Strongly emissive perovskite nanocrystal inks for high-voltage solar cells

Lead halide perovskite semiconductors have recently gained wide interest following their successful embodiment in solid-state photovoltaic devices with impressive power-conversion efficiencies, while offering a relatively simple and low-cost processability. Although the primary optoelectronic properties of these materials have already met the requirement for high-efficiency optoelectronic technologies, industrial scale-up requires more robust processing methods, as well as solvents that are less toxic than the ones that have been commonly used so successfully on the lab-scale. Here we report a fast, room-temperature synthesis of inks based on CsPbBr3 perovskite nanocrystals using short, low-boiling-point ligands and environmentally friendly solvents. Requiring no lengthy post-synthesis treatments, the inks are directly used to fabricate films of high optoelectronic quality, exhibiting photoluminescence quantum yields higher than 30% and an amplified spontaneous emission threshold as low as 1.5 μJ cm−2. Finally, we demonstrate the fabrication of perovskite nanocrystal-based solar cells, with open-circuit voltages as high as 1.5 V. Despite their impressive performance, more efforts are required to develop industrially scalable perovskite solar cells from less toxic solvents. Towards that aim, this study presents the use of colloidal nanoparticle inks for room-temperature fabrication of CsPbBr3 solar cells.

C olloidal semiconductor nanocrystals (NCs) enable solution processing and represent a powerful platform for tuning optical and electrical properties useful for optoelectronic devices [1][2][3] . Over the past decades, the main challenges have been the synthesis of NC solutions with tailored properties (bandgap, absorption, monodispersity) and the conversion of these solutions to high-quality NC films, with preservation of their properties in solution 1,2,4 . One major problem originates directly from the synthesis methods. The high-temperature (100-350 • C) syntheses often required to fabricate crystalline monodisperse and shapecontrolled NCs tend to employ ligands and solvents with long alkyl chains [5][6][7] . These bulky ligands and the residual chemicals from the synthesis prevent the production of dense films and form an insulating layer around the NCs after deposition onto a substrate 8 . Strategies have been developed to overcome these issues, such as the addition of conductive polymers 9,10 or post-synthesis ligand exchange [11][12][13] . Most surface passivation schemes reported so far have been based on solid-state ligand exchange, with long ligands used to stabilize the NCs in solution being replaced by much shorter ligands that ensure closer NC packing and thus higher electrical conductivity [11][12][13][14] . Progress in such schemes has allowed for a steady improvement of charge-carrier diffusion lengths in solid films (above 100 nm), enabling, for example, the fabrication of thick solar-cell devices while preserving efficient charge collection, leading to record efficiencies exceeding 10% 14,15 .
Halide perovskite semiconductors can merge the highly efficient operational principles of conventional inorganic semiconductors with the low-temperature solution processability of emerging organic and hybrid materials, offering a promising route towards cheaply generating electricity as well as light [16][17][18][19] . Perovskites not only show exceptional primary optoelectronic properties such as a direct bandgap 20 , small exciton binding energy 21 , low carrier recombination rates 22 , ambipolar transport 23 , and tunability of the bandgap covering a wavelength range from the near infrared 24 to the ultraviolet, but they are also very attractive for their ease of processability for mass production (for example, printing from solution) and for the wide availability of their chemical components. For the most efficient optoelectronic devices, the semiconductor films are currently processed from precursor solutions dissolved in organic solvents [25][26][27][28] . After deposition, the constituent ions self-assemble during crystallization directly upon the selected substrate when treated at temperatures below 120 • C. Such a process can form high-quality thin films, but it also has drawbacks. Principally, it couples both the thin-film morphology and the related optoelectronic properties/crystal quality in the same optimization step, making it more sensitive to processing conditions. The thin films are polycrystalline, but can exhibit varying morphologies determined by different factors. These include precursor ratio, solvent, processing additives, substrate roughness and surface energy, atmospheric/environmental conditions, annealing temperature, and treatment time. As a result, the thin films contain a significant density of structural and chemical defects, which introduce loss channels.
A key parameter for a rapid evaluation of the optoelectronic quality of the perovskite thin film is its photoluminescence quantum yield (PLQY). In solar cells, under steady-state solar illumination, electrons are photoexcited from the valence band into the conduction band, a process which splits the electron and hole quasi-Fermi levels. The extent of the splitting is determined by the charge density at which the recombination rate is equal to the carrier generation rate. Any additional non-radiative recombination process with a carrier lifetime shorter than the 10   radiative decay will thus reduce the steady-state charge density and the open-circuit voltage. Perovskite thin films hardly achieve 20% PLQY at excitation intensities that are relevant for photovoltaic applications, indicating a significant density of defect states 19,28,29 . On the other hand, colloidally synthesized CsPbBr 3 and MAPbBr 3 (MA = methylammonium) perovskite NCs currently can exhibit PLQYs close to 90% and can be synthesized in a wide variety of sizes and shapes [30][31][32][33][34][35][36][37][38] . However, the enhancement in PLQY comes at the expense of carrier extraction, because the bulky organic ligands used in the synthesis of these CsPbBr 3 NCs to passivate and stabilize their surface inhibit inter-particle connectivity. This prevents their use as proper 'inks' for the fabrication of dense conductive thin films 30,31,37 . Traditional ligand exchange procedures, as explained above, are of little help in this case, due to intrinsic lability of these NCs and the additional difficulty of purifying solutions of halide perovskite NCs from excess surfactants and unreacted precursors 39 . Thus using colloidal perovskite NCs as a starting material for the active layer is challenging. Nonetheless, during the preparation of this manuscript, Swamkar et al. reported a nanoscale phase separation stabilization of CsPbI 3 quantum dots (QDs) to low temperatures and showed their use as the active component in solar cells 40 .
Here we report a fast and room-temperature synthesis of CsPbBr 3 perovskite NC inks using short, low-boiling-point ligands and solvents. These inks have optical qualities close to those of NCs made with high-temperature hot-injection syntheses. The use of short ligands and solvents circumvents post-synthesis treatments and enables the production of thin films with high optoelectronic quality-that is, PLQYs greater than 30% and an amplified spontaneous emission (ASE) threshold as low as 1.5 µJ cm −2 . Importantly, the robustness of such properties are demonstrated by the fabrication of CsPbBr 3 NCs-based solar cells, with open-circuit voltages as high as 1.5 V.

Synthesis and characterization of CsPbBr 3 NC inks
The synthesis of the CsPbBr 3 NC inks, as depicted in Fig. 1a, is carried out by a simple and fast one-step injection (see also Supplementary Video 1). We use propionic acid (PrAc) to dissolve Cs 2 CO 3 and form a Cs + propionate complex which is diluted in a polar/apolar solvent mixture of isopropanol (IPrOH), hexane (HEX), and butylamine (BuAm) at room temperature. The dissolution reaction is exothermic and therefore does not require heating. Thus, unlike previous room-temperature syntheses 36,37 , no degassing was required for the precursors and the whole process was carried out under nitrogen. A separate solution, prepared in air by dissolving PbBr 2 in a similar mixture of chemicals (also at room temperature), was injected into the first solution. The NCs immediately nucleated and had already reached their maximum size 10 s after the injection (see below for additional details on the growth). At this point, the NCs were separated by centrifugation and redispersed in toluene, after which they could be used directly for device preparation. The size of the NCs could be controlled by varying the IPrOH to HEX ratio, as shown in Supplementary Fig. 1. Here, an increase of this ratio, and thus an increase of the polarity of the solution, led to larger crystalline domains. However, this route led to unstable colloidal solutions and was thus not investigated further. As the synthesis was performed under a nitrogen atmosphere (in the glovebox) at room temperature, it could be easily scaled-up to a gram-sized synthesis (900 ml, 1.9 g PbBr 2 , Fig. 1f), with no noticeable changes in the NC properties ( Supplementary Fig. 2).  31,40 . With propionic acid (PrAc), isopropanol (IPrOH), hexane (HEX), butylamine (BuAm), octadecene (ODE), oleic acid (OA) and oleylamine (OLAM).
Important differences between this reaction scheme and previous ones (for example, the one reported by Protesescu et al. 31 ) are that here we use only room-temperature reactions, and additionally the solvents and ligands of the current synthesis have much lower boiling points ( Table 1). As will be discussed later in more detail, upon spin-coating of a solution of these NCs on a substrate, all these solvents and ligands evaporated quickly at room temperature, ensuring a fast drying of the film and facilitating the fabrication of thick films by multiple layer depositions. Also, compared to most previous syntheses of halide perovskite NCs that involved two steps 30,36 , the present scheme is based on a single step. Finally, and equally important, the solvents used (IPrOH and HEX) are more environmentally friendly than the classical dimethyl formamide (DMF) used in the preparation of bulk perovskite films, and therefore are more amenable to scale-up in industrial processes. We could additionally replace hexane and toluene with the even less hazardous heptane (boiling point = 98 • C) and still prepare NCs with the same features and properties ( Supplementary Fig. 3).
As shown in Fig. 1b,c, the NCs prepared with this method tended to cluster in large aggregates. These aggregates were not formed due to the centrifugation of the NCs in the cleaning step, since similar clusters were also observed when the solution was investigated with transmission electron microscopy (TEM) directly after the injection of the PbBr 2 precursor (and thus before washing; see Supplementary Fig. 4). On the basis of highresolution TEM (HRTEM) (Fig. 1d), they contain crystalline NCs with domains of roughly 15-20 nm in size and atomic planes matching those of orthorhombic CsPbBr 3 ( Supplementary Fig. 5). The crystallinity of the NC domains and the orthorhombic crystal structure was further confirmed with X-ray diffraction (XRD) analysis 41 , as reported in Fig. 1e. The surface Cs:Pb:Br ratios were 0.90:1.00:3.00 according to X-ray photoelectron spectroscopy (XPS) analysis, indicating a slightly Cs-deficient CsPbBr 3 composition ( Supplementary Fig. 6) 42 . XPS was further used to determine the carbon content in drop-cast samples. In a drop-cast film of 8.5 nm cubic CsPbBr 3 NCs, synthesized with octadecene (ODE), oleic acid (OA) and oleylamine (OLAM), and washed twice at 12,000 r.p.m. (as described in previous works 31,43 ), 88 at.% of the sample surface consisted of carbon. This percentage is comparable to that reported for colloidal NCs in general (not necessarily based on halide perovskites) and synthesized under similar conditions 44 . For the NCs reported here, and washed only once for 2 min at 1,000 r.p.m., the surface carbon content significantly dropped to 22 at.%, which corresponds to a ∼25 fold decrease in the ratio of carbon to the overall CsPbBr 3 inorganic component. Homogeneous films (see scanning electron microscopy (SEM) image, Supplementary Fig. 7) could be easily prepared by spincoating and drop-casting of the NCs inks and required only about 10 min of drying under ambient conditions, whereas the 8.5 nm cubic CsPbBr 3 NCs required multiple high-speed centrifugation steps (to get rid of excess ligands) and long drying times under vacuum, as already reported 31,45 .

Optical characterization
The growth of the NCs during the synthesis was monitored over time by following their photoluminescence (PL), as shown in Supplementary Fig. 8. From the data collected it is evident that the reaction takes place within the first 10 s from the mixing of the two solutions (see Supplementary Video 1), and that even after 2 min no further growth was observed. The PL and optical absorption spectra recorded on the purified CsPbBr 3 NCs, both for samples dispersed in toluene and after deposition from such dispersions to form a solid film, are shown in Fig. 2a. The NCs in solution had a PL peak centred at ∼515 nm with a full-width-half-maximum (FWHM) of 24 nm, similar to that of cubic 8.5 nm NCs 31,43 . In the NC film, the optical absorption edge was redshifted by about 7 nm, namely from 515 nm to 522 nm (on the basis of the band edge), which may be attributed to a change in the local strain of the NCs in the different phase 46,47 . The PL from the film consistently followed the absorption edge, and retained the same Stokes shift (around 9 nm) and FWHM as in the solution. Similar PL dynamics were found in solutions and in films (Fig. 2b), in agreement with the fact that the PL quantum yield did not undergo a dramatic drop, from about 58 ± 6% in solution to 35 ± 4% in the film (excitation density around 250 µW cm −2 ). Note that the PLQYs from these films were slightly higher than those typically observed from films of cubic CsPbBr 3 NCs prepared with OA and OLAM (≈30%), which remained non-conductive 43 .
Measurements under femtosecond (fs)-excitation were carried out to investigate the presence of ASE from the NC films, which represents a good fingerprint of their optical quality 48 . ASE was readily observed for films fabricated by spin-coating of the NCs solution onto soda-lime glass substrates (200-300 nm thick, Fig. 2c). The ASE peak had a FWHM between 4.3 and 4.6 nm, and was redshifted by 3 nm from the PL maximum (λ PL = 522 nm). We observed a reduced ASE redshift compared to that reported for other CsPbBr 3 NCs 49-51 , which we tentatively ascribed to the increased Stokes shift observed in our spin-coated films (9 nm, compared to other synthetic approaches 37 , see Fig. 2a), since the optical gain arises in spectral positions where the optical reabsorption is reduced (Urbach tail) 49 . Figure 2d reports the emission intensity versus pumping fluence, from which an ASE threshold of around 2.44 µJ cm −2 could be extracted. It is known that the film morphology can induce optical feedback 52,53 (due to the relatively high refractive index of the NCs, n ≈ 2, with respect to that of air) and thus modify the ASE threshold and induce random lasing. For this reason, the ASE threshold was measured from ten different films in different positions prepared from two batches (total of 62 measurements). The threshold varied from 1.5 µJ cm −2 up to 38 µJ cm −2 , with a median value of 9 µJ cm −2 . To our knowledge, 1.5 µJ cm −2 is the lowest value reported for CsPbBr 3 NCs 49-51 and overall for perovskite thin films (a geometry that is more suitable for the fabrication of devices such as electrically pumped lasers). This evidence supports the high optical and electronic quality of the ink and its related thin films, most likely due to a low density of electronic defects. The ASE threshold found here is actually among the lowest reported to date, even if compared to that of other inorganic NCs commonly used in laser devices (CdSe nanoplatelets 54 , CdSe/CdS giant shells 52,53 , and CdSe/CdS dot-in-rods 55,56 ).

High-voltage solar cells
Wide-bandgap halides perovskites such as CsPbBr 3 can achieve high open-circuit voltages and are therefore of particular interest for multi-junction solar cells, visible light-emitting devices, and solar water splitting applications [57][58][59] . Moreover, the replacement of the organic component (methylammonium and/or formamidinium) in the lead halide framework with an inorganic cation such as caesium is sought as the most promising avenue for improving the thermal compositional stability 60,61 , which is important for solarcell operation. However, due to the relative insolubility of the CsBr precursor, it is difficult to achieve a high level of control over the thin-film morphology and optoelectronic properties using conventional approaches, and retain low processing temperatures. Thus, thin-film devices based on CsPbBr 3 have remained relatively unexplored to date. In testing the NCs inks reported here in solar-cell devices, we have selected a state-of-the-art architecturenamely, a stack comprised of fluorine doped tin oxide (FTO) coated glass, a compact layer of titanium dioxide (c-TiO 2 ) as the electron-extracting layer, a layer of 2,2 ,7,7 -tetrakis-(N ,N -p dimethoxyphenylamino)-9 -spirobifluorene (Spiro-MeOTAD) as the hole-transporting material and evaporated Au as the top contact. In Fig. 3a we report a sketch of the device stack and in Fig. 3b we report the energy levels of each layer. For the electrodes and the charge-extracting layers we have used values reported in the literature 62 , while for the NC thin films the energies of the conduction and valence band edges have been extracted by adding the optical bandgap to the position of the valence band maximum as determined by ultraviolet photoelectron spectroscopy (UPS) measurements (see Methods and Supplementary Fig. 9).
With a single spin-coating step of the suspended CsPbBr 3 ink on the c-TiO 2 substrate, a complete coverage of the substrate was obtained. The solar cell, in this case, exhibited a short-circuit value (J sc ) of 1.26 mA cm −2 , an open-circuit voltage (V oc ) of 0.87 V, and a fill factor (FF) of 0.65, leading to a power-conversion efficiency (PCE) of 0.72% (see Supplementary Table 1). The reported FF and V oc suggest a good electronic quality for such an extremely thin layer of perovskite; however, the photocurrent was low, as the film was very thin (its optical density was ∼0.3, see Supplementary Fig. 10). A thicker film could be prepared by performing sequential deposition cycles of the NC ink. This process, which was uniquely facilitated by the low boiling point of the ligands and solvents, enabled a fast drying of the film after each deposition step, with no need to anneal the film. In particular, the fact that we can deposit the perovskite from an apolar solvent means that the active layer is resistant to subsequent NC ink deposition cycles and other spincoating steps 14,15 .
By increasing the film thickness through sequential depositions, the optical density of the sample grew monotonically (see Supplementary Fig. 10) while the NCs maintained their structural properties, as proved by comparing the XRD pattern of these films to one obtained from a film deposited by simple drop-casting ( Supplementary Fig. 11). As reported in Fig. 3c, increasing the number of NC deposition cycles led to a direct increase of both J sc and V oc of the devices, while the FF remained almost constant (Supplementary Fig. 12 and Supplementary Table 2). An active layer with a thickness of 550 ± 50 nm ( Supplementary Fig. 13), prepared by nine sequential depositions, exhibited a PCE of 5.4%, which is comparable to the best-performing and fully optimized devices reported so far [58][59][60] . The solar cell had J sc and V oc values of 5.6 mA cm −2 and 1.5 V, as well as a FF of 0.62. The value for V oc is among the highest reported for perovskite halides, underlining the high quality of the active layer; for reference, the calculated maximum V oc and J sc are 2.05 V and 7.78 mA cm −2 under AM1.5 The use of TiO 2 as the electron-extracting layer in a flat junction architecture generally presents J -V characteristics that depend on the polarization history of the device, with a reduction of the photocurrent over time under polarization 62,63 . Interestingly, these devices show a rapid response of the photocurrent under polarization, resulting in electrically stable steady-state power output (see Supplementary Figs 12 and 13). Finally, it is worth mentioning that, although there are several factors that determine the V oc and their relative contribution is still a completely open issue in perovskite solar cells, preliminary literature reports seem to suggest that further improvements will be possible by carefully designing the charge-extracting layers 65 , thus highlighting the strong potential of our approach.

Conclusions
In summary, we have developed a scalable and environmentally friendly synthesis of inks consisting of a colloidal suspension of CsPbBr 3 NCs passivated with shorter ligands than previously reported syntheses. This has enabled the fabrication of compact, fast-drying films of NCs with low carbon content. The effective passivating ability of these ligands was responsible for the high PLQYs of the particles both in solution (58%) and, remarkably, after their deposition into films (35%), without severe degradation of charge-carrier conduction in the film. Simple spin-coated films of these inks exhibited a record low ASE threshold of 1.5 µJ cm −2 . The CsPbBr 3 perovskite inks were then used to fabricate a halide perovskite NC-based photovoltaic prototype device. With this formulation we demonstrated a fully air-processed, electrically stable, solar cell exhibiting PCEs exceeding 5% and an open-circuit voltage higher than 1.5 V. These values are among the highest achieved for wide-bandgap perovskite halides, and for Cs-based solar cells. Device fabrication. FTO-coated glass sheets were etched with zinc powder and HCl (2 M) to obtain the required electrode pattern. The substrates were sonicated in sequence with detergent (Alconox), distilled water, 2-propanol, acetone and 2-propanol again for 10 min, respectively. The substrates were then blown dry with N 2 and finally treated with oxygen plasma for 10 min to remove the last traces of organic residues. The TiO 2 precursor solution was prepared by mixing 6 µl of 2M HCl in 1 ml of 2-propanol to a titanium isopropoxide solution in 2-propanol (140 µl titanium isopropoxide in 1 ml of 2-propanol). This solution was then spin-coated at 2,000 r.p.m. for 60 s and sintered at 500 • C following Ball and colleagues 66 . The CsPbBr 3 suspension solution was spin-coated at 1,000 r.p.m. for 45 s. To increase the active layer thickness with subsequent deposition cycles the samples were left drying for 5 min at room temperature between each spinning. The samples were then transferred into a N 2 -filled glove box for the deposition of spiro-MeOTAD as the hole-transporting material (HTM). The HTM was spin-coated at 1,000 r.p.m. for 60 s. The solution was prepared by dissolving 75 mg of spiro-MeOTAD, 32 µl 4-tert-butylpyridine, 18.8 µl of a stock solution of 520 mg ml −1 LiTFSI in acetonitrile in 1 ml anhydrous chlorobenzene. After 12 h of exposure to dry air, a 75-nm-thick gold film was thermally evaporated through a shadow mask to prepare devices with a total area of about 0.0935 cm 2 .

Chemicals
Transmission electron microscopy. Conventional transmission electron microscopy (TEM) images were acquired on a JEOL JEM-1011 microscope equipped with a thermionic gun at 100 kV accelerating voltage. High-resolution TEM (HRTEM) imaging was performed on a JEOL JEM-2200FS microscope equipped with a Schottky gun operated at an 80-200 kV accelerating voltage, with a CEOS spherical aberration corrector in the objective lens enabling a spatial resolution of 0.9 Å, and an in-column -filter. The samples were prepared by drop-casting diluted NC suspensions onto 200-mesh carbon-coated copper grids for conventional TEM imaging, and 400-mesh ultrathin carbon-coated copper grids for HRTEM imaging, respectively. To avoid or reduce the diffusion of Pb induced by the electron beam, the images were taken by summing multiple fast acquisitions (≤0.2 s) after drift correction, by using a K2 direct detection camera (Gatan), allowing low doses (<5e − /pixel s).
X-ray and ultraviolet photoelectron spectroscopy. These analyses were performed on spin-cast films of CsPbBr 3 NC inks, respectively, to evaluate the chemical composition and to estimate the position of the valence band maximum (VBM) of the materials under investigation. The measurements were carried out with a Kratos Axis Ultra DLD spectrometer. For XPS measurements, high-resolution spectra of Cs 3d, Pb 4f , Br 3d and C 1s peaks were acquired at a pass energy of 10 eV using a monochromatic Al Kα source (15 kV, 20 mA). The UPS measurements were performed using a He I (21.22 eV) discharge lamp, on an area of 55 µm in diameter, at a pass energy of 5 eV and with a dwell time of 100 ms. The work function (that is, the position of the Fermi level with respect to the vacuum level) was measured from the threshold energy for the emission of secondary electrons during He I excitation. A −9.0 V bias was applied to the