Effect of Crystalline Microstructure Evolution on Thermoelectric Performance of PEDOT : PSS Films

Although organic polymer thermoelectric (TE) materials have witnessed explosive advances in the recent decade, the molecular mechanism of crystallization engineering of TE performance, even for the most successful polymer of poly(3,4ethylenedioxythiophene) : poly(styrene sulfonate) (PEDOT : PSS), is still far from clear. Here, we deepen the understanding of the role of annealing-induced crystalline microstructure evolution on TE performance of the PEDOT : PSS film with thickness of 10 μm, which is usually more effective than thin ones in applications. Annealed at optimized temperature of 220°C, the film displays a power factor of 162.5 times of that of the pristine film before annealing. The enhanced TE performance is associated with the changes of crystallographic and morphologic microstructures, including increased crystallinity and crystal grain size, a domain morphology transformation from granular to crystalline nanofibril, and reduced insulating PSS in the skin layer. These variances facilitate the carrier transport by a transition from 3D to 1D hopping, reduce the activation energy, and improve the carrier mobility. The mechanism of crystallization engineering reported here can be conceptually extended to other TE polymers and guides the future rational design of preparation principles for organic and composite TE materials.


Introduction
Thermoelectric (TE) materials are irreplaceable in their unique capabilities of harvesting low-grade heat and wide applications in TE generators (TEGs), local cooling, eskin and wearable devices, etc. [1][2][3][4]. In the recent decade, organic and organic/inorganic composite TE materials have witnessed drastic advances due to their diverse advantages, including solution processability, lightweight, rich source in earth, super flexibility, and highly adjustable molecular structure or component [5][6][7][8]. Because of their low thermal conductivities (κ, typically between 0.1 and 0.4 W m -1 K -1 ), their TE performance is often evaluated by the power factor (PF = S 2 σ) rather than the figure of merit (ZT = S 2 σT/κ), where σ, S, and T stand for the electrical conductivity, the thermopower or Seebeck coefficient, and the absolute temperature, respectively. To date, a variety of judicious strategies of both material preparation and flexible device assembly have been developed, and the corresponding TE performances are boosted [9][10][11][12]. Never-theless, in-depth research of the underlying molecular mechanism is rather limited. For example, although the crystalline structures of polymers definitely govern the diverse properties and applications of materials, the relationship between crystalline microstructure evolution and TE performance is still not clear.
Regretfully, the enhanced TE performance is simply explained by macromolecular conformation transition, polymer chain expansion into the ordered alignment, and removal of insulating acid of PSS. Little is known for the dependence of PEDOT : PSS TE performance with its crystalline microstructure evolution. The available few reports concentrate on the effect of crystallization on electrical conductivity (not TE performance) for PEDOT thin films (with thickness ≤ 100 nm) made by oxidative chemical vapor deposition (OCVD) [27][28][29] or spin casting [18][19][20]. Furthermore, note that relatively thick films of micron scale are usually more effective than thin ones in actual applications, for instance, organic solar cells [30] and harvesting the temperature gradient [31][32][33][34]. Therefore, a clear understanding toward the role of crystalline microstructure evolution (including the crystal grain size, the crystallinity, the crystal parameter, and the orientation) on TE performance of PEDOT : PSS films, especially the micron-thick, is strongly desired.
Here, we shed light on the annealing-induced crystallization engineering of TE performance for PEDOT : PSS films with thickness of 10 μm, obtained by filtering the commercialized product of Clevios PH1000 dispersion. First, the TE performance strongly depends upon the film annealing treatments, and the highest TE performance occurs at the optimized annealing temperature of 220°C. Then, the carrier transport mechanism is discussed based on the variable range hopping (VRH) model and mobility. Subsequently, the crystalline morphology is directly observed by atomic force microscopy (AFM), while the changes of the crystallinity and the crystallite orientation are monitored by two-dimensional grazing-incidence wide-angle X-ray scattering (2D GIWAXS). In addition, the surface-rich crystallization phenomenon and decreased PSS fraction in the skin region are discovered by X-ray photoelectron spectroscopy (XPS). Finally, the underlying mechanism of the role of crystalline microstructure evolution on TE performance of PEDOT : PSS films is discussed.

Experimental Section
2.1. Materials. The PEDOT : PSS aqueous solution (Clevios PH1000) was purchased from Heraeus Deutschland GmbH. The concentration of PEDOT : PSS was 1.3 wt%, and the mass ratio of PSS to PEDOT was 2.5. Poly(vinylidene fluoride) (PVDF) membranes (pore size: 0.2 μm; Jinteng Technology Co. Ltd.) were used directly. Deionized water was used in all of the experiments.

Freestanding Film Preparation and Subsequent
Annealing. First, PEDOT : PSS aqueous solution (2 mL) was directly vacuum filtered onto a porous PVDF membrane and then dried under vacuum at 60°C for 12 h, affording flexible and freestanding PEDOT : PSS nanofilms. The thickness of the film is around 10 μm. Then, the obtained PEDOT : PSS films were thermally annealed for 12 h at a desired temperature under nitrogen atmosphere in a tube furnace (NBD-O1200-50IT, NuoBaDi Materials Science and Technology Company).

Characterizations.
The morphologies of the films were observed by an atomic force microscope (AFM, Bruker Dimension ICON). Thermogravimetric analysis (TGA) measurements were carried out with a TGA 55 instrument under nitrogen gas flow at a heating rate of 10°C min -1 . Powder X-ray diffraction (XRD) measurements were carried out with a multifunction X-ray diffractometer (SmartLab) at a scanning speed of 10°min -1 . In the grazing-incidence wideangle X-ray scattering (GIWAXS) measurements, a Xeuss 2.0 GIWAXS instrument (Genix3D Cu) with a Pilatus 3R 300K detector by Xenocs Beijing Scientific Instrument Technology Co., Ltd., was employed. The samples were tested in vacuum, and the exposure time of 300 s was applied in integration of q xy and q z . Photoemission spectrum (XPS) analysis was conducted by multipurpose X-ray photoemission spectroscopy (Escalab Xi+).

Measurements of Thermoelectric Performances.
The room-temperature electrical conductivities and the Seebeck coefficients were measured by a commercial instrument of a Film Thermoelectric Parameter Test System (MRS-3RT, Wuhan Joule Yacht Science & Technology Co., Ltd.) using a quasi-steady-state mode. During the measurements, the film sample is rectangular in shape, with a dimension of 2 cm in length, 0.5 cm in width, and 10 μm in thickness. And the variable-temperature conductivities and the Seebeck coefficients for the PEDOT : PSS film samples were measured using a Film Thermoelectric Parameter Test System (MRS-3, Wuhan Joule Yacht Science & Technology Co., Ltd.) in the temperature range between 100 K and 300 K. The carrier mobility (μ) was measured by the HMS-3000 Hall Measurement System. Figure 1 displays the schematic illustration of the annealing PEDOT : PSS film (see experimental details in Supporting Information (available here)) and two possible π-π stacking orientations of PEDOT. In brief, thick films of the pristine PEDOT : PSS were obtained by vacuum-filtrating commercial Clevios PH1000 solution on the porous poly(vinylidene fluoride) (PVDF) membrane ( Figure S1). The film with thickness of~10 μm exhibits a high flexibility ( Figure S2), whose electrical conductivity and Seebeck coefficient are 1-3 S cm -1 and 13-15 μV K -1 , respectively. Then, the films were thermally annealed at a desired temperature. As a consequence, crystalline films of PEDOT : PSS were achieved, which were further employed to investigate the relationship between crystalline microstructure evolution and TE performance. Two typical crystalline microstructures are presented, where the π-π stacking orientations of PEDOT chains are parallel (edge-on) or perpendicular (face-on) with respect to the film surface. In order to avoid possible hightemperature degradation of PEDOT : PSS, the annealing temperature range is decided by the thermogravimetric analysis (TGA) test. In Figure S3, the first mass loss section before 150°C is ascribed to the removal of water molecules by evaporation. The second stage occurs between 250°C and 350°C, possibly arising from the PSS degradation via the 2

Results and Discussion
Energy Material Advances rupture of the sulfonate groups from styrene, and the onset degradation temperature is 275°C [35,36]. When the temperature is higher than 350°C, an obvious rupture of the polymer backbone or skeleton breakage takes place, resulting in further mass loss [37,38]. Thus, PEDOT : PSS is deduced to reveal an excellent thermal stability below 250°C, which is consistent with previous studies [39,40]. In the present study, the temperature range between 200 and 250°C is chosen for annealing treatment. Figure 2(a) presents the room-temperature TE performances, including the electrical conductivity, the Seebeck coefficient, and the power factor, as a function of annealing temperature for PEDOT : PSS films after being annealed in nitrogen atmosphere for 12 h. Distinctly, the TE performances strongly depend on the annealing temperature. First, all of the annealed films exhibit significantly enhanced electrical conductivities, being approximately two orders of magnitude higher than that of the pristine (1-3 S cm -1 ). Increased annealing temperature leads to a rapid improvement until 220°C and a subsequent reduction of the electrical conductivity. At the annealing temperature of 220°C, the PEDOT : PSS film shows the maximum room-temperature electrical conductivity of 596 ± 14 S cm -1 . The other important parameter, the Seebeck coefficient, reveals a similar trend to the electrical conductivity, i.e., first increase gradually and then reduce with annealing temperature. The largest Seebeck coefficient is 23:3 ± 0:7 μV K -1 for the film annealed at 220°C, being much larger than that of the pristine film (13-15 μV K -1 ) as well. As a consequence, the power factor exhibits the same trend in a curve shape as the electrical conductivity or the Seebeck coefficient. At the optimized annealing temperature of 220°C, the PEDOT : PSS film displays the highest power factor, reaching 32:5 ± 2:5 μW m -1 K -2 . In sharp contrast, the pristine film has a low power factor of only around 0.2 μW m -1 K -2 . In other words, by thermal annealing, the power factor for the PEDOT : PSS thick film is boosted, with a maximum of 162.5 times of that of the pristine film before annealing. Indeed, the power factor of 32.5 μW m -1 K -2 is higher than that of the pristine PEDOT : PSS films reported in many previous investigations [39,41].
In order to elucidate the underlying molecular mechanism toward the enhancement of TE performance by thermal annealing, the temperature dependence test of the electrical conductivity is conducted in the range of 100-300 K (Figure 2(b)). For all of the PEDOT : PSS films, their electrical conductivities improve with the increase of annealing temperature, confirming the semiconducting transport nature of PEDOT : PSS at low temperature. Moreover, at a fixed temperature, the electrical conductivities follow the sequence of annealing at 220°C > at 240°C > the pristine before annealing. To gain more insights into the carrier transport mechanism, the VRH model is employed, which has been extensively applied for polymer TE materials with semiconducting behaviors [42,43]. The equation is presented in the following: where σ 0 represents the electrical conductivity at the infinite temperature (0 K), T 0 stands for the characteristic temperature, T is the current temperature, and α equals 1/ð1 + DÞ (D is the dimensionality) [44]. If α is 1/4, the carriers display a three-dimensional (3D) transport mechanism. And if α is 1/2, a 1D carrier transport can be deduced. As evidenced by Figures 2(c) and 2(d), ln σ shows an essential linear relationship with (c) T −1/4 for the pristine film and (d) T −1/2 for the film annealed at 220 or 240°C. Therefore, we conclude that the carrier transport route varies from 3D to 1D direction after annealing, while the annealing temperature has little effect on the carrier transport dimension. Furthermore, the detailed parameters deduced from the VRH model are compared in Table 1. The room-temperature electrical conductivities for the annealed film samples are 596 ± 14 S cm -1 (220°C) and 445 ± 2 S cm -1 (240°C), respectively, much higher than that for the pristine film (1-3 S cm -1 ). Indeed, the pristine film exhibits a metal behavior at temperatures higher than 550 K, as presented in Figure S4. T 0 is an important parameter, characteristic of the activation energies of carrier hopping barriers. Here, the pristine thick film reveals a T 0 of 119.3 K, which is much lower than that of thin films prepared by spin coating, for instance, 4210 K for the film with 27 nm in thickness [42]. After annealing, the T 0 decreased to 14.3 K (220°C) and 17.6 K (240°C). Meanwhile, the σ 0 of the pristine (30.0 S cm -1 ) increases to

Energy Material Advances
To further elucidate the annealing process, the carrier mobility as a function of annealing temperature is measured by the Hall effect (Figure 2(e)). The carrier mobility of the pristine PEDOT : PSS film is very low, being only 9:37 × 10 −3 cm 2 V -1 s -1 . After annealing, the carrier mobilities are improved remarkably by at least one order of magnitude. Very interestingly, the curve shape of the carrier mobility in Figure 2(e) is nearly identical to that of the TE performance shown in Figure 2(a). In other words, carrier mobility is critical for TE performance. The whole curve can be essentially divided into two parts. When the annealing temperature is below 220°C, the carrier mobility increases dramatically with annealing temperature, hence a simultaneous enhancement of the electrical conductivity and the Seebeck coefficient. For the film annealed at 220°C, the carrier mobility reaches its maximum of 1.27 cm 2 V -1 s -1 , being approximately 136 times of that of the pristine film. And the maxima reach 596 ± 14 S cm -1 and 23:3 ± 0:7 μV K -1 , respectively. In other words, the optimized crystallization temperature is 220°C. A further increase of annealing temperature results in a gradually reduced mobility. Considering reduction of TE performance at annealing temperature over 220°C, three samples, including the pristine film and the films annealed at 220 or 240°C, have been chosen for the following study of the role of crystalline microstructure evolution on PEDOT : PSS TE performance.
The morphology feature changes induced by thermal annealing are directly observed by AFM images, as shown in Figure 3. Distinctly, the surface of the pristine film displays small PEDOT spherical nanoparticles (white contrast) connected by PSS (Figure 3(a)). In contrast, Figures 3(b) and 3(c) reveal that after the thermal annealing at 220°C or 240°C, a nanofibril-shaped morphology with distinct contrast is evident, and the granular particles of PEDOT : PSS agglomerates disappear. Moreover, the surface morphology is quantitatively estimated by the root mean square (RMS) from the topographic images (Figures 3(d)-3(f)). The pristine PEDOT : PSS film is relatively uniform and smooth, exhibiting a RMS of 1.19 nm (Figure 3(d)), which is similar to the previous report [45,46]. After annealing at 220°C or 240°C, the RMS increase distinctly to 1.48 nm (Figure 3(e)) and 1.28 nm (Figure 3(f)), respectively. This demonstrates that the surfaces become rough for the annealed films due to the increase of the crystallite domain sizes (white contrast). Additionally, the brightness increases, implying the improved crystallinity. Therefore, the AFM images provide a solid proof that the annealing procedure induces a morphology transition from nanospherical granule to wellcrystalline nanofibril, and both crystallinity and crystallite domain dimensions increase. Furthermore, the film annealed at 220°C displays the largest crystallite domain size and the roughest surface.
Although the ordered molecular packing of conducting polymers has been widely regarded as the cause to the simultaneous increase of S and σ [47,48], the proof of detailed structural information is limited, especially for thick films. Here, large-area mapping of 2D GIWAXS, a powerful technique to monitor the molecular stacking and orientation, has been employed. In Figures 4(a)-4(c), scattering of stacking of alternating PEDOT and PSS (q~0:3 Å −1 ) for all thick films along the q z axis exhibited much higher (100) intensities than that of q xy , confirming a dominant preferred orientation with π-π stacking parallel to the film plane, i.e., edge-on stacking. This is different from previous research of remarkable face-on stacking in thin films [49]. In addition, the scattering intensities follow the sequence of (a) ≪ (c) < (b), implying the order of the film crystallinity: the pristine film (a) ≪ the film annealed at 240°C (c) < the film annealed at 220°C (b). This agrees with the above AFM result (Figure 3), demonstrating that annealing enables a drastic improved crystallinity and the film annealed at 220°C exhibits the highest crystallinity. In addition to the 2D spectra, onedimensional (1D) GIWAXS patterns are collected by line scans. In Figure 4(d), the peaks at 0.3 Å -1 and 1.8 Å -1 are assigned to the stacking of alternating PEDOT : PSS and the π-π stacking (010), respectively [21,32]. Due to the deviance in crystallinity, the film annealed at 220°C displays the highest scattering intensities for both peaks, while the patterns for the pristine film are much weaker than either of the annealed ones. In the out-of-plane patterns (q z ) (Figure 4(e)), besides an amorphous halo of PSS, a series of sharp peaks of (100) and (200) reflections are characteristic of lamellar stacking of PEDOT and PSS, while the peak at q z ≈ 1:8 Å −1 corresponds to the preferred edge-on π-π stacking. The deviance in peak intensities agrees with Figure 4(d). Additionally, peak shifts can be recognized, where (100) and (200) reflections for the samples annealed at 220 and 240°C increase from 0.23 and 0.52 Å -1 to 0.25 and 0.56 Å -1 , respectively. Hence, the lamellar distances of alternating PEDOT and PSS can be calculated to be 27.3 and 25.1 Å for the samples annealed at 220 and 240°C, respectively. In case of low crystallinity for the pristine film and the relatively broad peak for the face-on stacking with low content, their stacking distances are not taken into account. Deduced from the (010) reflection, the lamellar π-π stacking distances (real space distances) are calculated to be 3.55 and 3.50 Å for the samples annealed at 220 and 240°C (Table S1), respectively, which is similar to the reported data (~3.5 Å) for the solution-processed PEDOT : PSS film [21]. And the crystal domain size for the film annealed at 220°C (4.7 nm) is a little bit larger than that at 240°C (4.4 nm), contributing to the enhanced thermoelectric performance as well.
The annealing also induces a variance in the PEDOT : PSS composition ratio, as evidenced by XPS analyses shown  Figure 5. The full spectra of the pristine film and the annealed (at 220 and 240°C) PEDOT : PSS films ( Figure 5(a)) are identical in a curve shape and clearly demonstrate the existence of the C, S, and O elements. The corresponding peak fittings of core level spectra (S 2p) are presented in Figures 5(b)-5(d). The S 2p peak can be divided into two sections: the peaks between 166 and 168 eV can be assigned to the sulfur atoms from the sulfonate groups in PSS, while the peaks in the range of 162-164 eV are relative to those from the thiophene groups in PEDOT [47,48]. Both peaks involve the contributions from a spinsplit doublet, i.e., S 2p 3/2 and S 2p 1/2 . Furthermore, the molar ratio of PEDOT : PSS can be estimated from the peak areas of PEDOT and PSS. After annealing, the PEDOT : PSS ratio reduces obviously from 1 : 2.8 (the pristine) to 1 : 2.1 (annealed at 240°C) and even 1 : 1.7 (annealed at 220°C). Note that all of the characterization techniques including XPS, AFM, and GIWAXS reveal the information on the film surface layer. In other words, the thermal annealing enables the motion of PEDOT molecules from the inner region to the outer film surface. The reason may be assigned to macromolecular crystallization, and preferred orientation usually tends to take place in the skin layer rather than the core region [47,48,50].
For semicrystalline polymers, the crystalline microstructure evolution is vital for carrier transport and hopping, hence affecting the TE performance. The hopping mainly means the "coarse-grained" way from one crystallite site to another, instead of one-to-neighbor atomic sites in the amorphous phase [27]. It is an effective method to reduce the carrier transport barriers between crystallized regions. Taking all evidences together, the annealinginduced PEDOT : PSS crystalline structure evolution is schematically displayed in Figure 6, which shed light on the mechanism understanding of the enhanced TE performance. The pristine PEDOT : PSS thick film reveals a low crystallinity, a small crystallite size, a mixture of relatively more edgeon crystallites and less face-on one, and a high content of PSS in the skin layer. These features contribute to a very low mobility (9:37 × 10 −3 cm 2 V -1 s -1 ), a relatively high hopping barrier (T 0 = 119:3 K), and a 3D transport mechanism of carrier hopping. As a consequence, the pristine film exhibits a very low TE performance (~0.2 μW m -1 K -2 ). The following factors of the annealing-induced crystalline microstructure evolution may contribute greatly to the enhanced TE performance for the annealed PEDOT : PSS thick films: (1) The thermal annealing induces significant increases of both crystallinity and crystallite grain size, as well as a domain morphology transformation from nanospherical granule to well-crystalline nanofibril. Due to the formation of the ordered crystalline structure, the "coil-like" PEDOT chains are expanded or stretched to straighter or "rod-like." The increased crystallite size and the occurrence of domain nanofibril morphology reduce the distances between adjacent crystallite grains and benefit the overlap between their total wavefunctions [21,50]. Thus, the carrier hopping among adjacent crystallites are facilitated by the increase of effective carriers' (2) Since the interchain coupling has a remarkable impact on the electronic transport in PEDOT : PSS films, the dominant preferred orientation of π-π edge-on stacking dramatically facilitates the carrier hopping and transport along the film plane direction, which is superior to the face-on stacking in thin films (≤100 nm), thus favoring the carrier hopping and transport of intercoupling macromolecular chains (3) The high fraction of PEDOT crystal domains and reduced insulating PSS in the film skin layer resulted from skin-layer preferred crystallization, helped the carrier hopping, and improved the TE performance. This phenomenon has not been observed in thin films (≤100 nm) of PEDOT : PSS These main factors will significantly reduce the hopping activation energies for carrier transportation, resulting in a remarkable increase of carrier mobility and TE performances. In addition, the increase of annealing temperature enhances the macromolecular segment motions of both the formation of the ordered crystalline structure and also the relaxation of disordered-to-ordered transition. In the present study, the optimized annealing temperature is 220°C, and the corresponding annealed film exhibits the highest mobility and TE performances. A further increase of annealing temperature leads to a reduced crystallinity, a reduced crystallite size, and an increased PSS content in the film skin region, hence reducing the carrier mobility and the TE performances.

Conclusions
In summary, we obtain the insights into a clear understanding of the mechanism of the role of annealing-  Figure 4: 2D GIWAXS patterns of (a) the pristine PEDOT : PSS film and the film being annealed at (b) 220 or (c) 240°C. Linearly offset 1D GIWAXS patterns of the pristine and annealed PEDOT : PSS films with (d) in-plane (q xy ) or (e) out-of-plane (q z ) scattering geometry. 7 Energy Material Advances induced crystalline microstructure evolution on carrier transport and TE performance for PEDOT : PSS films. The annealing yields obvious changes of crystallographic and morphologic structures in the films with thickness of 10 μm, including improved crystallinity, enlarged crystal domain dimension, and domain morphology transformation from granule to crystalline nanofibril. In addition, π-π edgeon stacking is the primary orientation, and the fraction of insulating PSS is reduced in the film skin layer. These microstructural variances lead to a mechanism transition from 3D to 1D for carrier transport, decreased activation energy, and improved mobility, hence facilitating the carrier hopping and transport. As a consequence, the TE performance enhances greatly upon annealing. Annealed at the optimized temperature of 220°C, the film displays a significantly enhanced power factor, being 162.5 times of that of the pristine film before annealing. Based on the mechanism proposed herein, it can be envisioned that by judicious crystalline microstructure evolution, the TE performance  Energy Material Advances for other polymers and polymer-based composites can be conveniently adjusted and remarkably enhanced. This work sheds light on the understanding of material chemistry and provides a guideline to fabricate high-performance organic and composite TE materials.

Data Availability
All data needed to evaluate the conclusions in the paper are presented in the paper and/or the Supplementary Materials. Additional data related to this paper may be requested from the authors (G.C.).