Introduction

Since the discovery of ferroelectricity in 1920 in Rochelle salt, enormous research interests have been dedicated to the study of ferroelectrics1. Prior to the renaissance of organic ferroelectrics in the past decade, research interests have always been centred on inorganic materials. In contrast to typical inorganic ferroelectric perovskites such as BaTiO3 (BTO) and Pb(Zr,Ti)O3 (PZT)2,3, the organic counterparts embrace a unique set of appealing attributes, including high flexibility, fracture-resistance, and solution-processability4,5. However, the construction of all-organic perovskites is constrained by the types of organic cations and anions that comply with Goldschmidt tolerance factor, which restricts discoveries to a few examples. Among these, the ferroelectric ones are even rarer6,7. For example, the recently discovered isostructural family of piperazinium2+ all-organic perovskites are mostly non-ferroelectric8,9. Achieving ferroelectricity in all-organic perovskites is challenging because polar molecular building blocks generally prefer to adopt an antiparallel arrangement in order to minimize the electrostatic energy of the system10,11,12,13. One recent breakthrough is the discovery of ferroelectric MDABCO-NH4-I3 (MDABCO = N-methyl-N′-diazabicyclo[2.2.2]octonium) perovskite14, whose spontaneous polarization was reported to be comparable to that of BTO. This organic ferroelectric surpasses many inorganic ferroelectrics in terms of parameters such as electrocaloric strength, Curie-Weiss constant and critical field for phase transition15. However, a major drawback of MDABCO-NH4-I3 is its low coercive field (<12 kV/cm). Given that a coercive field larger than 100 kV/cm is required for voltage switching in the range of 1–2 V for a 100 nm thick film16, the coercive field of MDABCO-NH4-I3 would require a thickness of at least 830 nm to enable operation at 1 V switching-voltage. This presents a roadblock for the miniaturization of the ferroelectric device. On the other hand, a coercive field higher than 1000 kV/cm (in PVDF-TrFE or HfO2) reduces the thickness of the ferroelectric channel to 10 nm, but the drawback is the device is susceptible to a high leakage current, or even dielectric breakdown17,18. It is therefore highly desirable to achieve a coercive field on the scale of 100 kV/cm, which demands new design strategies for all-organic ferroelectrics.

Herein, we demonstrate the modulation of the coercive field in organic ferroelectric perovskites via the reinforcement of hydrogen bonding. An all-organic perovskite crystal MDABCO-NH4-(PF6)3 was successfully synthesized, where strong N–H…F hydrogen bonds are formed by the N–H groups of MDABCO with the highly electronegative F (Fig. 1) in PF6. The coercive field of this all-organic compound reaches 110 kV/cm and its remnant polarization of 5.7 μC/cm2 agrees with our Berry phase calculated value of 6.5 μC/cm2. The coercive field can be enhanced further to 138 kV/cm by deuteration, which corroborates the fact that hydrogen bonding influences the coercive field. Our study highlights that coercive field in metal-free perovskites can be enhanced by regulating the strengths of intermolecular hydrogen bonds, suggesting that organic ferroelectrics have wide chemical tunability.

Fig. 1: Schematic of hydrogen-bond-assisted modulation of polarization and coercive field for all-organic ferroelectric perovskites.
figure 1

Red box shows ferroelectric polarization in inorganic perovskite due to ion displacement; Blue box shows ferroelectric polarization due to hydrogen bonds in organic ferroelectric PVDF; Black box shows a combination of both ion displacement and hydrogen bonds contribute to ferroelectric polarization in organic perovskite.

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Results and discussion

As a typical organic perovskite, MDABCO-NH4-I3 is composed of divalent amine (MDABCO) at the cuboctahedral centre (A-site), NH4+ cation at the octahedral centre (B site) and I− anions occupying the vertices of the octahedron (X site). As MDABCO exhibits the highest remnant polarization among known diamines, we chose it for the A-site cation14. To strengthen the hydrogen bond, we replaced I− with PF6− anions (Fig. 2a), the latter is compatible with NH4+ and MDABCO2+, yielding a perovskite structure with a tolerance factor of 0.914 (Table S1). MDABCO-NH4-(PF6)3 (denoted as MNP3 hereafter) crystal was grown by the slow evaporation of a mixture with a stoichiometric ratio of MDABCO-(PF6)2 and NH4PF6, where a single-crystal of 5 × 4 × 1 mm size could be obtained (Fig. 2b) Detailed synthesis procedure is described in Supplementary Information.

Fig. 2: Crystal structure of MDABCO-NH4-(PF6)3 and the polar displacement direction.
figure 2

a The A2+, B+ and X− components, and their relative positions in MDABCO-NH4-(PF6) unit cell; b as-grown 5 × 4 × 1 mm (l × d × h) size single-crystal. c Arrangement of different ions along the [111] polar axis, with arrows indicating the displacement of MDABCO molecules. The N–H…F hydrogen bonds (d < 2.4 Å) are highlighted with dotted orange lines. d Table of unit cell volume, distance in [111] direction (d111), displacement of A2+ cation (δA), and B+ cation (δB) for MNI3, MNP3 and D-MNP3.

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The structure of MNP3 was solved by single-crystal X-ray diffraction (SC-XRD), revealing that MNP3 adopts the polar trigonal R3 space group at room temperature (Table S2). The purity of the crystal can be judged from the good agreement between experimental and simulated powder X-ray diffraction (PXRD) patterns (Fig. S1). The overall structure is isostructural to MDABCO-NH4-I3 (MNI3), but with a bigger unit cell (a = b = c = 7.844 Å, α = β = γ = 84.857° for MNP3 and a = b = c = 7.259 Å, α = β = γ = 84.767° for MNI3) (Figs. S2, S3). The synthesized crystals are rhombohedral shaped and are 〈100〉 textured according to XRD (Figs. S4 and S5). There are multiple hydrogen bonds in MNP3, including N–H…F and C–H…F, with bond lengths ranging from 2.2 to 4.0 Å (Figs. S6 and S7). Particularly, the N–H bond of MDABCO is aligned in the [111] polarization direction of MNP3, and form hydrogen bonds with the three nearest fluorine atoms (Fig. 2c). The strong hydrogen bond is revealed from the bond distance, the N–H…F hydrogen bonds (2.4 Å) are much shorter than N–H…I in MNI3 (3.2 Å). The strong N–H…F bond causes MDABCO to be pulled towards [111] direction. As a result, the displacement of MDABCO in the [111] direction is 73.56 pm (4.98%) for MNP3, which is longer than MNI3 whose displacement is 65.56 pm (4.79%) (Fig. 2d).

To confirm the effect of hydrogen bond on ferroelectricity, we also synthesized deuterated MNP3 (D-MNP3) in which the proton in N–H bond of MDABCO is substituted by deuterium19,20,21. SC-XRD shows that unit cell dimension of D-MNP3, with a = b = c = 7.872 Å, and α = β = γ = 84.780°, is slightly larger and shows less octahedral tilt compared to MNP3. This increased unit cell is apparent in the PXRD (Fig. S8). The unit cell dimension (d111) stretches along the hydrogen-bond direction, i.e. [111] direction, which manifests the Ubbelohde effect22. This phenomenon is caused by the longer N–D bond than N–H, so that displacement of MDABCO (δA) is reduced from 73.56 pm to 72.56 pm by deuteration (Fig. 2d). The longer N–D bond is caused by the stronger D-F bonding, which reduces the electron density on D; the distance between the two nitrogen atoms in MDABCO is increased from 2.501 to 2.518 Å by deuteration.

A hallmark of ferroelectricity is the ferroelectric-to-paraelectric phase transition. Phase transition can be detected using differential scanning calorimetry (DSC) measurement, where a sharp endothermic peak can be seen at 311 K (Fig. 3a). MNP3 has a larger cuboctahedral unit cell than MNI3, thus MDABCO in MNP3 can rotate more freely, resulting in a lower phase transition temperature than MNI3 (448 K). The correlation between free space rotation and phase transition temperature is also demonstrated by the isostructural perovskite MDAB-Rb-I3, where its lower phase transition temperature (430 K) than MNI3 is correlated to its larger unit cell23.

Fig. 3: The ferroelectric-to-paraelectric transition of MNP3 and D-MNP3 crystal.
figure 3

a DSC and b temperature-dependent SHG intensity change of MNP3 and D-MNP3 revealing phase transition from ferroelectric (FE) phase to paraelectric (PE) phase, and that D-MNP3 has a higher Tc than MNP3. Temperature-dependent PXRD result of MNP3 (c) and D-MNP3 (d). e Solid-state 1H NMR spectrum of MNP3 at 293 K (FE phase) and 333 K (PE phase). Inset shows peak shift (6.88–6.84) of N–H in MDABCO indicating shrinkage of N–H bond. f Schematic illustration of MNP3 structure in ferroelectric (FE) phase and paraelectric (PE) phase. In the PE phase, molecules rotate freely.

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The phase transition temperature of D-MNP3 is 320 K, which is 19 K higher than MNP3. The effect of isotope substitution on phase transition temperature is a characteristic feature of hydrogen-bonded ferroelectrics such as KDP family19,24, glycine phosfite20, and supramolecular complexes4,25,26. This implies that all-organic perovskite share characteristics of organic ferroelectrics although the origin of its polarization arises from ionic displacement like inorganic perovskites. The higher phase transition temperature in deuterated perovskite originates from its stronger hydrogen bonds. A stronger hydrogen bond increases the barrier for the rotation of MDABCO in the PE phase similar to H/F substituted molecular ferroelectrics27. In addition, the much larger enthalpy change of D-MNP3 (94.5 J g−1) as compared with MNP3 (30.4 J g−1) is indicative of the stronger hydrogen bonding in D-MNP3 as compared to MNP3.

Ferroelectric-paraelectric phase transition was confirmed by a sharp change in second harmonic generation (SHG) intensity across the Curie temperature. The SHG signal originates from the non-centrosymmetric structure of materials, so it vanishes when centrosymmetry is restored in the paraelectric phase. Figure 3b shows the temperature-dependent SHG signal of MNP3 and D-MNP3 collected in the temperature range from 305 to 340 K. The SHG intensity of MNP3 gradually decreases with temperature and vanishes around 315 K due to the recovery of centrosymmetry in the paraelectric phase. Similarly, the SHG signal of D-MNP3 vanishes around 323 K when the crystal transits to the centrosymmetric phase, in line with the DSC results.

MNP3 and D-MNP3 undergo a phase transition from the ferroelectric R3 space group to the paraelectric cubic P432 space group similar to what has been reported for isostructural perovskite structures14,23. The paraelectric unit cell dimensions of MNP3 are a = b = c = 7.9557 Å, and α = β = γ = 90°, and for paraelectric D-MNP3, a = b = c = 7.9154 Å, and α = β = γ = 90° which confirmed by SC-XRD. To monitor structural changes during the phase transition process, temperature-dependent powder X-ray diffraction (PXRD) for both MNP3 and D-MNP3 were recorded. At the transition temperature of each sample, the XRD peaks of (100), (111) and (200) at 11.26°, 20.14° and 22.68° (Fig. S9) shift to lower angles of 11.12°, 19.36° and 22.4°, respectively, indicating that the unit cell dimension has expanded; the vanishing of other peaks indicate that a more symmetric cubic cell is achieved in paraelectric phase. It is observed that the (111) peaks due to ferroelectric phase at 20.14° and the paraelectric phase at 19.36° coexist without other intermediate states during phase transition (Fig. 3c 313 K and 3d 323 K), this suggests that the phase change is a discontinuous order-disorder type (first-order)28.

The fingerprints of ferroelectric-to-paraelectric phase transition can also be observed in the solid-state 1H NMR spectrum (Fig. 3e). The CH2 and CH3 peaks of MDABCO at 3.9 and 3.2 ppm become sharper in the paraelectric phase due to the free rotation of the MDABCO molecule. The peak of N–H of MDABCO at 6.88 is split into two due to a second-order effect of dipolar coupling to the quadrupolar nitrogen-14 nuclei29. This peak shift from 6.88 to 6.84 ppm that accompanies phase transition is indicative of a reduction of the N–H bond length because of the weakening of the hydrogen bonds in the paraelectric phase30. Furthermore, the P–F distance is also shortened in the paraelectric phase at 333 K as reflected from the solid-state NMR, where the J-coupling constant is reduced (Fig. S10). Therefore, solid-state NMR results validate that N–H and P–F bonds are elongated in the ferroelectric phase due to strong hydrogen bonds, whereas in the PE phase, both MDABCO and PF6 rotate freely and N–H and P–F bonds are shorter. The orientational disorder of molecules in the paraelelectric phase has been extensively studied previously by Jun Harada and others31,32,33,34,35. One characteristic of this rotation disorder is that under polarized light microscopy, the iridescence of the ferroelectric phase due to birefringence vanishes upon changing to the paraelectric phase, which is clearly confirmed for our MNP3 crystal (supplementary information Fig. S11).

Temperature-dependent dielectric permittivity was also used to detect the phase transition of the crystals as a ferroelectric-to paraelectric transition will result in an anomaly in the dielectric permittivity. Both MNP3 and D-MNP3 show an abrupt increase of dielectric permittivity around their phase transition temperature (Fig. S12 a and b). We also fitted the inverse of dielectric constant with the temperature at 100 kHz (Fig. S12 c and d) by Curie-Weiss law that is, \(\varepsilon =C/(T-{T}_{0})\) where \(\varepsilon\) is the dielectric constant, C is the Curie constant, T is the temperature, and T0 is the Curie temperature. Linear fitting of \(1/\varepsilon\) with the temperature at 100 kHz allows the Curie temperatures of MNP3 and D-MNP3 to be determined as 307 K and 316 K, and the Curie constants as 254 K and 382 K, respectively. The dielectric permittivity around the phase transition temperature is strongly dependent on frequency. A smaller dielectric permittivity at a higher frequency means the ferroelectric-to-paraelectric transition is the order-disorder type36. D-MNP3 shows a much higher dielectric permittivity than MNP3, which may be explained by the higher polarization of the former arising from its stronger H-bonding37.

The polarization (P) versus electric field (E) measurement as shown in Fig. 4a exhibits a hysteresis loop characteristic of a ferroelectric crystal. The remnant polarization Pr is 5.7 μC/cm2, and the coercive field Ec ranges from 54 to 110 kV/cm depending on the applied frequency. The coercive field of MNP3 is much higher than isostructural organic perovskite MNI3 (12 kV/cm), this is in line with the expectation that a stronger hydrogen bond leads to a higher coercive field. Compared with the relatively low coercive field of BTO (10 kV/cm) and PZT (76 kV/cm)38, and the exceedingly large field of 500 kV/cm for PVDF39, MNP3 exhibits a mid-range coercive field that translates to a threshold voltage of 1 V for a 90 nm thick film, which is compatible with the standard operating condition of thin-film ferroelectric random-access memory (FeRAM) devices. Furthermore, the coercive field of D-MNP3 is increased to 138 kV/cm at 50 Hz, this agrees with the fact that D-MNP3 has stronger hydrogen bonds. These results prove that we can control the coercive field of the ferroelectric materials by tuning the intermolecular hydrogen-bond strength. The smaller polarization value of D-MNP3 (4.6 μC/cm2) than MNP3 (5.7 μC/cm2) at 5 Hz can be attributed to two factors: first, D-MNP3 (482.09 Å3) has a larger unit cell volume than MNP3 (477.13 Å3); secondly, the A-site cation displacement of D-MNP3 (72.56 pm) is shorter than MNP3 (73.56 pm) due to the Ubbelohde effect. The polarization values at 5 Hz differ by about 23%, but it increases to 84% at 50 Hz where the polarization value of D-MNP3 is 2.7 μC/cm2 and MNP3 is 5.0 μC/cm2. Considering that the decline in polarization at high frequency is due to the inability of the switching dipoles to keep up with the oscillating field, this effect is more apparent in D-MNP3 due to its stronger hydrogen bond. (Fig. S13).

Fig. 4: Polarization-electric field (P-E) loop and ferroelectric domains.
figure 4

a P-E loop of MNP3 and D-MNP3 at 50, and 5 Hz. PFM phase images of b in-plane and d out-of-plane polarization. PFM amplitude images of c in-plane and e out-of-plane polarization. PFM phase (f) and amplitude (g) switching curve from MNP3 thin film.

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According to the 432F3 transition, MNP3 has four polar axis and eight polarization directions40. Since the polar axes are in between the in-plane (y-axis) and out-of-plane (z-axis) directions, ferroelectric domains whose polarization directions are not perpendicular to these can be observed in both in-plane and out-of-plane phase and amplitude images using piezoresponse force microscopy (PFM). In Fig. 4b, a 180° in-plane domain change can be seen clearly, which indicates two polarized domains in the y-axis direction (+y and −y). In the amplitude scan of the in-plane PFM image (Fig. 4c), there is a region that overlaps with some parts of the out-of-plane phase (Fig. 4d), which is indicative of canted polarizations with vectors in both −y and −z directions (Fig. S14 b green area). We could also identify areas (red in Fig. S14) with canted polarizations and vector components in −y and −z direction, and another two regions with canted polarizations and vector components along +y and −y region according to the in-plane phase image. Therefore, from the PFM images, we can observe four distinct polarization domains (Fig. S14 b and c), similar to the previously reported isostructural MNI314. Due to the high coercive field of MNP3, it was difficult to perform phase switching on a thick, single-crystal sample using PFM. Instead, ferroelectric switching was performed on a 1 μm-thick thin film of MNP3 prepared by a spray-coating method (Fig. S15). A clear 180-degree phase reversal (Fig. 4f) and butterfly curve (Fig. 4g) were observed by sweeping the tip voltage from −10 to 10 V, indicating that ferroelectric domains can be electrically switched.

To validate our experimentally measured polarization values and coercive field, we performed density functional theory (DFT) calculations of both MNI3 and MNP3. First of all, spontaneous polarization value was calculated by Berry phase calculation developed by King-Smith and Vanderbilt41,42. We constructed a √2 × √2 × 1 supercell based on the unit cell of the ferroelectric phase, and built the centrosymmetric reference phase via rotation, displacement and distortion of the components, after which two MDABCO molecules in the simulation cell are aligned antiparallel to each other (Fig. 5a). The calculated polarization of MNP3 is shown in Fig. 5b, where it varies continuously from 0 to 6.5 μC/cm2 along the dynamic path, thus the spontaneous polarization agrees well with the experimental value of 5.7 μC/cm2. Although the larger displacement of MDABCO in MNP3 than MNI3 (Fig. 2d) gives a larger dipole moment in the former, the larger unit cell volume of MNP3 than MNI3 offsets the increased polarization as the overall polarization is defined by the sum of dipole moment per unit volume. Furthermore, in MNI3, there is a large polarization contribution from the off-centre displacement of NH4+ at B site (δB = 0.6366 pm) because of its larger rB/rX ratio (0.664) compared to the typical octahedron ratio (0.414–0.592) predicted by Pauling’s rule. On the other hand, the NH4+ ion in MNP3 is located at the centre of the octahedron (δB = 0.0312 pm) with a rB/rX ratio of 0.570, thus there is no off-centre displacement.

Fig. 5: Berry phase calculation of spontaneous polarization in MNP3.
figure 5

a The dynamic path of ferroelectric phase transition in MNP3. b The polarization value of MNP3 was obtained by Berry phase calculations. The path connects the centrosymmetric reference phase (λ = 0) to the ferroelectric phase (λ = 1). c Electrostatic potential of MNI3 and MNP3, and the Mulliken charges of H atoms participating in hydrogen bonds.

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The displacement of MDABCO in the polarization direction with respect to the MNP3 framework is the primary driver of the ferroelectricity of MNP3. The large electrostatic potential differences between MDABCO-NH and PF6 increase the hydrogen-bond strength compared to MNI3. Mulliken charge calculations for MNP3 and MNI3 reveal that PF6− has a more pronounced inductive effect than I− on the H atom in MDABCO (Fig. 5c), thus the stronger N–H…F hydrogen retards the polarity reversal process of organic ferroelectric perovskites. Therefore, any reinforcement of hydrogen bonding will hinder the ferroelectric switching process, since multiple hydrogen bonds need to be successively broken and established during the rotation of MDABCO, as illustrated in Fig. 5a. This is the root cause of the enhanced coercive field for MNP3.

In conclusion, incorporating stronger hydrogen bonds in MDABCO-NH4-I3 by the substitution of PF6− for I− increase the coercive field by an order of magnitude from ~12 kV/cm to 110 kV/cm. This enables a threshold voltage of 1 V for a 90 nm thick film, which complies with the miniaturization requirement of FeRAM devices. Both experimental results and DFT calculations provide robust evidence for the mechanism of MDABCO rotation upon ferroelectric switching and phase transition, the dynamics of which are modulated by the intermolecular hydrogen bonding interactions between MDABCO2+ and PF6−. This work demonstrates that regulating intermolecular interactions between the cation and anion can be used for the engineering of the coercive field in organic ferroelectric materials.

Methods

Synthesis

Detailed synthetic methods for preparing precursors are provided in Supplementary Information. For growing MNP3 crystal, 209 mg (0.5 mmol) of MDABCO-(PF6)2 and 81.5 mg (0.5 mmol) of NH4PF6 were dissolved in acetonitrile until saturated, and the parallelogrammic single-crystals were grown after slow evaporation.

Characterization

All chemicals were purchased from Sigma–Aldrich without further purification. 1H and 13C nuclear magnetic resonance (NMR) was taken by AVII 400 MHz NMR spectrometer of Bruker. Thermogravimetric analyses (TGA) were performed under a nitrogen atmosphere with a heating rate of 10 °C/min using a TA Instruments Trios V3.1 thermogravimetric analyzer. Differential scanning calorimetry (DSC) scans were performed under a nitrogen atmosphere with a heating rate of 10 °C/min using Mettler-Toledo DSC. The dielectric constant was measured by CVU unit in Keithley-SCS4200 with the pelleted sample. Powder X-ray diffraction (PXRD) patterns were recorded on a Bruker D8 Focus Powder X-ray diffractometer using Cu Kα radiation (40 kV, 40 mA) at room temperature. Ferroelectric P-E curve was measured with Precision Multiferroic II Ferroelectric Test System of the Radiant Technologies with high voltage amplifier. Piezoresponse force microscopy (PFM) tests were performed on Bruker Dimension Icon Atomic Force Microscope with grown crystals or spray-coated samples on ITO.

Density functional theory (DFT) calculations

We performed the Berry phase calculations41,42 within the DFT framework as implemented in the Vienna ab initio simulation package (VASP)43,44. The exchange-correlation interactions were treated within the Perdew-Burke-Ernzerh (PBE) generalized gradient approximation45. To complement the deficiencies of DFT in treating dispersion interactions, the third-generation (D3) van der Waals corrections proposed by Grimme46 were employed. The plane-wave cutoff energy was set to 520 eV, and the k-point mesh to 3 × 3 × 4. The polarization was calculated using a supercell twice the size of the unit cell so that a centrosymmetric reference phase can be constructed. A convergence threshold of 0.01 eV/Å in force was reached in structural optimization. Electrostatic potential and Mulliken charge were calculated by DMol347,48 code in Materials Studio using a double numerical polarized basis set and PBE45 exchange-correlation functional.

Reporting summary

Further information on research design is available in the Nature Research Reporting Summary linked to this article.