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Organic–inorganic bismuth (III)-based material: a lead-free, air-stable and solution-processable light-absorber beyond organolead perovskites

Miaoqiang Lyu1, Jung-Ho Yun1, Molang Cai2, Yalong Jiao2, Paul V. Bernhardt3, Meng Zhang1, Qiong Wang1, Aijun Du2, Hongxia Wang2, Gang Liu4, and Lianzhou Wang1 (*) Nano Res., Just Accepted Manuscript • DOI: 10.1007/s12274-015-0948-y

http://www.thenanoresearch.com on November. 18, 2015

© Tsinghua University Press 2015

Just Accepted

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Nano Research DOI 10.1007/s12274-015-0948-y

www.theNanoResearch.com∣www.Springer.com/journal/12274 | Nano Research

64 Nano Res.

Organic-inorganic bismuth (III)-based

material: a lead-free, air-stable and

solution-processable light-absorber beyond

organolead perovskites

Miaoqiang Lyu,a Jung-Ho Yun,a Molang Cai, b

Ya-Long Jiao,b Paul V. Bernhardt,c Meng Zhang,a

Qiong Wang,a Aijun Du,b Hongxia Wang,b Gang

Liu,d and Lianzhou Wang a*

a. Nanomaterials Centre, School of

Chemical Engineering and AIBN, The

University of Queensland, St Lucia, Brisbane,

QLD 4072, Australia. Email:

[email protected]. b. School of Chemistry, Physics and

Mechanical Engineering, Science and

Engineering Faculty, Queensland University

of Technology, Brisbane, QLD 4001,

Australia.. c. School of Chemistry and Molecular

Biosciences, The University of Queensland,

Brisbane 4072, Australia d. Advanced Carbon Division, Institute of

Metal Research Chinese Academy of Sciences

(IMR CAS), 72 Wenhua Road, Shenyang,

China,110016

Organic-inorganic methylammonium bismuth iodide ((CH3NH3)3Bi2I9) is

developed and applied in solution-processable solar cells, which shows

advantages of non-toxicity and air-stability compared to lead/tin based

organo-perovskite materials.

Prof. Lianzhou Wang, http://www.nanomac.uq.edu.au/lianzhou-wang


2 Nano Res.

material, Lead-free

1 Introduction

Solar cells based on organolead halide perovskites have seen rapid progress as one of the most promising photovoltaic technologies, owing to their solution-processability, low-cost and high efficiencies in the past few years [1-4]. To date, over 20% efficiencies have been achieved based on lead-containing organic-inorganic halide perovskites [5, 6]. However, serious concerns on the use of toxic lead as well as the instability of these perovskites in humid conditions have triggered efforts towards developing non-toxic and stable organic-inorganic alternatives, which is critically important for their market acceptance in future commercialization [3, 7, 8]. To address the toxicity issue, for example, tin-based halide perovskites, such as CsSnI3, CH3NH3SnI3 and CH3NH3SnI3-xBrx, have been reported recently [9-11]. Although over 5% efficiencies have been achieved, those tin-based perovskites are not stable in air due to oxidation of Sn2+ to Sn4+ and this oxidation process occurs so rapidly that all device assembling and characterizing processes have to be done under an inert atmosphere [11]. The compound (CH3NH3)2CuCl4-xBrx was recently reported, showing satisfactory stability in the air with the presence of Cl. By adjusting the ratio between Cl and Br, the band-gap of this layered perovskite material can range from 2.48 eV to 1.80 eV. However, a low efficiency of less than 0.02% was achieved [12]. To address the moisture instability, Smith et al. reported a layered hybrid perovskite with markedly enhanced stability in ambient humidity [13]. Moreover, bromide-doping has been found to improve the stability of the hybrid perovskites [14]. However, all the above-mentioned reports are still based on

lead-containing perovskites and thus the toxicity problem still exists. Therefore, the toxicity of the

lead-containing perovskites and instability of their tin-based counterparts are still two open research challenges in this research field [7]. While organic-inorganic perovskites are among the largest families of hybrid materials, there are a wide range of other metal halide candidates exhibiting interesting properties. Organic-inorganic bismuth iodide based hybrid materials have attracted interest because of their potential semiconducting character, rich structural diversity and interesting electronic and optical properties, for which the 6s2 lone pair of the Bi3+ plays an important role in the various material properties [15-18]. Meanwhile, organic groups play a key role in determining connection manners of adjacent bismuth iodide octahedra in the crystal structure and hence influencing the electronic and optical properties of the bismuth hybrids. Combined with different organic groups, bismuth iodide octahedra can form mononuclear or polynuclear inorganic frameworks of corner-, edge- and face-sharing bismuth halogenoanions, leading to a wide range of structures and tunable electronic or optical properties [19]. Methylammonium bismuth iodide (CH3NH3)3Bi2I9 (MBI) comprises two face-sharing bismuth iodide octahedra and has been investigated previously, mainly focusing on its low-temperature phase transitions, dielectric properties and optical studies [17, 18, 20]. MBI can be considered as an analogue of the methylammonium lead/tin iodide by simply replacing the lead/tin cations with bismuth (III), but it does not adopt a perovskite structure. Motivated by the non-toxicity of the bismuth element, we explore the possibility of applying MBI as a new light-absorber alternative to organolead perovskite in the solution-processable solar cells. Previous studies on this material are very limited

and most of them focused on single crystals of MBI [17, 18, 20]. However, for practical applications in

Address correspondence to Lianzhou Wang, [email protected]


3Nano Res.

electric or optoelectric devices, techniques for preparing MBI films on certain substrates are more desirable, yet little has been published in this endeavour. Herein, we report a simple, low-temperature solution process to prepare MBI films. Detailed material and electronic properties are studied by both experimental and theoretical techniques. Moreover, we present the first demonstration of applying organic-inorganic methylammonium bismuth iodide in solution-processable solar cells. Even though the preliminary power conversion efficiency of the devices is low, the MBI shows its advantages of non-toxicity, stability in ambient air and humid environment and low-temperature solution-processibility, which may address both the toxicity issue of lead-based perovskites and instability concern of tin-based ones.

2 Experimental

Growth of single crystal (CH3NH3)3Bi2I9: Typically, 0.2384 g of MAI and 0.5897 g of BiI3 were added to 2 ml of anhydrous methanol and the solution was mixed by ultra-sonication for ~ 0.5 h. Clear saturated (CH3NH3)3Bi2I9 solution (0.5 ml) in methanol was transferred to another clean vial and the solvent was allowed to evaporate slowly overnight in the fume-hood at room temperature, leading to the formation of MBI single crystals on the inner wall of the vial. Fabrication of (CH3NH3)3Bi2I9 films: The (CH3NH3)3Bi2I9 precursor solution was prepared by mixing 0.2384 g of MAI and 0.5897 g of BiI3 in 1 ml of N, N-dimethylformamide, which was filtered by PTFE syringe filters (0.45 µm) before use. One-step spin-coating method was developed to deposit the MBI film on substrates (such as, FTO, glass slide and FTO covered with TiO2 mesoporous layers), in which 50 µL of (CH3NH3)3Bi2I9 solution was dispensed onto the substrate and spin-coated at 4000 RPM for 30s with a ramping rate of 4000 RPM/s, followed by heating at 100 ℃ for 10 min on a hotplate. A reddish film was achieved after heating treatment (Fig S1 in the Electronic Supplementary Material (ESM)). Assembly of solar cells: FTO (fluorine-doped tin oxide) coated glass slides (FTO22-7, Yingkou OPV

Tech New Energy Co. Ltd) were firstly pre-patterned by removing part of the conductive layer with 2 M HCl aqueous solution and zinc powder. The etched substrates were cleaned in acetone, ethanol and isopropanol, subsequently. After oxygen plasma treatment (RIE, ATX-600, 30 W, 10 min), a thin TiO2 blocking layer was deposited by spray pyrolysis at 500 ℃ using 0.2 M isopropanol solution of titanium diisopropoxide bis(acetylacetonate) as the precursor solution. For planar-structured devices, (CH3NH3)3Bi2I9 was directly deposited on the FTO/TiO2 blocking layer. For mesoporous-structured devices, a ~ 1.8-µm-thick TiO2 mesoporous layer was deposited onto the FTO/TiO2 blocking layer substrates by doctor blading. Detailed method for preparation of TiO2 mesoporous layer can be found in the Supporting Information section. After annealing treatment at 500 ℃ for 60 min, the TiO2 mesoporous layer covered FTO substrates were used for (CH3NH3)3Bi2I9 film deposition. Pristine P3HT (MW 54000-75000, Sigma-Aldrich, 15mg/ml in 1, 2-dichlorobenzene) was used as the hole-transporting layer, which was spin-coated at 2000 RPM for 30s after 60 soaking. Then, the coated substrates were further annealed at 100 ℃ for 15 min to remove the remaining solvent. Finally, a ~ 60 nm gold contact electrode was deposited by e-beam evaporation (Temescal FC-2000). The cell area was defined by a metal mask with an aperture area of 7.068 mm2.

3 Results and discussion

Single crystals of MBI have been synthesized and investigated previously [18]. However, detailed crystallographic data, as well as atomic coordinates, bond distances and bond angles were still missing, which are crucial for structural modelling and computational chemistry. To better understand the fundamental properties of the MBI, we developed a simple method to prepare MBI single crystals by slowly evaporating a methanol solution of BiI3 and CH3NH3I (1.5:1 in molar ratio). Although the solubility of BiI3 and CH3NH3I in methanol is not very high, slowly evaporation of its saturated solution gives MBI single crystals with a diameter ranging from 100 to 200 µm (Fig 1. (a)), which is large enough for single X-ray diffraction analysis. Also, it can be seen that the as-prepared MBI single


4 Nano Res.

crystal displays a regular hexagonal shape. According to the results from the single X-ray diffraction data, the crystal parameters of MBI can be derived. The main crystallographic data of MBI has been presented in the Table 1, and a more complete table as well as the information concerning the atomic coordinates, bond distances and bond angles are available in the Supporting Information (Table S1-S4). MBI crystals belong to the hexagonal system with space group of P 63/mmc. Its lattice parameters are in accordance to the report by Jakubas and co-workers [18]. The theoretical XRD spectrum of the MBI was calculated by Mercury 3.3 software (The Cambridge Crystallographic Data Centre (CCDC)) based on the results of single crystal X-ray diffraction measurement, as shown in Fig 1 (b). To verify the property of the MBI, we developed a facile, low-temperature and solution-processable technique for depositing the MBI film on fluorine-doped tin oxide (FTO) substrates or glass substrates, inspired by the fabrication technique of organolead perovskites.1 The precursor solution for MBI film deposition is composed of BiI3 and CH3NH3I (1.5:1 in molar ratio) in anhydrous N, N-dimethylformamide, which shows a dark red colour (Fig S1). The MBI film was prepared by spin-coating the precursor solution on substrates, followed by annealing treatment on a hot-plate at a low temperature (100 ℃ ). It is interesting to find that the organic-inorganic bismuth-based films can be deposited at a temperature as low as that of the organolead perovskites. Fig 2 shows the surface morphology of the MBI film fabricated on the FTO/TiO2 blocking layer (BL) substrate and glass substrates, respectively. Hexagonal shaped MBI crystals are uniformly distributed on the FTO / TiO2 BL substrate, showing irregular orientation due to relatively high roughness factor of the FTO conductive layer. On the contrary, regular hexagonal shaped MBI crystals can be observed on the glass slide, due to the low roughness factor of the glass surface. The thickness of the MBI crystals is around 500nm, while the diameter of the single plate is up to several micrometres. X-ray photoelectron spectroscopy (XPS) was performed to confirm the elemental composition and valence of Bi and I of the MBI film (Fig 3.). According to the survey spectrum of the MBI film, the detected elements include Bi, I, N, O, Sn and C (Fig S2). The O and Sn should be originated from

the FTO glass. High-resolution scanned XPS spectra for I 3d, Bi 4f, N 1s and C 1s are shown in Fig 3 (a-d). Peaks for Bi 4f7/2 (158.8 eV) and 4f5/2 (164 eV), and I 3d5/2 (618.8 eV) and 3d3/2 (630.3 eV) can be attributed to characteristic signals from the Bi3+ and I- species, respectively. Bismuth iodide adopts a face-sharing bi-octahedral cluster in the MBI crystal, which is in accordance to the previous report [18]. However, the [Bi2I9]3- bi-octahedral anions are isolated from randomly disordered CH3NH3+ cation. In order to investigate the fundamental electronic band structure of the MBI, an optimized crystal model was established based on the single crystal XRD results, aiming to fix the position of CH3NH3+ cation and representing the possibly lowest energy state of this crystal (Fig 4 (a) and (b)). Fig 4 (c) shows the calculated electronic band structure based on the aforementioned crystal model, suggesting that the MBI crystal has an indirect band gap of ~2.25 eV, comprising filled I-5p orbits and empty hybrid I-6p/Bi-6p orbits for the valence band maximum (VBM) and conduction band minimum (CBM), respectively. The calculated band gap is very close to the measured value (2.11 eV). The valence and conduction bands are well distributed in energy, which can be observed from the total density of states (DOS) in Fig 4 (d). The alignment of VBM and CBM with respect to the vacuum level for the MBI crystal was theoretically calculated as well based on a similar method to the previous report [22]. The calculated VBM and CBM are -5.70 eV and -3.45 eV, respectively. Also, Fermi energy level of the MBI crystal is predicted to be -5.25 eV based on the calculation results. In order to experimentally estimate the energetic position of the VBM with respect to the vacuum level, ultraviolet photoemission spectroscopy (UPS) was performed. The VBM position is usually determined by extrapolating the secondary electron cut-off (SECO) region, and the VBM of MBI is located around -5.90 eV (Fig 6). Therefore, the CBM of the MBI crystal can be calculated based on the VBM value and the optical band gap of the MBI, which is determined to be -3.78 eV. The experimental determined VBM, CBM with respect to the vacuum level are very close to that of the theoretical calculation results. Kelvin probe force microscopy (KPFM) was performed to determine the work function (or Fermi-level) of the MBI crystals. The MBI crystals


5Nano Res.

display dark red colour, as shown in the colourful microscope image equipped on the KPFM (Fig 5 (a)). A regular hexagonal shaped MBI crystal was also evidently confirmed by the scanned atomic force microscope (AFM) images in Fig 5 (b). Surface potential image was performed on the same crystal. The surface potential (i.e. contact potential difference (CPD)) disperses uniformly on the MBI crystal (Fig 5 (c)) and the averaged CPD is determined to be around -0.42±0.021 V from the cross-sectional profile (Fig 5 (d)). The work function of the MBI crystal can be calculated according to Equation 1 [23].

Equation 1 Here, φtip corresponds to the work function of a Pt-coated conductive cantilever probe, which is calibrated with highly oriented pyrolytic graphite (HOPG). Averaged φtip was determined to be 5.1±0.05eV. Hence the work function of the MBI sample is ~5.52±0.054 eV, which is comparable to that of the computational calculation (-5.25 eV). Based on the above-mentioned energy level study of the MBI crystals, an energy band diagram was summarized together with that of TiO2, P3HT and Au layers (Fig 6). Although the MBI crystal shows p-type conductive characteristic, its energy band level exhibits its possible potential as a light-absorber in solar cells with a device structure of FTO/TiO2 BL/TiO2 mesoporous layer (mpl)+MBI+P3HT/Au. Therefore, the possibility of applying MBI as a light-absorber in solution-processable solar cells was investigated. The thicknesses of the TiO2 BL and TiO2 mpl are around ~ 100 nm and ~ 1.8 µm, respectively (Fig 7 (a)) and the MBI material can infiltrate into the TiO2 mpl very uniformly, as is evident from its SEM morphology (Fig S5). The elemental composition of the MBI/TiO2 film was further confirmed by the Energy Dispersive Spectroscopy (EDS) characterization (Fig S6 and Table S5). Current-voltage (I-V) measurement was carried out and a preliminary efficiency of 0.190% was achieved with a photocurrent density (Jsc) of ~1.157 mA/cm2, an open-circuit voltage (Voc) of 0.354 V and a fill factor (FF) of 0.464. Since current-voltage hysteresis is one of the well-known issues existing in the lead-based perovskite solar cells, we also investigated the hysteresis of the thus-assembled

solar cells [7]. We measured the hysteresis of a device by obtaining I-V curves using both forward and reverse scan and interestingly, little I-V hysteresis was observed in the MBI film based solar cells. Moreover, we also demonstrated that a higher Voc of 0.510 V can be achievable when a planar structure (FTO/TiO2 BL/MBI film/P3HT/Au) was adopted (Fig S7). It is interesting and encouraging to discover that such a trivalent metal cation based organic-inorganic hybrid material shows photovoltaic effect and little I-V hysteresis, which provides us with a family of new possible candidates to replace lead/tin based perovskites in solution-processable solar cells with various trivalent metal cation according to the periodical table. The photocurrent of the MBI based solar cell was not very high, thus it is important to determine the contribution from the P3HT, which has been used as an active layer in polymer solar cells. We then assembled devices with a structure of FTO/TiO2 BL/TiO2 mpl+P3HT/Au and measured their I-V characteristic (Fig S8). The P3HT layer indeed contributes to the overall performance of the device. An efficiency of 0.061% was achieved based on the device structure of FTO/TiO2 BL/TiO2 mpl+P3HT/Au, while the efficiency of the device combined with MBI film is around 0.190%. The main contribution from the P3HT comes from the photocurrent density, which is around 41% of the total Jsc of the best-performing device. The incident photon-to-current conversion efficiency (IPCE, also written as EQE (External quantum efficiency)) results indicates that the device using P3HT only gives relatively low EQE value over 400 nm to 500 nm, while inclusion of the MBI film as the light-absorber significantly enhances the EQE of the device in the range from 375nm to 500 nm, confirming the contribution of the MBI to the overall photocurrent (Fig S9.). The photocurrent from the MBI film is relatively low in the assembled solar cells, which might result from its inherent indirect large band gap. In addition, the excitons binding energy of the MBI crystal is estimated to be larger than 300 meV, which is much larger than the lead-based organic-inorganic perovskite materials (circa 40 meV) [13, 20]. The large excitons binding energy may partially explain the low open-circuit voltage in our devices because of photovoltage loss in dissociating excitons into holes and electrons [7].


6 Nano Res.

Further, we studied the stability of the MBI film in the ambient air with an average humidity level of ~50%. The MBI film exhibited excellent stability over 40 days upon continuous exposure to the humidity and ambient air at room temperature, as is evident by the XRD patterns (Fig 7d). Time-dependent performance of the solar cells using the structure FTO/TiO2 BL/TiO2 mpl+MBI+P3HT/Au was also evaluated It is evident that the device showed very limited performance degradation even after storage in the air for 21 days (Fig S10 and Table S6). Thus it is quite encouraging to observe this behaviour, by simply changing the metal cation of the organic-inorganic hybrid material from tin/lead to bismuth, the moisture stability can be significantly improved. We believe the findings can inspire more research to further explore lead-free and stable organic-inorganic hybrid materials by using different metal cations. Although some research work has been performed based on the MBI single crystal, to our knowledge, rare attention was paid to develop MBI films on certain substrates. Apart from fabricating the MBI films on the TiO2 mesoporous layer, we also managed to deposit the films on bare glass slides in order to measure its conductivity and Hall Effect. It is worth noting that oxygen plasma is necessary as a pre-treatment technique to facilitate the spreading of the MBI precursor solution and enhance the film uniformity later on. Based on the MBI film fabricated on the glass slide, conductivity was measured by a Van der Pauw four-point probe and its average conductivity is around 0.0083 S/cm and resistivity is ~121.04 Ω∙cm. Hall effect measurement was performed as well to determine conducting type, carrier concentration and carrier mobility of the MBI film. According to the measurement, the Hall coefficients are positive, which confirms the positive sign of carriers (p-type) for the solution-processed MBI film. The p-type characteristic of the MBI film was further confirmed qualitatively by Seebeck effect measurement. In fact, the computational calculation as well as the experimental determination of energy levels also suggests the p-type characteristic of the MBI film (Fig 6), which is in consistence with that of the Hall Effect results. Based on the Hall coefficient, the carrier concentrations and carrier mobility can be determined based on Equation 2 and 3,

respectively.

Equation

2 Equation 3 In the equations, n is the carrier concentration, RH is the Hall coefficient, e is the elementary charge of an electron, Is is the electric current, d is the film thickness, Bz is the magnetic field, µ is the carrier mobility and ρ is the resistivity of the sample. The calculated carrier concentration of the MBI film is ~1016 cm-3, which is much lower than that of pristine CsSnI3 (1019 cm-3) and comparable with SnF2 doped CsSnI3 (1017 cm-3) [10, 25], another typical lead-free perovskite material for solar cells. However, the intrinsic carrier concentration of the MBI is 7 orders of magnitude higher than CH3NH3PbI3 (109 cm-3) [25]. The high background carrier densities in the light-absorber may contribute to the bulk recombination, which may in turn reduce the open-circuit voltage (Voc) of the solar cells based on this material [10]. Due to the high intrinsic background carrier densities, the reported Voc is only 0.01 V for the pristine CsSnI3

and 0.24 V for SnF2 doped CsSnI3, much lower than that of achieved by planar MBI film based solar cells (0.510 V) in our case [10]. Such a Voc tendency is consistent with the background carrier densities measured by the Hall Effect technique, suggesting that reduce the background carrier concentration may provide an alternative route to further improve the performance of solar cells based on the MBI. The carrier mobility of the MBI film is around 1 cm2V-1s-1, which is 2~3 magnitude lower than that of CH3NH3SnI3 (~2320 cm2V-1s-1) and p-type CsSnI3 (~520 cm2V-1s-1) [24]. The low carrier mobility in the MBI film may be responsible to the relatively low performance in MBI-based solar cells. Doping process, which has been demonstrated to be very effective in CsSnI3, might offer a solution to regulate the carrier concentration and carrier mobility for further improve the device performance of the MBI film based devices [10]. Apart from the high carrier concentration and low carrier mobility, the relatively large indirect band gap of the MBI should be one of the performance limitations. Compared with solar cells based on the CsSnI3, one of the typical lead-free


7Nano Res.

materials for solution-processable perovskite solar cells, the device performance was mainly contributed by its high photocurrent density (22.7 mA/cm2) due to its extremely low band gap (1.3 eV) [10, 26]. Therefore, reducing the band-gap by developing other bismuth based organic-inorganic materials may be a promising strategy to further improve its solar cell performance and to this end, the highly tunable band-gap correlated with its organic groups has already provided many possible candidates [15, 16, 27-29]. Considering the advantages of non-toxicity, stability in the ambient humid air and facile, low-temperature solution-processibility, the bismuth-based organic-inorganic hybrid materials presented here may open up a new alternative strategy towards lead-free, low cost solar cells.

4 Conclusions

In summary, we have presented the first demonstration of applying organic-inorganic methylammonium bismuth iodide in replace of lead/tin based perovskite materials in solution-processable solar cells. The organobismuth based material shows the advantages of non-toxicity, ambient stability and low-temperature solution-processibility, which may provide promising solution to address the toxicity and stability challenges in perovskite solar cells. The devices yielded a preliminary efficiency of ~0.190% with an open-circuit voltage of ~ 0.510 V in its planar structure. Reducing the background carrier concentration and its optical band-gap may be two effective strategies for further improving the device performance. Also, this is the first demonstration of trivalent metal cation based, solution-processable organic-inorganic hybrid light-absorber for solar cell application, which may inspire more research work to develop and apply organic-inorganic hybrid materials beyond divalent metal cations (Pb (II) and Sn (II)) for solar energy utility. Furthermore, the facile low-temperature solution methods developed in this work for preparing methylammonium bismuth iodide (MBI) both in single-crystal and films on the substrates

may facilitate further fundamental research and practical applications based on this material.

Appendix A: Supplementary material CCDC No. 1059280 includes the supplementary crystallographic data for methylammonium bismuth iodide (MBI), which can be available free of charge via http://www.ccdc.cam.ac.uk/conts/retrieving.html, or from the Cambridge Crystallographic Data Centre, 12 Union Road, Cambridge CB2 1 EZ, United Kingdom; or by email: [email protected].

Acknowledgements

Financial support from CRC-Polymers programs and ARC DPs and FT programs is acknowledged. This work was performed in part at the Queensland node of the Australian National Fabrication Facility and Australian Microscopy & Microanalysis Research Facility at the Centre for Microscopy and Microanalysis, The University of Queensland (UQ). M.Q. Lyu acknowledges the support from IPRS Scholarship. Prof. Jose Alarco is acknowledged for assistance in Hall Effect measurement and beneficial discussion on Hall data analysis. Mr. Shengli Zhang from QUT is acknowledged for Seebeck effect measurement and beneficial discussion. Dr. Xiaojing Zhou from UoN is acknowledged for UPS analysing. Dr. Ravi Chandra Raju Nagiri and Eunji Yoo are acknowledged for discussion on Hall Effect data analysis. H.W thanks the financial support from Australian research council (ARC) Future Fellowship scheme (FT120100674). Mr. Ripon Bhattacharjee from QUT is acknowledged for kind help on TiO2 substrate preparation.

Electronic Supplementary Material: Supplementary material (please give brief details, e.g., further details of the annealing and oxidation procedures, STM measurements, AFM imaging and Raman spectroscopy measurements) is available in the online version of this article at http://dx.doi.org/10.1007/s12274-***-****-* (automatically inserted by the publisher).

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[16] Mitzi, D.B., Brock, P. Structure and optical properties of several organic-inorganic hybrids containing corner-sharing chains of bismuth iodide octahedra. Inorg. Chem. 2001, 40, 2096-2104.

[17] Kawai, T., Ishii, A., Kitamura, T., Shimanuki, S., Iwata, M., Ishibashi, Y. Optical Absorption in Band-Edge Region of ( CH3NH3)3Bi2I9 Single Crystals. J. Phys. Soc. Jpn. 1996, 65, 1464-1468.

[18] Jakubas, R., Zaleski, J., Sobczyk, L. Phase transitions in (CH3NH3)3Bi2I9 (MAIB). Ferroelectrics 1990, 108, 109-114.

[19] Fisher, G.A., Norman, N.C. The Structures of the Group 15 Element(III) Halides and Halogenoanions. Adv. Inorg. Chem. 1994, 41, 233-271.

[20] Kawai, T., Shimanuki, S. Optical Studies of (CH3NH3) 3Bi2I9 Single Crystals. Phys. Stat. Sol. (b) 1993, 177, K43-K45.

[21] Murphy, A. Band-gap determination from diffuse reflectance measurements of semiconductor films, and application to photoelectrochemical water-splitting. Sol. Energy Mater. Sol. Cells 2007, 91, 1326-1337.

[22] Cai, Y., Zhang, G., Zhang, Y. W. Layer-dependent Band Alignment and Work Function of Few-Layer Phosphorene. Sci. Rep. 2014, 4, 6677.

[23] Melitz, W., Shen, J., Kummel, A.C., Lee, S. Kelvin probe force microscopy and its application. Surf. Sci. Rep. 2011, 66, 1-27.

[24] Ren, S., Chang, L.Y., Lim, S.K., Zhao, J., Smith, M., Zhao, N., Bulovic, V., Bawendi, M., Gradecak, S. Inorganic-organic hybrid solar cell: bridging quantum dots to conjugated polymer nanowires. Nano Lett. 2011, 11, 3998-4002.

[25] Stoumpos, C.C., Malliakas, C.D., Kanatzidis, M.G. Semiconducting tin and lead iodide perovskites with organic cations: phase transitions, high mobilities, and near-infrared photoluminescent properties. Inorg. Chem. 2013, 52, 9019-9038.

[26] Chen, Z., Wang, J.J., Ren, Y., Yu, C., Shum, K. Schottky solar cells based on CsSnI3 thin-films. Appl. Phys. Lett. 2012, 101, 093901.

[27] Hrizi, C., Chaari, N., Abid, Y., Chniba-Boudjada, N., Chaabouni, S. Structural characterization, vibrational and optical properties of a novel one-dimensional organic–inorganic hybrid based-iodobismuthate(III) material, [C10H7NH3]BiI4. Polyhedron 2012, 46, 41-46.

[28] Leblanc, N., Mercier, N., Allain, M., Toma, O., Auban-Senzier, P., Pasquier, C. The motley family of polar compounds (MV)[M(X5−xX′x)] based on anionic chains of trans-connected M(III)(X,X ′ )6 octahedra (M=Bi, Sb; X, X′=Cl, Br, I) and methylviologen (MV) dications. J. Solid State Chem. 2012, 195, 140-148.

[29] Liu, B., Xu, L., Guo, G.-C., Huang, J.-S. Three inorganic–organic hybrids of bismuth (III) iodide complexes containing substituted 1, 2, 4-triazole organic components with charaterizations of diffuse reflectance spectra. J. Solid State Chem. 2006, 179, 1611-1617.


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Fig. 1 Optical microscope image of single crystals of MBI; (b) Powder XRD patterns of MBI film: measured curve (red) and calculated curve from the single-crystal X-ray structure (black); (c) Reflectance spectrum of MBI film prepared on FTO/TiO2 mesoporous film and (d) Calculated spectrum according to the Kubelka-Munk equation.

Fig. 2 SEM images of MBI film deposited on FTO/TiO2 BL (a) (b), and a bare glass substrate (c) (d). Images (b) and (d) are high magnification view of (a) and (c), respectively. Scale bars on the image (a-d) are 20 µm, 2 µm, 20 µm and 5 µm, respectively.

Fig. 3 XPS spectra of a spin-coated MBI film on a FTO substrate: (a) I 3d, (b) Bi 4f, (c) N 1s and (d) C 1s.

Fig. 4 First-principle calculations. (a) (b) Three-dimensional crystal structures of the MBI viewed from different directions. (c) The calculated electronic band structure along highly symmetric points. (d) Total density of states (DOS) and contribution from selected orbitals for the MBI.


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Fig. 5 KPFM measurement of the MBI crystal. (a) Colored microscopic optical image of MBI crystals. (b) Top-view morphology of a MBI crystal by AFM. (c) Surface potential image measured by KPFM and (d) its corresponding cross-sectional profile of the surface potentials plotted along the direction marked with the red line (d). The scale bars for (a), (b) and (c) are 20 µm, 2 µm and 20 µm, respectively.

Fig. 6 Energy band levels of the device using (CH3NH3)3Bi2I9 (MBI) as the light-absorber. VBM value of MBI is determined by the UPS result and the CBM value of MBI is calculated based on the optical band-gap of MBI. The energy levels for TiO2, P3HT and Au are cited from previous references.24-26 MBI* corresponds to the computational calculated energy band level for the MBI. Two dash lines denoted on the band diagram of MBI and MBI* are assigned to the Fermi level of MBI crystals derived from KPFM and computational calculation, respectively.

Fig. 7 (a) SEM image of a cross-sectional morphology of the best-performing device; (b) J-V curve of the best performing device; (c) J-V curves of a device scanned from forward and reverse; (d) XRD patterns of the MBI film on the FTO substrate under different exposure time to the ambient air.

Table 1 Crystallographic data for MBI (CH3NH3)3Bi2I9) a

Empirical formula C3H18Bi2I9N3

Formula weight 1656.26

Wavelength, Å 0.71073 (Mo Kα)

Crystal system Hexagonal

Space group P 63/mmc

a= b, c, Å 8.4668(6), 21.614(2)

a, b, g, ° 90, 90, 120

Volume, Å3 1341.9 (2)

Z, ρcalcd (g cm-3) 2, 4.099

Final R indices [I>2sigma(I)]

b R1 = 0.0689, wR2 = 0.1672

R indices (all data) b R1 = 0.0752, wR2 = 0.1723 a Obtained by monochromated Mo Kα (λ = 0.71073 Å) radiation

b ,


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Electronic Supplementary Material

Organic-inorganic bismuth (III)-based material: a lead-free, air-stable and solution-processablelight-absorber beyond organolead perovskites

Miaoqiang Lyu,1 Jung-Ho Yun,1 Molang Cai,2 Yalong Jiao,2 Paul V. Bernhardt,3 Meng Zhang,1 Qiong Wang,1

Aijun Du,2 Hongxia Wang,2 Gang Liu,4 and Lianzhou Wang 1(*)

1Nanomaterials Centre, School of Chemical Engineering and AIBN, The University of Queensland, St Lucia, Brisbane, QLD 4072, Australia. 2 School of Chemistry, Physics and Mechanical Engineering, Science and Engineering Faculty, Queensland University of Technology, Brisbane, QLD 4001, Australia. 3School of Chemistry and Molecular Biosciences, The University of Queensland, Brisbane 4072, Australia. 4Advanced Carbon Division, Institute of Metal Research Chinese Academy of Sciences (IMR CAS), 72 Wenhua Road, Shenyang, China, 110016

Supporting information to DOI 10.1007/s12274-****-****-* (automatically inserted by the publisher)

Part 1. Detailed Experimental Details All experimental processes were performed in ambient air unless otherwise mentioned. All chemicals

were used as received in this work. Abbreviations: MAI= methylammonium iodide (CH3NH3I), MBI= methylammonium bismuth iodide ((CH3NH3)3Bi2I9), P3HT= Poly(3-hexylthiophene-2,5-diyl), DMF= N, N-dimethylformamide, TiO2 blocking layer = TiO2 BL.

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Preparation of MAI

Typically, 24 ml of methylamine (Sigma-Aldrich, 33wt% in absolute ethanol), 10 ml of hydroiodic acid (Sigma-Aldrich, 55wt% in water), and 100 ml of absolute ethanol were mixed at 0℃ and magnetically stirred for 2 h. MAI was crystallized by a rotary evaporator at 50℃ . A yellowish powder was finally achieved via rotary evaporation, which was further purified to white colour after re-crystallization in the absolute ethanol and diethyl ether. Finally, the MAI powder was dried in a vacuum oven overnight at 60 ℃ and stored in a glove-box (O2 level < 0.1ppm, H2O level <0.1ppm).

Deposition of BiI3 films

Bismuth iodide (0.5897 g) was added into 1 ml anhydrous DMF to form 1 mol/L solution. After sonication for 1 h, the solution was filtered using PTFE syringe filters (0.45 µm). The substrates used for BiI3 film deposition were preheated at 100 ℃ before spin-coating. The spin-coating conditions were 2000RPM for 30 s and the ramping rate was 200RPM/s. After spin-coating, the film was annealed at 150 ℃ for 10 min. The BiI3 solution in DMF shows reddish colour. The as-spin-coated film showed red colour, which changed to black upon heat treatment.

Preparation of (CH3NH3)3Bi2I9 solution

A (CH3NH3)3Bi2I9 solution (1 mol/L in DMF) was prepared by mixing 0.2384 g of MAI, 0.5897 g of BiI3 and 1 ml of anhydrous DMF. The mixed solution was treated by ultra-sonication for 30min and finally filtered by PTFE syringe filters (0.45 µm), finally achieving a clear dark reddish solution.

Growth of MBI single crystals

To prepare MBI single crystals, 0.2384 g of MAI, 0.5897 g of BiI3 were added to 2 ml of anhydrous methanol and the solution was mixed by ultra-sonication for ~ 0.5 h. CH3NH3)3Bi2I9 is less soluble in methanol than that in the DMF. However, the boiling point of methanol is much lower than that of DMF. Clear (CH3NH3)3Bi2I9 solution (0.5 ml) in methanol was transferred to another clean vial and the solvent was allowed to evaporate slowly overnight in the fume-hood at room temperature, leading to the formation of MBI single crystals on the inner wall of the vial.

Fabrication of TiO2 mesoporous layers

A TiO2 blocking layer (BL) was deposited firstly on the cleaned FTO substrate by an aerosol spray pyrolysis method at 500 ℃ . The precursor solution for the spray pyrolysis is 0.2 M isopropanol solution of titanium diisopropoxide bis(acetylacetonate) and the thickness of the TiO2 blocking layer was controlled to

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be 80-100 nm. A TiO2 mesoporous layer was deposited onto the FTO/TiO2 BL by doctor blading. Typically, the TiO2 paste was prepared by ultra-sonicating a mixture of commercial TiO2 nanopac paste (TiO2 paste, TP-18AT, NANOPAC Co., Ltd.) and anhydrous ethanol (weight ratio 1:0.5). The substrate was covered by 3M scotch tape to define the thickness and protect the unwanted area. After doctor blading, the samples were annealed in a furnace at 500 ℃ for 60 min with a ramping rate of 5 ℃ /min.

Device fabrication

FTO (fluorine-doped tin oxide) coated glass slides (FTO22-7, Yingkou OPV Tech New Energy Co. Ltd) were firstly pre-patterned by removing part of the conductive layer with 2 M HCl aqueous solution and zinc powder. The etched substrates were cleaned in acetone, ethanol and isopropanol, subsequently. After oxygen plasma treatment (RIE, ATX-600, 30 W, 10 min), a thin TiO2 block layer was deposited by spray pyrolysis at 500 ℃ using 0.2 M isopropanol solution of titanium diisopropoxide bis(acetylacetonate) as the precursor solution.

One-step spin-coating method was used to deposit the MBI film on the TiO2 mesoporous layers, in which 50 µL of CH3NH3)3Bi2I9 solution (1 mol/L in DMF) was dispensed onto the substrate and spin-coated at 4000 RPM for 30s with a ramping rate of 4000 RPM/s, followed by heating at 100 ℃ for 10 min on a hotplate. A reddish film was achieved after heating treatment. Pristine P3HT (MW 54000-75000, Sigma-Aldrich, 15mg/ml in 1, 2-dichlorobenzene) was used as the hole-transporting layer, which was spin-coated at 2000 RPM for 30s after 60 soaking. Then, the coated substrates were further annealed at 100 ℃ for 15 min to remove the remaining solvent. The devices were left in the glove box overnight. Finally, a ~ 60 nm gold contact electrode was deposited by e-beam evaporation (Temescal FC-2000). The cell area was defined by a metal mask with an aperture area of 7.068 mm2.

Characterizations

The current-voltage (I-V) plots were recorded using a solar simulator (AM 1.5, 100mW/cm2, Oriel) equipped with a Keithley model 2420 digital source meter. Powder X-ray diffraction (XRD) measurement was performed on a Bruker Advanced X-ray Diffractometer with Cu Kα from 10° to 70° (2θ). SEM (scanning electron microscope) images were recorded by a Field-Emission Scanning Electron Microscope (FE-SEM, JEOL 7100). Optical image of single MBI crystals were carried out on an optical microscope (Leica). Reflectance spectra of the MBI films were performed by a UV-visible light spectrometer (Shimadzu UV2450). For the stability test of the MBI films, the as-deposited MBI films on the bare FTO substrates were placed in the ambient air with a relative humidity level ranging from 30% to 80%. Then, XRD spectra were analysed periodically to confirm the crystal structure. X-ray photoelectron spectroscopy (XPS, Kratos Ultra) was performed using mono Al X-ray gun. The theoretical XRD spectrum of MBI was calculated by Mercury 3.3 software (The Cambridge Crystallographic Data Centre (CCDC)) based on the results of single crystal X-ray diffraction measurement. The Conductivity of the MBI films were characterized by a Van der Pauw four-point probe equipped with a computer controlled source meter (KeithLink LRS4-T, Taiwan) at room temperature. Hall effect measurement was performed by a cryogen free high field measurement system

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(Cryogenics Ltd., UK) equipped with Keithley 2400, in which a source current of 1 mA was used and the magnetic field was scanned from -5 Tesla to 5 Tesla. The incident photon-to-current conversion efficiency (IPCE) was measured via a Newport 1918-c power meter equipped with a 300 W Oriel xenon light source and an Oriel Cornerstone 260 ¼ m monochromators in DC mode. A standard Si reference was used before the IPCE measurement.

Single-crystal X-ray diffraction measurement

Single-crystal X-ray diffraction data of MBI single crystals were recorded by an Oxford Diffraction Gemini Ultra S CCD diffractometer with Mo-Kα radiation (λ 0.71073 Å). The crystal was cooled to 190 K with an Oxford Cryosystems Desktop Cooler. The structure was solved by direct methods with SHELXS[1] and refined by full matrix least-squares analysis with SHELXL97 within the WinGX package[2].

First-principles electronic band structure calculations

All the calculations were performed based on the density functional theory (DFT) using VASP code.[3, 4] The electron exchange and correlation was treated with generalized gradient approximation (GGA)/Perdew, Burke, and Ernzerhof (PBE) functionals.[5] The cut-off energy for plane-wave basis set was 500 eV. For the first Brillouin zone, the k-point grid [6] was set to be 5x5x2. The optimized structure contains 6 C, 36 H, 6 N, 4 Bi and 18 I atoms, respectively. All atoms were full relaxed and force and energy were converged to 0.01 eV/Å and 10-5 eV, respectively.

The band alignments were derived from surface calculations using of vacuum potential as a common reference and aligning the average electrostatic potential in crystals to the vacuum potential. Valence- and conduction-band positions relative to the average electrostatic potential are separately determined from bulk calculations. For the surface calculations we use a repeated-slab geometry and choose (CH3NH3)3Bi2I9 (0 0 1) surface with each slab containing 4 layers. A vacuum layer with a thickness at least 15 Å was set to minimize artificial interactions between two neighboring slabs. All atoms near the surface were full relaxed and residual force and energy were less than 0.01 eV/Å and 10-5 eV, respectively.

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Part 2. Additional Tables, Figures and discussions

Table S1. Crystallographic data for MBI (CH3NH3)3Bi2I9) a

Empirical formula C3H18Bi2I9N3

Formula weight 1656.26

Temperature 190(2) K

Wavelength 0.71073 Å

Crystal system Hexagonal

Space group P 63/mmc

a, c, Å 8.4668(6), 21.614(2)

Volume, Å3 1341.9(2)

Z, Calculated density (g cm-3) 2, 4.099

Absorption coefficient, mm-1 23.441

F(000) 1400

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Crystal size, mm3 0.36 x 0.14 x 0.091

Theta range for data collection, degree 2.92 to 29.01

Index ranges

-10 ≤ h ≤ 11

-10 ≤ k ≤ 8

-29 ≤ l ≤ 26

Reflections collected 3788

Independent reflections 694 [R(int) = 0.1235]

Completeness to theta = 25.00°, % 99.4

Absorption correction Analytical

Max. and min. transmission 0.18 and 0.036

Refinement method Full-matrix least-squares on F2

Data / restraints / parameters 495 / 2 / 294

Goodness-of-fit on F2 1.214

Final R indices [I>2sigma(I)] b R1 = 0.0689, wR2 = 0.1672

R indices (all data) b R1 = 0.0752, wR2 = 0.1723

Largest diff. peak and hole, e.Å-3 1.9 and -1.8

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a Obtained by monochromated Mo Kα (λ = 0.71073 Å) radiation

b ,

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Table S2. Atomic coordinates ( x 104) and equivalent isotropic displacement parameters (Å2x 103). U(eq) is defined as one third of the trace of the orthogonalized Uij tensor.

________________________________________________________________________________

x y z U(eq)

________________________________________________________________________________

Bi(1) 6667 3333 1555(1) 38(1)

I(1) 10038(3) 5019(1) 2500 53(1)

I(2) 3311(2) 1656(1) 812(1) 72(1)

N(1) 13333 6667 1110(30) 160(30)

C(1) 12890(50) 5780(100) 490(30) 100(40)

N(2) 10000 10000 2500 290(130)

C(2) 9250(60) 8500(110) 2040(40) 30(30)

________________________________________________________________________________

Table S3. Selected bond distances (Å) for CH3NH3)3Bi2I9

Bi(1)-I(2) 2.9391(18)

Bi(1)-I(1) 3.2061(19)

N(1)-C(1) 1.489(10)

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N(2)-C(2) 1.489(10)

Symmetry transformations used to generate equivalent atoms:

#1 -y+1,x-y,z #2 -x+y+1,-x+1,z #3 x,y,-z+1/2

#4 -y+2,x-y,z #5 -x+y+2,-x+2,z #6 -x+y+1,-x+2,z

#7 -y+2,x-y+1,z #8 -y+2,x-y+1,-z+1/2 #9 -x+y+1,-x+2,-z+1/2

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Table S4. Selected bond angles (°) for CH3NH3)3Bi2I9

Bond Angle (°)

I(2)-Bi(1)-I(2)#2 92.93(7)

I(2)#1-Bi(1)-I(1) 91.46(4)

I(2)-Bi(1)-I(1) 173.61(7)

I(1)-Bi(1)-I(1)#2 83.79(5)

I(2)#1-Bi(1)-I(1)#1 173.61(7)

Bi(1)-I(1)-Bi(1)#3 79.10(6)

Symmetry transformations used to generate equivalent atoms:

#1 -y+1,x-y,z #2 -x+y+1,-x+1,z

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Fig S1. (a) MBI solution in anhydrous DMF solution, showing a deep reddish colour; (b) A photo of a MBI film deposited onto a TiO2 mesoporous layer coated FTO substrate.

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Fig S2. X-ray photoelectron spectroscopy (XPS) of the (CH3NH3)3Bi2I9 film on a FTO glass.

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Fig S3. UPS spectrum of (CH3NH3)3Bi2I9 film on a FTO substrate. The valence band maximum (VBM) is estimated to be around 5.90 eV (inset graph) and hence the valence band energy is ~ -5.90 eV with respect to the vacuum level.

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Fig S4. Three-dimensional AFM image of the (CH3NH3)3Bi2I9 crystal on a platinum coated glass substrate

and this crystal was used for KPFM measurement.

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Fig S5. SEM images of a (CH3NH3)3Bi2I9 infiltrated TiO2 mesoporous film. (a) Low-magnification surface

morphology; (b) High-magnification surface morphology; (c) Surface morphology of the film showing the

comparison of the white and dark areas; (d) Back-scattering SEM image of the same area shown in (c), in

which the white area can be confirmed to be occupied with higher concentration of Bi due to a larger

atomic number of Bi compared with that of Ti.

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Fig S6. (a) SEM images of a (CH3NH3)3Bi2I9 infiltrated TiO2 mesoporous film; (b) Energy dispersive spectrum of the MBI infiltrated TiO2 mesoporous film collected on the marked point.

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Table S5 Quantified element distribution of the EDS result for the sample presented in Fig S6.

Element Mass% Atom%

C 4.56 12.40

N 4.87 11.37

O 23.76 48.51

Ti 27.05 18.45

Sn 6.62 1.82

I 22.47 5.78

Bi 10.67 1.67

Total 100.00 100.00

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Fig S7. (a) A cross-sectional SEM image of the planar structured device; (b) The I-V curve of the assembled

device measured under irradiation of simulated AM 1.5 sunlight with a light intensity of 100 mW/cm2.

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Fig S8. A representative I-V curve of the device with a structure of FTO/TiO2 BL/TiO2mpl+P3HT/Au, under irradiation of simulated AM 1.5 sunlight with a light intensity of 100 mW/cm2.

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Fig S9. The incident photon-to-current conversion efficiency (IPCE) results for solar cells using P3HT only

(red curve) and MBI+P3HT (black curve).

Fig S10. The time-dependent efficiencies for the best-performing device.

Table S6 Current-voltage summary of the time-dependent performance of the best-performing device.

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Test time Voc / V Jsc / mAcm-2 Fill factor Efficiency%

Day 1 0.283 0.566 0.495 0.080 Day 2 0.328 0.594 0.528 0.103 Day 3 0.359 1.005 0.508 0.186 Day 9 0.323 1.118 0.536 0.194

Day 14 0.362 0.976 0.487 0.174 Day 21 0.370 0.976 0.493 0.181

[1] G. M. Sheldrick, Acta Cryst. A, 2008, A64, 112-122.

[2] L. J. Farrugia, Journal of Applied Crystallography 1999, 32, 837-838.

[3] Kresse G. and Furthmüller J., Phys. Rev. B 1996, 54, 11169-11186

[4] Kresse G. and Furthmüller J., Computational Materials Science, 1996,6, 15-50.

[5] J. P. Perdew, K. Burke, M. Ernzerhof, Physical review letters 1996, 77, 3865.

[6] H. J. Monkhorst, J. D. Pack, Physical Review B 1976, 13, 5188.

Address correspondence to Lianzhou Wang, [email protected]

22

Silver Nanowires with Semiconducting Ligands for Low Temperature Transparent Conductors

Brion Bob,1 Ariella Machness,1 Tze-Bin Song,1 Huanping Zhou,1 Choong-Heui Chung,2 and Yang Yang1,*

1 Department of Materials Science and Engineering and California NanoSystems Institute,

University of California Los Angeles, Los Angeles, CA 90025 (USA)

2 Department of Materials Science and Engineering, Hanbat National University, Daejeon

305-719, Korea

Abstract

Metal nanowire networks represent a promising candidate for the rapid fabrication of transparent electrodes with high transmission and low sheet resistance values at very low deposition temperatures. A commonly encountered obstacle in the formation of conductive nanowire electrodes is establishing high quality electronic contact between nanowires in order to facilitate long range current transport through the network. A new system of nanowire ligand removal and replacement with a semiconducting sol-gel tin oxide matrix has enabled the fabrication of high performance transparent electrodes at dramatically reduced temperatures with minimal need for post-deposition treatments of any kind.

Keywords: Silver Nanowires, Sol-Gel, Transparent Electrodes, Nanocomposites

23

1. Introduction. Silver nanowires (AgNWs) are long, thin, and possess conductivity values on the same order of magnitude as bulk silver

(Ag) [1]. Networks of overlapping nanowires allow light to easily pass through the many gaps and spaces between nanowires, while transporting current through the metallic conduction pathways offered by the wires themselves. The high aspect ratios achievable for solution-grown AgNWs has allowed for the fabrication of transparent conductors with very promising sheet resistance and transmission values, often approaching or even surpassing the performance of vacuum-processed materials such as indium tin oxide (ITO) [2-6].

Significant electrical resistance within the metallic nanowire network is encountered only when current is required to pass between nanowires, often forcing it to pass through layers of stabilizing ligands and insulating materials that are typically used to assist with the synthesis and suspension of the nanowires [7, 8]. The resistance introduced by the insulating junctions between nanowires can be reduced through various physical and chemical means, including burning off ligands and partially melting the wires via thermal annealing [9, 10], depositing additional materials on top of the nanowire network [11-14], applying mechanical forces to enhance network morphology [15-17], or using various other post-treatments to improve the contact between adjacent wires [18-21]. Any attempt to remove insulating materials the network must be weighed against the risk of damaging the wires or blocking transmitted light, and so many such treatments must be reined in from their full effectiveness to avoid endangering the performance of the completed electrode.

We report here a process for forming inks with dramatically enhanced electrical contact between AgNWs through the use of a semiconducting ligand system consisting of tin oxide (SnO2) nanoparticles. The polyvinylpyrrolidone (PVP) ligands introduced during AgNW synthesis in order to encourage one-dimensional growth are stripped from the wire surface using ammonium ions, and are replaced with substantially more conductive SnO2, which then fills the space between wires and enhances the contact geometry in the vicinity of wire/wire junctions. The resulting transparent electrodes are highly conductive immediately upon drying, and can be effectively processed in air at virtually any temperature below 300 °C. The capacity for producing high performance transparent electrodes at room temperature may be useful in the fabrication of devices that are damaged upon significant heating or upon the application of harsh chemical or mechanical post-treatments.

2. Results and Discussion

2.1. Ink Formulation and Characterization

Dispersed AgNWs synthesized using copper chloride seeds represent a particularly challenging material system for promoting wire/wire junction formation, and often require thermal annealing at temperatures near or above 200 °C to induce long range electrical conductivity within the deposited network [22, 23]. The difficulties that these wires present regarding junction formation is potentially due to their relatively large diameters compared to nanowires synthesized using other seeding materials, which has the capacity to enhance the thermal stability of individual wires according to the Gibbs-Thomson effect. We have chosen these wires as a demonstration of pre-deposition semiconducting ligand substitution in order to best illustrate the contrast between treated and untreated wires.

Completed nanocomposite inks are formed by mixing AgNWs with SnO2 nanoparticles in the presence of a compound capable of stripping the ligands from the AgNW surface. In this work, we have found that ammonia or ammonium salts act as effective stripping agents that are able to remove the PVP layer from the AgNW surface and allow for a new stabilizing matrix to take its place. Figure 1 shows a schematic of the process, starting from the precursors used in nanowire and nanoparticle synthesis and ending with the deposition of a completed film. The SnO2 nanoparticle solution naturally contains enough ammonium ions from its own synthesis to effectively peel the insulating ligands from the AgNWs and allow the nanoparticles to replace them as a stabilizing agent. If not enough SnO2 nanoparticles are used in the mixture, then the wires will rapidly agglomerate and settle to the bottom as large clusters. Large amounts of SnO2 in the mixture gradually begin to increase the sheet resistance of the nanowire network upon deposition, but greatly enhance the uniformity, durability, and wetting properties of the resulting films. We have found that AgNW:SnO2 weight ratios ranging between 2:1 and 1:1 produce well dispersed inks that are still highly conductive when deposited as films.

The nanowires were synthesized using a polyol method that has been adapted from the recipe described by Lee et al. [22, 23] Silver nitrate dissolved in ethylene glycol via ultrasonication was used as a precursor in the presence of copper chloride and PVP to provide seeds and produce anisotropic morphologies in the reaction products. Synthetic details can be found in the experimental section. Distinct from previous recipes, we have found that repeating the synthesis two times without cooling down the reaction mixture generally produces significantly longer nanowires than a single reaction step. The lengths of nanowires produced using this method fall over a wide range from 15 to 65 microns, with diameters between 125 and 250 nm. This range of diameters is common for wires grown using copper chloride seeds, although the double reaction produces a number of wires with roughly twice their usual diameter. The morphology of the as-deposited AgNWs as determined via SEM is shown in Figure 2(a), higher magnification images are also provided in Figures 2(c) and 2(d).

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The SnO2 nanoparticles were synthesized using a sol-gel method typical for multivalent metal oxide gelation reactions. A large excess of deionized water was added to SnCl4·5H2O dissolved in ethylene glycol along with tetramethylammonium chloride and ammonium acetate to act as surfactants. The reaction was then allowed to progress for at least one hour at near reflux conditions, after which the resulting nanoparticle dispersion can be collected, washed, and dispersed in a polar solvent of choice. The material properties of SnO2 nanoparticles formed using a similar synthesis method have been reported previously [24], although the present recipe uses excess water to ensure that the hydrolysis reaction proceeds nearly to completion.

After mixing with SnO2 nanoparticles, films deposited from AgNW/SnO2 composite inks show a largely continuous nanoparticle layer on the substrate surface with some nanowires partially buried and some sitting more or less on top of the film. Representative scanning electron microscopy (SEM) images of nanocomposite films are shown in Figure 2(b). Regardless of their position relative to the SnO2 film, all nanowires show a distinct shell on their outer surface that gives them a soft and slightly rough appearance, as is visible in the higher magnification images shown in Figure 2(e) and 2(f). The SnO2 nanoparticles do a particularly good job coating the regions near and around junctions between wires, and frequently appear in the SEM images as bulges wrapped around the wire/wire contact points.

The precise morphology of the SnO2 shell that effectively surrounded each AgNW was analyzed in more detail using transmission electron microscopy (TEM) imaging. Figures 3(a) to 3(c) show individual nanowires in the presence of different ligand systems: as-synthesized PVP in Figure 3(a), inactive SnO2 in Figure 3(b), and SnO2 activated with trace amounts of ammonium ions in Figure 3(c). The as-synthesized nanowires show sharp edges, and few surface features. In the presence of inactive SnO2, which is formed by repeatedly washing the SnO2 nanoparticles in ethanol until all traces of ammonium ions are removed, the nanowires coexist with somewhat randomly distributed nanoparticles that deposit all over the surface of the TEM grid. When AgNWs are mixed with activated SnO2, a thick and continuous SnO2 shell is formed along the nanowire surface. In when sufficiently dilute SnO2 solutions are used to form the nanocomposite ink, nearly all of the nanoparticles are consumed during shell formation and effectively no nanoparticles are left to randomly populate the rest of the image.

As the AgNWs acquire their metal oxide coatings in solution, the properties of the mixture change dramatically. Freshly synthesized AgNWs coated with residual PVP ligands slowly settle to the bottom of their vial or flask over a time period of several hours to one day, forming a dense layer at the bottom. The AgNWs with SnO2 shells do not settle to the bottom, but remain partially suspended even after many weeks at concentrations that are dependent on the amount of SnO2 present in the solution.

A comparison of the settling behavior of various AgNW and SnO2 mixtures after 24 hours is shown in Figures 3(d) and 3(e). The ratios 8:4, 8:16, and 8:8 indicate the concentrations of AgNWs and SnO2 (in mg/mL) present in each solution. The 8:8 uncoupled solution, in which the PVP is not removed from the AgNW surface with ammonia, produces a situation in which the nanowires and nanoparticles do not interact with one another, and instead the nanowires settle as in the isolated nanowire solution while the nanoparticles remain well-dispersed as in the solution of pure SnO2. The mixtures of nanowires and nanoparticles in which trace amounts of ammonia are present do not settle to the bottom, but instead concentrate themselves until repulsion between the semiconducting SnO2 clusters is able to prevent further settling.

Our current explanation for the settling behavior of the wire/particle mixtures is that the PVP coating on the surface of the as-synthesized wires is sufficient to prevent interaction with the nanoparticle solution. The addition of ammonia into the solution quickly strips off the PVP surface coating and allowing the nanoparticles to coordinate directly with the nanowire surface. This explanation is in agreement with the effects of ammonia has on a solution of pure AgNWs, which rapidly begin to agglomerate into clusters and sink to the bottom as soon as any significant quantity of ammonia is added to the ink.

We attribute the stripping ability of ammonia in these mixtures to the strong dative interactions that

occur via the lone pair on the nitrogen atom interacting with the partially filled d-orbitals of the Ag atoms

on the nanowire surface. These interactions are evidently strong enough to displace the existing

coordination of the five-membered rings and carbonyl groups contained in the original PVP ligands and

allow the ammonia to attach directly to the nanowire surface. Since ammonia is one of the original

surfactants used to stabilize the surface of the SnO2 nanoparticles, we consider it reasonable that ammonia

coordination on the nanowire surface would provide an appropriate environment for the nanoparticles to

adhere to the AgNWs.

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Scanning Energy Dispersive X-ray (EDX) Spectroscopy was also conducted on nanoparticle-coated AgNWs in order to image the presence of Sn and Ag in the nanowire and shell layer. The line scan results are shown in Figure 3(f), having been normalized to better compare the widths of the two signals. The visible broadening of the Sn lineshape compared to that of Ag is indicative of a Sn layer along the outside of the wire. The increasing strength of the Sn signal toward the center of the AgNW is likely due to the enhanced interaction between the TEM’s electron beam and the dense AgNW, which then improves the signal originating from the SnO2 shell as well. It is also possible that there is some intermixing between the Ag and Sn x-ray signals, but we consider this to be less likely as the distance between their characteristic peaks should be larger than the detection system’s energy resolution.

2.2. Network Deposition and Device Applications

For the deposition of transparent conducting films, a weight ratio of 2:1 of AgNWs to SnO2 nanoparticles was chosen in order to obtain a balance between the dispersibility of the nanowires, the uniformity of coated films, and the sheet resistance of the resulting conductive networks. Nanocomposite films were deposited on glass by blade coating from an ethanolic solution using a scotch tape spacer, with deposited networks then being allowed to dry naturally in air over several minutes.

The as-dried nanocomposite films are highly conductive, and require only minimal thermal treatment to dry and harden the film. Without the use of activated SnO2 ligands, deposited nanowire networks are highly insulating, and become conductive only after annealing at above 200 °C. The sheet resistance values of representative films are shown in Figure 4(a). The capability to form transparent conductive networks in a single deposition step that remain useful over a wide range of processing temperatures provides a high degree of versatility for designing thin film device fabrication procedures.

Figure 5(a) shows the sheet resistance and transmission of a number of nanocomposite films deposited from inks containing different nanowire concentrations. The deposited films show excellent conductivity at transmission values up to 85%, and then rapidly increase in sheet resistance as the network begins to reach its connectivity limit. The optimum performance of these networks at low to moderate transmission values is a consequence of the relatively large nanowire diameters, which scatter a noticeable amount of light even when the conditions required for current percolation are just barely met. Nonetheless, the sheet resistance and transmission of the completed nanocomposite networks place them within an acceptable range for applications in a variety of optoelectronic devices. Figure 5(b) shows the wavelength dependent transmission spectra of several nanowire networks, which transmit light well out into the infrared region. The presence of high transmission values out to wavelengths well above 1300 nm, where ITO or other conductive oxide layers would typically begin to show parasitic absorption, is due to the use of semiconducting SnO2 ligands, which is complimentary to the broad spectrum transmission of the silver nanowire network itself.

Avoiding the use of highly doped nanoparticles has the potential to provide optical advantages, but can create difficulties when attempting to make electrical contact to neighboring device layers. In order to investigate their functionality in thin film devices, we have incorporated AgNW/SnO2 nanocomposite films as electrodes in amorphous silicon (a-Si) solar cells. Two contact structures were used during fabrication: one with the nanocomposite film directly in contact with the p-i-n absorber structure and one with a 10 nm Al:ZnO (AZO) layer present to assist in forming Ohmic contact with the device. The I-V characteristics of the resulting devices are shown in Figure 6(a).

The thin AZO contact layers typically show sheet resistance values greater than 2.5 kΩ/⧠, and so cannot be responsible for long range lateral current transport within the electrode structure. However, their presence is clearly beneficial in improving contact between the nanocomposite electrode and the absorber material, as the SnO2 matrix material is evidently not conductive enough to form a high quality contact with the p-type side of the a-Si stack. We hope that future modifications to the AgNW/SnO2 composite, or perhaps the use of islands of high conductivity material such as a discontinuous layer of doped nanoparticles will allow for the deposition of completed electrode stacks that provide both rapid fabrication and good performance.

Figure 6(b) contains the top view image of a completed device. The enhanced viscosity of the nanowire/sol-gel composite inks allows for films to be blade coated onto substrates with a variety of surface properties without reductions in network uniformity. In contrast with traditional back electrodes deposited in vacuum environments, the nanocomposite can be blade coated into place in a single pass under atmospheric conditions and dried within moments. We anticipate that the use of sol-gel mixtures to enhance wetting and dispersibility may prove useful in the formulation of other varieties of semiconducting and metallic inks for deposition onto a variety of substrate structures.

3. Conclusions

In summary, we have successfully exchanged the insulating ligands that normally surround as-synthesized AgNWs with shells of substantially more conductive SnO2 nanoparticles. The exchange of one set of ligands for the other is mediated by

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the presence of ammonia during the mixing process, which appears to be necessary for the effective removal of the PVP ligands that initially cover the nanowire surface. The resulting nanowire/nanoparticle mixtures allow for the deposition of nanocomposite films that require no annealing or other post-treatments to function as high quality transparent conductors with transmission and sheet resistance values of 85% and 10 Ω/⧠, respectively. Networks formed in this manner can be deposited quickly and easily in open air, and have been demonstrated as an effective n-type electrode in a-Si solar cells when a thin interfacial layer is deposited first to ensure good electronic contact with the rest of the device. The ligand management strategy described here could potentially be useful in any number of material systems that presently suffer from highly insulating materials that reside on the surface of otherwise high performance nano and microstructures.

4. Experimental Details

Tin oxide nanoparticle synthesis. Tin chloride pentahydrate was dissolved in ethylene glycol by

stirring for several hours at a concentration of 10 grams per 80 mL to serve as a stock solution. In a typical

synthesis reaction, 10 mL of the SnCl4·5H2O stock solution is added to a 100 mL flask and stirred at room

temperature. Still at room temperature, 250 mg ammonium acetate and 500 mg ammonium acetate were

added in powder form to regulate the solution pH and to serve as coordinating agents for the growing

oxide nanoparticles. 30 ml of water was then added, and the flask was heated to 90 °C for 1 to 2 hours in

an oil bath, during which the solution took on a cloudy white color. The gelled nanoparticles were then

washed twice in ethanol in order to keep trace amounts of ammonia present in the solution. Additional

washing cycles would deactivate the SnO2, and then require the addition of ammonia to coordinate with

as-synthesized AgNWs.

Silver nanowire synthesis. Copper(ii) chloride dihydrate was first dissolved in ethylene glycol at

1 mg/ml to serve as a stock solution for nanowire seed formation. 20 ml of ethylene glycol was then added

into a 100 ml flask, along with 200 µL of copper chloride solution. the mixture was then heated to 150 °C

while stirring at 325 rpm, and .35g of PVP (MW 55,000) was added. In a small separate flask, .25 grams of

silver nitrate was dissolved in 10 ml ethylene glycol by sonicating for approximately 2 minutes, similar to

the method described here.22 The silver nitrate solution was then injected into the larger flask over

approximately 15 minutes, and the reaction was allowed to progress for 2 hours. After the reaction had

reached completion, the various steps were repeated without cooling down. 200 µL of copper chloride

solution and .35g PVP were added in a similar manner to the first reaction cycle, and another .25g silver

nitrate were dissolved via ultrasonics and injected over 15 minutes. The second reaction cycle was allowed

to progress for another 2 hours, before the flask was cooled and the reaction products were collected and

washed three times in ethanol.

Nanocomposite ink formation. After the synthesis of the two types of nanostructures is complete,

27

the double washed SnO2 nanoparticles and triple-washed nanowires can be combined at a variety of weight

ratios to form the completed nanocomposite ink. The dispersibility of the mixture is improved when more

SnO2 is used, although the sheet resistance of the final networks will begin to increase if they contain

excessive SnO2. AgNW agglomeration during mixing is most easily avoided if the SnO2 and AgNW

solutions are first diluted to the range of 10 to 20 mg/ml in ethanol, with the SnO2 solution being added

first to an empty vial and the AgNW solution added afterwards. The dilute mixture was then be allowed to

settle overnight, and the excess solvent removed to concentrate the wires to a concentration that is

appropriate for blade coating.

Film and electrode deposition. The completed nanocomposite ink was deposited onto any desired

substrates using a razor blade and scotch tape spacer. The majority of the substrates used in this study were

Corning soda lime glass, but the combined inks also deposited well on silicon, SiO2, and any other

substrates tested. Electrode deposition onto a-Si substrates was accomplished by masking off the desired

cell area with tape, and then depositing over the entire region. The p-i-n a-Si stacks and 10 nm AZO

contact layers were deposited using PECVD and sputtering, respectively.

ACKNOWLEDGMENTS The authors would like to acknowledge the use of the Electron Imaging Center for Nanomachines

(EICN) located in the California NanoSystems Institute at UCLA.

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Figure 1. Process flow diagram showing the synthesis of AgNWs and SnO2 nanoparticles followed

by stirring in the presence of ammonium salts to create the final nanocomposite ink. Transparent

conducting films were produced by blade coating the completed inks onto the desired substrate.

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Figure 2. (a,c,d) SEM images of as-synthesized AgNWs at various magnifications. (b,e,f) SEM

images of nanocomposite films, showing the tendency of the SnO2 nanoparticles to coat the entire

outer surface of the AgNWs, increasing their apparent diameter and giving them a soft appearance.

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Figure 3. Schematic diagrams and TEM images of (a) a single untreated AgNW, (b) an AgNW in the

presence of uncoupled SnO2 (all ammonium ions removed), and (c) an AgNW with a coordinating

SnO2 shell. Scale bars in images (a), (b), and (c) are 300 nm, 400 nm, and 600 nm, respectively. (d,e)

Optical images of AgNW and SnO2 nanoparticle dispersions mixed in varying amounts (d) before and

(e) after settling for 24 hours. The numbers associated with each solution represent the AgNW:SnO2

concentrations in mg/ml. The uncoupled solution contains AgNWs and non-coordinating SnO2

nanoparticles, and shows settling behavior similar to the pure AgNW and pure SnO2 solutions. (f)

Normalized Ag and Sn EDX signal mapped across the diameter of a single nanowire, with the inset

showing the scanning path across an isolated wire.

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Figure 4. Sheet resistance versus temperature for films deposited using (red) AgNWs that have been

washed three times in ethanol and (blue) mixtures of AgNW and SnO2 with weight ratio of 2:1. The

annealing time at each temperature value was approximately 10 minutes. The large sheet resistance

values of the bare AgNWs when annealed below 200 °C is typical for nanowires fabricated using

copper chloride seeds, which clearly illustrate the impact of SnO2 coordination at low treatment

temperatures.

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Figure 5. (a) Sheet resistance and transmission data for samples deposited from solutions of varying

nanostructure concentration. Each of these samples were fabricated starting from the same

nanocomposite ink, which was then diluted to a range of concentrations while maintaining the same

AgNW to SnO2 weight ratio. (b) Transmission spectra of several transparent conducting networks

chosen from the plot in plot (a).

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Figure 6. (a) I-V characteristics of devices made with AgNW/SnO2 rear electrodes with (blue) and

without (red) a 10 nm AZO contact layer. The dramatic double diode effect is likely a result of a

significant barrier to charge injection at the electrode/a-Si interface. (b) Top view SEM image of the

AgNW/SnO2 composite films on top of the textured a-Si absorber. (c) Schematic cross section of the

a-Si device architecture used in solar cell fabrication. The thickness of the thin AZO contact layer is

exaggerated for clarity.

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