cross-section of the printed objects.

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The resins prepared of MPOEA and GPA at different mass ratios form clear and uniform liquids, indicating their good miscibility (Fig. 1d). As a proof-of-concept, complex geometries were fabricated using the optimized resin formulations (GP6M4) on an LCD printer (Fig. 1e–h). All printed components exhibit good dimensional accuracy, morphological fidelity, and consistency in multi-directional forming. Scanning electron microscope (SEM) image of the fractured surface of the printed specimens (Fig. 1h) reveals uniform layer thickness and a densely bonded interlayer interface, indicating outstanding printability of the GP6M4 resin.

Physical and mechanical properties

As shown in Fig. 2a, the consistently high gel content (>95%) and low swelling ratio (~1.0) across all compositions indicate near-complete crosslinking and high crosslink density. The gel contents of the printed samples increase with lower GPA loading, while the swelling ratio slightly decreases accordingly. During network formation, the participation of GPA molecules in the crosslinked network immobilizes remaining C = C bonds through steric hindrance and restricted mobility, rendering further reactions kinetically unfavorable due to topological constraints29,30. The addition of MPOEA comonomers alleviates steric limitations, enhancing the resin’s crosslinking efficiency.

Fig. 2: Physical, mechanical, and thermal properties of printed thermosets.
figure 2

a Gel content and swelling ratio of thermosets. b Tensile stress-strain curves of thermosets. c Flexural stress-deflection curves of thermosets. d Tensile toughness of thermosets. e Image of GP6M4 sample during bending test. f Storage modulus of GP5M5, GP6M4, and GP7M3. g Tan δ curves of GP5M5, GP6M4, and GP7M3. h TGA and DTG curves of GP5M5, GP6M4, and GP7M3. i TGA and DTG curves of GP6M4 under N2 and O2 atmospheres.

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Consequently, MPOEA incorporation increases crosslink density, offering superior structural rigidity and strength while reducing elongation at break and flexural deflection (Fig. 2b, c and Supplementary Table 1). By tuning the ratio of GPA and MPOEA, the resulting polymers’ mechanical behavior can be manipulated from elastomeric to rigid-plastic. Notably, tensile toughness displays a non-monotonic trend, peaking at the GP8M2 composition (Fig. 2d). This optimal toughness arises from steric hindrance of GPA, which moderates crosslinking efficiency to create a looser network. This structural feature facilitates energy dissipation through segmental motion, balancing toughness and modulus. This remarkable flexibility stems from the entangled fatty acid chains in MPOEA that form an energy-dissipating phase within the crosslinked network11. Therefore, the sample exhibited favorable flexural resistance and was nearly indestructible throughout the bending test (Fig. 2e).

The temperature-dependent storage modulus (E′) and tan δ profiles of GP5M5, GP6M4, and GP7M3 polymers are shown in Fig. 2f, g. According to the classical rubber elasticity theory31,32, the crosslink density (ρc) of the polymer is proportional to the storage modulus in the rubbery plateau region (E′T) through the relationship: ρc = E′T/3RT, where T represents the absolute temperature at Tg + 40 °C, and R is the ideal gas constant (8.314 J/(mol·K)) (Supplementary Table 1). The stiffness of the polymer is reflected by the E′ in the glassy state. As shown in Fig. 2f, increasing the GPA/MPOEA mass ratio from 5:5 to 7:3 results in a significant reduction of E′ in the glassy state, accompanied by a corresponding decrease in ρc. Concurrently, the glass transition temperature (Tg) decreases from 107.43 °C to 81.03 °C with higher GPA content. Notably, the peak area of the tan δ curves reflects the mobility of molecular segments. The increased height and broadening of the tan δ peaks at higher GPA concentrations indicate enhanced segmental mobility resulting from a reduced crosslink density26,33. In such systems, a lower number of crosslinking points allows greater molecular chain motion, thereby facilitating the emergence of β-relaxation peaks34. This behavior is attributed to the unrestricted movement of side groups, branches, or local moieties within loosely crosslinked networks.

The thermal decomposition behavior of GP7M3, GP6M4, and GP5M5 polymer was investigated by thermogravimetric analysis (TGA) (Fig. 2h). All samples exhibit multiple weight loss peaks (Tmax) within the temperature range of 300‒500 °C. With the increasing GPA content, the gradual reduction in gel content compromises the initial thermal stability of GP7M3, resulting in a lower pyrolysis onset temperature compared to GP5M5 and GP6M4. However, GP7M3 demonstrates superior catalytic char-forming capability due to its higher phosphorus content from the enriched phosphate groups35. During combustion, GP7M3 generates larger amounts of polyphosphoric acid (HPO3)n, a strong Lewis acid that efficiently catalyzes polymer dehydration and carbonization. This catalytic process promotes the formation of a dense and continuous char layer, ultimately leading to the highest char residue for GP7M3 at 800 °C.

Figure 2i depicts the pyrolysis behavior of GP6M4 polymer under different gaseous atmospheres. Under a nitrogen atmosphere, the polymer primarily undergoes thermal decomposition, exhibiting two distinct pyrolysis peaks in the TGA curves. In contrast, under an oxygen atmosphere, the polymer experiences thermo-oxidative decomposition, characterized by three distinct pyrolysis peaks. The first peak (N2/O2) corresponds to the decomposition of phosphate groups and the onset of the carbonization process. The thermal degradation of phosphorus-containing compounds catalyzes dehydration and crosslinking of the resin matrix, promoting the formation of a dense carbonaceous char layer. Consequently, this stage is accompanied by the release of volatile species such as H2O and PO· radicals, resulting in initial mass loss. The second peak (N2/O2) arises from the pyrolysis of the crosslinked network. At elevated temperatures, the polymer backbone undergoes chain scission, generating macromolecular fragments and combustible gases such as CO and CH4. The char layer formed during the earlier stage acts as a physical barrier, slowing down the decomposition rate and leading to a distinct secondary mass loss stage. The third peak appears only under O2 atmosphere, which is associated with the oxidation of the char layer. The thermally stable carbonaceous residue formed in the earlier stages is further oxidized in the presence of oxygen, releasing gaseous oxides such as CO and CO2. This oxidation-induced mass loss is exclusive to oxidative conditions, accounting for the significantly higher residual char yield under nitrogen36.

Flame-retardant property

The flame resistance of the polymer was evaluated through limiting oxygen index (LOI) and vertical burning (UL-94) tests (Fig. 3a). GP5M5 exhibits an LOI value of 25.15%, classifying it as a combustible material. In the UL-94 test, GP5M5 shows momentary flame extinguishment after the first ignition but fails to self-extinguish after the second ignition (Fig. 3b), and thus does not qualify for any flammability rating. In contrast, GP6M4 demonstrates significantly enhanced flame retardancy, with an LOI of 30.44% and a V-0 rating in the UL-94 test. During twice ignitions, the flame self-extinguishes within 1 s, and the combustion residues remain intact without detachment (Fig. 3b). Further increasing the GPA loading raises the phosphorus content in the polymer, leading to a pronounced improvement in flame retardancy. The LOI values increase significantly from 25.15% (GP5M5) to 39.48% (GP10M0), confirming that phosphorus incorporation effectively enhances the flame resistance of the polymer (Fig. 3a)37,38,39,40,41,42.

Fig. 3: Flame retardancy of thermosets.
figure 3

a LOI values and UL-94 rating of the thermosets. b Video snapshots of GP5M5 and GP6M4 during UL-94 tests. c Comparison of the peak heat release rate (PHRR) and total heat release rate (THR) of GP6M4 with those in the literature. d Thermal stability tests for 3D-printed components.

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The flame-retardant property of GP6M4 polymer was evaluated by cone calorimetry (Supplementary Fig. 2). GP6M4 demonstrates outstanding flame retardancy, exhibiting a total heat release (THR) of 35.98 MJ/m² and a peak heat release rate (PHRR) of 265.99 kW/m². These values are markedly lower than those for other phosphorus-based flame-retardant systems reported in the literature (Fig. 3c). PHRR and THR are critical parameters for fire hazard assessment. A higher PHRR indicates more rapid heat release and a greater risk of rapid fire spread, while a higher THR corresponds to increased total energy output and longer fire duration. The low PHRR and THR values of GP6M4 reveal its dual advantages in minimizing both fire intensity and persistence. This outstanding performance originates from the multi-action flame-retardant mechanism of the phosphorus-based additives43. Phosphorus compounds catalyze the formation of a dense, cohesive char layer that insulates the underlying material and impedes heat transfer. Phosphorus-derived radicals scavenge high-energy species in the flame, interrupting combustion chain reactions.

To visually validate its protective efficacy under realistic thermal threats, we leveraged its outstanding 3D printing adaptability to precisely fabricate electronics prototypes featuring complex geometric architectures. Under simulated extreme conditions of localized thermal runaway ignition, these prototypes were subjected to continuous exposure to an alcohol lamp flame (approximately 800–1000 °C) for 1.5 min. An infrared thermal imager was employed to record the dynamic evolution of their surface temperature fields in real-time (Fig. 3d). The thermal imaging data clearly reveal a significant transient temperature decrease or plateau on the material surface during the initial combustion phase (10–30 s). This phenomenon is directly attributed to the efficient mechanism of the phosphorus-based flame retardant: Upon thermal decomposition, it undergoes strong endothermic reactions and releases non-combustible gases (such as PO· radicals), rapidly diluting the concentration of flammable gases and oxygen. Despite enduring prolonged high-temperature flame impingement for 1.5 min, with the final surface temperature reaching ~500 °C, all prototypes maintained highly intact carbon layer structures. Acting as a physical barrier between the electronics and the external fire source, the GP6M4 significantly impeded heat transfer towards the internal components while resisting melting, dripping, and secondary ignition, thereby effectively mitigating thermal runaway propagation.

Flame-retardant mechanisms

To analyze the gas-phase flame retardant mechanism of the polymer, the volatile pyrolysis products of GP6M4 were analyzed using TGA-FTIR (Fig. 4a and Supplementary Fig. 3). The FTIR spectrum of pyrolysis products at the maximum decomposition rate reveals characteristic absorption bands corresponding to H2O (3570 cm−1), hydrocarbons (2990 cm−1), CO2 (2360 cm−1), carbonyl compounds (1723 cm−1), and P–O–C/P = O groups (1131 cm−1)44. These species primarily originate from the pyrolysis of the fatty acid carbon chains in PO and the phosphate groups in PA. During thermal degradation, the P‒O‒C bonds begin to cleave at 200 °C (20 min), generating phosphorus-containing radicals (Fig. 4b). These radicals migrate into the flame zone and disrupt combustion chain reactions by scavenging H· and OH· radicals, which also leads to the release of H2O (Fig. 4c). Notably, the long fatty acid chains remain intact at this stage (Fig. 4d) due to the higher dissociation energy of C‒C bonds compared to P‒O‒C bonds45. Upon further heating to 350 °C (33.4 min), all functional groups exhibit pronounced pyrolysis peaks. Between 33.4 and 40.8 min, the synchronized intensity trends of the P‒O‒C/P = O bonds and H2O signals indicate simultaneous radical scavenging and polyphosphate-catalyzed dehydration and carbonization processes. At 40.8 min, the P‒O‒C/P = O peak intensities decline, while hydrocarbon and carbonyl signals reach their maximum, suggesting thermal degradation of the previously formed char layer. This degradation exposes the C‒H bonds in the fatty acid chains to pyrolysis. Full-range and magnified spectra (Fig. 4a–d) show the complete disappearance of phosphate ester, H2O, and carbon-related signals after 40.8 min, confirming the conclusion of dehydration, pyrolysis, and carbonization. The subsequent detection of CO2 (Fig. 4e) is attributed to the further thermal decomposition of exposed chain segments and residual chars.

Fig. 4: Gas-phase flame retardant analysis of Thermosets.
figure 4

a FTIR spectra of the thermal decomposition products of GP6M4 at different temperatures; Curves of the absorbance of the main pyrolysis products of GP6M4 varying with temperature and time: (b) P‒O‒C/P = O, (c) H2O, (d) hydrocarbons, and (e) CO2.

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Raman spectroscopy (Fig. 5a) reveals two characteristic vibration bands in the char residue of GP6M4, i.e., the D-band at 1366 cm−1, associated with disordered carbonaceous structures, and the G-band at 1600 cm−1, corresponding to graphitic domains. The intensity ratio ID/IG of 0.8 indicates a relatively high degree of graphitization. This structural ordering results from poly(metaphosphate)-catalyzed dehydration of the polymer, which promotes the formation of compact carbon layers that effectively impede oxygen diffusion, heat transfer, and gas permeation46. X-ray photoelectron spectroscopy (XPS) analysis further elucidates the chemical composition and elemental states of the char (Supplementary Fig. 4a). Deconvolution of the C1s spectrum (Fig. 5b) reveals three distinct peaks at 284.8 eV (C–H/C–C bonds), 286.5 eV (C–O–C/C–OH bonds), and 288.8 eV (C = O/O–C–O bonds)42. The O1s spectrum (Fig. 5c) exhibits peaks at 530.77 eV (P = O bonds from decomposed phosphate groups), 532.07 eV (P–O/C–O moieties), and 533.2 eV (C–OH, O–C–O, and COOR groups). P2p analysis (Supplementary Fig. 4b) identifies two primary peaks at 133.6 eV (P–O bonds) and 135.9 eV (P = O bonds)38. These structural features demonstrate that the residual char consists of aliphatic carbon derived from MPOEA pyrolysis and condensed phosphocarbonaceous phases formed from GPA decomposition. The synergistic interaction between these components imparts superior condensed-phase flame retardancy, primarily through enhanced thermal insulation and effective mass transfer barrier effects.

Fig. 5: Condensed-phase flame retardant analysis of thermosets.
figure 5

a Raman spectrum of GP6M4 char. b Deconvoluted C 1 s spectra from XPS. c Deconvoluted O1s spectra from XPS. d SEM image of GP6M4 surface. e Percentage of elements of GP6M4. f EDS image of elemental distribution of GP6M4 resin. g SEM image of GP6M4 residual chars. h Percentage of elements of GP6M4 chars. i EDS image of elemental distribution of GP6M4 chars.

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The elemental distribution and composition of GP6M4 before and after combustion were examined via SEM-EDS. The printed material exhibits a smooth surface with homogeneously distributed elements, and no evidence of phosphorus aggregation (Fig. 5d). Quantitative analysis (Fig. 5e, h) reveals a significant increase in the relative phosphorus content from 3.56% prior to combustion to 15.46% post-combustion. EDS elemental mapping (Fig. 5f, i) further confirms this phosphorus enrichment during combustion, transitioning from sparse to a dense spatial distribution of phosphorus. This transformation highlights the dynamic redistribution of phosphorus species during thermal exposure.

This achievement of the polymer with superior flame-retardant property is attributed to the synergistic effects of both gas-phase and condensed-phase mechanisms (Fig. 6). One hand, radicals such as PO· and HPO₂· are generated from the thermal degradation of phosphate groups, and effectively quench reactive H· and OH· radicals47, thereby interrupting combustion chain reactions. On the other hand, polyphosphoric acid catalyzes dehydration and carbonization of the polymer, promoting the formation of a dense, three-dimensional char structure. This structure features a continuous surface carbon layer coupled with an internal porous network (Fig. 5g). The multiscale barrier contributes to flame suppression through two primary pathways, i.e., physical isolation and chemical catalysis. The tortuous char microstructure impedes the transfer of heat and oxygen, while accelerated char formation sequesters combustible volatiles, limiting their release. These mechanisms effectively disrupt the fire triangle including heat, oxygen, and fuel, thereby suppressing the self-sustaining combustion cycle and delivering high-efficiency flame retardancy.

Fig. 6
figure 6

Proposed flame-retardant mechanisms.

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Application in printing electronics

Molecular-level diffusion between adjacent printed layers results in a continuous phase matrix structure, facilitating unhindered charge migration without aggregation. This ensures homogeneous charge distribution throughout the electrode layer, providing a solid foundation for high-efficiency single-electrode TENG. Therefore, fluorinated ethylene propylene (FEP) served as the external tribo-layer in contact with the printed GP6M4 (GPM) as the electrode layer to fabricate a single-electrode triboelectric generator (GPM-F) (Scheme 1d). Leveraging the synergistic effects of triboelectrification and electrostatic induction, the system effectively converts mechanical energy into electrical energy48. When FEP comes into contact with GPM, equal magnitudes of positive and negative charges are generated on their respective surfaces during contact. Upon separation, electrons flow from the ground to the GPM electrode layer through electrostatic induction, neutralizing the triboelectric charges. Complete charge equilibrium is reached at the tribo-layer/electrode interface once full separation occurs. Subsequent contact between FEP and GPM reverses the electrostatic induction, generating an opposing current as electrons flow back to the ground. Repetition of this cycle produces alternating current (AC)49,50. As demonstrated in Supplementary Movie 1, biomechanical energy generated by human joint motion is effectively converted into electrical energy through mechanical-to-electrical transduction, with the resulting current output generating sufficient intensity to illuminate the “TENG” alphabetic display.

The electrical output characteristics of the GPM-F TENG was evaluated (Fig. 7a–c). At a 1 Hz excitation frequency, the GPM-F achieves peak electrical outputs of 10.1 V open-circuit voltage (VOC), 0.05 μA short-circuit current (ISC), and 13.06 nC short-circuit transferred charge (SCq). Frequency-dependent analysis (0.5‒4 Hz, Fig. 7d) reveals a positive correlation between vibration frequency and VOC magnitude. This enhancement originates from accelerated FEP/GPM contact-separation cycles, which promote more efficient triboelectric charge accumulation while minimizing charge loss per cycle. As shown in Fig. 7e, the open-circuit voltage of GPM-F increases with expanding contact area. This correlation occurs because larger contact areas facilitate greater charge transfer between the two triboelectric layers. According to the principle of the triboelectric effect, increased contact area allows for more extensive charge accumulation, generating higher potential differences (i.e., elevated open-circuit voltage) upon separation. The ability of the GPM-F TENG to power practical electronics was further quantified by evaluating its charge storage capability. As shown in Fig. 7f, the device can effectively charge commercial capacitors ranging from 1 to 10 μF. The charging rate and final voltage exhibit a clear dependence on the capacitor value, with smaller capacitors being charged to a higher voltage more rapidly. This experiment directly demonstrates the substantial energy output of the TENG and its viability for managing and storing harvested mechanical energy. In practical terms, the energy harvesting capability of GPM-F meets the low voltage, current, and power density requirements of microelectronics integrated into firefighting suits51. Fig. 7g demonstrates a simulation where a firefighter wears customized GPM-F modules integrated at strategic locations (forearms, upper arms) to capture mechanical energy from operational movements (e.g., climbing, equipment handling). The converted electrical energy powers smoke-penetrating warning alarms and activates position-indicating lights on the suit. This dual functionality not only enhances safety by providing real-time alerts in smoke-filled environments but also enables rapid location identification of firefighters during emergency scenarios.

Fig. 7: Electrical characterization of TENGs for energy harvesting.
figure 7

a The open-circuit voltage of GPM-F at a frequency of 1 Hz. b Short-circuit transferred charge of GPM-F at 1 Hz. c Short-circuit current of GPM-F at 1 Hz. d Open-circuit voltage of GPM-F under different friction frequencies (0.5‒4 Hz). e The stable voltage of GPM-F operating continuously at a frequency of 4 Hz for 400 s. f The relationship between the open-circuit voltage of GPM-F and the contact area at a frequency of 2 Hz. g Demonstration of self-powered wearable TENGs and the induced currents generated at different parts.

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The practical power supply capability of the GPM-F TENG was further evaluated through a complete ‘harvest-store-use’ cycle. The device successfully charged a 4.7 μF capacitor to 3.5 V, with the stored energy subsequently used to illuminate an LED (Fig. 8a). Furthermore, capitalizing on the design freedom of LCD 3D printing, we fabricated customized TENG with complex geometries, capable of directly driving commercial LED in a parallel configuration with an output of ~5 V (Fig. 8b, c). The electrical output of the TENG was investigated as a function of external load resistance. The output power density exhibits a characteristic volcano-shaped curve, peaking at 130 μW/m² at a load resistance of 10⁸ Ω (Fig. 8d). This well-defined maximum confirms efficient impedance matching, a prerequisite for maximum power transfer. The underlying mechanism is elucidated in Fig. 8f. The voltage increases with load resistance to saturate at the open-circuit value, behaving as a constant voltage source, while the current density decreases monotonically, following Ohmic behavior. The power maximum thus logically occurs where their product is optimized, adhering to the maximum power transfer theorem. In addition, the TENG maintains a highly stable output voltage with negligible decay over 15,000 continuous cycles at 3 Hz (Fig. 8e).This exceptional durability stems from the robust crosslinked network of our material and the uniform, defect-free interlayer bonding achieved via 3D printing.

Fig. 8: Electrical output analysis and practical application demonstration of TENGs.
figure 8

a Real-time voltage profile during the charging process of a 4.7 μF capacitor and its subsequent application in lighting an LED. Demonstration of the GPM-F TENG powering LED arrays in b series and c parallel configurations, with the corresponding output voltages displayed. d Dependence of the output power density on the external load resistance. e The stable voltage of GPM-F operating continuously at a frequency of 3 Hz for 5000 s. f Relationships of the output voltage and current density with varying load resistance. g A multi-axis radar chart comparing the overall performance of this work against state-of-the-art TENG from literature across five key metrics: customizability, power output, sustainability, flame retardancy, and mechanical robustness.

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Crucially, the innovation of our work lies not in maximizing a single parameter, but in achieving a balanced and practical combination of key functionalities. A comprehensive performance comparison with leading TENG systems (Fig. 8g) highlights this core advantage. While some materials excel in pure electrical output, they often compromise on other vital aspects52,53,54,55,56,57. In contrast, our bio-based, flame-retardant platform is the only one that synergistically integrates high customizability, inherent safety, proven sustainability, robust mechanical integrity, and sufficient power generation. This unique profile makes it ideally suited for real-world applications, such as wearable technology and intelligent systems, where overall material virtues and operational reliability are as critical as electrical performance.

Conclusions

This study successfully developed a bio-based thermoset polymer that combines exceptional mechanical properties, flame-retardant property, and TENG functionality using LCD 3D printing technology. The resin was synthesized from renewable PA and PO, with most formulations achieving a quantified bio-based content exceeding 60%. The photosensitive resin exhibited excellent processability, with a rapid curing rate (<30 s) and low viscosity (<42 mPa·s), for 3D printing requirements. By optimizing the GPA/MPOEA mass ratio, the GP6M4 polymer demonstrated superior mechanical strength and flame resistance, including high tensile and flexural properties, a Tg of 101.62 °C, an LOI value of 30.1%, and UL-94 V0 rating. It also achieved excellent electrical performance with a stable AC output of 11.8 V under 4 Hz excitation for TENG application. The resulting TENG demonstrated a peak power density of 130 μW/m2, exceptional operational stability over 15,000 cycles, and the capability to power commercial LED and efficiently charge capacitors, completing a practical ‘harvest-store-use’ energy cycle. These remarkable properties were attributed to the high gel content (98.91%), excellent C = C conversion (93.25%), and a unique triple-synergistic flame-retardant mechanism, i.e., condensed-phase barrier (dense char layer blocking heat/oxygen), gas-phase dilution (H₂O/CO₂ release suppressing combustibles), and chemical quenching (PO·/HPO₂· radicals interrupting chain reactions). Notably, PA not only contributed to enhanced flame retardancy due to its high phosphorus content and acid catalysis, but also underscored the material’s eco-friendly advantages. This study demonstrates a viable strategy for the development of sustainable, flame-retardant, 3D-printable thermosetting materials for high-performance TENG.

Experimental Section

Materials

PO (melting point: 18 °C; acid value: 0.16 mg KOH/g) was procured from Shanghai Dingfen Chemical Technology (China). PA (70% aqueous solution), GMA (97%, containing 100 ppm MEHQ stabilizer), and methacrylic anhydride (MAA, 94%, containing 0.2% Topanol stabilizer) were purchased from Shanghai Macklin Biochemical Technology (China). Cetyltrimethylammonium bromide (CTAB, 99%), ethyl acetate (99.5%), sodium methoxide (97%), diethanolamine (99%), 4-dimethylaminopyridine (99%), and phenylbis(2,4,6-trimethylbenzoyl)phosphine oxide (BAPO, 97%) were obtained from Shanghai Aladdin Biochemical Technology (China). Anhydrous sodium sulfate (99%) was supplied by Tianjin Jinhui Taiya Chemical Reagent (China). Hydroquinone (99%) was sourced from Shanghai Yi En Chemical Technology (China).

Synthesis of MPOEA

MPOEA was synthesized via amidation and esterification reactions according to the reported protocol10. Diethanolamine (12.5 g) and sodium methoxide (0.16 g) were mixed in a three-neck flask and stirred in an oil bath at 80 °C for 30 min. Subsequently, PO (20 g) was added, and the temperature was elevated to 120 °C for a 4-hour reaction. The resulting product was purified via ethyl acetate extraction followed by rotary evaporation to obtain PO-based diethanolamide (POEA). For the subsequent esterification step, POEA (10 g), MAA (15 g), 4-dimethylaminopyridine (0.2 g), and hydroquinone (0.05 g) were homogenously mixed in a three-neck flask. The mixture was maintained at 65 °C in an oil bath for 5 h, followed by purification through saturated sodium bicarbonate extraction and rotary evaporation, ultimately yielding the target compound MPOEA.

Synthesis of GPA

GPA was synthesized via a nucleophilic ring-opening reaction between PA and GMA. PA (6.60 g) and GMA (12.79 g) were charged into a 500 mL three-necked flask equipped with a mechanical stirrer. Subsequently, the phase-transfer catalyst CTAB (0.19 g) and the inhibitor hydroquinone (0.019 g) were added. The flask was immersed in an oil bath at 85 °C, and the mixture was stirred continuously (250 rpm) until complete dissolution. The reaction system was then fitted with a reflux condenser and allowed to proceed for 2 h to obtain the crude product. After cooling to room temperature, the mixture was transferred to a separatory funnel and extracted with ethyl acetate. The organic phase was collected, dehydrated with anhydrous sodium sulfate under vigorous stirring, and then decanted. The supernatant was concentrated by rotary evaporation at 60 °C (0.08 MPa) to remove the solvent, yielding GPA as the final product in 76% yield. FTIR spectra (Supplementary Fig. 5a, b, KBr particle, cm−1): 942 w, 1013 m, 1045 m, 1118 w, 1175 vs, 1297 s, 1377 m, 1454 w, 1635 s, 1712 vs, 2957 m, 3300-3400 b. 1H NMR spectrum (Supplementary Fig. 5c): 1.9 ppm (‒CH3), 4.2 ppm (‒CH2‒), 5.6 ppm and 6.1 ppm (‒CH = CH2). 31P NMR spectrum (Supplementary Fig. 5d): 0.3 ppm (P‒OH), 0.8 ppm (P‒O‒C).

3D printing of bio-based polymers

Photocurable inks were prepared by homogeneously blending MPOEA and GPA at varying mass ratios (5:5, 4:6, 3:7, 2:8, 1:9, and 0:10), designated as GP5M5, GP6M4, GP7M3, GP8M2, GP9M1, and GP10M0, respectively. BAPO (2 wt% of total mixture mass) was incorporated as the photoinitiator. The printing process was conducted on an Anycubic Photon Mono 4 LCD photocuring printer with the following parameters: layer thickness of 0.05 mm, exposure time of 15 s per layer, and UV irradiation wavelength of 365 nm. After printing, the specimens were post-cured under a 36 W UV lamp (365 nm) for 8 min to ensure complete polymerization.

Characterization

The chemical structure of GPA was characterized by FTIR and nuclear magnetic resonance (NMR) spectroscopy. FTIR analysis was performed using a Nicolet iS20 FTIR Spectrometer (Thermo Fisher Scientific, USA) in transmittance mode, with 32 scans collected over the wavenumber range of 400–4000 cm−1. 1H and 31P NMR spectra were recorded on a Bruker Avance NEO 400 MHz NMR spectrometer (Germany) using deuterated chloroform (CDCl3, δ(1H) = 7.26 ppm) as the solvent. All measurements were performed at a magnetic field strength corresponding to 400/500 MHz operating frequency.

The viscosities and photo-rheology of the resins were characterized to evaluate their 3D printability. The viscosity of the resins was measured at room temperature using a rotational rheometer (UCHEN, China). Real-time photo-rheology was performed on a Haake Mars 40 rheometer (Thermo Fisher Scientific, Germany) to monitor the evolution of the storage modulus (G′) and loss modulus (G”) under UV irradiation, which reflects the gel point of network formation. All measurements were conducted at room temperature using a plate-cone geometry in oscillatory mode.

The gel content and swelling ratio of the printed samples were determined according to the ASTM D 2765-16 standard using a solvent extraction method. The sample was weighed (denoted as w1), wrapped in filter paper, and placed in a Soxhlet extractor connected to a 500 mL round-bottom flask. Subsequently, 250 mL of toluene was added to the flask as the solvent, and the extraction system was refluxed at 130 °C for 12 h. After extraction, the swollen sample was removed, blotted to remove surface liquid, and weighed (wg). The swollen sample was then dried to constant mass (w2) in a vacuum oven at 110 °C. The drying, cooling, and weighing steps were repeated until the difference between two consecutive measurements was less than 0.2 mg. The gel content and swelling ratio were calculated using the following formulas:

$${{\rm{Gel}}}\,{{\rm{c}}}{{\rm{ontent}}}=({w}_{1}/{w}_{2})\times 100 \%$$
$${{\rm{Swelling}}}\,{{\rm{r}}}{{\rm{atio}}}=1+[{\rho }_{1}({w}_{g}-{w}_{2})/{\rho }_{2}{w}_{2}]\times 100 \%$$

where ρ1 (g/cm3) was the density of the resin and ρ2 (g/cm3) was the density of the solvent.

Tensile properties were evaluated in accordance with ASTM D638-10 by using dumbbell-shaped specimens (total length: 75 mm, end width: 12.5 mm, narrow section width: 4 mm, gauge length: 25 mm, thickness: 3 mm). Flexural properties were assessed following ASTM D790-10 using rectangular specimens (80 mm × 10 mm × 3 mm). All mechanical tests were conducted on an Instron 3365 computer-controlled universal testing machine (USA) at a crosshead speed of 10 mm/min at room temperature. Seven replicates were tested for each group.

Dynamic mechanical properties of samples (30 mm × 10 mm × 2 mm) were characterized using a Q800 dynamic mechanical analyzer (DMA, TA Instruments, USA) in single cantilever mode. Tests were conducted with a frequency of 1 Hz at the temperature sweep from −50 °C to 150 °C, and a heating rate of 5 °C/min. The temperature-dependent profiles of storage modulus (E′), loss modulus (E′′), and damping parameter (tan δ) were recorded. The Tg was defined as the peak value of the tan δ curve.

Thermal stability was evaluated using a Netzsch STA 449 F5 simultaneous thermal analyzer (Germany). Approximately 15 mg of sample was heated from room temperature to 800 °C at a heating rate of 5 °C/min under a nitrogen atmosphere.

The flammability of samples were evaluated through LOI, UL-94, and cone calorimetry tests. LOI values were determined in accordance with ASTM D2863-17a using an FTT0077 Oxygen Index Tester (Fire Testing Technology, UK) with specimen dimensions of 80 mm × 10 mm × 4 mm under controlled oxygen/nitrogen atmospheres. UL-94 rating was assessed with a Vouch 5402 Horizontal/Vertical Burn Tester (Suzhou Yangyi Volcheck Testing Technology, China) employing specimens measuring 125 mm × 13 mm × 4 mm. Combustion behavior was quantified following ASTM E1354 via an FTT0007 Cone Calorimeter (Fire Testing Technology, UK) using 100 mm × 100 mm × 3 mm specimens, with testing conducted at a radiant heat flux of 35 kW/m2.

The morphological characteristics of the resin and combustion char residues were examined using a ZEISS Gemini SEM 300 scanning electron microscope (Zeiss, Germany) in secondary electron imaging mode. Prior to imaging, samples were sputter-coated with gold layer, with analyses conducted at an accelerating voltage of 5 kV. XPS analysis of the char residues was performed on an ESCALAB 250 XPS spectrometer (Thermo Scientific, USA) equipped with an Al Kα monochromatic X-ray source (hv = 1253.6 eV). The survey spectra were acquired at a pass energy of 100 eV with 8 accumulations, followed by high-resolution elemental scans at a pass energy of 30 eV and energy step size of 0.05 eV. The Raman spectra of the chars were acquired at room temperature using a Horiba LabRAM HR Evolution spectrometer (Horiba, Japan) in back scattering geometry, with excitation provided by a 514 nm argon-ion laser line.

The electrical output of the single-electrode triboelectric generator (GPM-F), fabricated by bonding a copper electrode to a 3D-printed GP₆M₄ friction layer (5 × 5 × 0.3 cm³), was systematically characterized. A low-noise electrometer (Keithley 6514) was employed to measure open-circuit voltage, short-circuit current, and transferred charge. The power generation capability was quantified via impedance matching analysis across load resistances. Operational stability was verified through 15,000 contact-separation cycles at 3 Hz, showing negligible performance decay. Practical energy storage and delivery functions were demonstrated by charging commercial capacitors (1–10 μF) and directly powering various LED configurations, with operational voltages recorded under parallel connection.