Importance of Critical Molecular Weight of Semicrystalline n-Type

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Importance of Critical Molecular Weight of Semicrystalline n-Type Polymers for Mechanically-Robust, Efficient Electroactive Thin Films Joonhyeong Choi, Wansun Kim, Donguk Kim, Seonha Kim, Junsu Chae, Siyoung Q. Choi, Felix Sunjoo Kim, Taek-Soo Kim, and Bumjoon J. Kim Chem. Mater., Just Accepted Manuscript • DOI: 10.1021/acs.chemmater.8b05114 • Publication Date (Web): 23 Apr 2019 Downloaded from http://pubs.acs.org on April 27, 2019

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Chemistry of Materials

Importance of Critical Molecular Weight of Semicrystalline n-Type Polymers for Mechanically-Robust, Efficient Electroactive Thin Films

Joonhyeong Choi,†,‡ Wansun Kim,†,§ Donguk Kim,‡ Seonha Kim,‡ Junsu Chae,‡ Siyoung Q. Choi,‡ Felix Sunjoo Kim,|| Taek-Soo Kim,*,§ and Bumjoon J. Kim*,‡

‡

Department of Chemical and Biomolecular Engineering, §Department of Mechanical

Engineering, Korea Advanced Institute of Science and Technology (KAIST), Daejeon, 34141, Republic of Korea ||

School of Chemical Engineering and Materials Science, Chung-Ang University (CAU), Seoul,

06974, Republic of Korea

*

(T.-S. Kim) E-mail: [email protected]

*

(B. J. Kim) E-mail: [email protected]

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ACS Paragon Plus Environment

Chemistry of Materials 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

ABSTRACT Mechanical properties of conducting polymers are an essential consideration in the design of flexible and stretchable electronics, but the guideline for the material design having both high mechanical and electrical properties remains limited. Here we provide an important guideline for the design of mechanically-robust, electroactive polymer thin films in terms of the molecular weight of the polymers. These studies based on a highly-efficient, representative n-type conjugated polymer (P(NDI2OD-T2)) revealed a marked enhancement in mechanical properties across a narrow molecular weight range, highlighting the existence of a critical molecular weight that can be exploited to engineer films that balance processability, mechanical, and electronic properties. We found the thin films formed from high molecular weight polymers (i.e., number-average molecular weight (Mn) ~ 163 kg mol-1) to exhibit superior mechanical compliance and robustness, with a 114-fold enhanced strain at fracture and a 2820-fold enhanced toughness, as compared to those of low molecular weight polymer films (Mn = 15 kg mol-1). In particular, we observed a jump in the mechanical properties between the Mn = 48 and 103 kg mol-1, yielding a 26-fold enhanced strain at fracture and a 160fold enhanced toughness. The significant improvement of tensile properties indicates the presence of a critical molecular weight at which entangled polymer networks start to form, as supported by the analysis of the thermal and crystalline properties, specific viscosity and microstructure. Our work provides useful guidelines for the design of conjugated polymers with recommendations for the best combinations of mechanical robustness and electrical performance for flexible and stretchable electronics.

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1. INTRODUCTION With the rapid advancements in flexible electronic devices, e.g., smart wearable sensors, the development of intrinsically stretchable electronic materials is becoming increasingly important.1-2 Conjugated polymers are considered as one of the most promising candidate materials for flexible electronics owing to their light weight, flexibility, and tunable optical and electronic properties. However, the rigid polymer backbone and high crystallinity of conjugated polymers, which often are prerequisites for good charge transport, may result in inferior ductility and toughness.3-7 Most of high-performance conjugated polymers (i.e., donoracceptor (D-A) type copolymers) consist of different aromatic groups and fused rings in the polymer backbone, which induces higher tensile modulus and reduces flexibility.8,9 Therefore, achieving superior mechanical and electrical properties simultaneously with conjugated polymers is a grand challenge in the field of organic electronics. While researchers have started to investigate the mechanical properties of conjugated polymers, these efforts have focused on p-type polymers, mainly poly(3-hexylthiophene) (P3HT),10-13 with almost no reports to date on their n-type counterparts. Considering the importance of high-performance n-type polymers for the development of organic complementary circuits in wearable and mechanically-robust devices, it is essential to improve the stability of n-type conjugated polymers against mechanical stresses to a level suitable for commercialization.14 N-type conjugated polymers have been extensively utilized as electron acceptors for developing highly efficient organic field-effect transistors (OFETs) and all-polymer solar cells (all-PSCs) owing to their broad light absorption ability, chemical tunability and long-term stability.15-25 We recently reported the importance of n-type polymers for enhancing the mechanical properties of the active blend layer in all-PSCs.20,26 N-type conjugated polymers are usually integrated into OFETs and all-PSCs as sub-100 nm films. Geometrical confinement of polymers within thin films (thickness ≤ 100 nm) has a considerable influence on the 3



mechanical properties (i.e., plastic deformation zone size and entanglement density), resulting in different mechanical behavior compared to that of bulk.27-32 Thus, it is crucial to understand the mechanical properties of n-type conjugated polymers in thin films. However, the difficulty in handling and directly characterizing the mechanical response of sub-100 nm thin films has limited such studies.14,33,34 The molecular weight of conjugated polymers plays a key role in determining the material properties (i.e., electrical, structural, mechanical and thermal properties) and the optoelectronic device performance. Previous studies on regioregular P3HT based systems have shown that the crystallinity, film morphology, and charge transport properties are strongly affected by the molecular weight.35,36 Also, for other types of low bandgap polymers, controlling the molecular weight of the polymers is known to influence the electrical performance and polymer blend morphologies in PSCs.37-41 Based on studies of commodity polymers, it can be presumed that the molecular weight of conjugated polymers has critical impacts on their mechanical behavior in thin films. For example, the molecular weight and the film-thickness dependent tensile properties of polystyrene has been studied in the sub-100 nm thin films.28,30,31,33 While it is generally assumed that high molecular weight polymers are needed to achieve both the electrical and mechanical properties required for mechanicallyrobust conducting thin films, polymers should be designed with optimal molecular weights just exceeding the critical molecular weight, rather than aiming to maximize molecular weight. The mechanical properties of the polymer film are not expected to significantly vary once the molecular weight of the polymer surpasses the critical molecular weight, as this marks the onset of entangled chain networks.13,42-46 Rather, excessively high molecular weight polymers are often difficult to process into thin film devices due to their low solubility and extremely high viscosity. Nevertheless, there are limited reports describing the mechanical properties of semicrystalline conjugated polymer thin films as a function of molecular weight, or that determine 4


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the critical molecular weight. To the best of our knowledge, studies of the molecular weight effects on the mechanical properties are reported only for P3HT.10-13,47,48 Notably, P3HT is a rather special case, in that it is used in electronic devices as a homopolymer and can be synthesized by living polymerization methods. In contrast, most of the recently developed conjugated polymers are D-A copolymers, often with more than two chemically distinct monomers per chain and complex structures. Further, these are synthesized by different polymerization methods (e.g., step-growth polymerization) than P3HT (e.g., quasi-living polymerization). Given the potential of fundamental studies on the intrinsic mechanical properties of n-type D-A conjugated polymers to inform the design of flexible electronics from these materials and other high-performance p/n types of D-A conjugated polymers, we sought to investigate the thermal, mechanical, morphological, and optoelectronic properties of a model high-performance D-A conjugated polymer. In this work, we chose n-type conjugated polymer (poly-[[N, N’ –bis(2-octyldodecyl)naphthalene-1,4,5,8-bis-(dicarboximide)-2,6-diyl]-alt-5,5’-(2,2’-bithiophene)]) (P(NDI2ODT2)) as a model system to investigate the effect of molecular weight on mechanical properties of conjugated polymers in thin films. As one of the best preforming n-type semicrystalline conjugated polymers reported to date, (P(NDI2OD-T2)) has been extensively used in OFETs and PSCs.49-52 The optical and electrical properties of (P(NDI2OD-T2)) are well known,53-58 but little is known about the mechanical properties in thin films. Therefore, our work aims to gain comprehensive understanding of the structure-property relationships, i.e., the optical properties, thermal properties, specific viscosity, microstructural properties and mechanical behaviors of (P(NDI2OD-T2)) as a function of molecular weight. Findings for the model (P(NDI2OD-T2)) system could guide the design of other conjugated polymers for implementation in thin films. First, P(NDI2OD-T2) batches were synthesized with controlled number-average molecular weights (Mn), ranging Mn = 15 , 20, 48, 103 to 163 kg mol-1 with 5



similar dispersity (Ð ). The crystallinity and specific viscosity values did not correlate linearly with Mn, but rather a sudden jump in the properties of the thin films was seen between 48 and 103 kg mol-1, producing 26- and 160-fold increases in the strain at fracture and toughness, respectively, despite the mere 2-fold increase in Mn. Such marked enhancements in mechanical properties in the films from the Mn 103 and 163 kg mol-1 are attributed to entangling of the polymer chains above a critical molecular weight. Moreover, the mechanical robustness of the thin films containing high-Mn polymers was preserved after thermal annealing, which is important to produce optimal performance in OFET devices.

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2. EXPERIMENTAL SECTION Characterizations. 1H-NMR spectra were measured on a Liquid 400 NB NMR spectrometer. Polymer molecular weights were estimated by SEC using o-dichlorobenzene (oDCB) as the eluent relative to polystyrene standards at 80 ℃ (flow rate: 1 mL min-1.) on a system equipped with a Waters 1515 Isocratic HPLC pump, a temperature control module, and a Waters 2414 refractive index detector. UV-visible absorption spectra were obtained on a UV1800 spectrophotometer (Shimadzu Scientific Instruments) at room temperature. The thermal properties were measured using a TA instruments DSC 25 with heating and cooling rates of 10 ℃ min-1. The melting temperature and melting enthalpy were reported from the second heating cycle. Specific viscosity was measured using an Anton Paar MCR 302 parallel plate rheometer, applying a constant shear rate (100 s-1) with a measuring time of 15 s per point at room temperature. The measured viscosity values of solvent (toluene and chlorobenzene) were within 3% of the literature values.59 The pristine and fractured surface morphologies and the thickness of the P(NDI2OD-T2) thin films were investigated by AFM (Park Systems XE-100) in noncontact mode. GIXS measurements were performed at beamline 3C in the Pohang Accelerator Laboratory (South Korea). GIXS samples were prepared by spin coating the polymers onto a Si substrate. X-rays with a wavelength of 1.1179 Å were used. An incidence angle of ~0.14° was chosen to allow for complete penetration of X-rays into the film.

Double Cantilever Beam (DCB) Test. The thickness of the P(NDI2OD-T2) thin films was controlled to 100 ± 15 nm by varying the solution concentration and the spin-coating speed depending on the molecular weight of P(NDI2OD-T2). P(NDI2OD-T2) thin films coated onto glass substrates were placed in a thermal evaporation chamber and held under high vacuum (< 10-6 Torr) for > 1 h before evaporating ~90 nm of Au. Following Au deposition, DCB specimens were fabricated by cutting the 25.4 mm × 25.4 mm samples into 8 mm × 25.4 7



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mm rectangles. Next, a glass substrate template (8 mm × 25.4 mm) was attached to the prepared sample using epoxy (Epo-Tek 353ND, consisting of bisphenol F and imidazole, Epoxy Technology) in order to prepare sandwiched structures with the following configuration: glass/Epoxy/Au/P(NDI2OD-T2) thin films/glass. The epoxy was cured at room temperature for 72 h in a dry box to exclude any thermal influence on the P(NDI2OD-T2) thin films. The high-precision DCB test equipment (Delaminator Adhesion Test System; DTS Company, Menlo Park, CA) was used, and all tests were performed under controlled conditions (~30% relative humidity, RH, at 25 ℃). The prepared DCB specimens were subject to multiple loading/crack-growth/unloading cycles at a constant displacement rate of 0.5 µm s-1 during the test, and the load versus displacement curves were recorded continuously. In the load versus displacement curves, the debond length (a), and the applied strain energy release rate (G) were calculated as

𝑎= ( 𝐺=

𝐶𝐸 ′ 𝐵ℎ3 8

12𝑃2 𝑎2 𝐸 ′ 𝐵2 ℎ 3

1/3

)

− 0.64ℎ, ℎ 2

(1 + 0.64 𝑎) ,

(1)

(2)

where C is the specimen compliance, du/dP, and u is the total displacement of the beam end. Here, P is the applied load, E' is the plane-strain modulus of the beam, B is the sample width, and h is the half height of the substrate. The critical fracture energy, Gc, is the critical value of G at the point of the critical load, Pc, at which the slope of the load-displacement curve starts to decrease. Pseudo Free-Standing Tensile Test. Tensile testing of the P(NDI2OD-T2) thin films having different Mn was performed on a water surface to separate the P(NDI2OD-T2) thin films from the glass substrate and support the specimen during the test. Each P(NDI2OD-T2) solution of different Mn was prepared under the same conditions as described in the DCB test 8


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specimen preparation. These prepared solutions were spin-coated on polystyrene sulfonate (PSS)-coated glass substrates. The PSS layer was used to float the P(NDI2OD-T2) thin films on a water surface as the PSS layer dissolves readily in water, allowing separation of the P(NDI2OD-T2) thin films from the glass substrate. The specimens were patterned in the shape of a dog-bone by a femtosecond laser. After floating the patterned specimen, each specimen was gripped by the PDMS-coated Al grips on the specimen gripping areas. The floated specimens of P(NDI2OD-T2) thin films were tested at a strain rate of ~0.8 × 10-3 s-1 until fracture, and the stress-strain curves were obtained. All tests were performed under controlled conditions (~30% relative humidity, RH, at 25 ℃).

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3. RESULTS AND DISCUSSION Synthesis of a Series of P(NDI2OD-T2) with Different Molecular Weights

Figure 1. (a) Synthesis of P(NDI2OD-T2) samples with different Mn; (b) UV-vis absorbance spectra of P(NDI2OD-T2) samples with different Mn in chloroform at 0.01 g L-1.

P(NDI2OD-T2) polymers with different Mn were synthesized by Stille coupling of 4,9dibromo-2,7-bis(2-octyldodecyl)benzo[lmn][3,8]-phenanthroline-1,3,6,8-tetraone and 5,5′bis(trimethylstannyl)-2,2′-bithiophene (Figure 1a), following a previously reported procedure.39 After finding conditions that yield high molecular weight polymers, the reaction time, temperature and the amount of end-capping agent were varied to control Mn. The detailed synthetic procedures are presented in Supporting Information. The Mn values were estimated to be 15.2, 20.5, 48.3, 103.5 and 163.2 kg mol-1 (P-15k, P-20k, P-48k, P-103k and P-163k) by size exclusion chromatography (SEC) relative to polystyrene standards using odichlorobenzene as the eluent at 80 °C to mitigate aggregation in the solution state (Table 1).

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Material Properties of P(NDI2OD-T2) with Different Mn. Table 1. Characteristics of P(NDI2OD-T2) with different Mn. λmax (nm)b

Tm (°C)c

Tc (°C)c

ΔHm (J g-1)c

1.75

621

282.2

267.1

4.0

20.5

2.46

654

309.1

292.2

7.6

P-48k

48.3

2.06

695

316.5

297.5

17.0

P-103k

103.5

2.28

703

325.1

303.6

11.1

Batch

Mn (kg mol-1)a

P-15k

15.2

P-20k

Ð

a

P-163k 163.2 2.42 700 323.7 295.7 9.2 b Estimated by SEC relative to polystyrene standards. Estimated from UV-vis absorption spectra of the polymer solutions. cDetermined from DSC analysis. a

UV-vis absorbance spectra of P(NDI2OD-T2) as a function of molecular weight were obtained in chloroform at 0.01g L-1 to examine the effects of molecular weight on polymer aggregation behavior (Figure 1b). Both high- and low-energy bands ascribed to π-π* transition (nearly 390 nm) and intramolecular charge transfer transition (500 – 800 nm), respectively, were observed in all samples. However, P(NDI2OD-T2) exhibited significantly different absorption behavior in the solution state as a function of Mn; namely, the absorption region corresponding to intramolecular charge transfer red-shifted with increasing Mn. The greatest shift of this peak into the near-infrared region was seen for the high-Mn P-103k and P-163k samples, both with an intense absorption and a shoulder at around 710 nm and 810 nm, respectively. These distinct absorption properties as a function of Mn are ascribed to different chain aggregation tendency in the solution state. The high-Mn polymers P-103k and P-163k show evidence of aggregation even at the low concentrations used for the UV-Vis spectroscopy studies due to coiling or cluster-like chain conformation of long polymer chains.39,58 On the contrary, the lower Mn batches showed relatively less aggregation due to weak interactions between the shorter polymer chains.

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Figure 2. Thermal analysis of P(NDI2OD-T2) as a function of molecular weight: (a) Second heating cycles of DSC curves; (b) summary of melting temperature and melting enthalpy. The crystallinity of P(NDI2OD-T2) as a function of Mn was evaluated by differential scanning calorimetry (DSC). The melting temperature (Tm), crystallization temperature (Tc), and melting enthalpy (ΔHm) of P(NDI2OD-T2) are tabulated as a function of molecular weight in Table 1. Figure 2a shows the thermogram of the second heating cycle, and the Tm and ΔHm are plotted in Figure 2b as a function of Mn. As the Mn increased from 15 to 103 kg mol-1, the Tm increased from 282.2 to 325.1 °C, and slightly decreased again to 323.7 °C for the 163 kg mol-1 sample, which agrees well with previous findings.53 The ΔHm values also showed a strong molecular weight dependence, with ΔHm increasing significantly from 4.0 to 17.0 J g-1 as Mn increased from 15 to 48 kg mol-1. These results suggest enhanced crystallinity of the lower Mn ~ 48 kg mol-1, which is ascribed to the smaller contour length of the P-48k polymer forming chain extended crystals with less entangled structures.37,55 On the contrary, further increases in Mn from 48 to 163 kg mol-1 induced a decline of ΔHm to ~ 9.2 J g-1, since the longer 103 and 163 kg mol-1 polymer chains can easily fold, kink and entangle with each other, hindering crystallization kinetics and resulting in disordered structures. Considering that the Tm is dependent on the extended chain length and crystal lamellar thickness,13 we expect that the microstructure of the P(NDI2OD-T2) polymers evolves from fully extended-chain structures 12


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into an entangled, interconnected two-phase “fringed-micelle” morphology between Mn= 48 and 103 kg mol-1. The non-monotonic trend of crystallinity of the NDI-based polymers with Mn is in accordance with previous work.13,37,55,60

Figure 3. Specific viscosity of P(NDI2OD-T2) with different Mn values in (a) toluene and (b) chlorobenzene (5 g L-1) at room temperature.

To better understand the effects of molecular weight on the intermolecular interactions and chain aggregation of P(NDI2OD-T2), we measured the specific viscosity (ηsp) as a function of Mn in toluene and chlorobenzene according to a previously reported procedure.59 It is well known that polymer chains begin to entangle when their Mn is above the critical molecular weight, resulting in a marked increase in melt and solution viscosity. Therefore, a quantitative assessment of the viscosity of polymers as a function of molecular weight is one of most useful measurements to directly identify the existence of polymer entanglement. Since large quantities of polymer are needed to estimate melt viscosity, which would be the most ideal factor for understanding the viscoelastic properties of the polymers, we instead measured polymer viscosity in the solution state at 5 g L-1. The specific viscosity is expressed as ηsp = (η-ηs)/ηs, where, η is the measured viscosity of the polymer solution and ηs is the solvent viscosity. Figure 3 shows the ηsp values as a function of molecular weight measured at room temperature, and the η and ηs values obtained in toluene and chlorobenzene are summarized in Tables S1 13



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and S2. In toluene, the ηsp of the P(NDI2OD-T2) increased marginally from 0.2 for P-15k to 1.4 for P-48k. However, a dramatic increase of ηsp was observed between the P-48k and P103k, from 1.4 for P-48k to 52.3 for P-103k, yielding a 37-fold increase in ηsp with a mere 2fold increase in Mn. This finding indicates the existence of much stronger interactions between polymer chains, and is attributed to evolution of the longer polymer chains into entangled structures as well as the formation of binary hooking contacts one polymer onto two or more polymer chains,61,62 resulting in significantly enhanced viscosity in the solution state. The ηsp of P-163k further increased to 295.7, which is 5 times greater than that of P-103k. The ηsp values in chlorobenzene follow the same trend. The largest jump in ηsp in chlorobenzene from 0.7 to 12.4 was again observed between P-48k to P-103k, with P-163k exhibiting the highest ηsp of 48.2. These results further support the presence of a critical molecular weight between P-48k and P-103k. The relatively lower ηsp values in chlorobenzene are attributed the greater solubility of P(NDI2OD-T2) in chlorobenzene compared to toluene.58

OFET Performance of P(NDI2OD-T2) with Different Mn Table 2. OFET performance of P(NDI2OD-T2) as a function of Mn. Batch

μe, avg (cm2 V-1 s-1)

μe, max (cm2 V-1 s-1)

Ion/Ioff

Vth (V)

P-15k

0.25 ± 0.01

0.28

>104

12.5

P-20k

0.26 ± 0.10

0.40

>103

24.1

P-48k

0.33 ± 0.14

0.52

>104

13.9

P-103k

0.21 ± 0.03

0.24

>105

4.4

P-163k

0.12 ± 0.03

0.16

>105

4.5

To investigate the effect of molecular weight on charge transport in thin films formed from P(NDI2OD-T2), we fabricated OFET devices with a top-gate/bottom-contact geometry 14


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as described in detail in the Supporting Information. The electrical characteristics of the OFET devices of P(NDI2OD-T2) are summarized in Table 2 as a function of molecular weight, and the corresponding transfer and output characteristics are provided in Figure S1 and S2. Guided by previous results showing the device performance to improve with thermal annealing,63 the polymer films were annealed at 220 ˚C for 30 min. While all samples ranging from 15 and 103 kg mol-1 provided sufficiently high electron mobilities (μe > 0.2 cm2 V-1 s-1) for implementation in devices, the average electron mobility (μe) of P(NDI2OD-T2) increased with molecular weight from 0.25 (P-15k) to 0.33 cm2 V-1 s-1 (P-48k). Within this Mn range, the higher μe values measured with increasing Mn are mainly attributed to an increased population of well-ordered crystalline domains, which are beneficial for efficient charge transport.53,55 With a further increase of Mn to 103 kg mol-1, the μe values decreased to 0.21 cm2 V-1 s-1, presumably due to decreased crystallinity. To better understand the Mn dependence on OFET performance, we performed atomic force microscopy (AFM) measurements to investigate the surface topography of the P(NDI2OD-T2) films, which were prepared by same conditions as used for OFET fabrication. Distinct surface morphologies were observed as a function of Mn: As the Mn increased from 15 to 163 kg mol-1, the surface of the polymer film became much smoother with root-mean-square (RMS) surface roughness (Rq) values decreasing from 6.4 to 0.5 nm (Figure S3). The morphology changed most noticeably between P-48k and P-103k; the P-48k thin film exhibited fibrillar ordered domains, different from the considerably blurred and disordered grains seen in the P-103k and P-163k films. Although the fibrillar structures can be beneficial for electron transport, it should be emphasized that the increased amorphous character of the P-103k and P-163k films is crucial to withstanding mechanical stress. Ordered crystal structures, such as those observed in P-48k films, without tie molecules interconnecting the crystalline domains are intrinsically weak and highly susceptible to crack propagation, as will be described next. 15



Mechanical Properties of P(NDI2OD-T2) Thin Films with Different Mn

Figure 4. (a) Result of cohesion energy measurements on P(NDI2OD-T2) thin films as a function of Mn. (b-c) Schematics of the specimen fabrication and the double cantilever beam test set up. AFM images of the P(NDI2OD-T2) thin films as prepared (d-g) and after the DCB test (h-k).

Next, we investigated the mechanical properties of P(NDI2OD-T2) thin films as a function of Mn in terms of the cohesion energy and the tensile properties. First, the cohesion energy of P(NDI2OD-T2) films with different Mn was measured by a double cantilever beam (DCB) test to understand the mechanical resistance of the thin films (Figure 4a-c). The DCB test enables the quantitative measurement of the adhesion and cohesion of various thin films including those formed from conjugated polymers.26,64,65 To separate the effects of thickness on cohesion energy, the thickness of the P(NDI2OD-T2) films was controlled to 100 ± 15 nm for the mechanical tests since the cohesion energy can be expected to increase with thickness due to the greater plastic zone. Fractures occurred cohesively during the fracture energy 16


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measurements, and the fractured surfaces were imaged by AFM. The cohesion energy of P(NDI2OD-T2) films increased with Mn; cohesion energies of 1.06 ± 0.15 J m-2 and 5.07 ± 0.52 J m-2 were measured for the P-15k and P-163k, respectively, as shown in Figure 4a and Table S3. Moreover, we note the sharp increase in cohesion energy observed again between the P-48k and P-103k samples from 1.62 ± 0.04 J m-2 to 4.66 ± 0.53 J m-2. To gain a deeper understanding of the trend in the cohesion energy, the film morphologies of the P(NDI2OD-T2) film before and after the fracture test were compared by AFM (Figure 4d-k). As shown in Figure S3, the formation of fibrillar structures was suppressed for higher Mn polymers, yielding smoother surfaces. Before the fracture tests, decreasing Rq values were observed with increasing molecular weight, from 4.1 for the P-15k film to 0.3 nm for the P163k film. We observed a difference in the surface morphology of the fractured films as a function of Mn, which supports the trend in the cohesion energy. In case of the low-Mn films (P-15k, P-20k and P-48k), brittle cracking occurred by chain pullout.11,48 As a result, the fractured surfaces of low-Mn films possess similar surface morphologies as the films prior to fracture, as seen in Figure 4h,i. In stark contrast, when the Mn increased from 48 to 163 kg mol-1, the ordered domains become less distinct and the interfaces between the ordered domains blurred, consistent with the greater content of amorphous regions, which may include tie molecules and entanglements; accordingly, the P-103k and P-163k films were elongated and torn significantly during the cohesive fracture, leaving a distinct plastic wake on the fractured surfaces of P-103k and P-163k films, as shown in Figure 4j,k. The roughness values (Rq) of the P-103k films increased from 1.2 nm for the film surface before fracture to 94.7 nm for the fractured surface, and those of the P-163k films increased from 0.3 nm to 93.5 nm after the fracture, indicating significant mechanical resistance to fracture. The plastic wakes on the fracture surfaces of the polymer films on the Au (upper) side further support the occurrence of bridging polymer chains and asperities (i.e., plastic deformation) during the cohesive fracture 17



(Figure S4) The cohesion energy was then measured after annealing under the same conditions as used for OFET fabrication (220 °C for 30 min). Thermal annealing resulted in a slight reduction in the cohesion energy of the P(NDI2OD-T2) films (Figure 4a and Table S3), but the overall trend is similar to that observed before annealing. For the low-Mn films, the cohesion energy after annealing was measured to be 0.75 ± 0.02, 0.80 ± 0.10, and 1.35 ± 0.27 J m-2 for the P15k, P-20k, and P-48k films, respectively. Consistent with results before annealing, for the thermally annealed high-Mn films (P-103k and P-163k), the cohesion energy was measured as 3.65 ± 0.42 and 4.68 ± 0.41 J m-2, respectively, significantly higher than that measured for the low-Mn films. However, in case of the low-Mn P(NDI2OD-T2) films, the lower cohesion energies likely result from the lack of tie molecules and entanglements to provide interconnections between domains and release the strain energy. In addition, large aggregates and significantly high peak-to-valley roughness values (Rpv) were observed by AFM for the annealed P-48k films; these aggregates are postulated to act as stress concentrators and weaken the cohesion (Figure S5), and thereby lead to brittle fracture. The cohesion energy of the highMn films also slightly decreased after thermal annealing due to the increased inhomogeneity, but the presence of substantial tie molecules is thought to have maintained the interconnected networks within the film without severe aggregation, leading to high cohesion energy even after thermal annealing.

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Figure 5. (a) Schematics of the specimen fabrication and testing processes for the tensile testing of P(NDI2OD-T2) thin films. (b) Representative stress-strain curves, (c) strain at fracture, and (d) elastic modulus as a function of Mn. Tensile testing was performed in order to understand the stretchability of P(NDI2ODT2) thin films as a function of Mn (Figure 5a) by the pseudo free-standing method.26,66,67 The strain at fracture and elastic modulus are presented in Figure 5c,d, and other data, including the tensile strength and toughness, are presented in Tables S4 and S5. We note that the film thicknesses of all the films are fixed as 100 ± 15 nm because the tensile properties of the polymer thin films can significantly vary depending on the film thicknesses30-32 (Also, see the tensile results of poly(methyl methacrylate) thin films in Figure S6 and Table S6). The strain at fracture showed markedly different behaviors in terms of Mn (i.e., between P-48k and P103k). From the lowest to the highest molecular weight, the strain at fracture increased by a factor of 114 from 0.3 ± 0.1 % for the P-15k films to 34.2 ± 5.8 % for the P-163k films, for an approximately 10-fold increase in Mn. In particular, we observed a jump in the strain at fracture between the P-48k and P-103k films, yielding a 26-fold increase from 1.1 ± 0.2 % (P-48k) to 28.4% ± 3.5 % (P-103k) (Figure 5b and 5c). It is expected that the high-Mn films (P-103k and 19



P-163k) could be much more stretchable when integrated with highly deformable substrates such as PDMS. Since conjugated polymers that can bear more than a few percent of strain without loss of electrical function are regarded to be promising stretchable electroactive materials,3,6,20,68 the greater strains at fracture of P-103k and P-163k are advantageous for flexible and stretchable electronics. Changes in the ductility occurring between the P-48k and P-103k films also indicate the presence of a critical molecular weight at which entangled polymer networks start to form. Accordingly, the toughness values of the high-Mn films (P103k and P-163k) are distinct from those of the low-Mn films (P-15k, P-20k, and P-48k); the toughness values of the P-163k films (10703.6 ± 1265.2 kJ m-3) were 2,820 times greater than those of the P-15k films (3.8 ± 1.0 kJ m-3) (Table S4). Similarly, the sharp transition in the toughness values, from 34.1 ± 7.1 to 5467.9 ± 679.5 kJ m-3, occurred between the P-48k and P-103k films. The stress-strain curves (Figure 5b) revealed the ductility of the films, with significant plastic deformation observed for films formed from polymer with Mn at or over 103 kg mol-1, whereas almost instant brittle fracture was observed in P-48k films, as shown in optical microscopy images acquired before and after the tensile tests and from the Supplementary Movies of the test. Additionally, the elastic modulus of the P(NDI2OD-T2) films increased gradually from 0.40 ± 0.03 GPa to 0.84 ± 0.16 GPa as the Mn increased from 15 to 103 kg mol-1, and then remained constant for the higher Mn sample from P-163k (Figure 5d). This is in good agreement with previous studies that showed the elastic modulus of regioregular P3HT films to depend on the molecular weight before saturating above a certain molecular weight.11,69 In the low-Mn regime, the elastic modulus of the P(NDI2OD-T2) film increased with Mn, allegedly due to enhanced intermolecular interactions, but when the polymer chains were entangled, the elastic modulus predictably saturated since the elastic modulus is determined by short range interactions in the initial elastic regime. As in cohesion testing, tensile testing was performed on thermally annealed films and 20


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the results are shown in Figure 5c,d and Table S5. Annealing was expected to increase the crystallinity of the films, resulting in corresponding increases in elastic modulus and decreases in the strain at fracture. However, the elastic moduli did not change significantly for the highMn films (P-103k and P-163k), but decreased for low-Mn films (P-15k, P-20k, and P-48k). The strain at fracture was found to be slightly reduced overall relative to that measured before annealing, but the trend in terms of Mn is similar: high-Mn films exhibited significantly higher strain at fracture, 25.0 ± 1.8 % for P-103k films and 28.6 ± 1.2 % for P-163k films compared to that of low-Mn films (0.4 ± 0.2 % for P-15k films and 1.0 ± 0.1 % for P-48k films). The decreased elastic modulus of the low-Mn films may be attributed to the increased inhomogeneity of the low-Mn films after annealing as observed in the AFM and optical microscopy images (Figure S5 and S7). Since the increased roughness indicated inhomogeneity over the thickness of the thin films, stress would be concentrated in the thinner region of the films under tension, resulting in lower elastic moduli and strain at fracture.11 However, for the high-Mn films, the surfaces were relatively flat with similar roughnesses as those of unannealed films, likely due to the slower reptation dynamics of the longer polymer chains. Therefore, thermal annealing did not reduce the elastic modulus of the high-Mn films. Further, the superior strain at fracture and toughness of the high-Mn films was maintained even after thermal annealing (Table S5). Taken together, the cohesion energy and tensile tests indicate that the use of P(NDI2OD-T2) with Mn above the critical molecular weight is essential for imparting the high mechanical robustness and resilience required in stretchable electronics.

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Structural Properties of P(NDI2OD-T2) Thin Films with Different Mn

Figure 6. (a) Line cut profiles of GIXS in the in-plane direction, (b) the coherence lengths calculated from (100) and (001) peaks depending on Mn values.

To further clarify the correlation between the mechanical and structural properties of the thin films, the molecular packing structure of P(NDI2OD-T2) in thin films was investigated using grazing incidence X-ray scattering (GIXS) measurements. The films were prepared under the same conditions used to prepare samples for mechanical testing and OFETs. Figure S8 and Figure 6a show the 2D-GIXS patterns of P(NDI2OD-T2) films as a function of Mn, and their corresponding in-plane line cut profiles (out-of-plane line cut profiles are provided in Figure S9). Analysis of the GIXS data, i.e., fitting of the (100), (001) and (010) scattering peaks, is presented in Table S7. As the Mn increased, the (100) lamellar spacing of the P(NDI2OD-T2) polymers increased gradually from 23.8 Å for the P-15k films to 24.5 Å for the P-103k films, and then remained at 24.4 Å for the P-163k films. This finding suggests that the lamellar packing of the polymers loosened with increasing molecular weight as entanglements of the longer polymer chains hindered dense packing, as noted for other systems.60,70,71 Likewise, the π-π stacking distance of the P(NDI2OD-T2) polymers, corresponding to the (010) peak, was also observed to increase gradually from 3.87 Å for the 15k films to 3.95 Å for the 163k films, 22


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also indicating looser intermolecular packing structures. Next, to compare the relative sizes of the crystalline domains in the thin films, the coherence lengths (Lc) as a function of Mn were calculated from the (100), (001) and (010) peaks, respectively, using Scherrer equation (Figure 6b and Table S7).72-74 An increase of the Lc,100 values from 18.7 nm (P-15k) to 26.4 nm (P-48k) was observed, followed by a decrease to 23.3 nm for the P-103k films and a further reduction to 14.3 nm for the P-163k films. Analogously, the Lc,001 and Lc,010 values of the P(NDI2OD-T2) films showed the same trend in terms of Mn. For example, the Lc, 001 increased in the regime between P-15k and P-48k (i.e., 9.6 nm for P-15k films, 14.1 nm for P-20k films, 19.1 nm for P-48k) and, then, reduced to 18.0 nm for the P-103k films and 12.8 nm for the P-163k films. In addition, it is noted that the intensity of the lamellar (200) peak increased with Mn up to 48 kg mol-1, but then decreased for the highest Mn films (P-163k). Additionally, the GIXS results of the thermally annealed thin films prepared under the same conditions as for the OFET fabrication exhibited a similar trend with those of the pristine thin films (Figure S10). The (100) coherence length of thermally annealed films (Lc,100ann) increased significantly with Mn in the low-Mn regime (P-15k (18.7  38.3 nm), P-20k (23.0  36.3 nm) and P-48k (26.4  35.7 nm)) compared with the high-Mn regime (P-103k (23.3  26.2 nm) and P-163k (14.3  21.0 nm)), attributed to the hindered kinetics of longer polymer chains, leading to less developed crystalline structures even after thermal annealing. These different trends in microstructural properties upon thermal annealing correspond well with the AFM and optical microscopy images showing severe aggregates only for low-Mn films from P-20k and P-48k (see Figure S3 and Figure S7). The low-Mn thin films with large Lc,100ann become stiffer upon thermal annealing, indicating deteriorating mechanical properties, which is supported by the reduced toughness and fracture energy measured for these films, as shown in Table S3 and S5, in contrast to the sustained higher fracture energies and strain at fracture of the thermally

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annealed thin films of P-103k and P-163k with higher contents of amorphous domains and entangled chains.

Figure 7. Schematic illustration of the proposed mechanical failure mechanism of P(NDI2ODT2) thin films as a function of Mn values (low-Mn region : 15-48 kg mol-1 and high-Mn region : 103-163 kg mol-1). Based on the combined results of the thermal properties, viscosity measurements, crystalline structures of thin films and mechanical tests, a schematic illustration was proposed for the fracture mechanism of P(NDI2OD-T2) thin films depending on the Mn, as shown in Figure 7. In the low-Mn region (P-15k to P-48k), more ordered and larger crystalline domains formed with increasing molecular weight. However, due to the short polymer chains and absence of inter-domain connectivity, the films are vulnerable to mechanical stress and exhibit substantially lower cohesion energy and strain at fracture than the high-Mn films. On the contrary, in the high-Mn region (P-103k to P-163k), the long polymer chains generate more tie molecules and interchain entanglements, leading to less developed crystalline domains, as shown in Figure 7. As a result, the connectivity between the ordered domains is dramatically enhanced by the tie molecules and entanglements in the high-Mn films, and these connected polymer networks can release substantial strain energy with greater plastic deformation during cohesive fracture. Furthermore, the superior mechanical properties of high-Mn films were 24


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maintained even after thermal annealing, in contrast to the deteriorated mechanical properties of low-Mn films, in which the more developed brittle crystalline domains persist after thermal annealing in the absence of interconnection.

4. CONCLUSIONS In conclusion, the intrinsic thermomechanical, structural, and electronic properties of the representative semi-crystalline conjugated copolymer, P(NDI2OD-T2), were thoroughly investigated as a function of molecular weight. Experiments reveal a critical molecular weight in the range of 48 and 103 kg mol-1 that produces abrupt transitions in thermal, mechanical and viscoelastic properties. This two-fold increase in Mn increased the strain at fracture and toughness by factors of 26 and 160, respectively, without reducing electrical properties. The superior mechanical properties are attributed to the presence of tie molecules and entanglements that can release strain energy with large plastic deformation during crack propagation. Furthermore, the superior mechanical robustness of high-Mn thin films was preserved through thermal annealing, which is essential for optimization of OFET performance. Consequently, the use of conjugated polymers that just exceed the critical molecular weight is crucial to simultaneously optimize electrical performance, mechanical robustness and solution processability for flexible and stretchable organic electronics.

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ASSOCIATED CONTENT Supporting Information. Materials and methods, detail experimental procedures, and additional data. This material is available free of charge via the Internet at http://pubs.acs.org.

AUTHOR INFORMATION Corresponding Authors *

(B. J. Kim) E-mail: [email protected]

*

(T.-S. Kim) E-mail: [email protected]

Author Contributions †

(J.C. and W.K.) These authors contributed equally.

Notes The authors declare no competing financial interest.

ACKNOWLEDGMENTS This research was supported by the National Research Foundation Grant (NRF2017M3A7B8065584, 2012M3A6A7055540 and 2016R1A5A1009926), and the Korea Institute of Energy Technology Evaluation and Planning (KETEP) and the Ministry of Trade, Industry & Energy (MOTIE) of the Republic of Korea (No. 20183010014470), provided by the Korean Government. We thank Profs. R. A. Letteri, G. E. Stein and B. Ma for helpful discussions.

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Importance of Critical Molecular Weight of Semicrystalline n-Type

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