Role of Molecular Weight on the Mechanical Device Properties of

Article pubs.acs.org/Macromolecules

Role of Molecular Weight on the Mechanical Device Properties of Organic Polymer Solar Cells Christopher Bruner and Reinhold Dauskardt* Department of Materials Science and Engineering, Stanford University, 496 Lomita Mall, Durand Building, Stanford, California 94305-2205, United States S Supporting Information *

ABSTRACT: For semiconducting polymers, such as regioregular poly(3-hexylthiophene-2,5-diyl) (rr-P3HT), the molecular weight has been correlated to charge carrier field-effect mobilities, surface morphology, and gelation rates in solution and therefore has important implications for long-term reliability, manufacturing, and future applications of electronic organic thin films. In this work, we show that the molecular weight rr-P3HT in organic solar cells can also significantly change the internal cohesion of the photoactive layer using micromechanical testing techniques. Cohesive values ranged from ∼0.5 to ∼17 J m−2, following the general trend of greater cohesion with increasing molecular weight. Using nanodynamic mechanical analysis, we attribute the increase in cohesion to increased plasticity which helps dissipate the applied energy. Finally, we correlate photovoltaic efficiency with cohesion to assess the device physics pertinent to optimizing device reliability. This research elucidates the fundamental parameters which affect both the mechanical stability and efficiency of polymer solar cells. to ∼2−5 J m−2.8−10 In this research, we demonstrate that active layer cohesion values as high as ∼17 J m−2 can be achieved by increasing the Mw of rr-P3HT and adjusting active layer thickness. This dramatic increase is explained through a greater degree of plastic deformation in the BHJ layer during cohesive failure. For rr-P3HT, increased plasticity is realized through interchain π−π stacking for the crystalline segments increasing the number of van der Waals bonds needed to be broken for chain disentanglement. In addition, more interlamellar amorphous regions with chain folds and tie molecules allow for polymer chain entanglements which would increase frictional molecular pullout during cohesive failure and may even require polymer chain scission.11−13 This provides an opportunity to improve and optimize an important mechanical property that affects the reliability of the solar cell during fabrication and in-service. The cohesive energy, Gc, of the BHJ layer, consisting of rrP3HT and phenyl-C61-butyric acid methyl ester (PC60BM) as an electron acceptor, for each OSC was measured using the four-point bend (FPB) technique, and the results are summarized in Figure 1a.8,10 We note that in a separate study we measured Gc values of 1.9 and 0.5 J m−2 for the homo rr-

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t is well-known that the molecular weight of a polymer can significantly affect its physical, thermomechanical, and electrical properties. For semiconducting polymers, such as regioregular poly(3-hexylthiophene-2,5-diyl) (rr-P3HT), which are used as active layer materials in bulk heterojunction (BHJ) solar cells, the molecular weight (Mw) has been correlated to carrier field-effect mobilities, surface morphology, and gelation rates in solution.1−4 This has important implications for power conversion efficiency (PCE), large-scale manufacturing, and future large-scale applications of polymer-based organic electronics.5,6 For example, state-of-the-art roll-to-roll manufactured rolled sheets of organic solar cells (OSCs) allow for automated placement and replacement in large-scale solar parks.7 Maintaining high values of internal mechanical cohesion along with optimized PCE’s are naturally important for these applications. In this work, we show that the Mw of rr-P3HT in OSCs can significantly change the internal cohesion of the active layer even for relatively small increases in Mw. By measuring the PCE of the resulting OSCs fabricated with rrP3HT of varying Mw, we show how optimized combinations of PCE and cohesion can be achieved. Recent reports on the mechanical reliability of BHJ layers in OSCs have revealed relatively low cohesion and/or adhesion values ranging from ∼1 to 2 J m−2.8−10 By optimizing processing conditions, annealing temperature and time, and thin-film stacking order, cohesion and adhesion could increase © XXXX American Chemical Society

Received: October 25, 2013 Revised: December 19, 2013

A

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being constrained by the stiff top and bottom glass substrates. In addition, studies have shown that the ultimate tensile stress, σUTS, for films of rr-P3HT exhibits an increase from 4 to 10 MPa for Mw of 22 and 102 kDa, respectively.14 The strain to failure was also observed to increase from 13% to 145%, respectively, indicative of increased plastic deformation under tension. The more pronounced sawtooth-like features in the load vs displacement plot (Figure 1b) for OSCs made with ≥53 kDa rr-P3HT are indicative of a time-dependent viscoelastic relaxation process at the crack tip. As loading occurs, a cohesive viscoelastic zone develops at the crack tip. At a critical load, the cohesive zone ruptures and the crack propagates rapidly with a corresponding drop in the applied load (Figure 1c). As the load drops, the crack slows and eventually arrests. On continued loading, the cohesive zone re-forms (along with crack tip blunting) until the instability is again reached and the crack propagates. The time-dependent viscoelastic properties together with the loading rate employed lead to the characteristic sawtooth features. This mechanism has been observed in other polymer systems such as poly(methyl methacrylate) (PMMA).15,16 This is in contrast to OSCs made with 28 kDa rr-P3HT, where we see a relatively flat load plateau indicating pure cohesive failure with no significant plasticity. Nanodynamic mechanical analysis (nDMA) was employed on films of rr-P3HT:PC60BM to provide an indication of the viscoelastic behavior by measuring the storage modulus, E′, loss modulus, E″, and loss tangent, tan δ, for the varying BHJ films at room temperature. From Figure 2, we see that for all Mw BHJ

Figure 1. (a) Comparison of cohesive energy of the BHJ layer (rrP3HT:PC60BM) vs BHJ layer thickness and rr-P3HT Mw. Data for cohesive energy of OSCs with P3HT Mw of 28 kDa was gathered from previously published studies.8 (b) Representative load vs displacement for FPB testing of OSCs. (c) Side view of crack propagating through OSC during FPB testing. A relatively larger plastic zone will inhibit crack growth due to relaxation of applied stress to the BHJ layer whereas a small plastic zone will contribute little to stress relaxation allowing for facile crack propagation.

P3HT (MW ∼ 28 kDa) and PC60BM films, respectively.8 For all devices, mechanical failure occurred within the BHJ layer as confirmed by elemental analysis of the cohesive surfaces with X-ray photoelectron spectroscopy (XPS) which revealed an elemental composition of ∼90% C, ∼8% S, and ∼2% O (Table S3, Supporting Information). It is clear from Figure 1a that at the same BHJ layer thickness the Mw of rr-P3HT had a marked effect on the cohesive energy of the layer. Moreover, particularly at the higher Mw, the cohesive energy was observed to increase markedly with increasing BHJ layer thickness. The significant change in cohesive energy of the BHJ layer with Mw was most likely due to significant plasticity and viscoelasticity at the crack tip. Plastic deformation would also explain the thickness dependence on cohesive energy that we observed for OSCs made with ≥53 kDa rr-P3HT. By increasing the thickness of the BHJ layer, the plastic zone surrounding the crack tip in the BHJ layer can expand to a larger volume before

Figure 2. (a) Storage and loss modulus of films made of rr-P3HT and PC60BM over varying cyclic load frequency taken using nDMA. (b) Tan δ for various films over cyclic load frequency. B

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layers E′ is ∼14−16 GPa, E″ is ∼1−1.5 GPa, and tan δ is ∼0.10. This clearly indicates the viscoelastic nature of the P3HT:PC60BM films, but the measurements do not reveal significant differences with Mw. This is not surprising given the relatively small increases in Mw. Polymers such as polyethylene exhibit changes in tan δ only over large shifts in Mw (0.12 for Mw ∼ 125 kDa to 0.20 for Mw ≥ 3000 kDa).17 Additionally, nDMA is performed in compression whereas cohesive failure occurred in tension, reaching very high strain to failure at the crack tip. Increased molecular connectivity and associated viscoelasticity with increasing Mw would be more apparent under these conditions as revealed by molecular dynamics of polymers under tension.18,19 However, it is interesting to note the increase in storage modulus as we decrease the Mw of rrP3HT. This may be indicative of the more crystalline nature of lower Mw rr-P3HT.20,21 To understand the effect of Mw on viscoelasticity and plasticity which ultimately affect cohesion, we need to understand how it affects the polymer on the molecular level. We know that for rr-P3HT, as Mw increases, the polymer chains are more likely to fold back on themselves, forming crystalline lamellae via π−π stacking.22 Additionally, bridging between ordered crystalline domains of neighboring polymers via tie molecules is more likely to occur, thus increasing polymer chain interconnectivity.21 This bridging occurs more readily with increasing polymer Mw.21,22 Studies utilizing transmission electron microscopy (TEM) and X-ray diffraction (XRD) have confirmed that for rr-P3HT:PC60BM blends an increase in rr-P3HT Mw promotes interconnectivity at the molecular scale between polymer and crystalline domains for many different processing conditions.2,4,22−25 By increasing the degree of polymer chain folding and bridging, the layer becomes more resilient to chain pullout via greater degree of van der Waals interactions. Indeed, studies have shown that low Mw (28 kDa) but with a steady decrease in PCE. This is due to a sharply declining short-circuit current (Jsc) (Table S4, Supporting Information). This has been attributed to charge carrier recombination processes. Because charge carriers have relatively short lifetimes due to a low dielectric constant of the organic polymer matrix, increasing the BHJ layer thickness means fewer carriers can be extracted before recombination.29,30 Below this optimum threshold, with decreasing BHJ layer thickness, we see a decline in PCE due to, again, a decrease in Jsc. However, the decrease in Jsc is due to lower photon harvesting in such thin films reducing overall charge carrier concentration.31 These results provide important insight into the trade-offs between making mechanically robust OSCs with high PCE which must be optimized prior to large-scale manufacturing. In conclusion, we have demonstrated how tuning the Mw of rr-P3HT can greatly alter the cohesive energy in OSCs. In general, increasing the Mw increases intermolecular bonding and raises the probability for chain entanglement allowing for a greater degree of plastic deformation before cohesive failure can occur. Additionally, device performance was investigated, and it was determined that the most mechanically robust OSCs did not necessarily correspond to the most efficient. However, C

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(DTS Delaminator, DTS). The load vs displacement curve for the FPB specimens was then taken and the critical load for stable crack growth, Pc, was then extracted. This allowed us to calculate the cohesive energy, Gc, as the critical value of the strain energy release rate:32,33

Gc =

21Pc 2L2 16b2h3E′

(1)

where E′, b, and h are the specimens plain strain modulus, width, and half-height, respectively. A schematic is given in the Supporting Information (Figure S3). For unstable crack growth (sawtooth pattern observed), Gc was calculated by

Gc =

(2)

where Gi is the Gc value at initiation (crest of sawtooth) and Ga is the Gc value at arrest (bottom of sawtooth).34,35 The cohesion values reported are taken from the average of at least three OSCs. Surface Analysis. The BHJ layer thickness was measured with a profilometer (Dektak 150+ Surface Profiler, Veeco). A scratch was induced on the portion of the OSCs without Ca/Al layer to give a combined thickness of PEDOT:PSS and BHJ layer. The average thickness of the PEDOT:PSS layer was subtracted from this measurement. X-ray photoelectron spectroscopy (XPS) (PHI VersaProbe XPS Microprobe) was performed on cohesive surfaces to identify failure pathways. X-ray beam spot was 100 × 100 μm2 with detection angle of 35°. Elemental composition was determined from resultant spectrum. Noncontact atomic force microscopy (AFM) (XE-70, Park Systems) was used to characterize surface topography of the cohesive surfaces and to obtain root-mean-squared roughness (Rq). Nano-Dynamic Mechanical Analysis. Nano-dynamic mechanical analysis (nDMA) was performed using a Berkovich diamond tip with a nanoindentor (TI 750 Ubi, Hysitron) on specially prepared films of P3HT:PC60BM (∼400 nm, Table S2). Tests utilized a quasistatic load of 50 μN with a dynamic load of 5 μN to get a displacement amplitude of ∼1−2 nm. This lead to a contact depth of less than 100 nm. Frequency sweep tests were ran from 10 to 200 Hz. Tests were performed in ambient conditions, and at least 30 tests were performed on each sample. Current−Voltage Measurements. Photovoltaic performance was assessed via current−voltage measurements taken in a glovebox where the devices were fabricated. OSCs were annealed at 85 °C for 30 min before current−voltage measurements were taken. A source meter (Keithley 2400) and solar simulator (Spectra-Physics 911601000) were used to take current−voltage readings under 1 sun (AM 1.5G) using an NREL certified KG-5 filtered silicon diode.

Figure 4. (a) Cohesion strength of BHJ layer of OSCs with corresponding device PCE. (b) BHJ layer thickness with corresponding device PCE.

relatively high PCEs could be achieved with increases in cohesive energy of the BHJ layer. This presents significant opportunities that will impact large-scale manufacturing as studies have shown that device technologies with relatively low cohesive/adhesive energy (90% (Table S1, Supporting Information). All OSCs were fabricated on top of 50 nm of poly(3,4-ethylenedioxythiophene):poly(styrenesulfonate) (PEDOT:PSS) covered indium tin oxide (ITO, 120 nm) glass substrates. Devices for mechanical testing were built on 50 mm × 50 mm × 0.7 mm ITO/glass substrates. All other devices used 20 mm × 20 mm × 0.7 mm ITO/glass substrates. The BHJ layer, consisting of rr-P3HT and phenyl-C61-butyric acid methyl ester (PC60BM, Solenne), was spun-cast from solutions with the solvent chlorobenzene at 65 °C and allowed to dry for ∼12 h. Solution concentrations and spin-casting parameters are in Table S2 (Supporting Information). Electrodes made of 7 nm of Ca and 100 nm of Al were thermally evaporated onto the BHJ layer. The active area for current−voltage tested OSCs was 0.1 cm2. Adhesion/Cohesion Testing. Solar cells used for cohesion testing were further processed to produce four point bend (FPB) specimens as previously described.8,10 The FPB specimens were loaded under displacement control with a displacement rate of 0.25 μm/s and moment arm, L, of 6.5 mm using a micromechanical testing machine

Polymer batch information, BHJ layer spin-coating parameters, XPS data, solar cell performance, optical/AFM image for fracture surface of 28 kDa P3HT-based organic, compliance vs crack length for four point bend test, and four point bend schematic. This material is available free of charge via the Internet at http://pubs.acs.org.

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected] (R.D.). Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was supported by the Center for Advanced Molecular Photovoltaics (CAMP) under the King Abdullah University of Science and Technology (KAUST) under award KUS-C1-015-21. D

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(32) Dauskardt, R.; Lane, M.; Ma, Q.; Krishna, N. Eng. Fract. Mech. 1998, 61 (1), 141−162. (33) Charalambides, P. G.; Cao, H. C.; Lund, J.; Evans, A. G. Mech. Mater. 1990, 8 (4), 269−283. (34) Mai, Y. W.; Atkins, A. G. J. Mater. Sci. 1975, 10 (11), 2000− 2003. (35) Hoagland, R. G.; Rosenfield, A. R. Int. J. Fract. 1974, 10 (2), 299−302.

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