High-Temperature Antiferroelectric of Lead Iodide Hybrid Perovskites

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A High-Temperature Antiferroelectric of Lead Iodide Hybrid Perovskites Shiguo Han, Xitao Liu, Yi Liu, Zhiyun Xu, Yaobin Li, Maochun Hong, Junhua Luo, and Zhihua Sun J. Am. Chem. Soc., Just Accepted Manuscript • DOI: 10.1021/jacs.9b05124 • Publication Date (Web): 31 Jul 2019 Downloaded from pubs.acs.org on July 31, 2019

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A High-Temperature Antiferroelectric of Lead Iodide Hybrid Perovskites Shiguo Han,†,§,# Xitao Liu,†,# Yi Liu,†,§ Zhiyun Xu,† Yaobin Li,† Maochun Hong,† Junhua Luo† and Zhihua Sun†,* † State

Key Laboratory of Structural Chemistry, Fujian Institute of Research on the Structure of Matter, Chinese Academy of Sciences, Fuzhou, Fujian 350002, China § University

of Chinese Academy of Sciences, Chinese Academy of Sciences, Beijing 100039, P. R. China

Supporting Information Placeholder ABSTRACT: Antiferroelectric, characterized by the natural

polarization-electric field (P-E) double hysteresis loops, has been developed as a promising branch for energy storage. Here, we present the first antiferroelectric in the booming family of lead iodide hybrid perovskites, (BA)2(EA)2Pb3I10 (1, where BA = n-butylammonium and EA = ethylammonium), which exhibits one of the highest Curie temperatures (~363 K) for the majority of known molecular systems. Strikingly, its high-temperature antiferroelectricity, triggered by an antipolar alignment of adjacent dipoles, is convinced by the characteristic double P-E hysteresis loops, thus enabling remarkable energy storage efficiencies in the range of 65% ~83%. This merit is almost comparable to many inorganic counterparts, suggesting great potentials of 1 for energy storage. Another fascinating attribute is that 1 also acts as a room-temperature biaxial ferroelectric with spontaneous polarization of 5.6 μC∙cm-2. As far as we know, this study on the high-temperature antiferroelectric, along with roomtemperature biaxial ferroelectricity, is unprecedented for the versatile lead-iodide hybrid perovskites, which sheds light on the design of new electric-ordered materials and facilitates their application of high-performance devices.

Electroactive materials that are capable of energy storage and conversion have long been an important topic for the academic research, and their applications are indispensable as the basic elements for advanced electronic devices.1 In this portfolio, antiferroelectric, which demonstrates the adjacent dipoles initially aligned in an antiparallel arrangement,2 has taken an indispensable status. What makes antiferroelectrics fascinating is the natural polarization versus electric field (PE) double hysteresis loops, stemmed from the switching of antipolar dipoles through applying strong electric field.3 This attribute is distinguishing from the linear dielectrics and ferroelectrics, and renders antiferroelectrics extraordinary capacities for electric energy storage, including high energy storage density and fast discharging rate, etc.4 Since the original identification of antiferroelectricity in PbZrO3, 5 despite substantial endeavors, the amount of existing antiferroelectric materials remains quite scarce, for which

the mainstay is limited to inorganic oxides 6 (e.g. PbZrO3, PbHfO3 and AgNbO3) and some molecular compounds,7 such as NH4H2PO4, Cu(HCOO)2∙4H2O, [H-55dmbp][Hca], and (3pyrrolinium)CdBr3, etc. In this context, it is highly urgent to explore new antiferroelectric systems as the essential ingredients for future fabrication of high-performance energy storage devices.8 Lead iodide hybrid perovskites, most notably MAPbI3 (MA = CH3NH3+), have recently emerged as the revolutionary class of electroactive materials for photoelectric and photovoltaic applications.9 Among them, multilayered two-dimensional (2D) hybrid perovskites display unique physical attributes of structure tunability, phase stability, quantum and dielectric confinement effects.10 In contrast to the single-cation counterparts, the underlying merit of 2D multilayered system is their acentric and polar nature.11 Particularly, the high degree of structural flexibility allows for an admission of dynamic organic cations between the inorganic perovskite sheets, of which the reorientation controlled by electric field provides great opportunities to design new electric-ordered materials.12 For instance, ferroelectricity has been achieved in some 2D hybrid perovskites, such as (BA)2(MA)x-1PbxBr3x+1 13 and [S-1-(4-chlorophenyl)ethylammonium]2PbI4, 14 triggered by the dynamic ordering of organic cations. Nevertheless, to date, antiferroelectric of lead iodide hybrid perovskites still remains unexplored, since it requires the antipolar alignment of adjacent dipoles that can be switched by applying electric field,3a, 15 along with the typical double P-E hysteresis loops. Consequently, continuous efforts should be devoted to subtly tailoring antiferroelctrics in the family of lead iodide hybrid perovskites, which have been developed as a promising class of electroactive materials. Here, we report the first high-temperature antiferroelectric of lead iodide hybrid perovskite, (BA)2(EA)2Pb3I10 (1), which adopts the 2D trilayered motif with EA cation locating in the inorganic perovskite sheets. The distinct antiferroelectric property of 1 is convinced by the double P-E hysteresis loops, accompanied by the high Curie temperature (Tc = 363 K). Such attribute enables high energy storage efficiencies ranging from 65% to 83%, comparable to some inorganic oxides, revealing its great potential for energy storage.

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Besides, 1 also exhibits room-temperature biaxial ferroelectricity with large spontaneous polarization (Ps) of 5.6 μC cm-2. This study on the high-temperature antiferroelectric of lead iodide hybrid perovskites, along with multiaxial ferroelectricity, paves an avenue to exploring new electric-ordered materials and promotes their potentials for future device application. Crystals of 1 were obtained from hydroiodic solutions by the temperature-cooling method, and the phase purity was confirmed by powder XRD pattern (Figures S1 and S2). The preliminary DSC traces display two pairs of reversible thermal peaks at 322/315 K (T1) and 363/360 K (T2) in the heating/cooling mode, revealing an occurrence of two phase transitions for 1 (Figure 1a). This successive phase transition behavior is also confirmed by distinctive dielectric anomalies (Figure 1b), of which the large peak-like anomalies around T2 involve with dramatic structural changes. 3b Our ensuing studies on bulk physical properties disclose that three states for 1 can be defined as paraelectric phase (PEP, T > T2), antiferroelectric phase (AFEP, T1 < T < T2) and ferroelectric phase (FEP, T < T1), respectively.

Figure 1. (a) DSC traces for 1. (b) Temperature dependence of dielectric constant (εˊ) measured along the c-axis direction upon cooling. Inset: the enlarged view of εˊ around T1. Structure analyses reveal that 1 crystallizes in space group Cmc21 at 298 K (point group mm2, Table S1). As depicted in Figure 2a, the basic architecture for 1 consists of [Pb3I10]∞ trilayered frameworks and organic BA cations riding in the interlayer space, constructing the 2D crystallographic oriented perovskite motif.10 It is notable that the cornersharing PbI6 octahedra of [Pb3I10]∞ trilayers adopt a slightlydistorted geometry (Figure S3), as disclosed by the deviation of Pb-I bonds and I-Pb-I angles (Table S2). This characteristic accounts for the formation of distorted square-like cavities enclosed by terminal halides, which can incorporate EA+ cation in the [Pb3I10]∞ sheets. For most cases of hybrid perovskites, the cations in the inorganic sheets are small-size MA+, FA+ and Cs+; it is very difficult to accommodate larger cations.16 However, for 1, high degree of inorganic framework distortion allows for the alloying of EA+ cation into perovskite cavities (Figure S4). Such organic cations are closely linked to PbI6 octahedra via N-H∙∙∙I hydrogen bonds, and the CH3CH2- moieties orientate along the c-axis. Meanwhile, bulky BA+ moieties exhibit the same orientation along its c-axis, with -NH3 groups linked to inorganic sheets through N-H∙∙∙I hydrogen bonds (Figure S5). This parallel arrangement of dipoles gives rise to its polar alignment. Owing to the long alkyl chain of CH3(CH2)3- moiety, dynamic motions of the adjacent organic BA+ cations easily occur, which provide the driving force to phase transition around T2.

Figure 2. Crystal structures of 1 at different phases. (a) The -oriented 2D perovskite motif at FEP viewed along b-axis. The arrows denote parallel array of dipoles related to dynamic orientation. (b) Packing diagram at AFEP. The antiparallel array of dipoles accounts for its antiferroelectric order. (c) Packing structure at PEP. Yellow-dashed line denotes the crystallographic mirror plane.

At AFEP, crystal structure of 1 belongs to space group Pbca (point group mmm). Figure 2b depicts that its perovskite motif is still identified as an octahedral tilting architecture, which results in the emergence of local polarization. Different from its FEP, however, an antiparallel alignment is assumed to satisfy the requirement of crystallographic symmetry for antipolar order. For instance, both BA+ and EA+ cations in the unit cell have two orientations (Figure 2b). Symmetry-related moieties in the neighboring slab displace in exactly opposite directions with the same magnitude. Meanwhile, the tilting of PbBr6 octahedra also occurs, as verified by the off-center displacements of Pb ions (Figure S3). This bipartite array of electric dipoles yields an unstable antipolar mode with higher free energy than that of its competing FEP,17 which cancels out the net polarization of the unit cell at AFEP. Further heating beyond T2, 1 transforms to tetragonal space group I4/mmm (point group 4/mmm). The characteristic is that both BA+ and EA+ cations become disordered with eight disordered orientational populations, and locate on the crystallographic mirror plane (Figure 2c); the inorganic PbI6 octahedra also feature a highly-symmetric configuration. This centrosymmetric packing of 1 eliminates its macroscopic polarization at PEP. Above structure analyses suggest that order-disordering of organic BA+ and EA+ cations renders a driving force to phase transitions of 1, while the emergence of polarization relates to the reorientation of organic parts and tilting motion of inorganic components. Symmetry breaking of 4/mmmFmm2 at T1 belongs to one of the 88 species for ferroelectrics and reveals its biaxial nature.18 Concretely, the number of polarization directions can be defined as n = Np/Nf, where Np and Nf denote the number of symmetry operation elements at PEP and FEP, respectively. For 1, phase transition from PEP (Np = 16) to FEP (Nf =4) results in four polarization directions.18b Figure S6 depicts the polarized direction is confined along the c-axis at FEP, coinciding with [110]direction at PEP. Due to four equivalent [110]-directions at PEP, the biaxial nature can be verified for 1, which is also confirmed by the electrical measurements (Figure S7).

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S10 and S11). This figure-of-merit is comparable to many inorganic oxides,6, 22

Figure 3. Physical properties related to symmetry breaking in 1. (a) Variable-temperature SHG effect. Inset: SHG signals of 1 and KDP at room temperature. (b) Temperature-dependent Ps of 1 obtained from the pyroelectric measurement. Inset: pyroelectric current recorded in the cooling mode.

Variable-temperature second harmonic generation (SHG) and pyroelectric properties of 1 were measured to verify its symmetry breaking. Figure 3a displays that its SHG signal is unresponsive above T1, coinciding with its centrosymmetric structures (for both Pbca at AFEP and I4/mmm at PEP). Upon cooling below T1, SHG signal emerges and the intensity is measured to be 0.4 times that of KDP at room temperature (the inset Figure 3a), revealing its acentric structure. This strong temperature dependence of SHG activities confirms the broken symmetry of 1. Pyroelectric results suggest the emergence of ferroelectric polarization below T1 and the estimated Ps value is 5.6 μC/cm2 (at 298 K, Figure 3b), which is comparable to other hybrid ferroelectrics.19 Besides, the temperature dependence of Ps shows a coincident tendency with that of the SHG activities, well following the Landau theory equation.20 Hence, the variable-temperature behavior of Ps and SHG is reminiscent of its potential ferroelectricity. As the convincing evidence, the P-E hysteresis loops of 1 were measured to verify its ferroelectric and antiferroelectric behaviors. Figure 4 depicts the typical J-E curves and P-E hysteresis loops for ferroelectrics. At 298 K, the Ps value of 1 is measured as 5.5 μC/cm2, which agrees with that obtained from pyroelectric current. This figure is comparable to some other hybrid perovskite ferroelectrics,21 such as (pyrrolidinium)MnCl3 (5.4 μC/cm2) and (trimethylchloromethyl-ammonium)MnCl3 (4.0 μC/cm2), but smaller than that of organic perovskite (22 μC/cm2). Another interesting merit is that 1 behaves as the biaxial ferroelectric with narrow band gap of 1.9 eV, and the polarization can be switched with electric field parallel to crystallographic b- and c-axes (Figures S7 and S8). What makes 1 more intriguing is the emergence of high-temperature antiferroelectricity between T1 and T2. At AFEP, the typical double P-E hysteresis loops and J-E traces of 1, unlike that recorded below T1, behave as the proof to its antiferroelectricity. The forward switching electric field (AFE-to-FE transition, EAFE-FE) and backward switching electric field (FE-to-AFE transition, EFEAFE) are ~32 and ~10 kV/cm, respectively. On account of energy criterion, this field-induced transition mainly results from their free energy difference.17 According to the double P-E loops, temperature-dependent energy storage efficiencies were estimated in the range of 65%~83%. For instance, the AFE attribute of 1 enables an energy storage density of 0.15 J/cm3 at 355 K, and energy storage efficiency (η = Wreco/(Wreco+Wloss) is estimated as 83% (at 60 kV/cm, Figures

Figure 4. Ferroelectric and antiferroelectric properties of 1. (a, b) The current versus electric field (J-E) traces collected at FEP and AFEP, respectively. (c) P-E hysteresis loops measured at FEP. (d) The double P-E hysteresis loops for electric energy storage.

such as Ag(Nb0.8Ta0.2)O3 (77%), Sr0.98Ca0.02TiO3 (72.3%) and (Pb0.95Sr0.05)ZrO3 (68%). Particularly, the Curie temperature for the antiferroelectricity of 1 reaches up to 363 K, at which thermal and dielectric properties exhibit dramatic changes. As far as we know, this is one of the highest figures for the molecular AFE systems,23 including NH4H2PO4 (148 K), (CH3)2NH2]Zn(HCOO)3 (156 K), NH4H2AsO4 (216 K), (C4H5NH3)CdBr3 (246 K), [H-55dmbp][Hca] (318 K), and (iBA)2CsPb2Br7 (353 K). In terms of structure variability and tunability, this work on high-temperature antiferroelectric of lead iodide halide perovskites will be inspiring to the precise molecular design and optimizing antiferroelectric performances for electric energy storage. In summary, we have reported the first high-temperature antiferroelectric of lead iodide hybrid perovskites, showing a high Curie temperature and large energy storage efficiency. Such intriguing attributes makes it a promising candidate for electric energy storage. Another underlying merit is that 1 is also a room-temperature biaxial ferroelectric with narrow bandgap (~1.9 eV). For photovoltaic “superstar” of lead iodide hybrids (particularly MAPbI3), the experimental evidence of bulk ferroelectricity has been a contradictory issue.9d It is believed that 1 might be a target model approaching that of MAPbI3, which can be used to elucidate the underlying role of ferroelectricity to photovoltaic performances. This study paves a pathway to the exploration of new electric-ordered materials and facilitate their electronic device application. ASSOCIATED CONTENT

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Supporting Information Experimental method, Figures S1-S11 and Tables S1-S5. CCDC 1912877 - 1912879 contain structure data for 1. AUTHOR INFORMATION Corresponding Author [email protected] Author Contributions S. Han and X. Liu contributed equally to this work. Notes The authors declare no competing financial interests. ACKNOWLEDGMENT This work was supported by the National Natural Science Foundation of China (NSFC 21622108, 21875251, 21833010 and 21525104), NSF of Fujian Province (2018H0047), the Strategic Priority Research Program of Chinese Academy of Sciences (XDB20010200). The referees are greatly appreciated for their excellent suggestions to improve the quality of manuscript. REFERENCES (1) (a) Chu, B.; Zhou, X.; Ren, K.; Neese, B.; Lin, M.; Wang, Q.; Bauer, F.; Zhang, Q. M. A Dielectric Polymer with High Electric Energy Density and Fast Discharge Speed, Science 2006, 313, 334; (b) El-Kady, M. F.; Strong, V.; Dubin, S.; Kaner, R. B. Laser Scribing of HighPerformance and Flexible Graphene-Based Electrochemical Capacitors, Science 2006, 335, 1326; (c) Larcher1, D.; Tarascon, J-M. Towards Greener and More Sustainable Batteries for Electrical Energy Storage, Nature Chem. 2015, 7, 19. (2) Blinc, R.; Zeks, Soft Modes in Ferroelectrics and Antiferroelectrics, North-Holland, 1974. (3) (a) Kittel, C. Theory of Antiferroelectric Crystals, Phys. Rev. 1951, 82, 729; (b) Lines, M. E.; Glass, A. M. Principles and Applications of Ferroelectrics and Related Materials; Oxford University Press, New York, 1977. (4) Liu, Z.; Lu, T.; Ye, J.; Wang, G.; Dong, X.; Withers, R.; Liu, Y. Antiferroelectrics for Energy Storage Applications: a Review, Adv. Mater. Technol. 2018, 3, 1800111. (5) (a) Sawaguchi, E.; Maniwa, H.; Hoshino, S. Antiferroelectric Structure of Lead Zirconate, Phys. Rev. 1951, 83, 1078; (b) Shirane, G.; Sawaguchi, E.; Takagi, Y. Dielectric Properties of Lead Zirconate. Phys. Rev. 1951, 84, 476. (6) (a) Mischenko, A. S.; Zhang, Q.; Scott, J. F.; Whatmore, R. W.; Mathu, N. D. Giant Electrocaloric Effect in Thin-Film PbZr0.95Ti0.05O3, Science 2006, 311, 1270; (b) Zhao, L.; Liu, Q.; Gao, J.; Zhang, S.; Li, J.-F. Lead-Free Antiferroelectric Silver Niobate Tantalate with High Energy Storage Performance, Adv. Mater. 2017, 29, 1701824; (c) Li, J.; Li, F.; Xu, Z.; Zhang, S. Multilayer Lead-Free Ceramic Capacitors with Ultrahigh Energy Density and Efficiency, Adv. Mater. 2018, 30, 1802155. (7) (a) Schmidt, V. H. Review of order-disorder models for KDP-family crystals, Ferroelectrics, 1987, 72, 157; (b) Lasave, J.; Koval, S.; Dalal, N.S.; Migoni, R. L. Origin of Antiferroelectricity in NH4H2PO4 from First Principles, Phys. Rev. Lett. 2007, 98, 267601; (c) Li, P.-F.; Liao, W.Q.; Tang, Y.-Y.; Ye, H.-Y.; Zhang, Y.; Xiong, R.-G. Unprecedented Ferroelectric-Antiferroelectric-Paraelectric Phase Transitions Discovered in an Organic-Inorganic hybrid perovskite, J. Am. Chem. Soc. 2017, 139, 8752. (8) (a) Horiuc hi, S.; Kagawa, F.; Hatahara, K.; Kobayashi, K.; Kumai, R.; Murakami, Y.; Tokura, Y. Above-Room-Temperature Ferroelectricity and Antiferroelectricity in Benzimidazoles, Nature Commun. 2012, 3, 1308; (b) Kobayashi, K.; Horiuchi, S.; Ishibashi, S.; Murakami, Y.; Kumai, R. Field-Induced Antipolar-Polar Structural Transformation and Giant Electrostriction in Organic Crystal, J. Am.

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