Solution-Processed Lithium-Doped ZnO Electron Transport Layer for

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Solution Processed Lithium Doped ZnO Electron Transport Layer for Efficient Triple Cation (Rb, MA, FA) Perovskite Solar Cells Md Arafat Mahmud, Naveen Kumar Elumalai, Mushfika Baishakhi Upama, Dian Wang, Arman Mahboubi Soufiani, Matthew Wright, Cheng Xu, Faiazul Haque, and Ashraf Uddin ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.7b09153 • Publication Date (Web): 14 Sep 2017 Downloaded from http://pubs.acs.org on September 15, 2017

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ACS Applied Materials & Interfaces

Solution Processed Lithium Doped ZnO Electron Transport Layer for Efficient Triple Cation (Rb, MA, FA) Perovskite Solar Cells Md Arafat Mahmud, Naveen Kumar Elumalai*, Mushfika Baishakhi Upama, Dian Wang, Arman Mahboubi Soufiani, Matthew Wright, Cheng Xu, Faiazul Haque and Ashraf Uddin* School of Photovoltaic and Renewable Energy Engineering, University of New South Wales, Sydney, NSW 2052, Australia

*Corresponding authors: NKE ([email protected]) & AU ([email protected]) Abstract: The current work reports the Lithium (Li) doping of low-temperature processed zinc oxide (ZnO) electron transport layer (ETL) for highly efficient, triple-cation based MA0.57FA0.38Rb0.05PbI3 (MA: methyl ammonium, FA: formamidinium, Rb: Rubidium) perovskite solar cells (PSCs). Lithium intercalation in the host ZnO lattice structure is dominated by interstitial doping phenomena, which passivates the intrinsic defects in ZnO film. In addition, interstitial Li doping also downshifts the Fermi energy position of L-ZnO ETL by 30 meV, which contributes to the reduction of electron injection barrier from photoactive perovskite layer. Compared to the pristine ZnO, the power conversion efficiency (PCE) of the PSCs incorporating Lithium doped ZnO (L-ZnO) is raised from 14.07%

to 16.14%. The superior

performance is attributed to the reduced current leakage, enhanced charge extraction characteristics and mitigated trap-assisted recombination phenomena in L-ZnO devices, thoroughly investigated by means of electrochemical impedance spectroscopy (EIS) analysis. L-ZnO PSCs also exhibit lower photo-current hysteresis than ZnO devices, which is investigated with regard to electrode polarization phenomena of the fabricated devices.

Keywords: intersitital Li doping, trap state passivation, ZnO ETL, perovskite solar cell, electron injection barrier, electrode polarization

ACS Paragon Plus Environment


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1. Introduction: Low-temperature processed perovskite solar cells (PSCs) have emerged as promising candidate in the photovoltaic industry for large-scale commercial production with flexible substrates via roll-to-roll process. 1-6 In particular, PSCs, incorporating solution processed ZnO thin films as electron transport layer (ETL) offer a facile, cost-effective device fabrication pathway. 6-11 However, since ZnO exhibits relatively open, hexagonal close-packed lattice structure, where native Zn atoms occupy only half of the tetrahedral sites, all its octahedral sites remain empty. 12 The unoccupied octahedral ZnO sites can act as defect states within ZnO bandgap, even lying as deep as 0.2 eV below the conduction band.

13

The presence of such

deep trap sites in ZnO ETL causes the light-generated electrons from perovskite conduction band to hop amid these trap sites,

14-15

before they finally get extracted by the electrode. This electron hopping

phenomena in ZnO film can result in trap-mediated carrier recombination at perovskite/ZnO interface, 14, 16-17

which is a major process limitation of such low temperature processed ETL. To overcome this short-

coming of ZnO ETL, trap-state passivation of ZnO film has been reported by introducing a secondary organic ETL layer [PCBM ([6,6]-phenyl C61 butyric acid methyl ester)]

16

between ZnO and perovskite

layer as a mean of interfacial engineering or by ZnO surface modification with metal oxide doping 11, 17-18 in Burstein-Moss mechanism. 19 In particular, ZnO surface modification with extrinsic dopant 11, 17-18 is an interesting approach, since secondary organic layer like PCBM can either aggregate or diffuse during the perovskite annealing process,

20

which effectively eliminates the benefit associated with the trap-state

passivation of ZnO ETL. In this regard, it is intriguing to explore the surface modification of ZnO ETL along with extrinsic lithium (Li) dopant, which (Li dopant) has been reported to reduce defect states of ZnO film, applied either as gate contact in thin film transistor (TFT) 21 or as immobilized photocatalyst. 22 However, the so far reported ZnO deposition and related Li doping methods are performed through chemical vapour deposition (CVD) 21 or spray pyrolysis techniques, 22 which necessiate high-temperature processing and are not suitable for flexible substrates. 23 To date, no study has been undertaken to investigate the compatibility of extrinsic Li doping in low temperature, solution processed ZnO ETL for PSCs, although Li doping has been reported with other conventional ETLs (TiO2 or SnO2), either in mesoporous

24-25

or planar 0

3, 26

PSC device structures. In

particular, extrinsic Li dopant incorporation in low temperature (185 C) processed SnO2 ETL

3

has

demonstrated that it does not necessitate high-temperature processing to ensure the intercalation of Li ions into metal oxide lattice structure, as reported in earlier literatures.

24-26

However, the reported Li treated

SnO2 ETL 3 still falls short to meet up the stringent requirement of sub-150 0C temperature processing 23 to make the overall process compatible with flexible substrates. In this respect, implementation of Li treated ZnO ETL for perovskite device is an intriguing prospect, since ZnO ETL can be processed below


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150 0C temperature, 14, 17 which can meet up the temperature prerequisite for flexible substrate. Adding to the merits, ZnO ETL also exhibits an order of higher thin film conductivity (0.85 mS/cm) 27-28 compared to that of SnO2 ETL (0.04 mS/cm), 3 which (conductivity of ZnO) is expected to enhance with extrinsic Li doping due to increased carrier concentration.

29

Thus, Li treated ZnO ETL bids fair to be an interesting

prospect as a possible replacement for pristine ZnO ETL due to its enhanced conductive property 27-28 and conceivable beneficial role in trap state passivation of pristine ZnO film in Burstein-Moss effect. Besides, it is also intriguing to investigate the charge transport property,

31-32

30

inherent electrode

polarization and concomitant photo-current hysteresis characteristics 33 and device aging phenomena 34 of such low temperature processed devices. In this study, we have presented extrinsic Li doping of sol-gel ZnO ETL via a low-temperature (140 0C) route for fabricating highly efficient [highest power conversion efficiency (PCE): 16.14%] triple cation based MA0.57FA0.38Rb0.05PbI3 (MA: methyl ammonium, FA: formamidinium, Rb: Rubidium) PSC. Lithium intercalation in host ZnO lattice structure has been found to be dominated by interstitial doping phenomena, which passivate the intrinsic defects in ZnO ETL film. The photovoltaic performance and photo-current hysteresis phenomena in pristine and Li treated ZnO ETL based PSCs have been systematically investigated and compared. Li doped ZnO device exhibits about 16% higher average PCE than conventional ZnO ETL based devices. Adding to the merits, Li doped ZnO device also exhibits suppressed photo-current hysteresis phenomena compared to pristine ZnO based PSCs. In our work, enhanced device performance and suppressed hysteresis phenomena with Li doped ZnO devices have been investigated via several material and device characterization techniques including dopant intercalation, microstrain and dislocation density analysis from X-Ray diffraction (XRD) spectral fitting, material workfunction and electron injection barrier determination with ultraviolet photoelectron spectroscopy

(UPS),

non-radiative

carrier

recombination

investigation

with

steady

state

photoluminescence (PL) measurement, as well as charge transport and trap-assisted recombination analysis by electrochemical impedance spectroscopy (EIS) analysis. 2. Experimental Section 2.1. Fabrication Process At first, patterned ITO/glass substrates were washed with cleaning concentret (Hellmanex III), deionized (DI) water, acetone and propan-2-ol with successive cycles of 10 minutes. Sol-gel ZnO ETL was deposited on cleaned ITO/glass substrate in an identical maner, as described in our previous works.

14, 35

For Li treated ZnO ETL, 0.1 M bis(trifluoromethane)sulfonimide lithium (Li-TFSI) salt in acetonitrile solvent was spin-casted over the ZnO film at a spin-rate of 4000 rpm for 30 s and the film was dried at 140 0C for 30 minutes. For further referencing in our manuscript, the Li treated ZnO ETL will be named



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as L-ZnO ETL. For both the ZnO and L-ZnO ETL devices, the precursor solution for triple cation based perovskite (MA0.57FA0.38 Rb0.05PbI3) was prepared by mixing PbI2, MAI, FAI and RbI powder with appropriate molar ratio in N, N–dimethylformamide (DMF) solvent. The prepared perovskite precursor solution was vigorously stirred for 24 hours at 70 0C temperature on a hotplate and then filtered with a 0.45 µm polytetrafluoroethylene (PTFE) filter before spin-coating. The perovskite solution was spincasted at 3000 rpm in gas-assisted method using a N2 spray gun. Then, the substrates were annealed at 100 0C temperature for 10 minutes by restricted volume solvent annealing (RVSA) process thermal decomposition of perovskite on basic ZnO surface. precursor

solution

(73.3

mg/ml)

of

37

36

to inhibit

For the hole transport layer (HTL), a

2,2',7,7'-Tetrakis[N,N-di(4-methoxyphenyl)amino]-9,9'-

spirobifluorene (Spiro-OMeTAD) in chlorobenzene was prepared and it was doped with 28.8 µL 4-tertbutylpyridine (4-TBP) and 17.5 µL Li-TFSI (520 mg/ml in acetronitrile) to increase its p-type conductivity. The HTL precursor solution was spin-coated at a spin-rate of 3000 rpm for 30s. Afterwards, a 100 nm thick Ag layer was deposited by the thermal evaporation process (evaporation rate: 2 Å /s, vacuum level: 1x10-6 mbar). During the evaporation, a metal, shadow mask was used to fix the device area to be 4.5 mm2. 2.2. Characterization Process The photovoltaic peformance of the fabricated devices were determined with a Keithley 2400 source meter under AM 1.5G simulated light, calibrated with a national renewable energy laboratory (NREL)certified reference solar cell. During the measurement, a black, non-reflective aperture mask (having the same aperture area as the device active area) was placed over the devices to ensure there is no overestimation in the device short-circuit current density. XPS characterization was carried out using ESCALAB250Xi instrument (Thermo Scientific, UK) with a background pressure of 2x10-9 mbar using a mono-chromated AlKα (hυ = 1486.68 eV) anode (120 W, 13.8kV, 8.7 mA). UPS measurement was conducted in ESCALAB250Xi instrument (Thermo Scientific, UK) by the application of He I photon line (hυ = 21.22 eV) from a He discharge lamp under a background pressure of 4x10-10 mbar. X-ray diffraction (XRD) characterization was performed with CuKα radiation using PANalytical Empyrean Thin-Film XRD machine. The related step-size in the step-scanning process of XRD was 0.02 degree. The optical characterization of ETL films was performed with a UV–vis-NIR spectrometer (Perkin Elmer–Lambda 950). The surface morphology imaging was done with Carl Zeiss AURIGA SEM and the surface topogoly was analyzed with Bruker Dimension ICON SPM AFM apparatus. Steady state PL measurement was performed using a 1/4 meter monochromator (CornerstoneTM 260). A 409 nm, 60 mW continuous wave (CW) laser, along with an optical density 2 (OD2) filter acted as an excitation source for this machine. The PL was traced by a built-in silicon charge-coupled device (CCD) camera (AndorTM). A low pass filter (442 nm) was utilized to filter out the laser lines from the detected signal. The integration


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time for the PL spectra was 2 s. An Autolab PGSTAT-30 impedance analyzer was used to conduct the EIS characterization and Mott-Schottky analysis. A low oscillating amplitude [root mean square (RMS) value: 20 mV] of alternating current (AC) was chosen during the analysis to maintain a linear response. The frequency range was varied from 1 Hz to 106 Hz. 3. Results and Discussion 3.1 Material and Optical Characterization of ETL Films Deposited L-ZnO film has been investigated with XPS elemental analysis to ensure the incorporation of Li atoms in spin-coating process for the surface modification of pristine ZnO film. Fig. 1(A) shows the high resolution XPS spectra of Li1s for ZnO and L-ZnO, respectively. As observed from Fig. 1(A), Li1s spectra of L-ZnO film exhibit a distinct peak at 55.45 eV binding energy, which ensures the presence of Li atoms in it. 25 As expected, no such peak was observed for Li1s spectra of pristine ZnO film (Fig. 1(A)). To determine the atomic percentage of individual elements, we can refer to the XPS survey spectra of ZnO and L-ZnO ETL, presented in Fig. S1(A) and S1(B), respectively. Corresponding peak binding energies, peak width and elemental atomic percentage for ZnO and L-ZnO ETL films have been listed in Table S1 and Table S2, respectively. As observed from Table S2, the atomic percentage of Li1s in L-ZnO film is 11.26%. As observed from the XPS survey spectra (Fig. S1(A)), L-ZnO film does not demonstrate any traceable amount of sulphur or fluorine atom from Li-TFSI precursor solution, which is consistent with its air-annealed fabrication method, inducing Li-TFSI oxidation. 25-26 To investigate the intercalation of Li atom 25, 38 in ZnO host lattice structure, we have inspected the deconvoluted Gaussian subpeaks of O1s XPS spectra for ZnO and L-ZnO ETL films [Fig. 1(B) and Fig. 1(C)]. Both ZnO and L-ZnO ETL films exhibit a main (O1s A) and a shoulder peak (O1s B), pertaining to O1s XPS spectra. However, the peak locations of O1s A and O1s B shift to higher binding energy for L-ZnO film, compared to ZnO film (Fig. 1(B)-1(C) and Table S1-S2). The area ratio of O1s A to O1s B peak [area ratio, S1:S2, as marked in Fig. 1(B) and 1(C)] is relatively low (S1:S2=1.324) for L-ZnO compared to ZnO film (S1:S2=1.311), which indicates that L-ZnO exhibits more pronounced O1s B shoulder peak 25 than pristine ZnO film. The more pronounced shoulder peak at higher binding energy in L-ZnO film ascertains the interaction of Li with oxygen atoms

25

and thus the intercalation of Li

38

into ZnO lattice structure even with sub 150 0C

temperature processing. To determine whether the Li intercalation phenomena in L-ZnO film are dominated by interstitial or substitutional doping

29, 39

of ZnO host lattice structure, we have performed the XRD characterization of

both ZnO and L-ZnO films, as presented in Fig. S2. From the XRD data, both (002) and (101) characteristic ZnO diffraction peaks 40 are visible in ZnO and L-ZnO films. No secondary Li phase could



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be traced from the XRD pattern of L-ZnO film, denoting the crystalline components of L-ZnO film consist of ZnO phase alone. 22, 39 However, the crystallinity and the crystallite dimension of the two films are different, as hinted by the variation in relative normalized peak intensity and peak width (full width at half-maximum, FWHM) of the respective films from the XRD peak analysis (Table S3). To reasonably compare the XRD spectral fitting data of ZnO and L-ZnO films, the experiment setup and the film thickness of the two films were kept similar, so that the corresponding volume fractions considered for the two films remain identical during XRD characterization. Fig. 1(D) depicts the difference in film crystallinity (from relative normalized peak intensity) and crystallite dimension of ZnO and L-ZnO films, pertaining to (002) and (101) diffraction peaks from Table S3. From Fig. 1(D), L-ZnO film demonstrates relatively high crystallinity and larger crystallite dimension compared to ZnO film, pertaining to major (002) wurtzite diffraction peak. Scanning electron microscopy (SEM) images of the two ETL films (Fig. S3) also reveal that the grain size of L-ZnO film is higher than ZnO film, which conforms to its (L-ZnO film’s) higher individual crystallite dimension from XRD spectral fitting. The spinel shapes visible in SEM image of L-ZnO film also hint towards the formation of Li2ZnO2 compound in it, which contributes to its enhanced crystallinity

21-22

compared to pristine ZnO film. From atomic force microscopy (AFM)

measurement of the two ETL films (Fig. S4), the average and the root mean square (RMS) surface roughness of L-ZnO film both increase (Table S4) with the enlargement in its crystallite dimension. Rough surface topography attained with L-ZnO film ensures enhanced development of grain structure, 21 which is consistent with its relatively enhanced crystallinity compared to ZnO film. The increased crystallite dimension with L-ZnO film from XRD, SEM and AFM characterization is an important observation. In general, Li ions (Li+) exhibit smaller ionic radius (68 pm) than Zn ions (Zn2+) (74 pm).

21

Thus, the increase in crystallite size of L-ZnO film denotes that the intercalation phenomena

of Li are dominated by the interstitial placement of Li+ ions in ZnO lattice sites rather than the substitutional exchange of Li with native Zn atoms, which (substitutional exchange) would otherwise lower the crystallite dimension of L-ZnO.

21, 29

The interstitial doping of Li in L-ZnO film can be

explained with a schematic, presented in Scheme-1. Scheme-1(A) illustrates the ZnO lattice structure, showing the intrinsic defect sites (interstitials or vacancies). Li intercalation in ZnO host lattice structure can take place by interstitial doping (Scheme-1(B)) or by substitutional exchange with Zn atom (Scheme1(C)) by the following Frenkel’s reactions: 41 = +

(1)

= +ℎ

(2)


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Where, , , and ℎ refer to Li+ ions on intersitital ZnO site (Scheme-1(B)), Li atom on Zn lattice site in the substitutional position (Scheme-1(C)), electron and hole, respectively. In general, the energy of interstitial doping of Li+ ion (Eq. 1) is much lower than the substitutional exchange of Li atom in Zn lattice matrix (Eq. 2). 42 Thus, Li intercalation in L-ZnO is dominated by interstitial doping phenomenon, which can passivate the interstitial trap sites. Besides, interstitial doping increases the electron carrier concentration of L-ZnO film according to Eq. 1, which is consistent with the enhanced conductivity with lithium doped ZnO film, as reported in earlier literatures. 27-28 In this work, we have also investigated the effect of Li intercalation with regard to oxygen vacancy filling in ZnO. From the deconvoluted O1s XPS spectra of ZnO and L-ZnO films (Fig. 1(B) and 1(C)), the integrated area ratio, S =

for both the films are nearly identical (ZnO: 56.97%, L-ZnO:

56.78%). Identical S values for ZnO and L-ZnO films indicate that chemical state of oxygen shows nearly same stoichiometry for the two films by filling the oxygen vacancy This is also consistent with the identical Zn3s peaks

26

43

during the air-annealing process.

of the two ETL films from the XPS survey curves

(Fig. S1), which indicate that the oxygen vacancy filling phenomena by the partial reduction 25 of Zn2+ ion to Zn+ ion are almost similar in ZnO and L-ZnO films. Therefore, intrinsic defect passivation in L-ZnO film is dominated by interstitial doping with Li+ ions in host ZnO lattice structure. We have probed into the effect of interstitial doping on the microstrain 44 and dislocation density 45 of LZnO ETL film. Fig. 2(A) illustrates how the microstrain and the dislocation density values for ZnO and L-ZnO film vary, with regard to major XRD diffraction peaks. Corresponding numeric values are listed in Table S3. From Fig. 2(A), microstrain value is lower for L-ZnO film, related to (002) diffraction peak. In general, microstrain provides a quantitative measure of the defect states of the nano-crystalline structure and is mathematically expressed as: 44

ԑ=

(⁄)

(3)

where, ԑ, and denote refer to microstrain, Bragg diffraction angle and full-width at half maximum

(FWHM) of the characteristic XRD peak at 2 diffraction angle, respectively. Thus, the lower microstrain value with L-ZnO film is a signatory of interstitial defect or trap state passivation in it (LZnO)

22

by the intercalation of Li+ ions, which is also consistent with enhanced crystallinity of L-ZnO

film. From Fig. 2(A), L-ZnO film also demonstrates lower dislocation density pertaining to (002) diffraction peak. Usually, dislocation density refers to the crystalline imperfection, specially at the grain boundaries and mathematically can be expressed with the following formula: 45

=

!


(4)


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here, d and n stand for crystallite size from XRD spectral fitting and a unity constant for the lowest possible crystalline dislocation, respectively. Thus, the lower dislocation density value in L-ZnO film also implicates that the Li doping reduces the dislocations or cracks at the grain boundaries of pristine ZnO film, which is also visible from the SEM images of the respective films, presented in Fig. S3. We have also investigated the transmittance patterns of the deposited ETL films [Fig. 2(B)] on ITO/glass substrate. In a normal structured (n-i-p) perovskite device,

14, 17

the ETL is film is expected to be highly

transmitting so that overlying perovskite film can absorb maximum sunlight. From Fig. 2(B), the average visible (370-740 nm) transmittance (AVT)

46-47

value of L-ZnO ETL film is slightly higher than that of

ZnO (ZnO: 86.82%, L-ZnO: 87.19%) which demonstrates its (L-ZnO film’s) excellent optical compatibility as ETL for a normal structured PSC. The optical bandgap of L-ZnO film is slightly higher compared to pristine ZnO film (ZnO: 3.31 eV, L-ZnO: 3.32 eV), as conceived from the respective Tauc plots [Fig. S5(A) and Fig. S5(B)]. The slight blue-shift of the bandgap in L-ZnO film can be explained with Burstein-Moss effect. 19 In Burstein-Moss phenomena, an extra energy is needed to excite additional valence band electron (from interstitial Li+ ions) to higher energy states in the conduction band, according to Pauli’s principle, a doubly occupied state is impossible

49

48

since,

and thus the optical bandgap of L-

ZnO film is widened. Hence, the widening of optical bandgap in L-ZnO film also indicates higher conductivity of L-ZnO film due to higher electron carrier concentration from interstitial Li+ ions. 29 To examine the effect of interstitial Li doping on the material workfunction of L-ZnO film, we have conducted UPS measurement of ZnO and L-ZnO films atop ITO/glass substrate. Fig. 2(C) depicts the evolution of secondary electron edges for ZnO and L-ZnO ETL films on ITO/glass substrate from UPS characterization. From Fig. 2(C), the workfunctions of ZnO and L-ZnO on ITO/glass substrate are 3.79 eV and 3.82 eV, respectively. The downshift in Fermi level of L-ZnO film is due to the presence of interstitial Li+ ions, acting as acceptor states and is consistent with previous literatures. 3, 43 The conducive role of Fermi level shift in L-ZnO film has been explained in the subsequent subsection in terms of the modulation of the electron injection barrier 48 from photo-active perovskite layer. 3.2 Material and Optical Characterization of Perovskite Film MA0.57FA0.38Rb0.05PbI3 perovskite films were fabricated in identical manner on both the ETL films. Fig. 3(A) depicts the XRD patterns of perovskite films on both ETL films. For both the perovskite films, we can observe (110), (202), (220), (310), (224) and (314) characteristic perovskite XRD peaks.

50-51

Fig.

3(B) presents the relative normalized peak intensity and crystallite dimension of the perovskite films from XRD spectral fitting. Corresponding numerical values are listed in Table S5 for quantatitive comparison. As observed from Fig. 3(B) and Table S5, the material crystallinity and crystallite dimension of the


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perovskite film atop L-ZnO film is relative high than perovskite on ZnO film, pertaining to all the characteric peaks except (220). To probe further into it, we have also calculated the values of microstrain and dislocation density for the perovskite films on top of ZnO and L-ZnO films (Fig. 3(C)). From Fig. 3(C), MA0.57FA0.38Rb0.05PbI3 perovskite overlying L-ZnO ETL film demonstrates lower microstrain and dislocation density for all the major diffraction peaks except for (220) peak. The observed variation in the bulk property of identical perovskite layer on two dissimilar ETL or HTL film has also been reported in previous literautres.

52-53

However, to investigate the role of underneath ETL layers on the bulk property

of overlying MA0.57FA0.38Rb0.05PbI3 perovskite films, we have conducted steady state PL measurement of perovskite/ETL bilayers on a glass substrate (Fig. 3(D)). The PL measurement was conducted from both the perovskite and the glass/ETL sides. As comprehended from Fig. 3(D), perovskite on L-ZnO bilayer film demonstrates higher PL intensity than perovskite/ZnO bilayer, irrespective of the excitation side. Higher PL intensity of perovskite/L-ZnO bilayer when illuminated from the perovskite side denotes the bulk property of MA0.57FA0.38Rb0.05PbI3 perovskite is enhanced 54 on top of L-ZnO film, compared to the compositionally similar (perovskite) layer deposited on ZnO film. The enhanced bulk charge-carrier recombination property of perovskite/L-ZnO film from (perovskite side) PL measurement is also consistent with higher crystallinity, lower microstrain and lower dislocation density values of the same bilayer film obtained from XRD spectral fitting. From Fig. 3(D), perovskite/L-ZnO also exhibits higher PL intensity than perovskite/ZnO bilayer film when excited from the glass/ETL side, which indicates that the perovskite/ETL interfacial charge transfer property is also superior in perovskite/L-ZnO bilayer film due to reduced non-radiative carrier recombination phenomena.

54

This potential reduction in non-

radiative carrier recombination phenomena in perovskite/L-ZnO bilayer can be attributed to the interstitial trap-state passivation in L-ZnO film by the intercalation of Li+ ions in ZnO lattice structure. We have also analyzed the surface morphology and topography of the fabricated perovskite films on top of ZnO and L-ZnO ETL films. Fig. 4(A)-4(B) and 4(C)-4(D) present the surface morphology images of MA0.57FA0.38Rb0.05PbI3 perovskite films on top of ZnO and L-ZnO ETL films, respectively. MA0.57FA0.38Rb0.05PbI3 perovskite atop L-ZnO film exhibits slightly higher grain size compared to the same (perovskite) on ZnO, which is consistent with the higher individual crystallite dimension of the former film from XRD spectral fitting. Relatively large grain growth of perovskite film on L-ZnO film is correlated with the crack-free, compact surface morphology of the underlying L-ZnO film (Fig. S3), which ensures higher grain boundary mobility for homogeneous perovskite nuclei distribution 55 on it (LZnO), leading to enalargement of perovskite grain. Conversely, the crack-induced grain boundaries of ZnO film (Fig. S3) leads to perovskite nucleation at its grain boundary cavities, which limits the grain boundary mobility of perovskite grain, 55 resulting in relatively small grain distribution of perovskite. Fig. 4(E)-4(F) and 4(G)-4(H) show the two and three dimensional (2D and 3D) AFM images of perovskite



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films on top of ZnO and L-ZnO ETL films, respectively. The average and RMS surface roughness of the fabricated perovskite films are listed in Table S6. From Table S6, the surface roughness of perovskite on L-ZnO film becomes slightly higher compared to perovskite/ZnO film along with the increase in grain size. 3.3 Device Performance We have fabricated PSCs to comprehend the performance of ZnO and L-ZnO films as electron selective layers. The full device structure is: ITO/ZnO or L-ZnO/MA0.57FA0.38Rb0.05PbI3 perovskite/SpiroOMeTAD/Ag (Fig. 5(A)). For further referencing in our work, we will term the two PSCs as ZnO device and L-ZnO device. Table 1 presents the average and the best photovoltaic performance of the ZnO and LZnO PSCs at forward bias to short circuit (FB-SC) and short circuit to forward bias (SC-FB) scan directions. As comprehend from Table 1, L-ZnO devices exhibit nearly 16% higher average efficiency (ZnO PSC: 13.68%, L-ZnO PSC: 15.82%) compared to ZnO devices. The reproducibility data of both the PSCs have been presented with statistical box-charts in supporting information (Fig. S6). The corresponding numeric values of J-V parameters have also been listed in Table S7 and Table S8, respectively. The J-V curves of the champion ZnO and L-ZnO devices have been shown in Fig. 5(B) and 5(C), respectively. L-ZnO device demonstrates a highest PCE of 16.14% at FB-SC direction, whereas, the corresponding value for ZnO device is 14.07%. Fig. 5(D) and 5(E) present the stabilized PCE and stabilized current density of the champion ZnO and L-ZnO PSCs at maximum power point (MPP)

56

along with time. The corresponding numeric values are tabulated in Table S9. From Table S9, L-ZnO device shows stabilized PCE and stabilized current density (at MPP) values of 14.82% and of 19.56 mA/cm2, respectively, whereas, the corresponding values for ZnO device are about 12.87% and 19.07 mA/cm2, respectively. Here, it is noteworthy that the stabilized PCE (at MPP) values are about 91%-92% of the efficiency value attained in FB-SC scan direction, which conforms to the observation of previous perovskite studies. 57-58 The enhanced photovoltaic performance with L-ZnO devices can be attributed to a number of aspects identified from our earlier characterizations. About 16% higher average PCE in L-ZnO devices come along with about 5.2% higher average VOC (ZnO: 966.47 mV, L-ZnO: 1015.98 mV) values, about 1.5% higher average JSC (ZnO: 21.78 mA/cm2, L-ZnO: 22.11 mA/cm2) values and most prominently about 8.4% higher fill factor (ZnO: 65%, L-ZnO: 70.43%) values compared to ZnO devices. The enhancement in average VOC values with L-ZnO devices can be correlated with four possible mechanisms. Firstly, the intercalation of Li+ ions passivates the interstitial trap states of ZnO lattice, which reduces the non-radiative carrier recombination at perovskite/ETL interface in L-ZnO device, as


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observed from PL measurement (Fig. 3(D)) presented in our work. Reduced non-radiative carrier recombination contributes to higher VOC values

54

in L-ZnO devices compared to ZnO PSCs. Secondly,

the bulk property of MA0.57FA0.38Rb0.05PbI3 perovskite is enhanced on top of L-ZnO film, as indicated by the higher crystallinity, lower microstrain and dislocation density values of perovskite/L-ZnO bilayer film from XRD spectral fitting and PL measurement. Enhanced perovskite bulk property indicates reduced intrinsic defect density in it, recombination phenomena

59

59

which can lower the trap-assisted Shockley-Read-Hall (SRH)

in L-ZnO device. In general, SRH recombination rate is expressed as:

"# = $

$% $& '( ) * +! ,

(5)

% ( - )$& (**- )

Where, . , .* , 01 , 2, 3, 2 , 2 425 3 refer to capture co-efficients of electrons and holes by trap levels, density of trap states, intrinsic charge carrier, electron and hole density, equilibrium density of electron and holes, respectively. Since, the trap density, electron and hole capture co-efficients for perovskite/LZnO are lower due to higher perovskite crystallinity, the SRH recombination phenomena are anticipated to be less in L-ZnO device according to Eq. 5. Lower SRH recombination ensures higher average VOC values

59

with L-ZnO device. Thirdly, enhanced VOC values in L-ZnO devices can also be attributed to

crack-free, compact surface morphology of L-ZnO film, which blocks the shunting path

60

between the

overlying perovskite layer and ITO electrode and indicates towards comparatively lower leakage current 61

in L-ZnO device. To investigate the leakage current phenomena in the fabricated devices, we have

conducted the dark J-V measurement of both ZnO and L-ZnO devices. Fig. 5(F) shows the dark J-V curves of the two device under study. As comprehended from Fig. 5(F), L-ZnO device exhibits lower leakage current compared to ZnO PSC under reverse bias. Hence, the reverse dark saturation current density (J0) in L-ZnO device is relatively lower than that in ZnO PSC. In general, the VOC of a solar cell is mathematically expressed as: 62

67$ =

89 :

=

ln =>? @

(6)

here, k, T and q stand for Boltzmann constant, absolute temperature in Kelvin scale and elementary charge, respectively. As evident from Eq. 6, lower J0 and higher JSC value (from Table 1) in L-ZnO device lead to a rise in average VOC values in it compared to ZnO device. Lastly, perovskite film on top of LZnO film demonstrates relatively high grain size than the same (perovskite) on top of ZnO ETL films (Fig. 4). It has been reported that the donor-accetor pair (DAP) distance in large grain-sized perovskite is larger with reference to that in a smaller grain-sized perovskite film. 63 Moreover, the acceptor and donor binding energies have also been reported to reduce with the increase in grain size of perovskite film. Generally, DAP recombination-mediated photon energy is mathematically expressed as: 63


63


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Page 12 of 30

:!

ℏB = CD − (CF + CH ) + IJ

(7)

@ JK L

where, CD , CF , CH , MN , MO and R denote material bandgap, binding energies of acceptor and donor, free space permittivity, relative dielectric constant and the ionized DAP inter-distance, respectively. Since, for perovskite on L-ZnO film, the DAP inter-distance is expected to be higher and the acceptor and donor binding energies are lower due to larger grain size, the DAP recombination photon energy for perovskite/L-ZnO film should be lower from Eq. 7. Thus, we hypothesize the lower DAP recombination in L-ZnO device contributes to the enhancement in VOC values.

63

perovskite grain has also been reported in earlier perovskite literatures

55

VOC enhancement with larger due to reduced recombination,

which is consistent with the observation of our current work. The enhancement in JSC for L-ZnO device can be ascribed to the downshift in Fermi level position (from UPS characterization) of L-ZnO ETL film and concomitant influence on the electron injection barrier 16, 48 from photo-active perovskite layer. Efficient electron transfer from perovskite conduction band (~3.95 eV)

3

to ITO/ETL cathode depends on the electron injection barrier or energetic offset between the

perovskite conduction band and the Fermi position of ITO/ETL cathode. 3, 16 For ZnO device, the electron injection barrier is 0.16 eV, while the corresponding value (0.13 eV) for L-ZnO device is 30 meV lower than ZnO device, due to the downshift of the Fermi position in L-ZnO Thus, the lower electron injection barrier in CA ETL provides more efficient electron extraction capability 3 and higher JSC values in L-ZnO device, compared to ZnO device. The most significant boost in average PCE of L-ZnO devices originates form their higher average FF value compared to ZnO devices. As observed from Table 1, L-ZnO devices demonstrate about 14% lower average series resistance, RS value (ZnO: 7.25 Ω.cm2, L-ZnO: 6.23 Ω.cm2) and nearly 5.5% higher average shunt resistance, RSh value (ZnO: 1082 Ω.cm2, L-ZnO: 1140 Ω.cm2) compared to ZnO devices. Generally, FF of a solar cell can be mathematically expressed by the following empirical formula: 62

PP = PPN Q 1 −

.L> L?T

L!

+ U.L>! V Q1 − ?T

PPN =

WXY N.Z [[@ L?T WXY

L>\

] 1 −

.L> L?T

L!

+ U.L>! ^V ?T

WXY _ (WXY N.Z) WXY

heree, " , "$` , abc and "d denote series resistance, the characteristic resistance "$` = e

(8)

(9) ef? , =>?

normalized open circuit voltage (abc = ef? , 69 and n refer to thermal voltage and ideality factor of the g

cell) and shunt resistance of the solar cell, respectively. As comprehended from Eq. 8, reduced RS and


Page 13 of 30

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enhanced RSh value in L-ZnO devices lead to augmented FF value in them than ZnO devices. Reduction of series resistance, RS in L-ZnO device is related with enhanced crystallinity of perovskite film 64 atop LZnO ETL film, as observed from XRD spectral fitting and PL measurement. Lower series resistance in LZnO device is also connected with the increased conductivity of L-ZnO film due to enhanced carrier concentration via interstitial Li doping mechanism, which is also evident by the higher dark injection current (dark forward current) 64 for an applied bias greater than 0.8V in the forward bias region 65 for LZnO device compared to ZnO device (Fig. 5(F)). Reduction in series resistance value for L-ZnO device is also consistent with an earlier study 66 which reports that the reduced number of grain boundaries in large grain sized perovskite (perovskite on L-ZnO ETL in our case) lowers the series resistance of a PSC owing to lower inter-particle contact resistance. 66 On the other hand, enhanced shunt resistance value in L-ZnO device is coherent with its reduced leakage current,

64

compared to ZnO device from the dark current-

voltage measurement (Fig. 5(F)). The increase in average shunt resistance of L-ZnO devices also implicates towards mitigated carrier recombination phenomena

67

in them. In this regard, we have

calculated the flat-band potentials (VFB) of ZnO and L-ZnO PSCs from Mott-Schottky analysis, which (VFB) enable us to probe into the trap-mediated charge-carrier recombination phenomena of the devices. 67 Fig. 6(A) and 6(B) present the Mott Schottky curves of ZnO and L-ZnO devices, respectively at a frequency of 10 KHz in dark condition. As observed from Fig. 6(A) and 6(B), L-ZnO device demonstrates relatively high flat-band potential (ZnO device: 0.93V, L-ZnO device: 1.04V) compared to ZnO device. High flat-band potential in L-ZnO device indicates lower trap-assisted recombination phenomena

14, 67

in it due to the intrinsic trap state passivation of L-ZnO ETL by interstitial Li doping.

Suppressed recombination phenomena in L-ZnO device result in enhanced shunt resistance, FF and PCE value in it compared to ZnO PSC. 3.4 Charge Transport Characteristics To discern the charge transport characteristics of ZnO and L-ZnO PSCs, their EIS analysis was conducted, which (EIS) can reveal detailed information regarding the charge extraction and carrier recombination mechanism of the PSCs. 67 Fig. 6(C) shows the Nyquist curves of ZnO and L-ZnO devices at an applied bias of 950 mV in dark condition. A suitable electronic circuit model

8, 14

(the inset of Fig.

6(C)) was used to fit the experimental data from EIS analysis. The electronic circuit model includes a series resistance, RSE in conjunction with 2 parallel R-C elements. From the equivalent circuit in Fig. 6(C), RSE, RC, RRec, Cµ and CC stand for the wire and metal contact resistance, contact resistance at the interfaces between perovskite and neighbouring charge selective contacts, recombination resistance and capacitance and bulk capacitance, respectively. Table 2 tabulates the calculated values of RSE, RC, RRec, Cµ and CC by the iterative convergence method, pertaining to the Nyquist curve, presented at Fig. 6(C).



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Page 14 of 30

As comprehended from Table 2, the interfacial contact resistance, RC in L-ZnO device is nearly 21.2% lower compared to ZnO devices (ZnO device: 33.12 Ω.cm2, L-ZnO device: 26.10 Ω.cm2). Lower RC value in L-ZnO device confirms the superior charge transport property

14, 68

from MA0.57FA0.38Rb0.05PbI3

perovskite to L-ZnO ETL, which is consistent with reduced electron injection barrier and higher perovskite crystallinity in L-ZnO device. Reduced contact resistance (RC) of L-ZnO device from EIS characterzation also conforms to its smaller series resistance (RS) value 68 from light J-V curve and lower leakage

64

and higher dark injection current

65

from dark J-V curve. Apart from lower contact resistance,

L-ZnO PSC also demonstrates about 38.5% higher recombination resistance, RRec (ZnO device: 64.35 Ω.cm2, L-ZnO device: 89.10 Ω.cm2) than ZnO device, As, the external power supply, alone, acts as the source of charge carriers under dark EIS analysis, the higher recombination resistance for L-ZnO PSC is correlated with the suppressed charge recombination phenomena at the trap or dislocation states

30, 68

located in the perovskite-bulk or at the interfaces between perovskite and L-ZnO. Higher recombination resistance value (RRec) of L-ZnO device from EIS analysis also conforms to its larger shunt resistance (RSh) from J-V measurement.

68

To investigate more into the the recombination phenomena, we have

charcterized the capacitive spectra (Fig. 6(D)) of ZnO and L-ZnO PSCs. Usually, the intermediate frequency (1-10 kHz frequency) response of a PSC is attributed to the trap-mediated charge accumulation phenomena at the interfaces between perovskite and neighbouring charge selective layers, whereas, while the high frequency (>10 kHz) capacitance is correlated with the bulk perovskite properties. 31, 33 From Fig. 6(D), L-FDT PSC demonstrates lower capacitance values at the intermediate region compared to ZnO device. Lower intermediate frequency capacitance reaffirms that the trap-mediated recombination process at the interfacial areas between perovskite and ETL/HTL films are more suppressed

31

in L-ZnO device.

This is also congruent with the higher flat-band potential of L-ZnO PSC from the Mott-Schottky analysis. Moreover, the high frequency capacitive responses of ZnO and L-ZnO devices (the inset of Fig. 6(D)) are slightly different, which conforms to the variation in bulk MA0.57FA0.38Rb0.05PbI3 perovskite property 31, 33 in two devices, although the perovskite fabrication techniques were identical in both cases. 3.5 Hysteresis Behaviour From Fig. 5(B) and Fig. 5(C), it is apparent that both the ZnO and L-ZnO devices demonstrate hysteresis behaviour to some extent. For quantitative comparision of their photocurrent-hysteresis characteristics, we have calculated the correponding hysteresis index (HI)

7, 56

values from the following mathematical

expression: hi =

=jkl>? )ebcm,=>?ljk )ebcm, =jkl>? )ebcm,


(10)

Page 15 of 30

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where, n[o$ )6pqm2, and n$[o )6pqm2, refer to the values of current density at a bias of 6pqm2 for FB-SC and SC-FB scan directions, respectively. The corresponding values are tabulated in Table S10. As observed from Table S10, L-ZnO PSC exhibits lower HI values (CC device: 0.29, CA device: 0.03) compared to ZnO devices, indicating quantitatively that the hysteresis phenomena is less

7

for L-ZnO

devices. To probe into the dissimilarity in photo-current hysteresis of the two devices, we can re-analyze the capacitive responses of the fabricated devices (Fig. 6(D)). In general, the low frequency region of the capacitive spectra reveals significant information regarding the electrode polarization process 33 in PSCs, that governs the photo-current hysteresis phenomena in them. For PSCs, electrode polarization is associated with the slow kinetics of migrating ions 33 to the external electrodes, which modulate the local electric field of PSC.

33

Electrode polarization is demonstrated by large capacitance values in PSCs,

especially at low frequency region in the capacitive spectra.

33

The increased low frequency capacitance

modulates the time taken by a PSC for attaining the steady-state in each scan direction, resuting in hysteresis characteristics. 33 As observed from Fig. 6(D), L-ZnO PSCs demonstrate smaller low-frequency (0.1-1 Hz) capacitance compared to ZnO devices. Relatively small low-frequency capacitive response in L-ZnO PSC indicate suppressed electrode polarization process in them and explain lower photo-current hysteresis phenomena 33 with reference to ZnO PSCs. 3.6 Device Stability A systematic, 3 weeks-long degradation study of the fabricated PSCs was conducted to comprehend the cell stability.34 Un-encapsulated ZnO and L-ZnO devices were preserved in a N2 filled glovebox following the standard rules, reported in previous perovskite stability experiments in dark condition. 66, 69 The photovoltaic characterization was conducted regularly at room temperature while maintaining a relative humidity (RH) level of 35%-40%. Fig. S7(A), S7(B), S7(C) and S7(D) illustrate the normalized PCE, VOC, JSC and FF values of ZnO and L-ZnO devices, respectively from the day-0 up to the 21st day. Day-wise J-V parameters (PCE, VOC, JSC, and FF) of ZnO and L-ZnO devices are listed in Table S11 and Table S12, respectively. From Fig. S7, both ZnO and L-ZnO PSCs demonstrate comparable normalized device stability (ZnO PSC: 86.96%, L-ZnO: 87.26%) at the end of degradation study. Slightly high normalized device stability with L-ZnO PSC comes along with relatively high normalized JSC value (ZnO PSC: 93.20%, L-ZnO PSC: 98.41%) in it, compared to ZnO device. To comprehend the higher normalized JSC value in L-ZnO PSC, we have investigated the content of hydroxide groups on the surface of two ETL films using XPS elemental analysis; since, the hydroxide groups on ZnO film have been reported to provoke thermal-induced perovskite degradation.

37

Fig. S8(A) and Fig. S8(B) show the

deconvoluted Gaussian subpeaks and envelop curves of high resolution C1s XPS spectra for ZnO and LZnO ETL, respectively. Corresponding peak binding energies, peak center, associated chemical bond and



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Page 16 of 30

elemental atomic percentage for the two ETL films have been listed in Table S13. From Fig. S8, for both the ETL films, the C1s B and C1s D peaks are associated with hydroxyl/phenol group (C-OH) and carboxylate carbon (O=C-OH), respectively,70 which correspond to the presence of hydroxide group on the ETL surfaces. As observed from Table S13, the contents of hydroxyl (ZnO: 0.88%, L-ZnO: 0.80%) and carboxylate carbon (ZnO: 0.88%, L-ZnO: 0.82%) group are relatively low in L-ZnO ETL film, compared to ZnO film, which contributes to the relatively high normalized JSC value and slightly high normalized PCE value in L-ZnO PSC. However, since L-ZnO device exhibits higher initial PCE, it also demonstrates higher PCE than ZnO PSCs after 3 weeks, as observed from Table S11 and Table S12. 4. Conclusion In conclusion, we have reported Li treated ZnO (L-ZnO) ETL for fabricating highly efficient triple cation based MA0.57FA0.38Rb0.05PbI3 PSCs in a low temperature process. Li intercalation in host ZnO lattice structure is dominated by interstitial doping phenomena, which passivate the intrinsic defects in ZnO ETL. Interstitial Li doping also downshifts the Fermi position of L-ZnO ETL by 30 meV, which contributes to the reduction of electron injection barrier from photo-active perovskite layer. Perovskite film grown on L-ZnO film demonstrates enhanced crystallinity, reduced microstrain and dislocations and higher grain size compared to the same (perovskite) deposited on pristine ZnO film. The enhanced bulk perovskite property in L-ZnO device has been found to substantially suppress the non-radiative chargecarrier recombination in it (L-ZnO device), as compared to ZnO device. Furthermore, L-ZnO device also demonstrates mitigated current leakage phenomena and superior charge extraction property at perovskite/ETL interface due to reduced trap-assisted recombination phenomena. Overall, L-ZnO PSCs demonstrate about 16% higer average PCE than ZnO devices. Adding to the merits, L-ZnO PSC shows lower photo-current hysteresis than ZnO devices, which has been found to be correlated with reduced electrode polarization process. Thus, the reported L-ZnO ETL bids fair to be a promising electron selective contact for large-scale production of efficient, stable and flexible substrate based perovskite solar cells, well-suited to the roll-to-roll device processing and device scalability. 5. Associated Content: Supporting Information X-Ray photoelectron spectroscopy data for ZnO and L-ZnO ETL film, XRD spectral fitting data for ZnO and L-ZnO ETL film, surface roughness data of ETL films from AFM measurement, XRD spectral fitting data for perovskite films on ETLs, J-V parameters for ZnO and L-ZnO PSCs, stabilized current density and PCE of ZnO and L-ZnO PSCs, hysteresis indices of the two PSCs, day-wise J-V parameters of ZnO


Page 17 of 30

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and L-ZnO PSCs from degradation study, XPS survey spectra and XRD pattern of ZnO and L-ZnO films, SEM and AFM image of two ETL films, Tauc plots for material bandgap determination, histogram showing the variation in J-V parameters of ZnO and L-ZnO PSCs, day-wise normalized J-V parameters from degradation study and high resolution C1s XPS spectra of ZnO and L-ZnO ETL films. 6. Acknowledgement: The authors would like to express their sincere gratitude to Future Solar Technologies Pty. Ltd. for the financial assistant provided for this study. The authors would also like to gratefully acknowledge the never-ending supports from the staffs of Photovoltaic and Renewable Energy Engineering School (SPREE), Electron Microscope Unit (EMU) and Solid State and Elemental Analysis Unit (SSEAU) under Mark Wainwright Analytical Center, University of New South Wales (UNSW). 7. References: (1) Tan, H.; Jain, A.; Voznyy, O.; Lan, X.; García de Arquer, F. P.; Fan, J. Z.; Quintero-Bermudez, R.; Yuan, M.; Zhang, B.; Zhao, Y.; Fan, F.; Li, P.; Quan, L. N.; Zhao, Y.; Lu, Z.-H.; Yang, Z.; Hoogland, S.; Sargent, E. H. Efficient and Stable Solution-Processed Planar Perovskite Solar Cells via Contact Passivation. Science 2017, 355, 722-726. (2) Li, Y.; Meng, L.; Yang, Y.; Xu, G.; Hong, Z.; Chen, Q.; You, J.; Li, G.; Yang, Y.; Li, Y. High-Efficiency Robust Perovskite Solar Cells on Ultrathin Flexible Substrates. Nat. Commun. 2016, 7, 10214. (3) Park, M.; Kim, J.-Y.; Son, H. J.; Lee, C.-H.; Jang, S. S.; Ko, M. J. Low-Temperature Solution-Processed Li-Doped SnO2 as an Effective Electron Transporting Layer for High-Performance Flexible and Wearable Perovskite Solar Cells. Nano Energy 2016, 26, 208-215. (4) Yin, X.; Chen, P.; Que, M.; Xing, Y.; Que, W.; Niu, C.; Shao, J. Highly Efficient Flexible Perovskite Solar Cells using Solution-Derived NiOx Hole Contacts. ACS Nano 2016, 10, 3630-3636. (5) Shin, S. S.; Yang, W. S.; Noh, J. H.; Suk, J. H.; Jeon, N. J.; Park, J. H.; Kim, J. S.; Seong, W. M.; Seok, S. I., High-Performance Flexible Perovskite Solar Cells Exploiting Zn2SnO4 Prepared in Solution Below 100[Thinsp][Deg]C. Nat. Commun. 2015, 6, 7410. (6) Liu, D.; Kelly, T. L. Perovskite Solar Cells with a Planar Heterojunction Structure Prepared Using Room-Temperature Solution Processing Techniques. Nat. Photonics 2014, 8, 133-138. (7) Ginting, R. T.; Jung, E.-S.; Jeon, M.-K.; Jin, W.-Y.; Song, M.; Kang, J.-W. Low-Temperature Operation of Perovskite Solar Cells: with Efficiency Improvement and Hysteresis-Less. Nano Energy 2016, 27, 569576. (8) Liu, D.; Yang, J.; Kelly, T. L. Compact Layer Free Perovskite Solar Cells with 13.5% Efficiency. J. Am. Chem. Soc. 2014, 136, 17116-17122. (9) Mahmud, M. A.; Elumalai, N. K.; Upama, M. B.; Wang, D.; Haque, F.; Wright, M.; Xu, C.; Uddin, A. Controlled Nucleation Assisted Restricted Volume Solvent Annealing for Stable Perovskite Solar Cells. Sol. Energy Mater. Sol. Cells 2017, 167, 70-86. (10) Mahmud, M. A.; Elumalai, N. K.; Upama, M. B.; Wang, D.; Puthen-Veettil, B.; Haque, F.; Wright, M.; Xu, C.; Pivrikas, A.; Uddin, A. Controlled Ostwald Ripening Mediated Grain Growth for Smooth Perovskite Morphology and Enhanced Device Performance. Sol. Energy Mater. Sol. Cells 2017, 167, 87101.



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FIGURES AND TABLES

Table 1. Average and best photovoltaic performance with ITO/ZnO or L-ZnO/MA0.57FA0.38Rb0.05PbI3 perovskite/Spiro-OMeTAD/Ag device structures (ZnO or L-ZnO devices) at both FB-SC (forward bias to short circuit) and SC-FB (short circuit to forward bias) directions at a scan rate 0.05 V/s. The average values have been presented with corresponding standard deviation values of ten samples from a random batch Device

ZnO

L-ZnO

Average/ Best

Open Circuit Voltage, VOC (mV)

Short Circuit Current Density, JSC (mA/cm2)

Fill Factor, FF (%)

Efficiency (%)

Series Resistance, RS (Ω.cm2)

Shunt Resistance RSh (Ω.cm2)

Average

966.47 ±7.38

21.78 ±0.23

65 ±0.47

13.68 ±0.24

7.25 ±0.25

1082±63

Best (FB-SC)

977.50

22.11

65.08

14.07

7.24

1089

Best (SC-FB)

892.18

21.24

58.83

11.15

12.38

1737

Average

1015.98±8.13

22.11 ±0.10

70.43 ±0.32

15.82 ±0.18

6.23 ±0.08

1140±33

Best (FB-SC)

1026.26

22.22

70.77

16.14

6.30

1140

Best (SC-FB)

940.94

22.04

63.38

13.15

7.79

954

Table 2. Fitted values of different electronic parameters from Nyquist plots of ZnO and L-ZnO PSCs (Fig. 6(C)) at a bias of 950 mV under dark Device

RSE (Ω.cm2)

RC (Ω.cm2)

RRec (Ω.cm2)

CC (nF/cm2)

Cµ (nF/cm2)

ZnO

2.33

33.12

64.35

20.22

105.56

L-ZnO

2.41

26.10

89.10

18.89

132.67


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Fig. 1 High resolution XPS spectra of Li1s for ZnO and L-ZnO ETL, Deconvoluted Gaussian subpeaks of O1s XPS spectra for (B) ZnO and (C) L-ZnO ETL, (D) Relative normalized peak intensity and crystallite size of ZnO and L-ZnO ETL films from the XRD spectral peak analysis



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Scheme-1: Fig. (A) shows wurtzite ZnO lattice structure, demonstrating the intrinsic defects like oxygen vacancy or zinc interstitial sites, which can act as charge trapping sites during the operation of perovskite solar cell. Fig. (B) demonstrates the interstitial doping phenomena by the intercalation of Li+ ions. In interstitial doping mechanism, Li+ ions passivate the oxygen vacancies or zinc interstitial sites. Since, foreign ions are incorporated in host ZnO lattice structure, crystallite dimension increases in interstitial doping mechanism. Fig. (C) depicts the substitutional doping process by Li intercalation in ZnO. In this process, zinc atoms in the ZnO lattice matrix are substituted by Li+ ions. Since, Li has lower ionic radius than Zn, such substitution reduces the crystallite dimension of ZnO lattice


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Fig. 2 (A) Microstrain and dislocation density of ZnO and L-ZnO ETL films from the XRD spectral fitting, (B) Transmittance curves of ZnO and L-ZnO films on top of ITO/glass substrate, (C) Evolution of secondary electron edge for pristine ZnO and L-ZnO films on ITO/glass substrate from ultraviolet photoelectron spectroscopy (UPS) measurement



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Fig. 3 (A) XRD patterns of MA0.57FA0.38Rb0.05PbI3 perovskite films fabricated on top of ZnO and L-ZnO ETL films showing the major diffraction peaks and the corresponding crystal orientation (The * signs denote the peaks for ITO/glass substrate), (B) Relative normalized peak intensity and crystallite size of MA0.57FA0.38Rb0.05PbI3 films on top of ZnO and L-ZnO ETL films from the XRD spectral peak analysis, (C) Microstrain and dislocation density of perovskite/ZnO and perovskite/L-ZnO ETL films from the XRD spectral fitting, (D) Steady state photoluminescence (PL) spectra of perovskite/ZnO and perovskite/L-ZnO bilayer films on glass substrate from both perovskite and glass/ETL sides. PL measurement was done on nine different spots of the substrates to ensure consistency of the measurement


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Fig. 4 Top view Scanning Electron Microscopy (SEM) images of MA0.57FA0.38Rb0.05PbI3 perovskite on top of (A)(B) ZnO and (C)-(D) L-ZnO ETL films Two and three dimensional atomic force microscopy (AFM) images of MA0.57FA0.38Rb0.05PbI3 perovskite on top of (E)-(F) ZnO and (G)-(H) L-ZnO ETL films



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Fig. 5 (A) General schematic representation of fabricated PSCs (ZnO device and L-ZnO device), J-V curves (both FB-SC and SC-FB direction at 0.05 V/s) of the best performing (B) ZnO device and (C) L-ZnO device, Stabilized current density and stabilized PCE as a function of time for the best performing (D) ZnO device and (E) L-ZnO device and (F) dark J-V of ZnO and L-ZnO PSCs


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Fig. 6 Mott Schottky curves of (A) ZnO device and (B) L-ZnO device at 10 KHz frequency under dark, (C) Nyquist plots of ZnO and L-ZnO devices at 950 mV bias under dark (The inset shows the equivalent electronic circuit used to fit the experimental data to extract various electronic parameters of the device), (D) Capacitive response of the fabricated PSCs as a function of frequency at zero external bias under dark (the inset shows the amplified view of the high frequency region)



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Solution-Processed Lithium-Doped ZnO Electron Transport Layer for

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