Reports
remains largely unexplored. Only hole–conductor–free mesoscopic PSCs using carbon as a back contact have shown so far promising stability under long term light soaking and long term heat exposure, but their certified PCE remains relatively low at 12.8% (12, 13). Wei Chen,1, 2* Yongzhen Wu,1* Youfeng Yue,1 Jian Liu,1 Wenjun
Regardless of their architec-Zhang,2 Xudong Yang,3 Han Chen,3 Enbing Bi,3 Islam Ashraful,1
tures, all high efficiency PSCs so
Michael Gr?tzel,4? Liyuan Han1,3?
far employed small areas, their
12
Photovoltaic Materials Unit, National Institute for Materials Science, Tsukuba, Ibaraki 305-0047, Japan. Michael Gr?tzel device size being often <0.1 cm2 Centre for Mesoscopic Solar Cells, Wuhan National Laboratory for Optoelectronics, Huazhong University of Science and (Table S1) (14). As such a small Technology, Wuhan, China. 3State Key Laboratory of Metal Matrix Composites, Shanghai Jiao Tong University, 800 Dong
device size is prone to induce Chuan Road, Minhang District, Shanghai 200240, China. 4Laboratory of Photonics and Interfaces, Department of
measurement errors, an obliga-Chemistry and Chemical Engineering, Swiss Federal Institute of Technology, Station 6, CH-1015 Lausanne, Switzerland.
tory minimum cell area of >1 *These authors contributed equally to this work.
cm2 is required for certified PCEs
?Corresponding author. E-mail: michael.graetzel@epfl.ch (M.G.); han.liyuan@nims.go.jp (L.H.)
to be recorded in the standard
The recent stunning rise in power conversion efficiencies (PCEs) of “Solar Cell Efficiency Tables” perovskite solar cells (PSCs) has triggered worldwide intense research. edited by public test centers, However, high PCE values have often been reached with poor stability at an such as National Renewable En-2
illuminated area of typically less than 0.1 cm. We used heavily doped ergy Laboratory (NREL) in the inorganic charge extraction layers in planar PSCs to achieve very rapid US and the National Institute of carrier extraction even with 10-20 nm thick layers avoiding pinholes and Advanced Industrial Science and eliminating local structural defects over large areas. This robust inorganic Technology (AIST) in Japan (15).
2
nature allowed for the fabrication of PSCs with an aperture area >1 cm It has been recommended that showing a power conversion efficiency (PCE) >15% certified by an the record efficiencies should be accredited photovoltaic calibration laboratory. Hysteresis in the current-recorded with cell size of ideally
voltage characteristics was eliminated; the PSCs were stable: >90% of the 1 to 2 cm2 or larger to allow initial PCE remained after 1000 hours light soaking. comparison of competing tech-nologies (16–19). Although a few
works reported attempts of fab- Organic–inorganic metal halide perovskite solar cells
(PSCs) have attracted large attention due to the meteoric ricating centimeter-scale PSCs, for example by using vacu-rise in their solar to electric power conversion efficiency um evaporation system (20) or modified two-step approach (PCE) over the last few years (1). In particular, (21) to produce large area MAPbI3 films, the PCEs obtained methylammoniun (CH3NH3PbI3, denoted as MAPbI3) and for these devices reached only 10.9%~12.6%. Apart from the formamidinium lead iodide (CH(NH2)2PbI3) emerged as a small device areas, the widely recognized hysteresis and sta-highly attractive solar light harvesting materials because of bility issues of PSCs have raised doubts on the reliability of their intense broad–band absorption, high charge carrier previously claimed high efficiencies (22, 23).
The poor reproducibility and lack of uniformity of PSCs mobility, low–cost precursor materials and simple solution
processing (2, 3). Their ambipolar semiconducting render it challenging to obtain high efficiencies with large characteristics further enable variable device architectures, devices. It is difficult to control the formation of cracks and ranging from mesoscopic to planar structures with n–i–p or pinholes in the selective carrier extraction layers over large p–i–n layouts (4). Mesoporous TiO2-based PSCs have areas. As small size PSCs typically show a wide spread in reached the highest performance (5–7); their certified PCE their PCEs, previous work has focused on improving the attaining presently 20.1% (8). However, there is growing uniformity of perovskite layer by varying its deposition interest in inverted (p–i–n) planar device architectures methods (3, 6, 10). However, fewer studies have aimed at typically employing a MAPbI3–PCBM ([6,6]–phenyl–C61–identifying selective extraction layers for photogenerated butyric acid methyl ester) bilayer junction, because of their charge carriers placed over the current collector to prevent simple fabrication and relatively small hysteresis (9–11). A their recombination at its surface (24–26); event though key question that remains open to date is the true power such selective contacts have turned out to be equally im-conversion efficiency and stability of planar PSCs as none of portant to developing efficient solar light harvesters (27). the devices has been certified so far and their stability The dilemma with optimizing of such charge carrier extrac-tion layers in solar cells is that the film should be thin to
Efficient and stable large-area perovskite solar cells with inorganic charge extraction layers
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mimimize resistive losses while at the same time it should cover the whole collector area in a contiguous and uniform manner. Meeting these requirement becomes increasingly difficult as the device area increases.
Here we present a strategy that addresses simultaneous-ly the scale up and stability issues facing current PSC em-bodiments. We develop heavily p–doped (p+) NixMg1–xO and n–doped (n+) TiOx contacts to extract selectively photogen-erated charge carriers from an inverted planar MAPbI3–PCBM film architecture. We implement the p+ and n+ dop-ing by substituting Ni(Mg)2+ ions and Ti4+ ions on the NixMg1–xO lattice and TiOx matrix by Li+ and Nb5+ ions, re-spectively. The resulting dramatic increase in the electrical conductivity enables 10-20 nm thick oxide layers to be used for selective extraction of one type of charge carriers while improving their electronic blocking effect for the other type by reducing the density of pinholes and cracks over large areas. Accordingly, the series resistance (Rs) of the oxides decreased and the shunt resistance (Rsh) greatly increased with respect to the undoped ones, allowing excellent fill fac-tor (FF) with values exceeding 0.8 and hysteresis–free be-havior to be achieved. With this strategy, we successfully fabricated large size (>1 cm2) PSCs with an efficiency of up to 16.2%. A PCE of 15% was certified by a public test center (Calibration, Standards and Measurement Team at the Re-search Center for Photovoltaics, National Institute of Ad-vanced Industrial Science and Technology, AIST, Japan). This is listed as the first official efficiency entry for PSCs in the most recent edition of the “Solar Cell Efficiency Tables” (28). Moreover, the devices based on these stable p+ and n+ doped inorganic metal oxides charge extraction layers show high stability rendering them attractive for future practical deployment of PSCs.
We fabricated PSCs with an inverted planar device archi-tecture (Fig. 1A); a cross sectional scanning electron micros-copy (SEM) image of the device is shown in Fig. 1B. We first deposited the NiO based hole extraction layer onto fluorine-doped tin oxide (FTO) glass using spray pyrolysis. The pre-cursor solution was composed of nickel (II) acetylacetonate alone or together with doping cations (Mg2+ from magnesi-um acetate tetrahydrate, Li+ from lithium acetate) in super–dehydrated acetonitrile/ethanol mixture. The MAPbI3 per-ovskite layer with thickness of ~300 nm was deposited by a reported method (6), which was followed by the deposition of a thin PCBM layer (80 nm) via spin–coating its chloro-benzol solution (20 mg ml–1) at 1000 rpm for 30 s. An n–type TiOx based electron extraction layer with and without Nb5+ doping was further deposited on the PCBM by spin–coating a diluted methanol solution of titanium isopropoxide (or mixed with niobium ethoxide), followed by controlled hy-drolysis and condensation (14). Finally, the device was com-pleted by thermal evaporation of a 100 nm thick Ag cathode. The band alignments of relevant functional layers are shown in Fig. 1C, based on the energy levels determined by ultraviolet (UV) photoelectron spectroscopy (UPS) and ul-traviolet–visible (UV–Vis) absorption spectroscopy meas-urements (fig. S1) (14). The uniformity of the perovskite and PCBM layers was examined by cross sectional SEM observa-tion (fig. S2) (14). The full XPS spectra of the NiO and TiOx based charge carrier extraction layers are shown in fig. S3 (14), revealing the designated compositions for the target materials. The close–up observation on the morphology of the charge carrier extraction layers are depicted in fig. S4 (14), while their pin–hole densities were examined by elec-trical measurement as discussed below.
The stoichiometric form of NiO is a wide band gap semi-conductor with a very low intrinsic conductivity of 10?13 S cm–1 (29). Self–doping by introducing Ni3+ acceptors into the NiO crystal lattice renders the crystals more conductive, depending on the film deposition techniques and conditions (11, 30–32). The room temperature specific conductivity of our NiO films from Hall effect measurements was 1.66 × 10?4 S cm–1. This value is much lower than the typically used p–type contact layer of PEDOT: PSS that shows a conductiv-ity of 1 to 1000 S cm–1 (33). The low conductivity of NiO will lead to a high Rs resulting in a low FF of the solar cells (34). Substitutional doping by Li+ is an effective way to increase the p-conductivity of NiO (35). Values of heavily (p+)–doped NiO films can reach 1 to 10 S cm–1 at room temperature un-der optimal conditions (36). For our Li+ doped NixMg1–xO films the conductivity is 2.32 × 10?3 S cm–1, ~12 times greater than that of the undoped reference.
A Mg2+ content of 15 mol% was alloyed in the Li+ doped nickel oxide film, to compensate the undesirable positive shift of its valence band (EVB) caused by Li+ incorporation into the lattice (fig. S1) (14, 35, 37, 38). As the Li+ content was adjusted to 5 mol%, the doped oxide has the formula of Li0.05Mg0.15 Ni0.8O if one assumes that the molar ratio of the three different cations in the spray pyrolyses solution is maintained in the mixed oxide. This co–doping strategy is feasible because the mismatch of the ionic radii of Li+ (0.76?), Mg2+ (0.71?) and Ni2+ (0.69?) is quite small, confer-ring good lattice stability to the LixMgyNi1–x–yO ternary ox-ides. We compared conductivity of NiO and Li0.05Mg0.15Ni0.8O films by using contact–current mode of a scanning probe microscope (SPM) and show the results in Fig. 2, A and B. At a bias potential of 1.0 V, the electric current increased by a factor of ~10 (from ~0.3 nA to ~3 nA) upon replacing un-doped NiO by a Li0.05Mg0.15Ni0.8O film. The XPS spectra in fig. S6 (14) reveal that the doping increased the relative con-tent of Ni3+ acceptors in the Li0.05Mg0.15Ni0.8O sample. These findings are consistent with reports on Li+ doped NiO films in (32) and explains the increase in carrier concentration from 2.66 × 1017 cm–3 of the undoped NiO film to 6.46 × 1018 cm–3 for Li0.05Mg0.15Ni0.8O film that we derived from Hall ef-fect measurements.
The electron specific n–type TiOx contact used commonly for organic PV is normally fabricated by hydrolysis and con-densation of titanium alkoxides at temperature below 150°C (39, 40), where the crystallization of TiO2 is slow. In order to
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prevent heat-induced degradation of the perovskite layer and the adjacent interfaces, we kept the annealing tempera-ture of TiOx films below 70°C. While such TiOx films have been used extensively in organic PVs, details on their struc-ture and mechanism of electric conduction have so far not been elucidated (41). One commonly recognized problem is that the amorphous nature of TiOx leads to extremely low specific conductivities that are in the range of 10?8–10?6 S cm–1 (42). Nb5+ doping has proved to be effective for enhanc-ing the conductivity of crystalline anatase TiO2 films to ~104 S cm–1, enabling its use as a transparent conducting oxide similar to conventional indium tin oxide (43). By analogy, this dopant is expected to improve also the conductivity of the amorphous TiOx via substitution of Ti4+ by Nb5+ which is expected to create donor centers. From the current-voltage (I–V) curves obtained by SPM measurement shown in Fig. 2B, the conductivity of TiOx film was estimated to increase from about 10?6 S cm–1 to 10?5 S cm–1 upon adding 5 mol% Nb5+ to the precursor solution. By resolving the XPS spectra in fig. S6 (14), the relative content of Ti3+ in comparison to Ti4+, i.e., the donor species in the TiOx film responsible for its n–type conductivity, has increased via Nb5+ doping.
We derived the optimal thickness of the NiO- and TiOx-based charge-extraction layers from the electrical measure-ments shown in fig. S7 (14). A complete layer with no pin-holes of NiO and TiOx requires at least thickness of 20 nm and 10 nm, respectively, regardless of the presence of do-pants. These minimum thicknesses should depend on the underlayers’ (FTO for NiO, or PCBM for TiOx) surface chem-istry and morphology, as well as the fabrication methods used for the NiO and TiOx films. We compared small solar cells with size of 0.09 cm2 and varied the NiO layer thick-ness from 10, 20 to 40 nm, keeping that of the Ti0.95Nb0.05Ox fixed at 10 nm. Conversely, we fixed the Li0.05Mg0.15Ni0.8O layer thickness at 20 nm and varied that of the TiOx films from 5 to 30 nm. For each condition, 20 cells were made and measured to establish any underlying trends. As shown in fig. S8 (14), although high performance can be occasional-ly achieved from devices with very thin charge carrier ex-traction layers (~10 nm NiO or ~5 nm TiOx), most devices showed lower PCEs because of lower FF and open–circuit voltage (Voc) (fig. S9) (14), which can be attributed to the presence of substantial levels of pinholes. The reproducibil-ity of device performance was greatly enhanced as the thickness of charge extraction layers increased, while the optimal performance was attained with 20 nm NiO and 10 nm TiOx film, in agreement with the electrical measure-ment. Further increasing the film thickness of the two charge extraction layers can lead to a large efficiency decline caused by increased internal resistance, larger optical loss, or both (fig. S10) (14). Thus, we fixed the thickness of NiO and TiOx at 20 nm and 10 nm, respectively for the following studies of doping effect on device performance.
Figure 3A shows the effect of doping the NiO and TiOx charge extraction layers on the photocurrent density–voltage (J–V) curves of PSCs measured under simulated AM 1.5 sunlight with forward scanning direction. The short cir-cuit current (Jsc), Voc, FF, and PCE data are listed in Table S2 (14). Both doping of NiO and TiOx reduced Rs and improved the FF, and to a lesser extent Jsc and Voc. The TiOx electron extraction layers mainly affect the shape of J–V curves in the forward bias range from 0.7 to 1.0 V, where a Schottky-barrier type contact between PCBM and Ag strongly re-stricted efficient electron collection (fig. S9A) (14, 44). The Nb5+ doping of TiOx reduced the interfacial electron transfer resistance and facilitated electron transport, increasing the photocurrent especially in the 0.7 to 1.0 V forward bias re-gion. The NiMg(Li)O-based hole extraction layer promoted ohmic contact formation at the FTO–MAPbI3 interface by decreasing the barrier height through the staircase energy level alignment shown in Fig. 1C. P+–doping increased the electrical conductivity by decreasing the charge transport resistance and hence enhancing hole extraction.
Doping of both NiO and TiOx improves the cell perfor-mance by increasing the values of FF and Voc to 0.827 and 1.083 V, respectively, leading to PCE of 18.3% for this planar PSC with MAPbI3. In comparison to PEDOT:PSS based PSCs, the Voc increased by ~100 mV, indicating that with Li0.05Mg0.15Ni0.8O, the holes can be extracted at a higher en-ergy level (10). Furthermore, the FF of 0.83 is one of the highest values for reported PSCs (8, 9, 23), demonstrating the key role of the charge extraction layers in minimizing resistive losses and improving the photovoltaic perfor-mance.
To gain further insight into the reasons for the perfor-mance enhancement by the doping, we characterized the charge carrier extraction, transportation and recombination by nanosecond time–resolved photoluminescence (PL) decay using a picosecond laser flash as excitation source and by measuring transient photocurrent/photovoltage decays on the microsecond scale. The charge extraction involved in our cells include the electron transfer from the MAPbI3 ab-sorber layer to PCBM/TiOx and hole transfer to NiO as well as the carrier transport in the TiOx and NiMg(Li)O layers. The perovskite/PCBM interface has been demonstrated to be very efficient for electron extraction (9–11). Doping of the TiOx extraction layer is unlikely to have a direct impact on the electron injection rate because of it physical separation from the MAPbI3 by 80 nm-thick PCBM layer. Nevertheless, it greatly accelerates the electron extraction by decreasing the electron transport time as shown in Fig. 3C. Figure 3B shows the PL decays of the MAPbI3 films on different sub-strates, including a glass slide, and NiO and Li0.05Mg0.15Ni0.8O deposited on FTO glass. The MAPbI3 itself showed a long PL lifetime of >100 ns, indicating slow carrier recombination in the perovskite layer (10). When contacted with the p–type hole extraction layers, the PL lifetimes were shortened to a similar degree for both doped and undoped NiO–MAPbI3 contacts. Thus, doping has a negligible influence on the hole injection.
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We derived the charge transport and recombination time cm2 device was certified (fig. S13) (14).
To demonstrate the superiority of the solution processi-constants (τt and τr) from the transient photocurrent and
photovoltage decays measured at short and open circuit, ble Ti0.95Nb0.05Ox charge carrier extraction layer, two refer-respectively (Fig. 3C). The τt decreased by ~5 fold from 4.41 ences, i.e., Ca (4nm)/Ag, LiF (1.5nm)/Ag that were deposited μs for undoped cell to 0.88 μs for doped cell, confirming the by thermal evaporation, were compared with our best inter-much faster charge transport through the doped charge car-facial condition. As shown in Fig. 4C, without sealing, the rier extraction layers compared to the undpoed ones. How-Ti(Nb)Ox based PSC shows the best stability since its PCE ever, the τr of a doped cell is substantially longer than that only decreased by ~5% of its initial value after 1 week. The of an undoped one (84.8 vs. 50.5 μs, Fig. 3D), which we at-Ca/Ag based PSC showed the fastest degradation, which lost tribute to slower interfacial charge recombination since the its initial PCE by >30% after 1 day and by ~45% within 1 very rapid carrier collection prevents charge accumulation week. This difference is attributed to the fast oxidization of at the interface of the pervovskite with the charge extraction very reactive Ca, leading to dramatic loss on Jsc and FF. The
LiF/Ag based PSC lost 15% of its initial PCE within 1 week. layer. The doping-induced difference in charge
transport/recombination kinetics should be the main reason The extremely thin LiF layer (< 2.5 nm, as required by effi-cient tunneling) (46) and the high sensitivity of LiF to mois-responsible for their performance enhancement.
The hysteresis of our cells was examined by using differ-ture is likely to be responsible for the corresponding cell’s ent scan rates and directions. By decreasing the step width inferior stability. It is possible that the stability of Ca or LiF from 70 mV to 5 mV, the PCEs determined from the forward based devices can be improved if they are thoroughly sealed. scan slightly increased from 18.14% to 18.35%. However, the However, the requirement on sealing quality will be much reverse scan PCEs decreased substantially from 22.35% at 70 more critical in comparison to the air–stable interface of mV/step (fig. S11A and table S3 (14)) to 18.40% at 5 mV/step. Ti0.95Nb0.05Ox (39).
The Ti0.95Nb0.05Ox layer also shields the perovskite from The steady power outputs measurements (fig. S11D) (14)
indicate that the PCEs obtained in forward scans and at the intrusion of humidity. It assumes a similar role in OPVs small step widths (5–10 mV) are near the real performance. (39). We exposed bare MAPbI3, MAPbI3/PCBM and MAP-The Voc and PCE obtained at a fast reverse scan, i.e., 1.273 V bI3/PCBM/Ti(Nb)Ox, to ambient air under room light for 3 and 22.35% at 70 mV/step, are largely overestimated. With weeks. A striking difference in color degradation associated the step widths of 5–10 mV, the PCE deviations between with perovskite decomposition became clearly visible (fig. forward and reverse scans are very small, i.e., within 0.3% in S14) (14). Thus it appears that the hydrophobic nature of absolute PCE values reflecting negligible hysteresis. A histo-PCBM may protect the perovskite from reaction with water, gram comparing the difference in the PCEs obtained from while the coating of Ti(Nb)Ox could further enhance the scanning in the forward and reverse bias directions is stability.
Figure 4D also shows the long–term stability of PSCs us-shown in fig. S11C (14), supporting the absence of hysteresis
for the optimized device architecture. In stark contrast, for ing the optimized inorganic charge extraction layers. The undoped charge extraction layers, a pronounced hysteresis silver back contact was protected by a covering glass which was observed even for slow scan rates (fig. S11E) (14), which was separated from the front FTO glass by a UV-activated is likely to arise from unbalanced charge accumulation at glue used as a sealant. The cells maintained 97% of their the two interfaces (45). Thus, the Li0.05Mg0.15Ni0.8O and initial PCE after keeping them in the dark for 1000 hours. Ti0.95Nb0.05Ox charge extraction layer create a robust low im-Exposing the cells for 1000 hours at short–circuit condition pedance interface that can mitigate the J–V hysteresis un-to full sunlight from a solar simulator, resulted in a PCE
degradation of less than 10%. This degradation is consistent der a routine scanning condition.
We fabricate cells with active area >1 cm2 as a first step with the general tendency among 10 devices, as shown in toward scale up of the photovoltaic devices. Figure 4A fig. S15 (14), indicative of their good long-term stability. Dur-shows the J–V curve of such cell with aperture area of 1.02 ing this time, an electric charge of around 72’000 C (4.49 × cm2. It shows excellent performance, with Jsc, Voc, and FF 1023 electrons) passed through the device. This result shows reaching values of 20.21 mA cm–2, 1.072 V and 0.748, respec-that the planar cell structure and the metal oxides extrac-tively, corresponding to a PCE of 16.2%. Hysteresis for these tion layers, as well as the organo–metal halide perovskite large area devices is also small (fig. S12) (14). The corre-material, are robust enough to sustain continued current sponding IPCE (Fig. 4B) shows a broad plateau with maxi-flow under light exposure for 1000 hours. A further increase mum value of 90.1% over practically the whole visible range. in the PCE without sacrificing stability could be expected The integrated Jsc from IPCE matches well with the meas-from varying the composition of the pervoskite, e.g., replac-ured value. Compared to small size cells (0.09 cm2), a ~10% ing part of the methylammonium cations in the MAPbI3 decrease in PCE was observed in large size cells (1.02 cm2), pervovskite by formamidinium (47) and a small fraction of which is mainly caused by the large sheet resistance of the the iodide by bromide anions. FTO. We sent one of our best large cells to a public test cen-REFERENCES AND NOTES
ter (AIST, Japan) for certification. A PCE of 15.0% for a 1.017 1. G. Hodes, Perovskite-based solar cells. Science 342, 317–318 (2013). Medline
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