Generating Nanoscopic Patterns in Conductivity within a Poly(3

Publication Date (Web): September 21, 2018. Copyright © 2018 American Chemical Society. *E-mail [email protected] (B.Z.)., *E-mail [email protected]...
2 downloads 0 Views 8MB Size
Download PDF
Article Cite This: Macromolecules XXXX, XXX, XXX−XXX

pubs.acs.org/Macromolecules

Generating Nanoscopic Patterns in Conductivity within a Poly(3hexylthiophene) Crystal via Bias-Controlled Scanning Probe Nanolithography Binghua Wang,† Bin Zhang,*,† Changyu Shen,† Jingbo Chen,*,† and Günter Reiter‡ †

School of Materials Science & Engineering, Zhengzhou University, Zhengzhou 450002, People’s Republic of China Institute of Physics, University of Freiburg, 79104 Freiburg, Germany

Macromolecules Downloaded from pubs.acs.org by KAOHSIUNG MEDICAL UNIV on 09/21/18. For personal use only.

‡

S Supporting Information *

ABSTRACT: Patterning the electronic properties of semiconducting polymers on surfaces at the nanometer scale has a significant impact on their application in nanophotonics and nanoelectronics. Surprisingly, little attention has been paid to methods employing the phase transition from crystal to melt and thermally activated oxidation of conjugated polymer crystals, both accompanied by dramatic changes in electrical conductivity. Here, we propose a novel concept based on bias-controlled scanning probe lithography (SPL) which permits localized oxidation of poly(3-hexylthiophene) (P3HT) crystals and enables the on-demand formation of nanoscale, nonconductive structures without visible changes in topography. The approach relies on Joule heating induced local melting and oxidation of P3HT, resulting in nanoscopic nonconductive structures. Besides locally switching off electrical conductivity, this approach even allows to tune the electrical conductivity by adjusting the applied bias. Thus, we can generate patterns with nanometer spatial resolution and controllable conductivity within crystals of conjugated polymer through bias-controlled physical and chemical changes of the polymer.

1. INTRODUCTION For many years, conjugated polymers have garnered substantial research interest because of their mechanical flexibility, facile solution-based manufacturing processes, and their impressive optoelectronic properties.1−3 Among this class of polymers, due to its optimal balance between solution processability and optoelectronic properties, poly(3-hexylthiophene) (P3HT) is widely applied in various organic electronic devices such as field effect transistors, light-emitting diodes, and solar cells.4−7 However, when P3HT is exposed to oxygen, light, and moisture combined with high temperatures, its crystalline order and chemical structures are easily destructed, causing severe changes in electrical conductivity.8,9 In general, thermally activated oxidation is an undesirable effect in most cases as it leads to a deterioration of electrical properties of P3HT.10,11 However, it has not yet been explored if thermally activated oxidation combined with the crystal−melt transition (reorganization and degradation) can be exploited to engineer spatially resolved patterns within P3HT crystals. Patterning the properties of semiconducting polymers on surfaces at the nanometer scale is crucial for the development of fields such as nanoelectronics, nanophotonics, and organic electronics.12−16 In this context, scanning probe lithographic (SPL) methods have attracted a great deal of attention due to their capability of nanometer scale control and the corresponding potential applications.17,18 With SPL, nanopatterns are generated by high field gradients (mechanical, © XXXX American Chemical Society

electric, thermal, etc.) between the scanning probe tip and the surface of a sample.19,20 Farina et al.21 have reported a nontopographic patterning method which was achieved by the mechanical friction between the AFM tip and the sample surface. They proposed that the mechanical interaction produces an increase of local molecular disorder, inducing a localized reduction of the semiconductor conductivity. Hot probe tips have been employed to fabricate nanopatterns in a polymer film, e.g., in poly(methyl methacrylate) or poly(styrene-r-benzocyclobutene) films.22−25 However, such thermally activated SPL is difficult for two reasons: (1) a relatively large power is consumed because only a small percentage (95%, purity ∼99.995%) was purchased from Luoyang Microlight Material Technology Co., Ltd., and used without further purification.32 P3HT was dissolved in chlorobenzene by heating the solution for 6 h to 70 °C in darkness. Afterward, the solution was allowed to cool to room temperature and was subsequently filtered through a 0.22 μm polytetrafluoroethylene syringe filter. The substrates used were highly doped p-type (100) silicon wafers (resistivity ∼0.001 Ω·cm), cleaned with UV light for 1 h.33 Samples for optical microscopy (OM), atomic force microscopy (AFM), and grazing incidence wide-angle X-ray scattering (GIWAXS) experiments were prepared by spin-coating solutions onto silicon wafers with a KW-4A spin-coater (Institute of Microelectronics, Chinese Academy of Sciences, China).34 Spin-coating was performed for 30 s at a speed of 3000 rpm for solution with a concentration of 1.0 wt %.35,36 The thickness of the resulting P3HT films was about 30 nm (measured by an Alpha-SE ellipsometer, J.A.Woollam, USA). The solutions and films were prepared in a dry nitrogen atmosphere. Crystallization Temperature Protocol and Washing Procedure. The temperature protocol adopted for melt crystallization experiments is schematized in Figure S1 and consisted of the following steps. Specimens for isothermal crystallization were first heated from room temperature to 240 °C at a rate of 30 °C/min and held there for 1 min. Then, these samples were quenched to the corresponding isothermal crystallization temperature Tc at a rate of 40 °C/min and kept there for different crystallization time (tc) under a nitrogen atmosphere. Here, it should be pointed out that in melt crystallization experiments samples were stored in a dark chamber and under nitrogen protection. To facilitate the access of P3HT nanofibers by SPL, it was necessary to remove amorphous portions B

DOI: 10.1021/acs.macromol.8b01465 Macromolecules XXXX, XXX, XXX−XXX

Article

Macromolecules taken in the attenuated total reflection (ATR) mode on a Nicolet iS50 FTIR spectrometer (Thermo Fisher Scientific, USA), equipped with an ATR accessory. Samples for FTIR measurements were prepared by drop-casting concentrated P3HT solutions onto a bare silicon substrate (double-side polishing, resistivity >2000 Ω·cm) followed by drying in the dark under vacuum. The samples (of thickness about 200 nm) were placed on the sample holder after crystallization and annealed in dry nitrogen. All spectra were recorded in the range from 450 to 4000 cm−1 by accumulating 32 scans at a resolution of 4 cm−1. A JEM-2100 transmission electron microscope (TEM) (JEOL, Japan) operating at 200 kV was used to measure bright-field TEM images and SAED patterns. The accumulated electron dose on the specimen for taking one micrograph was set at an exposure of 200 e− nm−2. The very low dose is to prevent electron damage from the image. Samples for TEM were prepared by spin-coating P3HT solution on carboncoated mica.39 Then, after crystallization or annealing, these P3HT films were carefully floated off onto a pure water surface. Subsequently, samples were transferred to a copper grid supported by a piece of filter paper to remove excess water. The sample was then shadowed with a few nanometers of amorphous carbon for electron diffraction observations.

heating. Such a situation is similar to the research by Lyuksyutov and co-workers.40 They proposed that current flow through the polymer will induce the Joule heating under the AFM tip, which increases the temperature within the polymer film and creates a localized region of molten polymer. In our case, at sufficiently high bias voltage (see Figure 1f) the local temperature (TJH) induced by Joule heating may be above the melting point (Tm) of P3HT, and the crystalline structure may melt, accompanied by degradation induced by thermally activated oxidation. Thus, the required energy to rise the local temperature above Tm is a key parameter for controlling the above-described physical and chemical changes. In the following, we present a detailed analysis the electrical power required to melt a P3HT crystal during SPL via the electrothermal process of Joule heating caused by the current through the P3HT crystal. In the PF-TUNA mode, we determined the contact area between the conducting tip and the sample assuming that the indentation of the hemispherical tip (radius, r ≈ 7 nm) was ca. 1 nm.41 The heated volume on the P3HT nanofiber was taken as a cylinder of height h (thickness of nanofiber, ≈11 nm) and radius 2r. The value of radius was considered to be 2r due to the spreading out of the current beneath the AFM tip.41 To simplify, we consider that this cylinder exchanges heat only laterally (optimistic case without heat flow toward substrate and tip). The heating power can be estimated as P = 2πrhφ.42 Here, the heat flow φ is calculated by φ = k(ΔT/Δx), with the thermal conductivity k = 0.18 W k−1 m−1.43 ΔT/Δx is the temperature gradient around the heated area, and Δx ≈ 2r. Then, the required minimum power (P) to reach the melting temperature (Tm = 220 °C, i.e., ΔT = 195 °C) is determined as 1.2 × 10−2 mW. In our case, it is difficult to evaluate directly the heating power based on the measured current due to the change in resistivity during the phase transition from crystalline to molten P3HT. Ginger et al. have noted the discrepancy of carrier mobility between planar device and conductive atomic force microscopy measurements at the same bias.41,44,45 The hypothesis has been proposed to explain the discrepancy such as tip-induced injection enhancement44 and nanoscale confinement effects.45 On P3HT thin films, the current densities at 1 and 7 V were estimated to be ≈107 and ≈108 A/m2, respectively.41 Because the bias voltage during scanning is typically small, e.g., +1 V, the corresponding low current density will cause only a small increase in temperature (P ≈ 1.2 × 10−4 mW, ΔT ≈ 2 °C; see the Supporting Information for details), which is not enough to change any physical or chemical property of the nanofiber. However, when a higher bias voltage is applied, e.g., VP = +7 V, a large current density will give rise to a large amount of Joule heat within the P3HT crystal (P ≈ 2.4 × 10−2 mW), which is capable of causing a significant rise in temperature within the nanofiber. It should be pointed out that even after taking into account the heat loss (≈90%), ΔT ≈ 380 °C still sufficient be reached at such power. When the Joule heating temperature (TJH) in the ambient atmosphere is high enough to result in the disappearance of conductivity, we suppose that we were able to induce localized physical changes (e.g., melting and rearrangement) and chemical degradation (e.g., thermally activated oxidation) within P3HT nanofibers (Figure 1b). As a result, in the regions exposed to high bias voltage the nanofiber was molten and oxidized (polymers shown in red

3. RESULTS AND DISCUSSION We created a conductivity pattern within an individual P3HT nanofiber, grown by melt crystallization in a thin P3HT film at 185 °C for 2 s. The mean height and the width of these P3HT nanofibers were 11 ± 1 nm and 90 ± 5 nm, respectively. Using TEM and GIWAXS measurements (see Figures S3−S5), we determined the packing and orientation of the P3HT chains within the nanofibers as edge-on. For this orientation, the side chains were oriented along the substrate normal, while πstacking between neighboring polymers was in the plane of the film (Figure 1a). Regular and well-defined packing of P3HT molecules with the individual nanofibers led to a rather constant lateral size (width) along their long axis. Details of the preparation of the samples and the characterization of the conductivity of the initial nanofibers can be found in the Experimental Section. The capability of patterning through bias-controlled SPL is demonstrated in Figure 1c−h. Using a conducting AFM tip, we were able to measure the current from the substrate (highly doped silicon on a metallic sample holder) through the P3HT nanofiber. Scanning across the sample, the morphology and conductivity of the initial nanofiber were characterized at a bias voltage of +1 V, as shown in Figure 1c,d. As can be seen in Figure 1d, along the whole nanofiber the current was rather uniform. Patterns in conductivity were fabricated and imaged in situ by well-controlled, periodically abrupt changes in the dc bias voltage (VP) from +7 to +2 V, as shown in Figures 1e, 1f, and 1j (left). After such patterning, the conductivity through the P3HT crystal in the vertical (height) direction along the P3HT nanofiber showed significant periodic variations, as shown in Figure 1h. The regions where VP = +7 V was applied showed almost no conductivity, separating conductive structures with a mean length of ≈200 nm along the P3HT nanofiber. Consequently, the current nearly disappeared in the region exposed to VP = +7 V. By contrast, the surface topography did not shown any detectable changes. Regions exposed to VP = +2 V showed no changes in conductivity or topography. The current−voltage (I−V) characteristics measured at two different positions on the patterned nanofiber are shown in Figure 1i, clearly indicating that the resistance of spot B (VP = +7 V) was much larger than that of spot A (VP = +2 V). It is suggested that the current across the P3HT crystal induced by the applied bias voltage caused localized Joule C

DOI: 10.1021/acs.macromol.8b01465 Macromolecules XXXX, XXX, XXX−XXX

Article

Macromolecules

Figure 2. Understanding the formation of patterns in conductivity. (a, b) Height and current images of the initial P3HT nanofiber. (c, d) Height and current images of the same P3HT nanofiber after patterning by bias-controlled SPL. (e, f) Appropriate washing in hot anisole allowed to selectively remove the conductive regions (crystalline parts scanned with VP = +2 V). (g, h) Height and current images of two initial P3HT nanofibers, the left being scanned under nitrogen and the right one being scanned in ambient atmosphere at various values of VP, not affecting topography (i) but leading to differences in conductivity (j). The size of the scale bar is 300 nm. The images in (a−h) were scanned in an ambient atmosphere.

Figure 3. Creating patterns in conductivity with different periodicities. Initial nanofibers, crystallized for 2 s at 185 °C, were alternatingly scanned at VP = +7 V and VP = +2 V, leading to patterns in conductivity with different periodicities: (a, e) 300 nm, (b, f) 200 nm, (c, g) 130 nm, and (d, h) 60 nm. The size of the scale bar is 400 nm. All images were scanned in an ambient atmosphere.

nanofiber were washed away entirely. However, the nonconductive parts of the nanofiber mostly remained. Resulting from local thermally activated oxidation46 insoluble structures were generated, as shown in Figure 2e. This observation is consistent with previous research by Gabriel et al. on P3HT samples, which became partially insoluble in hot solvents when annealed in ambient atmosphere at high temperatures (e.g., 250 and 350 °C).47 FTIR spectra of P3HT samples annealed at 280 °C for 3 min under different atmospheres (air and N2 atmospheres) are shown in Figure S6. These spectra exhibited the characteristic absorption feature corresponding to the CH2 stretching vibration of P3HT at 2850−2950 cm−1. Intriguingly, the

color in Figure 1b) which caused significant changes in conductivity. For an in-depth understanding of the mechanisms giving rise to the conductivity change after scanning at high VP, we carefully washed (see Experimental Section) a patterned sample, as shown in Figure 2a−f. The topography and current image of the initial nanofiber are shown in Figures 2a and 2b, respectively. After patterning through scanning at different bias voltages, the resulting domains with a mean length of ≈320 nm were clearly visible in the current image (Figure 2d). Again, no change was detectable in the topography image (Figure 2c). Subsequently, these nanostructures were washed for 10 s in anisole at 70 °C. As expected, the conductive regions of the D

DOI: 10.1021/acs.macromol.8b01465 Macromolecules XXXX, XXX, XXX−XXX

Article

Macromolecules

Figure 4. Adjusting the electrical conductivity. (a) Topography and (b) current image of the initial nanofiber. (c, d) Patterning was performed at different bias voltages (VP = +4, 5, and 6 V). (e) Topography and (f) current image of the nanostructures after patterning. The current images in (b) and (f) were acquired under VP = +1 V. The size of the scale bar is 400 nm. The images in (a−f) were scanned in ambient atmosphere. (g−i) Laterally averaged current along the nanofiber corresponding to different VP. The green dotted line indicates the mean current value of the nanofiber before SPL patterning.

the nanofiber under nitrogen imply that the molten and subsequently rapidly and nonisothermally recrystallized P3HT nanofiber consisted of crystalline structures of lower perfection. Accordingly, a higher reduction in conductivity could be achieved through a combination of the exposure to air and annealing at higher temperatures well above Tm. Thus, the question arises if we can fine-tune the electrical conductivity of conjugated polymer crystals through a control of the applied bias voltage and the thermally activated oxidation process. Moreover, Figure 3 shows a series of nanostructures with different periodicities of the conductivity patterns fabricated through bias-controlled SPL. By alternating VP periodically in the direction perpendicular to the fast-scan direction at a writing speed of 2 nm s−1, we achieved patterns with a period ranging from hundreds of nanometers to dozens of nanometers. In addition, by use of the concept of bias-induced SPL, a resolution limit about 50 nm, the value of AFM tip diameter, was achieved. Figure 4 shows that it is possible to tune the conductivity finely by controlling VP. Significant changes in the conductivity contrast were observed in the current image of Figure 4f when varying VP from +4 to +5 V and up to +6 V. For VP = +4 V, no conductivity contrast can be seen in Figure 4f. This suggests that the TJH of the nanofiber was still below the melting temperature Tm. Correspondingly, the value of the laterally averaged current along the nanofiber (light red line in Figure 4g) was similar to the value of the mean current of ∼125 fA measured on the nanofiber before SPL patterning (green dotted line in Figure 4g). The area scanned with VP = +5 V showed a change in conductivity, the corresponding average value of the current diminished to about 50 fA (see Figure 4h). When increasing VP to 6 V, the average current eventually decreased to a value close to zero (red line in Figure 4i),

infrared spectrum of oxidized sample had distinct additional absorbance features at around 1170 and 1675 cm−1. The strong absorption at 1675 cm−1 assigned to an aromatic ketone may indicate a reaction of alkoxy radicals with hydroxyl groups.10 In addition, the presence of the 1170 cm−1 peak, which is assigned to sulfone groups, also indicated that P3HT molecules were oxidized.10 It has also been confirmed that the mechanism of thermally activated oxidation involved radical oxidation of hexyl side-chains and the degradation of thiophene rings.48 The destruction of the macromolecular backbone resulted in a loss of π-conjugation, provoking the deterioration of the conductivity in areas scanned at high VP. Similar results were found in our studies on samples annealed on a hot stage under different surrounding gas atmospheres (see Figure S6), providing also evidence for the insolubility of oxidized P3HT. To examine the influence of the surrounding gas atmospheres during bias-controlled SPL, we have performed another experiment, as shown in Figure 2g−j. The initial topography and current image of two nanofibers are shown in Figures 2g and 2h, respectively. As in the previously shown experiments, one of the nanofibers was scanned in ambient atmosphere (air), whereas the other nanofiber was scanned the same way but under pure nitrogen gas. For the nanofiber exposed to air (Figure 2j) alternating conductive/nonconductive patterns were observed, as long as the TJH under the conductive tip exceeded Tm of P3HT. By contrast, scanning the P3HT crystal at different VP in nitrogen (Figure 2j) resulted in significantly different changes in the conductivity. At VP = +7 V, the current image showed only a reduction in conductivity of the scanned area. Only at VP = +9 V, the conductivity of the corresponding region disappeared. These changes in conductivity caused by scanning E

DOI: 10.1021/acs.macromol.8b01465 Macromolecules XXXX, XXX, XXX−XXX

Article

Macromolecules

Outstanding Young Talent Research Fund of Zhengzhou University (1521320004), and Startup Research Fund of Zhengzhou University (1512320001).

indicating that the P3HT nanofiber was transformed into an electric insulator. The poor electrical conductivity of a P3HT nanofiber at VP = +5 and +6 V results from only partial melting and thus only partial thermal oxidation at TJH close to Tm of P3HT. Similar results were obtained when a P3HT sample was annealed on a hot stage. A strong decrease in electrical conductivity with increasing annealing temperature (Ta) was observed when Ta was above 200 °C (see Figure S7). These experiments demonstrate the capacity of SPL to generate complex conductivity patterns in a continuous P3HT crystal by tuning the bias voltage during scanning.

■

(1) Liu, Y.; Cole, M. D.; Jiang, Y.; Kim, P. Y.; Nordlund, D.; Emrick, T.; Russell, T. P. Chemical and Morphological Control of Interfacial Self-Doping for Efficient Organic Electronics. Adv. Mater. 2018, 30 (15), 1705976. (2) Homyak, P. D.; Liu, Y.; Harris, J. D.; Liu, F.; Carter, K. R.; Russell, T. P.; Coughlin, E. B. Systematic Fluorination of P3HT: Synthesis of P(3HT-co-3H4FT)s by Direct Arylation Polymerization, Characterization, and Device Performance in OPVs. Macromolecules 2016, 49 (8), 3028−3037. (3) Noriega, R.; Rivnay, J.; Vandewal, K.; Koch, F. P.; Stingelin, N.; Smith, P.; Toney, M. F.; Salleo, A. A general relationship between disorder, aggregation and charge transport in conjugated polymers. Nat. Mater. 2013, 12 (11), 1038−1044. (4) Baran, D.; Ashraf, R. S.; Hanifi, D. A.; Abdelsamie, M.; Gasparini, N.; Röhr, J. A.; Holliday, S.; Wadsworth, A.; Lockett, S.; Neophytou, M.; et al. Reducing the efficiency-stability-cost gap of organic photovoltaics with highly efficient and stable small molecule acceptor ternary solar cells. Nat. Mater. 2017, 16 (3), 363−369. (5) Ren, Z.; Zhang, X.; Li, H.; Sun, X.; Yan, S. A facile way to fabricate anisotropic P3HT films by combining epitaxy and electrochemical deposition. Chem. Commun. 2016, 52 (73), 10972− 10975. (6) Seibers, Z. D.; Le, T. P.; Lee, Y.; Gomez, E. D.; Kilbey, S. M. Impact of Low Molecular Weight Poly (3-hexylthiophene) s as Additives in Organic Photovoltaic Devices. ACS Appl. Mater. Interfaces 2018, 10 (3), 2752−2761. (7) Shen, X.; Hu, W.; Russell, T. P. Measuring the Degree of Crystallinity in Semicrystalline Regioregular Poly(3-hexylthiophene). Macromolecules 2016, 49 (12), 4501−4509. (8) Rahimi, K.; Botiz, I.; Stingelin, N.; Kayunkid, N.; Sommer, M.; Koch, F. P. V.; Nguyen, H.; Coulembier, O.; Dubois, P.; Brinkmann, M.; Reiter, G. Controllable Processes for Generating Large Single Crystals of Poly (3-hexylthiophene). Angew. Chem., Int. Ed. 2012, 51 (44), 11131−11135. (9) Tremel, K.; Ludwigs, S. Morphology of P3HT in thin films in relation to optical and electrical properties. In P3HT Revisited-From Molecular Scale to Solar Cell Devices; Springer: 2014; pp 39−82. (10) Manceau, M.; Rivaton, A.; Gardette, J.-L.; Guillerez, S.; Lemaître, N. The mechanism of photo-and thermooxidation of poly (3-hexylthiophene)(P3HT) reconsidered. Polym. Degrad. Stab. 2009, 94 (6), 898−907. (11) Kim, T.-H.; Song, S. H.; Kim, H.-J.; Oh, S.-H.; Han, S.-Y.; Kim, G.; Nah, Y.-C. Effects of oxidation potential and retention time on electrochromic stability of poly (3-hexyl thiophene) films. Appl. Surf. Sci. 2018, 442, 78−82. (12) Choi, J.; Gunkel, I.; Li, Y.; Sun, Z.; Liu, F.; Kim, H.; Carter, K. R.; Russell, T. P. Macroscopically ordered hexagonal arrays by directed self-assembly of block copolymers with minimal topographic patterns. Nanoscale 2017, 9 (39), 14888−14896. (13) Wang, L.; Boutilier, M. S.; Kidambi, P. R.; Jang, D.; Hadjiconstantinou, N. G.; Karnik, R. Fundamental transport mechanisms, fabrication and potential applications of nanoporous atomically thin membranes. Nat. Nanotechnol. 2017, 12 (6), 509−522. (14) Forth, J.; Liu, X.; Hasnain, J.; Toor, A.; Miszta, K.; Shi, S.; Geissler, P. L.; Emrick, T.; Helms, B. A.; Russell, T. P. Reconfigurable Printed Liquids. Adv. Mater. 2018, 30 (16), 1707603. (15) Russell, T. P.; Chai, Y. 50th anniversary perspective: putting the squeeze on polymers: a perspective on polymer thin films and interfaces. Macromolecules 2017, 50 (12), 4597−4609. (16) Kim, H.; Kang, B.-G.; Choi, J.; Sun, Z.; Yu, D. M.; Mays, J.; Russell, T. P. Morphological Behavior of A2B Block Copolymers in Thin Films. Macromolecules 2018, 51 (3), 1181−1188.

4. CONCLUSION In our study, we have demonstrated that bias-controlled SPL represents a straightforward, single-step, in situ method for nanopatterning the conductivity of P3HT crystals, leaving its topography unperturbed. Nanoscale Joule heating allowed to raise the local temperature within P3HT nanofibers above its nominal melting temperature, allowing for a localized thermally activated oxidation, which generated nonconductive section on a P3HT nanofiber. A continuous transformation from a conductive to a nonconductive behavior of a P3HT crystal can be achieved by varying the applied bias voltage. With this approach based on a combination of Joule heating and thermally activated oxidation, we can control and pattern charge transport properties within conjugated polymer crystals at a nanometer length scale.

■

ASSOCIATED CONTENT

S Supporting Information *

. The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.macromol.8b01465. Experiments; temperature and time protocols; AFM height and current images of P3HT nanofibers before and after the washing process; TEM images and SAED patterns of P3HT nanofibers before and after annealing in air atmosphere with high temperature; 2D GIWAXS patterns of P3HT thin films annealing at different temperatures in different atmospheres; AFM images and 2D GIWAXS patterns of P3HT nanofibers before and after the washing process; FTIR spectra of P3HT thin films after annealing in different atmospheres; and the conductivity of P3HT nanofibers measured by PFTUNA after annealing at different temperatures in an air atmosphere (PDF)

■

REFERENCES

AUTHOR INFORMATION

Corresponding Authors

*E-mail [email protected] (B.Z.). *E-mail [email protected] (J.C.). ORCID

Bin Zhang: 0000-0002-8293-1321 Günter Reiter: 0000-0003-4578-8316 Notes

The authors declare no competing financial interest.

■

ACKNOWLEDGMENTS The authors are grateful to the National Science Foundation of China (Nos. 51773182 and 11372284), China Postdoctoral Science Foundation (Nos. 2016M592302 and 2018T110741), F

DOI: 10.1021/acs.macromol.8b01465 Macromolecules XXXX, XXX, XXX−XXX

Article

Macromolecules (17) Gartside, J. C.; Arroo, D. M.; Burn, D. M.; Bemmer, V. L.; Moskalenko, A.; Cohen, L. F.; Branford, W. R. Realization of ground state in artificial kagome spin ice via topological defect-driven magnetic writing. Nat. Nanotechnol. 2018, 13 (1), 53−58. (18) Wang, D.; Russell, T. P. Advances in Atomic Force Microscopy for Probing Polymer Structure and Properties. Macromolecules 2018, 51 (1), 3−24. (19) Ryu, Y. K.; Garcia, R. Advanced oxidation scanning probe lithography. Nanotechnology 2017, 28 (14), 142003. (20) Jo, A.; Joo, W.; Jin, W.-H.; Nam, H.; Kim, J. K. Ultrahighdensity phase-change data storage without the use of heating. Nat. Nanotechnol. 2009, 4 (11), 727−731. (21) Farina, M.; Ye, T.; Lanzani, G.; di Donato, A.; Venanzoni, G.; Mencarelli, D.; Pietrangelo, T.; Morini, A.; Keivanidis, P. E. Fast ultrahigh-density writing of low-conductivity patterns on semiconducting polymers. Nat. Commun. 2013, 4 (1), 2668. (22) Vettiger, P.; Despont, M.; Drechsler, U.; Durig, U.; Haberle, W.; Lutwyche, M. I.; Rothuizen, H. E.; Stutz, R.; Widmer, R.; Binnig, G. K. The “Millipede”More than thousand tips for future AFM storage. IBM J. Res. Dev. 2000, 44 (3), 323−340. (23) Binnig, G.; Despont, M.; Drechsler, U.; Haeberle, W.; Lutwyche, M.; Vettiger, P.; Mamin, H.; Chui, B.; Kenny, T. W. Ultrahigh-density atomic force microscopy data storage with erase capability. Appl. Phys. Lett. 1999, 74 (9), 1329−1331. (24) Altebaeumer, T.; Gotsmann, B.; Pozidis, H.; Knoll, A.; Duerig, U. Nanoscale shape-memory function in highly cross-linked polymers. Nano Lett. 2008, 8 (12), 4398−4403. (25) Zimmermann, S. T.; Balkenende, D. W. R.; Lavrenova, A.; Weder, C.; Brugger, J. Nanopatterning of a Stimuli-Responsive Fluorescent Supramolecular Polymer by Thermal Scanning Probe Lithography. ACS Appl. Mater. Interfaces 2017, 9 (47), 41454−41461. (26) Garcia, R.; Knoll, A. W.; Riedo, E. Advanced scanning probe lithography. Nat. Nanotechnol. 2014, 9 (8), 577−587. (27) Dagata, J. A.; Schneir, J.; Harary, H. H.; Evans, C.; Postek, M. T.; Bennett, J. Modification of hydrogen-passivated silicon by a scanning tunneling microscope operating in air. Appl. Phys. Lett. 1990, 56 (20), 2001−2003. (28) Garcia, R.; Martinez, R. V.; Martinez, J. Nano-chemistry and scanning probe nanolithographies. Chem. Soc. Rev. 2006, 35 (1), 29− 38. (29) Li, Y.; Maynor, B. W.; Liu, J. Electrochemical AFM “dip-pen” nanolithography. J. Am. Chem. Soc. 2001, 123 (9), 2105−2106. (30) Wang, C.; Ke, X.; Wang, J.; Liang, R.; Luo, Z.; Tian, Y.; Yi, D.; Zhang, Q.; Wang, J.; Han, X.-F.; et al. Ferroelastic switching in a layered-perovskite thin film. Nat. Commun. 2016, 7, 10636. (31) Berson, J.; Burshtain, D.; Zeira, A.; Yoffe, A.; Maoz, R.; Sagiv, J. Single-layer ionic conduction on carboxyl-terminated silane monolayers patterned by constructive lithography. Nat. Mater. 2015, 14 (6), 613−621. (32) Wu, J.; Yue, G.; Xiao, Y.; Ye, H.; Lin, J.; Huang, M. Application of a polymer heterojunction in dye-sensitized solar cells. Electrochim. Acta 2010, 55 (20), 5798−5802. (33) Zhang, B.; Chen, J.; Zhang, H.; Baier, M. C.; Mecking, S.; Reiter, R.; Mülhaupt, R.; Reiter, G. Annealing-Induced Periodic Patterns in Solution Grown Polymer Single Crystals. RSC Adv. 2015, 5 (17), 12974−12980. (34) Zhang, B.; Chen, J.; Baier, M. C.; Mecking, S.; Reiter, R.; Mülhaupt, R.; Reiter, G. Molecular-Weight-Dependent Changes in Morphology of Solution-Grown Polyethylene Single Crystals. Macromol. Rapid Commun. 2015, 36 (2), 181−189. (35) Wang, B.; Tang, S.; Wang, Y.; Shen, C.; Reiter, R.; Reiter, G. n.; Chen, J.; Zhang, B. Systematic Control of Self-Seeding Crystallization Patterns of Poly (ethylene oxide) in Thin Films. Macromolecules 2018, 51 (5), 1626−1635. (36) Zhang, B.; Chen, J.; Freyberg, P.; Reiter, R.; Mülhaupt, R.; Xu, J.; Reiter, G. n. High-Temperature Stability of Dewetting-Induced Thin Polyethylene Filaments. Macromolecules 2015, 48 (5), 1518− 1523.

(37) Zhang, B.; Wang, B.; Chen, J.; Shen, C.; Reiter, R.; Chen, J.; Reiter, G. Flow-Induced Dendritic β-Form Isotactic Polypropylene Crystals in Thin Films. Macromolecules 2016, 49 (14), 5145−5151. (38) Desbief, S.; Hergué, N.; Douhéret, O.; Surin, M.; Dubois, P.; Geerts, Y.; Lazzaroni, R.; Leclere, P. Nanoscale investigation of the electrical properties in semiconductor polymer-carbon nanotube hybrid materials. Nanoscale 2012, 4 (8), 2705−2712. (39) Zhang, B.; Chen, J.; Liu, B.; Wang, B.; Shen, C.; Reiter, R.; Chen, J.; Reiter, G. Morphological Changes of Isotactic Polypropylene Crystals Grown in Thin Films. Macromolecules 2017, 50 (16), 6210−6217. (40) Lyuksyutov, S. F.; Vaia, R. A.; Paramonov, P. B.; Juhl, S.; Waterhouse, L.; Ralich, R. M.; Sigalov, G.; Sancaktar, E. Electrostatic nanolithography in polymers using atomic force microscopy. Nat. Mater. 2003, 2 (7), 468−472. (41) Reid, O. G.; Munechika, K.; Ginger, D. S. Space charge limited current measurements on conjugated polymer films using conductive atomic force microscopy. Nano Lett. 2008, 8 (6), 1602−1609. (42) Gidon, S.; Lemonnier, O.; Rolland, B.; Bichet, O.; Dressler, C.; Samson, Y. Electrical probe storage using Joule heating in phase change media. Appl. Phys. Lett. 2004, 85 (26), 6392−6394. (43) Duda, J. C.; Hopkins, P. E.; Shen, Y.; Gupta, M. C. Thermal transport in organic semiconducting polymers. Appl. Phys. Lett. 2013, 102 (25), 251912. (44) Douhéret, O.; Lutsen, L.; Swinnen, A.; Breselge, M.; Vandewal, K.; Goris, L.; Manca, J. Nanoscale electrical characterization of organic photovoltaic blends by conductive atomic force microscopy. Appl. Phys. Lett. 2006, 89 (3), 032107. (45) Lin, H.-N.; Lin, H.-L.; Wang, S.-S.; Yu, L.-S.; Perng, G.-Y.; Chen, S.-A.; Chen, S.-H. Nanoscale charge transport in an electroluminescent polymer investigated by conducting atomic force microscopy. Appl. Phys. Lett. 2002, 81 (14), 2572−2574. (46) Fornari, R. P.; Blom, P. W.; Troisi, A. How Many Parameters Actually Affect the Mobility of Conjugated Polymers? Phys. Rev. Lett. 2017, 118 (8), 086601. (47) Rodrigues, A.; Castro, M. C. R.; Farinha, A. S. F.; Oliveira, M.; Tomé, J. P. C.; Machado, A. V.; Raposo, M. M. M.; Hilliou, L.; Bernardo, G. Thermal stability of P3HT and P3HT:PCBM blends in the molten state. Polym. Test. 2013, 32 (7), 1192−1201. (48) Ljungqvist, N.; Hjertberg, T. Oxidative degradation of poly(3octylthiophene). Macromolecules 1995, 28 (18), 5993−5999.

G

DOI: 10.1021/acs.macromol.8b01465 Macromolecules XXXX, XXX, XXX−XXX