Control of Crystallite Orientation in Diketopyrrolopyrrole-Based

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Control of Crystallite Orientation in Diketopyrrolopyrrole-Based Semiconducting Polymers via the Tuning of Intermolecular Interactions Sung Y. Son, Gang-Young Lee, Sangwon Kim, WonTae Park, Sang Ah Park, Yong-Young Noh, and Taiho Park ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.8b20297 • Publication Date (Web): 19 Feb 2019 Downloaded from http://pubs.acs.org on February 19, 2019

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ACS Applied Materials & Interfaces

Control of Crystallite Orientation in Diketopyrrolopyrrole-Based Semiconducting Polymers via the Tuning of Intermolecular Interactions Sung Yun Son,† Gang-Young Lee,†,§ Sangwon Kim,† Won-Tae Park,† Sang Ah Park,† Yong-Young Noh,*,† and Taiho Park*,† †Department

of Chemical Engineering, Pohang University of Science and Technology

(POSTECH), Pohang, Gyeongbuk, 37673, Republic of Korea. KEYWORDS: semiconducting polymers, diketopyrrolopyrrole, crystallite orientations, intermolecular interactions, charge transport

ABSTRACT Considerable previous studies have focused on the notion that semiconducting polymers with edge-on dominant orientation are advantageous for horizontal charge transport while polymers with face-on dominant orientation are advantageous for vertical charge transport, since the crystallite orientation determines the ππ stacking direction, which in turn affects the interchain charge transport direction. Here, we report that the crystallite orientation is dependent on the

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intermolecular interactions in the semiconducting polymer. In this study we control the intermolecular interactions in a donoracceptor (DA) semiconducting polymer via side chain engineering. To perform side chain engineering, we use two different polymers, in the first instance with side chains only on A units (PDPP-B), and then in the second instance, side chains on both D and A units (PDPP-C8). We observe that PDPP-C8 is characterized by weaker intermolecular interactions due to the additional side chains on D units. An morphology analysis reveals that PDPP-B and PDPP-C8 films have microstructures, which are characterized by edge-on and faceon dominant orientation, respectively. Therefore, we demonstrate that our strategies effectively control intermolecular interactions and, consequently, the crystallite orientation. Finally, we compare the vertical and horizontal mobilities of both polymer films. These results show that crystallite orientation has significant influence on charge transport behaviors.

1. Introduction Semiconducting polymers have attracted significant attention in the field of organic electronics due to their tunable optoelectronic properties, solution processability, and their potential applications to soft electronic devices.1-9 In the past two decades, significant progress has occurred with the development of various semiconducting polymers, especially in the fields of organic fieldeffect transistors (OFETs) and organic photovoltaics (OPVs).10-18 Among the various semiconducting polymers, donoracceptor (DA) semiconducting polymers are characterized by outstanding performance, e.g., diketopyrrolopyrrole(DPP)-based OFETs have high carrier mobilities (i.e., > 10 cm2/Vs),19-22 while benzothiadiazole(BT)-based OPVs have high power conversion efficiencies (i.e., > 10%).23-25 These characteristics are attributable to alternating


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between electron-rich and electron-deficient units along the backbone, which intrinsically produces push–pull driving forces that affect the polymer and induces intrachain charge transport.26 Moreover, the high backbone planarity, which is characteristic of DA semiconducting polymers, can enhance ππ interactions between adjacent polymer backbones, thus facilitating interchain charge transport.27,28 Numerous previous studies have analyzed the thin film morphology of semiconducting polymers to obtain an improved understanding of charge transport mechanisms in semiconducting polymers.29-32 These studies have also attempted to gain insight into molecular design strategies, with the aim of producing high-mobility semiconducting polymers. Thanks to these efforts, we now know that various factors, such as crystallinity, crystallite size, chain packing distance, and crystallite orientation, significantly affect the charge transport in semiconducting polymers and, consequently, device performance.33 Among these factors, crystallite orientation can significantly influence the direction of interchain charge transport, because interchain charge transport occurs through ππ stacking between adjacent polymer backbones and the ππ stacking direction in crystallites is determined by the crystallite orientation.34,35 A popular notion is the idea that semiconducting polymers with edge-on dominant orientation are advantageous for horizontal charge transport while polymers with face-on dominant orientation are advantageous for vertical charge transport (Figure 1).10,36,37


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Figure 1. Schematic representation of horizontal and vertical interchain charge transport via ππ stacking between crystallite polymer chains that adopt edge- and face-on orientations, respectively. It has been generally known that crystallites orientation can be determined by the interplay between polymer intermolecular interaction and polymer-substrate interaction.38 However, Chen et al. and Kim et al. reported that polymer-substrate interaction is a minor factor while chain aggregation propensity in solution determines a film’s chain orientation, respectively.39,40 They suggested that polymers with a high propensity to aggregate (i.e. strong intermolecular interaction) tend to adopt edge-on orientation. Conversely, polymers that do not readily aggregate (i.e. weak intermolecular interaction) easily interact with a substrate and are more likely to adopt face-on orientation. Based on the previous results, it appears to us that the tuning of intermolecular interactions between polymer chains is key to controlling crystallite orientation. The most commonly used method to control intermolecular interactions in DA semiconducting polymers is the modification of side chain bulkiness on either electron rich (D) or electron deficient (A) units. Although it is an effective and straightforward method, side chain modifications essentially require that side chains be carefully chosen to tune intermolecular interactions and to maintain a balance between polymer solubility and its electrical properties at the same time. Considering this, we assumed that


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introduction of additional short side chains to the polymer backbones can be a much easier approach to enable effective changes in the intermolecular interaction that causes drastic changes in the chain orientation without significant decrease in solubility and electrical properties. Therefore, in this study, we introduce longer side chains onto electron deficient (A) units to control solubility and shorter side chains onto electron rich (D) units to tune the intermolecular interactions. As shown in Scheme 1, we synthesized two different DPP-based polymers; one with only octyldodecyl chains on the DPP units (PDPP-B), the other with octyldodecyl chains on the DPP units and octyloxy chains on the benzene units (PDPP-C8). We then compared the polymer chain aggregation behaviors and chain orientation. FET and space charge limited current (SCLC) mobilities were measured and correlated with their chain orientations.

Scheme 1. Synthetic routes for PDPP-B and PDPP-C8. 2. Results and Discussion 2.1. Synthesis and characterization Scheme 1 shows the synthetic routes for PDPP-B and PDBB-C8. Both polymers were prepared via the Suzuki cross-coupling polymerization. The detailed synthesis procedures are described in the Supporting Information. Polymer chemical structures were confirmed using 1H NMR


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spectroscopy and their spectra are well consistent with the presented chemical structures (Figure S1 and S2). Table 1. Number average molecular weight, polydispersity, thermal, optical, and electronic properties of the PDPP-B and PDPP-C8 polymers.

polymer

Mn (kg/mol)

ĐMa)

Td (°C)b)

Tm

Tc

εabs

ΔEg

(°C)c) (°C)d) (M–1cm–1) (eV)e)

EHOMO

ELUMO

(eV)f)

(eV)g)

PDPP-B

52.2

2.3

394

295

261

106,000

1.55

–5.40

–3.85

PDPP-C8

48.2

2.2

386

240

221

98,000

1.43

–5.25

–3.82

a)Dispersity

calculated using ĐM = Mw/Mn; b)Decomposition temperature (5% weight loss) determined by TGA; b)Melting temperature at the endothermic peak in the DSC thermogram heating curve; c)Recrystallization temperature at the exothermic peak in the DSC thermogram cooling curve; d)Optical bandgap calculated from the absorption onsets in the film state; e)EHOMO was estimated from Eonset from the first oxidation potential relative to ferrocene (Fc), Fc/Fc+ redox system the ionization potential (IP) value was -4.8 eV; f)ELUMO = EHOMO + Eg. As shown in Table 1, the molecular weights and dispersity of both polymers are quite similar, which is necessary to avoid any effects that molecular weight may have on chain aggregation properties. The gel permeation chromatography (GPC) profiles for PDPP-B and PDPP-C8 are shown in Figure S3. 2.2. Thermal, optical, and electronic properties Polymer thermal properties were investigated via thermogravimetric analysis (TGA) and differential scanning calorimetry (DSC). Although both polymers exhibited similar thermal stabilities, which were analyzed using TGA (Figure S4), their melting (Tm) and recrystallization temperatures (Tc), obtained via DSC measurements, were distinctly different (Figure 2a,b). Table 1 summarizes the thermal properties of the polymers.


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Figure 2. DSC thermograms of a) PDPP-B and b) PDPP-C8. c) Absorption spectra for the polymer solutions. d) Normalized film absorption spectra. The films were annealed at 200C for 10 min. Normalized temperature dependent UV-vis absorption spectra for e) PDPP-B and f) PDPP-C8. In the PDPP-B DSC thermograms, Tm and Tc appear at 295 and 261°C, respectively, which corresponds to endothermic and exothermic processes, respectively. When compared with PDPPB, PDPP-C8 is characterized by lower Tm and Tc, which indicate decreased intermolecular interactions between polymer chains in the PDPP-8 crystalline domains. We attribute the decrease in intermolecular interactions to the additional octyloxy chains on the benzene units, which hinder robust chain packing. Next, optical properties for the PDPP-B and PDPP-C8 solutions and films were investigated using UV-Vis absorption spectroscopy to study chain conformation and packing modes. Table 1 presents a detailed summary of their absorption properties. In the polymer solution absorption spectra, absorption peaks mainly appear between 600 and 800 nm, which may indicate strong intrachain charge transfer between the electron rich (D) and


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electron deficient (A) units (Figure 2c).41 This is supported by optimized geometries and electron density models for PDPP-B and PDPP-C8, which were calculated using density functional theory (DFT) models at the B3LYP/6-311G* level (Figure S5). Since both polymers have high backbone planarity, electrons are delocalized along the polymer backbones in both the highest occupied molecular orbital (HOMO) and the lowest unoccupied molecular orbital (LUMO) states. Note that PDPP-C8 shows a smaller dihedral angle between the thiophene and benzene rings (Φ2) than PDPP-B, which can be attributed to the SO interaction between the thiophene ring and methoxy side chain.42 In the polymer solution spectra, the maximum absorption coefficient (εabs) for the two polymers is similar. The PDPP-C8 absorption spectrum is slightly red-shifted compared with the spectrum for PDPP-B. Electron donating effects along the octyloxy chains increase PDPP-C8’s HOMO level. This tendency also occurred in the film state, thus,PDPP-C8 exhibited a reduced optical band gap (ΔEg) (Figure 2d and Table 1). Cyclic voltammetry (CV) analysis was performed to measure polymer energy levels (Figure S6). The HOMO energy levels were estimated from the oxidation onsets in the cyclic voltammograms. The PDPP-B and PDPP-C8 HOMO levels were –5.40 and –5.25 eV, respectively (Table 1), using the ferrocene energy level at –4.8 eV as an internal standard. The LUMO energy levels were derived from the HOMO levels and ΔEg. The PDPP-B and PDPP-C8 LUMO levels were –3.85 and –3.82 eV, respectively. Finally, PDPP-B and PDPP-C8 solutions were characterized using temperature-dependent UVvis absorption spectroscopy to investigate chain conformation and aggregation properties. At 30C, both polymer solutions show distinct A0–0 peaks when compared with their A0–1 peaks


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(Figure 2e,f). In the PDPP-B solution UV–vis spectra, the A0–0 peak intensities decreased only slightly upon heating while the intensities significantly decreased in the UV–vis spectra for the PDPP-C8 solution. Since the A0–0 peaks indicate intrachain charge delocalization, this suggests that PDPP-C8 chains become easily isolated from chain aggregates and adopt coild conformation at higher temperatures leading to shortening of the conjugation length.43 Conversely, PDPP-B chains tend to maintain self-organized structures and consequently their extended conformation even at higher temperatures. 2.3. Microstructure study To investigate the effects of intermolecular interaction intensity on crystallite orientation, polymer films were characterized using grazing incidence wide angle X-ray scattering (GIWAXS). As shown in Figure 3, as-cast films of both polymers displayed weak scattering peaks, which indicates that they have low crystallinity. To induce chain ordering, we annealed the films at 200C for 10 min (annealed film). As a result, strong scattering peaks appeared in the annealed films for both polymers. The distinct, out-of-plane (h00) peaks in the GIWAXS pattern for the PDPP-B annealed film indicate regular and periodic lamellar packing, as well as edge-on dominant crystallite orientation because out-of-plane lamellar peaks generally originate from crystallites that adopt edge-on orientation. Distinct scattering peaks in both out-of-plane and in-plane directions appeared in the GIWAXS pattern for the PDPP-C8 annealed film. The strong, out-of-plane (010) peaks indicate that crystallites in the PDPP-C8 annealed film predominantly adopted face-on orientation. Thus, these results suggest that edge-on orientation is thermodynamically favorable for PDPP-B that has


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stronger intermolecular interaction, while face-on orientation is thermodynamically favorable for PDPP-C8 that has weaker intermolecular interaction.

Figure 3. GIWAXS patterns for as-cast, annealed, and melt-annealed PDPP-B and PDPP-C8 films. We annealed both polymer films at temperatures above Tm (300C for PDPP-B and 250C for PDPP-C8) for 10 min and observed the resulting GIWAXS patterns for the films (melt-annealed films). The PDPP-B melt-annealed film was characterized by a similar pattern when compared with the PDPP-B annealed film. On the other hand, the GIWAXS pattern for the PDPP-C8 melt-annealed film had a completely different pattern compared with its annealed film. The melt-annealed film is only characterized by strong out-of-plane (h00) scattering peaks, which indicate edge-on dominant orientation. This drastic change in crystallite orientation in PDPP-C8 implies that intermolecular interactions


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between chains may significantly increase during annealing at temperatures above its Tm and cooling below its Tc. Since the chain packing distance is closely related to intermolecular interactions, we extracted the out-of-plane and in-plane GIWAXS spectra from the GIWAXS patterns to evaluate lamellar packing distances in each film (Figure S7). In all films, PDPP-C8 had longer packing distances than PDPP-B, which is attributable to the additional octyloxy chains on benzene units in PDPPC8 (Table S1). Out-of-plane lamellar distance in the PDPP-C8 melt-annealed film (21.74 Å) is shorter than its distance for the annealed film (21.81 Å), while the PDPP-B films were characterized by the opposite trends (a distance of 19.71 and 20.59 Å for the annealed and meltannealed films, respectively). Such decrease in the lamellar packing distance in PDPP-C8 indicate the presence of enhanced intermolecular interaction upon melt-annealing. To monitor changes in the PDPP-C8 microstructure, we characterized thin films by varying the temperature and holding time at 250C using a GIWAXS, equipped with a sample stage for in-situ heating. For the out-of-plane spectra, the scattering peak intensity (100) is very weak from RT to 200C. Intensity begins to increase at 250C, which is above the Tm for PDPP-C8. Intensity increases further as the holding time increases to up to 24h (Figure 4a). On the other hand, in-plane (100) peak intensity increases from RT to 150C (Figure 4b). This is associated with the appearance of the out-of-plane (010) peak at 150C, because they both generally originate from lamellar and ππ stacking in crystallites with face-on orientation. Then, the in-plane (100) peak intensity gradually decreases as temperature increases to 250C. Finally, the in-plane (100) peaks nearly vanish during


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holding at 250C. Meanwhile, a (010) peak appeared at 250C and slightly increased with holding time, which is also associated with the distinct out-of-plane (100).

Figure 4. a) Out-of-plane and b) in-plane GIWAXS spectra for PDPP-C8 as a function of temperature and holding time at 250C. c) Schematic illustrations showing changes in the crystallite orientations for PDPP-B and PDPP-C8 when increasing the annealing temperature. Based on these results, we suggest that PDPP-C8 translational chain motion becomes more active and thus chain ordering occurs as the temperature increases to 150C, which leads to crystallite formation with face-on orientation (Figure 4c). As the temperature increases above Tm, intense and long-range chain motion is induced, allowing chains to vigorously interact with each other. As a result, the chain packing distance shortens and crystallites increasingly adopt edge-on


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orientation. Conversely, crystallites predominantly adopt edge-on orientation in PDPP-B, due to its intrinsically stronger intermolecular interactions. We also observed the annealed and melt-annealed film surface morphology for both polymers using atomic force microscopy (AFM). We found that the morphologies between the annealed and melt-annealed films are similar for both polymers (Figure S8). This suggests that changes in the crystallite orientation during melt-annealing do not affect the surface morphology. 2.4. Charge transport analysis Based on differences in the microstructural features of both polymers, the activation energies (ΔEA) for charge transport were measured in a FET configuration. This was accomplished by varying the annealing temperature to investigate the effect of crystallite orientation on charge transport. The ΔEA values were assessed using the Arrhenius equation and are summarized in Table S2. The temperature-dependent transfer curves are shown in Figure S9.

Figure 5. ΔEA for charge transport in PDPP-B and PDPP-C8 films as a function of the annealing temperature.


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As shown in Figure 5, the ΔEA values for the PDPP-B films do not change significantly depending on the annealing temperature. This is explained by the occurrence of edge-on dominant orientation that is maintained throughout the entire range of annealing temperatures, even above Tm. Meanwhile, for the PDPP-C8 films, we obtain a maximum ΔEA value with film annealing at 150C, whereupon PDPP-C8 predominantly adopts face-on orientation. Afterward, the value decreases with the annealing temperature until it reaches a minimum, whereupon the PDPP-C8 melt-annealed film predominantly adopts edge-on orientation. This result supports the popular notion that edge-on orientation is favorable for horizontal charge transport in semiconducting polymers. Finally, we measured the SCLC and FET mobilities of PDPP-B and PDPP-C8 films annealed at 200°C for 10 min to investigate the effects of crystallite orientation on charge mobilities because such films were characterized by significant differences in their crystallite orientation. Currentvoltage curves for the SCLC and FET devices are shown in Figure S10 and S11, respectively. Table 2. SCLC and FET mobilities of the PDPPB and PDPP-C8 films annealed at 200°C for 10 min. SCLC mobility

FET mobility

(10-4 ×cm2/Vs)a)

(cm2/Vs)b)

PDPP-B

0.095

0.41 ± 0.05

PDPP-C8

1.4

0.011 ± 0.001

polymer

a)The

SCLC mobilities are from single measurement. b)The FET mobilities are averages of 6 devices. The PDPP-C8 annealed film with face-on dominant orientation has higher SCLC mobility, but lower FET mobility when compared with the PDPP-B annealed film with edge-on dominant orientation (Table 2). This result clearly shows that the introduction of additional octyloxy side


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chains on PDPP-C8 leads to a distinct difference in charge transport behaviors between PDPP-B and PDPP-C8. 3. Conclusion We successfully synthesized two different DPP-based polymers with side chains on only DPP units, and then with side chains on both DPP and benzene units, to compare intermolecular interaction intensities and the resulting crystallite orientation. PDPP-C8, with additional side chains, has weak intermolecular interaction and face-on dominant orientation in its annealed film, while PDPP-B has strong intermolecular interaction and edge-on dominant orientation in its annealed film. Annealing above Tm induces improvement in intermolecular interaction for PDPPC8. As a result, the PDPP-C8 melt-annealed film showed drastic orientation changes compared with its annealed film. Based on the stark differences in crystallite orientation between the PDPP-B and PDPP-C8 annealed films, we showed that confirmed the popular relationships between crystallite orientation and charge transport by measuring SCLC and FET mobilities. Therefore, this study demonstrates that our strategies can effectively control intermolecular interactions and crystallite orientation in DA semiconducting polymers, which can be modulated depending on the device structure. ASSOCIATED CONTENT Supporting Information. The following files are available free of charge. Experimental details and additional figures. (PDF) AUTHOR INFORMATION Corresponding Author


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*E-mail: [email protected] (Y.-Y.N). *E-mail: [email protected] (T.P.). Present Addresses § Korea Conformity Laboratories, Seoul, 08503, Republic of Korea. Author Contributions The manuscript was written through contributions of all authors. All authors have given approval to the final version of the manuscript. Notes The authors declare no competing financial interest. ACKNOWLEDGMENT The GIWAXS measurement was performed at a synchrotron radiation source on beamline 3C at the Pohang Accelerator Laboratory (PAL), Pohang, Korea. This work was supported by the Technology Development Program to Solve Climate Changes of the National Research Foundation of Korea (NRF) (Code No. 2015M1A2A2056216) and the Center for Advanced Soft Electronics under the Global Frontier Research Program (Code No. 2012M3A6A5055225 and 2013M3A6A5073183). REFERENCES (1) Sirringhaus, H.; Tessler, N.; Friend, R. H. Integrated Optoelectronic Devices Based on Conjugated Polymers. Science 1998, 280, 1741-1744.


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Control of Crystallite Orientation in Diketopyrrolopyrrole-Based

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