The as-synthesized p-type bismuth telluride ingot (home-made) was ball-milled into fine powders (< 45 μm) for subsequent slurry preparation. The composite slurry was primarily formulated with an active material, a home-made binder (synthesized according to literature^[21], denoted as MF-BPA; MF-BPA is a melamine-formaldehyde resin cross-linked with bisphenol A, which provides high thermal stability and stepwise carbonization behavior) and a conductive agent (multi-walled carbon nanotubes, denoted as MWCNTs). In brief, for the synthesis of MF-BPA, paraformaldehyde (12 equiv.) was dissolved in distilled water or ethanol under base catalysis (NaOH, pH 8.5-9.5) at 60 °C, followed by sequential addition of melamine (2 equiv.) at 80-85 °C for 30-45 min and bisphenol A (3 equiv.) at 80-90 °C for 1.5-2 h. The reaction was neutralized to pH 7.0-7.5 and collected without solvent removal, affording a viscous pale yellow to amber resin solution (50-70 wt% solids) functionalized with hydroxyl, methylene, and ether bridges.

A preliminary 1.5 g slurry was prepared with a mass ratio of Bi_0.5Sb_1.5Te₃:MF-BPA:MWCNTs = 7:2:1. A series of slurries were subsequently prepared by maintaining a constant Bi_0.5Sb_1.5Te₃:MF-BPA mass ratio of 7:2 while varying the MWCNTs doping concentrations (0, 5, 10, 15, and 20 wt%).

Absolute ethanol and ethylene glycol monobutyl ether were utilized as co-solvents, with the latter serving as a retarder to control the solvent evaporation rate. After introducing the BYK-2150 dispersant, the mixture underwent high-energy blending for 3 min to ensure uniform dispersion. The well-dispersed slurry was then coated onto a surface-treated PI substrate (RAYITEK Hi-Tech Film Company, Ltd., Shenzhen, 100 μm thickness) using a doctor blade with a 100 μm clearance, resulting in large Bi_0.5Sb_1.5Te₃/MWCNTs composite films. The thin film was punched into a round shape (12 mm diameter) and subsequently placed in a FJH instrument (model: HTS-7026D, Shenzhen Zhongke Jingyan Technology Co., Ltd., China) operating in pulse mode. Nine pulse operations were performed under a peak temperature of 450 °C, with each pulse width set at 200 ms.

The FJH peak temperature of 450 °C was selected via systematic optimization. Gradient screening experiments with fixed pulse width and pulse count were carried out at four peak temperatures: 350, 400, 450 and 500 °C. Temperatures below 400 °C cannot drive adequate atomic diffusion and binder carbonization, leading to marginal enhancement in electrical conductivity. Partial decomposition of the PI substrate occurred at 500 °C. Consequently, 450 °C achieved the best trade-off between microstructural reconstruction and substrate integrity. The 200 ms pulse width was optimized to maximize energy input into the film and reduce bulk heat accumulation on the substrate.

The crystallographic structures and phase compositions of the Bi_0.5Sb_1.5Te₃ powders, pristine films, and Bi_0.5Sb_1.5Te₃/MWCNTs composite films were characterized by X-ray diffraction (XRD) using an X’Pert Powder diffractometer (Malvern Panalytical, Netherlands) equipped with Cu Kα radiation (λ = 0.154 nm). Molecular vibrational information and the structural defect states of the incorporated MWCNTs were analyzed by Raman spectroscopy (Invia Reflex, Renishaw, UK, 532 nm laser excitation). The elemental distribution of Bi, Sb, Te, and C within the samples was characterized using energy-dispersive X-ray spectroscopy (EDS). Thermogravimetric analysis (TGA) was performed (TA Instruments Q500, USA) to evaluate the thermal decomposition behavior of the binder systems. The rheological properties of the slurries were measured (Anton Paar MCR 302, Austria) to assess their suitability for doctor-blade coating. Cross-sectional and surface morphologies were examined by scanning electron microscopy (SEM; Hitachi SU-8000, Japan).

Thermoelectric properties, including the S and σ, were simultaneously measured using an MRS-3 system (Wuhan JouleYacht Technology Co., Ltd., China) under low vacuum from 300 to 400 K. S was determined via the dynamic method, where linear regression of pulse-induced ΔV vs. ΔT curves was employed to eliminate parasitic thermopower and instrumental drift. Electrical conductivity was measured using the standard direct current (DC) four-probe technique to exclude contact and lead resistances. The PF was calculated as S²σ to evaluate the overall transport performance of the Bi_0.5Sb_1.5Te₃-based composite films. Room-temperature Hall effect measurements, including Hall coefficient (R_H), sheet/bulk resistivity (ρ), carrier concentration (p) and Hall mobility ( $$ \mu_{H}=\frac{R_{H}}{\rho} $$), were performed on an Accent HL5500 Hall System (Nanometrics, USA) with the Van der Pauw configuration. The four-contact Van der Pauw geometry eliminates strict requirements for sample shape.

Figure 1 presents a detailed schematic of the Bi_0.5Sb_1.5Te₃-based slurry and film preparation workflow. Bi_0.5Sb_1.5Te₃ powder serves as the thermoelectric matrix material, MF-BPA is introduced as a novel composite binder, and MWCNTs are added as a conductive network. These components are precisely proportioned in a mixed solvent of ethanol and 2-butoxyethanol. The mechanical shearing and collision action generated by ball milling achieves fine particle size reduction and uniform dispersion of the solid particles. The resulting homogeneous slurry system provides a solid material foundation for subsequent high-quality film casting or screen printing. The lower part of Figure 1 illustrates the FJH for the post-treatment of thin-film samples. The punched film pellets (12 mm diameter) are fixed in a custom metal fixture and placed under vacuum to prevent oxidation. The system employs a temperature pulse mode, performing 9 cyclic pulse heating operations at 450 °C with a fixed pulse width of 200 ms (heating duration for every single pulse, left bottom of Figure 1). Compared to conventional time-consuming tube furnace annealing, this millisecond-level ultrafast thermal shock process offers significant advantages in microstructural control. It can rapidly eliminate residual stress and partial organic volatiles within the film layer, while the extremely high heating and cooling rates effectively suppress excessive grain growth of Bi_0.5Sb_1.5Te₃ crystallites. This thermal processing mechanism, while improving material densification, may retain abundant nanoscale grain boundaries, which are expected to strongly scatter low-frequency phonons and further reduce lattice thermal conductivity. Notably, the peak temperature of the FJH process can be flexibly adjusted by tuning pulse voltage, pulse width, and the number of pulses. This tunability enables the process to be adapted to different thermoelectric material systems, binder formulations, and substrates with varying thermal tolerance, extending its applicability to a broader range of flexible electronic material systems.

The decomposition characteristics of the binder directly affect the purity and charge carrier transport network of the thermoelectric thin film after annealing. The TGA curves in Figure 2A compare the thermal decomposition behavior of three binder systems [sodium carboxymethyl cellulose (CMC), poly(vinylidene fluoride) (PVDF), and MF-BPA]. The results show that conventional binders CMC and PVDF undergo dramatic mass loss between approximately 100 °C and 200 °C, with almost no residual carbon retention. In contrast, the MF-BPA adopted in this study exhibits outstanding high-temperature stability and stepwise decomposition characteristics, retaining approximately 25% residual mass at 500 °C^[21]. This high thermal stability of MF-BPA means that during the rapid Joule heating process, it will not instantaneously gasify and cause film layer structural collapse or macroscopic pore formation. Instead, its moderate carbonization residue may form an in situ continuous conductive carbon skeleton between Bi_0.5Sb_1.5Te₃ particles, synergistically optimizing the interfacial charge transport channels of the composite material together with the carbon nanotubes^[22].

The rheological behavior of the slurry is a core parameter determining film quality and printing resolution. Figure 2B shows the viscosity versus shear rate curves (flow curves) for different binder systems. The MF-BPA system not only exhibits the highest apparent viscosity (stable at approximately 0.35 Pa·s at low shear rates) but also demonstrates typical and stable shear-thinning (non-Newtonian) fluid characteristics. In comparison, the CMC system has extremely low viscosity that drops sharply with increasing shear rate. The superior rheological properties demonstrate that MF-BPA is an ideal choice for preparing heavy metal compound slurries. Its high viscosity effectively suppresses gravitational sedimentation of high-density Bi_0.5Sb_1.5Te₃ powder during storage, ensuring excellent suspension stability, while the shear-thinning behavior ensures good flowability and printability of the slurry when external shear force is applied (such as tape casting or screen printing), thereby guaranteeing the surface flatness and internal uniformity of the thermoelectric film^[23].

Figure 3A and B shows the macroscopic morphology of the composite slurry before and after film formation on a flexible PI substrate via doctor blade coating. The wet film surface in Figure 3A is smooth and continuous, indicating that the slurry possesses excellent flowability and wettability. The solid thick film after drying in Figure 3B shows no obvious macroscopic cracks or curling delamination, confirming that the slurry system has excellent film-forming capability and structural adhesion. Figure 3C shows the circular film samples (12 mm diameter) punched after drying and subjected to FJH treatment. This process demonstrates the excellent macroscopic processability of the composite material - capable of meeting customized requirements for components of different geometric dimensions, laying a reliable foundation for subsequent thermoelectric property measurements.

Figure 3D presents the XRD patterns of composite materials with different MWCNTs mass fractions (5 to 20 wt%). The results show that the characteristic diffraction peaks of all composite samples are in high agreement with the standard card of Bi_0.5Sb_1.5Te₃ (JCPDS#49-1713, matrix phase), with no obvious impurity peaks observed. This indicates that the introduction of MWCNTs did not alter the intrinsic crystal structure of the thermoelectric matrix, and the two form a physical composite as the primary interaction, avoiding the occurrence of harmful interfacial side reactions, thereby preserving the superior charge carrier transport characteristics of the matrix material.

Figure 3E shows the Raman spectrum of the composite material, clearly displaying the D band (~1,350 cm^-1) and G band (~1,580 cm^-1) characteristic of carbon materials. The calculated intensity ratio I_D/I_G is approximately 1.02, indicating a certain number of structural defects on the MWCNT surfaces. The spectrum also presents a 2D band at ~2,700 cm^-1, which together with the D and G bands constitutes the fingerprint features of sp²-hybridized carbon materials^[24]. The coexistence of the strong G peak and the broad multi-layer 2D peak confirms the multi-walled nature of the incorporated carbon nanotubes, ruling out the possibility of pure amorphous carbon or single-layer graphene^[25].

Figure 4A and B shows surface-view SEM images of composite thermoelectric materials with 10 and 20 wt% MWCNTs additions, respectively. These two sets of images intuitively reveal the significant evolution of the internal microstructure and heterogeneous interface of the composite material as the second-phase filler concentration increases. At 10 wt% addition, MWCNTs present a uniformly dispersed state, with the tubular structure sparsely but continuously penetrating and bridging on the surface and grain boundaries of micron-scale Bi_0.5Sb_1.5Te₃ matrix particles. When the MWCNTs concentration increases to 20 wt%, the microstructure undergoes dramatic changes. The high-density carbon nanotubes form a highly entangled complex network like a “bird’s nest”, almost completely encapsulating and covering the Bi_0.5Sb_1.5Te₃ matrix particles. Figure 4C confirms that through the specific fabrication process (slurry coating combined with annealing), a flexible thermoelectric thick film with a specific thickness (~13 μm) and good uniformity over a larger macroscopic range was successfully obtained.

It is particularly noteworthy that ultrafast Joule heating appears to be the main mechanism driving pore reduction and dense particle fusion, because it creates localized, transient thermal fields that accelerate neck formation and densification^[26], while MWCNT loading mainly affects particle packing/dispersion and can even agglomerate at higher concentrations^[27]. Unlike traditional long-time isothermal annealing, millisecond-level pulsed Joule heating can generate intense non-equilibrium thermal gradients inside the film. This instantaneous local high-energy burst (450 °C, 200 ms pulse width per heating cycle) rapidly activates atomic diffusion on particle surfaces, which is hypothesized to initiate a local Ostwald ripening-like mass transport process^[28]. Driven by such atomic migration, partial micron-scale pores generated during tape casting may be filled, and obvious necking and interfacial welding can form between discrete Bi_0.5Sb_1.5Te₃ particles, enabling fast reconstruction of the film microstructure. Together with undecomposed MF-BPA and MWCNTs, a structurally robust and electrically continuous three-dimensional composite network is formed. More critically, since thermal energy is injected instantaneously in the form of extremely short pulses, the global heat cannot accumulate significantly toward the underlying substrate, thereby substantially reducing thermal exposure to the underlying flexible PI substrate and preserving its structural integrity while achieving densification of the thermoelectric active layer^[29].

The corresponding cross-scale microstructural morphology [Figure 4D-F], the distribution of MWCNTs in the Bi_0.5Sb_1.5Te₃ matrix undergoes a significant interface reconstruction process driven by sintering thermodynamics. Before sintering [Figure 4D], due to the spatial obstruction of solvents and undecomposed binders, MWCNTs are mainly loosely distributed around the matrix particles in a free or locally randomly entangled state, with weak interfacial bonding. With the instantaneous injection of Joule heating pulses [Figure 4E], the binder (MF-BPA) undergoes in situ carbonization and volume shrinkage, accompanied by particle rearrangement driven by local thermal stress, causing the nanotubes to gradually adhere to the matrix particle surfaces. After sintering is complete [Figure 4F], MWCNTs tightly and uniformly wrap the Bi_0.5Sb_1.5Te₃ particle surfaces, significantly reducing the original physical gaps, and successfully constructing a continuous and highly dense three-dimensional interconnected conductive network inside the composite system [as shown in the physical model in Figure 4G].

In the composite material system of MWCNTs and Bi_0.5Sb_1.5Te₃, MWCNTs, as a highly conductive 1D nanomaterial, significantly alter the thermoelectric transport properties of the matrix upon introduction. On one hand, MWCNTs interconnect within the Bi_0.5Sb_1.5Te₃ matrix to form an effective 3D conductive network, greatly reducing the overall electrical resistance of the material and significantly improving electrical conductivity with increasing doping amount^[30]. On the other hand, the heterogeneous interfaces formed between MWCNTs and Bi_0.5Sb_1.5Te₃ are proposed to have an energy filtering effect on charge carriers, potentially capable of selectively scattering low-energy charge carriers, thereby contributing to an enhancement of the S^[31].

Further quantitative analysis shows that the I_D/I_G = 1.02 of the MWCNTs indicates that there are relatively abundant structural defects on their tube wall surfaces. This is particularly critical for thermoelectric thin films prepared by wet-process techniques such as doctor-blade coating. On one hand, appropriate surface defects significantly improve the dispersibility and rheological properties of carbon nanotubes in the slurry, ensuring microstructural uniformity and macroscopic mechanical flexibility of the doctor-bladed film. On the other hand, the defect sites on MWCNTs surfaces can form tight interfacial coupling with the Bi_0.5Sb_1.5Te₃ matrix, constructing many heterogeneous interfaces. These interfaces not only act as strong phonon scattering centers to effectively reduce lattice thermal conductivity but also build continuous high-speed charge carrier transport channels inside the film by optimizing the interphase contact potential, thereby achieving synergistic regulation of composite film electrothermal transport.

As shown in Figure 5A, in terms of σ, all samples exhibit a monotonically decreasing trend with increasing temperature. At 300 K, as the MWCNTs doping amount increases from 0% to 20%, σ gradually increases from 1,000 S·m^-1 for the pristine Bi_0.5Sb_1.5Te₃ film (without MWCNTs) to 4,200 S·m^-1 (20 wt% MWCNTs). When the temperature rises to 400 K, σ generally decreases but still maintains the gradient order of 20 wt% MWCNTs > 15% > 10% > 5% > 0%, with the 20 wt% MWCNTs sample still reaching 3,500 S·m^-1. This decreasing behavior originates from the reduction in carrier mobility due to enhanced phonon scattering with increasing temperature in semiconductors. The significant increase in σ with MWCNTs doping amount is the basic requirement for electrical transport. However, the Hall mobility of the MWCNT-free pristine film is extremely low at 1.227 cm²·V^-1·s^-1; this low value is mainly limited by intense charge carrier scattering originating from abundant grain boundaries and micropores inside the tape-casted film.

The large leap in electrical conductivity of the composite film in Figure 5A is the inevitable result of the combined action of the 3D carbon nanotube percolation network and the Joule heating interface fusion effects^[28]. For untreated doctor-bladed films, the dense original grain boundaries and organic insulating residues inside constitute extremely high potential barriers, producing strong scattering of charge carriers. However, the local high-temperature melting effect induced by ultrafast pulsed Joule heating greatly improves this transport bottleneck^[28]. On one hand, the instantaneous thermal field effectively clears the insulating potential barriers on particle surfaces; on the other hand, atomic rearrangement and fusion at grain boundaries significantly reduce the interfacial scattering resistance for charge carriers crossing particle boundaries. Therefore, the Joule heating-driven matrix grain boundary optimization, complemented by the cross-grain boundary conductive “highways” built by MWCNTs, fundamentally removes the limitation of extremely low carrier mobility in solution-processed films^[28], enabling the 20 wt% doped composite material to exhibit an outstanding macroscopic electrical conductivity of up to 4,200 S·m^-1.

In the composite system, the significant improvement in overall electrical conductivity under conditions of severely limited intrinsic carrier mobility is fundamentally due to the highly conductive MWCNTs building a three-dimensional continuous percolation network inside the matrix. By virtue of their large aspect ratio and excellent intrinsic electrical conductivity, MWCNTs establish direct physical contact and low-resistance charge transport channels between adjacent matrix particles. When the MWCNTs addition reaches the percolation threshold, the dominant transport path of charge carriers changes, enabling direct transport along the carbon nanotube network, thereby effectively bypassing the high-resistance matrix grain boundary regions. The construction of this continuous microscopic conductive network fundamentally overcomes the limitation of film grain boundary potential barriers on carrier transport, ultimately achieving a large improvement in the macroscopic electrical conductivity of the composite material^[30,35,36].

The S in Figure 5B shows a monotonically increasing trend with increasing temperature. At 300 K, the S values of all doped samples are approximately 170-180 μV/K, reaching a maximum of 214 μV/K at 400 K (5 wt% MWCNTs). Appropriate MWCNTs doping (~5 wt%) may improve the S value by a proposed energy filtering effect of MWCNTs-Bi_0.5Sb_1.5Te₃ heterogeneous interfaces, increasing the proportion of high-energy charge carriers, thereby optimizing the S value^[37]. Excessive addition (15 and 20 wt%) leads to a slight decrease in S due to the bipolar effect caused by excessively high carrier concentration, partially offsetting this advantage^[38].

Figure 5C shows PF variation versus temperature for samples with different MWCNTs filling contents. Incorporation of MWCNTs significantly enhances the base material’s PF; at the same test temperature, sample PF overall increases with rising MWCNTs filling ratio. For 20% and 15% MWCNTs filled samples, PF first rises then falls as temperature increases, both peaking around 380 K. For 10% MWCNTs filled sample, PF stays stable at 300-340 K, then decreases when temperature exceeds 340 K; for 5% MWCNTs filled sample, PF drops slightly at 300-340 K, then continuously increases when temperature exceeds 340 K; the pure base material without MWCNTs has generally low PF, with no obvious variation in 300-400 K.

The optimal PF of 146 μW·m^-1·K^-2 (380 K, Bi_0.5Sb_1.5Te₃/20 wt% MWCNTs) in this study significantly outperforms many reported non-vacuum wet-process coated Bi_0.5Sb_1.5Te₃ films^[39]. For most unoptimized formulations without specialized high-energy post-treatment, their PFs are typically below 50 μW·m^-1·K^-2[40,41], while some optimized printed samples can exceed this value. This demonstrates the enormous progress of the 3D conductive network construction strategy in overcoming grain boundary potential barriers. Compared to common flexible organic-inorganic composites such as PEDOT:PSS/MWCNT systems (PF typically 50-130 μW·m^-1·K^-2), and highly ordered (Bi,Sb)₂Te₃/single-wall CNT composite films prepared by in situ crystal growth (PF up to ~1,680 μW·m^-1·K^-2)^[42], or traditional high-vacuum magnetron sputtered films and hot-pressed bulk alloys (PF up to 2,000-4,000 μW·m^-1·K^-2)^[43], the slurry wet-process technique represented by doctor-blade coating has become the core development direction in the flexible thermoelectric field due to its superior scalability and extremely low manufacturing cost^[44]. The introduction of MWCNTs may not only exert an energy filtering effect but also act as a 1D “nano-skeleton” endowing the film with excellent crack resistance and bending flexibility^[45]. Systematic quantitative evaluation of bendability and long-term mechanical reliability will be further performed in our follow-up work. This design concept of sacrificing partial intrinsic electrical properties in exchange for large-area conformal adhesion capability effectively reduces interfacial contact thermal resistance at the system level, providing an extremely valuable engineering reference for the continuous manufacturing of next-generation low-cost flexible thermoelectric components^[46].

Figure 5D provides a lateral comparison of this work (doctor-blade coating combined with ultrafast Joule heating pulse annealing) with other typical thick-film fabrication processes (such as brush coating, dispenser printing, and tape casting) in the literature in terms of thermoelectric performance parameters. Objectively speaking, since this work employs millisecond-level ultrafast thermal processing at 450 °C, the crystallinity and absolute electrical conductivity of the film are inevitably lower than those of screen-printed or tape-cast thick films that have undergone high-temperature sintering for several hours. At room temperature, the optimized film in this work delivers a competitive electrical conductivity of 4,200 S·m^-1, with a S of ~180 μV/K and showing comparable performance among solution-processed flexible thick films. This performance gap is mainly attributed to the “double-edged sword” effect of the ultrafast heating thermodynamic process. Although millisecond-level Joule heating pulses effectively suppress excessive grain growth and maintain high-density heterogeneous interfaces (favorable for maintaining high S and reducing thermal conductivity), they also lead to the inability to achieve high densification inside the film (as shown by the hierarchical porous structure in the cross-sectional SEM images). The large number of micron-scale pores and amorphous carbonization residues at interfaces constitute strong electron scattering sources, greatly limiting the further improvement of macroscopic carrier mobility.

Comprehensive performance comparison shows that, relying on the constructed 3D percolation network and the energy filtering effect induced by heterogeneous interfaces [Figure 6A and B], this work not only successfully maintains an excellent - even superior to some traditional chemical deposition and magnetron sputtering processes - S, but also demonstrates advantages in manufacturing cost, process compatibility, and large-area scalable manufacturing that physical vapor deposition techniques cannot match. However, limited by the inevitable grain boundary residues, amorphous phases, and carrier mobility losses caused by interfacial contact resistance in the composite system, the current σ and overall PF still have a certain order-of-magnitude gap compared to extreme processes such as PLD, as shown in Figure 6C^[47-49]. Future research may focus on exploring advanced post-processing techniques (such as flash sintering) to achieve deeper healing of grain boundary defects, while combining heterogeneous band engineering for fine matching and decoupling regulation of carrier concentration, to further narrow the performance gap with vacuum deposition techniques and ultimately realize the continuous manufacturing of truly high-performance, low-cost flexible thermoelectric components. Furthermore, future work should address the integration of electrode layers and other device components to realize fully functional thermoelectric generators, as the scalable fabrication of complete devices remains an important direction for practical applications^[50].