Influence of Cl incorporation in perovskite precursor on crystal growth

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Influence of Cl incorporation in perovskite precursor on crystal growth and storage-stability of perovskite solar cells Hui Zhang, Yifan Lv, Jinpei Wang, Huili Ma, Zhengyi Sun, and Wei Huang ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.8b19390 • Publication Date (Web): 17 Jan 2019 Downloaded from http://pubs.acs.org on January 17, 2019

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ACS Applied Materials & Interfaces

Influence of Cl incorporation in perovskite precursor on crystal growth and storage-stability of perovskite solar cells

Hui Zhang,†, Yifan Lv†, Jinpei Wang†, Huili Ma†, Zhengyi Sun†, Wei Huang †,‡

Dr. H. Zhang, Y. Lv, J. Wang, Dr. H. Ma, Dr. Z. Sun, Prof. W. Huang † Key

Laboratory of Flexible Electronics (KLOFE) and Institute of Advanced

Materials (IAM), Jiangsu National Synergetic Innovation Center for Advanced Materials (SICAM), Nanjing Tech University (NanjingTech), 5 Xinmofan Road, Nanjing 210009, P.R. China

Prof. W. Huang ‡ Shaanxi

Institute of Flexible Electronics (SIFE), Northwestern Polytechnical

University (NPU), 127 West Youyi Road, Xi’an 710072, P.R. China

Corresponding:

Hui Zhang (E-mail: [email protected])

Keywords: perovskite solar cells; compositional engineering; grain boundaries; halide doping; non-radiative recombination

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Abstract Solar cells based on organic-inorganic hybrid lead halide perovskites are very promising for their high performance and solution process feasibility. Elemental engineering on perovskite composition is a facile path to obtain high quality crystal for efficient and stable solar cells. It was found that partially substituting the I- with Cl- in the perovskite precursor promoted crystal growth with grain size larger than layer thickness, and facilitated the generation of a self-passivation layer of PbI2. While the residual Cl- ions were suspected to diffuse to the hole transport layer consisting of ubiquitously spiro-OMeTAD, the formation of highly bounded ionic pairing of Cl- with oxidized state of spiro-OMeTAD led to insufficient charge extraction and severely reversible performance degradation. This issue was effectively alleviated upon Brdoping owing to the generation of Pb-Br bonds in the lattice that strengthened the phase stability by improving the binding energy between each unit. The binary halide (Br/Cl-) doped perovskites resulted in a champion PCE of 20.2% with improved long-term storage stability.


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Introduction Organic-inorganic hybrid lead halide perovskites are one of the most promising semiconductors owing to their excellent optoelectronic properties, facile fabrication process and broad application potentials. Since their first attempt as light harvesting materials in photovoltaic devices,1 solar cells based on perovskites have achieved remarkable progress and the certified record power conversion efficiency has exceeded 23% in just a few years.2 However, state-of-the-art perovskite solar cells exist a common challenge on achieving simultaneously high efficiency and long lifetime due to the presence of non-radiative recombination centers throughout the solution processed multi-crystal perovskite film, which hampers their development towards industrial applications.

The non-radiative charge recombination in perovskite solar cells have been observed to occur mostly at regions near the grain boundaries, pinholes and crystallographic defects.3 Many efforts have been focused to eliminate these recombination centers by enlarging perovskite grain size or passivating defect states in post-treatments. According to the classical LaMer curve,4 triggering fast nucleation and prolonging crystal growth process are effective ways to obtain large sized crystals. For examples, accelerating nucleation by anti-solvent dripping,5 high temperature casting,6 vacuum flashing,7 light illuminating8 etc., and retarding crystal growth through the formation of additive coordinated intermediates,9

solvent engineering,10

vapor induced

recrystallization11 etc. By these, perovskite films with grain size in the range of micrometer to millimeter can be obtained. On the other hand, post-treatments have been carried out to passivate the bulk and surface defects, such as self-forming highly conductive layer of MAI12 or PbI213 at the grain boundaries to facilitate charge transport, introducing Lewis acids14, Lewis base,15-16 ammonium halide17-18 to bind antisite defects (PbI3-, PbI42-) or under-coordinated ions (Pb2+, I-), and high optoelectronic quality perovskite films with prolonged charge diffusion length and enhanced photoluminescence quantum efficiency (PLQE) can be achieved.19 In addition, interface engineering and the use of proper charge transport layers was



favorable for the reduction of interfacial defects and hence boost the device performance and stability.

20,21

In most cases, the aforementioned methods have to be

combined to substantially suppress the non-radiative recombination centers. However, these measures either complicate the fabrication process that is likely to increase the production cost or bring extrinsic chemicals that induce insufficient charge transport and extraction, limiting their application in photovoltaics. To increase the long term stability and retain the high performance of perovskite solar cells, some new strategies need to be developed to finely regulate large crystal growth and passivate structural defects without bringing in extra techniques and chemicals.

Since the microstructure and crystal quality of perovskite thin films are determined by the crystallization dynamics, it is important to rationally design the nucleation and crystal growth process during perovskite formation. However, the conventional onestep method is hard to guarantee the perovskite quality with a high reproducibility because it is deposited by spin-coating the pre-mixed organic and inorganic precursors, and the perovskite formation associated with mass transport, chemical reaction and crystallization take place instantaneously. While in two-step sequential deposition method,22 briefly, the inorganic lead halide (e.g. PbI2 ) layer is firstly spin-coated on the substrate and then followed by an intercalation process of organic compound (e.g. MAI), where the nucleation is dependent on the crystallinity and morphology of the asspun PbI2 in the first step, and the nucleus start to grow bigger by consuming the surrounding precursors or coalescing with each other during the second step. Therefore, the crystallization dynamics and the ultimate microstructure of the as fabricated perovskite can be determined separately in the two-step method.

On the other hand, compositional engineering,23 through which different anions or cations with suitable size are incorporated into the lattice to form mixed perovskites, is able to change the crystallization dynamics,24 strengthen the chemical bond between crystal units25 or passivate the intrinsic defects.26 For examples, alloying cations such as Formamidinium (FA+), Methylammonium (MA+), Cs+, Rb+ to form binary, triple or


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quadruple cations in 3D perovskites, resulting in optimized optoelectronic properties and improved thermodynamic stability against heat, moisture and irradiation due to the optimized Goldschmidt tolerance factor.27 And also, partially substituting I- anions by small halides (Br- or Cl-) was verified to be able to slow down the crystal growth28 and suppress the formation of crystallographic defects,29 allowing for large sized and well orientated crystal growth with enhanced PLQE30 and photo-stability.31 In addition, the complex mixed perovskites can be facially prepared directly through doping the ions into the precursor solution without introducing extra techniques, and the reported efficient and stable devices are achieved mostly from the mixed perovskites. It seems that the elemental composition engineering combined in a two-step deposition method is one of the most effective approaches to simultaneously improve device performance and stability. However, the structure-property relationships to correlate the device performance with perovskite elemental composition, and how the dopants, especially for the Cl-, that are included in the precursor affect the film formation and device performance is still not clear. It was known that the Cl incorporation was benefit for the high quality perovskite crystal growth, even though the Cl signal is hardly detected in the annealed perovskite films, the existence of residual Cl in the film was wildly believed.24 Owing to their high diffusivity and reactivity, the influence of the residual Cl- dopants might have some dramatic impacts on the device performance and stability, but that has been rarely explored.

Herein, we chose a mixed perovskite of FA0.85MA0.15PbI3 as reference, and elemental composition engineering was performed by partially substituting the I- anions with small halide ions (Br-, Cl-). Systematical comparisons have been made on crystal growth mechanism, thin film morphologies and photovoltaic properties of the device fabricated on perovskites with different halide doping via a two-step sequential deposition method. It has been found that with the assistant of halide doping, the perovskite crystal growth can be well controlled, and the as fabricated perovskite thin films possess larger grain size with substantially enhanced photoluminescence quantum efficiency. A champion power conversion efficiency exceeding 20% with enhanced



storage stability can be successfully achieved from the binary halides (Br-/Cl-) doped perovskites. Halide doping assisted perovskite crystal growth The schematic illustration of the halide doping assisted perovskite growth via two-step sequential deposition method was depicted in Figure 1. The PbI2 film was firstly prepared, and then the organic precursor solution of FAI doped by 15 mol% MAI, MACl, MABr or binary MACl / MABr, respectively, in isopropyl alcohol (IPA) was spin-coated onto the PbI2 layer and dried to form a two-layer stack of solid inorganic and organic film. A post annealing was applied to facilitate the formation of corresponding mixed hybrid perovskites of FA0.85MA0.15PbI3, FA0.85MA0.15PbI3-yBry, FA0.85MA0.15PbI3-zClz, and FA0.85MA0.15PbI3-y-zBryClz which herein-after denoted as PVSK-I, PVSK-Br, PVSK-Cl and PVSK-BrCl, respectively. The MA+ was introduced in the small bandgap perovskite of FAPbI3 to enhance the formation of black -FAPbI3 phase instead of yellow -FAPbI3 phase. Noted that the final composition of the mixed perovskite was very different from that of the ratios in the precursor solution especially for the films fabricated by the two-step deposition method, we still kept the initial mole ratio of FA:MA in perovskite composition for simplification. The chemical composition of the as prepared perovskite PVSK-BrCl film was analyzed by X-ray photoelectron spectroscopy (XPS) (Figure S1). As estimated from the integrated XPS spectra, the mole ratio of Br/I at the perovskite surface was around 4%. Although no Cl element was detected, we keep the formula of PVSK-Cl and PVSK-BrCl for Cl- doped perovskites due to their crucial role in the crystallization process.

In two-step sequential deposition method, the chemical reaction between PbI2 and organic compound occurs as the organic precursor penetrates into the layered PbI2, and nucleus are generated along with the reaction and grow to a larger size. In pure iodide perovskites, the competition of anisotropic growth in different direction induces some significant strains within the crystal bulk that lead them to split into small units32-33 and set the upper limit of the grain size to the layer thickness, resulting in small crystals with numerous pinholes and grain boundaries throughout the film (Figure 1-IV).


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Figure 1. Schematic illustration of perovskite formation with different halide doping via a two-step sequential deposition method. Sequentially deposition of (I) PbI2 and (II) organic precursor, (III) intercalation process and perovskite formation through (IV) pure iodide, (V) Br- or (VI) Cl- doping.

Partially substituting I- with small halides (Cl-, Br-) in the organic precursor is profitable for the growth of high quality perovskites. Their functions can be understood from two aspects. First, the doped halides can change the nucleation dynamics. Due to the difference in ionic radii, Cl- (1.67 Å) < Br- (1.84 Å) < I- (2.07 Å),34 the small halide ions are more diffusive and are able to promote faster mass transport and deeper penetration of organic precursor into the layered PbI2.22 The infiltrated halides are prone to form various intermediate phases in terms of FAPbCl3, MAPbCl3, FAPbBr3, MAPbBr3, that trigger heterogeneous nucleation and seed the crystal growth. Second, it is known that the Gibbs free energy for the crystal growth and coalescence is dependent on the surface tension of each grains which is a function of surface charge σ2

density, as it is expressed: 8 𝛾 = 𝛾0 ― 2𝑐0, where γ0 is the surface tension at zero charge, σ is the excess surface charge and c0 is the double-layer capacitance. The doped halide ions can accumulate at the grain surface and boundaries, resulting in an excess negatively charged surface, which can reduce the surface tension and facilitate the crystal coalescence at the expense of small crystals.

In particular, it has been pointed out that it is energetically disfavored for the Cl incorporation into perovskite crystal lattice compared to Br and I.35 Most of the doped



Cl- are freely diffusive during intercalation process and then released as gaseous MACl at the post-annealing procedure. The downward infiltration and upward escape of MACl associated with an dissolve-recrystallization process36 facilitate mass transport and prolong the crystal growth process, allowing for nucleus to adjust their orientations and small crystals to coalesce with each other to minimize the total Gibbs free energy in the resultant multi-crystalline film. Therefore, perovskites with larger grains and preferred orientation are likely to be obtained via Cl- doping (Figure 1-VI). By contrast, the introduction of Br- can also enhance the crystal morphology to some extent (Figure 1-V). Owing to its relatively large size and insufficient volatility of MABr compared to Cl-, the Br- have less significant influence on crystallization dynamics and hence thin film morphologies. However, it was found that the complexation of Pb2+ with Br- was 7 times greater than that with I-.37 The doped Br- was more involved in the interaction with Pb2+ and form highly strengthened Pb-Br bond in the lattice, which was able to enhance perovskite phase stability.

As shown in Figure 2a-d, perovskite thin film morphologies were studied by scanning electron microscopy (SEM). The pristine PVSK-I film was consisted of randomly distributed small perovskite crystals along with numerous pinholes, and the grain size (~400 nm) was smaller than the layer thickness (~500 nm) (Figure 2a). Upon Br- doping, although no significant change on grain size was observed, the surface coverage of the film was clearly enhanced with minimized pinholes (Figure 2b). By contrast, the crystal quality of Cl- doped perovskite was significantly improved with obviously enhanced surface coverage, homogeneously distributed crystals and enlarged grain size (Figure 2 c&d). The averaged grain size (above 1 m) in PVSK-Cl and PVSK-BrCl films was very similar, indicating the same amount of Cl- was incorporated into the film. The grain size larger than the layer thickness (~500 nm) enabled charges transport through single grains instead of boundaries, predicting tremendously eliminated charge recombination pathways. Noticed that some white phases were appeared around the grain boundaries in Cl- doped perovskites (Figure 2 c&d), which was ascribed to the self-formed PbI2 rich region.38


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Figure 2. Thin film characterizations. Scanning electron microscopy (SEM) images of perovskites of a) PVSK-I, b) PVSK-Br, c) PVSK-Cl and d) PVSK-BrCl, (e-f) X-ray diffraction (XRD) patterns, (g) UV– vis absorption spectra, (h) steady state photoluminescence (PL) and (i) time resolved photoluminescence spectra of perovskites with different halide doping.

X-ray diffraction (XRD) characterizations were performed to confirm the conversion of PbI2 to perovskite. Figure 2e showed spectral XRD patterns of the perovskites with different halide doping. As demonstrated from the diffraction patterns, all of the precursors were successfully transformed to perovskites. The typical diffraction peaks at 14.0, 28.1, 31.5 and 42.8 were assigned to the (110), (220), (310), and (116) perovskite crystal planes, respectively. The intensity of (110) and (220) peak increased as different halides were doped into the precursor following the order of: PVSK-I < PVSK-Br < PVSK-Cl < PVSK-BrCl, indicating the improved crystallinity and orientation. It was noted that the incorporation of Br- into the crystal lattice caused the shift of the (110) peak towards higher 2θ degrees (Figure 2f), suggesting the Br- doping induced lattice shrinkage due to the formation of short Pb-Br bond in the lattice. More



shift extent was observed in PVSK-BrCl compared to PVSK-Br, hinting that the doped Cl- facilitated the Br- infiltration and incorporation into the PbI2 lattice. In addition, the PbI2 diffraction peak at 11.8 was likely to appear in the Cl- doped perovskites in terms of PVSK-Cl and PVSK-BrCl (Figure 2e and in Figure S2a), which confirmed the presence of PbI2 and agreed with aforementioned SEM images (Figure 2 c&d). It seems that the doping of Cl- is favorable for the self-formation of PbI2, while the reason for the PbI2 generation is not very clear but might be ascribed to the dissociation of intermediates of MAPbIxCl3-x phase as the escape of gaseous MACl. Initial assessment on the stability of perovskite thin films was made by monitoring their XRD patterns as they were exposed to ambient air, a clear PbI2 peak appeared in the PVSK-I films after one-hour exposure, while the spectral pattern of halide doped perovskites varied only a little (Figure S2b), indicating pinholes within perovskite film accelerating the degradation of perovskite to PbI2.

As shown in absorption and emission spectra (Figure 2g&h), the absorption on-set of PVSK-I thin film was located at the wavelength of 820 nm, giving an optical bandgap (Eg,optical) of 1.51 eV. Compared to the pristine PVSK-I, the Br- doped perovskites (PVSK-Br and PVSK-BrCl) showed a clear blue-shift (~10 nm) on the absorption onset and PL emission peak, indicating the successful incorporation of small amount of Br into the perovskite lattice39 that broadened their bandgaps to 1.53 eV. And again, PVSK-BrCl showed slightly more blue shift in PL emission peak compared to PVSKBr, confirming the Cl- doping was benefited for the Br- alloying into the perovskite lattice. Even though Cl- doped perovskites showed no obvious bandgap changes, a significant enhancement in the PL intensity (Figure 2h) was observed owing to the eliminated non-radiative electron-hole recombination centers in the Cl- incorporated perovskites. To quantify the quality of the perovskite films, the characterization of external photoluminescence quantum efficiency (PLQE) was performed. The PLQE is defined as the ratio of the number of radiatively emitted photons and the number of incident photons, which represents the quantity of nonradiative recombination channels in the film. The PLQE were measured at high excitation fluence (equal to 40 suns) when


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all charge trap states were filled by photogenerated electrons. As shown in Table 1, the PLQE was relatively low for PVSK-I films and increased as halide doping following the same order of PVSK-I < PVSK-Br < PVSK-Cl < PVSK-BrCl as the crystal quality. The achieved maximum PLQE value of 19.6% in PVSK-BrCl films indicated the substantial reduction of non-radiative decay pathways in the film.

To inspect the charge recombination dynamics, time-resolved photoluminescence (PL) measurements were conducted on perovskites with different halide doping, as shown in Figure 2i. By fitting the data with a two-exponential decay function, the recombination lifetimes extracted from the pristine PVSK-I thin film were 16.7 ns and 54.5 ns. With halide doping, the lifetimes were increased to 26.7 ns and 78.8 ns, 37.6 ns and 137.2 ns, 62.7 ns and 183.8 ns for the PVSK-Br, PVSK-Cl and PVSK-BrCl, respectively. The longer-lived charge carrier lifetime indicates a lower defect concentration and longer charge diffusion length, which was in agreement with that the enlarged grain size was beneficial in healing defects in perovskite thin films. Table 1 Optoelectronic properties of perovskite thin films with different halide doping and their photovoltaic performance of the corresponding solar cells Types of

Eg,optical

PLQE

JSC

VOC

FF

PCE

HI

(V)

(%)

(%)

(%)

perovskite

(eV)

(%)

(mA/cm2)

PVSK-I

1.51

5.5

21.9 (23.4)

1.02 (1.06)

67.9 (73.5)

15.2 (18.2)

13.6

PVSK-Br

1.53

12.1

23.0 (23.9)

1.07 (1.09)

67.7 (69.7)

16.8 (18.1)

10.9

PVSK-Cl

1.51

16.4

23.9 (23.6)

1.07 (1.09)

74.2 (77.9)

19.1 (19.9)

7.2

PVSK-BrCl

1.53

19.6

23.2 (24.0)

1.09 (1.10)

75.8 (76.4)

19.2 (20.2)

6.6

The parameters in the bracket derived from the champion device.

Perovskite solar cells To investigate the influence of halide doping on photovoltaic performance, solar cells were fabricated in a planar device configuration of ITO/TiO2/perovskite/SpiroOMeTAD/Au, where the TiO2 and Spiro-OMeTAD (2,2’,7,7’-tetrakis(N,N-di-pmethoxyphenylamine)9,9’-Spirobifluorene) were employed as electron and hole



transport layer, respectively. Vertical thin film morphologies and layer thickness of the device were presented in the cross sectional SEM images (Figure 3a). The photovoltaic parameters for averaged and champion (in the bracket) cell were summarized in Table 1, and their typical current density-voltage (J-V) curves were plotted in Figure 3b. The averaged short-circuit current density (JSC), open-circuit voltage (VOC), fill factor (FF) and power conversion efficiency (PCE) of 21.9 mA/cm2, 1.02 V, 67.9%, 15.2%, respectively, were obtained for the devices fabricated on pure iodide perovskite PVSK-I. When the perovskite was doped by Br-, the device showed enhanced performance with simultaneously improved JSC, VOC, FF and PCE to 23.0 mA/cm2, 1.07 V, 67.7% and 16.8 %, respectively, due to the suppressed pinholes. The device performance was further enhanced by employing Cl- doped perovskites (both PVSK-Cl and PVSK-BrCl) with averaged PCE exceeding 19%, which was mainly ascribed to the enlarged grain size and self-formed passivation layer of PbI2 in Cl- doped perovskites. And also, it was pointed out that residual Cl- could reduce the electron-hole coupling at grain boundaries that slowed down charge recombination40 and induced band bending at the TiO2 interfaces that improved charge collection efficiency.41 Comparing to cells based on PVSK-Cl, the device fabricated on PVSK-BrCl yielded a higher VOC and lower JSC due to the broadened bandgap and narrowed absorbing wavelength range, in agreement with the EQE measurements in Figure 3c. We integrated the EQE spectra over the AM 1.5G solar spectrum, and the calculated JSC values were in good agreement with the values determined from J-V measurements. The binary halides (Br-/Cl-) doping benefited from both advantages of Cl- and Br- doping, resulted in a champion cell with PCE of 20.2% from the cell based on PVSK-BrCl.

Statistical deviation of the photovoltaic parameters of these four types of devices were summarized from 16 cells fabricated in different batches (Figure 3d), and it was clearly seen that the reproducibility of the device was higher if the perovskites owing higher homogeneity of the crystal quality. The steady-state measurements were carried out in Figure 3e to check the device performance at the point of maximum output power, the stabilized efficiency for PVSK-I, PVSK-Br, PVSK-Cl and PVSK-BrCl was 11.5%,


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15.7%, 18.9% and 19.0%, respectively. The devices fabricated on halide doped perovskites showed enhanced initial stability under operational condition.

Figure 3. Characterization of perovskite solar cells. a) Cross sectional SEM images of the solar cells, b) current density- voltage (J-V) curves, c) external quantum efficiency (EQE), d) statistical deviations of the photovoltaic parameters, e) steady state measurements and f) open circuit voltage (Voc) as functions of light intensity of the solar cells fabricated on PVSK-I (○), PVSK-Br (◇), PVSK-Cl (△) and PVSK-BrCl (□), respectively.

To investigate how charges recombine within the device, the logarithmic relationship between VOC and light intensity was analyzed. In an ideal case, the slope (n) should be equal to kT/q, where all carriers decay through band-to-band transitions, and reaches to 2 in the case of the Shockley-Reed-Hall recombination mechanism dominates. As shown in Figure 3f, the Voc of solar cells fabricated on the PVSK-I, PVSK-Br, PVSKCl and PVSK-BrCl thin films against light intensity showed a slope of 2.46, 1.79, 1.26 and 1.23, respectively. The slope > 2 in the device based on PVSK-I indicated the presence of high density of defect states that opened another type of recombination channels wherein charges recombined via more than one defect. As the enhancement of the perovskite crystal quality via compositional engineering, the non-linear recombination pathways could be closed up. In Br- doped perovskite solar cells, the Shockley-Reed-Hall recombination dominated due to the large area of intrinsic grain boundaries which acted as trap-assisted recombination channels, resulting in a slope



close to 2. The trap states were significant eliminated as Cl- doping in perovskite (PVSK-Cl and PVSK-BrCl) that allowed for the high quality crystal growth and the majority of charges decayed through band-to-band transitions. Hysteresis phenomenon in perovskite solar cells, which was attributed to the combination of ions migration and interfacial charge recombination in the device,42 provided another hint on the crystal quality. Similar with some other TiO2 based planar perovskite solar cells, the hysteresis phenomenon occurred in our devices, and it was quantified as hysteresis index (𝐻𝐼 = |PCE𝑅𝑒𝑣𝑒𝑟𝑠𝑒𝑑 ― PCE𝐹𝑜𝑟𝑤𝑎𝑟𝑑| PCE𝑅𝑒𝑣𝑒𝑟𝑠𝑒𝑑

) as shown in Table 1. The HIs were reduced by half from 13.6% in

PVSK-I to 6-7% in PVSK-Cl and PVSK-BrCl owing to the reduced grain boundaries, which were believed to be the main channels for the ions migration.42

Therefore, there appears to be a clear correlation between the device performance and the crystal quality, that is, crystals with more defects in terms of pinholes and grain boundaries are less efficient on performance, lower reproducibility and worse operational stability, and vice versa.

Self-storage stability To access the impact of the halide doping on self-storage stability of the perovskite solar cells, the photovoltaic performance of aforementioned devices was monitored during the storage in a nitrogen filled glove box without light illumination. As shown in Figure 4a, only a slight performance variation (within 15%) occurred in all devices besides the one based on PVSK-Cl, which suffered from a dramatic PCE decay by more than 85% in 120 hours. Noted from the shape changes on the J-V curves (Figure S3a), the performance degradation was ascribed to the increase of series resistance which lead to a significant decrease in JSC, FF and PCE (Figure S3b). While it is interesting to find that when the devices were exposed to oxygen, the degraded cell recovered to almost (more than 95% of) its initial performance as shown in Figure 4a. Differently, the devices based on PVSK-I were trended to decay due to the high density of grain boundaries and pinholes which acted as main channels for oxygen penetration to


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accelerate the perovskite degradation.

Figure 4. Storage stability. a) Time evolution of power conversion efficiency of device based on perovskites with different halide doping when kept in the nitrogen glove box or oxygen exposure without light illumination, and b) 3D molecular structures of the bonded Spiro-OMeTAD-TFSI and SpiroOMeTAD-Cl during storage and O2 exposure.

Considering the significant performance degradation and recovery occurred in the device based on PVSK-Cl by the action of oxygen exposure, the introduced Cl- and spiro-OMeTAD was suspected for the reversible performance behavior. As similar with the reported literatures43, although no chloride signal was detected in PVSK-Cl films, we agreed that a trace amount of Cl- (under the detecting limit) could be retained in the films. Additionally, it is known that the pristine spiro-OMeTAD is a poor hole transport material, and the doping of lithium bis((trifluoromethyl)sulfonyl)amide (Li-TFSI) and 4-tert-butylpyridine (tBP) together with oxygen exposure have been proven as an efficient prerequisite to enhance its conductivity.

In order to distinguish if the dopants or the oxidized state of spiro-OMeTAD was accountable, the self-storage stability of the devices based on PVSK-Cl were compared by employing Li-TFSI and tBP doped spiro-OMeTAD without oxygen exposure treatment (Figure S4), and some other organic HTLs instead of spiro-OMeTAD such as

poly[bis(4-phenyl)(2,4,6-trimethylphenyl)amine]

(PTAA)

and

poly(3-

hexylthiophene-2,5-diyl) regioregular (P3HT) which need to be doped by Li-TFSI and tBP but oxygen exposure is not required to enhance their conductivities (Figure S5). In addition, an inverted device configuration of ITO/NiO/PVSK-Cl/PCBM/Ag was also fabricated to test the stability (Figure S6). Even though the achieved performance of



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the above devices were lower than that of using oxidized spiro-OMeTAD as HTL, no significant degradation was observed within 120 hours when stored in the same condition (Figure S7). We also test the stability of perovskite solar cells based on MAPbI3-xClx, which was fabricated using one step deposition method by mixing MAI and PbCl2 in a mole ratio of 3:1. Similarly, a fast performance degradation during storage was observed when employing doped Spiro-OMeTAD as HTL, which could be effectively alleviated by using doped PTAA and P3HT instead of Spiro-OMeTAD (Figure S7). Therefore, the influence of the dopants (Li-TFSI and tBP) and the residual Cl- ions self on the reversible performance degradation, and the different device stability behavior arise from different ratio of cations in the perovskite composition was excluded, then we attributed it to the interaction between the residual Cl- ions and the oxidized state of spiro-OMeTAD.

As demonstrated by Abate and Wang et al.,44-45 the oxidation of spiro-OMeTAD follows a two-step reaction mechanism: first, an equilibrium is established between spiro-OMeTAD, oxygen molecule and oxidized spiro-OMeTAD (spiro-OMeTAD+O2-) in Equation (1); second, the addition of Li-TFSI moves the equilibrium forward and generate some weakly bounded ionic species of spiro-OMeTAD+TFSI- in Equation (2), resulting in effective generation of mobile holes on the organic matrix and hence high conductivity. Spiro-OMeTAD + O2  Spiro-OMeTAD+O2-

(1)

Spiro-OMeTAD+O2- +Li-TFSI  Spiro-OMeTAD+TFSI- + Li+O2-

(2)

Spiro-OMeTAD+TFSI- + Li+O2-+Cl-  Spiro-OMeTAD+Cl-+ LiCl+ O2-

(3)

The residual Cl- in PVSK-Cl perovskites has a large impact on the oxidation state of spiro-OMeTAD. As calculated from density functional theory (DFT), the ionic pair association strength of Spiro-OMeTAD+ with Cl- is 0.54 eV higher than that with TFSI(Figure 4b). And also, the binding energy of small halide (Cl-) with Li+ ions (LiCl) is much higher than Li-TFSI,46-47 the presence of Cl- in HTL may change the reaction direction of Equation (2) and move the equilibrium in Equation (1) backward to non-


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oxidized spiro-OMeTAD. It was recognized that the distribution of residual Cl- in perovskite film were depth dependence, and most of them were initially located at the TiO2/perovskite interface and no Cl was detected on the surface.48 Therefore, the residual Cl- showed no influence on the fresh device performance. Owing to the high diffusivity, the free Cl- ions in PVSK-Cl were able to diffuse within the device and reach at spiro-OMeTAD layer. The formation of highly bounded Spiro-OMeTAD+Clspecies in Equation (3) led to an insufficient hole generation on the organic matrix and therefore a dramatic reduction in conductivity and hence device performance. When the devices were exposed to oxygen again, the excessive oxygen molecules could slowly re-oxidize the spiro-OMeTAD and recover its high conductivity. In comparison, the mobility of free Cl- in PVSK-BrCl is lower than that in PVSK-Cl owing to the as formed Pb-Br bond in the crystal lattice, and no obvious performance degradation has been observed in the device based on PVSK-BrCl. Overall, although the Cl- doping is benefit for the large crystal growth, the residual Cl- can disturb the equilibrium and deoxidize the spiro-OMeTAD and hence shorten the stability of the devices that employing spiro-OMeTAD as HTL. To simultaneously boost the device efficiency and long term stability, the introduced Cl- should be either fixed in the lattice through compositional engineering (e.g. Br incorporation) that enhance the binding energy of each unit in the crystal lattice or used in the device employing other HTL, e.g. PTAA, P3HT etc., or inverted device structure using NiO instead of spiro-OMeTAD.

Conclusion The crystal quality of perovskite, in terms of crystal size, grain boundary and phase stability, is essential to the performance and long term stability of perovskite solar cells. It was found that compositional engineering, e.g., partially substituting the I- with small halide of Br- or Cl-, can significantly promote the formation of high quality perovskites. In particular, the Cl- doping could modify the crystallization dynamics in terms of facilitating fast nucleation and prolonging crystal growth process that enabled the perovskite growth with enlarged grain size (above 1m in ~500 nm thick films) and preferred crystal orientation. Although the residual Cl- ions had a negligible impact on



the initial performance of the resultant photovoltaic devices, they were suspiciously more mobile and could diffuse to the ubiquitously used HTL of spiro-OMeTAD. The highly bounded ionic pairing of Cl- with oxidized spiro-OMeTAD+ significantly reduced the conductivity of spiro-OMeTAD and caused a fast performance degradation. The above problem was effectively avoided by binary halides (Br-/Cl-) doping owing to the incorporation of Pb-Br bonds in the lattice that enhanced the crystal stability by improving the binding energy of each unit. The as formed binary halides doped perovskites PVSK-BrCl with substantially eliminated non-radiative recombination centers (PLQE of 19.6%) resulted in a champion PCE of 20.2% in the corresponding cells with enhanced storage stability. Therefore, carefully engineering on perovskite composition is an efficient way to enhance the crystal quality towards high efficiency and stable perovskite solar cells. Additionally, the development of efficient hole transport materials instead of Spiro-OMeTAD or the use of inverted device structure would be another possibility to enhance the device stability.

Experimental: Device fabrication: The commercial ITO patterned glass substrates were cleaned in cleaning agent, deionized water, acetone, and isopropanol under sonication for 10 min in sequence, then treated by UV-ozone plasma for 10 min. The 50 nm electron transport layer (ETL) of titanium diisopropoxide bis(acetylacetonate) (Sigma-Aldrich) stabilized TiO2 nanoparticle solution (4-8 nm diameter, 20 wt% in water, PlasmaChem GmbH) was coated on ITO substrate at 6000 r.p.m., and annealed at 150 °C for 30 min in air. The perovskite layer was then deposited by a sequential two-step spin coating method; first, 1.2 M of PbI2 (99%, Sigma-Aldrich) in anhydrous N,N-Dimethylformamide (DMF) (99.8%, Sigma Aldrich) was spin coated onto the ETL at 1500 r.p.m. for 30 s and annealed at 100 °C for 10 min to remove the remaining solvent; Second, after the PbI2 coated substrates cooling to room temperature (25 °C), 0.3 M of CH(NH2)2I (98%, Sigma Aldrich) in 1 ml isopropanol with 15 mol% of CH3NH3I, CH3NH3Br, CH3NH3Cl or mixed CH3NH3Br and CH3NH3Cl, respectively, was spin coated onto the PbI2 at 2500 r.p.m. for 45 s. The thermal annealing was carried out at 150 °C for 20 min with


Page 18 of 26

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the presence of small amount (5 L) DMF in a closed environmental to form perovskite of PVSK-I, PVSK-Br, PVSK-Cl and PVSK-BrCl, respectively. The hole transport layer (HTL)

was

prepared

by

dissolving

72.3

mg

(2,2’,7,7’-tetrakis(N,N-di-p-

methoxyphenylamine)9,9’-Spirobifluorene) (Spiro-OMeTAD, Sigma Aldrich), 28.9 mL 4-tert-butylpyridine (99.9%, Sigma-Aldrich) and 17.5 mL of a stock solution of 520 mg/ml lithium bis(trifluoromethylsulphonyl)imide in acetonitrile (99.9%, SigmaAldrich) in 1 mL chlorobenzene (99.9%, Sigma-Aldrich). The ~200 nm thick HTL was deposited by spin coating the solution at 3000 r.p.m. for 45 s. For the Spiro-OMeTAD layer oxidation, the uncompleted devices were stored in a dry air desiccator for 12 hours before the electrode deposition. Finally, the Au (100 nm, 1 Å·s-1) was deposited by thermal evaporation under a pressure of

Influence of Cl incorporation in perovskite precursor on crystal growth

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