Molecular doping enabled scalable blading of efficient hole-transport-layer-free perovskite solar cells

The efficiencies of perovskite solar cells (PSCs) are now reaching such consistently high levels that scalable manufacturing at low cost is becoming critical. However, this remains challenging due to the expensive hole-transporting materials usually employed, and difficulties associated with the scalable deposition of other functional layers. By simplifying the device architecture, hole-transport-layer-free PSCs with improved photovoltaic performance are fabricated via a scalable doctor-blading process. Molecular doping of halide perovskite films improved the conductivity of the films and their electronic contact with the conductive substrate, resulting in a reduced series resistance. It facilitates the extraction of photoexcited holes from perovskite directly to the conductive substrate. The bladed hole-transport-layer-free PSCs showed a stabilized power conversion efficiency above 20.0%. This work represents a significant step towards the scalable, cost-effective manufacturing of PSCs with both high performance and simple fabrication processes.

S ince their first report in 2009, halide perovskite solar cells (PSCs) have attracted enormous interest as emerging lightto-electricity conversion optoelectronic devices with both high efficiency and ease of solution processability [1][2][3] . To date, the power conversion efficiencies (PCEs) of solution-processed PSCs have been certified at well above 20% [4][5][6] . However, all topperforming PSCs were fabricated by a spin-coating method which is unsuitable for high throughput and scalable module production. Several scalable film deposition techniques have been developed for PSC fabrication, such as doctor-blading 7-10 , spray deposition 11 , slot-die coating 12 , solution printing 13 , and electrodeposition 14 . In particular, doctor-blade coating with a controlled substrate temperature has gained most success in depositing highquality perovskite films with large grains and controllable optoelectronic properties 15 . Moreover, blade coating has great potential to be scaled up for roll-to-roll fabrication. Encouragingly, the PCEs of PSCs with bladed perovskite layers have steadily increased to above 19%, approaching the levels of their spin-coated counterparts 13,16 .
A next step towards the scalable fabrication of PSCs is to coat the charge transport layers by scalable processes. However, the existing multiple-layer device structures, either p-i-n or n-i-p structures with a hole transport layer (HTL), perovskite, and electron transport layer (ETL), impose a great challenge on the scalable fabrication of PSCs in a cost-effective manner [17][18][19][20] . First, high-performance organic hole-transporting materials, such as 2,2′,7,7′-Tetrakis[N,N-di(4-methoxyphenyl)amino]-9,9′-spirobifluorene (spiro-OMeTAD), are much more expensive than perovskite materials. Second, for the devices with p-i-n structure, it is difficult to blade-coat perovskite films onto a hydrophobic HTL, e.g., poly(bis(4-phenyl)(2,4,6-trimethylphenyl)amine) (PTAA), while a hydrophilic HTL, e.g., NiO x , is difficult to form on plastic substrates at low temperature, nor can easily be coated at optimal thickness and electronic quality to achieve the highest device efficiency. Hence, we believe that the elimination of the HTL from the PSC device architecture can be an effective way to solve this challenge. By reducing one step of the standard coating process, at least one third of the manufacture cost and time can be saved, in addition to savings of material cost for the HTL.
The reported ambipolar charge transport characteristics of halide perovskite materials endow them with great potential to construct PSCs with simplified device architecture, such as HTLfree PSCs with either conventional or inverted structure [21][22][23] . Yet, a critical challenge for HTL-free PSCs is the ineffective hole collection by indium tin oxide (ITO) because of mismatched work functions between ITO and the perovskite, which generally causes significant efficiency loss [23][24][25] . In addition, the construction of HTL-free devices by the blade-coating method faces more challenges due to rougher and more non-uniform perovskite films, than do those prepared by spin-coating with the antisolvent method.
Here we report a molecular-doping strategy for a perovskite layer to doctor-blade HTL-free PSCs with a simplified device structure of ITO/perovskite/ETL/Cu. Doping methylammonium lead iodide (MAPbI 3 ) films with 2,3,5,6-tetrafluoro-7,7,8,8-tetracyanoquinodimethane (F4TCNQ), a strong electron-withdrawing molecule, leads to a modified ITO/MAPbI 3 interface with favorable band bending, which facilitates the extraction of photoexcited holes from perovskite to ITO electrode. Combining additive engineering and solvent annealing, high-quality MAPbI 3 films are obtained with full surface coverage, micrometer-sized grains, and improved crystallinity. By employing a bladed and doped MAPbI 3 film as the photoactive layer in a HTL-free device structure, a stabilized PCE over 20.0% with almost no J-V hysteresis is obtained.

Results
Perovskite films fabrication and molecular doping. The MAPbI 3 perovskite films were directly deposited onto blank ITO glass substrates by a previously reported doctor-blading method 9 , which is schematically illustrated in Fig. 1a. Before blade coating, the ITO substrates were cleaned by a UV-ozone treatment to improve their wettability to perovskite solution 23 . MAPbI 3 precursor dissolved in N,N-diethylformamide (DMF) containing ptype dopant (F4TCNQ) was dripped onto substrates which were heated at 150°C, followed by a quick blading process to spread the solution onto the substrates. After blading, the perovskite precursor solution dried in a few seconds, and the substrates were quickly removed from the hot plate and annealed at 100°C for varied durations. The bladed perovskite films showed thicknesses of 500 ± 20 nm from cross-sectional imaging (Fig. 1b) by scanning electron microscopy (SEM). To implement molecular doping of the perovskite layer, F4TCNQ with chemical structure shown in Fig. 1a was dissolved in DMF separately, and then added to the as-prepared MAPbI 3 precursor solution. Perovskite films made without or with F4TCNQ were denoted as MAPbI 3 or MAPbI 3 :F4TCNQ, respectively.
The doping of perovskite is expected to change the work function of the films. Kelvin probe force microscopy (KPFM) was used to determine the surface potential change of perovskite films upon F4TCNQ mixing. In addition to direct mixing of F4TCNQ and MAPbI 3 in precursor solution, MAPbI 3 :F4TCNQ was prepared by solid diffusion via heating F4TCNQ powder on MAPbI 3 and then removing the excess F4TCNQ. The topography and contact potential difference (CPD) images of pristine MAPbI 3 and MAPbI 3 :F4TCNQ films by either standard solution mixing or solid diffusion are shown in Fig. 1c-h. For better comparison, the three CPD images are flattened and presented in the same color scale, while the histograms of the absolute CPD values of the three types of films are plotted in Fig. 1i. The pure MAPbI 3 film has coarse grains and uniform CPD distribution (Fig. 1c, f). In contrast, solution mixed MAPbI 3 :F4TCNQ film, in which uniform doping is expected, has fine grains and greater heterogeneity of the CPD distribution ( Fig. 1d, g). For the sample with solid-diffused F4TCNQ, some small F4TCNQ particles were placed on top of a MAPbI 3 film, and the CPD imaging was conducted around the edge (the green color line as indicated in Fig. 1h) to include both neat and doped perovskite regions. Such method leads to a clear edge between diffused and undiffused regions that can be visualized in the CPD image ( Fig. 1h), thus the CPD heterogeneity is even higher. The change of absolute CPD caused by F4TCNQ doping is clearly demonstrated in Fig. 1i. A decrease of surface potential or in other words an increase of work function was observed for the MAPbI 3 :F4TCNQ films, irrespective of direct mixing or solid diffusion. The increase of heterogeneity for the MAPbI 3 :F4TCNQ films is also reflected in the broadening of the curves. This confirmed p-type doping of MAPbI 3 by F4TCNQ. Specifically, after the addition of electronwithdrawing F4TCNQ molecule to the MAPbI 3 , owing to the deep lowest unoccupied molecular orbital (LUMO) level of F4TCNQ, electrons in the valance band of MAPbI 3 can be thermally activated into the LUMO level of F4TCNQ, which leaves holes in MAPbI 3 and thus causes its p-doping (Fig. 1j). As a result, the hole concentration in the MAPbI 3 film increased. The carrier concentration change upon F4TCNQ doping has been evaluated by analyzing the dark capacitance vs. voltage (C-V) characteristics of the perovskite devices ( Supplementary Fig. 1). Charge density profiles were extracted from the C-V curves using Mott-Schottky analysis 26,27 . The carrier concentration in the perovskite layer was calculated to be 0.6 × 10 16 cm −3 in the neat MAPbI 3 film which increased to 2.4 × 10 16 cm −3 after the addition of 0.03 wt% F4TCNQ. This molecular p-type doping strategy enabled the fabrication of an HTL-free PSC with a simplified device configuration but still effective p-n heterojunction (i.e., p-doped perovskite and n-type ETL), and thus enhanced device performance.
Electronic properties of doped perovskite films. The electronic properties of the MAPbI 3 :F4TCNQ films were investigated by comparing their current density-voltage (J-V) characteristics with pristine MAPbI 3 films deposited on glass substrates in the dark condition. The lateral conductivity was measured across gold electrodes spaced 100 µm apart (Fig. 2a). As shown in Fig. 2b, the MAPbI 3 :F4TCNQ film exhibited a nearly tenfold increase in conductivity over the neat MAPbI 3 film, which confirmed that F4TCNQ doped MAPbI 3 . The charge transfer from MAPbI 3 to F4TCNQ is expected to cause quenching of photoluminescence (PL) from the MAPbI 3 layer, which was measured by using time-resolved photoluminescence (TRPL) decay for MAPbI 3 films deposited on glass substrates (Fig. 2c). The MAPbI 3 :F4TCNQ prepared on glass substrate showed a PL lifetime of 65 ns, which is fourfold shorter than the PL lifetime of pristine MAPbI 3 film (268 ns) on glass.
The relative large molecular size of F4TCNQ likely prevents its integration into the lattice of MAPbI 3 . Therefore, we speculate that F4TCNQ molecules most probably remain at the grain boundary (GB) regions of perovskite films and/or perovskite/ electrode interfaces. To determine the lateral distribution of F4TCNQ molecules, the local dark conductivity at grains and GBs was measured using conducting atomic force microscopy (c-AFM). Figure 2d illustrates the c-AFM setup, where the microscope tip was positioned at locations randomly chosen from topographic AFM images (Fig. 2e, g). The tip bias was scanned in the range +2.0 to −2.0 V at a rate of 0.14 V s −1 . For the neat MAPbI 3 film, the dark currents were similar at the grain and GB regions ( Fig. 2f and Supplementary Fig. 2a, b), which is in accordance with previous reporting 28 . In sharp contrast, the dark currents measured at the GB regions of the MAPbI 3 :F4TCNQ film were notably higher ( Fig. 2h and Supplementary Figs. 2c, d), elucidating the remarkably improved conductivity at the GB region after molecular doping of MAPbI 3 by F4TCNQ. At the regions close to GBs, the F4TCNQ molecules and MAPbI 3 undergo charge transfer, leading to an increase in carrier density in MAPbI 3 and thus the enhanced dark conductivity 29 .
Morphologies and crystallinities of perovskite films. Highquality perovskite films with large grains, high crystallinity and complete surface coverage are also crucial to achieving high-  performance HTL-free devices. The blending of F4TCNQ in the perovskite in a small percentage (ranged from 0.01 to 0.05 wt%) did not notably change the morphology and grain size of the perovskite films, as shown by the SEM images in Supplementary  Fig. 3, nor cause a drastic change of crystallinity, as shown by the XRD patterns of bladed MAPbI 3 films prepared with different amounts of F4TCNQ in Supplementary Fig. 4. Accordingly, methylammonium hypophosphite (MHP) and methylammonium chloride (MACl) were introduced as additives to the precursor solution, and co-solvents (dimethyl sulfoxide (DMSO)/ chlorobenzene (CBZ)) post-annealing process was also combined to achieve a high-quality and high-crystalline MAPbI 3 film (Supplementary Figs. 5, 6 and 7 for SEM images and XRD patterns of bladed MAPbI 3 films prepared with different amounts of MHP as well as using co-solvents annealing technique in the absence or presence of MACl) 9,10,16,30,31 . Figure 3a, b show representative SEM images of the bladed perovskite films after optimization, which featured a uniform morphology that extended hundreds of micrometers. The perovskite domains were densely packed together, and without observable voids at domain boundaries (Fig. 3a). It should be noted that these domains are not single crystalline grains. Each micrometer-scale convection domain was comprised of micro-sized grains ranging from 800 to 1700 nm (average above 1 μm, Fig. 3b).  Table 1, which indicated the optimal F4TCNQ concentration was 0.03 wt%. Figure 3d shows the current density vs. voltage (J-V) characteristics of PSCs based on the pristine MAPbI 3 films and MAPbI 3 :F4TCNQ films (with 0.03 wt% F4TCNQ), with photovoltaic parameters summarized in Table 1. The reference PSC without F4TCNQ (i.e., MAPbI 3 ) exhibited a poor PCE of just 11.0%, which is in accordance with that of most reported HTL-free PSCs 23,24 . Surprisingly, with the addition of F4TCNQ, notable concurrent enhancements of the J sc , V oc , and FF have been achieved. The champion cell based on MAPbI 3 :F4TCNQ film achieved the highest PCE of 20.2%, with a J sc of 22.7 mA cm −2 , a V oc of 1.10 V, and a FF of 0.81. To the best of our knowledge, this is the record efficiency for the bladed PSCs in the absence of HTLs. For comparison, the PSC based on the PTAA HTL and undoped MAPbI 3 layer has been fabricated, which showcases a PCE of 19.5%, along with a J sc of 22.1 mA cm −2 , a V oc of 1.12 V, and a FF of 0.79. This result is in good agreement with our previous reports 8,9,16 . It is worth noting that the demonstrated HTL-free device showed similar or even better photovoltaic performance, highlighting the effectiveness of such a doping strategy for constructing high-performance PSCs with not only a simpler structure, but also less deposition steps and lower material cost. The J-V curves of MAPbI 3 :F4TCNQ-based device were measured in both reverse and forward scan modes at a scan rate of 0.1 V s −1 , and almost no hysteresis was observed (Supplementary Fig. 8). Figure 3e shows the stabilized photocurrent density of this device as measured at 0.93 V, giving a stabilized PCE exceeding 20.0% under continuous light soaking. This high stabilized PCE output is among the best record for HTL-free PSCs, and is quite remarkable given that the device is prepared using a scalable fabrication method. Figure 3f shows the external quantum efficiency (EQE) of the champion device. The integrated J sc from the EQE measurement agreed well with the J sc values measured under AM 1.5 G one sun illumination (Fig. 3d). The fabrication procedure was repeated to confirm the reliability and reproducibility of the results. The PCE histogram was collected from 30 independent cells (Fig. 3g), revealing that more than 73% of the cells had PCEs above 18.0%, and more than 93% had PCEs higher than 17.0% under one sun illumination. A study of the stability for the devices with or without HTL has been conducted.  The devices were stored under ambient conditions (where humidity and temperature were about 20% and 25°C, respectively) without any encapsulation. As compared to the devices with HTL, the HTL-free device exhibits a slightly better stability, which maintained over 92% of its initial PCE after storage for 500 h (Supplementary Fig. 9). It is noted from Fig. 3d and Table 1 that the almost twofold efficiency enhancement in the F4TCNQ-doped devices primarily originates from the remarkable FF increase. Further analysis of the J-V curves in Fig. 3d revealed that F4TCNQ doping of the perovskite layer significantly reduced device series resistance (R s ) from 25.6 Ω cm 2 for undoped device to 2.6 Ω cm 2 for doped device. The influence of R s on device FF has been well established in diode based solar cells 32 , according to the classical equation based on diode equation: where I is the device output current, I L is the maximum light generated current, I 0 is saturated dark current, R s is series resistance, n is ideal factor, k is Boltzmann constant, T is temperature. Assuming a diode ideal factor of 1, the FFs for the undoped and doped devices were calculated to be 53% and 83% from Eq. (1), which are very close to measured values of 56% and 81%, respectively. This indicates the increased FF for the devices with doped MAPbI 3 is mainly caused by the reduction of the series resistance. To find out the origin of reduced series resistance by F4TCNQ doping, energy level diagrams are constructed and shown in Fig. 4 for the perovskite/ITO interface with or without F4TCNQ doping. Here the energy levels were extracted from the literature for the films made with same composition and similar method 33,34 . For undoped MAPbI 3 , the Fermi level of pristine MAPbI 3 (nearly intrinsic) is slightly lower than that of ITO (−4.7 eV) 34 . The Fermi level lines up during contact formation and causes a downward band-bending in MAPbI 3 , which forms a barrier for hole extraction from MAPbI 3 to ITO (Fig. 4a). This would also increase the possibility of charge recombination and increase series resistance at the ITO/ perovskite interface. In contrast, the incorporation of F4TCNQ in the MAPbI 3 films turns the MAPbI 3 to more p-type, particularly at the interface of ITO and MAPbI 3 , which should avert band-bending direction (Fig. 4b). This was confirmed by surface potential measurement that the F4TCNQ doping can increase work function of perovskite by~0.1 to 0.2 eV (Fig. 1i). As is shown in Fig. 4b, the F4TCNQ doping results in an upward band bending at the ITO/MAPbI 3 :F4TCNQ interface. In this case, the hole transfer from F4TCNQ-doped MAPbI 3 to ITO is thus facilitated, resulting in a reduced series resistance to increase FF 35 . The diode Eq. (1) predicts that R s variation has no influence on device V oc , while in our case, the F4TCNQ doping increased device V oc from 1.00 to 1.10 V. This discrepancy indicates there is another factor, in addition to R s , which leads to the V oc enhancement. To find out the reason, a transient photovoltage spectroscopy (TPV) measurement was also carried out to understand the charge carrier recombination dynamics. As shown in Supplementary Fig. 10, the MAPbI 3 :F4TCNQ-based PSC had a longer carrier lifetime (0.69 μs) than the reference device based on neat MAPbI 3 (0.30 μs) under one sun illumination. This result was consistent with the larger V oc for the device with doping. The observed carrier recombination lifetime enhancement could be a result of band-bending configuration change due to doping, though there may be other reasons that may also change the carrier recombination lifetime. For instance, for the device with undoped MAPbI 3 , the retarded hole extraction would increase the possibility of charge recombination, because they cannot be effectively extracted out.

Discussion
We demonstrated how molecular doping influences the series resistance and interfacial charge transfer within HTL-free PSC devices with or without F4TCNQ incorporated in the perovskite layer. In the absence of HTL, both the electrons and holes can transport to the ITO/perovskite interface and recombine there, leading to serious interfacial charge recombination and thus inferior charge collection 36 . In addition, the mismatched work functions between perovskite and ITO would induce an undesirable energy barrier that leads to inferior hole transfer. These situations are all very detrimental to device performance. Molecular doping of perovskite by F4TCNQ to some extent induces favorable interfacial band bending, which allows efficient transport of holes at ITO/perovskite interface (Fig. 4b) 35 .
Choosing an appropriate p-type dopant with proper molecular size and selected surface hydrophility with ITO substrate is promising to further improve the ITO/perovskite interfacial affinity. Owing to the larger molecular size of F4TCNQ than the interstitial sites in perovskite structure, the F4TCNQ concentrated at the grain boundary regions. The dopant assisted with charge dissociation and thus resulting in more balanced transfer of electrons and holes to their respective electrodes, which decreased the possibility of the charge carrier recombination. We found the charge trap density within MAPbI 3 :F4TCNQ-based PSC was even slightly smaller than its pristine MAPbI 3 counterpart ( Supplementary Fig. 11). It has been demonstrated that the trap sites within the organic semiconductors can be reduced due to partial filling by doping, typically at very low doping levels 37 .
In summary, an additive-assisted strategy for p-type molecular doping of solution-bladed perovskite films was demonstrated. F4TCNQ-doped MAPbI 3 films show increased electrical conductivity, especially at grain boundary regions, and increased charge carrier concentrations. The incorporation of F4TCNQ in perovskite layers could modify the ITO/perovskite interface with reduced series resistance, which could be attributed to the favorable interfacial band bending for facilitated hole transfer and extraction from perovskite to ITO. The simple but effective approach enabled the scalable fabrication of HTL-free PSC devices with a simplified device geometry. Molecular doping of perovskite film by F4TCNQ led to the considerable enhancement of photovoltaic performance from 11.0 to 20 31 . Finally, C 60 (20 nm, Nano-C), bathocuproine (BCP, 8 nm, Sigma-Aldrich), and Cu electrode (80 nm) were sequentially deposited on the perovskite layer by thermal evaporation to form a complete PSC device.
Device characterization. SEM images were obtained with a Quanta 200 FEG environmental SEM. X-ray diffraction (XRD) patterns were acquired by a Bruker D8 Discover Diffractometer with Cu Kα radiation (1.5406 Å). PL spectrum was measured with a Horiba iHR320 Imaging Spectrometer at room temperature. A 532 nm green laser (Laserglow Technologies) with an intensity of 100 mW cm −2 was used as the excitation source. TRPL was obtained using the DeltaPro with a pulsed laser source at 406 nm (Horiba NanoLED 402-LH; pulse width below 200 ps, 20 pJ per pulse,~1 mm 2 spot size), and the signal was recorded using time corrected TCSPC. c-AFM measurements were performed on an Asylum Research MFP-3D AFM using Pt-coated Si conductive probes (PPP-EFM, Nanosensors). The conductivity of perovskite films was measured by a two-electrode method on glass substrates. The J-V characteristics of the cells (voltage scanning rate 0.1 V s −1 ) and the steady photocurrent under maximum power output bias (0.93 V) were recorded with a Keithley 2400 source-meter under simulated AM 1.5 G irradiation by a Xenon lamp (Oriel 67005), which was calibrated by a silicon diode equipped with a Schott visible color KG5 glass filter. EQE was measured with a Newport QE measurement kit. The J-V and EQE measurements were conducted under ambient air conditions without encapsulation. A non-reflective shadow mask was used to define a 0.08 cm 2 active area of the PSCs. The series resistance (R s ) is derived from the slope of the J-V curve at the open-circuit voltage point. TPV decay was measured under 1 sun illumination and recorded by a 1 GHz Agilent digital oscilloscope.
Data availability. The data that support the findings of this study are available from the corresponding author on reasonable request.