Mesostructured Fullerene Electrodes for Highly Efficient n–i–p

Mesostructured Fullerene Electrodes for Highly Efficient n−i−p Perovskite Solar Cells Yufei Zhong, Rahim Munir, Ahmed Hesham Balawi, Arif D. Sheikh, Liyang Yu, Ming-Chun Tang, Hanlin Hu, Frédéric Laquai, and Aram Amassian* King Abdullah University of Science and Technology (KAUST), KAUST Solar Center (KSC), and Physical Sciences and Engineering Division (PSE), Thuwal 23955-6900, Saudi Arabia S Supporting Information *

ABSTRACT: Electron-transporting layers in today’s stateof-the-art n−i−p organohalide perovskite solar cells are almost exclusively made of metal oxides. Here, we demonstrate a novel mesostructured fullerene-based electron-transporting material (ETM) that is crystalline, hydrophobic, and cross-linked, rendering it solvent- and heatresistant for subsequent perovskite solar cell fabrication. The fullerene ETM is shown to enhance the structural and electronic properties of the CH3NH3PbI3 layer grown atop, reducing its Urbach energy from ∼26 to 21 meV, while also increasing crystallite size and improving texture. The resulting mesostructured n−i−p solar cells achieve reduced recombination, improved device-to-device variation, reduced hysteresis, and a power conversion efficiency above 15%, surpassing the performance of similar devices prepared using mesoporous TiO2 and well above the performance of planar heterojunction devices on amorphous or crystalline [6,6]-phenyl-C61-butyric acid methyl ester (PCBM). This work is the first demonstration of a viable, hydrophobic, and high-performance mesostructured electron-accepting contact to work effectively in n−i−p perovskite solar cells.

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typically elevated processing temperature of TiO2 restricts its usage to glass substrates, while its surface states can cause severe recombination problems.7−9 The ultraviolet lightsensitivity of TiO2 means that its photocatalytic activity and surface energy may also be subject to change and variability.10,11 To achieve efficient and stable printed devices in the future, finding suitable materials to partialy or fully replace TiO2 would be the first step. Organic ETLs, fullerene-based materials, or self-assembled monolayers, for instance, have been proposed as alternatives to take advantage of their lower trap state density, ease of processing, and hydrophobicity.8,12−19 Among current state-of-the-art organic ETLs, fullerene derivatives seem to gain favor in the perovskite community. Snaith and co-workers used a fullerene self-assembled monolayer as the bottom interlayer, which promoted more efficient charge extraction than when using TiO2 alone.12 However, the minuscule options of bottom ETLs available to n−i−p devices may in the near future become a bottleneck to further developments in the field. Instead, planar heterojunction p−i−n architecture devices which utilize a bottom HTL, among a myriad inorganic and organic materials available, has emerged as the predominant perovskite

merging photovoltaics have seen a drastic rise in their efficiency over the past decade. Among these, organohalide perovskite solar cells appear to be on a path to becoming “game changers” having surpassed a power conversion efficiency (PCE) of 20% while being amenable to solution-based manufacturing.1 The rapid developments in efficiency in these solar cells were achieved through improvements in the processing of polycrystalline perovskite thin films and by the development of suitable contacts for photovoltaic operational efficiency.2,3 The heterojunction between the perovskite light absorber and the selective charge extraction layers, known as electron- and hole-transporting layers (ETLs and HTLs, respectively), is crucial for obvious device operational reasons and for acting as the de facto substrate upon which the perovskite layer is grown. In its earliest and most successful implementation to date, the perovskite solar cell architecture is that of a n−i−p device whereby a mesostructured ETL was used as substrate to grow the perovskite, yielding the highest PCEs (∼20%) known to this field.1,4,5 The success of the mesostructured ETL is attributed in part to the need for a more distributed heterojunction enabling the extraction of electrons (minor carrier) which exhibit a shorter carrier lifetime.6 Tremendous effort has been devoted to improving n−i−p devices, yet no ETLs have been shown to outperform titanium dioxide (TiO2) to date. The © 2016 American Chemical Society

Received: September 19, 2016 Accepted: October 20, 2016 Published: October 21, 2016 1049

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Figure 1. Comparison of the mesoporous TiO2 (mp-TiO2) (a, g) and mesostructured PCBM (ms-PCBM) electron-transporting materials (ETMs) (b−f, h). The fabrication process of ms-PCBM is summarized in panels b−f. SEM images and contact angles (insets) are shown in panels g and h for mp-TiO2 and ms-PCBM, respectively.

device architecture.20,21 This architecture yields PCE up to 18% and often utilizes hydrophobic organic HTLs to improve perovskite thin-film growth, moisture sensitivity, and stability of devices, as well as their hysteresis.22−24 The p−i−n architecture also utilizes a fullerene ETL, which plays a crucial role in the device’s success. Huang et al. demonstrated that fullerene molecules infiltrate into grain boundaries of the perovskite film, which they appear to passivate.15 Other reports have also suggested that fullerene mixed directly into the perovskite may also work similarly.25,26 The vast majority of the p−i−n work has relied on amorphous fullerene ETLs, but a recent report has also demonstrated significant benefits in reducing the energetic disorder at the fullerene/perovskite interface by increasing the crystalline order of the fullerene ETL.15,16 Studies pertaining to the use of fullerenes both in the context of p−i−n and n−i−p devices clearly point to their suitability as contact materials. However, the full range of benefits of this material have not yet been fully realized, especially in the context of n−i−p devices, where the combinations of fullerene crystallization, mesostructuring as well as hydrophobicity may be of great benefit. Indeed, crystalline films of [6,6]-phenyl-C61butyric acid methyl ester (PCBM) are denser than their amorphous counterparts, resulting in a different polarization energy.27 PCBM, which is highly soluble in chlorinated solvents, can also be ultraviolet (UV)-cross-linked without damaging its semiconducting properties, rendering it insoluble and possibly resistant to the harsh chemical and thermal effects associated with fabrication of the organohalide perovskite film.28−30 In this Letter, we demonstrate for the first time a fullerenebased mesostructured organic bottom ETL (ms-PCBM) which

works very effectively in n−i−p organohalide perovskite (CH3NH3PbI3) solar cells. Photothermal deflection spectroscopy (PDS) measurements demonstrate reduced trap state density in the perovskite film grown on ms-PCBM as compared to films grown on mp-TiO2 in identical conditions. Microstructural investigations indicate that PCBM, whether amorphous or crystalline, promotes the growth of larger and more oriented perovskite domains, with even larger domains observed on ms-PCBM. Optical transmission attributes reduced parasitic absorption in the higher band gap msPCBM than in TiO2. These benefits combine to reduce recombination and improve the efficiency of solar cells by several power points to well above 15%, while reducing deviceto-device variability and hysteresis. This work presents for the first time a viable, high-performance, mesostructured, and hydrophobic alternative to state-of-the-art metal oxide ETMs so widely used in n−i−p perovskite solar cells. It also demonstrats the benefits of using crystalline organic materials as ETLs enhancing performance of devices, showing the first step to replace metal oxide ETL with organic materials and paving the way for fully printed devices in the future. In Figure 1a−f we illustrate the fabrication process of the mesostructured (ms-) PCBM electron transport layer and compare its morphology and contact angle of water to those of mesoporous (mp-) TiO2. In both cases, the mp-TiO2 and msPCBM are fabricated on top of a compact TiO2 (c-TiO2) layer. To make ms-PCBM, we adopted a bottom-up approach whereby we first spin-coat a layer of as-cast amorphous (a-) PCBM on c-TiO2 (Figure 1b).31 Upon annealing, the a-PCBM layer crystallizes (c-PCBM) and becomes our seed layer for the next steps (Figure 1c). We laminate a blended layer of PCBM 1050

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Figure 2. (a) Energy diagram of relevant materials used in n−i−p solar cell devices. (b) Specular transmittance of glass/c-TiO2/mp-TiO2 (black) and glass/c-TiO2/ms-PCBM (red). J−V curves of devices based on (c) mp-TiO2 and (d) ms-PCBM as ETM. (e) EQE spectra of mpTiO2 and ms-PCBM based devices (current density integrations are displayed as open squares). (f) Histogram of PCEs (average of forward and reverse scans) of devices based on mp-TiO2(black) and ms-PCBM ETMs (red).

angle measurements shown in the insets of Figure 1g,h reveal a hydrophobic surface for ms-PCBM (∼89°) in contrast to a very hydrophilic mp-TiO 2 surface (∼7°). As intended, the appearance of the cross-linked ms-PCBM ETM remains intact upon exposure to processing solvents (Figure S1), including chloroform, whereas the non-cross-linked or incompletely cross-linked ms-PCBM dissolves. This proves the ms-PCBM ETM is capable of handling subsequent perovskite thin-film processing and solar cell device fabrication.28−30 In Figure 2a we show the energy diagrams for the ETMs, CH3NH3PbI3, and Spiro-MeOTAD HTL as determined by using ultraviolet photoelectron spectroscopy (UPS) and (low-

and polystyrene (PS) on top of c-PCBM (Figure 1d) and annealed the bilayer stack so as to diffuse the PCBM toward the interface and initiate the secondary nucleation of c-PCBM nanostructures within the PS matrix (Figure 1e). The c-PCBM was cross-linked for stabilization via UV light exposure, and the PS matrix was dissolved by rinsing the sample in acetone (Figure 1f), thus finalizing the fabrication of the ms-PCBM ETM. The scanning electron microscopy (SEM) image of msPCBM exhibits larger pores delineated by nanorod-like PCBM crystals oriented isotropically in the plane of the substrate, as shown in Figure 1h, whereas mp-TiO2 (Figure 1g) is characterized by pores typically smaller than 100 nm. Contact 1051

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(see standard deviation of device figures of merit in Figure S3e−g). The results thus far point to significant advantages in using ms-PCBM as opposed to mp-TiO2. Indeed, as we will demonstrate and discuss below, these differences stem from a combination of factors, including the surface energy of hydrophobic PCBM films which can help reduce the number of nucleation sites, leading to larger crystalline domains, the crystalline order of PCBM, as well as from mesostructuring of the PCBM layer.16,23 To help deconvolute these effects, we have fabricated planar heterojunction n−i−p solar cells based on c-TiO2, c-PCBM, and cross-linked a-PCBM films, for which the device characteristics are shown in panels a, b, and c of Figure 3, respectively. The primary observation is that planar heterojunction devices perform poorly compared to their mesostructured or mesoporous counterparts, providing further evidence that mesostructuring of the ETL in the n−i−p configuration is beneficial so far as the mesostructured ETM shortens the distance for the diffusion and collection of electrons (minority carriers) from the perovskite light harvester.6 Comparison of the planar heterojunction devices among themselves reveals interesting and intriguing differences in efficiency and hysteresis for devices fabricated on the different ETLs, with c-PCBM showing highest PCE and lowest hysteresis as compared to a-PCBM and c-TiO2, while the latter two behave very similarly. PCBM has been used to passivate surface trap states of TiO2 and also to passivate the traps near the perovskite/PCBM interface.8,15,25,33 The presence of PCBM at the interface with the perovskite can therefore reduce recombination at the perovskite interface. The comparison between a-PCBM and c-PCBM is also very revealing. As Huang and co-workers previously reported, increasing the order of the PCBM ETL in p−i−n devices enhances the performance of perovskite solar cells by reducing the energetic disorder of the PCBM at the interface with the perovskite layer.16 Our work in the n−i−p configuration shows significant benefits of c-PCBM versus a-PCBM and suggests that the benefit of reduced energetic disorder is also true in the n−i−p architecture, where it can now be implemented successfully thanks to cross-linking. We delve a little deeper into the device physics and compare the recombination behaviors in the devices based on ms-PCBM and mp-TiO2 as well as c-TiO2, c-PCBM, and a-PCBM. To do so, we have performed light intensity-dependent photocurrent and photovoltage measurements. It is known that JSC should follow a power law dependence with changing light intensity (I), according to JSC ∝ Iα.34 As shown in Figure S4a (see results of planar-based devices in Figure S4b), α = 0.92 and 0.98 are observed for devices using mp-TiO 2 and ms-PCBM, respectively. As α is close to unity, weak bimolecular recombination is the dominant loss mechanism. The slightly higher α in devices based on ms-PCBM indicates less carrier loss than in devices based on mp-TiO2 in short circuit conditions. The relationship between the VOC and I is given by δVOC = kT/q for trap-free conditions. Deviation from a slope of unity indicates a competition between trap-assisted and bimolecular recombinations.1,28,35 In Figure 3d, we observe a larger slope (1.76) in devices with mp-TiO2 compared to ms-PCBM (1.34). This indicates that devices based on mp-TiO2 suffer more carrier loss due to trap-assisted recombination at open-circuit conditions. Note that we also performed the same measure-

energy) inverse photoemission spectroscopy (IPES).27,32 As the prevailing bottom ETL/ETM, mp-TiO2 exhibits a suitable conduction band minimum (CBM, 3.6 eV) and valence band maximum (VBM, 6.9 eV) to promote electron transfer while also acting as a good hole blocker. The electronic affinity (EA) of crystalline PCBM is the same as the CBM of TiO2, indicating that it is a suitable electron acceptor for the perovskite layer while also exhibiting a sufficiently deep ionization energy (IE) to block hole transfer. These energetic properties explain in part why PCBM can be a successful ETL/ETM for perovskite solar cells. The large band gap of PCBM is potentially interesting as well. We therefore compare the transmittances of the mp-TiO2 and ms-PCBM. We observe important differences in transmission in the ultraviolet and blue regions where TiO2 is more absorbing and is likely causing more light scattering than ms-PCBM, both of which should be decreased when using the latter in solar cells. The perovskite solar cells were fabricated by spin-coating the perovskite formulation in a single step on top of these ETMs and subsequently depositing Spiro-MeOTAD and Au/Ag on top as hole transport layer and electrode, respectively. Panels c and d of Figure 2 show the J−V curves for the best devices when using mp-TiO2 and ms-PCBM as ETLs, respectively (steady state parameters shown in Figure S2). The device parameters are listed in Table 1. The devices using ms-PCBM Table 1. Device Figures of Merit for n−i−p Solar Cells Using mp-TiO2 and ms-PCBM as the Bottom ETM

mpTiO2 mpTiO2 msPCBM msPCBM

scan

VOC (V)

JSC (mA cm−2)

FF (%)

PCEaverage (%)

PCEmax (%)

reverse

1.01

20.21

59

12.1 ± 0.9

13.2

forward

0.98

19.50

54

10.3 ± 1.0

11.7

reverse

1.06

20.60

65

14.2 ± 0.5

15.4

forward

1.04

20.43

63

13.4 ± 0.7

14.9

showed improvement in all figures of merit, including opencircuit voltage (VOC), short-circuit current (JSC), and fill factor (FF), giving rise to a PCE (rev. scan) up to 15.4% (ms-PCBM) versus 13.2% given by using mp-TiO2. We also observed reduced hysteresis in ms-PCBM-based devices (hysteresis, 0.5%) compared to that of mp-TiO2-based devices (hysteresis, 1.5%). In addition, we tested device stability under light soaking and found ms-PCBM-based devices show reduced light soaking effect (see Figure S3a,b). The EQE spectrum shown in Figure 2e suggests the slightly higher JSC came from a higher EQE response of the ms-PCBM-based cell between 300 and 500 nm. To further substantiate these results, we have summarized in Figure 2f the PCE distribution of 140 devices. Devices made with ms-PCBM clearly show a higher average PCE of 14.2% and a much narrower spread of device performance than mpTiO2, which shows an average PCE of 12.1%. We have also plotted the distribution of hysteresis of these devices in Figure S3c,d and confirm that hysteresis is substantially smaller in msPCBM devices. These significant differences stemming from using ms-PCBM versus mp-TiO2 as the ETM are mostly apparent in the FF (59% for mp-TiO2 vs 65% for ms-PCBM) and in the Voc (0.98 V for mp-TiO2 vs 1.06 V for ms-PCBM), which indicate reduced recombination in the ms-PCBM case 1052

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Figure 3. J−V curves of planar heterojunction n−i−p devices using (a) c-TiO2, (b) c-PCBM, and (c) a-PCBM as ETL. Light intensitydependent measurement (VOC versus I) for n−i−p devices based on (d) mp-TiO2 and ms-PCBM ETMs, as well as (e) planar c-TiO2, a-PCBM, and c-PCBM ETLs.

scattering features is shown in Figure 4e, while the azimuthal distribution of the radially integrated (110) perovskite diffraction ring is plotted in Figure 4f. The measurements reveal full conversion to the perovskite phase on either ETM. Importantly, the pole figures reveal a stronger (110) diffraction intensity as well as much more out-of-plane texturing of the perovskite (110) diffraction on ms-PCBM, whereas there is an absence of any texturing of the perovskite on mp-TiO2. A similar texturing effect has also been observed in p−i−n planar heterojunction solar cells using on CuSCN as HTL.38 The full width at half-maximum of the (110) diffraction feature is narrower (0.29 nm−1) on ms-PCBM, equivalent to an average crystallite size of ∼314 nm (Scherrer formula), and broader on mp-TiO2 (0.42 nm−1), equivalent to a crystallite size of ∼209 nm, consistent with SEM observations. The formation of larger grains and crystallites in general is an indication of fewer defects, namely, grain boundaries, which can commensurately reduce trap states and improve the optoelectronic properties of the perovskite film, resulting in less recombination losses in solar cells.2,23,39 The observation of larger crystallites on ms-PCBM than on mp-TiO2 may be due to the larger pores on the surface of the former, which allow the perovskite crystallites to grow to a larger size than on mp-TiO2, where perovskite domains tend to be confined by the smaller pores. Alternatively, the hydrophobicity of the ms-PCBM may also be affecting the microstructure of the perovskite film.23 To investigate these aspects, we revert back to the planar heterojunction devices where the perovskite film is formed on c-TiO2, c-PCBM, and aPCBM. As expected, the contact angle measurements reveal very similar differences between the planar PCBM-based and cTiO2 surfaces, in line with their mesostructured counterparts

ment on planar ETL devices (Figure 3e) and observed similar trends, suggesting trap passivation benefits exist in both planar and mesostructured PCBM devices and c-PCBM outperforms both c-TiO2 and a-PCBM in this respect. The device-level analysis of performance and recombination for mesostructured and planar n−i−p devices points to a potential impact of the choice of ETL on the optoelectronic properties of the polycrystalline perovskite film itself. We investigate the quality of the semiconductor film by probing its tail states by measuring the sub-band gap absorption using photothermal deflection spectroscopy (PDS) (Figure 4a).36,37 The Urbach tail related to trap states in the perovskite film grown on ms-PCBM is steeper than for films grown on mpTiO2. The absorption tails are analyzed to extract the Urbach energy of the polycrystalline perovskite films grown on mpTiO2 (25.7 meV) and ms-PCBM (20.8 meV), clearly revealing a significantly less defective perovskite film on ms-PCBM than on mp-TiO2. We have also performed PDS measurements on perovskite films grown on planar ETLs and obtained similar results, confirming the intrinsic advantage of using PCBM as bottom ETL instead of TiO2 (see Figure S4). To further explain the improved optoelectronic properties of the polycrystalline perovskite layer, we have examined the morphology and microstructure of the films formed atop these ETMs. As shown in Figure 4b,c, the perovskite film grown on ms-PCBM is characterized by larger domain size ranging from 150 to 450 nm, whereas the one grown on mp-TiO2 shows a larger distribution of domain sizes from 50 to 400 nm (cross section shown in Figure S5). The microstructure of these films was probed via two-dimensional grazing incidence wide-angle X-ray scattering (GIWAXS) (Figure 4e, and see data for mpTiO2 in Figure S6). The azimuthally integrated intensity of 1053

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Figure 4. (a) PDS measurements of perovskite films grown on mp-TiO2 and ms-PCBM. Lines are linear fits to ln(α) according to the Urbach equation in sub-band gap region of absorption. SEM images (top view) of perovskite films grown on (b) mp-TiO2 and (c) ms-PCBM. (d) Two-dimensional GIWAXS image of the perovskite film grown on ms-PCBM showing the (110) diffraction ring. (e) Integrated intensity plots along qr showing differences in the intensity of the perovskite (110) diffraction. (f) Pole figure (qr= 10 nm−1) showing differences in the texture of the (110) diffraction.

benefits also reduce the hysteresis as well as the device-todevice performance spread. This work presents the most successful organic ETL to be implemented to date for n−i−p based perovskite solar cells, taking advantage of the remarkable potential of organic semiconductors and carbon nanomaterials, such as fullerenes and their derivatives.

(see contact angle in Figure S7). The SEM images of the perovskite films fabricated on c-TiO2, c-PCBM, and a-PCBM are shown in panels a, b, and c of Figure S8, respectively. They reveal a very similar trend in crystallite size, with more uniformly large features observed on both c-PCBM and aPCBM and a distribution of small and large crystals observed on c-TiO2. These results provide a strong indication that the major differences in the perovskite microstructure are due to differences in the surface properties of the PCBM, namely, its surface energy, rather than to differences due to crystallinity of PCBM alone. We observe larger grains on top of ms-PCBM than on a-PCBM or c-PCBM, indicating there is also a microstructural benefit to growing the perovskite film on a hydrophobic and mesostructured ETM rather than on a planar one. We have demonstrated a novel mesostructured electrontransporting material by crystallizing, mesostructuring, and cross-linking molecular fullerene. The resulting ETM is hydrophobic and works very effectively as bottom ETM in n−i−p perovskite solar cells. The ms-PCBM ETM combines the benefits of a hydrophobic growth surface and crystallinity with those of improving the microstructure and semiconducting properties of the polycrystalline perovskite film itself, as demonstrated by reduced Urbach energy. These benefits are responsible for decreasing recombination in the solar cells and lead to significant enhancement of the PCE up to 15%; these

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsenergylett.6b00455. Experimental details, characterization of ETL materials (SEM, GIWAXS, contact angle), and device performance using planar ETLs (device data, light intensity measurement, etc.) (PDF)

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Notes

The authors declare no competing financial interest. 1054

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ACKNOWLEDGMENTS This work was supported by the King Abdullah University of Science and Technology. A.H. Balawi and F. Laquai thank K. Vandewal and M. Baier for contributing to the setup for photothermal deflection spectroscopy (PDS). Part of this work was performed at D-line at the Cornell High Energy Synchrotron Source (CHESS) at Cornell University. CHESS is supported by NSF and NIH/NIGMS via NSF Award DMR1332208. Dr. Detlef-M. Smilgies and Dr. Ruipeng Li from CHESS are thanked for their assistance with beamline setup for the GIWAXS measurements.

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DOI: 10.1021/acsenergylett.6b00455 ACS Energy Lett. 2016, 1, 1049−1056