Characterization of Low-Frequency Excess Noise in CH3NH3PbI3

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Characterization of Low-frequency Excess Noise in CH3NH3PbI3-based Solar Cells Grown by Solution and Hybrid Chemical Vapor Deposition Techniques Qian Shen, Annie Ng, Zhiwei Ren, Huseyin Cem Gokkaya, Aleksandra B. Djurisic, Juan Antonio Zapien, and Charles Surya ACS Appl. Mater. Interfaces, Just Accepted Manuscript • Publication Date (Web): 02 Nov 2017 Downloaded from http://pubs.acs.org on November 2, 2017

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ACS Applied Materials & Interfaces

Characterization of Low-frequency Excess Noise in CH3NH3PbI3-based Solar Cells Grown by Solution and Hybrid Chemical Vapor Deposition Techniques Qian Shen1, Annie Ng1, Zhiwei Ren1, Huseyin Cem Gokkaya1, Aleksandra B. Djurišić2, Juan Antonio Zapien3 and Charles Surya*1,4 1

Department of Electronic and Information Engineering, the Hong Kong Polytechnic

University, Hong Kong SAR 2

Department of Physics, the University of Hong Kong, Pokfulam, Hong Kong SAR

3

Department of Physics and Materials Science, City University of Hong Kong, Hong Kong S.A.R. 4

School of Engineering, Nazarbayev University, 53 Kabanbay batyr ave., Astana, 010000, Kazakhstan

*

Corresponding author e-mail: [email protected]

Keywords: low-frequency noise, perovskite solar cells, defect passivation, stability, chemical vapor deposition

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Abstract: In this paper, detailed investigations of low-frequency noise (LFN) characteristics of hybrid chemical vapor deposition (HCVD)- and solution-grown CH3NH3PbI3 (MAPI) solar cells are reported. It has been shown that LFN is a ubiquitous phenomenon observed in all semiconductor devices. It is the smallest signal that can be measured from the device, hence systematic characterization of the LFN properties can be utilized as a highly sensitive nondestructive tool for the characterization of material defects in the device. It has been demonstrated that the noise power spectral densities (PSDs) of the devices are critically dependent on the parameters of the fabrication process, including the growth ambient of the perovskite layer and the incorporation of the mesoscopic structures in the devices. Our experimental results indicated that the LFN arises from a thermally activated trapping and detrapping process resulting in the corresponding fluctuations in the conductance of the device. The results show that the presence of oxygen in the growth ambient of the HCVD process and the inclusion of an mp-TiO2 layer in the device structure are two important factors contributing to the substantial reduction in the density of the localized states in the MAPI devices. Furthermore, the lifetimes of the MAPI perovskite based solar cells (PSCs) are strongly dependent on the material defect concentration. The degradation process is substantially more rapid for the devices with higher initial defect density compared to the devices prepared in optimized conditions and structure that exhibit substantially lower initial trap density.


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1. Introduction The present total worldwide power consumption is estimated to be ∼18 TW which is expected to grow by as much as one-third by 2035. Continued dependence on fossil fuels for power generation will result in irreversible damages to the environment with devastating consequences. According to the International Energy Agency, nearly half of the net increase in electricity generation will come from renewables.1 Hence, the development of costeffective and highly efficient renewable energy sources will be critical for meeting the future global energy demands for the 21st Century. Among the various photovoltaic (PV) materials, organic-inorganic perovskite thin films have drawn significant attention in recent years due to its tremendous progress in the development of high efficiency solar cells. Since their first report in 2009, the power conversion efficiency (PCE) of PSCs had improved from 3.81% to a record 22.1% at the time of the compilation of this manuscript.2,3 The tremendous enhancement in the device performance is driven by the following key factors: i.) superior physical properties for the perovskite materials such as high absorption coefficients, extremely long carrier diffusion lengths4 and tunable bandgaps;5 ii.) development of highly effective material growth techniques for the deposition of high quality perovskite films, such the solution process,6 vapor assisted solution process,7 hybrid chemical vapor deposition (HCVD),8 solvent engineering,9 thermal evaporation,10,11 and the vacuum-assisted growth techniques12 … etc.; iii.) development of highly efficient device structures, for instance the planar structure, mesoporous structure3 and inverted structure13,14; and iv.) development of high-quality electron transport layer (ETL) and hole transport layer (HTL) for efficient collection and transport of photo-generated carriers.15-17 Recent theoretical and experimental studies had demonstrated the beneficial effects of defect passivation on the PCE of the PSCs. Theoretical analyses indicate that localized bandgap states in the perovskite films are highly detrimental to the device performance due to 3 ACS Paragon Plus Environment


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their impact on the recombination rates of the photo-carriers and have significant effects on the carrier mobility and diffusion length.18-20 Such defects are typically located at the grain boundaries in the perovskite layers and at the interface between the active layer and the ETL or HTL.21,22 Thus, the growth of pinhole-free perovskite films with large crystals and interface modification will be important for eliminating the defect states. It was observed by Ren et al.23 and Yin et al.24 that significant improvements in the device PCEs can be accomplished by annealing the perovskite material in dry oxygen which leads to the passivation of the material defects in the devices. A number of published theoretical works have elucidated the role of oxygen in defect passivation for the MAPI films.24,25 Material defect density is also linked to the long-term stability of the PSCs which is crucial for realizing the commercial potential of the devices. Work by Li et al.26 and Han et al.27 showed that the MAPI decomposes into CH3NH3I (MAI) and PbI2 in the presence of moisture, heat and UV radiation resulting in significant increase in the defect density of the perovskite material leading to the degradation in the devices performance. To date, there are little systematic investigations on the material defects in PSCs as a function of the process parameters. Ideally, the characterization of the trap density should be conducted directly on the devices to ascertain a positive correlation between the fabrication parameters, the trap density and the PV performance of the devices. However, most of the conventional trap density characterization techniques are either destructive and typically performed on thin films, such as the transmission electron microscopy and the scanning electron microscopy, or may cause degradations in the device characteristics as in the case of deep level transient spectroscopy. Low-frequency excess noise is a ubiquitous phenomenon and is observed in all semiconductor devices including the PSCs. During the measurement, we applied either a constant current bias or a constant voltage bias on the device under study using all-passive biasing sources. The output fluctuation signal was amplified by a low-noise amplifier and the 4 ACS Paragon Plus Environment

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corresponding PSD was measured in the ac mode. The noise PSD typically exhibits a γ functional form of S ( f ) ∝ 1/ f where γ ≈ 1. It has been demonstrated in various

semiconductor optoelectronic and electronic devices that the LFN in the device conductance originates from the random trapping and detrapping of the carriers by the localized defect states in the material. When a trap captures a carrier, its charge state will be modified resulting in the corresponding variation in the local band bending at the vicinity of the trap. This leads to a modulation in the local carrier density. Furthermore, the change in the charge state of the trap will also lead to a variation in the coulombic scattering rate and thereby resulting in the fluctuation in the carrier mobility. Hence, the trapping and detrapping processes lead to correlated fluctuations in both the carrier density and mobility. The random trapping and detrapping of the carriers by a single defect will result in the corresponding fluctuation in the conductance of the device in the form of a random telegraph noise with the noise PSD given by

=

(1)

where A is a proportionality constant related to the magnitude of the fluctuation arising from the trapping and detrapping of a single carrier and τ is the fluctuation time constant. If the traps in the device capture and emit carriers independently of each other, the overall current noise PSD, SI(f), can be expressed as

= , , ,

(2)

where NT is the trap density and I is the dc current applied to the device. Low-frequency noise had been widely studied in conventional optoelectronic devices, such as silicon solar cells28,29 and GaN light emitting diodes30 and is shown to be a versatile non-destructive tool for the characterization of device reliability. Recently, LFN characterizations have been conducted on different classes of thin film solar cells, such as polymer solar cells31-34 and dye-sensitized solar cells35 for investigating the degradation 5 ACS Paragon Plus Environment


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mechanisms of the devices. In this paper, we have utilized LFN characterization as an effective non-destructive tool for monitoring the trap density directly on the device structure. We report on the dependence of the defect state density on the processing parameters for the PSCs and their impacts on the performance of the devices. 2. Results and Discussion It is important to note that because of the complicated device structure for the PSCs, the exact origin of the noise is usually obscured due to the presence of a number of different layers and interfaces within the devices. Therefore, we have simplified the sample structure to FTO/MAPI/Au to facilitate the investigation of the underlying physical mechanism of the observed LFN from the perovskite film itself and, thereby pinpoint the impact of the perovskite growth techniques on the relative defect density of the perovskite layer. The MAPI layer in the resistive structure was grown under different experimental conditions as detailed below: i.) MAPI films grown by conventional two-step solution technique and without oxygen annealing (type A); ii.) MAPI films grown by conventional two-step solution technique with post-deposition oxygen annealing (type B); iii.) MAPI films grown by HCVD technique using pure nitrogen as the carrier gas (type C); and iv.) MAPI films grown by HCVD technique using N2/O2 (85%/15%) mixture as the carrier gas (type D). The I-V characteristics of the different types of samples in resistive structure are shown in Figure S1. Systematic characterizations of low-frequency voltage noise PSD, ! , were performed on the resistive structure over a wide range of temperatures and the obtained LFN PSDs are found to exhibit a spectral form given by ! ∝ 1/ % where γ ≈1. Typical experimental results of ! measured form the different types of samples in resistive structure at room temperature are shown in Figure 1 while the ! measured from different types of samples with systematic variation in temperatures are shown in Figure 2a. The data show that ! vary significantly with the growth conditions of the samples. It is noteworthy that a systematic dependence of γ on the temperature is observed as indicated in Figure 2b. This has significant 6 ACS Paragon Plus Environment

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implications on the underlying physical mechanism of the observed LFN. The fluctuation time constant, τ, for the Lorentzian in Equation 1 may arise from either a tunnelling or thermally activated process. For LFN dominated by the trapping and detrapping of carriers arising from the tunneling process, the fluctuation time constant is given by & = &' exp+, where + is the WKB parameter and z is the tunneling distance. The Lorentzian at an arbitrary frequency, ,, is shown to peak sharply under the condition where ,& = 1. Thus, one can

define a tunneling distance . = − exp,&' that identifies the physical location of the 0

“active” traps responsible for the observed LFN at frequency ,. Based on the tunneling model for LFN, the parameter γ is shown to depend on the spatial distribution of the localized states, NT, as summarized below for ≈ . : 2 = 1 for

678

2 > 1 for 2 < 1 for

9

= 0;

678

9

> 0; and

678

9

< 0.

6 :;

6 :;

6 :;

(3)

On the other hand, for LFN dominated by the thermally activated trapping and detrapping processes, the fluctuation time constant will be given by & = &' exp/@A B where E is the activation energy, @A is the Boltzmann constant, and T is the absolute temperature. It is shown that in this case the Lorentzian peaks sharply at . = −@A B ln,&' and where &' is the inverse phonon frequency of the order 10-12 s. Thus, & is strongly temperature dependent and γ is shown to depend on the energy distribution of NT at . as summarized below: 2 = 1 for

678

9

= 0;

6 :;

2 > 1 for

678

9

> 0; and

2 < 1 for

678

9

< 0.

6 :;

6 :;

(4) 7

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To differentiate the two processes, one observes that the WKB parameter, λ, and hence . are relatively insensitive to the temperature change for the tunnelling process whereas significant variation in . occurs arising from relatively minor changes in the temperature of the samples for the thermally activated process due to the strong temperature dependence of the thermally activated time constant. Thus, the basic mechanism that underlies the LFN process can be identified through detailed examination of the temperature dependence of γ. Experimental investigation of γ exhibits systematic temperature dependence as the temperature of the samples varied between room temperature and 90 K as shown in Figure 2b. This strongly suggests that a thermally activated process underlies the observed LFN in the perovskite material. Based on the thermal activation model for the LFN, it can be shown that G

E. F ≈ H

I

G J

=H

I

!

.

(5)

where C is a proportionality constant. It is noted that the constant C is related to the proportionality constant A in Eq. 2 and both constants are related to the magnitude of the fluctuation in the conductance of the resistive structure due to: i.) fluctuation in the local carrier concentration at the vicinity of the trap; ii.) fluctuation in the carrier mobility due to the modulation of the Coulombic scattering rate arising from the fluctuation in the density of the fixed charged states;36 and iii.) fluctuation in the current arising from the modulation of the energy barrier at various heterojunctions37 such as the TiO2/perovskite interface, metal/perovskite interface and the grain boundaries. Furthermore, the impacts of the transport mechanism on the conductance need to be taken into account as well. It is noteworthy that for an intrinsic perovskite material with low carrier concentration, space charge limited current mechanism may potentially be the dominant transport mechanism for the carriers. Despite the complicated picture underlying the quantitative evaluation of the conductance fluctuation for our samples, the exact magnitudes of the constants A and C should bear little consequence on the conclusion of our investigation as we are only interested in the relative changes in the trap 8 ACS Paragon Plus Environment

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density of the perovskite layer due to the various growth techniques. Experimental data in Figure 2 demonstrate significant reduction in the trap density due to post-deposition oxygen annealing (type B) for solution processed samples and the use of a carrier gas consisting of an N2/O2 (85%/15%) mixture (type D) in HCVD growth process. Recent studies by the authors25 showed that the composition of the carrier gas had little effect on the morphology of the films, which demonstrate similar grain size for the films grown in different compositions of the carrier gas under the same post-deposition cooling rate. It is unlikely that the reduction of the trap density arises from the surfactant effect of oxygen leading to the improved crystallization of the perovskite layer. Hence, the data strongly suggest that the reduction of the defect states density was achieved through the passivation of the localized states within the bulk of the perovskite material and at the grain boundary. Based on Equation 5, a normalized trap density, 7 , can be determined as a function of the energy and the results for the four different types of perovskite films are plotted in Figure 3. The results demonstrate significant reduction in the trap density for types B and D films compared to their counterparts (types A and C) which were not exposed to oxygen during the fabrication process. Further reduction in the trap density is seen in the type D film compared to the type B film. This is attributed to the fact that HCVD-grown films have demonstrated substantially enhanced grain size which greatly reduces the density of grain boundaries. The LFN results are consistent with recent studies on the photothermal deflection spectroscopies of solutionand HCVD-grown perovskite films.25 The variation of as a function of the activation energy has profound impact on the properties of the LFN PSD.

Since every single trap in the samples has an associated

activation energy, as indicated in Figure 3, and according to Eq. 1, each trap level will give rise to a LFN PSD in the form of a Lorentzian. If a peak exists in the trap density at an arbitrary activation energy, as in the case of Type C sample, then a “bump” may be observed in the LFN PSD (see Figure S1). The position of this “bump” will exhibit strong dependence 9 ACS Paragon Plus Environment


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on temperature due to the thermally activated nature of the fluctuation time constant, &. This phenomenon is observed in Figure 3 in which a peak in the 7 is found at energy ~ 0.35eV for the type C sample. Equation 1 stipulates the presence of a a Lorentzian bump in the noise PSD at ,&KLMH = 1, which is clearly observed in the voltage noise PSDs for the type C sample as shown in Figure S2. A strong temperature dependence for the corner frequency, ,N is observed. An Arrhenius plot of & is shown in Figure 4 with an activation energy, , of 0.37 eV which agrees well with the peak position of 7 in Figure 3. The model provides plausible explanation for the drastic changes in the values of 2 for type C sample in the temperature range between 160 K and 200 K. As ,N decreases with the reduction of the temperature and falls within the range of the measurement window of the signal analyzer, the noise PSD will be strongly influenced by the presence of the Lorenztian. This leads to the significant reduction in the value of 2 as Equation 1 stipulates that O/ O = 0 for ,P & P 1. Barone et al.38-41 had performed a number of in-depth investigations on the properties of LFN in organic and PSCs as a function of temperature. It was suggested that the trapping and detrapping phenomena, associated with the defect states, lead to fluctuations in the number of charge carriers and contribute, therefore, to a current fluctuation that can be measured by the external contacts. Under illumination the amplitude, Var [I], of the 1/f noise of the device in the solar cell structure can be defined as the sum of a dark (Var [Idark]) and a photo-induced (Var [Iph]) contribution. The quantitative trap density can be determined through numerical fitting of Var [I] as a function of the photo-current. The model was used to analyse the trap density of the perovskite as a function of the device temperature. It was demonstrated that 10 ACS Paragon Plus Environment

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significant increase in the noise magnitude can be observed at around 160 K. It is suggested that the drastic rise the noise power spectral density stipulates a corresponding increase in the trap density in the device due to a phase transition from the tetragonal phase to orthorhombic phase as the device temperature decreases from room temperature to ~160 K. This is in contrast with the LFN data obtained from our samples which indicate much more gradual changes in the noise PSDs as a function of temperature. Nevertheless, one observes increase in the noise PSDs for T < 150 K in all types of samples under investigation. It is believed that the differences in the temperature dependencies of the LFN properties reported by the two groups arise from the difference in the structure of the samples under study. In our case, the low temperature characterizations were performed on the simple resistive structure as opposed to the complete photovoltaic devices reported by Barone et al. In a heterojunction structure, the effects on the current due to the trapping and detrapping phenomena in the photovoltaic devices may be significantly amplified depending on the physical location of the trap relative to the metallurgical junction.37 Also, investigation of phase transition in perovskite thin films by Osherov et al.42 indicated a much more gradual change in the structural and optoelectronic properties of the film below 150 K. Thus, the gradual rise in the LFN observed from our samples is consistent with the gradual phase transition as observed by Osherov et al. It is noteworthy that the observation of a Lorentzian bump in type C sample is perfectly consistent with the gradual phase transition model as the data in Fig. 2a indicate just a small “bump” in the LFN PSD for the type C sample. We have also investigated the LFN properties of the PSCs fabricated under different processing conditions and device structures: i.)

Type I devices are the control consisting of planar structure and the MAPI layer was grown by solution technique without oxygen annealing process;

ii.)

Type II devices consist of planar structure and the MAPI layer was grown by solution technique with post-deposition oxygen annealing; 11 ACS Paragon Plus Environment


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iii.)

Type III devices consist of a planar device structure and the MAPI layer was grown by HCVD technique using pure nitrogen as the carrier gas;

iv.)

Type IV devices consist of a planar structure and the MAPI layer was grown by HCVD technique using a mixture of nitrogen and oxygen (N2:O2=85%:15%) as the carrier gas; and

v.)

Type V devices consist of a mesoporous structure and the MAPI layer was grown by HCVD technique using a carrier gas consisting of a mixture of nitrogen and oxygen (N2:O2=85%:15%).

The experimental room temperature /QP data for the devices are shown in Figure 5. From the data it is observed that the type I device exhibits the highest LFN PSD and is close to two orders of magnitude higher than the type II device. This clearly demonstrates substantial reduction in the trap density for type II PSC arising from the low-temperature post-deposition oxygen annealing. Type II and type III devices demonstrate roughly the same noise level despite using pure nitrogen as the carrier gas for the growth of the perovskite layer by HCVD technique in the type III device and hence no trap passivation process was performed. The low /Q P for the type III device is attributed to the reduction in the trap density due to significant enhancement in the grain size grown by the HCVD process as shown in Figure S3. With the use of the N2/O2 (85%/15%) carrier gas in the HCVD process, further reduction in /QP can be achieved for type IV device which demonstrates, once again, the effect of trap passivation due to exposure of the MAPI layer to dry oxygen. It is interesting to note that further reduction in the LFN is observed with the inclusion of a mesoporous layer in the device (type V). This is contrary to common expectation as the addition of the mesoporous layer greatly increases the interfacial area between the perovskite and TiO2. This would normally result in the increase in the interface states at the perovskite/TiO2 interface. The type V device exhibits a significant reduction in the LFN, suggesting that the perovskite/TiO2 interface is relatively benign. Furthermore, studies by Ng 12 ACS Paragon Plus Environment

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et al.25 on the crystallinity of the HCVD-grown perovskite films, the XRD technique demonstrated that perovskites deposited on mp-TiO2 substrates exhibit improvement in the crystallinity compared to the samples without the mp-TiO2 scaffold. The results are, thus, consistent with the observed reduction in the LFN PSD for the type V device compared to the type IV PSC. Our experimental results on LFN characterizations are consistent with the observed enhancements of the I-V characteristics of the devices as shown in Figure 6 and the averaged photovoltaic parameters are summarized in Table 1. Significant enhancements in the PCEs and FFs are observed among the oxygen passivated devices (types II, IV and V) compared to their unpassivated counterparts (types I and III). It is noteworthy that significant enhancement in the FF for type V device compared to the type IV device is due mainly to the improvement in extraction efficiency of photo-carriers by the mesoporous TiO2 structures. Thus, the lowering in the LFN observed in type V device may also partly be attributed to the efficient collection of the photo-generated carriers and thereby diminishing the probability of the carriers being captured by the localized states in the perovskite layer in addition to the enhancement in the crystallinity of the perovskite film. We have characterized the electrochemical impedance spectroscopies (EIS) on the five different types of devices and the fitting results are shown in Figure S4a-e. The values of the parameters calculated from the EIS data are summarized in Table S1. Based on the experimental data we have computed the selective contact resistances (Rsc) and the recombination resistances (Rrec) of the devices utilizing the equivalent circuit model43 as shown in Figure S4f. The obtained results for Rsc and Rrec are shown in Figure 7 in which Rs in the circuit model is the series resistance due to the electrodes while CPE-sc and CPE-rec are the capacitances with frequency dispersion related to Rsc and Rrec respectively. The CPE is used in the equivalent circuit model instead of the ideal capacitance as the spatial inhomogeneities induced by the defects at the interfaces are considered.43,44 In Figure 7a the 13 ACS Paragon Plus Environment


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computed values for Rsc are shown as a function of the voltage bias. The Rsc contributes to total series resistance43 and affects the FF of the solar cells. Substantial reductions in the Rsc are observed among the oxygen passivated devices particularly for types IV and V devices. This is attributed to the enhancement in the crystallinity of the perovskite films (types IV and V) as well as the enhancement in the collection efficiency (type V) of the device due to the incorporation of the TiO2 mesoporous structure. In Figure 7b, we observe significant increase in the Rrec among the oxygen passivated devices compared to the unpassivated counterparts. This is in excellent agreement with the experimental results on the LFN characterizations presented above which indicated significant reduction in the trap density due to oxygen passivation of the defect states and thereby suppressing the carrier recombination rate in the devices. The carrier lifetimes as a function of the biasing voltages determined from the results of EIS are summarized in the Table S1 and Figure S5. The obtained lifetimes of the devices in this work are in the same range as the reported lifetimes of the MAPI based solar cells using the same characterization method.44 The order of 10-3 to 10-2 s for the lifetime is determined for the optimized device (Type V) compared to the order of 10-4 to 10-3 s for the devices without the incorporation of oxygen and TiO2 mesoporous structure (Type I and Type III), indicating that passivation of trap states by oxygen leads to enhancements in both the electrical transport and the lifetimes of the carriers in the devices. Furthermore, enhanced carrier collection efficiency can be achieved by incorporation of mesoscopic TiO2 layer. The extracted lifetimes from EIS results for different types of devices are consistent with the trend in the relative defect densities of the perovskite films extrapolated from the noise spectra. This shows that the carrier recombination dynamics in the devices are significantly affected by the trap density in the perovskite materials. This is consistent with the work of Landi et al.41 who comprehensively discussed the recombination kinetics due to the trap states in the perovskite devices. The stability of the device is an important issue that must be addressed to realize the 14 ACS Paragon Plus Environment

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potential of PSCs. Since material defects had often been linked to the degradation of optoelectronic devices, it is important to investigate how the HCVD and the oxygen passivation processes may affect the lifetimes of the PSCs. Figure 8 illustrates the degradation in the PCE of the devices as a function of the shelf time. The data show that the devices with oxygen passivation are significantly more stable than those without. In particular, the type I device exhibited very rapid degradation such that the measurement was stopped after ~1500 hrs of shelf time. The results clearly indicate the benefits of oxygen annealing in enhancing the device lifetime. The incorporation of a mesoporous layer further improves the stability of solar cells. This is attributed to the efficient collection of the photo-generated carriers by the TiO2 mesoporous structure thereby significantly reducing the impact of the material defects on the operation of the device. From the experimental results presented in Table 1, the significant degradation in the PCE of type III device is associated with the decrease in VOC, JSC and FF for shelf time 0 ≤ S ≤ 1000 hrs. In contrast, the oxygen passivated devices (types II, IV and V) demonstrate much more gradual degradations in device properties. Furthermore, it is found that the film quality of the perovskite materials grown by different techniques also has strong impact on the degradation rate of the devices. It is noteworthy that the devices without oxygen annealing (Type I and Type III) exhibit much faster burn off rates as shown in Figure 8. It is believed that the perovskite films without oxygen annealing contain higher density of unpassivated defect states, which likely lead to the accumulation of photoexcitied carriers within the devices. These unpassivated defect states not only degrade the photovoltaic performance of the devices but also cause instability in the devices as the accumulated charges will readily react with the external species from the environment as well as the additives in the HTM.45-48 One can roughly differentiate the time degradation in the PCEs of the devices into two regimes. Regime I (0 < t ≤1000 hr) is characterized by high rate of degradation in the PCE 15 ACS Paragon Plus Environment


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and is ascribed as the burn-in phase of the device. In regime II (t >1000 hr), the degradation of the PCE is significantly slower. It is interesting to note that the rates for PCE degradations of the devices are closely correlated with the degradation in the normalized current noise PSD, /Q P , of the devices as indicated in Figure 9. The rapid degradation of the device performance during the burn-in phase is associated with dramatic increase in the current noise PSD, /Q P while the rate of increment of /Q P is substantially reduced in regime II. The high rates of degradation in the PCEs in regime I is, thus, attributed to the rapid generation of traps in the perovskite films. Based on the experimental results it is proposed that the high initial rates of degradation arise from the breaking of weak bonds in the material. When the weak bonds are eliminated, the rate of generation of traps will be substantially reduced resulting in the corresponding slow degradation in the PCE as shown in regime II in both Figures 8 and 9. These results clearly indicate that oxygen passivation results in significant improvement in the stability of the PSCs due to reduction in the trap density and the enhancement in the robustness of the device. This phenomenon is analogous to the other solar cells, such as the organic solar cells and amorphous silicon solar cells, for which several hundreds of hours to several weeks are typically observed as the burn-in period for these types of solar cells respectively.49,50. The burn-in degradation mechanism for these conventional solar cells is also attributed to the rapid formation of trap states due to the photo-induced reactions in the organic active layers or in the hydrogenated amorphous silicon. Similar to the PSCs, the generations of new trap states lead to the rapid initial degradation in performance during first 1000 hours which is followed by a phase with much slower degradation, indicating that the weak bonds or the reactive species have been depleted after the burn-in period.


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3. Conclusion Low-frequency noise characterizations were performed on PSCs to study the impact of defect states on the optoelectronic properties and the stability for both solution- and HCVDbased devices. Lower defect density is determined for the perovskite films with higher crystallinity prepared by HCVD technique. Oxygen annealing not only passivates the defect states of MAPI films grown by solution and HCVD techniques but also improves the stability of the devices. The mesoporous TiO2 layer in the device is beneficial in enhancing the extraction efficiency of photo-carriers and hence improve the stability of solar cell due to the reduction in the probability of the carriers being capture by the traps.

4. Experimental Section Device fabrication. The structure of the device is shown in Figure S6. A patterned glass/FTO substrate was first ultrasonically cleaned by detergent solution followed by rinsing in DI water, acetone and isopropanol. The ETL was deposited on patterned FTO-coated glass substrate, which consists of a compact TiO2 (c-TiO2) thin film on the FTO with a sheet resistance of 7-10Ω/□. The c-TiO2 layer was grown by spin coating a solution of titanium isopropoxide on the substrate followed by sintering at 450°C for 2 hours. For the mesoscopic devices, a mesoporous TiO2 (mp-TiO2) layer was deposited by spin coating a diluted solution of 18NRT paste in ethanol followed by a 450°C calcinating process for 2 hours to remove the organic components. The perovskite film was grown on top of the ETL either by all-solution process or HCVD technique. For both techniques, lead iodide, dissolved in DMF (462mg/ml), was first spin coated on the substrates at 1250 rpm and the perovskite films were formed after the reaction of PbI2 thin films with MAI in the solution phase or vapor phase. For solution process, the perovskite films were prepared in accordance with our previous work.23 For HCVD process, 17 ACS Paragon Plus Environment


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the PbI2 samples were transferred into a quartz tube furnace in which the MAI powder was maintained at 180°C to facilitate its sublimation and subsequent transport to the substrate located downstream using a carrier gas. It was shown that the optimal carrier gas composition consists of a nitrogen/oxygen (85%/15%) mixture. The sample was maintained at a constant substrate temperature for the crystallization of the MAPI thin films. The optimal growth temperature was shown, in a recent work by the authors, to be ∼165°C.25 Subsequent to the growth process, the substrate was allowed to cool down to room temperature at a rate of 0.7°C/min. Using such a slow cooling rate was shown to be crucial in improving the uniformity, reducing the density of pinholes and enhancing the grain size of the MAPI film.25 The samples were then taken out from the quartz tube and the HTL, consisting of a solution of spiro-MeOTAD (80mg/ml) in chlorobenzene with additives of Li-TFSI (17.5 µL from a stock solution of 520 mg/ml in acetonitrile) and 29 µL of tBP, was spin-coated on the perovskite layer at 3500 rpm. Finally, an Au electrode was deposited by thermal evaporation via a shadow mask with a cell active area of 0.06 cm2. Device and sample characterization. The device was illuminated by an Oriel Sol3A solar simulator equipped with and AM1.5 filter. The light intensity was calibrated carefully to 100mW/cm2 in an area of 2 inch × 2 inch prior to measurement. The I-V characteristics, both under illumination and in the dark, were measured by an Agilent B1500A semiconductor device analyzer. For LFN measurement the device was enclosed in a shielded room to eliminate extraneous noise. The device was powered by a battery-powered circuit. During the measurement, we applied either a constant current bias or a constant dc voltage bias on the device with a resistive structure or a solar cell structure respectively. The output fluctuation signal was amplified by a low-noise amplifier and the corresponding PSD was measured in the ac mode. The biasing condition is indicated in the figure caption. A fast Fourier transform (FFT) was performed on the amplified signal by an HP3651A dynamic signal analyzer to 18 ACS Paragon Plus Environment

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obtain the PSD of the noise. At least 500 FFT traces were averaged for each measurement of PSD. A

cryostat cooled with liquid nitrogen and a Lakeshore DRC 91C temperature

controller was used to stablize the temperature of device within 40 mK. The EIS measurement was performed in the frequency range from 1 Hz to 1MHz by using a Zahner elektrik model IM6 in dark under different voltage bias.

Supporting Information Available I-V characteristics of different types of the resistive films; The voltage noise PSDs for the type C structure; SEM images; Impedance spectra of different types of devices and device configuration. These materials are available free of charge via the Internet at http://pubs.acs.org.

Acknowledgements This work was supported by GRF grants (Grant No. PolyU 152045/15E), (Grant No. PolyU 152468/16E), RGC Theme-based Research Scheme (Grant number: HKU T23-713/11), a Poly U Central Allocation Grant (1ZVGH) and the Clarea Au Endowed Professorship.

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Figure 1 Typical room temperature voltage noise power spectral density, ! , for the resistive structures fabricated with different types of perovskite films. A constant current bias was applied to obtain a voltage of 0.5 V on all devices using an all-resistive current source.



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Figure 2 Temperature dependencies of: (a) ! measured by applying a constant current bias to obtain a voltage of 0.5 V on all devices with different types of perovskite films in the resistive structure; and (b) frequency exponent, 2, of the LFN PSD of the various resistive structures.


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Figure 3 Normalized trap density, 7 E. F, computed for the different types of resistive structures as a function of energy.



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Figure 4 The Arrhenius plot of the fluctuation time constant, &, indicating an activation energy of 0.37 eV.


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Figure 5 Room temperature /Q P for the different types of PSCs. A constant dc voltage bias of 0.8 V was applied on all types of the devices during the measurement.



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Figure 6 I-V characteristics for the 5 different types of PSCs.


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Figure 7 (a) Selective contact resistance, Rsc, for the different types of PSCs; and (b) recombination resistances, Rrec, for the different types of PSCs.



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Figure 8 Degradation of the PCEs of the different types of PSCs as a function of shelf time.


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Figure 9 Degradation of the /QP measured from the different types of PSCs as a function of the shelf time. A constant dc voltage bias of 0.8 V was applied on all types of the devices during the measurement.



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Table 1 Average photovoltaic properties of the 5 types of PSCs measured at different shelf times.

Type I

Type II

Type III

Type IV

Type V

Hours

VOC (V)

JSC (mA/cm2)

FF

0 1104 0 1104 3696 6912 0 1104 3696 6912 0 1104 3696 6912 0 1104 3696 6912

0.98±0.01 0.76±0.17 1.02±0 0.93±0.05 0.93±0.05 0.93±0.02 0.98±0 0.67±0.24 0.38±0.46 0.29±0.22 1.01±0.02 0.95±0.02 0.91±0.02 0.92±0.06 0.97±0.02 0.91±0.02 0.87±0.02 0.89±0.02

18.9±0.5 8.6±0.7 23.2±1.2 17.4±1.2 15.9±1.7 16.1±1.5 23.3±1.1 16.9±2.0 7.1±8.2 1.8±0.5 24.2±0.6 22.4±1.2 20.9±2.7 22.4±2.0 23.9±1.3 21.8±1.5 21.9±1.6 22.2±1.5

0.53±0.01 0.25±0.02 0.61±0.06 0.59±0.03 0.51±0.09 0.48±0.03 0.65±0.07 0.29±0.03 0.16±0.14 0.19±0.15 0.69±0.03 0.52±0.12 0.43±0.20 0.44±0.13 0.73±0.02 0.63±0.01 0.62±0.05 0.59±0.01

PCE (%) 9.9±0.2 1.6±0.3 14.4±1.5 9.4±0.3 7.4±1.0 7.1±0.8 14.8±1.1 3.3±1.2 1.1±1.8 0.1±0.1 16.9±1.0 11.1±2.5 8.6±4.8 9.4±3.9 16.9±0.9 12.6±0.7 11.9±0.3 11.5±0.5

Lifetime (s) 1.10E-04

1.81E-04

2.62E-04

6.93E-04

1.34E-03

-


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Characterization of Low-Frequency Excess Noise in CH3NH3PbI3

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