Liberating a hidden antiferroelectric phase with interfacial electrostatic engineering

Julia A Mundy^{1

2}, Bastien F Grosso³, Colin A Heikes⁴, Dan Ferenc Segedin^{2

5}, Zhe Wang⁶, Yu-Tsun Shao⁶, Cheng Dai⁷, Berit H Goodge^{6

8}, Quintin N Meier³, Christopher T Nelson⁹, Bhagwati Prasad¹, Fei Xue⁷, Steffen Ganschow¹⁰, David A Muller^{6

8}, Lena F Kourkoutis^{6

8}, Long-Qing Chen⁷, William D Ratcliff^{4

11}, Nicola A Spaldin³, Ramamoorthy Ramesh^{1

5

12}, Darrell G Schlom^{8

10

13}

Affiliations

¹ Department of Materials Science and Engineering, University of California, Berkeley, Berkeley, CA 94720, USA.
² Department of Physics, Harvard University, Cambridge, MA 02138, USA.
³ Department of Materials, ETH Zürich, Zürich CH-8093, Switzerland.
⁴ NIST Center for Neutron Research, National Institute of Standards and Technology, Gaithersburg, MD 20878, USA.
⁵ Department of Physics, University of California, Berkeley, Berkeley, CA 94720, USA.
⁶ School of Applied and Engineering Physics, Cornell University, Ithaca, NY 14853, USA.
⁷ Department of Materials Science and Engineering, The Pennsylvania State University, University Park, PA 16802, USA.
⁸ Kavli Institute at Cornell for Nanoscale Science, Ithaca, NY 14853, USA.
⁹ Materials Science and Technology Division, Oak Ridge National Laboratory, Oak Ridge, TN 37830, USA.
¹⁰ Leibniz-Institut für Kristallzüchtung, 12489 Berlin, Germany.
¹¹ Department of Materials Science and Engineering, University of Maryland, College Park, MD 20742, USA.
¹² Materials Science Division, Lawrence Berkeley National Laboratory, Berkeley, CA 94720, USA.
¹³ Department of Materials Science and Engineering, Cornell University, Ithaca, NY 14853, USA.

PMID: 35108054
PMCID: PMC8809685
DOI: 10.1126/sciadv.abg5860

Liberating a hidden antiferroelectric phase with interfacial electrostatic engineering

Julia A Mundy et al. Sci Adv. 2022.

. 2022 Feb 4;8(5):eabg5860.

doi: 10.1126/sciadv.abg5860. Epub 2022 Feb 2.

Authors

Affiliations

¹ Department of Materials Science and Engineering, University of California, Berkeley, Berkeley, CA 94720, USA.
² Department of Physics, Harvard University, Cambridge, MA 02138, USA.
³ Department of Materials, ETH Zürich, Zürich CH-8093, Switzerland.
⁴ NIST Center for Neutron Research, National Institute of Standards and Technology, Gaithersburg, MD 20878, USA.
⁵ Department of Physics, University of California, Berkeley, Berkeley, CA 94720, USA.
⁶ School of Applied and Engineering Physics, Cornell University, Ithaca, NY 14853, USA.
⁷ Department of Materials Science and Engineering, The Pennsylvania State University, University Park, PA 16802, USA.
⁸ Kavli Institute at Cornell for Nanoscale Science, Ithaca, NY 14853, USA.
⁹ Materials Science and Technology Division, Oak Ridge National Laboratory, Oak Ridge, TN 37830, USA.
¹⁰ Leibniz-Institut für Kristallzüchtung, 12489 Berlin, Germany.
¹¹ Department of Materials Science and Engineering, University of Maryland, College Park, MD 20742, USA.
¹² Materials Science Division, Lawrence Berkeley National Laboratory, Berkeley, CA 94720, USA.
¹³ Department of Materials Science and Engineering, Cornell University, Ithaca, NY 14853, USA.

PMID: 35108054
PMCID: PMC8809685
DOI: 10.1126/sciadv.abg5860

Abstract

Antiferroelectric materials have seen a resurgence of interest because of proposed applications in a number of energy-efficient technologies. Unfortunately, relatively few families of antiferroelectric materials have been identified, precluding many proposed applications. Here, we propose a design strategy for the construction of antiferroelectric materials using interfacial electrostatic engineering. We begin with a ferroelectric material with one of the highest known bulk polarizations, BiFeO₃. By confining thin layers of BiFeO₃ in a dielectric matrix, we show that a metastable antiferroelectric structure can be induced. Application of an electric field reversibly switches between this new phase and a ferroelectric state. The use of electrostatic confinement provides an untapped pathway for the design of engineered antiferroelectric materials with large and potentially coupled responses.

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Figures

**Fig. 1.. Energetics of BiFeO₃ ground states.**
(A) Structure and relative energy of the R3c ferroelectric ground state, antiferroelectric “*Pnma*-AFE” and *Pbam* states, and nonpolar *Pnma* state. Energies are given for structures with the in-plane lattice constants of the R3c phase. (B) Decomposition of the structure of the *Pnma*-AFE phase into its main distortions (of symmetries Σ, Λ, R, S, and T) relative to the ideal cubic $Pm \bar{3} m$ perovskite. Bismuth, iron, and oxygen are shown in red, blue, and orange, respectively. (iii) Double tilts where a pair of identical clockwise octahedral rotations is followed by a pair of identical counterclockwise rotations (T). Note that only one octahedron is visible per pair, as the other one is perfectly aligned. (C) Calculated energy as a function of in-plane lattice constant. The calculated dependence on the energy for the one polar and five nonpolar structures identified in our calculations. The calculated ground state in-plane pseudocubic lattice constant of R3c is 3.945 Å. The new *Pnma*-AFE structure is the lowest-energy nonpolar structure above 3.9 Å.

**Fig. 2.. Phase stability of BiFeO₃ heterostructures.**
(A) Phase stability when thin films of BiFeO₃ are confined between dielectric layers. The stability of the polar R3c phase (colored volume) in comparison to a nonpolar alternative (uncolored void), is plotted as a function of BiFeO₃ layer thickness, strain, and the dielectric constant of the confining dielectric. The phase stability can be tuned by the superlattice construction to form the *Pnma*-AFE phase. At a thickness of 60 Å (gray plane intersecting the polar volume) and 0% strain, the nonpolar-to-polar transition is predicted to occur in the DFT calculations at a surrounding dielectric constant of 273 (the intersection is projected on the bottom of the plot by a vertical dashed line at the center of the visible gray plane). Additional dashed lines are displayed to show the values on the three axes. (B) Phase-field simulations for an unstrained 60-Å-thick BiFeO₃ layer surrounded by a nonpolar material with variable dielectric constant (38). The polar phase stays stable for dielectric constants larger than 293, and a mixed state between R3c and *Pnma*-AFE is predicted to exist for surrounding dielectric constants ranging from 293 to 130; at lower dielectric constants, the *Pnma*-AFE phase is stable.

**Fig. 3.. Atomic-scale characterization of the *Pnma*-AFE phase of BiFeO₃ in confined (La_0.4Bi_0.6FeO₃)_n/(BiFeO₃)_n superlattices.**
(A) Annular dark-field (ADF) and EELS spectroscopic imaging showing the atomic concentrations of bismuth, iron, and lanthanum in red, blue, and green, respectively. (B) HAADF-STEM image of the (La_0.4Bi_0.6FeO₃)_n/(BiFeO₃)_n superlattice region shown in (A). The atomic-scale displacements in (B) are calculated showing an up-up/down-down picometer-scale distortion with a 45° axis. (C) Imaging of the sample shown in (B) using scanning diffraction measurements. The displacement of the electron beam due to the Lorentz force is shown, directly imaging the alternating electrical dipoles in the newly formed antiferroelectric. (D to F) Averaged images from the BiFeO₃ (BFO) layers from the sample in (B) along various crystallographic zone axes. The corresponding orientation of the *Pnma*-AFE unit cell is shown on each image with bismuth in red and iron in blue.

**Fig. 4.. DF-TEM images of (La_xBi_1−xFeO₃)₁₅/(BiFeO₃)₁₅ superlattices demonstrating structural tunability with altered properties of the surrounding dielectric layer.**
(A) The (La_0.4Bi_0.6FeO₃)₁₅/(BiFeO₃)₁₅ superlattice imaged in Fig. 3 (B and E) showing a coherent region of the [001]_pc-oriented *Pnma*-AFE polymorph of BiFeO₃. (B to D) A (La_0.3Bi_0.7FeO₃)₁₅/(BiFeO₃)₁₅ superlattice showing phase coexistence between the *Pnma*-AFE and R3c polymorphs of BiFeO₃. Scale bar, 15 nm. (E) Results of our phase-field model of the (La_0.3Bi_0.7FeO₃)₁₅/(BiFeO₃)₁₅ superlattice simulating the coexistence of ferroelectric (FE) and antiferroelectric (AFE) domains, as observed.

**Fig. 5.. Electric field–induced switching between the *Pnma*-AFE state and the ferroelectric state of BiFeO₃.**
(A) Application of an electric field alters the relative stability of the R3c and *Pnma*-AFE phases in heterostructures. For a given layer thickness/dielectric, the ferroelectric R3c phase can become more stable than the antiferroelectric *Pnma*-AFE phase (plotted for 15 f.u.). (B) Barrier for switching between the *Pnma*-AFE and R3c phases calculated using the nudged elastic band method. There is a 26 meV/f.u. activation barrier to switch between the phases. The reaction pathway is shown, tracing the damping of the Σ and Λ antipolar distortion modes of *Pnma*-AFE and increase of the polar Γ mode of R3c structure. (C to E) Polarization-electric field hysteresis loops for (La_xBi_1−xFeO₃)₁₅/(BiFeO₃)₁₅ superlattices with x = 0.5, 0.3, and 0.2, respectively. Tuning the lanthanum concentration in the dielectric layer converts the structure from an antiferroelectric phase as shown in (C) and (D) to a ferroelectric structure in (E).

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References

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Liberating a hidden antiferroelectric phase with interfacial electrostatic engineering

Affiliations

Liberating a hidden antiferroelectric phase with interfacial electrostatic engineering

Authors

Affiliations

Abstract

Figures

References

LinkOut - more resources

Full Text Sources