Facet-Dependent Control of PbI2 Colloids for over 20% Efficient

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Letter

Facet-Dependent Control of PbI2 Colloids for Over 20% Efficient Perovskite Solar Cells Chenxin Ran, Weiyin Gao, Nengxu Li, Yingdong Xia, Qi Li, Zhaoxin Wu, Huanping Zhou, Yonghua Chen, Minqiang Wang, and Wei Huang ACS Energy Lett., Just Accepted Manuscript • DOI: 10.1021/acsenergylett.8b02262 • Publication Date (Web): 28 Dec 2018 Downloaded from http://pubs.acs.org on January 1, 2019

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ACS Energy Letters

Facet–Dependent Control of PbI2 Colloids for Over 20% Efficient Perovskite Solar Cells Chenxin Ran,a,* Weiyin Gao,a Nengxu Li,c Yingdong Xia,d Qi Li,e Zhaoxin Wu,a, f Huanping Zhou,c,* Yonghua Chen,b,d,* Minqiang Wang,a Wei Huangb,d aDepartment

of Electronic Science and Technology, School of Electronic and Information

Engineering, Xi’an Jiaotong University, Xi’an 710049, Shaanxi, P. R. China bShaanxi

Institute of Flexible Electronics (SIFE), Northwestern Polytechnical University

(NPU), Xi'an 710072, Shaanxi, P. R. China. cBeijing Key Laboratory for Theory and Technology of Advanced Battery Materials, Department

of Materials Science and Engineering, College of Engineering, Peking University, Beijing 100871, P. R. China. dKey

Laboratory of Flexible Electronics (KLOFE) & Institution of Advanced Materials

(IAM), Jiangsu National Synergetic Innovation Center for Advanced Materials (SICAM), Nanjing Tech University (NanjingTech), Nanjing 211816, Jiangsu, P. R. China. ePhysical

Science Division, IBM TJ Watson Research Center, Yorktown Heights, NY

10598, USA. fCollaborative

Innovation Center of Extreme Optics, Shanxi University, Taiyuan

030006, China

*E–mail: [email protected] [email protected] [email protected]

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Abstract Organic–inorganic hybrid perovskite based solar cells (PSCs) have made impressive progress during the past few years, and record power–conversion efficiency over 23% has been achieved. However, for solution–processed perovskite film, one of the major challenges for the film quality improvement is the lack of the regulation strategy of the colloids in the precursor solution. Here, we demonstrate facet–dependent colloid engineering to adjust the quality of perovskite film and its photovoltaic performance. This is realized by the development of a facet passivation strategy to synthesize PbI2 (S–PbI2) with tunable morphology. Further, facet– dependent PbI2 colloids regulation is demonstrated, and the relations among morphology of PbI2 crystal, features of PbI2 colloids, morphology of PbI2 film, and photovoltaic performance of MAPbI3 are established. The PSCs based on 10% S–PbI2 exhibited a best PCE of 20.22 % with improved device reproducibility and hysteresis index compared to that using the commercial PbI2 (~18%).


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The emerging development of efficient electronic devices (e.g., solar cells, light– emitting diodes, photo–detectors, and transistors) based on organic–inorganic hybrid lead halide perovskite materials has been one of the most exciting research fields in the past few years.1–5 The perovskite materials have been identified to be the most promising candidate in the next generation photovoltaic devices. Specifically, perovskite solar cells (PSCs) has obtained newest record power conversion efficiency (PCE) of 23.3 %, and holds its great promising in the large–scale commercialization as silicon alternative due to its solution processibility, low processing cost, abundant raw materials, high light utilization efficiency, flexibility, and acceptable long–term stability.6–9 To boost the photovoltaic performance of PSCs, one of the most effective strategies is to improve the quality of the solution–processed perovskite film. Strategies, such as processing engineering,10–14 interface engineering,15–18 solvent engineering,19–22 composition engineering,23–25 additive engineering,26–29 and colloidal engineering,30–36 have been developed to modify the crystallinity, orientation, grain size, defect trap states, carrier transportation, and optical responses of the perovskite film. When fabricating perovskite film via solution method, the precursor solution is primarily prepared. It has been noticed that the characteristics of the colloids in precursor solution strongly determine the features of the perovskite film.30–36 In 2015, Yan et al. reported that adjusting the stoichiometric ratio of PbI2/MAI/MACl could tune the coordination degree and mode of PbI2/MAI colloids in the initial colloidal dispersion, leading to improved quality of MAPbI3 film using one–step spin–coating method.32 Thereafter, dimethylsulfoxide (DMSO),33,

34

hydrohalic

acids,35 and aging time36 are proposed to be effective in changing the characteristics of PbI2/MAI colloidal dispersion. In addition to one–step spin–coating film formation method, adjusting the PbI2 colloidal dispersion in sequential film fabrication process has also made some progress. Wu et al.30 used a small amount of H2O to make a homogenous PbI2 precursor by changing the dispersibility of the PbI2 colloids. It is found that the morphology, crystallinity, and orientation of the PbI2 film are shown to be tunable upon H2O addition. Later, Li et al.31 used acetonitrile (ACN) to reduce the 4

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average size of PbI2 colloids, though the size distribution of PbI2 colloids was still shown to be unchanged. Despite the progresses being made, the main determinant factors and the underlying mechanism for colloidal controlling are still lack of understanding. It has been demonstrated that the raw materials in the precursor solution of MAPbI3 perovskite actually form a colloidal dispersion rather than really dissolved solution, so the characteristics of raw material should also have an impact on the colloidal properties, which further determines the quality of perovskite film.37 This helps explaining the diversity results obtained from different labs even following the same film fabrication process, because different raw materials (from different manufacturers and/or different batches) could lead to different features of colloids in precursor. However, little has been known about how the features of the raw materials influence the properties of colloids in precursor, quality of perovskite film, and the photovoltaic performance of PSCs. In this work, a facet–dependent colloid engineering strategy is developed to demonstrate that how the raw materials colloids engineering may translate into the film features and whole device performance. For the first time, synthesized PbI2 (S– PbI2) with controlled morphology and size distribution is realized, and its growth mechanism is theoretically investigated via first principle calculation. It is found that hydrogen (H) atom could preferentially adsorb and chemically bond with I atom on (101) facet of PbI2 crystal and lower its surface energy. This facet passivation capability leads to diversity in growth rate of different facets, leading to tunable morphology of PbI2 crystal. It is further observed that the PbI2 colloid in the prepared precursor shows a tunable size distribution depending on the facet proportion of PbI2 crystal. As a result, the properties of the corresponding PbI2 films as well as MAPbI3 films can be facilely adjusted depending on different PbI2 raw materials. The n–i–p PSCs device using S–PbI2 as raw material shows a PCE of 20.22 % (under reverse scan) compared to that using commercial PbI2 (C–PbI2) of 18 %. More importantly, device based on S–PbI2 also exhibits superior device reproducibility and hysteresis index compared to that based on C–PbI2.

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Figure 1. (a–j) SEM images of S–PbI2 synthesized with different AA concentration: (a, f) 5% S–PbI2, (b, g) 10% S–PbI2, (c, h) 20% S–PbI2, (d, i) 40% S–PbI2, and (e, j) 80% S– PbI2. (k) Schematic diagram of the morphology evolution of S–PbI2 under increased AA concentration. (l) XRD patterns of S–PbI2 samples. (m) Intensity ratio evolution of (001) diffraction peak to (101) diffraction peak in XRD patterns.

It has been reported that the weak acid environment is necessary to obtain PbI2 phase (pH ≤ 6) instead of Pb(OH)I phase (pH ≥ 7), and a small amount of acetate acid (AA) has been used to create such weak acid environment.38 Following this procedure, we surprisingly find that by adjusting the AA concentration in the reaction 6 ACS Paragon Plus Environment

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solution, the morphology (both shape and size) of the prepared PbI2 crystal is shown to be different. Figure 1 and Figure S1 shows the SEM images and photographs of the as–prepared S–PbI2 samples, respectively. It can be observed that the morphology of the PbI2 crystal changes as the AA concentration increasing. Specifically, as the AA concentration increasing from 5 wt% to 80 wt% (named as 5% S–PbI2 – 80% S–PbI2), the proportion of (101) facet of PbI2 crystal increases, while the proportion of (001) facet decreases. It is worth noting that when the AA content is very low (5% S–PbI2), there are many unstructured PbI2 crystals existing (Figure 1a), and the well–shaped hexagonal tetradecahedron structure with neat surface is formed as the AA concentration increasing (Figure 1f–j). This observation indicates the important role of AA on the formation of well–shaped PbI2 crystal. In addition, it is found that the size of the PbI2 crystal dramatically decreases when the AA concentration increases (Figure S2). Table 1 collects the proportion of (001) and (101) facets as well as mean crystal size of the S–PbI2 samples. From the XRD measurement (Figure 1l, m), it can be observed that the intensities of (00l) peaks significantly reduce, while the intensities of (h0l) peaks increase as the AA concentration increasing. This is in agreement with the SEM images that show reducing proportion of top surface and increasing proportion of bevel surface of PbI2 crystal as AA concentration increasing. Moreover, due to the easy evaporability of AA, the prepared S–PbI2 shows no organic residue (Figure S3). The purity of raw material is mostly important to obtain the perovskite film with high quality, so the S–PbI2 prepared in our work can be directly applied in perovskite film fabrication. Table 1 Proportion of (001) and (101) facets and crystal size in different S–PbI2 samples. AA Conc. Sample (μm) (001)a (101)a Mean Size (μm)b (wt%) 5% S–PbI2 10% S–PbI2 20% S–PbI2 40% S–PbI2 80% S–PbI2

5 10 20 40 80

75.8% 66.7% 59.2% 44.7% 34.3%

24.2% 33.3% 40.8% 55.3% 65.7%

22.51 20.21 16.48 10.68 4.25

aCalculated

from the surface area of each facet of 10 typical particle in SEM image.

bCalculated

from every particles in SEM image. 7 ACS Paragon Plus Environment


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To disclose the mechanism of the morphology evolution of the S–PbI2 samples, the crystal growth model is adopted to describe the growth kinetics of PbI2 crystal. The Wulff’s theorem notes that the thermodynamically controlled crystal growth process tends to minimize the total surface free energy of the crystal, so the growth rate of a certain facet is proportional to its surface free energy per unit area (eV/Å2).39 This means that the facet with higher surface energy possesses a faster growth rate, which would cause the gradual disappearance of this facet during growth procedure, leading to the reduced proportion of this facet in the final crystal.40 It has been well– established that the presence of a capping agent can preferentially interact with a specific facet of a crystal, and the order of the surface energies of different facets could also be changed.41,

42

In our case, because the increased AA concentration is

observed to lead to the morphology variation of PbI2 crystal, it is reasonable to hypothesize that hydrogen (H) atom might function as capping agent and change the surface energy of PbI2 crystal. Table 2 Calculated surface energies () for the clean PbI2 (001) and (101) facets terminated by Pb layer (Pb–ter.) and I layer (I–ter.). For the (001) facet, only I termination is applicable. Calculated surface energies for the fully H covered (H) PbI2 (001) and (101) facets terminated by I atoms, and the single H binding energies (Eb). The last column lists the I–H bond length (LI–H). “(i)” and “(o)” denotes the inner and outer I sites, as depicted in Figure 2b.

Pb–ter.

I–ter.

H

(eV/Å2)

(eV/Å2)

(eV/Å2)

(001)

N/A

0.027

0.141

(101)

0.175

0.114

0.079

Facet

Eb (eV)

LI–H (Å)

–0.037 –0.252 –0.221 (o) (i)

1.710 1.633 (i)

1.723 (o)

To verify this assumption, density functional theory (DFT) (calculation details can be seen in Supporting Information) is carried out to calculate the surface energies of (101) and (001) facets of PbI2 crystal before and after H adsorption. The calculated surface energies of clean (001) and (101) facets of PbI2 crystal with and without H adsorption are listed in Table 2. It should be noted that that (001) facet is not able to be terminated by Pb, while (101) facet terminated by Pb shows much higher values (0.175 eV/Å2) than that terminated by I (0.114 eV/Å2). This result means Pb exposure 8 ACS Paragon Plus Environment

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is not favorable in PbI2 crystal, so the bonding of CH3COO– group with Pb2+ during PbI2 formation is not thermodynamically possible. Therefore, the case of PbI2 (001) and (101) facets terminated by I atoms is considered in this work. When no H atom is presented, the (001) facet of PbI2 crystal shows the lowest surface energy. Due to the nature of the layered structure of PbI2 crystal, the surface energy of the (001) facet is related to the interlayer Van der Waals (vdW) interaction, which is negligible without vdW–DF included (0.003 eV/Å2). Even if considering the vdW–DF exchange–correlation functional, which suggests stronger interlayer interaction, the calculated surface energy of the (001) facet is also much lower compared to that of the (101) facet. This result indicates the dominant exposure of (001) facets in naturally formed PbI2 crystal, which is in agreement with our experimental finding of the extra exposure of the (00l) facet when PbI2 is prepared using classic precipitation method from supersaturated PbI2 precursor in neutral aqueous solution (Figure S4).

Figure 2. Unrelaxed and relaxed structures of (a) (001) and (b) (101) facet of PbI2 crystal that are fully covered by H atoms viewed along the z–axis (perpendicular to the direction of the slab stacking) of the supercells. The x–direction of all the supercells is cut to emphasize the surface structure. “(i)” and “(o)” in part (b) denotes the inner and outer adsorbed H sites, respectively. The gray spheres represent the Pb atoms, purple for I atoms and cyan for H 9 ACS Paragon Plus Environment


atoms. (c) Illustration of the growth mechanism of PbI2 crystal under different H concentrations.

When H atom is presented, starting from the relaxed clean surface configurations, the I–H bond length is manually set to 1.650 Å in both the unrelaxed structures of (001) and (101) facets. As illustrated in Figure 2a, it is found that the (001) facet is less stable after fully covered by H atoms. The distance between the H–adsorbed I atom and the adjacent Pb atoms increases significantly from 3.248 Å to 3.674 Å upon the existence of the H atom. In addition, our Bader charge analysis reveals that the charge transfer between the H–adsorbed I atom and its neighboring Pb atoms almost vanishes in the relaxed configuration. The similar configurational relaxation is found in the single H adsorption (1/32 H coverage) case, where the binding energy (Eb) between the H and I atoms on (001) facet is also negligible as listed in Table 2. Besides, the H covered (001) facet has a much higher surface energy (0.141 eV/Å2) even compared to the clean (001) facet (0.027 eV/Å2). These results demonstrate that the Pb–I bonds are effectively “broken” upon H binding with I on (001) facet, which means the bonding of H on (001) facet is unfavorable. In other words, the H adsorption on the (001) surface can be regarded as a Van der Waals interaction rather than bonding, which is better to be considered as hydrogenated (001) surface. On the contrary, the (101) surface appear to “hold” the adsorbed H atoms in a more superior way, because it accommodates the H atoms by undergoing a more significant surface reconstruction without breaking any Pb–I bonds (Figure 2b). With the assistance of a stronger binding, the inner H atoms (site “(i)”) are pulled even closer to the surface. Therefore, all the adsorbed H atoms form an overall “zigzag” contour, which supports a more energetically favorable charge density distribution on the surface. The adsorbed I atom site holds 7.58 e– in this configuration, about 0.14 e– is from the H atom and the rest 0.44 e– is from the adjacent Pb–I bonding. The H covered (101) facet has a surface energy of 0.079 eV/Å2, which is lower than both the clean (101) surface (0.114 eV/Å2) and the fully hydrogenated (001) surface (0.141 eV/Å2). All these results show that H preferentially adsorbs on (101) facet of PbI2 crystal and 10 ACS Paragon Plus Environment

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significantly lowers its surface energy. As a result, when H concentration is high, according to Wulff’s theorem discussed above, H–covered (101) facet with lowered surface energy will expose more in the formed PbI2 crystal. These results perfectly explain the morphology evolution of PbI2 crystal as AA concentration increasing observed in Figure 1. The growth process of PbI2 crystal with hexagonal tetradecahedron structure is depicted in Figure 2c. For the crystal growth, the PbI2 monomer will first adsorb on all the PbI2 facets and then migrate to the edges of these facets to promote the facet growth. According to the calculation results above, under low H concentration, the surface energy of (101) facet is higher (i.e., higher growth rate of (101) facet), and thus its adjacent facet (001) will become dominate in the final crystal, shaping like plate. While under high H concentration, the (101) facet is “passivated” by the adsorbed H atoms, and the surface energy of (101) facet becomes lower (i.e., higher growth rate of (001) facet). Therefore, the (101) facet will be in higher proportion in the final crystal, shaping like frustum. It is worth noting that under high H concentration, the lateral growth of small PbI2 crystal is simultaneously restrained because of the H passivation of (101) facet, so the size of PbI2 crystal becomes smaller as well as AA concentration increasing. To demonstrate the facet–dependent variation on the features of PbI2 colloids, we characterize the S–PbI2 colloidal dispersions in DMF. The Tyndall effect of all the S–PbI2 dispersions (Figure 3b) indicates their colloidal nature. The dynamic light scattering (DLS) measurements (Figure 3c) demonstrate that the S–PbI2 colloidal dispersions show decreased average size and narrowed size distribution as AA content increasing. For example, 5% S–PbI2 dispersion shows average size at ~ 1000 nm with broad size distribution, while 80% S–PbI2 dispersion shows average size at ~ 420 nm with much narrower size distribution. When forming PbI2 colloids in DMF, layered PbI2 will be exfoliated by DMF molecule through (101) facet or defect sites on (001) facet due to the coordination interaction between nitrogen in DMF and Pb2+ in PbI2.43 In this case, PbI2 colloid flakelets composed of (001) planes could be formed. This can be evidenced by the DLS measurement that the PbI2 colloids show 11



two distinguished size distributions, which are attributed to the lateral dimension and longitudinal dimension of PbI2 (001) plate (Figure S5). Therefore, the size distribution of PbI2 colloids shows a high dependence on the proportion of different facets and crystal size of PbI2 raw material (Figure 3d, e). It is worth noting that the commercial PbI2 (C–PbI2) material (from Sigma Aldrich) shows average size at ~ 900 nm with relatively broad size distribution (Figure 3c), which is consistent with the results reported in the literatures.31,

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This broad size distribution of C–PbI2

colloidal dispersion is in accordance with the inhomogeneous crystal size of the C– PbI2 raw material (Figure S6), demonstrating again the dependence between raw material and colloids in dispersion.


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Figure 3. (a) Photographs, (b) Tyndall effect, and (c) DLS measurements of S–PbI2 colloidal dispersions in DMF. Schematic diagram of the formation of PbI2 colloidal dispersion and PbI2 film from (d) 5% S–PbI2 and (e) 80% S–PbI2. (f–j) SEM images, (k) grain size distribution, and (l) XRD of PbI2 films made from different S–PbI2 samples.

Series of S–PbI2 films made from different S–PbI2 precursors are then characterized, and a disciplinary evolution of the film morphology can be observed. The S–PbI2 films (Figure 3 f–j) show increased surface coverage with gradual disappeared voids on the film as the AA concentration increasing. This morphology evolution is shown to be closely depending on the size distribution of PbI2 colloids. For example, as depicted in Figure 3d, 5% S–PbI2 with larger average size and broader size distribution leads to a PbI2 colloids with larger average size and broader size distribution. This broad size distribution of PbI2 colloids leads to non–continuous morphology of 5% S–PbI2 film with distinct large voids, and the size distribution of the grains is also broad (Figure 3k). As AA content increasing, the continuity is improved and the grain size is reduced. 80% S–PbI2, with smaller average size and narrower size distribution of crystals (Figure 1j), forms a colloidal dispersion with reduced colloid size as well as narrowed size distribution (Figure 3e). This 80% S– PbI2 dispersion leads to denser PbI2 film composed of small PbI2 grains with narrow size distribution (Figure 3k). For C–PbI2, the deposited C–PbI2 film shows medium–sized voids on the film (Figure S7), which is consistent with observed morphology evolution in S–PbI2. The XRD patterns of these S–PbI2 films (Figure 3l) show only typical (00l) peaks that indicates the parallel orientation of PbI2, which verifies the exfoliation of PbI2 (00l) planes by DMF molecule when forming PbI2 colloidal dispersion. In addition, the decreased intensity of (00l) peaks indicates that the reduced size and narrowed size distribution of S–PbI2 colloids (from 5% S–PbI2 to 80% PbI2) lead to the suppressive (001) orientation. Meanwhile, the full width at half maximum (FWHM) of the (001) peaks tend to broaden. This is in agreement with the SEM observation of decreased grain size of S–PbI2 films. It is worth noting that owing to the (001) orientation, the root–mean–square (RMS) roughness of all the S– 13 ACS Paragon Plus Environment


PbI2 films is measured to be less than 10 nm (Figure S8). Moreover, the absorption spectra of S–PbI2 films (Figure S9) also show distinct differences. For example, 5% S–PbI2 film with larger voids shows higher absorption intensity. This is due to the light scattering and refraction effect of those micro–sized voids on the film.44 All these observations above can be explained as follows: the features of PbI2 crystal determine the properties of PbI2 colloid, which further determines the morphology evolution of PbI2 film. All in all, these results demonstrate a facile method to adjust the features of PbI2 film (morphology, orientation, and optical property) via raw material engineering of the PbI2 crystal. More importantly, this controllability of PbI2 film is of great strategic significance, because when making MAPbI3 film following sequential film deposition process, the morphology and orientation of PbI2 film have been reported to show decisive impact on the quality of the resultant MAPbI3 film.45 This strong correlation is the basis for the regulation and control of the properties of MAPbI3 perovskite film through raw material engineering.


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Figure 4. (a–e) SEM images of MAPbI3 made from different S–PbI2 films. Schematic diagram of the conversion from PbI2 grains into MAPbI3 grains in (f) 5% S–PbI2 and (g) 80% S–PbI2 films. (h) Grain size distribution and (i) XRD pattern of MAPbI3 films.

The MAPbI3 films converted from S–PbI2 films are sequentially prepared by the deposition of MAI layer on top of the S–PbI2 films followed by post–annealing process. The SEM images in Figure 4a–e shows the morphology of these MAPbI3 films. It can be found that compact films composed of densely packed MAPbI3 grains with size of several hundred nanometers are formed. The average size of the MAPbI3 grains decreases and the size distribution narrows (Figure 4h, from 5% MAPbI3 to 80% MAPbI3). This morphology variation can be explained by the variation of available space for volume expansion from PbI2 to MAPbI3. It has been reported that the volume of PbI2 will expand when the intercalation of MAI molecules into the PbI2 15



planes to form MAPbI3 perovskite structure. This volume expansion originates from the crystallographic reconstruction of the edge–sharing octahedra of PbI2 structure (density 6.16 g/cm3) to corner–sharing octahedra of MAPbI3 perovskite structure (density 4.16 g/cm3), which involves ~ 2.0 times of volume expansion.46–48 In this case, 5% S–PbI2 film with larger voids possesses larger available space allowing for volume expansion compared to the 80% S–PbI2 one with negligible voids (Figure 4f, g). According to the grain size statistics analysis in Figure 4h, the porous 5% S– PbI2 film leads to MAPbI3 film with larger grains, while the compact 80% S–PbI2 film with negligible voids leads to MAPbI3 film with smaller grains. The tunable space for volume expansion is beneficial to control the compressive strain between the grains, and this residual strain in the perovskite film has been shown to plays key role in determining the photovoltaic performance of the film.49–51 However, MAPbI3 film made from 5% S–PbI2 film shows fluctuant surface morphology (Figure S10), which is due to the overlarge space for volume expansion. The morphology of different MAPbI3 films can be also identified by AFM images (Figure S11), where the RMS roughness surface roughness reduces as the AA concentration increases. Except for the different morphologies, all these MAPbI3 films show similar XRD peaks (Figure 4i), which indicates the formation of the typical tetragonal phase MAPbI3 crystal with space group I4/mcm.48 Besides, the optical properties (UV–Vis and PL spectra in Figure S12) of these MAPbI3 films are observed to be well– matched, indicating the optical identity of these films.


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Figure 5. (a) The PCE distribution of the PSCs with different PbI2 as raw material. (b) The external quantum efﬁciency spectrum together with integrated JSC for perovskite solar cell based on 10% S–PbI2. (c) The J–V curve with reverse (1.2 V~–0.2 V) & forward (–0.2 V~1.2 V) scan of the best device based on 10% S–PbI2. (d) The stabilized output holding the voltage at the maximum power point (0.94 V) for the 10% S–PbI2 based PSC device.

We further investigate the photovoltaic performance of the PSCs using different S–PbI2 as raw materials. The devices architecture is based on regular n–i–p structure of ITO/SnO2/perovskite/spiro–OMeTAD/Au. More details are described in experimental section. Figure 5a shows the PCE distribution of the PSCs with C–PbI2 and all S–PbI2 samples. It is found that the C–PbI2 based PSCs yield average PCE about 18% with wide statistical distribution. With respect to the S–PbI2 based PSCs, an apparent improvement of both PCE and device reproducibility can be found as the AA concentration increasing, where 10% S–PbI2 based PSCs perform best. However, further increasing the AA concentration would lead to the PSCs with decreased PCE. Figure S13 displays the detailed characteristics of these devices, where slight difference is observed in the short–circuit current density (Jsc), but the open circuit voltage (Voc) and fill factor (FF) hold the same tendency with PCE. We attributed this phenomenon to the quality of the perovskite films induced by different raw materials. 17 ACS Paragon Plus Environment


As discussed above, although 5% S–PbI2 film with large voids allows volume expansion, the relatively rough surface hampers the formation of planar film and might increase the defects density along grain boundaries (GBs). To quantitatively confirm that, we developed space–charge–limited current (SCLC) measurements by sandwiching the different perovskite films between indium tin oxide (ITO) and gold (Au), respectively, and characterized the evolution of the space–charge–limited current with different biases. The defects densities of different samples are shown in Figure S14 and Figure S15. Note that C–PbI2 based sample has the highest trap density (8.95 ± 0.92 × 1017 cm–3) while this value of 5% S–PbI2 is 7.97 ± 0.67 × 1017 cm–3, 20% S–PbI2 is 7.24 ± 0.61 × 1017 cm–3. PSCs fabricated by 10% S–PbI2 has the lowest trap density (4.96 ± 0.80 × 1017 cm–3), which primely accords with PCE distribution of these PSCs. When increasing the AA concentration, the void size of S– PbI2 decrease and denser perovskite film can be formed. However, the limited available space on PbI2 film will lead to small grains of perovskite film, accompanying with strong compressive strain during volume expansion of the grains, which leads to the relatively poor performance. Thus, 10% S–PbI2 based perovskite film possesses the best film quality, yielding the highest PCE in the resulting PSCs. More importantly, it shows higher device reproducibility as well as the lower hysteresis compared with C–PbI2 one. Figure 5b shows the spectrum of the external quantum efficiency (EQE) of 10% S–PbI2 based PSCs along with the integrated JSC by calculating the overlap integral between EQE spectrum and standard AM1.5 solar emission. The high EQE over 90% can be achieved in the range of 370 – 400 nm and 450 – 550 nm, and EQE over 80% in the whole range of visible light region. The integrated Jsc value agrees well with the average measured results, confirming the accuracy of devices performance. Figure 5c shows the current density–voltage (J–V) curves of the best 10% S–PbI2 based PSC (forward and reverse scan directions, scan speed set at 33 mV/s). The best performing device exhibits negligible hysteresis (hysteresis index of 0.018 compared to 0.113 for C–PbI2 device) (Figure S16), with a Voc of 1.1 V, Jsc of 23.31 mA/cm2, FF of 78.21 %, and PCE of 20.22 % by reverse scan, and a Voc of 1.109 V, Jsc of 23.38 mA/cm2, FF of 77.26 %, and PCE of 18 ACS Paragon Plus Environment

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19.86 % by forward scan. In addition, a stabilized PCE of 19.91% is shown in Figure 5d, which is achieved by holding the bias voltage at 0.94V at the maximum power point. In summary, we develop a facile strategy to synthesize PbI2 crystal with controllable facet exposure, based on which the facet–dependent colloid engineering towards the features of MAPbI3 film and its photovoltaic performance is demonstrated. Theoretical calculation reveals that H atom can preferentially adsorb and chemically bonded with I atom on the (101) facet of PbI2 crystal and lower its surface energy, enabling the controllability of the morphology of PbI2 crystal. By using the different PbI2 crystal with tunable morphology, a facet–dependent evolution on the features of PbI2 colloids and spin–coated PbI2 films can be realized. Taking advantageous of the robust controllability of the characteristics of the S–PbI2 films, the photovoltaic performance of MAPbI3 film can be facilely adjusted. The perovskite film using 10% S–PbI2 as raw material yields a superior PCE of 20.22 % compared to that using C–PbI2 (18 %). More importantly, PSCs based on S–PbI2 also show superior device reproducibility and hysteresis index. This work, for the first time, addresses the significant role of raw material (i.e., PbI2) on the photovoltaic performance of the PSCs. This universal facet–dependent colloid engineering strategy could provide important practical regulative method for the fabrication of high quality perovskite film, and holds great promising in further boosting the PCE of the PSCs in the future.

Supporting Information The Supporting Information is available free of charge on the ACS Publications website at DOI: XXXX Experimental section, and all figures referenced in the Letter.

Notes The authors declare no competing financial interest. 19 ACS Paragon Plus Environment


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Acknowledgements This work was financially supported by the Natural Science Foundation of China (Grants

51802253),

National

Key

R&D

Program

of

China

(Grant

No.

2016YFB0400702), the China Postdoctoral Science Foundation (Grant No. 2017M613137), the Fundamental Research Funds for the Central Universities (Grant No. xjj2018018), and the support from the China Scholarship Council (CSC). This work was also financially supported by the National Basic Research Program of China, Fundamental Studies of Perovskite Solar Cells (Grant 2015CB932200), the Natural Science Foundation of China (Grants 11574248, 51602149, 61705102, 51722201, 51672008, and 91733301), Natural Science Foundation of Jiangsu Province, China (Grants BK20161011 and BK20161010), Young 1000 Talents Global Recruitment Program of China, Jiangsu Specially–Appointed Professor program, “Six talent peaks” Project in Jiangsu Province, China. This work was also supported by National Key Research and Development Program of China Grant No. 2017YFA0206701, Natural Science Foundation of Beijing, China (Grant No. 4182026), and Young Talent Thousand Program. The SEM work was performed at International Center by Dielectric Research (ICDR), Xi’an Jiaotong University, Xi’an, China.

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Figure 1. (a-j) SEM images of S-PbI2 synthesized with different AA concentration: (a, f) 5% S-PbI2, (b, g) 10% S-PbI2, (c, h) 20% S-PbI2, (d, i) 40% S-PbI2, and (e, j) 80% S-PbI2. (k) Schematic diagram of the morphology evolution of S-PbI2 under increased AA concentration. (l) XRD patterns of S-PbI2 samples. (m) Intensity ratio evolution of (001) diffraction peak to (101) diffraction peak in XRD patterns. 199x209mm (300 x 300 DPI)



Figure 2. Unrelaxed and relaxed structures of (a) (001) and (b) (101) facet of PbI2 crystal that are fully covered by H atoms viewed along the z-axis (perpendicular to the direction of the slab stacking) of the supercells. The x-direction of all the supercells is cut to emphasize the surface structure. “(i)” and “(o)” in part (b) denotes the inner and outer adsorbed H sites, respectively. The gray spheres represent the Pb atoms, purple for I atoms and cyan for H atoms. 199x120mm (300 x 300 DPI)


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Figure 3. (a) Photographs, (b) Tyndall effect, and (c) DLS measurements of S-PbI2 colloidal dispersions in DMF. Schematic diagram of the formation of PbI2 colloidal dispersion and PbI2 film from (d) 5% S-PbI2 and (e) 80% S-PbI2. (f-j) SEM images, (k) grain size distribution, and (l) XRD of PbI2 films made from different S-PbI2 samples. 199x274mm (300 x 300 DPI)



Figure 4. (a-e) SEM images of MAPbI3 made from different S-PbI2 films. Schematic diagram of the conversion from PbI2 grains into MAPbI3 grains in (f) 5% S-PbI2 and (g) 80% S-PbI2 films. (h) Grain size distribution and (i) XRD pattern of MAPbI3 films. 199x223mm (300 x 300 DPI)


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Figure 5. (a) The PCE distribution of the PSCs with different PbI2 as raw material. (b) The external quantum efﬁciency spectrum together with integrated JSC for perovskite solar cell based on 10% S-PbI2. (c) The J–V curve with reverse (1.2 V~-0.2 V) & forward (-0.2 V~1.2 V) scan of the best device based on 10% S-PbI2. (d) The stabilized output holding the voltage at the maximum power point (0.94 V) for the 10% S-PbI2 based PSC device. 199x135mm (300 x 300 DPI)


Facet-Dependent Control of PbI2 Colloids for over 20% Efficient

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