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. 2018 Mar 26;3(3):3522–3529. doi: 10.1021/acsomega.8b00026

Elucidating the Impact of Thin Film Texture on Charge Transport and Collection in Perovskite Solar Cells

Chuanpeng Jiang †,‡, Yuling Xie ‡, Richard R Lunt §, Thomas W Hamann ‡, Pengpeng Zhang †,*
PMCID: PMC6641401  PMID: 31458603

Abstract

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Organic–inorganic halide perovskites have emerged as one of the most promising materials for photovoltaic applications. Because of the polycrystalline nature of perovskite thin films, it is crucial to investigate the impact of microstructures on device performance. In this study, we employ ramp-annealing to tailor the texture of perovskite thin films via controlling the nucleation of perovskite grains. Electrochemical impedance spectroscopy studies further suggest that the thin film texture impacts not only the charge collection at the contact but also the carrier transport in the bulk perovskite layer. The combination of the two effects leads to enhanced performance in devices constructed of preferentially oriented perovskite thin films.

1. Introduction

Recently, organic–inorganic hybrid perovskite solar cells have drawn a tremendous amount of attention because of their low-cost, facile fabrication, and high efficiency.1−6 Perovskite films can be effectively formed simply by spin-coating precursor solution onto substrates, followed by a low-temperature annealing treatment.3,7 A high light absorption coefficient has been established in this class of material, where ∼90% of light (wavelength: 350–750 nm) can be absorbed in a 400 nm thick film.7,8 Additionally, the low exciton binding energy on the order of tens of millielectron volt indicates that free charge carriers are predominantly generated under illumination.9−12 However, beyond the light harvesting capability, carrier transport and charge collection also collectively contribute to the efficiency of photovoltaic devices.13−17

Because of the polycrystalline nature of perovskite thin films, it is thus critical to investigate the impacts of microstructures, such as grain size, crystallinity, and texture structure, on the electrical properties of thin films.18−23 More interestingly, recent studies have suggested that the texture of the perovskite thin film may play a more significant role than the size of perovskite grains in determining the device efficiency.20,24 Thus far, time-resolved photoluminescence (tr-PL) is the most commonly adapted technique for investigating the influence of the crystal orientation and/or the texture structure on charge carrier dynamics in perovskite solar cells. Although variations in the carrier lifetime are observed for perovskite thin films with different textures,15,25−28 it is challenging to disentangle the texture effects on charge transport and collection processes because the trap states in the bulk and at the surface would both influence the carrier lifetime determined from the tr-PL measurements.29

In this paper, we conduct electrochemical impedance spectroscopy (EIS) measurements on perovskite devices with thin film crystallinity and texture controlled by ramp-annealing. The device performance is strongly correlated to the texture of perovskite thin films. EIS measurements that are capable of differentiating the charge transport and collection processes according to their responses to light intensity, applied bias, and frequency further reveal that the more textured perovskite thin film exhibits suppressed surface recombination at the contacts, which, along with the longer carrier diffusion length, ultimately leads to the enhanced short-circuit current (Jsc) and fill factor (FF).

2. Results and Discussion

2.1. Ramp-Annealing Treatment and the Crystalline Structure of Thin Films

Figure S1a shows the scanning electron microscopy (SEM) image of the CH3NH3PbI3 thin film deposited on PEDOT:PSS/ITO using the conventional one-step solution process,7 where the annealing is performed at 100 °C for 10 min which we term as the fast-annealing treatment. Beyond the relatively high density of cracks, the θ–2θ X-ray diffraction (XRD) scan of such prepared films in the Bragg–Brentano configuration (blue curve in Figure 1b) further reveals the polycrystalline nature of the film with perovskite grains oriented along the (110), (310), (112) planes, and so forth. To improve the surface morphology and to gain a better control over the film crystallinity and texture, we adapt the ramp-annealing treatment in our fabrication process. The schematics of the annealing profile are sketched in Figure S3a, which consists of three steps, that is, low temperature annealing, ramping, and final annealing at 100 °C. The rationale behind this treatment is twofold: (i) to induce heterogeneous nucleation of the perovskite phase at the interface between the film (solvent–solute) and the substrate and (ii) to remove solvent molecules when the film is still largely amorphous, thus minimizing cracking. The microstructure of the perovskite thin film is determined by the nucleation and growth processes. In comparison to the homogeneous nucleation that occurs in the bulk of the solvent–solute film, nucleation at the film/substrate interface can potentially lead to columnar structures with preferentially oriented grains guided by the interaction between the perovskite and the substrate (see Figure S3 for schematics). Because the activation energy of homogeneous nucleation is larger than that of the heterogeneous nucleation, it is desirable to first perform the thermal annealing at a relatively low temperature where only heterogeneous nucleation can occur,30,31 followed by the ramping up of the temperature to 100 °C to complete the crystallization/growth process.

Figure 1.

Figure 1

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XRD patterns of perovskite thin films (110 nm) annealed (a) at different low temperatures and (b) at 40 °C followed by different ramping rates to 100 °C; the ratio of the (110) and (310) peak intensities are listed on the side for comparison; J–V characteristics of the devices with the perovskite layer (c) annealed at temperatures corresponding to those in (a) during the low temperature step followed by 100 °C annealing (with a fixed ramping rate of 10 °C/min) and (d) processed at the identical conditions as those presented in (b). For comparison, XRD and device performance of the fast-annealed sample are also included in (b,d), respectively.

To obtain a rationally tuned texture structure in perovskite thin films, the parameters of the ramp-annealing procedure are further analyzed. We first explore the low temperature annealing step. Figure 1a shows the XRD patterns of the CH3NH3PbI3/PEDOT:PSS/ITO samples annealed for 10 min at 40, 50, and 60 °C. Grains oriented along the (110) plane are the most predominant, followed by those along the (310) plane. We then analyze the ratio of the (110) and (310) peak intensities to illustrate the texture of the film. The largest ratio is achieved in samples annealed at 40 °C, suggesting that the film processed at this condition has developed a more textured structure with a preferred grain orientation. In contrast, at higher annealing temperatures (50 or 60 °C), it is likely that the nucleation is induced both inside the film and at the interface, leading to a more randomly oriented crystal structure. The other critical parameter that influences the film quality is the ramping rate of the temperature from 40 to 100 °C. As shown in Figure 1b, the optimized value is 10 °C/min, which we attribute to the balance between the kinetic-limited and thermodynamic-limited processes. If the rate is too low, that is, 1 °C/min, homogeneous nucleation may be induced during the long duration of the temperature ramping up process, whereas at the high rate (100 °C/min), the growth of the interface-nucleated grains can be hindered by mass transport. The XRD patterns of samples treated with other ramping rates are illustrated in Figure S4.

2.2. Correlation between Device Performance and Thin Film Texture

To examine how the crystalline quality and texture of the perovskite film impact the device performance, J–V curves are measured on devices constructed of ITO/PEDOT:PSS/CH3NH3PbI3 (110 nm)/PCBM/BCP/Ag. In this structure, CH3NH3PbI3 is the active layer which absorbs light and creates the photocarriers that are subsequently collected by the anode (ITO) and cathode (Ag) through the hole-transport (PEDOT:PSS) and electron-transport (PCBM/BCP) layers, respectively. In Figure 1c,d and Table S1, we show the J–V characteristics of the devices with the perovskite layer processed at the varying annealing temperatures or ramping rates corresponding to those exploited in the films displayed in Figure 1a,b, respectively. These data strongly suggest that the trend of the device performance closely follows that of the active layer crystallinity as revealed in the XRD patterns, demonstrating that the texture of the perovskite film is essential in determining the device efficiency.

It is worth mentioning that for films that are not thermally treated (black curve in Figure 1a), an intermediate phase emerges at small 2θ angles (<10°), corresponding to a large lattice spacing along the out-of-plane direction. This intermediate phase is likely to consist of CH3NH3I–PbI2–DMSO (solvent molecule).32−34 Although solvent molecules can be removed during the fast-annealing treatment, the volume shrinkage in films that are rapidly undergoing the transition to the crystalline phase will result in cracks as those observed in the SEM image in Figure S1a. Ramp-annealed samples, on the other hand, exhibit smooth and compact morphology with minimized density of cracks, as shown in Figures S1b–d and S2. This is presumably owing to the removal of solvent molecules during the low temperature annealing step when the film is still largely amorphous.35 Nevertheless, even though the cracks present on the fast-annealed sample increase the surface roughness, the film is still continuous without deep pinholes or shunting paths, as shown in the cross-sectional SEM images in Figure S1f.

In the following discussion, we focus on the analysis and comparison between the fast-annealed sample and the ramp-annealed one treated under the optimized annealing condition (unless specified), corresponding to the blue and red curves in Figure 1b,d, respectively, as they represent the two ends of the spectrum in terms of film crystallinity/texture and device performance. The ramp-annealed samples outperform the fast-annealed ones in both the Jsc and the FF, which leads to an improved power conversion efficiency (PCE) from 6.3 to 8.7%. The typical J–V curves under the forward and reverse scans are plotted in Figure S5, demonstrating a negligible hysteresis for both types of the devices. The histogram displayed in Figure S6 further summarizes the performance of more than 40 devices from 10 batches of samples prepared on different days, where we consistently observe the enhancement of Jsc and FF, and therefore improved PCE, in ramp-annealed samples.

The Jsc can also be calculated from the spectral response of the cell using the equation Jsc = q∫ϕ(λ) × EQE(λ) dλ, where ϕ(λ)is the photon flux, and EQE, the external quantum efficiency, represents the ratio between the number of collected charge carriers and the number of incident photons at a given wavelength. Because EQE measures how efficiently the cell converts incident light into electrical energy, it is determined by the product of the light harvesting efficiency, the carrier transport efficiency, and the charge collection efficiency, that is, ηLH(λ) × ηCT(λ) × ηCC(λ). Figure 2a displays the corresponding EQE data of the two devices, where the ramp-annealed one shows consistently higher quantum efficiency over the entire wavelength range. However, as illustrated in Figure 2b, the absorbance spectrum obtained from the UV–vis measurement does not show any noticeable difference between the fast- and ramp-annealed CH3NH3PbI3/PEDOT:PSS/ITO samples, indicating that ηLH(λ)is not the origin of the EQE enhancement.

Figure 2.

Figure 2

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(a) EQE curves corresponding to the fast- and ramp-annealed devices with the 110 nm thick perovskite active layer, and the calculated Jsc from EQE (12.8 and 10.1 mA/cm2) matches well with the Jsc extracted from J–V curves (13.1 and 10.2 mA/cm2). (b) Absorbance spectra of fast- and ramp-annealed CH3NH3PbI3 (110 nm)/PEDOT:PSS/ITO samples.

2.3. Impacts of Thin Film Texture on Charge Collection and Transport

It is important to point out that the textured structure of the perovskite layer may impact the carrier transport efficiency (ηCT(λ)) because of the enhancement of the long-range structural coherence in the film.36,37 The textured structure could also improve the charge collection efficiency (ηCC(λ)) because the energy level alignment and charge transfer at the contacts can be highly dependent on the surface orientation/termination of the perovskite layer.15,38 To clarify the underlying mechanism of the improved EQE and Jsc in the ramp-annealed devices and to pinpoint the effects of the thin film texture, we further perform EIS measurements, which can potentially disentangle processes occurring in the bulk and at the contact. As shown in the Nyquist plot under 1 sun illumination (Figure S7a), two arcs are present which are well-separated in frequency. The associated capacitance and resistance of each arc is extracted by fitting the impedance data to the equivalent circuit shown in Figure 3a. To assign the physical processes that correspond to the two arcs, EIS measurements are performed as a function of applied bias and under different light intensities at both short- and open-circuit conditions.

Figure 3.

Figure 3

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(a) Equivalent circuit for EIS data fitting and the dependence of capacitances (b) and resistances (c) on light intensity under the open-circuit condition. (d) Calculated interface recombination time constants under the short- and open-circuit conditions, respectively, at different light intensities for both fast- and ramp-annealed devices (110 nm perovskite layer).

The high frequency arc is associated with a capacitance, C2, on the order of 1 × 10–7 F cm–2, for both ramp-annealed and fast-annealed samples (Figure 3b). The magnitude of C2 is fairly constant with respect to applied bias (Figure S7c) and incident light intensities (Figure 3b), which is consistent with prior reports which have established this capacitance as the bulk dielectric capacitance of the perovskite layer.39,40 The capacitance, C1, which is associated with the low frequency arc, has previously been assigned to ionic or electrical accumulation layers at the perovskite interface with charge selective contacts, for example, accumulated holes at the cathode.40,41 While C1 is generally found to increase with light intensity, the value of C1 for the ramp-annealed samples is approximately 2 orders of magnitude larger than that of the fast-annealed ones (Figure 3b).

As shown in Figure 3c, the resistance associated with the bulk capacitance, R2, can be a few times larger than the resistance associated with the interface capacitance, R1, for both ramp-annealed and fast-annealed samples. Both resistances exhibit a similar behavior of decreasing with the increasing light intensity (Figure 3c) as well as with applied bias (Figure S7c), which is opposite to the trend observed for the interface capacitance (C1). As recently noted by Bisquert and co-workers,41 these observations suggest that the resistances are connected to a common process of recombination at the interface. In this case, the interface recombination kinetics are described by the time constant given by τs = (R1 + R2)C1. Figure 3d shows that at the open circuit, τs,oc is consistently 2 orders of magnitude larger for the ramp-annealed samples, indicating significantly faster recombination at the interface for the fast-annealed samples. Given this fact, it seems surprising that the fast-annealed cells produce the same open-circuit photovoltage, Voc, as the ramp-annealed samples. There has been a recent debate centered on the physical meaning of the low frequency response (R1 and C1).41−43 Therefore, to test the connection of τs,oc and Voc, J–V curves are also measured as a function of incident light intensity. Figure 4a shows a plot of Voc versus intensity for both ramp-annealed and fast-annealed samples. Fits of these trends to the diode equation reveal a nearly ideal diode quality factor, γ, of 1.4 for the ramp-annealed samples, whereas γ is 5.6 for the fast-annealed samples, suggesting that they are controlled by different recombination mechanisms. A close examination of the Voc versus intensity relation for the fast-annealed sample further reveals that the quality factor decreases with light intensity, and a γ of 5.6 is a simple average. This can be attributed to the filling of trap states or recombination centers associated with the interfaces by photogenerated carriers, resulting in the gradual drop of the quality factor to a value close to that of the ramp-annealed sample under sufficient illumination (>0.2 sun).44 The diode quality factors further accounts for the larger FF which is consistently measured for the ramp-annealed compared to the fast-annealed samples.

Figure 4.

Figure 4

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(a) Light Intensity dependence of Voc for fast- and ramp-annealed devices (110 nm perovskite). (b) Jsc and PCE dependence on the perovskite layer thickness in the fast- and ramp-annealed devices.

The Jsc is found to increase linearly with light intensity for both samples (Figure S8). The interfacial recombination time constants are also determined at short circuit, where τs,sc is found to be larger than τs,oc for both samples at all light intensities (Figure 3d). Unlike at the open circuit, however, the time constants of the ramp-annealed and fast-annealed samples are fairly close at the short circuit (see Figure S7d for resistances and capacitances at varying intensity under the short circuit). This suggests that interfacial recombination may not account for charge collection losses. It was recently suggested that the collection efficiency can be calculated from the ratio of the recombination resistance (R1 + R2) at the open circuit and the short circuit according to ηcc = 1 – Roc/Rsc.41 This analysis produces nominally identical values of ηcc = 0.985 for the ramp-annealed samples and ηcc = 0.978 for the fast-annealed samples. Thus, differences in charge collection efficiencies cannot completely account for the discrepancy in Jsc observed. Instead, if this photocurrent discrepancy mainly originates from the bulk diffusion/recombination processes through differences in the carrier diffusion length, it should be reflected in the device performance as a function of the perovskite layer thickness.45−47 By varying the concentration of the precursor solution, film thickness can be controlled between 110 and 450 nm, as listed in Table S2. Indeed, the highest PCE for the fast- and ramp-annealed samples is achieved at the thickness of 200 nm (7.0%) and 290 nm (12.1%), respectively, as illustrated in Figure 4b, suggesting a longer carrier diffusion length associated with the latter.

Because the Jsc and PCE variations between the ramp- and fast-annealed samples are most pronounced at the perovskite layer thickness of 290 nm, we further perform corresponding EIS measurements at this thickness which might provide more information on the recombination parameters under the short-circuit condition. The associated capacitances and resistances, as well as the derived recombination time constants are plotted in Figure 5. In contrast to the comparable τs,sc in thin devices, τs,sc of the 290 nm thick ramp- and fast-annealed samples differs by 2 orders of magnitude, as shown in Figure 5c, which originates from the larger R1 and R2 in the former (Figure 5b). Recently, it has been argued that the applied bias in EIS measurements can lead to ionic transport and polarization of the interfaces in devices with the regular structure, which consequently interferes with the EIS measurements and changes the mechanism of recombination from bulk- to surface-dominant.48 The same ionic processes will also lead to a gradual increase of Rs (defined in Figure 3a) over time and introduce hysteresis in J–V curves.48,49 Nevertheless, as shown in Figures S5 and S7b, no obvious J–V hysteresis or change in Rs is observed in our experiments, indicating that the ion migration/polarization issue is not as severe in our inverted-structured devices. This can be attributed to the surface passivation effect of fullerene molecules,50 as well as the avoidance of chemical reactions as these occur at the interfaces between the perovskite and TiO2/Spiro-OMeTAD in devices with the regular structure.48 Therefore, the larger R1 and R2 as measured in the ramp-annealed sample under the short-circuit condition (Figure 5b) provide a strong indication that beyond the surface recombination at contacts, the bulk recombination process is also suppressed in the ramp-annealed devices, which ultimately leads to the longer carrier diffusion length, as illustrated in Figure 4b. Note that similar to the thin devices, significant variation in τs,oc, mainly contributed by the modulation in C1, that is, interface capacitance, is also observed between the ramp- and fast-annealed 290 nm thick samples, as illustrated in Figures S9b and 5c.

Figure 5.

Figure 5

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Dependence of capacitances (a) and resistances (b) on light intensity under the short-circuit condition for both ramp- and fast-annealed devices with the 290 nm thick perovskite layer; and (c) calculated interface recombination time constants under the short- and open-circuit conditions at different light intensities.

Although there are theoretical predictions that defects in the perovskite layer mainly contribute to the shallow traps and thus are not detrimental to the device performance,51−53 our study suggests that it is still crucial to improve the crystallinity and texture of the perovskite layer to boost the device performance. On the one hand, the recombination at the contacts is expected to be the dominant loss mechanism in perovskite solar cells.14,54,55 Instead of introducing the interfacial layer at the contact54−56 or passivating the surface trap states with small molecules57,58 which could complicate the device fabrication process, our study demonstrates that the recombination kinetics at the interfaces can be effectively suppressed simply by the ramp-annealing treatment via controlling the surface orientations or terminations of perovskite grains. On the other hand, it is likely that the preferential crystal orientation in the ramp-annealed sample yields an increase in the density of low-angle grain boundaries in the polycrystalline perovskite thin film, which, as compared to the large-angle grain boundaries, exhibit better carrier transport properties with minimized bulk carrier recombination.

3. Conclusions

In this paper, we demonstrate that the texture of the perovskite thin film influences both the surface recombination at the contacts and the carrier diffusion length in the bulk. The combination of the two effects leads to enhanced performance in devices constructed of preferentially oriented perovskite thin films. These findings could aid in the simple design and fabrication of planar-structured high-efficiency perovskite solar cells.

4. Methods

4.1. Device Fabrication

PEDOT:PSS solution (Heraeus Technology Group) was spun onto patterned ITO/glass substrates (Xinyan Technology Ltd.) which were precleaned in acetone and isopropanol at 4000 rpm for 40 s. The film was then annealed at 150 °C for 20 min in a nitrogen-filled glovebox. PbI2 (2.88 g) and CH3NH3I (MAI, 0.636 g) were dissolved in a 5 mL mixture of dimethylformamide and dimethyl sulfoxide and stirred at 60 °C for 12 h. The precursor solution was then spun onto the PEDOT:PSS layer at 1000 rpm for 15 s and 5000 rpm for another 30 s. For the fast-annealing process, the sample was put on the hotplate and annealed at 100 °C for 10 min. The ramp-annealing sample was first heated at a low temperature (40–60 °C) for 10 min, followed by the ramping and annealing at 100 °C (10 min). PCBM (2 wt % in chlorobenzene) was spun onto the perovskite film at 1500 rpm for 30 s. The resulting samples were transferred to a vacuum deposition chamber (Angstrom Engineering), where BCP (5 nm) and Ag (100 nm) were thermally deposited with shadow masks under a deposition pressure of ∼3.0 × 10–6 Torr.

4.2. Characterization

Solar cell efficiencies are characterized in air using a Keithley 2420 source meter and a Newport solar simulator under 100 mW cm–2 illumination measured with an NREL calibrated mc-Si detector with a KG5 filter. Filters are applied to adjust the light intensities from 0.0092 to 1 sun for the J–V measurements at varied light intensities. The EQE measurements were carried out on a setup comprising a Xe lamp, a monochromator, a current–voltage preamplifier, and a lock-in amplifier. The light spectrum was determined with a monocrystalline photodetector calibrated by the National Institute of Standard and Technology. Devices were corrected for spectral mismatch (M) with values of approximately 0.98 < M < 1.01. The film morphology of the perovskite layer was studied by Hitachi S-4700II field emission SEM with the acceleration voltage of 15 kV. XRD patterns of the polycrystalline perovskites were characterized by a Bruker DAVINCI diffractometer using Cu Kα radiation (Rigaku D/MAX operating at 45 kV and 40 mA). Absorbance spectra were calculated from the transmission and reflection measured using a dual-beam Lambda 800 UV/vis spectrometer. EIS was conducted on an Autolab potentiostat coupled with Nova electrochemical software using a 50 mV amplitude perturbation and sweeping frequencies from 0.5 to 400 kHz. Capacitance and resistance data were fit using ZView software (Scribner Associates).

Acknowledgments

We acknowledge the financial support from Michigan State University Strategic Partnership Grant.

Supporting Information Available

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsomega.8b00026.

  • SEM images and XRD patterns, schematics of nucleation and crystallization, J–V characteristics and histogram of J–V parameters, impedance data analysis, light intensity dependence of Jsc, and tables of J–V parameters (PDF)

The authors declare no competing financial interest.

Supplementary Material

ao8b00026_si_001.pdf (1.4MB, pdf)

References

  1. Snaith H. J. Perovskites: The Emergence of a New Era for Low-Cost, High-Efficiency Solar Cells. J. Phys. Chem. Lett. 2013, 4, 3623–3630. 10.1021/jz4020162. [DOI] [Google Scholar]
  2. Best Research-Cell Efficiencies by National Renewable Energy Laboratory. http://www.nrel.gov/pv/assets/images/efficiency_chart.jpg (accessed December 2017).
  3. Lee M. M.; Teuscher J.; Miyasaka T.; Murakami T. N.; Snaith H. J. Efficient hybrid solar cells based on meso-superstructured organometal halide perovskites. Science 2012, 338, 643–647. 10.1126/science.1228604. [DOI] [PubMed] [Google Scholar]
  4. Li X.; Bi D.; Yi C.; Decoppet J.-D.; Luo J.; Zakeeruddin S. M.; Hagfeldt A.; Gratzel M. A vacuum flash-assisted solution process for high-efficiency large-area perovskite solar cells. Science 2016, 353, 58–62. 10.1126/science.aaf8060. [DOI] [PubMed] [Google Scholar]
  5. Arora N.; Dar M. I.; Hinderhofer A.; Pellet N.; Schreiber F.; Zakeeruddin S. M.; Grätzel M. Perovskite solar cells with CuSCN hole extraction layers yield stabilized efficiencies greater than 20%. Science 2017, 358, 768. 10.1126/science.aam5655. [DOI] [PubMed] [Google Scholar]
  6. Shin S. S.; Yeom E. J.; Yang W. S.; Hur S.; Kim M. G.; Im J.; Seo J.; Noh J. H.; Seok S. I. Colloidally prepared La-doped BaSnO3 electrodes for efficient, photostable perovskite solar cells. Science 2017, 356, 167. 10.1126/science.aam6620. [DOI] [PubMed] [Google Scholar]
  7. Docampo P.; Ball J. M.; Darwich M.; Eperon G. E.; Snaith H. J. Efficient organometal trihalide perovskite planar-heterojunction solar cells on flexible polymer substrates. Nat. Commun. 2013, 4, 2761. 10.1038/ncomms3761. [DOI] [PubMed] [Google Scholar]
  8. Sun S.; Salim T.; Mathews N.; Duchamp M.; Boothroyd C.; Xing G.; Sum T. C.; Lam Y. M. The origin of high efficiency in low-temperature solution-processable bilayer organometal halide hybrid solar cells. Energy Environ. Sci. 2014, 7, 399–407. 10.1039/c3ee43161d. [DOI] [Google Scholar]
  9. D’Innocenzo V.; Grancini G.; Alcocer M. J. P.; Kandada A. R. S.; Stranks S. D.; Lee M. M.; Lanzani G.; Snaith H. J.; Petrozza A. Excitons versus free charges in organo-lead tri-halide perovskites. Nat. Commun. 2014, 5, 3586. 10.1038/ncomms4586. [DOI] [PubMed] [Google Scholar]
  10. Galkowski K.; Mitioglu A.; Miyata A.; Plochocka P.; Portugall O.; Eperon G. E.; Wang J. T.-W.; Stergiopoulos T.; Stranks S. D.; Snaith H. J.; Nicholas R. J. Determination of the exciton binding energy and effective masses for methylammonium and formamidinium lead tri-halide perovskite semiconductors. Energy Environ. Sci. 2016, 9, 962–970. 10.1039/c5ee03435c. [DOI] [Google Scholar]
  11. Miyata A.; Mitioglu A.; Plochocka P.; Portugall O.; Wang J. T.-W.; Stranks S. D.; Snaith H. J.; Nicholas R. J. Direct measurement of the exciton binding energy and effective masses for charge carriers in organic-inorganic tri-halide perovskites. Nat. Phys. 2015, 11, 582–587. 10.1038/nphys3357. [DOI] [Google Scholar]
  12. Yang Z.; Surrente A.; Galkowski K.; Bruyant N.; Maude D. K.; Haghighirad A. A.; Snaith H. J.; Plochocka P.; Nicholas R. J. Unraveling the Exciton Binding Energy and the Dielectric Constant in Single-Crystal Methylammonium Lead Triiodide Perovskite. J. Phys. Chem. Lett. 2017, 8, 1851–1855. 10.1021/acs.jpclett.7b00524. [DOI] [PubMed] [Google Scholar]
  13. Zhao T.; Shi W.; Xi J.; Wang D.; Shuai Z. Intrinsic and Extrinsic Charge Transport in CH3NH3PbI3 Perovskites Predicted from First-Principles. Sci. Rep. 2016, 7, 19968. 10.1038/srep19968. [DOI] [PMC free article] [PubMed] [Google Scholar]
  14. Lee Y. H.; Luo J.; Son M.-K.; Gao P.; Cho K. T.; Seo J.; Zakeeruddin S. M.; Grätzel M.; Nazeeruddin M. K. Enhanced Charge Collection with Passivation Layers in Perovskite Solar Cells. Adv. Mater. 2016, 28, 3966–3972. 10.1002/adma.201505140. [DOI] [PubMed] [Google Scholar]
  15. Bae S.; Park J.-S.; Han I. K.; Shin T. J.; Jo W. H. CH3NH3PbI3 crystal orientation and photovoltaic performance of planar heterojunction perovskite solar cells. Sol. Energy Mater. Sol. Cells 2017, 160, 77–84. 10.1016/j.solmat.2016.10.019. [DOI] [Google Scholar]
  16. Baena J. P. C.; Steier L.; Tress W.; Saliba M.; Neutzner S.; Matsui T.; Giordano F.; Jacobsson T. J.; Kandada A. R. S.; Zakeeruddin S. M.; Petrozza A.; Abate A.; Nazeeruddin M. K.; Grätzel M.; Hagfeldt A. Highly efficient planar perovskite solar cells through band alignment engineering. Energy Environ. Sci. 2015, 8, 2928–2934. 10.1039/c5ee02608c. [DOI] [Google Scholar]
  17. Kim Y. C.; Yang T.-Y.; Jeon N. J.; Im J.; Jang S.; Shin T. J.; Shin H.-W.; Kim S.; Lee E.; Kim S.; Noh J. H.; Seok S. I.; Seo J. Engineering interface structures between lead halide perovskite and copper phthalocyanine for efficient and stable perovskite solar cells. Energy Environ. Sci. 2017, 10, 2109–2116. 10.1039/c7ee01931a. [DOI] [Google Scholar]
  18. Tsai C.-M.; Wu G.-W.; Narra S.; Chang H.-M.; Mohanta N.; Wu H.-P.; Wang C.-L.; Diau E. W.-G. Control of preferred orientation with slow crystallization for carbon-based mesoscopic perovskite solar cells attaining efficiency 15%. J. Mater. Chem. A 2017, 5, 739–747. 10.1039/c6ta09036b. [DOI] [Google Scholar]
  19. Shargaieva O.; Lang F.; Rappich J.; Dittrich T.; Klaus M.; Meixner M.; Genzel C.; Nickel N. H. Influence of the Grain Size on the Properties of CH3NH3PbI3 Thin Films. ACS Appl. Mater. Interfaces 2017, 9, 38428–38435. 10.1021/acsami.7b10056. [DOI] [PubMed] [Google Scholar]
  20. Docampo P.; Hanusch F. C.; Giesbrecht N.; Angloher P.; Ivanova A.; Bein T. Influence of the orientation of methylammonium lead iodide perovskite crystals on solar cell performance. APL Mater. 2014, 2, 081508. 10.1063/1.4890244. [DOI] [Google Scholar]
  21. Huang F.; Pascoe A. R.; Wu W.-Q.; Ku Z.; Peng Y.; Zhong J.; Caruso R. A.; Cheng Y.-B. Effect of the Microstructure of the Functional Layers on the Efficiency of Perovskite Solar Cells. Adv. Mater. 2017, 29, 1601715. 10.1002/adma.201601715. [DOI] [PubMed] [Google Scholar]
  22. Nie W.; Tsai H.; Blancon J.-C.; Liu F.; Stoumpos C. C.; Traore B.; Kepenekian M.; Durand O.; Katan C.; Tretiak S.; Crochet J.; Ajayan P. M.; Kanatzidis M.; Even J.; Mohite A. D. Critical Role of Interface and Crystallinity on the Performance and Photostability of Perovskite Solar Cell on Nickel Oxide. Adv. Mater. 2017, 30, 1703879. 10.1002/adma.201703879. [DOI] [PubMed] [Google Scholar]
  23. Long M.; Zhang T.; Zhu H.; Li G.; Wang F.; Guo W.; Chai Y.; Chen W.; Li Q.; Wong K. S.; Xu J.; Yan K. Textured CH3NH3PbI3 thin film with enhanced stability for high performance perovskite solar cells. Nano Energy 2017, 33, 485–496. 10.1016/j.nanoen.2017.02.002. [DOI] [Google Scholar]
  24. Giesbrecht N.; Schlipf J.; Oesinghaus L.; Binek A.; Bein T.; Müller-Buschbaum P.; Docampo P. Synthesis of Perfectly Oriented and Micrometer-Sized MAPbBr3 Perovskite Crystals for Thin-Film Photovoltaic Applications. ACS Energy Lett. 2016, 1, 150–154. 10.1021/acsenergylett.6b00050. [DOI] [Google Scholar]
  25. Li S.-S.; Chang C.-H.; Wang Y.-C.; Lin C.-W.; Wang D.-Y.; Lin J.-C.; Chen C.-C.; Sheu H.-S.; Chia H.-C.; Wu W.-R.; Jeng U. S.; Liang C.-T.; Sankar R.; Chou F.-C.; Chen C.-W. Intermixing-seeded growth for high-performance planar heterojunction perovskite solar cells assisted by precursor-capped nanoparticles. Energy Environ. Sci. 2016, 9, 1282–1289. 10.1039/c5ee03229f. [DOI] [Google Scholar]
  26. Chen H.; Zheng X.; Li Q.; Yang Y.; Xiao S.; Hu C.; Bai Y.; Zhang T.; Wong K. S.; Yang S. An amorphous precursor route to the conformable oriented crystallization of CH3NH3PbBr3 in mesoporous scaffolds: toward efficient and thermally stable carbon-based perovskite solar cells. J. Mater. Chem. A 2016, 4, 12897–12912. 10.1039/c6ta06115j. [DOI] [Google Scholar]
  27. Niu G.; Yu H.; Li J.; Wang D.; Wang L. Controlled orientation of perovskite films through mixed cations toward high performance perovskite solar cells. Nano Energy 2016, 27, 87–94. 10.1016/j.nanoen.2016.06.053. [DOI] [Google Scholar]
  28. Ye J.; Zheng H.; Zhu L.; Zhang X.; Jiang L.; Chen W.; Liu G.; Pan X.; Dai S. High-temperature shaping perovskite film crystallization for solar cell fast preparation. Sol. Energy Mater. Sol. Cells 2017, 160, 60–66. 10.1016/j.solmat.2016.10.022. [DOI] [Google Scholar]
  29. Wu B.; Nguyen H. T.; Ku Z.; Han G.; Giovanni D.; Mathews N.; Fan H. J.; Sum T. C. Discerning the Surface and Bulk Recombination Kinetics of Organic–Inorganic Halide Perovskite Single Crystals. Adv. Energy Mater. 2016, 6, 1600551. 10.1002/aenm.201600551. [DOI] [Google Scholar]
  30. Bassiri-Gharb N.; Bastani Y.; Bernal A. Chemical solution growth of ferroelectric oxide thin films and nanostructures. Chem. Soc. Rev. 2014, 43, 2125–2140. 10.1039/c3cs60250h. [DOI] [PubMed] [Google Scholar]
  31. Schwartz R. W.; Schneller T.; Waser R. Chemical solution deposition of electronic oxide films. C. R. Chim. 2004, 7, 433–461. 10.1016/j.crci.2004.01.007. [DOI] [Google Scholar]
  32. Jeon N. J.; Noh J. H.; Kim Y. C.; Yang W. S.; Ryu S.; Seok S. I. Solvent engineering for high-performance inorganic-organic hybrid perovskite solar cells. Nat. Mater. 2014, 13, 897–903. 10.1038/nmat4014. [DOI] [PubMed] [Google Scholar]
  33. Zuo L.; Dong S.; De Marco N.; Hsieh Y.-T.; Bae S.-H.; Sun P.; Yang Y. Morphology Evolution of High Efficiency Perovskite Solar Cells via Vapor Induced Intermediate Phases. J. Am. Chem. Soc. 2016, 138, 15710–15716. 10.1021/jacs.6b09656. [DOI] [PubMed] [Google Scholar]
  34. Rong Y.; Venkatesan S.; Guo R.; Wang Y.; Bao J.; Li W.; Fan Z.; Yao Y. Critical kinetic control of non-stoichiometric intermediate phase transformation for efficient perovskite solar cells. Nanoscale 2016, 8, 12892–12899. 10.1039/c6nr00488a. [DOI] [PubMed] [Google Scholar]
  35. Kang R.; Kim J.-E.; Yeo J.-S.; Lee S.; Jeon Y.-J.; Kim D.-Y. Optimized Organometal Halide Perovskite Planar Hybrid Solar Cells via Control of Solvent Evaporation Rate. J. Phys. Chem. C 2014, 118, 26513–26520. 10.1021/jp508015c. [DOI] [Google Scholar]
  36. Shi D.; Adinolfi V.; Comin R.; Yuan M.; Alarousu E.; Buin A.; Chen Y.; Hoogland S.; Rothenberger A.; Katsiev K.; Losovyj Y.; Zhang X.; Dowben P. A.; Mohammed O. F.; Sargent E. H.; Bakr O. M. Low trap-state density and long carrier diffusion in organolead trihalide perovskite single crystals. Science 2015, 347, 519–522. 10.1126/science.aaa2725. [DOI] [PubMed] [Google Scholar]
  37. Choi J. J.; Yang X.; Norman Z. M.; Billinge S. J. L.; Owen J. S. Structure of methylammonium lead iodide within mesoporous titanium dioxide: active material in high-performance perovskite solar cells. Nano Lett. 2014, 14, 127–133. 10.1021/nl403514x. [DOI] [PubMed] [Google Scholar]
  38. Yin J.; Cortecchia D.; Krishna A.; Chen S.; Mathews N.; Grimsdale A. C.; Soci C. Interfacial Charge Transfer Anisotropy in Polycrystalline Lead Iodide Perovskite Films. J. Phys. Chem. Lett. 2015, 6, 1396–1402. 10.1021/acs.jpclett.5b00431. [DOI] [PubMed] [Google Scholar]
  39. Almora O.; Zarazua I.; Mas-Marza E.; Mora-Sero I.; Bisquert J.; Garcia-Belmonte G. Capacitive Dark Currents, Hysteresis, and Electrode Polarization in Lead Halide Perovskite Solar Cells. J. Phys. Chem. Lett. 2015, 6, 1645–1652. 10.1021/acs.jpclett.5b00480. [DOI] [PubMed] [Google Scholar]
  40. Guerrero A.; Garcia-Belmonte G.; Mora-Sero I.; Bisquert J.; Kang Y. S.; Jacobsson T. J.; Correa-Baena J.-P.; Hagfeldt A. Properties of Contact and Bulk Impedances in Hybrid Lead Halide Perovskite Solar Cells Including Inductive Loop Elements. J. Phys. Chem. C 2016, 120, 8023–8032. 10.1021/acs.jpcc.6b01728. [DOI] [Google Scholar]
  41. Zarazua I.; Han G.; Boix P. P.; Mhaisalkar S.; Fabregat-Santiago F.; Mora-Seró I.; Bisquert J.; Garcia-Belmonte G. Surface Recombination and Collection Efficiency in Perovskite Solar Cells from Impedance Analysis. J. Phys. Chem. Lett. 2016, 7, 5105–5113. 10.1021/acs.jpclett.6b02193. [DOI] [PubMed] [Google Scholar]
  42. Contreras-Bernal L.; Salado M.; Todinova A.; Calio L.; Ahmad S.; Idígoras J.; Anta J. A. Origin and Whereabouts of Recombination in Perovskite Solar Cells. J. Phys. Chem. C 2017, 121, 9705–9713. 10.1021/acs.jpcc.7b01206. [DOI] [Google Scholar]
  43. Pockett A.; Eperon G. E.; Sakai N.; Snaith H. J.; Peter L. M.; Cameron P. J. Microseconds, milliseconds and seconds: deconvoluting the dynamic behaviour of planar perovskite solar cells. Phys. Chem. Chem. Phys. 2017, 19, 5959–5970. 10.1039/c6cp08424a. [DOI] [PubMed] [Google Scholar]
  44. Stranks S. D.; Burlakov V. M.; Leijtens T.; Ball J. M.; Goriely A.; Snaith H. J. Recombination Kinetics in Organic-Inorganic Perovskites: Excitons, Free Charge, and Subgap States. Phys. Rev. Appl. 2014, 2, 034007. 10.1103/physrevapplied.2.034007. [DOI] [Google Scholar]
  45. Nelson J.The Physics of Solar Cells; Imperial College Press: U.K., 2003. [Google Scholar]
  46. Stranks S. D.; Eperon G. E.; Grancini G.; Menelaou C.; Alcocer M. J. P.; Leijtens T.; Herz L. M.; Petrozza A.; Snaith H. J. Electron-Hole Diffusion Lengths Exceeding 1 Micrometer in an Organometal Trihalide Perovskite Absorber. Science 2013, 342, 341–344. 10.1126/science.1243982. [DOI] [PubMed] [Google Scholar]
  47. Adhyaksa G. W. P.; Veldhuizen L. W.; Kuang Y.; Brittman S.; Schropp R. E. I.; Garnett E. C. Carrier Diffusion Lengths in Hybrid Perovskites: Processing, Composition, Aging, and Surface Passivation Effects. Chem. Mater. 2016, 28, 5259–5263. 10.1021/acs.chemmater.6b00466. [DOI] [Google Scholar]
  48. Correa-Baena J.-P.; Turren-Cruz S.-H.; Tress W.; Hagfeldt A.; Aranda C.; Shooshtari L.; Bisquert J.; Guerrero A. Changes from Bulk to Surface Recombination Mechanisms between Pristine and Cycled Perovskite Solar Cells. ACS Energy Lett. 2017, 2, 681–688. 10.1021/acsenergylett.7b00059. [DOI] [Google Scholar]
  49. Richardson G.; O’Kane S. E. J.; Niemann R. G.; Peltola T. A.; Foster J. M.; Cameron P. J.; Walker A. B. Can slow-moving ions explain hysteresis in the current-voltage curves of perovskite solar cells?. Energy Environ. Sci. 2016, 9, 1476–1485. 10.1039/c5ee02740c. [DOI] [Google Scholar]
  50. Shao Y.; Xiao Z.; Bi C.; Yuan Y.; Huang J. Origin and elimination of photocurrent hysteresis by fullerene passivation in CH3NH3PbI3 planar heterojunction solar cells. Nat. Commun. 2014, 5, 5784. 10.1038/ncomms6784. [DOI] [PubMed] [Google Scholar]
  51. Yin W.-J.; Shi T.; Yan Y. Unique Properties of Halide Perovskites as Possible Origins of the Superior Solar Cell Performance. Adv. Mater. 2014, 26, 4653–4658. 10.1002/adma.201306281. [DOI] [PubMed] [Google Scholar]
  52. Walsh A.; Scanlon D. O.; Chen S.; Gong X. G.; Wei S.-H. Self-Regulation Mechanism for Charged Point Defects in Hybrid Halide Perovskites. Angew. Chem., Int. Ed. 2015, 54, 1791–1794. 10.1002/anie.201409740. [DOI] [PMC free article] [PubMed] [Google Scholar]
  53. Kim J.; Lee S.-H.; Lee J. H.; Hong K.-H. The Role of Intrinsic Defects in Methylammonium Lead Iodide Perovskite. J. Phys. Chem. Lett. 2014, 5, 1312–1317. 10.1021/jz500370k. [DOI] [PubMed] [Google Scholar]
  54. Hou Y.; Chen W.; Baran D.; Stubhan T.; Luechinger N. A.; Hartmeier B.; Richter M.; Min J.; Chen S.; Quiroz C. O. R.; Li N.; Zhang H.; Heumueller T.; Matt G. J.; Osvet A.; Forberich K.; Zhang Z.-G.; Li Y.; Winter B.; Schweizer P.; Spiecker E.; Brabec C. J. Overcoming the Interface Losses in Planar Heterojunction Perovskite-Based Solar Cells. Adv. Mater. 2016, 28, 5112–5120. 10.1002/adma.201504168. [DOI] [PubMed] [Google Scholar]
  55. Marin-Beloqui J. M.; Lanzetta L.; Palomares E. Decreasing Charge Losses in Perovskite Solar Cells Through mp-TiO2/MAPI Interface Engineering. Chem. Mater. 2016, 28, 207–213. 10.1021/acs.chemmater.5b03902. [DOI] [Google Scholar]
  56. Wang L.; Moghe D.; Hafezian S.; Chen P.; Young M.; Elinski M.; Martinu L.; Kéna-Cohen S.; Lunt R. R. Alkali Metal Halide Salts as Interface Additives to Fabricate Hysteresis-Free Hybrid Perovskite-Based Photovoltaic Devices. ACS Appl. Mater. Interfaces 2016, 8, 23086–23094. 10.1021/acsami.6b07368. [DOI] [PubMed] [Google Scholar]
  57. Noel N. K.; Abate A.; Stranks S. D.; Parrott E. S.; Burlakov V. M.; Goriely A.; Snaith H. J. Enhanced Photoluminescence and Solar Cell Performance via Lewis Base Passivation of Organic–Inorganic Lead Halide Perovskites. ACS Nano 2014, 8, 9815–9821. 10.1021/nn5036476. [DOI] [PubMed] [Google Scholar]
  58. Abate A.; Saliba M.; Hollman D. J.; Stranks S. D.; Wojciechowski K.; Avolio R.; Grancini G.; Petrozza A.; Snaith H. J. Supramolecular Halogen Bond Passivation of Organic–Inorganic Halide Perovskite Solar Cells. Nano Lett. 2014, 14, 3247–3254. 10.1021/nl500627x. [DOI] [PubMed] [Google Scholar]

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