Heterojunction Photovoltaics Using GaAs Nanowires and Conjugated

Dec 20, 2010 - Shaw , P. E.; Ruseckas , A.; Samuel , I. D. W. Adv. Mater. 2008, 20 ..... Sema Ermez , Eric J. Jones , Samuel C. Crawford , and Silvija...
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LETTER pubs.acs.org/NanoLett

Heterojunction Photovoltaics Using GaAs Nanowires and Conjugated Polymers Shenqiang Ren,† Ni Zhao,‡ Samuel C. Crawford,† Michael Tambe,† Vladimir Bulovic,‡ and Silvija Gradecak*,† †

Department of Materials Science and Engineering and ‡Department of Electrical Engineering and Computer Science, Massachusetts Institute of Technology, Cambridge, Massachusetts 02139, United States

bS Supporting Information ABSTRACT: We demonstrate an organic/inorganic solar cell architecture based on a blend of poly(3-hexylthiophene) (P3HT) and narrow bandgap GaAs nanowires. The measured increase of device photocurrent with increased nanowire loading is correlated with structural ordering within the active layer that enhances charge transport. Coating the GaAs nanowires with TiOx shells passivates nanowire surface states and further improves the photovoltaic performance. We find that the P3HT/ nanowire cells yield power conversion efficiencies of 2.36% under white LED illumination for devices containing 50 wt % of TiOx-coated GaAs nanowires. Our results constitute important progress for the use of nanowires in large area solution processed hybrid photovoltaic cells and provide insight into the role of structural ordering in the device performance. KEYWORDS: III-V Nanowires, conjugated polymers, bulk heterojunction solar cell, self-assembly, and molecular ordering

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hotovoltaic devices based on solution processable conjugated polymers are attractive for the production of low cost solar cells.1,2 To obtain high efficiencies of exciton dissociation and high photocurrent, it is desirable to have an interpenetrating network of electron-donor and electron-acceptor components within the device, creating what is often referred to as a bulk heterojunction (BHJ).3 This can be achieved using polymer blends4 or mixtures of conjugated polymers with C60 derivatives.3,5 Blended organic/ inorganic hybrid solar cell BHJs offer further enhancement in device performance because of the relatively high electron mobility, large dielectric constant, and superior physical and chemical stability of the nanostructured inorganic components.6 The hybrid cells have been fabricated using inorganic electron transporting materials such as TiO27 and ZnO nanoparticle networks8 or ordered TiO2 mesoporous films.9,10 The nanoparticle phase can act as both an electron transporter and a complementary photon absorber if narrower bandgap materials such as CdSe,11 PbS,12 or PbSe13 are used. Indeed, these polymer/nanocrystal solar cells demonstrated higher efficiencies than those based on the corresponding singlecomponent films. However, the hybrid cells are still limited by the inefficiency of the hopping charge transport through the discontinuous percolation pathways in the BHJ films.14 An ideal organic/inorganic film would provide continuous pathways for charge transport in the through-film direction and maintain solution processability and therefore low manufacturing cost. BHJ photovoltaic devices employing semiconductor nanowires (NWs) have the potential for further enhancement of electron transport due to the single crystalline and onedimensional structure of NWs. High aspect ratio NWs provide a direct electron pathway that does not rely on electron hopping, r 2010 American Chemical Society

in contrast to nanoparticle BHJ cells. Furthermore, the internanowire spacing can be controlled such that it is comparable to the exciton diffusion length of a typical polymer absorber.15 Hybrid cells have been fabricated by infiltrating a conjugated polymer into the channels between as-grown NW arrays to facilitate electron conduction to the anode.16-19 Such arrays have met with only partial device success due to the low NW packing density, difficulty of infiltrating the conjugated polymer and poor crystallinity of the polymer within the narrow channels. In this letter, we develop organic/inorganic solar cell devices by blending the uniformly sized GaAs NWs with a conjugated polymer poly(3-hexylthiophene) (P3HT) in a single solvent to form a uniform film consisting of dispersed NWs in a polymer matrix. This approach takes advantage of the high aspect ratio, high electron mobility, and complementary photon absorption of GaAs NWs, while preserving solution processability. By controlling the concentration of GaAs NWs dispersed in P3HT, we tailor the P3HT ordering and NW alignment, thus enhancing both electron and hole transport. Finally, we show that a thin shell of amorphous TiOx coated onto the GaAs NWs passivates the NW surface and serves as a tunneling barrier to efficiently reduce surface recombination, further enhancing the device performance. The GaAs NWs were grown on GaAs(111)B substrates using metal-organic chemical vapor deposition (MOCVD) and 25 nm Au colloidal nanoparticles as NW growth seeds to yield uniformly sized NWs, as described in our previous work.20 NW powder was Received: August 25, 2010 Revised: November 10, 2010 Published: December 20, 2010 408

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prepared from as-grown NWs (see Supporting Information) and used for the device fabrication, as follows. First, poly(3,4ethylenedioxythiophene)-poly(styrenesulfonate) (PEDOT-PSS) was spun cast onto a 12  12 mm2 glass substrate with prepatterned ITO electrodes. The NW powder was mixed with ethanol/1,2 dicholoro-benzene (1,2-DCB), sonicated, and filtered to obtain a clear solution that was then blended with filtered P3HT in 1,2-DCB solution in different weight ratios. In this manuscript, we use the simplified term GaAsx%, where x refers to the weight percentage of GaAs NWs in the P3HT/GaAs active layer. The blended solution was spun to form a 160 nm thick film and annealed overnight in 1,2-DCB solvent. The coated devices were then heat-treated at 175 °C for 10 min and quickly cooled to room temperature. Finally, the 10 nm thick bathocuproine (BCP) hole blocking layer and the top Mg/Ag electrode were evaporated. The final device area was defined by the overlap between the top and bottom electrodes. Current-voltage (J-V) characteristics of the devices were measured in a nitrogen atmosphere glovebox with a Keithley 6487 source meter. The light response was measured under illumination from a white LED with a broad emission spanning wavelengths from 400 to 675 nm and light output intensity of approximately 70 mW/cm2 (Supporting Information, Figure S1). The wavelength-dependent quantum efficiency data was measured using a lock-in amplifier with the sample under illumination from monochromatic light generated by a Xenon arc lamp coupled to a monochromator and chopped at 40 Hz. A calibrated silicon photodiode was used to measure the incident monochromatic light intensity. Transmittance and absorbance spectra of the device active layer were measured with a Cary 5E

UV-vis-NIR dual-beam spectrophotometer. A JEOL 2010 FEG analytical transmission electron microscope (TEM) operated at 200 kV was used for structural analysis. Device surface morphology was investigated using a Digital Instruments Dimension 3000 atomic force microscope (AFM) operated in tapping mode. Figure 1 shows a schematic of the resulting device and its simplified flat band energy. In this cell, the PEDOT-PSS layer creates a smooth surface on the ITO and serves as a hole transporting layer together with the P3HT, which also acts as the light-absorbing material. GaAs NWs with an electron affinity about 0.7 eV greater than that of P3HT were used to extend the photoabsorption of the active layer to the near-infrared and to form an electron transport pathway toward the Mg/Ag top electrode. A thin layer of BCP was used between the metal electrode and the active layer to block the diffusion of excitons and holes to the metal electrode, as well as to reduce damage to the active layer from metal evaporation.21 Devices with x = 0, 20, 30, 40, and 50% were prepared and measured under illumination from a 70 mW/cm2 white LED (to minimize NW aggregation, experiments were performed with a maximum loading of x = 50% GaAs NW). Figure 2 presents a summary of operating characteristics for these cells. At relatively low NW loading (x < 30%), the nominal short circuit current density (Jsc) does not change dramatically compared to a P3HT-only device. Once a threshold loading point (x ≈ 30%) is reached, the Jsc starts to increase, which is consistent with the formation of GaAs NW percolation networks that would enhance charge extraction out of the device. A similar trend is observed for the open circuit voltage (Voc) and fill factor (FF), which gradually rise to their full value of 0.59 V and 40%, respectively, at x = 50% (Figure 2b). The J-V response of the devices in the dark (Figure 2c) also indicates that NW loading concentration affects the leakage current and the turn-on voltage. The GaAs50% device has negligible leakage current and a relatively high turn-on voltage in comparison to devices with lower NW loading. To elucidate the role of GaAs nanowires in the devices, we measure the photoabsorption spectra of the pure P3HT thin film, P3HT/GaAs NW thin films, and GaAs NWs (Figure 3a). The pure P3HT thin film after annealing shows a well-ordered absorption behavior: the π-π* transitions were observed at 525 and 560 nm wavelengths, whereas the shoulder at 610 nm corresponds to an interchain absorption. Upon blending with GaAs NWs, the P3HT absorption shoulder at 610 nm is initially suppressed for low loading concentrations, but fully recovers for NW loadings of x > 20%. The intensity of this interchain absorption shoulder is correlated with the degree of order in the

Figure 1. Structure and energy band diagram of the hybrid P3HT/ GaAs device. (a) Schematic of the PEDOT-PSS/P3HT-GaAs/BCP/ Mg-Ag device stack. Light shines through the ITO and is absorbed by P3HT. Charges are separated at the P3HT/NW interface such that electron and hole transport occurs through the NWs and P3HT, respectively, toward the corresponding electrodes. (b) Proposed flat energy band diagram of the device, indicating energy level positions with respect to the vacuum level.

Figure 2. Hybrid solar cell characteristics based on the weight percentage of nanowires in the P3HT/GaAs blends, measured under 70 mW/cm2 white LED illumination. (a) J-V curves of illuminated P3HT/GaAs hybrid cells for nanowire loading concentrations of x = 0, 20, 30, 40, and 50%. (b) A summary of the operating characteristics for these cells. Jsc is shown on the left axis, and Voc and FF are shown on the right axis. (c) Dark J-V curves of the P3HT/GaAs cells for the same loading concentrations as in (a). 409

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Figure 3. (a) Absorption spectra of a P3HT thin film, GaAs nanowires (25 nm diameter) in 1,2-DCB solution, GaAs20%, and GaAs50% blends. (b) Comparison of EQE data for P3HT hybrid cells with different NW loading. Inset image shows S factor as a function of the nanowire loading.

P3HT phase,22 so the observed absorption behavior indicates an increase in molecular ordering in the P3HT phase as the NW loading is increased, which is discussed later in more detail. The photoabsorption results are supported by external quantum efficiency (EQE) measurements (Figure 3b). The EQE spectral response of pure P3HT is consistent with the measured absorption spectrum and shows an absorption edge at 660 nm. Any device response at longer wavelengths can therefore be attributed to the absorption by GaAs nanowires and/or GaAs-P3HT aggregates. Thus, the relative current contribution from this low-energy portion of EQE spectrum directly reflects the amount of charge carriers due to the GaAs NW absorption. The NW contribution can be quantified using a parameter S defined as the ratio of the photocurrent corresponding to the absorption from 660 to 900 nm to the overall photocurrent of the cell:23 R 900 EQEðλÞIðλÞdλ S ¼ R 660 ð1Þ 900 350 EQEðλÞIðλÞdλ

Figure 4. (a) Cross-sectional TEM image of a P3HT-GaAs active layer. BCP/Mg and PEDOT-PSS layers are schematically shown as the top and the bottom layer, respectively, but were not present in the TEM sample. The right bright-field TEM image shows a single GaAs nanowire and the conformal P3HT coating. Scale bars are 50 nm. (b) The protruding nanowire percentage as a function of GaAs nanowire concentration extracted from AFM surface topology (inset, scale bar is 100 nm). (c) In-plane GIXS measurement for annealed films with different nanowire loadings. Inset, schematic of vertical orientation of P3HT chains along nanowire surface.

and photoabsorption spectra indicate possible P3HT ordering with increased NW loading, which enhances charge transport and therefore increases the Jsc. To further explore the role of the NWs in the improved device performance, we have performed structural analysis using TEM, AFM, and grazing incidence X-ray scattering (GIXS) on the drying-mediated films. From TEM cross-section images (Figure 4a), we observed that the GaAs NWs tend to align vertically, rather than horizontally, within the P3HT matrix. This alignment is further confirmed by AFM top-view measurements, which show that that the overall percentage of NWs reaching the anode increases with the NW weight fraction, indicating that alignment increases with NW concentration (Figure 4b). Moreover, from the AFM height profile (Supporting Information, Figure S2), the average distance between protruding NWs in the GaAs50% blend was measured to be approximately 35 nm, which is considerably closer to the exciton diffusion length in P3HT (12 nm)24 compared to the inter-NW distance in other previous inorganic NW/organic BHJ cells.25 The observed self-alignment could be induced due to the kinetically controlled orientational

Here, I(λ) is the wavelength-dependent incident photon flux under AM1.5G solar illumination. The parameter S increases (Figure 3b, inset) with NW loading and reaches its maximum of approximately 3.8% for x = 50%. These relatively low S values indicate that the absorption by GaAs NWs does not contribute significantly to the overall device photocurrent and that other mechanisms must play a role in the overall improvement of device performance with increased NW loading. Taken together, the results presented so far indicate the following: (1) the parameter S increases with the NW loading, but it alone is too small to account for dramatic improvement in device performance; (2) a threshold concentration of NWs is necessary for any significant changes to be observed; (3) EQE 410

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order or competition between entropic and enthalpic driving forces through drying-mediated self-assembly.26-29 In addition, the interaction energy between the NWs and the underlying PEDOTPSS might also play a role in NW vertical alignment.30,31 Figure 4c shows the in-plane GIXS patterns of P3HT/GaAs films with various NW concentrations. In the case of the pristine P3HT thin film, the most pronounced peaks are (100) and (010) reflections, which correspond to the measured lamellar layer structure (16.2 Å) and out-of-plane π-π interchain stacking (3.8 Å), respectively. These values agree well with the known unit cell parameters of P3HT (a = 16.2, b = 3.8, and c = 7.8 Å).32 The GIXS results indicate that two different orientations of crystalline lamellae, (010) flat-on (lying) and (100) edge-on (standing),33,34 are induced with respect to the substrate in our pristine P3HT films. As the loading of GaAs NWs into the P3HT matrix increases, the intensity of the (100) Bragg peak gradually increases, while the (010) reflection disappears. When the NW weight ratio reaches 50%, the intensity of the (100) Bragg peak dramatically increases, and the higher order (200) and (300) peaks appear as well. These in-plane GIXS measurements can be understood as follows. For low concentrations (x < 20%), the NWs disrupt the P3HT chain crystallization, which is consistent with our photoabsorption measurements. As the amount of added self-aligned NWs increases, the molecular orientation of the P3HT becomes an edge-on structure with the [100] axis normal to the NW surface (Figure 4c, inset). A high NW concentration (x = 50%) enhances the chain alignment along the NW surface and the overall crystallinity of the P3HT, which is even greater than that of the pristine P3HT film. The overall effect is that NWs act as templates for increasing the π-conjugation of the P3HT phase, an effect resulting from the onedimensional NW geometry and the π-π stacking interaction at the interface between the P3HT backbone and the NWs. The structural ordering of NWs and P3HT directly correlates with the observed rise in Jsc, Voc, and FF with increased NW loadings (Figure 2). The rise in Jsc with increased loading of GaAs is consistent with the fact that nanowires function as a percolation network and direct electron transport pathway, and the NWinduced P3HT molecular ordering enhances the hole transport. The increase of Voc with NW loading can be understood from the dark current output of these cells (Figure 2c). The decrease of leakage current with increased loading is consistent with the vertical phase segregation of GaAs NWs and P3HT. Crosssectional TEM images and AFM surface morphology (Figure 4a,b) both suggest that the NWs segregate to the upper portion of the P3HT/NW active layer, yielding a bottom P3HT-rich region, which functions as an electron blocking layer. The upper GaAs-rich region acts as hole blocking layer. As NW loading increases, more interfacial area is present to extract electrons from P3HT to GaAs, which reduces charge carrier losses from bimolecular recombination in the P3HT phase.35,36 This leads to a large built-in electric field and a high Voc. As a result, the overall device performance is enhanced with increased NW loading, and the maximum power conversion efficiency of 1.98% under 70 mW/cm2 white LED illumination was achieved for the highest loading (x = 50%). It is important to emphasize that this enhancement is due solely to the structural self-assembly within the active region, without any additional device optimization, and that similar self-assembly processes may play a role in other organic/inorganic devices involving one-dimensional materials, including carbon nanotubes37,38 and other semiconductor nanowires.

Figure 5. (a) Bright-field TEM image of an amorphous TiOx layer coating on a GaAs NW and the corresponding EDS spectrum (upper right corner). Copper and carbon signals in the EDS spectrum originate from the TEM grid. The lower left inset shows a bright-field TEM image of GaAs NWs before TiOx coating. Scale bars are 20 nm. (b) HRTEM images and the corresponding Fourier transform of the self-assembled P3HT lamellar layer surrounding a GaAs-TiOx core-shell nanowire. Scale bars are 5 nm.

The device performance could be further enhanced by engineering the interface between P3HT and NWs to facilitate charge separation with optimized offset of the donor and acceptor lowest unoccupied molecular orbital (LUMO) states. We therefore prepared GaAs-TiOx core-shell NWs by coating the surface of GaAs NWs with a thin TiOx shell using sol-gel chemistry. The role of the TiOx layer is to passivate the nanowire surface and enhance the electron transport (TiOx has a lower conduction band edge energy than GaAs).39 In addition, the TiOx/P3HT interface is known to readily split excitons generated in P3HT.16 Figure 5 shows a bright-field TEM image of a GaAs-TiOx core-shell hybrid structure with a TiOx thickness of about 3 nm (for comparison, NW structures before TiOx coating are shown in the inset). The EDS spectrum (Figure 5a, inset) confirms the chemical composition of the core-shell nanowires, while the TEM image shows that the TiOx shell is conformal, smooth and completely covers the NW surface. After the TiOx coating, the core-shell NWs were blended with P3HT for further device fabrication. High-resolution TEM images (Figure 5b) show highly ordered lamellar domains of P3HT at the polymer/NW interface. The π-π stacking distance was measured to be 0.36 nm based on the corresponding Fourier transform (Figure 5b, inset), which agrees with the theoretical value32 and confirms that the P3HT π-π stacking is preserved after the TiOx coating. A J-V plot of a typical GaAs-TiOx hybrid cell is characterized by Jsc = 7.16 mA/cm2, Voc = 0.57 V, FF = 41%, and power conversion efficiency of 2.36% under 70 mW/cm2 white LED illumination (Figure 6a), which is 20% higher than the unpassivated GaAs NW cell. The EQE of the GaAs-TiOx core-shell cell is enhanced in both the P3HT and GaAs absorption range compared to an unpassivated hybrid cell with the same NW loading (Figure 6b). Interestingly, devices made with only 40 wt % GaAs-TiOx showed similar Jsc values to the ones prepared with 50 wt % GaAs. We note that the EQE from 660 to 800 nm is relatively low even in the presence of the TiOx shell due to the small band offset between the highest occupied molecular orbital (HOMO) of P3HT and the valence band of GaAs phase. Although the TiOx shell does not contribute significantly to the absorption, the S factor increased to 6.1% by incorporating the TiOx shell, which is likely due to the passivation of the GaAs nanowire surface states that can cause exciton quenching. 411

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In conclusion, a new type of hybrid BHJ solar cell formed by blending GaAs NWs with P3HT in solution was demonstrated. The morphology and alignment were addressed directly in the spin-coating and subsequent annealing procedures, in contrast to infiltrating polymer into NW arrays. The self-alignment of NWs in the upper portion of the active layer and the consequent formation of a P3HT-rich bottom layer enhance charge carrier extraction. Above a certain NW loading threshold, the NW structures facilitate P3HT molecular ordering and improve charge transport. The enhancement of P3HT molecular ordering is due to the one-dimensional morphology of the GaAs NWs and can be extended to other systems with one-dimensional components. To further enhance device performance, we modified the organic/inorganic interface for effective charge separation. We have shown that coating GaAs NWs with 3 nm of TiOx improves the power conversion efficiency of our P3HT/NW solar cells by 20%, resulting in an efficiency of 2.36% for these devices. Future efforts will concentrate on combining low-bandgap conjugated polymers with bandgap-matched 1D NW structures.

’ ASSOCIATED CONTENT Figure 6. (a) J-V measurements of the core-shell GaAs-TiOx nanowire (50 wt %) hybrid cell in the dark and under 70 mW/cm2 white LED illumination. The inset shows the energy diagram of the bulk band offsets for the active components involved. (b) Comparison of EQE data for the GaAs-TiOx core-shell nanowire (50 wt %) and GaAs nanowire (50 wt %) blend devices. (c,d) J-V characteristics under 850 nm infrared laser irradiation for (c) GaAs nanowire hybrid cell and (d) GaAs-TiOx core-shell nanowire hybrid cell. Black and red curves are the dark and illuminated states, respectively.

To measure effects of the surface passivation on the photoresponse of GaAs NWs before and after TiOx shell coating, we used a 850 nm infrared laser with an output of 120 mW/cm2 to excite the GaAs NWs within the devices. Here, we reasonably assume that the photocurrent comes only from absorption in the GaAs NWs and that effects from P3HT can be neglected under these illumination conditions. Figure 6c,d shows the comparison of electrical output characteristics under 850 nm illumination for GaAs NW core-only and GaAs-TiOx core-shell NW hybrid cells, respectively. GaAs-TiOx devices show an 8-fold increase in Jsc compared to the core-only hybrid cell, confirming that the surface passivation plays a significant role in the overall device performance. The core-shell device also demonstrates a negligible leakage current and relatively high FF under these illumination conditions. The low Jsc and Voc of the NW core-only device stems from interfacial recombination, as evidenced by the large dark current and high diode ideality factor (Figure 6c). The surface states of NWs lead to greater exciton quenching and therefore negligible EQE at wavelengths above 660 nm (Figure 3b). This is further supported by measuring the light intensity dependence of the photocurrent (Supporting Information, Figure S4), which shows that the Jsc increases linearly at low illumination intensities indicating that the main loss mechanism of the photocurrent is geminate electron-hole pair recombination. However, at high light intensities with greater densities of photogenerated carriers sublinear behavior is possible because of bimolecular recombination and/or space-charge buildup due to low hole mobility in P3HT.40 These findings are significant because high surface-to-volume ratio in the nanowires is expected to exacerbate recombination, making surface passivation important for future NW-based PV devices.

bS

Supporting Information. Experimental methods including GaAs nanowire growth, core-shell coating, device fabrication and testing, and structural analysis. SEM and AFM surface topology analysis and the J-V characterization of GaAs-TiOx devices under 850 nm laser illumination. This material is available free of charge via the Internet at http://pubs.acs.org.

’ AUTHOR INFORMATION Corresponding Author

*E-mail: [email protected].

’ ACKNOWLEDGMENT This work was supported by Eni SpA under the Eni-MIT Alliance Solar Frontiers Program. The authors acknowledge access to Shared Experimental Facilities provided by the MIT Center for Materials Science Engineering supported in part by MRSEC Program of National Science Foundation under award number DMR - 0213282. ’ REFERENCES (1) Halls, J. J. M.; Walsh, C. A.; Greenham, N. C.; Marseglia, E. A.; Friend, R. H.; Moratti, S. C.; Holmes, A. B. Nature 1995, 376, 498. (2) Gaynor, W.; Lee, J. Y.; Peumans, P. ACS Nano 2010, 4, 30. (3) Yu, G.; Gao, J.; Hummelen, J. C.; Wudl, F.; Heeger, A. J. Science 1995, 270, 1789. (4) Huang, J. S.; Li, G.; Yang, Y. Adv. Mater. 2008, 20, 415. (5) Shaheen, S. E.; Brabec, C. J.; Sariciftci, N. S.; Padinger, F.; Fromherz, T.; Hummelen, J. C. Appl. Phys. Lett. 2001, 78, 841. (6) Huynh, W. U.; Dittmer, J. J.; Alivisatos, A. P. Science 2002, 295, 2425. (7) Arango, A. C.; Carter, S. A.; Brock, P. J. Appl. Phys. Lett. 1999, 74, 1698. (8) Beek, W. J. E.; Wienk, M. M.; Kemerink, M.; Yang, X. N.; Janssen, R. A. J. J. Phys. Chem. B 2005, 109, 9505. (9) Liu, J. S.; Kadnikova, E. N.; Liu, Y. X.; McGehee, M. D.; Frechet, J. M. J. J. Am. Chem. Soc. 2004, 126, 9486. (10) Chang, J. A.; Rhee, J. H.; Im, S. H.; Lee, Y. H.; Kim, H.-J.; Seok, S. I.; Nazeeruddin, M. K.; Gratzel, M. Nano Lett. 2010, 10, 2609. (11) Huynh, W. U.; Dittmer, J. J.; Libby, W. C.; Whiting, G. L.; Alivisatos, A. P. Adv. Funct. Mater. 2003, 13, 73. 412

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