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Article Cite This: Chem. Mater. 2018, 30, 2102−2111

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Comparative Study of the Mechanical Properties of All-Polymer and Fullerene−Polymer Solar Cells: The Importance of Polymer Acceptors for High Fracture Resistance Wansun Kim,†,∥ Joonhyeong Choi,‡,∥ Jae-Han Kim,§,⊥ Taesu Kim,‡ Changyeon Lee,‡ Seungjin Lee,‡ Mingoo Kim,‡ Bumjoon J. Kim,*,‡ and Taek-Soo Kim*,† †

Department of Mechanical Engineering, Korea Advanced Institute of Science and Technology (KAIST), Daejeon 34141, Republic of Korea ‡ Department of Chemical and Biomolecular Engineering, Korea Advanced Institute of Science and Technology (KAIST), Daejeon 34141, Republic of Korea § Korea Atomic Energy Research Institute, Daejeon 34057, Republic of Korea ⊥ Department of Mechanical Engineering, Korea Advanced Institute of Science and Technology (KAIST), Daejeon 34141, Republic of Korea S Supporting Information *

ABSTRACT: High fracture resistance of polymer solar cells (PSCs) is of great importance to ensure long-term mechanical reliability, especially considering their potential in roll-to-roll printing processes and flexible devices. In this paper, we compare mechanical properties, such as the cohesive fracture energy, elastic modulus, and crack-onset strain, of allpolymer solar cells (all-PSCs) and fullerene-based solar cells (PCBM− PSCs) based on the same, representative low-bandgap polymer donor (PTB7-Th) as a function of acceptor content. The all-PSCs exhibit higher fracture energy (2.45 J m−2) than PCBM−PSCs (0.29 J m−2) at optimized device conditions. Additionally, a 15-fold higher crack-onset strain is observed in all-PSCs than in PCBM−PSCs. Dramatically different mechanical compliances observed for all-PSCs and PCBM− PSCs are investigated in detail by analysis of the blend morphologies as a function of acceptor content (either P(NDI2HD-T) or PCBM acceptors). The superior fracture resistance of all-PSCs is attributed to the more ductile characteristics of the polymer acceptor and the large degree of plastic deformation during crack growth, in contrast to the brittle nature of PCBM and the weak interaction between the polymer-rich phase and highly aggregated PCBM-rich domains. Therefore, this work demonstrates that replacing a small-molecule acceptor (i.e., PCBM) with polymeric materials can be an effective strategy toward mechanically robust PSCs.

1. INTRODUCTION

mechanical failure to occur before photochemical degradation during an outdoor field test.23 Since solar cell devices are inevitably exposed to large mechanical deformation and stress in the manufacturing process and under operating conditions, the long-term mechanical reliability of devices should be guaranteed. All-polymer solar cells (all-PSCs), in which fullerene derivatives are replaced with polymers as the electron acceptor, have emerged as successful candidates that simultaneously achieve high stability and efficiency.19,24−36 Recently, we compared the mechanical-, thermal-, and photostabilities of polymer donor/polymer acceptor blends with polymer/PCBM blends, and found all-polymer blends to exhibit superior mechanical strength, flexibility, and thermal-stability as

The demand to realize wearable and flexible-power-source devices has promoted rapid advances in the bulk-heterojunction (BHJ) polymer solar cell (PSC) field due to their light weight, roll-to-roll processability, and flexibility. Extensive studies have been conducted on fullerene-derivative-based (i.e., phenyl-C61 or C71-butyric acid methyl ester (PCBM)) solar cells (PCBM− PSCs) as electron acceptors, resulting in power conversion efficiencies (PCEs) exceeding 10%.1−9 In addition to high PCE, stability must be guaranteed for the commercialization of PSCs, yet fullerene derivatives have been reported to cause deterioration of device performance under thermal, photo, and mechanical stresses.10−19 In particular, the brittleness of fullerene derivatives can cause mechanical failure of PSCs, and therefore represents a major obstacle to their application in wearable and flexible devices. Other researchers have emphasized the importance of device reliability under external mechanical forces.16,17,20−22 For example, Krebs et al. reported © 2018 American Chemical Society

Received: January 13, 2018 Revised: March 5, 2018 Published: March 8, 2018 2102

DOI: 10.1021/acs.chemmater.8b00172 Chem. Mater. 2018, 30, 2102−2111

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Figure 1. (a) Device structure of PSCs and molecular structures of the polymer donor (PTB7-Th), fullerene acceptor (phenyl-C71-butyric acid methyl ester (PCBM)), and polymer acceptor (poly[[N,N′-bis(2-hexyldecyl)-naphthalene-1,4,5,8-bis(dicarboximide)-2,6-diyl]-alt-5,5′-thiophene] (P(NDI2HD-T)). (b) Schematic illustration of the DCB test and specimen, in which the PSC is sandwiched between two glass substrates using an epoxy. (c) Schematics of specimen preparation and a photograph of tensile test. Dog-bone-shaped specimens were prepared by femtosecond laser patterning on a glass/PSS/active layer.

compared to the PCBM blends.19,37,38 These superior mechanical- and thermal-stabilities of all-PSCs can be mainly attributed to the higher ductility, propensity for chain entanglement, and lower diffusion coefficients of the polymer acceptor.19,39,40 In the recently published paper by Balar et al., the mechanical properties, including the fracture energy, of allPSCs in a fixed donor−acceptor ratio (1:1) were studied in terms of the morphology change at different amounts of solvent-additive.41 However, to date, there have been limited studies of the mechanical properties of all-PSCs, and the relationship between the mechanical and morphological properties of all-PSCs and PCBM−PSCs is still not wellunderstood. Fracture energy is regarded as the macroscopic work required to break bonds and form new fractured surfaces, and is measured in terms of the critical strain energy release rate (Gc in J m−2). Premature fracture can occur in the form of adhesive or cohesive failure without appropriate fracture resistance; thus, higher fracture energy is essential for maintaining mechanical reliability. However, PSC devices consisting of various layers (i.e., electrode, buffer layer, and active layer) undeniably have weak interfaces between materials, and these interfaces can be easily separated under mechanical stress during processing, handling, packaging, and operation. In addition, deformability under tension is important for PSCs considering their attractive potential in flexible and stretchable devices where they are exposed to repetitive bending or stretching during operation. Therefore, it is essential to understand the fracture behavior of PSCs by assessment of mechanical properties, such as the fracture energy, elastic modulus, and crack-onset strain, in conjunction with active layer morphology and device performance. Dauskardt and Lipomi et al. reported that the mechanical properties of PSCs are significantly influenced by various factors (i.e., molecular structure and weight, thermal annealing, processing conditions, relative humidity, polymer−fullerene composition, etc.).16,42−51 While previous efforts have established important guidelines for developing mechanically robust

PSC devices, they have mostly focused on poly(3-alkylthiophenes)-based PCBM−PSCs that have limited PCE value lower than 4−5%. Despite the superior PCEs obtained from low-bandgap polymer-based PSCs, there have been very few reports on the mechanical properties of such devices.52,53 Herein, we investigate the mechanical properties, including cohesive fracture energy, elastic modulus, and crack-onset strain, of all-PSCs and PCBM−PSCs based on the same and well-studied low-bandgap polymer donor, poly[[4,8-bis[5-(2ethylhexyl)thiophene-2-yl]benzo[1,2-b:4,5-b′]dithiophene-2,6diyl]-[3-fluoro-2-[(2ethylhexyl)carbonyl]thieno-[3,4b]thiophenediyl]] (PTB7-Th; Figure 1a). First, the fracture energy was compared for all-PSCs and PCBM−PSCs using the double-cantilever beam (DCB) test (Figure 1b), which is a well-established method for measuring the fracture energy of multilayer thin-film devices.54−56 The all-PSCs exhibited substantially higher cohesive fracture energy (2.45 J m−2) than the PCBM−PSCs (0.29 J m−2) at optimized device conditions. Next, the elastic modulus and crack-onset strain of the active layers of all-PSCs and PCBM−PSCs were measured by tensile testing method (Figure 1c). Compared to PCBM− PSCs, the all-PSCs exhibited substantially higher ductility, with a 15-fold enhancement in crack-onset strain. These dramatic differences in the fracture and tensile properties were unveiled in more detail by correlating the mechanical results with the morphology of active layers for all-PSCs and PCBM−PSCs. Lastly, on the basis of the combined results of the DCB and tensile tests, decohesion mechanisms for all-PSCs and PCBM− PSCs are proposed. The superior fracture resistance of all-PSCs is mainly attributed to the ductility of the polymer acceptor, in contrast to the brittle characteristics of PCBM and its BHJ film. Through systematic investigation of the effect of the acceptor type by tuning the acceptor content, our observations suggest that the replacement of PCBM with a polymeric acceptor would be a powerful and effective strategy to prolong device lifetimes for applications that require stretchable and mechanically robust devices. 2103

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Figure 2. (a) Cohesive fracture energy (Gc) of all-PSC and PCBM−PSC as a function of acceptor content (wt%). The Gc values for the highestperforming all-PSC (43 wt% of P(NDI2HD-T)) and PCBM−PSC (60 wt% of PCBM) were 2.45 and 0.29 J m−2, respectively. (b) Schematics of cohesive fracture paths in the active layer of all-PSCs and PCBM−PSCs.

described in our previous report.60 The number-average molecular weight (Mn) and dispersity (Đ) of the PTB7-Th and P(NDI2HD-T) polymers were 23 kg mol−1 and 2.2, and 38 kg mol−1 and 2.3, respectively, relative to polystyrene standards by size exclusion chromatography. To determine the effect of the acceptor material types (i.e., small molecule or polymer) on the mechanical properties and fracture behavior of devices, we varied the acceptor content in the active layer of both all-PSCs and PCBM−PSCs with normal-type architectures (ITO/ PEDOT:PSS/active layer/LiF/Al). To elucidate the influence of the acceptor types on the mechanical properties of BHJ films, we did not use solvent-additive for fabricating devices that can impact the intrinsic mechanical properties of BHJ films.52 The device efficiencies of both all-PSCs and PCBM−PSCs fabricated with 25, 50, and 75 wt% acceptor in the active layer are summarized in Tables S1 and S2. The acceptor contents of the highest-performing (optimized condition) all-PSCs and PCBM−PSCs were 43 and 60 wt%, respectively. Figure 2a presents the cohesive fracture energy as a function of acceptor content for all-PSCs and PCBM−PSCs prepared under the same device conditions, as well as for PTB7-Th, P(NDI2HDT), and PCBM pristine films. The representative load− displacement curves of the DCB test for all-PSCs and PCBM−PSCs under optimized device conditions are included in the Supporting Information (Figure S1). The active layers remaining on the fractured surfaces were observed by an optical microscope, and this suggested that cohesive fracture occurred within the active layer for both allPSCs and PCBM−PSCs (Figure S2). For the all-PSCs, the debonded surfaces of the ITO side and the Al side both had the blue color of the all-PSC active layer, while both sides of the PCBM−PSCs had the brown color of the PCBM−PSC active layer. In addition, X-ray photoelectron spectroscopy (XPS) measurement of the debonded surfaces revealed that the elemental compositions of the fractured surfaces of both the ITO and Al sides of the devices were comparable to that of the active layer (83.5−85.1 atom % carbon, 6.1−7.4 atom % oxygen, and 5.2−7.4 atom % sulfur) and that they were similar to each other (Table S3). However, the fracture path moved closer to the Al layer of the PCBM−PSCs when the acceptor content exceeded 50% (Figure 2b). The change in the fracture path occurring in the PCBM−PSCs was further verified by depth profiling of the fractured specimens for the PCBM− PSCs (Figure S5). This will be discussed in more detail in the

2. EXPERIMENTAL SECTION DCB Test. The DCB test method was used to measure the fracture energy of the PSCs. The high-precision DCB test equipment (Delaminator Adhesion Test System; DTS Company, Menlo Park, CA) consisted of a linear actuator, a load cell, and loading grips. Normal-type all-PSCs and PCBM−PSCs were prepared as described in the Supporting Information. DCB specimens were fabricated by sandwiching the PSC structures with glass substrates using an epoxy (Epo-Tek 353ND, consisting of bisphenol F and imidazole; Epoxy Technology). The epoxy was cured at room temperature for 72 h in a dry box to exclude any thermal influence on the active layer. All tests were performed under controlled conditions (∼30% RH at 25 °C). During the test, multiple loading/crack-growth/unloading cycles were performed to measure the debond lengths and the fracture energy. The debond length, a, and the applied strain energy release rate, G, were calculated by

⎛ CE′Bh3 ⎞1/3 a=⎜ ⎟ − 0.64h ⎝ 8 ⎠

(1)

12P a ⎛⎜ h⎞ 1 + 0.64 ⎟ 2 3⎝ a⎠ E′B h

(2)

2 2

G=

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, where the slope of the load−displacement curve starts to decrease during the test. Tensile Testing of the Active Layer. Tensile testing of the active layers was performed on a water surface that is an efficient platform to separate the active layers from the substrate and support the specimen during the test.19,52,57−59 The blend solutions of the all-PSCs and PCBM−PSC active layers were prepared under the same conditions as the DCB test samples and were spin-coated on polystyrenesulfonatecoated (PSS-coated) glass substrates. Dog-bone-shaped specimens were prepared using a femtosecond laser pattering technique. Each specimen was gripped by attaching the PDMS-coated Al grips on the specimen gripping areas using van der Waals adhesion.

3. RESULTS AND DISCUSSION Fracture Energy of All-PSCs and PCBM−PSCs. The chemical structures of polymer donor (PTB7-Th) and two different acceptors {poly[[N,N′-bis(2-hexyldecyl)-naphthalene1,4,5,8-bis(dicarboximide)-2,6-diyl]-alt-5,5′-thiophene] (P(NDI2HD-T)) and phenyl-C71-butyric acid methyl ester (PCBM)} are shown in Figure 1a. P(NDI2HD-T) was synthesized by Stille coupling reaction following the procedure 2104

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Figure 3. AFM phase images of all-PSC active layers with (a) 25, (b) 43, (c) 50, and (d) 75 wt% P(NDI2HD-T) and PCBM−PSC active layers with (e) 25, (f) 50, (g) 60, and (h) 75 wt% PCBM, respectively.

following section. The cohesive fracture energies (Gc’s) of the pristine films were measured to be 2.57 ± 0.41 J m−2 for the polymer donor (PTB7-Th), 2.58 ± 0.5 J m−2 for the polymer acceptor (P(NDI2HD-T)), and 0.25 ± 0.09 J m−2 for PCBM. The thickness of active layers for all-PSCs and PCBM−PSCs ranges from 90 to 110 nm. While there have been limited reports for the Gc value of low-bandgap polymer-based PSC systems (i.e., PTB7-Th), the measured Gc values of the pristine donor (PTB7-Th) and acceptor (P(NDI2HD-T)) films are in the range of the previously reported values for pristine semiconducting polymers, such as the well-studied P3HT, 2− 5 J m−2.53,61 In the case of the pristine PCBM films, our results are slightly lower than the previous reported value of 0.5 J m−2,61 but the difference might be primarily attributed to the different device structure of the test specimens and testing method (four-point bending test versus DCB test). As shown in Figure 2a, very similar Gc values of ≈ 2.6 J m−2 were observed for the all-PSCs regardless of the polymer acceptor content and corresponded well with the values of the pristine films, suggesting that blending of two polymers does not significantly alter the ductility or cohesion of the blend films. In stark contrast, a gradual decrease was observed in the cohesive fracture energy of PCBM−PSCs with increasing PCBM content up to 50 wt%, from 2.57 ± 0.41 J m−2 (pristine PTB7-Th) to 1.15 ± 0.13 J m−2 (50 wt% PCBM). At higher PCBM contents, the Gc of PCBM−PSCs decreased sharply to 0.29 ± 0.09 J m−2 at 60 wt% PCBM (the highest-performing device), which is the same as that of the pristine PCBM film. Notably, the cohesive fracture energy of the highest-performing/optimal all-PSCs (2.45 ± 0.3 J m−2) was 9-fold higher than that of the optimal PCBM−PSCs (0.29 J m−2), which highlights the strong fracture resistance of all-PSCs. Over the past few years, different methods, such as annealing, UV radiation, and moisture treatment, have been employed to further improve the low fracture energy of PCBM− PSCs.47−50,62 However, these methods often lower the PCE of solar cells and significantly accelerate the degradation of the device performance. Therefore, our results suggest that the fracture energy of PSCs can be increased substantially without additional treatments by using a polymeric, rather than a smallmolecule (i.e., PCBM), electron acceptor. Blend Morphology Characterization. To further probe the markedly different trends in Gc as a function of acceptor content for all-PSCs and PCBM−PSCs, the blend morphologies were investigated as a function of acceptor content by atomic force microscopy (AFM). For all-PSCs, a finely

separated blend morphology was observed irrespective of the acceptor content, as shown in Figure 3a−d. In addition, we conducted resonant soft X-ray scattering (RSoXS) experiments, which can provide quantitative information about the phase separation of polymer blends.63−65 Figure S3 shows the RSoXS profiles for all-PSCs with different polymer acceptor contents obtained with a resonant photon energy of 287.5 eV, where the largest scattering contrast was observed. The PTB7-Th:P(NDI2HD-T) blend films with different P(NDI2HD-T) contents had very similar profiles with no discernible peaks in the q range 0.004−0.025 Å−1 (corresponding to a length scale of about 20−160 nm), suggesting that the PTB7-Th and P(NDI2HD-T) form blend morphologies without the formation of massively phase-separated domains. These features coincide well with the AFM results and with our previous work.60 In contrast, the blend morphologies of PCBM−PSCs were strongly dependent on PCBM content (Figure 3e−h). The PTB7-Th:PCBM film with 25 wt% of PCBM exhibited a relatively homogeneous morphology, whereas pronounced PCBM aggregation was observed in blend films with 50 wt% PCBM. In films with greater than 50 wt% PCBM, macrophase separation and large PCBM aggregations were observed. These PCBM-dependent morphological changes were also observed by transmission electron microscopy (TEM; Figure S4). The dramatic contrast was observed in the morphology of the all-PSCs and PCBM−PSCs, as exhibited in the Gc value and in the trend of Gc with increasing acceptor content. The higher Gc value for all-PSCs could be mainly attributed to the use of a polymer acceptor having better ductility and interfacial adhesion than the PCBM.19,66−68 The trend of similar Gc values irrespective to the increasing acceptor content correlates well with the results of the morphological characterization of all-PSCs. The relatively homogeneous morphologies of allPSCs imply the presence of intermixing in the amorphous domains of the polymer donor and acceptor, which is attributed in part to the lower diffusion kinetics of polymers.37 The wellmixed regime in the all-polymer blends was observed irrespective to the acceptor content, and hence, we expected the all-PSC blends to exhibit high and similar ductility with the pristine PTB7-Th and P(NDI2HD-T) films. In sharp contrast, the distinct PCBM aggregates and clusters in the PCBM-rich domain of the PCBM−PSCs provide an easier debonding crack pathway between the aggregated PCBM-rich domains and the polymer-rich phase, significantly decreasing the cohesive fracture energy due to the weak interaction between the polar PCBM clusters and the relatively less polar PTB7-Th 2105

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Figure 4. AFM images of the debonded surfaces in (a, b) all-PSCs and (c, d) PCBM−PSCs under optimized device conditions. Rq indicates the rootmean-square roughness of each surface.

Figure 5. Results of tensile testing of all-PSC and PCBM−PSC active layers. (a) Representative stress−strain curves and (b) optical microscopy images of all-polymer and polymer−fullerene active layers under optimized PCE conditions. (c) Elastic modulus and (d) crack-onset strains of allpolymer and polymer−fullerene active layers as a function of acceptor content.

film covered by a polymer-rich skin layer.61,75−78 As a result, for PCBM contents above 50 wt%, the fracture path of the PCBM−PSCs was located closer to the Al electrode side along the weakest path between the polymer-rich phase and the PCBM-rich domains. The lower fracture energy of PCBM− PSCs at these compositions is mainly ascribed to the large amount of PCBM aggregates, which act as debonding sites under mechanical stress (Figure 2).52 Furthermore, SEM images of the debonded surfaces of the ITO and the Al sides of the devices at 60 wt% PCBM clearly indicate that a fracture occurred along the highly aggregated PCBM-rich domains (Figure S6), consistent with our previous observations of the fractured surfaces of PCBM/polymer blend films.52 Debonded Surface Characterization. The debonded surfaces of the fractured specimens prepared under conditions yielding optimal device performance were examined by AFM to investigate the fracture behavior of all-PSCs and PCBM−PSCs. As shown in Figure 4a, b, the high root-mean-square roughness (Rq; 33−35 nm) of the debonded surfaces of the all-PSCs

domains.45,52,69 Even though relatively high PCBM contents are required to generate percolating pathways for electron transport and thereby enhance device performance,70−74 these compositions lead to a severe deterioration in mechanicalstability. In other words, it is apparent that a trade-off exists between electronic and mechanical properties in PCBM−PSCs. For further elucidation of the change in the fracture path of PCBM−PSCs within the active layer, XPS depth profile analysis was conducted on the Al side of the fractured specimen for PCBM−PSCs with 25 and 60 wt% PCBM (Figure S5). The remaining active layer on the Al side after the fracture test was much thinner for 60 wt% PCBM relative to that for 25 wt% PCBM (as made evident by shorter etching time to reach the Al electrode), which suggested that the crack path was located closer to the upper layers (i.e., to the LiF/Al side) with increasing PCBM content. Because of the relatively low surface energy of PTB7-Th compared with that of PCBM, some of PTB7-Th tends to segregate at the film surface to minimize the surface energy difference, resulting in the BHJ 2106

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Figure 6. Schematics of decohesion mechanism of (a) all-PSCs and (b) PCBM−PSCs.

elastic modulus of the all-PSCs (0.68 ± 0.04 GPa) was 4-fold lower than that of PCBM−PSCs (2.64 ± 0.12 GPa). The tensile properties, including elastic modulus and crackonset strain, are presented as a function of acceptor content in Figure 5c, d. For the pristine films of PTB7-Th and P(NDI2HD-T), the elastic modulus and crack-onset strain were determined to be 0.88 ± 0.07 GPa and 13.1 ± 0.8%, respectively, for the polymer donor (PTB7-Th) and 0.62 ± 0.07 GPa and 15.8 ± 1.2%, respectively, for the polymer acceptor (P(NDI2HD-T)). The polymer donor (PTB7-Th) and polymer acceptor (P(NDI2HD-T)), which contain long and flexible alkyl side chains, are likely to produce ductile films.43 The BHJ films of all-PSCs with acceptor contents from 25% to 75% exhibited similar elastic moduli and crack-onset strains as those measured for pristine films of PTB7-Th and P(NDI2HD-T). The PTB7-Th and P(NDI2HD-T) exhibited similar ductile tensile properties, and the BHJ film had relatively small domains of polymer donor and acceptor with the presence of intermixed domain that consisted of amorphous fractions of PTB7-Th and P(NDI2HD-T). Therefore, blending of the two polymers did not significantly alter the ductility of the BHJ films of the all-PSCs. The ductile BHJ films of all-PSCs dissipate a substantial amount of strain energy upon the applied strain, and this resulted in markedly higher mechanical compliance and crack-onset strain. These results agree well with the trends of the cohesive fracture energy. On the contrary, the BHJ films of the PCBM−PSCs became stiffer and more brittle with increasing PCBM content. The BHJ films of PCBM−PSCs showed a maximum elastic modulus of 2.79 ± 0.15 GPa and the lowest crack-onset strain of 0.7 ± 0.2% in the blend containing 75 wt% PCBM; however, a very low crackonset strain of 1.3 ± 0.1% was measured in blends with 50 wt% PCBM as well. These trends as a function of the PCBM content correlate well with the sharp decrease of the Gc for PCBM−PSCs with increasing the PCBM content and its saturation when the PCBM content exceeds 60 wt%. The high elastic modulus of the BHJ films of PCBM−PSCs is attributed to the stiff PCBM-rich phases because of the propensity of PCBM to aggregate.42 However, the high elastic modulus of the PCBM-rich phase may promote stress concentration along the highly aggregated PCBM-rich domains in the strongly phaseseparated BHJ films and lead to the brittle failure without significant plastic deformation.52 As a result, higher elastic modulus of PCBM−PSCs would not necessarily lead to stronger cohesion due to the brittleness of the BHJ films. Decohesion Mechanism of All-PSCs and PCBM−PSCs. On the basis of the combined results of the DCB and tensile tests, proposed decohesion mechanisms of the all-PSCs and

relative to that of the PCBM−PSCs (3−4 nm) clearly indicates the occurrence of plastic deformation during fracture, consistent with the higher Gc values of all-PSCs. Given that the amorphous parts of the polymer donor (PTB7-Th) and polymer acceptor (P(NDI2HD-T)) in all-PSCs were demonstrated to form some intermixing as shown in AFM and RSoXS results, tie chains and interchain entanglements were likely to occur in the BHJ films.79 Accordingly, the polymer chains in the mixed domains of the films can bridge and need to be disentangled (i.e., chain pullout) during cohesive failure, increasing resistance to the fracture.79 In addition, because fracture occurred near the center of the BHJ layer of all-PSCs, the plastic wake zone at the crack tip was not constrained by the stiff adjacent layer (LiF/Al) and instead spread in the BHJ layer, resulting in higher plasticity. Therefore, the greater plastic deformation in the BHJ layer of all-PSCs during cohesive failure likely resulted in the rougher debonded surfaces and higher cohesive fracture energy observed for these devices. In striking contrast, it is apparent, from the smooth debonded surfaces of the PCBM−PSCs with Rq values of 3−4 nm in Figure 4c, d, that plastic deformation was negligible in the PCBM−PSCs. As mentioned in the above discussion, the BHJ films of the PCBM−PSCs with high PCBM content were composed of large PCBM-rich domains covered by a skin layer of polymerrich phase at the surface of the film. Hence, when tensile loading normal to the crack plane is applied during the DCB test, the cohesive fracture occurs easily along the weakest path between the PCBM-rich domains and the polymer-rich phase. As a result, the smooth debonded surfaces were observed for the PCBM−PSCs (Figure 4c, d). Tensile Properties of All-PSCs and PCBM−PSCs BHJ Films. For an investigation of the correlation between the cohesive fracture energy and other mechanical properties such as the elastic modulus and crack-onset strain for all-PSCs and PCBM−PSCs, tensile testing of the active layers was performed as a function of acceptor content. The surface of water was utilized to separate the active layers from the substrate and to perform the tensile test (Experimental Section). The representative stress−strain curves and optical microscopy images acquired during the tensile test (Figure 5a, b) show dramatically different fracture behaviors of all-PSCs and PCBM−PSCs under optimized device conditions. The BHJ films of all-PSCs exhibited ductile crack growth, whereas the BHJ films of PCBM−PSCs showed brittle properties with the formation of a sharp crack through the entire film under tensile strain. Under optimized device conditions, the crack-onset strain of the all-PSCs (15.5 ± 1.2%) was 15-fold higher compared with that of PCBM−PSCs (1.1 ± 0.2%), and the 2107

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PCBM−PSCs are illustrated in Figure 6. The ductile active layer of the all-PSCs undergoes significant plastic deformation as polymer chains are stretched and pulled out, leaving behind a plastic wake on the fractured surface. On the plastic wake, extrinsic mechanisms, such as bridging of polymer chains and contact of asperities, may contribute to the toughening of the all-PSCs.80,81 For all-PSCs, there is likely a significant fraction of amorphous domains where some of the polymer donor and polymer acceptor are intermixed, thereby allowing the formation of finely separated polymer domains. This morphological feature along with the use of polymer donor and acceptor having similar ductile properties may enhance the mechanical compliance of the BHJ films and yield ductile blend films with similar mechanical properties and fracture behavior irrespective of the acceptor content.19,66,67 Accordingly, in the BHJ films of all-PSCs, limited deformation can occur prior to fracture alleviating locally concentrated stresses and allowing the plastic wake zone at the crack tip to spread within the BHJ layer, resulting in a significant amount of plastic deformation during cohesive failure.49,61 These results highlight the importance of all-PSCs for commercial applications in which catastrophic failure would be unacceptable. In contrast, in the active layer of the PCBM−PSCs (Figure 6b), the presence of stiff PCBM aggregates with weak interaction with the polymerrich phase is postulated to cause brittle fracture without significant plastic deformation, resulting in lower fracture resistance. As the PCBM content in the blends increased, more PCBM clusters were observed, and the BHJ films of PCBM−PSCs became stiffer and more brittle. In addition, the plastic wake zone at the crack tip is likely to be confined to a very small region. Accordingly, the PCBM−PSC devices with optimal device conditions yield very low cohesive fracture energy and crack-onset strain. The sharp decrease in the fracture energy of PCBM−PSCs with increasing PCBM significantly undermines the mechanical reliability of the devices and is one of the main obstacles to the commercialization of PCBM−PSCs.

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.chemmater.8b00172. Additional experimental procedures and figures including representative load−displacement curves, optical images, RSoXS profiles, TEM images, XPS depth profiles, and SEM images (PDF)

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AUTHOR INFORMATION

Corresponding Authors

*E-mail: [email protected]. *E-mail: [email protected]. ORCID

Bumjoon J. Kim: 0000-0001-7783-9689 Taek-Soo Kim: 0000-0002-2825-7778 Author Contributions ∥

W.K. and J.C. contributed equally.

Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This research was supported by the Basic Science Research Program (2015R1A1A1A05001115, 2016R1E1A1A02921128, 2015M1A2A2057509) and Wearable Platform Materials Technology Center (2016R1A5A1009926) funded by the National Research Foundation under the Ministry of Science, ICT and Future Planning.

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REFERENCES

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4. CONCLUSION In summary, we compared the fracture behavior and mechanical properties of all-PSCs and PCBM−PSCs for two different acceptor types as a function of acceptor content by measuring cohesive fracture energy, elastic modulus, and the crack-onset strain. We observed that the all-PSCs showed much higher Gc values of 2.45 J m−2 than PCBM−PSCs (0.29 J m−2) under optimized device conditions. While the all-PSCs had similar Gc values of ≈ 2.6 J m−2 irrespective of polymer acceptor content, the Gc values of PCBM−PSCs were strongly dependent on the amount of PCBM. More importantly, allPSCs exhibited 15-fold superior crack-onset strains compared to PCBM−PSCs. We demonstrated that the contrast in mechanical compliances correlated well with different morphological behaviors. Even though the PTB7-Th:P(NDI2HD-T) blends consisted of two different polymers, they did not show large phase-separated domains irrespective to the acceptor content, which yielded ductile blend films along with the use of polymer donor and acceptor having similar ductile properties. Contrary to the all-polymer blends, cracks propagated easily in the PCBM blends along the weak path between the polymerrich phase and the highly aggregated brittle PCBM-rich domains. Taken together, these results suggest that substituting small-molecule PCBM acceptors with polymeric acceptors is an effective strategy to realize mechanically robust PSCs. 2108

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