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Water-based nanoparticulate solar cells using a diketopyrrolopyrrole donor polymer† Ben Vaughan,ab Evan L. Williams,*c Natalie P. Holmes,a Prashant Sonar,c Ananth Dodabalapur,d Paul C. Dastoora and Warwick J. Belcher*a Organic photovoltaic devices with either bulk heterojunction (BHJ) or nanoparticulate (NP) active layers have been prepared from a 1 : 2 blend of (poly{3,6-dithiophene-2-yl-2,5-di(2-octyldodecyl)-pyrrolo[3,4-c]pyrrole-1,4-dione-alt-naphthalene}) (PDPP-TNT) and the fullerene acceptor, ([6,6]-phenyl C71-butyric acid methyl ester) (PC70BM). Atomic force microscopy (AFM) and scanning electron microscopy (SEM) have been used to investigate the morphology of the active layers of the two approaches. Mild thermal treatment of the NP film is required to promote initial joining of the NPs in order for the devices to function, however the NP structure is retained. Consequently, whereas gross phase segregation of the active layer occurs in the BHJ device spin cast from chloroform, the nanoparticulate approach retains
Received 5th November 2013, Accepted 6th December 2013
control of the material domain sizes on the length scale of exciton diffusion in the materials. As a result, NP devices are found to generate more than twice the current density of BHJ devices and have a
DOI: 10.1039/c3cp55037k
substantially greater overall efficiency. The use of aqueous nanoparticulate dispersions offers a promising approach to control the donor acceptor morphology on the nanoscale with the benefit of
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environmentally-friendly, solution-based fabrication.
Introduction In recent times major advances in the performance of bulk heterojunction organic photovoltaic (BHJ OPV) devices have been achieved through the development of new classes of donor materials.1 Consequently, device power conversion efficiencies (PCEs) have now exceeded 10% and continue to increase rapidly.2 Significantly, OPV device efficiencies have doubled twice in the last 10 years and limits of 20–24% are now being proposed as an achievable limit in single junction devices.3 One of the most popular recent approaches to the development of high performance polymers for OPV applications is the design of donor–acceptor (D–A) copolymers.4,5 In these polymers the energy levels, and thus band gap of the polymer, may be tuned by modifications of the D and A units and the absorption of the polymer subsequently red-shifted and broadened.6 Polymers based upon diketopyrrolopyrrole (DPP) as the acceptor unit have a
Centre for Organic Electronics, University of Newcastle, Callaghan, NSW, 2308, Australia. E-mail: [email protected] b CSIRO Energy Technology, P.O. Box 330, Newcastle, 2300, Australia c Institute of Materials Research and Engineering (IMRE), Agency for Science, Technology, and Research (A*STAR), 3 Research Link, Singapore 117602. E-mail: [email protected] d Microelectronics Research Center, University of Texas at Austin, Austin, TX, 78758, USA † Electronic supplementary information (ESI) available. See DOI: 10.1039/ c3cp55037k
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proven particularly interesting and successful.7–10 This interest is largely due to the polymer’s broad optical absorption and the excellent planarity of the unit resulting in strong intermolecular p–p interactions and high charge carrier mobility.11 Unfortunately these interactions also typically lead to highly crystalline materials making the processing of said polymers and the formation of optimal BHJ active layer morphologies difficult.12 Small molecule BHJ OPV devices based on DPP-based donors have also shown promise and device PCEs of 4.4% have been achieved.8,13 A number of methods have been explored to overcome the processing problem associated with the use of DPP-based polymers. These include the synthetic modification of oligomers to form diketopyrrolopyrrole–oligothiophene:fullerene dyads and likewise the formation of thiophene-based D–A copolymers wherein the DPP acceptor constitutes only a minor (up to 15%) fraction of the polymer to ensure processability.14,15 A second approach involves the use of mixed solvent processing and the addition of processing additives to enhance the film morphology of unmodified DPP-based polymer.16 Using these techniques, DPP-based polymer:fullerene BHJ OPV devices with well blended morphologies have been achieved and device PCEs as high as 6.7% have been observed; mirroring the small molecule results.17 A further refinement through the use of solubilizing side chains on furan-flanked DPP (rather than thiophene) copolymers has led to very finely blended films which give devices with PCEs in excess of 6%.18 Importantly though, all of these methods achieve the optimal D–A blend morphology through the use of toxic or high
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boiling point solvents or by introducing considerable additional complexity to the polymer synthesis. Additionally, there are questions regarding the feasibility of employing mixed solvent systems and precise control of the solvent evaporation rate in large-area, roll-to-roll processing which would be employed in large scale, industrial fabrication of OPV devices. The recent development of aqueous dispersions of organic semiconductor nanoparticles offers a new way of addressing the issue of polymer processability for the fabrication of organic solar cells.19 This approach holds great appeal since it is compatible with organic semiconductor materials that have been developed for and utilized in organic solar cells fabricated by more traditional, organic solventbased fabrication techniques. In other words, no additional material design or synthesis is necessary to change the solubility properties of the materials. Additionally, the characteristic morphology of blended semiconductor nanoparticles fabricated in this approach is core–shell with phase segregation determined by the relative surface energies of the component materials.20,21 Consequently, this approach typically results in the formation of nanoparticles which demonstrate a fullerene-rich core and polymer-rich shell morphology.21,22 This intrinsic structure offers an elegant way to control the blend morphology on the nanometer length scale, driving the formation of domain sizes in the active layer close to the ideal for nanoparticles in the 30 nm diameter range. We describe the fabrication and characterization of both BHJ and NP OPV devices made from 1 : 2 blends of (poly{3,6-dithiophene-2-yl-2,5-di(2-octyldodecyl)-pyrrolo[3,4-c]pyrrole-1,4-dione-altnaphthalene}) (PDPP-TNT) and the fullerene acceptor, ([6,6]-phenyl C71-butyric acid methyl ester) (PC70BM). Previous PDPP-TNT: PC70BM BHJ devices have required a mixed solvent system (chloroform : o-dichlorobenzene, 4 : 1 by volume) to produce an advantageous active layer morphology.23 Here, we show that whereas deposition of the active layer from pure chloroform results in a coarse, phase segregated morphology with domain sizes close to the micron scale, the use of PDPP-TNT:PC70BM blend nanoparticles controls the active layer morphology on the nanometer scale. Consequently, device performance increases from an optimized PCE of B1.0% to a PCE of B2.0% through an increase of Jsc. Thus, this work demonstrates that aqueous-based NP dispersions can indeed be employed as a means to preorganize the OPV active layer blend morphology without resorting to synthetic modification of the polymer or the use of toxic or high boiling point solvents. Due to the fact that the donor and acceptor species are forced into close proximity and the blend morphology is fixed during the formation and solidification of the nanoparticles, the blend morphology is insensitive to the solvent evaporation rate during the film formation process. This trait offers promise for more facile translation from lab-scale optimization to large-scale fabrication.
Experimental PDPP-TNT was synthesised as described by Sonar et al.24 An aqueous dispersion of semi-conducting PDPP-TNT : PC70BM nanoparticles (1 : 2 blend) was prepared using the
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miniemulsion technique as outlined previously.20 Dynamic light scattering (Zetasizer Nano-ZS, Malvern Instruments, UK) was used to measure the distribution of particle sizes in the aqueous dispersion and gave a peak 1 particle size of 35 nm. Poly(3,4-ethylenedioxythiophene) poly(styrenesulfonate) (PEDOT: PSS) (Clevios P) films were spin coated (4000 rpm) on pre-cleaned patterned indium tin oxide (ITO) glass slides and annealed at 140 1C for 30 minutes to eliminate water in the films. The PDPP-TNT: PC70BM nanoparticle layer was deposited by spin coating 35 ml of the dispersion (1750 rpm for 1 minute) in air to give a film of 115 5 nm thickness. Following the deposition of the nanoparticle layer the films were dried at 110 1C, 130 1C or 150 1C for 5, 10, 30 or 60 minutes. Bulk heterojunction layers were spin cast from a PDPPTNT : PC70BM blend solution (1 : 2, 12 mg ml 1 in CHCl3) at 2000 rpm for 1 minute in a nitrogen-filled glovebox to give a 125 5 nm thick active layer. The films were then transferred into a vacuum chamber for cathode evaporation. The calcium/ aluminium (Ca/Al) electrodes were evaporated on the active layers in vacuum (2 10 6 Torr). The thickness of the Ca and Al layers were measured to be about 20 nm and 100 nm, respectively, using a quartz crystal monitor. After evaporation, fabricated devices were tested without further annealing. Current density–voltage ( J–V) measurements were conducted using a Newport Class A solar simulator with an AM1.5 spectrum filter. The light intensity was measured to be 100 mW cm2 by a silicon reference solar cell (FHG-ISE) and the J–V data were recorded with a Keithley 2400 source meter. Individual devices were masked (illuminated area of 3.9 mm2) to eliminate edge effects. External quantum efficiency measurements were recorded when the devices were irradiated with light from a tungsten halogen lamp passed through an Oriel Cornerstone 130 monochromator, using an Ithaco Dynatrac 395 analogue lock-in amplifier and Thorlabs PDA55 silicon diode to collect the reference signal, and a Stanford Research Systems SR830 DSP digitising lock-in amplifier to measure device current. An ultraviolet-visible absorption spectrophotometer (UV-vis, Varian Cary 6000i) was used to study the absorption of PDPPTNT:PC70BM nanoparticulate and bulk heterojunction films. For the UV-vis characterisation, films were spin coated onto quartz slides. Differential scanning calorimetry was performed on a Shimadzu DSC-60A, samples were scanned at 10 1C min 1 from 30 1C to 350 1C. Samples of the pristine PDPP-TNT and a 1 : 2 blend of PDPP-TNT : PC70BM nanoparticles (prepared by drop casting the nanoparticle ink directly into the aluminium pan and allowing to air dry) were recorded. Glass transition points (Tgs, taken at the onset) were observed at 84.4 1C for the pristine polymer and 87.9 1C for the polymer:fullerene nanoparticles. Atomic force microscopy (AFM) images were collected using an Asylum Research Cypher in AC mode. SEM was performed on a Zeiss Sigma ZP FESEM at accelerating voltages of 1–3 kV, and magnification ranges of 10 000 to 100 000. Samples were prepared for scanning electron microscopy (SEM) by spin coating 2.6 ml of the nanoparticulate dispersion, diluted 1/10 with water, (3000 rpm, 1 minute, low acceleration of 112 rpm s 1) onto a conductive silicon substrate.
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Results and discussion A series of 1 : 2 PDPP-TNT : PC70BM BHJ and NP OPV devices were prepared and characterised in order to determine the effect that preorganisation of the active layer morphology via the NP approach has on film morphology and device performance. BHJ blend films were prepared from a chloroform solution and used as-spun since heat treatment of these films was shown to have little effect upon device performance; NP films were dried before cathode deposition. Similar thermal treatment has been known to promote particle joining and to enhance subsequent device performance.22 In order to optimise this drying step, temperatures of 110 1C, 130 1C and 150 1C and drying times of 5 min, 10, min, 30 min and 60 min were investigated. Fig. 1 shows the normalised thin film UV-vis spectra of 1 : 2 blends of PDPP-TNT : PC70BM as-spun from chloroform and from the nanoparticulate dispersion. Also shown are the spectra of thin films of both the pristine component materials. It is clear that the blend spectra are composed of the distinguishing features of both the polymer and fullerene. AFM images of thin (B120 nm) films of both a BHJ film spin-coated from chloroform and a NP as-spun thin blend film are shown in Fig. 2 and a number of features are apparent. The BHJ film shows large B200–500 nm diameter domains which extend above the surface of the bulk film by approximately 30 nm. Previous work has shown that the regions between these domains are sulphur-rich indicating that the film is comprised of fullerene-rich plateaus surrounded by polymer-rich valleys.23 Similar domain growth is observed for films containing PPVbased materials and is due to the relative solubilities of the two components.25 In contrast, AFM images of the deposited PDPP-TNT:PC70BM NP films shows that they retain their discrete
Fig. 1 UV-vis spectra of an as spun PDPP-TNT : PC70BM 1 : 2 nanoparticulate film (solid line), PDPP-TNT : PC70BM 1 : 2 bulk heterojunction thin film (thick dashed line), pure PDPP-TNT thin film (light dotted line) and pure PC70BM thin film (light dashed line). All spectra have been normalized to the highest peak in the 725–750 nm range except for the PC70BM film which has been normalized at l = 300 nm.
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Fig. 2 AFM images of an as-spun PDPP-TNT : PC70BM 1 : 2 blend film spincast from chloroform (a) and an as-spun PDPP-TNT : PC70BM 1 : 2 nanoparticulate film (b). Scale bars are 500 nm.
nanoparticulate nature and size, with NPs of approximately 35 nm diameter (consistent with the measured DLS results for the aqueous dispersion) observed as a continuous film. The NP approach can be seen to offer an initial advantage over the traditional single-solvent system, providing finer scale mixing of the donor and acceptor and greater interfacial area. Fig. 3 shows the J–V characteristics for the as-spun PDPPTNT : PC70BM 1 : 2 NP OPV devices and those heated at 110 1C, 130 1C and 150 1C for up to one hour. It is clear that all of the devices require a heating step to become functional, with devices heated at all temperature and time combinations showing vastly superior performance in both the J–V and EQE plots. Heating the devices at 110 1C for even prolonged periods has little effect on device performance beyond the initial performance boost. By contrast, heating the devices at 150 1C leads to a systematic decrease in both the Voc and Jsc, and therefore the efficiency, of the devices. Heating at 130 1C bridges the results of the two outlying temperatures and initially results in an increase in current (and efficiency) and then a systematic decrease in both Voc and Jsc (and efficiency). Thus, from the data shown in Fig. 3a–c there are broadly three states for these devices within the explored regime of heating conditions. As-spun devices barely perform and display very poor device parameters. Moderate heating (110 1C for any length of time or 130 1C for less than half an hour) results in dramatic device improvement giving devices with Jsc B5–6 mA cm 2, Voc B 0.77 V, and PCE between 1.5 and 2.0%. Significantly, all of these devices show current densities (and efficiencies) much superior to the corresponding BHJ device. These observations are consistent with the measured Tgs of 85 1C for the pristine polymer and 88 1C for the polymer:fullerene blend nanoparticles (see ESI†), which indicates that for all three temperatures investigated substantial mobility of the components would be expected. However, excessive annealing of the devices (130 1C for >10 min and 150 1C for any length of time) results in a reduction in performance and gives rise to devices with Jsc B 2–3 mA cm 2, Voc B 0.60 V, and PCE between 0.5 and 1.0%. All devices display FFs of approximately 0.4, which is B30% lower than the BHJ devices. We attribute this reduction to the surfactant present in the nanoparticulate active layer; we have separately observed a systematic increase in FF for PDPP-TNT : PC70BM 1 : 2 nanoparticulate OPV devices as the amount of surfactant used in the NP dispersion is decreased (data not shown).
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Fig. 4 SEM images of as-spun (a), heated at 110 1C for 10 minutes (b), heated at 130 1C for 10 minutes (c), and heated at 150 1C for 10 minutes (d), PDPP-TNT : PC70BM 1 : 2 nanoparticulate films. All scale bars are 200 nm.
Fig. 3 Current density versus voltage curves for PDPP-TNT : PC70BM 1 : 2 nanoparticulate OPV devices heated at 110 1C (a), 130 1C (b) and 150 1C (c) for 0 minutes (dotted line), 5 minutes (short dashes), 10 minutes (medium dashes), 30 minutes (long dashes) and 60 minutes (solid line).
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The observed device performance may be explained by the observation of changes in the morphology of the active layer upon heating. Fig. 4 shows scanning electron microscopy (SEM) images of an as-spun PDPP-TNT : PC70BM 1 : 2 nanoparticulate film, as well as films that have been heated at 110 1C, 130 1C and 150 1C for 10 minutes. The image of the as-spun film (Fig. 4a) shows that the PDPP-TNT:PC70BM NPs are ovoid, rather than spherical, in nature. The axis dimensions of the NPs (major axis B42 nm, minor axis B23 nm) compare well with the particle size determined from DLS for the NP suspension (35 nm) and indicate that the particles have an aspect ratio of 9 : 5. Upon heating of the film at 110 1C for 10 minutes the particles appear to swell slightly in size and join together (Fig. 4b). The joining of the particles is further pronounced upon heating of the film at 130 1C for 10 minutes but discrete particles are still clearly retained (Fig. 4c). By contrast, heating the film at 150 1C for 10 minutes results in the particles joining completely and the particulate structure of the film vanishing (Fig. 4d). In addition, apparently crystalline aggregates appear on the film surface, consistent with gross phase segregation of the fullerene from the polymer matrix.21 These images strongly support the device results, showing that the improvement in device performance upon initial heating of the films results from joining of the NPs whilst retaining the overall nanoparticulate structure. The improved connectivity within the device necessarily results in much improved transport and therefore improved device performance. Heating the film beyond 130 1C (or extended heating at lower temperatures) leads to gross phase segregation of the NP film. Both charge generation and transport are now adversely affected and device performance consequently drops. From Fig. 3, we see that heating the devices at 130 1C results in the optimal device performance. Table 1 compares the device characteristics of the 1 : 2 PDPP-TNT : PC70BM BHJ OPV device and a series of 1 : 2 PDPP-TNT : PC70BM NP OPV devices dried at 130 1C for different times, all of similar thickness.
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Table 1 Comparison of device characteristics for the 1 : 2 PDPPTNT : PC70BM BHJ device and 1 : 2 PDPP-TNT : PC70BM NP OPV devices heated at 130 1C for varying times. The series resistance (Rs) and shunt resistance (Rsh) values were calculated by taking the inverse slope of the current–voltage curve near Jsc and Voc, respectively
BHJ Device NP Device (as spun) NP Device (5 min dry) NP Device (10 min dry) NP Device (30 min dry) NP Device (60 min dry)
Voc (V)
Jsc (mA cm 2)
FF
0.82 0.38 0.77 0.76 0.61 0.58
2.03 0.55 5.28 6.09 3.12 2.94
0.63 0.25 0.42 0.43 0.41 0.42
PCE (%)
Rs (O)
1.05 606 0.05 17 000 1.71 827 1.99 532 0.78 1160 0.70 1260
Rsh (O) 79 600 17 900 8250 7130 10 600 12 300
DPP-TNT-based BHJ devices fabricated from a single solvent (chloroform) exhibit a PCE significantly lower than that observed for active layers deposited via a mixed solvent system.24 The Voc and FF of devices prepared from the single and mixed solvent systems are essentially identical,23 indicating that the difference in performance lies solely in the generated photocurrent. Indeed, the Jsc of the chloroform processed device is a mere one quarter of that of the Jsc observed using the mixed solvent system.24 In light of the coarse phase segregation in the chloroform processed active layer observed here and previously,23 the poor performance of the subsequent device can be explained by inefficient charge separation of excitons and therefore high geminate recombination in the active layer. Considering the device parameters for the NP OPV devices heated at 130 1C, there are a number of observations. Firstly, the as-spun device shows reduced characteristics for all parameters and can best be described as barely functioning. This observation is consistent with the findings of previous studies of NP OPV devices and thin films which show that initial heating of the NP film promotes joining of the NPs and consequently improves device performance.22 This observation is supported by the abrupt large decrease in Rs that occurs upon initial heating of the NP film. After the initial 5 minutes, subsequent heating results in a systematic decrease in the Voc of the device. FF remains invariant across the heating regime, however, Jsc and PCE are maximized after 10 minutes of annealing and then drop rapidly. The EQE plots for devices heated at 130 1C (Fig. 5) are in good agreement (5%) with the measured Jsc values, reinforcing the observation that the changes in device performance upon heating arise primarily as a result of changes in current generation. In addition, the shape of the EQE plots shows that the relative contribution of the polymer varies greatly with thermal treatment. In optimally treated devices the long wavelength peak (B730 nm, polymer absorption peak) is significantly higher than at B500 nm (minimum of polymer absorption). For extended treatment, the EQE contribution at 500 nm is larger than at 730 nm. The change in shape of the spectral response suggests that excitons created upon photoexcitation of the polymer are more efficiently dissociated and transported to the electrodes after optimal thermal treatment. This improvement could be due to improved exciton diffusion and/or a rearrangement of the blend components within the
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Fig. 5 EQE spectra for the as-spun PDPP-TNT : PC70BM 1 : 2 nanoparticulate OPV device (dotted line) and those heated at 130 1C 5 minutes (short dashes), 10 minutes (medium dashes), 30 minutes (long dashes) and 60 minutes (solid line).
nanoparticle. Previous work has shown that PDPP-TNT is amorphous at room temperature but rearranges to become increasingly crystalline following annealing above 100 1C;24 consistent with our observation of a Tg at 85 1C. Given that the NP shell is thought to be polymer-rich,20 the increase in device performance from 5–10 minutes can best be explained by the increased polymer crystallinity resulting in improved charge transport.24 Subsequent annealing, however, results in expulsion of PCBM from the ordered polymer matrix (Fig. 4d), as has been observed in the P3HT:PCBM system,21 resulting in suboptimal blend composition and a corresponding reduction in photocurrent generated from the PDPP-TNT polymer. Fig. 6 compares the J–V and EQE plots for the most efficient PDPP-TNT : PC70BM 1 : 2 nanoparticulate OPV device and the corresponding BHJ device spin coated from chloroform. It is clear from these plots that the NP device substantially outperforms the BHJ device. The EQE spectrum of the NP device shows significant contribution from the long wavelength region, attributed to polymer absorption. Previous studies have shown the thickness of the BHJ device to influence the shape of the EQE spectrum;23 in this study, attempts were made to maintain a comparatively equal thickness for the BHJ and NP devices. As mentioned, the large-scale phase segregation (hundreds of nanometers) observed for the BHJ film (Fig. 2a) is non-optimal for charge generation in these devices.26 Indeed, although the BHJ device exhibits a Voc larger than the optimal NP device it has a current density comparable to that of the overannealed NP device and consequently a much reduced efficiency. By comparison, the NP film retains a phase segregated morphology in which the domain sizes are constrained to tens of nanometers and thus are much better suited to exciton dissociation. As a result, the photocurrent of the optimised NP device is more than double that of the BHJ device.
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or solvent additives are typically employed to promote mixing and optimize the blend morphology. The NP approach offers a means to control the initial, intrinsic blend morphology independent of the evaporation time during spin coating. This control and the solvent-free nature of the aqueous dispersions are highly beneficial traits for the development of large-scale fabrication of OPV devices.
Acknowledgements Special thanks to the University of Newcastle Electron Microscopy and X-ray Unit. Thanks to Mr Nicolas Nicolaidis for his assistance with EQE measurements. The University of Newcastle and the Australian Renewable Energy Agency (ARENA) are gratefully acknowledged for PhD scholarships (NH). The University of Newcastle is gratefully acknowledged for a Visiting Fellowship (ELW). Authors ELW and PS acknowledge the Visiting Investigatorship Programme (VIP), project number 0721100037, of the Agency for Science, Technology and Research (A*STAR), Republic of Singapore. This work was performed in part at the Materials node of the Australian National Fabrication Facility, a company established under the National Collaborative Research Infrastructure Strategy to provide nano and microfabrication facilities for Australia’s researchers.
References
Fig. 6 Current density versus voltage curves (a) and EQE plots (b) for a 1.0% efficient PDPP-TNT : PC70BM 1 : 2 bulk heterojunction device deposited from chloroform (solid line) and the optimal 2.0% efficient PDPP-TNT : PC70BM 1 : 2 nanoparticulate OPV device heated at 130 1C for 10 minutes (dotted line).
Conclusions We have demonstrated that the use of NP films in PDPPTNT:PC70BM OPV devices provides control over the active layer morphology on the scale of typical exciton diffusion lengths in organic semiconductor materials. The resultant constrained morphology allows for improved exciton harvesting within the device and therefore an increase in photocurrent generation over comparable BHJ devices. Furthermore, we have shown that careful heating of the PDPP-TNT:PC70BM NP active layer is required to optimize device performance. Initial heating is required to promote connectivity of the NPs in order to improve charge transport through the film. However, over-annealing results in gross phase segregation within the NP film and a subsequent reduction in device performance. The NP approach shows great promise in donor–acceptor systems where largescale phase segregation is observed and mixed solvent systems
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Water-based nanoparticulate solar cells using a diketopyrrolopyrrole donor polymer† Ben Vaughan,ab Evan L. Williams,*c Natalie P. Holmes,a Prashant Sonar,c Ananth Dodabalapur,d Paul C. Dastoora and Warwick J. Belcher*a Organic photovoltaic devices with either bulk heterojunction (BHJ) or nanoparticulate (NP) active layers have been prepared from a 1 : 2 blend of (poly{3,6-dithiophene-2-yl-2,5-di(2-octyldodecyl)-pyrrolo[3,4-c]pyrrole-1,4-dione-alt-naphthalene}) (PDPP-TNT) and the fullerene acceptor, ([6,6]-phenyl C71-butyric acid methyl ester) (PC70BM). Atomic force microscopy (AFM) and scanning electron microscopy (SEM) have been used to investigate the morphology of the active layers of the two approaches. Mild thermal treatment of the NP film is required to promote initial joining of the NPs in order for the devices to function, however the NP structure is retained. Consequently, whereas gross phase segregation of the active layer occurs in the BHJ device spin cast from chloroform, the nanoparticulate approach retains
Received 5th November 2013, Accepted 6th December 2013
control of the material domain sizes on the length scale of exciton diffusion in the materials. As a result, NP devices are found to generate more than twice the current density of BHJ devices and have a
DOI: 10.1039/c3cp55037k
substantially greater overall efficiency. The use of aqueous nanoparticulate dispersions offers a promising approach to control the donor acceptor morphology on the nanoscale with the benefit of
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environmentally-friendly, solution-based fabrication.
Introduction In recent times major advances in the performance of bulk heterojunction organic photovoltaic (BHJ OPV) devices have been achieved through the development of new classes of donor materials.1 Consequently, device power conversion efficiencies (PCEs) have now exceeded 10% and continue to increase rapidly.2 Significantly, OPV device efficiencies have doubled twice in the last 10 years and limits of 20–24% are now being proposed as an achievable limit in single junction devices.3 One of the most popular recent approaches to the development of high performance polymers for OPV applications is the design of donor–acceptor (D–A) copolymers.4,5 In these polymers the energy levels, and thus band gap of the polymer, may be tuned by modifications of the D and A units and the absorption of the polymer subsequently red-shifted and broadened.6 Polymers based upon diketopyrrolopyrrole (DPP) as the acceptor unit have a
Centre for Organic Electronics, University of Newcastle, Callaghan, NSW, 2308, Australia. E-mail: [email protected] b CSIRO Energy Technology, P.O. Box 330, Newcastle, 2300, Australia c Institute of Materials Research and Engineering (IMRE), Agency for Science, Technology, and Research (A*STAR), 3 Research Link, Singapore 117602. E-mail: [email protected] d Microelectronics Research Center, University of Texas at Austin, Austin, TX, 78758, USA † Electronic supplementary information (ESI) available. See DOI: 10.1039/ c3cp55037k
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proven particularly interesting and successful.7–10 This interest is largely due to the polymer’s broad optical absorption and the excellent planarity of the unit resulting in strong intermolecular p–p interactions and high charge carrier mobility.11 Unfortunately these interactions also typically lead to highly crystalline materials making the processing of said polymers and the formation of optimal BHJ active layer morphologies difficult.12 Small molecule BHJ OPV devices based on DPP-based donors have also shown promise and device PCEs of 4.4% have been achieved.8,13 A number of methods have been explored to overcome the processing problem associated with the use of DPP-based polymers. These include the synthetic modification of oligomers to form diketopyrrolopyrrole–oligothiophene:fullerene dyads and likewise the formation of thiophene-based D–A copolymers wherein the DPP acceptor constitutes only a minor (up to 15%) fraction of the polymer to ensure processability.14,15 A second approach involves the use of mixed solvent processing and the addition of processing additives to enhance the film morphology of unmodified DPP-based polymer.16 Using these techniques, DPP-based polymer:fullerene BHJ OPV devices with well blended morphologies have been achieved and device PCEs as high as 6.7% have been observed; mirroring the small molecule results.17 A further refinement through the use of solubilizing side chains on furan-flanked DPP (rather than thiophene) copolymers has led to very finely blended films which give devices with PCEs in excess of 6%.18 Importantly though, all of these methods achieve the optimal D–A blend morphology through the use of toxic or high
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boiling point solvents or by introducing considerable additional complexity to the polymer synthesis. Additionally, there are questions regarding the feasibility of employing mixed solvent systems and precise control of the solvent evaporation rate in large-area, roll-to-roll processing which would be employed in large scale, industrial fabrication of OPV devices. The recent development of aqueous dispersions of organic semiconductor nanoparticles offers a new way of addressing the issue of polymer processability for the fabrication of organic solar cells.19 This approach holds great appeal since it is compatible with organic semiconductor materials that have been developed for and utilized in organic solar cells fabricated by more traditional, organic solventbased fabrication techniques. In other words, no additional material design or synthesis is necessary to change the solubility properties of the materials. Additionally, the characteristic morphology of blended semiconductor nanoparticles fabricated in this approach is core–shell with phase segregation determined by the relative surface energies of the component materials.20,21 Consequently, this approach typically results in the formation of nanoparticles which demonstrate a fullerene-rich core and polymer-rich shell morphology.21,22 This intrinsic structure offers an elegant way to control the blend morphology on the nanometer length scale, driving the formation of domain sizes in the active layer close to the ideal for nanoparticles in the 30 nm diameter range. We describe the fabrication and characterization of both BHJ and NP OPV devices made from 1 : 2 blends of (poly{3,6-dithiophene-2-yl-2,5-di(2-octyldodecyl)-pyrrolo[3,4-c]pyrrole-1,4-dione-altnaphthalene}) (PDPP-TNT) and the fullerene acceptor, ([6,6]-phenyl C71-butyric acid methyl ester) (PC70BM). Previous PDPP-TNT: PC70BM BHJ devices have required a mixed solvent system (chloroform : o-dichlorobenzene, 4 : 1 by volume) to produce an advantageous active layer morphology.23 Here, we show that whereas deposition of the active layer from pure chloroform results in a coarse, phase segregated morphology with domain sizes close to the micron scale, the use of PDPP-TNT:PC70BM blend nanoparticles controls the active layer morphology on the nanometer scale. Consequently, device performance increases from an optimized PCE of B1.0% to a PCE of B2.0% through an increase of Jsc. Thus, this work demonstrates that aqueous-based NP dispersions can indeed be employed as a means to preorganize the OPV active layer blend morphology without resorting to synthetic modification of the polymer or the use of toxic or high boiling point solvents. Due to the fact that the donor and acceptor species are forced into close proximity and the blend morphology is fixed during the formation and solidification of the nanoparticles, the blend morphology is insensitive to the solvent evaporation rate during the film formation process. This trait offers promise for more facile translation from lab-scale optimization to large-scale fabrication.
Experimental PDPP-TNT was synthesised as described by Sonar et al.24 An aqueous dispersion of semi-conducting PDPP-TNT : PC70BM nanoparticles (1 : 2 blend) was prepared using the
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miniemulsion technique as outlined previously.20 Dynamic light scattering (Zetasizer Nano-ZS, Malvern Instruments, UK) was used to measure the distribution of particle sizes in the aqueous dispersion and gave a peak 1 particle size of 35 nm. Poly(3,4-ethylenedioxythiophene) poly(styrenesulfonate) (PEDOT: PSS) (Clevios P) films were spin coated (4000 rpm) on pre-cleaned patterned indium tin oxide (ITO) glass slides and annealed at 140 1C for 30 minutes to eliminate water in the films. The PDPP-TNT: PC70BM nanoparticle layer was deposited by spin coating 35 ml of the dispersion (1750 rpm for 1 minute) in air to give a film of 115 5 nm thickness. Following the deposition of the nanoparticle layer the films were dried at 110 1C, 130 1C or 150 1C for 5, 10, 30 or 60 minutes. Bulk heterojunction layers were spin cast from a PDPPTNT : PC70BM blend solution (1 : 2, 12 mg ml 1 in CHCl3) at 2000 rpm for 1 minute in a nitrogen-filled glovebox to give a 125 5 nm thick active layer. The films were then transferred into a vacuum chamber for cathode evaporation. The calcium/ aluminium (Ca/Al) electrodes were evaporated on the active layers in vacuum (2 10 6 Torr). The thickness of the Ca and Al layers were measured to be about 20 nm and 100 nm, respectively, using a quartz crystal monitor. After evaporation, fabricated devices were tested without further annealing. Current density–voltage ( J–V) measurements were conducted using a Newport Class A solar simulator with an AM1.5 spectrum filter. The light intensity was measured to be 100 mW cm2 by a silicon reference solar cell (FHG-ISE) and the J–V data were recorded with a Keithley 2400 source meter. Individual devices were masked (illuminated area of 3.9 mm2) to eliminate edge effects. External quantum efficiency measurements were recorded when the devices were irradiated with light from a tungsten halogen lamp passed through an Oriel Cornerstone 130 monochromator, using an Ithaco Dynatrac 395 analogue lock-in amplifier and Thorlabs PDA55 silicon diode to collect the reference signal, and a Stanford Research Systems SR830 DSP digitising lock-in amplifier to measure device current. An ultraviolet-visible absorption spectrophotometer (UV-vis, Varian Cary 6000i) was used to study the absorption of PDPPTNT:PC70BM nanoparticulate and bulk heterojunction films. For the UV-vis characterisation, films were spin coated onto quartz slides. Differential scanning calorimetry was performed on a Shimadzu DSC-60A, samples were scanned at 10 1C min 1 from 30 1C to 350 1C. Samples of the pristine PDPP-TNT and a 1 : 2 blend of PDPP-TNT : PC70BM nanoparticles (prepared by drop casting the nanoparticle ink directly into the aluminium pan and allowing to air dry) were recorded. Glass transition points (Tgs, taken at the onset) were observed at 84.4 1C for the pristine polymer and 87.9 1C for the polymer:fullerene nanoparticles. Atomic force microscopy (AFM) images were collected using an Asylum Research Cypher in AC mode. SEM was performed on a Zeiss Sigma ZP FESEM at accelerating voltages of 1–3 kV, and magnification ranges of 10 000 to 100 000. Samples were prepared for scanning electron microscopy (SEM) by spin coating 2.6 ml of the nanoparticulate dispersion, diluted 1/10 with water, (3000 rpm, 1 minute, low acceleration of 112 rpm s 1) onto a conductive silicon substrate.
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Results and discussion A series of 1 : 2 PDPP-TNT : PC70BM BHJ and NP OPV devices were prepared and characterised in order to determine the effect that preorganisation of the active layer morphology via the NP approach has on film morphology and device performance. BHJ blend films were prepared from a chloroform solution and used as-spun since heat treatment of these films was shown to have little effect upon device performance; NP films were dried before cathode deposition. Similar thermal treatment has been known to promote particle joining and to enhance subsequent device performance.22 In order to optimise this drying step, temperatures of 110 1C, 130 1C and 150 1C and drying times of 5 min, 10, min, 30 min and 60 min were investigated. Fig. 1 shows the normalised thin film UV-vis spectra of 1 : 2 blends of PDPP-TNT : PC70BM as-spun from chloroform and from the nanoparticulate dispersion. Also shown are the spectra of thin films of both the pristine component materials. It is clear that the blend spectra are composed of the distinguishing features of both the polymer and fullerene. AFM images of thin (B120 nm) films of both a BHJ film spin-coated from chloroform and a NP as-spun thin blend film are shown in Fig. 2 and a number of features are apparent. The BHJ film shows large B200–500 nm diameter domains which extend above the surface of the bulk film by approximately 30 nm. Previous work has shown that the regions between these domains are sulphur-rich indicating that the film is comprised of fullerene-rich plateaus surrounded by polymer-rich valleys.23 Similar domain growth is observed for films containing PPVbased materials and is due to the relative solubilities of the two components.25 In contrast, AFM images of the deposited PDPP-TNT:PC70BM NP films shows that they retain their discrete
Fig. 1 UV-vis spectra of an as spun PDPP-TNT : PC70BM 1 : 2 nanoparticulate film (solid line), PDPP-TNT : PC70BM 1 : 2 bulk heterojunction thin film (thick dashed line), pure PDPP-TNT thin film (light dotted line) and pure PC70BM thin film (light dashed line). All spectra have been normalized to the highest peak in the 725–750 nm range except for the PC70BM film which has been normalized at l = 300 nm.
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Fig. 2 AFM images of an as-spun PDPP-TNT : PC70BM 1 : 2 blend film spincast from chloroform (a) and an as-spun PDPP-TNT : PC70BM 1 : 2 nanoparticulate film (b). Scale bars are 500 nm.
nanoparticulate nature and size, with NPs of approximately 35 nm diameter (consistent with the measured DLS results for the aqueous dispersion) observed as a continuous film. The NP approach can be seen to offer an initial advantage over the traditional single-solvent system, providing finer scale mixing of the donor and acceptor and greater interfacial area. Fig. 3 shows the J–V characteristics for the as-spun PDPPTNT : PC70BM 1 : 2 NP OPV devices and those heated at 110 1C, 130 1C and 150 1C for up to one hour. It is clear that all of the devices require a heating step to become functional, with devices heated at all temperature and time combinations showing vastly superior performance in both the J–V and EQE plots. Heating the devices at 110 1C for even prolonged periods has little effect on device performance beyond the initial performance boost. By contrast, heating the devices at 150 1C leads to a systematic decrease in both the Voc and Jsc, and therefore the efficiency, of the devices. Heating at 130 1C bridges the results of the two outlying temperatures and initially results in an increase in current (and efficiency) and then a systematic decrease in both Voc and Jsc (and efficiency). Thus, from the data shown in Fig. 3a–c there are broadly three states for these devices within the explored regime of heating conditions. As-spun devices barely perform and display very poor device parameters. Moderate heating (110 1C for any length of time or 130 1C for less than half an hour) results in dramatic device improvement giving devices with Jsc B5–6 mA cm 2, Voc B 0.77 V, and PCE between 1.5 and 2.0%. Significantly, all of these devices show current densities (and efficiencies) much superior to the corresponding BHJ device. These observations are consistent with the measured Tgs of 85 1C for the pristine polymer and 88 1C for the polymer:fullerene blend nanoparticles (see ESI†), which indicates that for all three temperatures investigated substantial mobility of the components would be expected. However, excessive annealing of the devices (130 1C for >10 min and 150 1C for any length of time) results in a reduction in performance and gives rise to devices with Jsc B 2–3 mA cm 2, Voc B 0.60 V, and PCE between 0.5 and 1.0%. All devices display FFs of approximately 0.4, which is B30% lower than the BHJ devices. We attribute this reduction to the surfactant present in the nanoparticulate active layer; we have separately observed a systematic increase in FF for PDPP-TNT : PC70BM 1 : 2 nanoparticulate OPV devices as the amount of surfactant used in the NP dispersion is decreased (data not shown).
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Fig. 4 SEM images of as-spun (a), heated at 110 1C for 10 minutes (b), heated at 130 1C for 10 minutes (c), and heated at 150 1C for 10 minutes (d), PDPP-TNT : PC70BM 1 : 2 nanoparticulate films. All scale bars are 200 nm.
Fig. 3 Current density versus voltage curves for PDPP-TNT : PC70BM 1 : 2 nanoparticulate OPV devices heated at 110 1C (a), 130 1C (b) and 150 1C (c) for 0 minutes (dotted line), 5 minutes (short dashes), 10 minutes (medium dashes), 30 minutes (long dashes) and 60 minutes (solid line).
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The observed device performance may be explained by the observation of changes in the morphology of the active layer upon heating. Fig. 4 shows scanning electron microscopy (SEM) images of an as-spun PDPP-TNT : PC70BM 1 : 2 nanoparticulate film, as well as films that have been heated at 110 1C, 130 1C and 150 1C for 10 minutes. The image of the as-spun film (Fig. 4a) shows that the PDPP-TNT:PC70BM NPs are ovoid, rather than spherical, in nature. The axis dimensions of the NPs (major axis B42 nm, minor axis B23 nm) compare well with the particle size determined from DLS for the NP suspension (35 nm) and indicate that the particles have an aspect ratio of 9 : 5. Upon heating of the film at 110 1C for 10 minutes the particles appear to swell slightly in size and join together (Fig. 4b). The joining of the particles is further pronounced upon heating of the film at 130 1C for 10 minutes but discrete particles are still clearly retained (Fig. 4c). By contrast, heating the film at 150 1C for 10 minutes results in the particles joining completely and the particulate structure of the film vanishing (Fig. 4d). In addition, apparently crystalline aggregates appear on the film surface, consistent with gross phase segregation of the fullerene from the polymer matrix.21 These images strongly support the device results, showing that the improvement in device performance upon initial heating of the films results from joining of the NPs whilst retaining the overall nanoparticulate structure. The improved connectivity within the device necessarily results in much improved transport and therefore improved device performance. Heating the film beyond 130 1C (or extended heating at lower temperatures) leads to gross phase segregation of the NP film. Both charge generation and transport are now adversely affected and device performance consequently drops. From Fig. 3, we see that heating the devices at 130 1C results in the optimal device performance. Table 1 compares the device characteristics of the 1 : 2 PDPP-TNT : PC70BM BHJ OPV device and a series of 1 : 2 PDPP-TNT : PC70BM NP OPV devices dried at 130 1C for different times, all of similar thickness.
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Table 1 Comparison of device characteristics for the 1 : 2 PDPPTNT : PC70BM BHJ device and 1 : 2 PDPP-TNT : PC70BM NP OPV devices heated at 130 1C for varying times. The series resistance (Rs) and shunt resistance (Rsh) values were calculated by taking the inverse slope of the current–voltage curve near Jsc and Voc, respectively
BHJ Device NP Device (as spun) NP Device (5 min dry) NP Device (10 min dry) NP Device (30 min dry) NP Device (60 min dry)
Voc (V)
Jsc (mA cm 2)
FF
0.82 0.38 0.77 0.76 0.61 0.58
2.03 0.55 5.28 6.09 3.12 2.94
0.63 0.25 0.42 0.43 0.41 0.42
PCE (%)
Rs (O)
1.05 606 0.05 17 000 1.71 827 1.99 532 0.78 1160 0.70 1260
Rsh (O) 79 600 17 900 8250 7130 10 600 12 300
DPP-TNT-based BHJ devices fabricated from a single solvent (chloroform) exhibit a PCE significantly lower than that observed for active layers deposited via a mixed solvent system.24 The Voc and FF of devices prepared from the single and mixed solvent systems are essentially identical,23 indicating that the difference in performance lies solely in the generated photocurrent. Indeed, the Jsc of the chloroform processed device is a mere one quarter of that of the Jsc observed using the mixed solvent system.24 In light of the coarse phase segregation in the chloroform processed active layer observed here and previously,23 the poor performance of the subsequent device can be explained by inefficient charge separation of excitons and therefore high geminate recombination in the active layer. Considering the device parameters for the NP OPV devices heated at 130 1C, there are a number of observations. Firstly, the as-spun device shows reduced characteristics for all parameters and can best be described as barely functioning. This observation is consistent with the findings of previous studies of NP OPV devices and thin films which show that initial heating of the NP film promotes joining of the NPs and consequently improves device performance.22 This observation is supported by the abrupt large decrease in Rs that occurs upon initial heating of the NP film. After the initial 5 minutes, subsequent heating results in a systematic decrease in the Voc of the device. FF remains invariant across the heating regime, however, Jsc and PCE are maximized after 10 minutes of annealing and then drop rapidly. The EQE plots for devices heated at 130 1C (Fig. 5) are in good agreement (5%) with the measured Jsc values, reinforcing the observation that the changes in device performance upon heating arise primarily as a result of changes in current generation. In addition, the shape of the EQE plots shows that the relative contribution of the polymer varies greatly with thermal treatment. In optimally treated devices the long wavelength peak (B730 nm, polymer absorption peak) is significantly higher than at B500 nm (minimum of polymer absorption). For extended treatment, the EQE contribution at 500 nm is larger than at 730 nm. The change in shape of the spectral response suggests that excitons created upon photoexcitation of the polymer are more efficiently dissociated and transported to the electrodes after optimal thermal treatment. This improvement could be due to improved exciton diffusion and/or a rearrangement of the blend components within the
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Fig. 5 EQE spectra for the as-spun PDPP-TNT : PC70BM 1 : 2 nanoparticulate OPV device (dotted line) and those heated at 130 1C 5 minutes (short dashes), 10 minutes (medium dashes), 30 minutes (long dashes) and 60 minutes (solid line).
nanoparticle. Previous work has shown that PDPP-TNT is amorphous at room temperature but rearranges to become increasingly crystalline following annealing above 100 1C;24 consistent with our observation of a Tg at 85 1C. Given that the NP shell is thought to be polymer-rich,20 the increase in device performance from 5–10 minutes can best be explained by the increased polymer crystallinity resulting in improved charge transport.24 Subsequent annealing, however, results in expulsion of PCBM from the ordered polymer matrix (Fig. 4d), as has been observed in the P3HT:PCBM system,21 resulting in suboptimal blend composition and a corresponding reduction in photocurrent generated from the PDPP-TNT polymer. Fig. 6 compares the J–V and EQE plots for the most efficient PDPP-TNT : PC70BM 1 : 2 nanoparticulate OPV device and the corresponding BHJ device spin coated from chloroform. It is clear from these plots that the NP device substantially outperforms the BHJ device. The EQE spectrum of the NP device shows significant contribution from the long wavelength region, attributed to polymer absorption. Previous studies have shown the thickness of the BHJ device to influence the shape of the EQE spectrum;23 in this study, attempts were made to maintain a comparatively equal thickness for the BHJ and NP devices. As mentioned, the large-scale phase segregation (hundreds of nanometers) observed for the BHJ film (Fig. 2a) is non-optimal for charge generation in these devices.26 Indeed, although the BHJ device exhibits a Voc larger than the optimal NP device it has a current density comparable to that of the overannealed NP device and consequently a much reduced efficiency. By comparison, the NP film retains a phase segregated morphology in which the domain sizes are constrained to tens of nanometers and thus are much better suited to exciton dissociation. As a result, the photocurrent of the optimised NP device is more than double that of the BHJ device.
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or solvent additives are typically employed to promote mixing and optimize the blend morphology. The NP approach offers a means to control the initial, intrinsic blend morphology independent of the evaporation time during spin coating. This control and the solvent-free nature of the aqueous dispersions are highly beneficial traits for the development of large-scale fabrication of OPV devices.
Acknowledgements Special thanks to the University of Newcastle Electron Microscopy and X-ray Unit. Thanks to Mr Nicolas Nicolaidis for his assistance with EQE measurements. The University of Newcastle and the Australian Renewable Energy Agency (ARENA) are gratefully acknowledged for PhD scholarships (NH). The University of Newcastle is gratefully acknowledged for a Visiting Fellowship (ELW). Authors ELW and PS acknowledge the Visiting Investigatorship Programme (VIP), project number 0721100037, of the Agency for Science, Technology and Research (A*STAR), Republic of Singapore. This work was performed in part at the Materials node of the Australian National Fabrication Facility, a company established under the National Collaborative Research Infrastructure Strategy to provide nano and microfabrication facilities for Australia’s researchers.
References
Fig. 6 Current density versus voltage curves (a) and EQE plots (b) for a 1.0% efficient PDPP-TNT : PC70BM 1 : 2 bulk heterojunction device deposited from chloroform (solid line) and the optimal 2.0% efficient PDPP-TNT : PC70BM 1 : 2 nanoparticulate OPV device heated at 130 1C for 10 minutes (dotted line).
Conclusions We have demonstrated that the use of NP films in PDPPTNT:PC70BM OPV devices provides control over the active layer morphology on the scale of typical exciton diffusion lengths in organic semiconductor materials. The resultant constrained morphology allows for improved exciton harvesting within the device and therefore an increase in photocurrent generation over comparable BHJ devices. Furthermore, we have shown that careful heating of the PDPP-TNT:PC70BM NP active layer is required to optimize device performance. Initial heating is required to promote connectivity of the NPs in order to improve charge transport through the film. However, over-annealing results in gross phase segregation within the NP film and a subsequent reduction in device performance. The NP approach shows great promise in donor–acceptor systems where largescale phase segregation is observed and mixed solvent systems
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