Topotactic Reduction toward a Noncentrosymmetric Deficient

Oct 30, 2017 - Topotactic Reduction toward a Noncentrosymmetric Deficient Perovskite Tb0.50Ca0.50Mn0.96O2.37 with Ordered Mn Vacancies and Piezoelectr...
8 downloads 4 Views 6MB Size
Download PDF
Article Cite This: Chem. Mater. 2017, 29, 9840-9850

pubs.acs.org/cm

Topotactic Reduction toward a Noncentrosymmetric Deficient Perovskite Tb0.50Ca0.50Mn0.96O2.37 with Ordered Mn Vacancies and Piezoelectric Behavior Hao Zhang,†,‡ Song Gao,‡ Qinghua Zhang,§ Jingen Wu,‡ Jie Liang,⊥ Cheng Dong,§ Lin Gu,§ Shuxiang Dong,‡ Junliang Sun,† Fuhui Liao,† Jianhua Lin,† Ruqiang Zou,*,‡ and Guobao Li*,† †

College of Chemistry and Molecular Engineering, Peking University, Beijing 100871, P. R. China Beijing 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 § Beijing National Laboratory for Condensed Matter Physics, Institute of Physics, Chinese Academy of Sciences, Beijing 100190, P. R. China ⊥ School of Space and Environment, Beihang University, Beijing 100191, China ‡

S Supporting Information *

ABSTRACT: Low-temperature reduction of perovskite Tb0.5Ca0.5MnO3−x yields novel crystal-structured noncentrosymmetric compound Tb0.50Ca0.50Mn0.962.33+O2.37, which unusually crystallizes in cubic lattice I23 (a ∼ 15.27 Å) based on a 4ap × 4ap × 4ap expansion relative to the simple cubic perovskite unit cell. Rietveld refinements and HAADF-STEM images are used for the structure determination, revealing a rare-typed metal-anion coordination framework which consists of corner-shared tetrahedra and pyramids, and edge-shared bipyramids and octahedra. 2/64 B-site Mn-ordered vacancies are observed for the first time acting as the apex and body center of the I lattice in reduced systems. Room-temperature piezoelectricity is detected, with a quasistatic d33 value of ∼0.32 pC N−1 and inverse d33 value of ∼10.5 pm V−1. This phase primarily exhibits antiferromagnetic ordering below TN ∼ 70 K, with ferromagnetic responses resulted from spin-canting below 40 K. This work provides a new way toward synthesizing unconventional acentric materials, in the absence of second-order Jahn−Teller active “distortion centers”.

■

INTRODUCTION

centering displacements in their coordination polyhedra, as illustrated by second-order Jahn−Teller (SOJT) distortion.10−16 However, the drawback is that these symmetrybreaking cations are not compatible with d-spin behaviors, which are the basis of magnetic or spintronic devices because dspin behaviors require metal ions to have unpaired d electrons coupled with each other through exchange interaction. To achieve the coexistence of NCS structure and d-spin related properties, many approaches have been made, such as introducing “long-pair” ns2 cations (Pb2+, Bi3+),17−22 creating charge order,23,24 manipulating spiral spin arrangements,25,26 or more recently utilizing cooperative tilting distortion.27−29 Compared with various approaches, low-temperature topotactic reactions,30 which specifically target at preparing new structural compounds, is another promising way toward new NCS compounds especially with d-spin behavior. In the topotactic reaction, the oxygen ions could be extracted from

Searching for new crystal-structured solid-state materials has attracted much interest since a new crystal structure could intrinsically lead to undiscovered physical and electronic properties. These new materials, if crystallizing in a noncentrosymmetric (NCS) structure, would be of particular significance in fundamental research of SHG (second harmonic generation) devices, piezo-/ferroelectrics, multiferroics, and even catalysis.1−9 However, the discovery and preparation of new NCS compounds remain a challenge. First, the conventional high-temperature synthesis route favors only thermostable phases that prevents us from designing or controlling the structural features of final products; second, even if new structural features are achieved, the close packing of cation and anion strongly favors the formation of a highly symmetric arrangement in which the unlike-charge electronic attraction could be maximized, and the like-charge electronic repulsion could be minimized. The common strategy toward noncentrosymmetry by hightemperature fabrication is to introduce d0 transitional-metal cations (Ti4+, Zr4+, Nb5+, W6+) which would undergo off© 2017 American Chemical Society

Received: September 28, 2017 Revised: October 30, 2017 Published: October 30, 2017 9840

DOI: 10.1021/acs.chemmater.7b04115 Chem. Mater. 2017, 29, 9840−9850

Article

Chemistry of Materials

topotactic reaction can be a powerful technique in searching for unconventional NCS materials with a new crystal structure and d-spin-related properties.

or intercalated into specific crystallographic sites in precursors, leading to usual topological metal-anion coordination frameworks. One advantage of this strategy is to enable d-spin metalanion polyhedra to distort into acentric ones during structural rearrangements. This allows us to get NCS properties directly in dn compounds, without incorporating any d0 cations. In addition, topotactic reactions would further tune the d-spin/ state of transitional-metal centers which interact with each other through extended metal-anion coordination frameworks to give rise to a new electronic structure and related properties. A lot of excellent work has been made in the topotactic field,31−38 but only few of them achieved noncentrosymmetry in the lattice. For reduced Ba4CaFe3O9.5,39 its complex cationordered structure has a low-symmetry arrangement of CaO6, FeO5, and FeO4 polyhedra on the B-cation sites that breaks the structural inversion symmetry and leads to SHG activity. In oxidized Ba2YFeO5.5,40 some of the centric FeO4 tetrahedra are oxidized to acentric FeO5 pyramids that bring polar behavior to the whole lattice, leading to a pyroelectric property. In both cases, the major driving force for symmetry-breaking distortion is attributed to the lattice strain generated by cations with different ionic radii. One reason that most topochemical phases fail to achieve noncentrosymmetry may be the fact that only alkaline metals (Ca2+, Sr2+, Ba2+) and transitional metals are primarily concerned during materials preparation, yet lanthanide, especially post-lanthanide metals, was rarely considered. Although different B-site metal cations in host perovskites can intrinsically lead to different metal-anion frameworks, the A-site cations are of equal importance in determining the final crystal structure. Lanthanide cations, when partially occupying the Asite of the lattice, would considerably affect the whole metalanion framework in reduced structure since they require a higher coordination number than alkaline metal cations. One example is the reduced La1−xCaxMnO2+δ,41 in which its lattice has to adopt chainlike MnO4 tetrahedral layers to satisfy the coordination requirements of A-site La3+ cations. This also inspires us to utilize post-lanthanum (Tb in this work) to search for new NCS compounds. In this report, a new structured noncentrosymmetric compound Tb0.496(2)3+Ca0.504(10)2+Mn0.963(14) 2.334(3)+O2.372(1) (nominal Tb0.50Ca0.50Mn0.962.33+O2.37) with d-spin magnetic properties was successfully prepared from host perovskite Tb0.5Ca0.5MnO3 by low-temperature reduction method. The obtained Tb0.50Ca0.50Mn0.96O2.37 exhibits a 4ap × 4ap × 4ap lattice expansion on the basis of a simple cubic perovskite unit cell, crystallizing in a noncentrosymmetric space group I23, a ∼ 15.27 Å. The metal-anion coordination frameworks in the structure evolve in a very unusual way that the original cornershared MnO6 octahedra in host perovskite are reduced into tetrahedra, pyramids, bipyramids, and edged-shared octahedra. Ordered vacancies of Mn cations are unexpectedly observed on the apical sites and body-centered sites of the lattice. To our knowledge, this is the first time to observe an ordered Mn vacancy in deficient perovskite systems. Room-temperature piezoelectricity was successfully measured, with a quasi-static d33 value of ∼0.32 pC N−1 and inverse d33 value of ∼10.5 pm V−1. When made into multilayered pellets, the entire pellet reached 3.5 pC N−1 that could even match with α-SiO2 of d11 ∼ 2.31 pC N−1. The mechanism of piezoelectricity is also discussed. At low temperature, the phase exhibits antiferromagnetic ordering below TN ∼ 70 K, with increasing ferromagnetic responses below 40 K, which results from spincanting behavior. This work proves that the low-temperature

■

EXPERIMENTAL SECTION

Preparation of Tb0.5Ca0.5MnO3 perovskites as Precursor. Perovskite Tb0.5Ca0.5MnO3 was synthesized by citric method. Appropriate stoichiometric ratios of Tb4O7 (AR), CaCO3 (AR 99.5%), and MnCO3 (AR 99.5%) were first dissolved separately in nitric acid and then mixed together. About 2 equivmol of citric acid was then added. The solution was dried while stirring to form gel, in which all the ions were well-distributed. Later, this prepared gel was calcined under 1300 °C for 48 h with 3 cycles of intermittent regrinding. X-ray powder diffraction data and Rietveld refinement proved that the resulting powder was pure and well-crystalline. Synthesis of Tb0.50Ca0.50Mn0.96O2.37 by Low-Temperature Reduction. Reduction was performed using NaH (>95%) as a solid-state reducing agent.41 A 5 g portion of Tb0.5Ca0.5MnO3 phase was thoroughly ground with 2 equivmol of NaH in an agate mortar in an argon-filled glovebox (O2 and H2O < 0.3 ppm). The resulting mixture was then sealed in silica tubes under vacuum and heated at 190 °C for 48 h. The sample was then taken out, reground in the glovebox, and resealed in tubes before being heated for 7 further periods of 48 h at 210 °C. Finally, the sample was washed with methanol several times under a nitrogen atmosphere to remove sodium-containing phases (NaOH and NaH) and then dried under vacuum. Characterization. High-resolution powder X-ray diffraction (HRXRD) data were collected on a PANalytical Empyrean diffractometer with Cu Kα1 (λ = 1.5407 Å) radiation (2θ range, 5− 120°; step, 0.013°; scan speed, 6 s step−1) at 50 kV and 40 mA. TOF neutron powder diffraction data were collected on a WISH diffractometer (ISIS neutron source). Rietveld refinements of XRD and neutron diffraction data were performed by Fullprof software.42 Selected-area electron diffraction (SAED) was carried out on a JEM 2100F instrument with an accelerating voltage of 200 kV. The RED (rotation electron diffraction) data collection and processing were performed using the RED data collection and processing software, respectively, and the three-dimensional reciprocal lattice was reconstructed from obtained SAED frames visualized by the RED data processing software,43 from which the unit-cell parameters were determined. The reflection conditions were deduced, especially from two-dimensional main-zone slices cut from the reconstructed threedimensional reciprocal lattice. The diffraction intensities were extracted but cannot be used for single-crystal structural determination because the intensities I(hkl) used as |F(hkl)|2 were incorrect due to electron energy loss. High-angle annular dark field scanning transmission electron microscopy (HAADF-STEM) images were collected on an ARM-200CF (JEOL, Tokyo, Japan) instrument operated at 200 keV and equipped with double spherical aberration (Cs) correctors. The attainable resolution of the probe defined by the objective prefield is 78 pm. The ABF (angular bright-field) images were also collected on this TEM. To facilitate the observation we inverted the ABF contrast, and atomic columns are now shown in white spots. The stoichiometry of the sample was confirmed by inductively coupled plasma-atomic emission spectrometer (ICP, Prodigy 7) and iodometry titration. TGA and DSC measurements were performed on a NETZSCH STA449C instrument. Magnetic properties were measured on a SQUID MPMS device. For electric properties, since the reduced sample cannot be calcined under high temperature, a small quantity of epoxy resin was added into the powder (about 1 drop of epoxy for ∼0.2 g of sample) and mixed thoroughly to press into Tb0.50Ca0.50Mn0.962.33+O2.37/polymer pellets (thickness, ∼0.3 mm; diameter, ∼5 mm; area, ∼0.2 cm2). The epoxy resin is not piezoelectric so it will not disturb measurements. These pellets were further kept in air for 24 h for solidification. After that, these hardened pellets were covered up and down with silver paste and Pt wires as electrodes for electric measurements. Piezoelectric and ferroelectric measurements were carried out on a quasi-static d33 meter 9841

DOI: 10.1021/acs.chemmater.7b04115 Chem. Mater. 2017, 29, 9840−9850

Article

Chemistry of Materials (ZJ-3D; Institute of Acoustics, Beijing, China.) and TF analyzer 2000. Dielectric and pyroelectric measurements were carried out on a SQUID PPMS with WK 6500B and KEITHLEY 6517A devices, respectively.

relationship with its parent phase. As shown in section 6 in the SI, the Tb0.50Ca0.50Mn0.96O2.37 is stable at room temperature, but will begin reoxidation above 150 °C. At 480 °C, the sample will be finally reoxidized back to its parent phase Tb0.50Ca0.50Mn0.96O2.95 with perovskite structure, though with Mn vacancy. This facile reoxidation suggests a topochemical conversion; that is to say, the reduced new phase Tb0.50Ca0.50Mn0.96O2.37 should be structurally related to the perovskite arrangements. The DSC and related XRD measurements (see section 7 in the SI) under Ar atmosphere reveal that the sample can sustain up to 600 °C. However, at 820 °C, the sample will totally decompose to CaMnO244,45 and Tb2O3. Crystal Structure Determination. At the beginning, the HRXRD pattern indexing suggests an I-, P-, or even R-unit cell with various possible space groups and cell lengths. Therefore, TEM RED (rotation electron diffraction) techniques are specially used to identify its unit cell and space group (Figure 1a−c.), with the help of TEM SAED analysis (Figure 1d,e.).

■

RESULTS Preparation of Tb0.50Ca0.50Mn0.96O2.37. During sample preparation, the precursor Tb0.5Ca0.5MnO3 was reduced repeatedly, and after each cycle of reduction reaction, the resulting mixture was taken out for XRD collection. These XRD patterns are provided in Figure S1 in the Supporting Information. As the reaction progresses, the reflection peaks of the precursor are gradually decreasing while a new set of reflection peaks are becoming more and more prominent, implying new crystal structure features in the obtained phase. This reduction reaction at 210 °C was further repeated several times until no variation in reflection peaks was observed in the XRD pattern. The XRD pattern of our final product is provided in Figure S2 in the SI. With the help of TEM techniques (discussed later in the next section), all reflection peaks can be indexed well by cubic I lattice (a ∼ 15.27 Å), and no impurities are found. The product has a large difference in XRD pattern compared with its perovskite precursor. First, the first reflection peak has considerably shifted from 2θ ∼ 20.4°, d(011) ∼ 4.3 Å (of precursor), to 2θ ∼ 8.1°, d ∼ 10.8 Å (of product), implying large lattice expansion in the phase. Second, during reduction, the precursor’s strongest peak (121) (d ∼ 2.67 Å) gradually merges with its adjacent (200) and (002) peaks to yield the strongest peaks (044)/(404)/(440) (d ∼ 2.7 Å) of the product, demonstrating that the new lattice adopts higher symmetry and may still retain a perovskite topological framework. The stoichiometry of elements Tb, Ca, and Mn of the product was obtained by the inductive coupling plasma (ICP) measurements, while the total oxygen contents were confirmed by iodometry titration. The valence of Tb was also confirmed to be 3+ by XPS.56,57 The detailed information on these measurements is provided in section 3 of the SI. The oxygen stoichiometries derived from these analysis data indicate an average manganese oxidation state of Mn2.334(3)+. Therefore, the chemical formula is Tb0.496(2)3+Ca0.504(10)2+Mn0.963(14) 2.334(3)+O2.372(1), that is, nominal Tb 0.50 Ca 0.50Mn0.962.33+O2.37. Our Tb0.50Ca0.50Mn0.96O2.37 is a unique and final phase that can be obtained under the current reaction condition. This is because of the following: First, further reduction at 210 °C will cause no difference in XRD pattern, proving that the current phase is stable and that no more structural changes would happen under the current condition. Second, if we raise the temperature to 230 °C, Tb0.50Ca0.50Mn0.96O2.37 will directly decompose to rock-salt type CaMnO244,45 and Tb1−xCaxMnO2+δ, which exhibits a structure similar to that of reported La1−xCaxMnO2+δ41 (see section 4 in the SI). Third, although the precursor Tb1−xCaxMnO3 is actually a solid solution in the entire range of 0 ≤ x ≤ 1, the ratio of Tb/Ca needs to be approximately 1:1 to obtain this new phase. Otherwise, increasing the Ca components in the precursor (such as Tb0.3Ca0.7MnO3) will yield impurity CaMnO2 (see section 5 in the SI) after the reduction, while increasing Tb components (such as Tb0.6Ca0.4MnO3) will directly hinder the reduction. TGA measurements were conducted in air to specifically confirm the oxygen deficiency of the product and its phase

Figure 1. TEM data. (a−c) Rotation electron diffraction (RED) data. (d, e) Selected area electron diffraction (SAED) data of the Tb0.50Ca0.50Mn0.96O2.37.

Several possible types of lattice could be derived as cubic Im3̅m, Im3,̅ and I23; tetragonal I4/mmm; and orthorhombic I222 and Immm, with cell length in all three dimensions around 15.27 Å. Both the HRXRD and TEM data could be well indexed on all these unit cells. Due to I-lattice systematic extinction, the reflection (100) with d ∼ 15.3 Å is missing, while the reflection (011) [or (101), (110)] takes its place as the first observable peak with large d-spacing ∼ 10.8 Å. Since there is no observable peak separation in the HRXRD profile especially at high diffraction angle, the cubic I lattice is preferred and is attempted for later structural determination. It is clear in SAED data (Figure 1d,e.) that a series of reflections (044), (404), and (440) are very distinctive. The cell length of ∼15.27 Å is just approximately 4 times longer than a cubic perovskite unit cell (a ∼ 3.8 Å). All of these observations imply that, during reduction, the Tb0.50Ca0.50Mn0.96O2.37 undergoes 4-fold expansion in all crystallographic a-, b-, and c-axes on the basis of a cubic perovskite unit cell. On the basis of the above, we initially built a 4-fold expanded structural model with I23, a ∼ 15.27 Å (later we will discuss the other space group) that contains 64 cubic perovskite units, but 9842

DOI: 10.1021/acs.chemmater.7b04115 Chem. Mater. 2017, 29, 9840−9850

Article

Chemistry of Materials

Figure 2. Structure evolution from low-temperature reduction. (a) Structure of orthorhombic distorted Tb0.5Ca0.5MnO3 as precursor and simple cubic perovskite unit cell. (b) After reduction, heavy cations’ arrangements of Tb0.50Ca0.50Mn0.96O2.37. (c) Final crystal structure of Tb0.50Ca0.50Mn0.96O2.37.

apical and body-centered site (marked by green circles) of the lattice. This helical arrangement could result from a 3- or 3̅rotation axis. Carefully scrutinizing the atom information on these images, we found that the symmetry operation m and 2rotation could not be differentiated. However, the diagonal mirror plane m (like in Im3m ̅ ) and 4-rotation axis (like in I4/ mmm) could be precluded because, if they exist, the [111] zone helical arrangements would not be allowed. The space group Im3̅ is centric while its subgroup I23 is acentric, yet our diffraction data did not provide a clear choice between them. However, since room-temperature piezoelectric responses were successfully detected (discussed in next section), the noncentrosymmetric space group I23 was primarily adopted for further structure determination. The second step is to locate oxygen ions to complete the whole structure. Room-temperature XRD and NPD data were jointly refined. The coordinates of heavy ions were fixed, and a difference-Fourier procedure was repeated. Every time after the procedure, at least one or two oxygen ions could be identified, with a residual density of 0.2−0.3. These oxygen ions were then added into the structure and further refined. The differenceFourier procedure was repeated several times until all oxygen ions could be located, and no additional oxygen peaks could be found. After the procedure, all ions’ coordinates were set to refine together. The convergence could still be readily reached, and all coordinate parameters were stable. All obtained ⟨Mn− O⟩ bond lengths in the structure are among a reasonable range 1.78−2.35 Å. The inverted ABF images as shown in Figure 3d could provide information about the oxygen location (marked by red circles). We can see that the projection of oxygen arrangements is in good agreement with the image. However, the strange thing is that no bonding oxygen ions could be found around the apical (0, 0, 0) and also bodycentered (1/2, 1/2, 1/2) Mn cations, as shown in Figures 2b and 3c, marked by a green sphere. At first we thought that the problem is about the symmetry, and we tried again with lower symmetry I3 and even I222, but still no oxygen ions could be found around this Mn ion. Later, when we checked the atoms’ occupancies, we found that the occupancy of Mn at the apical and body-centered sites could be refined to 0.14, but at other sites the Mn occupancies were still approximately 1. Meanwhile, the atom displacement Uiso at the apical and bodycentered sites could also be unreasonably high (more than 60) if refined. This anomaly indicates that the reduction reaction had already extracted the Mn ion away from these specific sites, leaving 2/64 Mn vacancy. (There are totally 64 Mn ions in the unit cell. Apical (0, 0, 0) plus body-centered (1/2, 1/2, 1/2) Mn ions amounts to 2/64 vacancy in one cell.) Since the Mn ions were extracted, it is not surprised to see that there are no bonding oxygen ions around them. To clarify this, we first reset the Mn ion occupancy to 0.14 and refined the coordination

only heavy ions were set, as shown in Figure 2a,b. The Tb/Ca ions are placed at 8c, 24f sites that are equal to the A-site in perovskite, and the Mn ions are placed at 2a, 6b, 12d, 12e, 24f sites that are equal to the B-site. The structure model could also be seen as a 3a × 3b × 2c expansion on the basis of distorted or t h o rh o mb i c pe r ov s k i t e un i t ce l l o f p r ecu r s or Tb0.5Ca0.5MnO3−x.46 (Pbnm; a = 5.334 Å, b = 5.458 Å, c = 7.467 Å). The strategy for structure determination (the final structure is shown in Figure 2c) is introduced here in detail. The first step is to locate all cations in the lattice. XRD data were first used for the cation lattice refinements; after a few cycles of refinements their location could be confirmed, and the convergence could be readily reached (Rwp ∼ 15.2), when space group was Im3̅ or I23. However, when the space group was Im3̅m, the convergence could not be reached, and the cations could not be settled. One possible reason is that the symmetry Im3̅m is too high to describe the structure. HAADFSTEM images were collected to further confirm the cation lattice. As shown in Figure 3a−c, the refined cation structure in

Figure 3. (a−c) HAADF-STEM images of obtained Tb0.50Ca0.50Mn0.96O2.37 projected along [100] and [111]. (d) Inverted ABF images revealing possible oxygen sites (red circle). Helical cation arrangements (white circles) can be seen from parts b and c, with Mn vacancies (green circles).

I23 and Im3̅ could agree well with the HAADF-STEM images projecting along both [100] and [111]. The refined cation lattice (in Figure 3a) is considerably different from that of perovskite Tb0.5Ca0.5MnO3−x, indicating novel metal-anion coordination geometries. If using Im3̅m, the refined position of Tb/Ca would not fit the images. In Figure 3b,c of the [111] zone, helical arrangements were clearly observed around each 9843

DOI: 10.1021/acs.chemmater.7b04115 Chem. Mater. 2017, 29, 9840−9850

Article

Chemistry of Materials parameter and atom displacements U of all other Tb/Ca, Mn, and O ions. After this, the refinements of Mn occupancy at (0, 0, 0) and (1/2, 1/2, 1/2) were performed again and fully yielded a negative value of −0.034. This demonstrates that no Mn ions were on these sites. Considering this, it is not difficult to understand why there seem to be empty channels around the apical sites (marked by green circles) penetrating the whole lattice along the [111] direction, as viewed by [111] STEM images in Figure 3b,c. One could see the image contrast around the apical site (green circle) is relatively lower than elsewhere, and this is due to the Mn vacancy. Moreover, our refined structure model theoretically gives a stoichiometry of Tb0.501Ca0.499Mn0.969O2.406, which is very close to the ICP and iodometry chemical analysis result Tb0.496(2)3+Ca0.504(10)2+Mn0.963(14)2.334(3)+O2.372(1), especially on its Mn and O stoichiometric components. This provides more evidence for the existence of ordered Mn vacancies. It also should be noted that, during refinements, some ions exhibited high Uiso displacements: Tb/Ca(1) Uiso ∼ 0.047, Tb/ Ca(3) Uiso ∼ 0.021, Mn(2) Uiso ∼ 0.043, O(1) Uiso ∼ 0.051, O(2) Uiso ∼ 0.036, and O(7) Uiso ∼ 0.042. The Tb/Ca(1) and O(7) are around the apical/body-center sites (0, 0, 0). Their high Uiso should be attributed to apical Mn vacancies because the vacancies leave them more free space for thermal vibration. For precise description of the structure, anisotropic atom displacements Uaniso were used for these ions, and highly anisotropic displacements were obtained. After this, anisotropic Uaniso displacements were also applied to other ions for better structure description. After convergence, the difference-Fourier procedure was performed again, and all residual peaks were below 3%. The observed, calculated, and difference plots from the XRD and neutron refinements are shown in Figure 4. The final structure model has been shown before in Figure 2c, and will be further discussed in Figure 5. More detailed information about refinement and structure parameters is summarized in section 8 in the SI. All ⟨Mn−O⟩ bond lengths are within 1.75−2.35 Å, as shown in Table 1, with the BVS calculation. For clarity, Tb/Ca are omitted, and the structure features of Tb0.50Ca0.50Mn0.96O2.37 are described by two stacked layers, that is, layer 1 at z ∼ 0 and layer 2 at z ∼ 0.25, as shown in Figure 5. Other layers at z ∼ 0.5 and z ∼ 0.75 could be obtained by I translation of layer 1 and 2-rotation operation of layer 2. Each layer shares different coordination networks composed of different types of ⟨Mn−O⟩ polyhedra. Totally, there are five types of coordination polyhedra, and their central Mn cations are located at the 24f site for Mn(1), 6b site for Mn(2), 12d site for Mn(3), 12e site for Mn(4), and 8c site for Mn(5). In layer 1 (Figure 5a), the green circle represents apical Mn vacancy. Parts A and B are used for the layer description. As shown in part A, each vacancy is surrounded by Mn(3)O4 tetrahedra (reminiscent of the anion-deficient layers in brownmillerite structures) and Mn(1)O5 distorted pyramids. The O(7) are the closest ions to each Mn vacancy (green sphere), with a distance of ∼3.218 Å. Each vacancy is surrounded by 12 O(7) anions. Here, eight O(7) anions are already shown in the plot, another four omitted O(7) are from up and down Mn(3) tetrahedra. Each Mn(3)O4 tetrahedron is linked by four Mn(1)O5 distorted pyramids through cornershared O(7) and O(5), while each Mn(1)O5 pyramid is linked by two Mn(3)O4 tetrahedra, one Mn(4)O5 bipyramid (in module B) and two Mn(5)O6 octahedra (in module C). As for the distortion in Mn(1)O5 pyramids, two diagonal O(6), (7)

Figure 4. Rietveld refinements against room-temperature neutron powder diffraction data and high-resolution XRD data simultaneously. The observed data are represented by a red circle, simulated data is represented by a green line, and difference curves are represented by a blue line.

would move up and another two O(2), (4) down from the pyramid base. The Uaniso values of Mn(1), O(7), and O(2) are relatively high. The Mn(1) cation displaces mainly toward O(7) or O(6), while the O(7) vibrates primarily in the Mn(3)O4 equatorial plane. In this condition, there is a strong possibility that the Mn(3)O4 tetrahedron would rotate around its axial axis because of O(7) equatorial displacements, and the Mn(1)O5 pyramids would undergo a large distortion due to relatively high displacements of these ions. 9844

DOI: 10.1021/acs.chemmater.7b04115 Chem. Mater. 2017, 29, 9840−9850

Article

Chemistry of Materials

Figure 5. Structure description of Tb0.50Ca0.50Mn0.96O2.37, with space group I23, a ∼ 15.27 Å. The Mn ions on different crystallographic sites are depicted by different colors: gray for Mn(3), blue for Mn(1), orange for Mn(4), yellow for Mn(2), and cyan for Mn(5). The structure can be basically described by two layers: (a) layer 1 at z = 0; (b) layer 2 at z ∼ 0.25. Patterns A, B, and C are extracted to describe coordination frameworks for each layer. Local geometries around each Mn cation with bond length and displacements are also plotted here.

Table 1. ⟨Mn−O⟩ Bond Length and BVS Calculation ⟨A⟩ Mn(1) 24f

Mn(2) 6b Mn(3) 12d Mn(4) 12e

Mn(5) 8c a

⟨B⟩ O(2) O(4) O(5) O(6) O(7) O(1) O(3) O(5) O(7) O(1) O(3) O(6) O(2) O(4)

× × × × ×

4 2 2 2 2

×2 ×3 ×3

distance ⟨A−B⟩, Å

BVS calculation

total valencea

1.966(1) 1.909(2) 2.100(2) 2.245(1) 2.040(3) 2.234(3) 2.298(1) 2.016(2) 2.136(4) 2.129(2) 2.292(3) 2.117(1) 2.072(2) 1.989(1)

0.602 0.725 0.432 0.276 0.511 0.301 0.253 0.543 0.393 0.400 0.258 0.390 0.453 0.584

2.546 2.546 2.546 2.546 2.546 1.710 1.710 1.872 1.872 1.838 1.838 1.838 3.111 3.111

has driven the O(1) away from it, which breaks its 3-rotation symmetry in the equatorial plane, with different ⟨Mn−O⟩ bond length and larger angular ⟨O(3)−Mn(4)−O(1)⟩ ∼ 133.95°. The two axial O(6) anions also deviate slightly from the axial axis toward O(3). The anisotropic Uaniso of Mn(2) is unconventionally high in one dimension. However, this is reasonable because its two neighbored Mn(4) are the only two cations for Mn(2) to share with, and either Mn(4) would tend to pull Mn(2) closer to itself to form stronger bonds in an edge-shared manner. This competition causes the Mn(2) to displace toward either side of the two neighbored Mn(4). The O(1) also has large Uaniso in the equatorial plane because they need to move correspondingly to the Mn(2) vibration. In layer 2 (Figure 5b), part C (at x = 0−0.5, y = 0−0.5, z = 0.25) is specially selected for structure description. Other parts in this layer could be obtained by 2-rotation operation of this part. In total, there are four Mn(5)O6 octahedra in this layer; each of them is corner-shared with six Mn(1)O5 distorted pyramids (another two are from up and down layers) by three O(2) with bond length of ∼2.072 Å and three O(4) with shorter bond length of ∼1.989 Å. In acentric space group I23, these O(2) and O(4) are inequivalent, and O(2) exhibits much higher thermal displacement than O(4). This difference in thermal vibration may further enhance the local polarization of the Mn(1) polyhedron and contributes to the piezoelectric responses. Other polyhedra like Mn(1), Mn(3), and Mn(4) polyhedra have been discussed before in layer 1.

Calculated average valence of Mn: 2.2705+.

In part B, it is very unusual that each Mn(2) octahedron would share its two edges and corners with Mn(4)O5 distorted trigonal bipyramids to perform chainlike geometries. The distance between Mn(2) and Mn(4) is 3.048 Å, and they link in an edge-shared manner. The bond angular ⟨O(1)−Mn(4)− O(1)⟩ is ∼92.1°, and ⟨O(1)−Mn(2)−O(1)⟩ is ∼88.6°. The Mn(2)O6 is still a typical octahedron, yet the Mn(4)O5 is not a typical trigonal bipyramid because its coedged linking manner 9845

DOI: 10.1021/acs.chemmater.7b04115 Chem. Mater. 2017, 29, 9840−9850

Article

Chemistry of Materials

Figure 6. (a) Inverse piezoelectric measurements and (b) banana-typed electric hysteresis loop of Tb0.50Ca0.50Mn0.96O2.37 at room temperature. (c−f) Dielectric measurements of Tb0.50Ca0.50Mn0.96O2.37 single pellet from 2 to 380 K with different ac frequencies.

The BVS (Table 1) result yields an averaged oxidation state of Mn2.271+, which is close to the chemical analysis result Mn2.334+. However, we notice that the valence values of Mn(2), Mn(3), and Mn(4) are calculated as less than 2+, and the valence of Mn(5) is a little more than 3+. If the valence is considered to be 2+ for Mn(2), Mn(3), and Mn(4) ions, 2.5+ for Mn(1) ions, and 3+ for Mn(5) ions, the final averaged oxidation state will be Mn2.323+, which is almost equal to the chemical analysis result Mn2.334+. Therefore, the BVS results are reasonable. Piezoelectric Properties. Room-temperature piezoelectric responses were detected by a quasi-static d33 meter (ZJ-3D; Institute of Acoustics, Beijing, China). As shown in section 9 in the SI, after being polarized under an electric field of 300 V, one single pellet (thickness, ∼0.3 mm; area, ∼0.2 cm2) only exhibits a d33 of 0.1 pC N−1, which is at the limit of instrument detection. For confirmation of the piezoelectricity, multilayered pellets were specially prepared to amplify the piezoelectric signal, as illustrated in sections 9 and 10 in the SI. In total, 11 single pellets were connected together by silver paste (also as electrodes), and their interstitial surfaces were connected by Pt wires at intervals. All these stacked pellets could release electrons simultaneously under the same pressure. After being polarized under 300 V, these multilayered pellets as a whole exhibit a much higher d33 value of ±3.5 pC N−1, which already exceeds the d11 of ∼2.31 pC N−1 in α-SiO2. The averaged d33 for each pellet is therefore calculated to be ±0.32 pC N−1. This averaged value is 3 times higher than the d33 ∼ 0.1 pC N−1

obtained from one single pellet measurement. The reason is that the stacked pellets are harder for better stress exertion during the test, and correspondingly they would release more electrons. Inverse piezoelectric measurements and polarization hysteresis measurements at room temperature for a single pellet were also carried out using a TF Analyzer 2000, as shown in Figure 6. When voltages were applied to the materials, strains that are parallel to the polarization directions were obtained attributed to the materials’ macroscopic deformations. The estimated inverse d33 value is 10.5 pm V−1. For ferroelectric measurements, banana-typed47 hysteresis loops (Figure 6b, also section 11 in the SI) were observed during the test, proving that the sample is not a real ferroelectric. In fact, this type of loops comes from the current leakage of the material. In current−voltage curves (Figure S12b in the SI), no charging peaks can be seen, indicating no electric polarization reversed during the test. On the basis of these observations, the sample should be piezoelectric but not ferroelectric. More measurements under higher voltage also demonstrate this point (Figure S12c,d in the SI). Dielectric measurements (Figure 6) were also carried out, but unfortunately no dielectric anomaly peaks could be observed, indicating that no observable phase transition happened from 2 to 380 K. As the frequency increased, both the dielectric constants and loss curves would shift toward hightemperature bands. This is related to the Maxwell−Wagner relaxation where charges trapped at interfaces (grain 9846

DOI: 10.1021/acs.chemmater.7b04115 Chem. Mater. 2017, 29, 9840−9850

Article

Chemistry of Materials

even at 2 K, indicating ferromagnetic responses mixed inside. Magnetization-field (M−H) isotherms were also collected at several selected temperatures, as shown in Figure 7. The inset shows that the M−H curves become S-shaped below 70 K, which should be the onset temperature of antiferromagnetic ordering. Hysteresis loops can be observed at 42 K and especially at 3.5 K, with coercive fields Hc of ∼0.4 T and remnant magnetic moments of ∼0.39 μB per formula. The magnetic moments can reach ∼2.12 μB per formula unit under the magnetic fields of 6 T. This hysteresis loop is typical of spin-canting ferromagnetic features. On the basis of all the above observations, we propose that the obtained phase is primarily antiferromagnetic below TN ∼ 70 K, with mixed spincanting ferromagnetic responses below 40 K.

boundaries) of electrically inhomogenous materials would give rise to Debye-like relaxation processes under an ac measuring voltage. This effect is typical of polycrystalline samples.48,49 The high dielectric loss at high temperature is due to leakage currents arising from oxygen vacancy in crystal and porosity in pellets. Pyroelectric measurements were also performed (section 12 in the SI), but no pyrocurrent peak could be observed which represents a pyro-to-para-electric phase transition. Above 200 K, the pyrocurrent curve evolves and reaches its maximum around 300 K. This is caused by thermal release of electrons trapped at grain boundaries, as mentioned above. However, since our sample can be stable up to 600 °C (873 K) in Ar, it is very possible that the phase transition may occur above 380 K. Further study on high-temperature dielectric and pyroelectric measurement may be needed. However, considering the high electric current leakage and poor resistivity at high temperature, this may be not possible to perform. In summary, our new phase Tb0.50Ca0.50Mn0.96O2.37 should be piezoelectric but not ferroelectric. Magnetic Properties. Zero-field-cooled (ZFC) and fieldcooled (FC) dc magnetization data were collected in an applied field of 100 Oe for Tb0.50Ca0.50Mn0.96O2.37 in the temperature range 5 < T/K< 300, as shown in Figure 7. The data collected

■

Figure 7. M−T and M−H magnetic property measurements of Tb0.50Ca0.50Mn0.96O2.37.

DISCUSSION Reduction of perovskite Tb0.5Ca0.5MnO3−x with sodium hydrate yields novel deficient oxide Tb0.50Ca0.50Mn0.96O2.37, which unexpectedly exhibits a 4ap × 4ap × 4ap lattice expansion on the basis of a cubic perovskite unit with 2/64 Mn ordered vacancies and oxygen rearrangements. The symmetry has been lifted from orthorhombic Pbnm (of precursor) to noncentrosymmetric cubic I23. Such an expansion in lattice with breaking of inversion symmetry was not observed in other reduced phases. During reduction, about 2/64 Mn cations were gradually extracted through ionic diffusion, leaving ordered vacancies on I-centered and apical sites in the new lattice. The deintercalation of oxygen ions and the associated rearrangement of anion lattice further converted the perovskite phase into expanded Tb0.50Ca0.50Mn0.96O2.37 phase which exhibits novel structural features. Since two structure types share no common cation lattice, and cation extraction is also involved, this reduction would not be a dynamically rapid process at low temperature, which explains why the reaction has to be repeated at least 7 times to obtain a pure phase. The obtained Tb0.50Ca0.50Mn0.96O2.37 adopts a new type of ordered pattern of oxygen ions and anion vacancies in I lattice structure, which consists of MnO4 tetrahedra, distorted MnO5 pyramids, bipyramids, and MnO6 octahedra linked in cornershared or edge-shared manners. The anion extraction in this phase is not limited to the oxygen on the MnO6 equatorial plane, and the rearrangement of the oxygen anion lattice and coordination linkage adopts complex ordering schemes. The geometry frameworks, that one Mn(2) octahedron shares its edges with two pyramids, are quite unique among all anion-deficient perovskite phases, as shown in Figure 8. We propose that the formation of this geometry should be related to the tendency of forming rock-salt-typed MnO sheets, as in CaMnO2,44 which is more thermodynamically stable because it

in the range 140 < T/K < 300 can be readily fitted to the Curie−Weiss law [χ = C/(T − θ)] to yield experimental C = 9.491 cm3 K mol−1 and negative θ = −55.6 K, indicating antiferromagnetic behavior. The experimental C = 9.491 cm3 K mol−1 is in good agreement with theoretical C = 9.498 cm3 K mol−1 calculated from our structure model Tb0.496(2)Ca0.504(10)Mn0.963(14)2.334(3)+O2.372(1), which also demonstrates that our structure model is valid. Below 70 K, the ZFC curve first splits from the FC curve, then reaches its maximum at 42 K and further decreases rapidly, indicating primary antiferromagnetic ordering inside the structure. The FC curve continues to increase and seems to be not saturated

Figure 8. Local geometry around Mn(2) in layer 1 at z = 0, in comparison with MnO sheets in CaMnO2. 9847

DOI: 10.1021/acs.chemmater.7b04115 Chem. Mater. 2017, 29, 9840−9850

Article

Chemistry of Materials

Figure 9. In the asymmetric I23 group, both the oxygen anions and Mn(1), Mn(5) cations are allowed to deviate from the original special crystallographic sites to break the local centrosymmetry of the polyhedral. The Mn ions are depicted by different colors: gray for Mn(3), blue for Mn(1), orange for Mn(4), yellow for Mn(2), and cyan for Mn(5).

can be directly synthesized under high temperature.45 As shown in Figure 8, if the central O(1) also bonds with adjacent Mn(1), Mn(3) ions in their equatorial planes, a MnO sheetlike framework could be formed, just as in CaMnO2. However, three factors exclude the possibility for extra bonding between O(1) and Mn(1)/Mn(3): First, the averaged distances between O(1) and adjacent Mn(1), Mn(3) are, respectively, 2.837 and 2.569 Å, which are longer than conventional ⟨Mn−O⟩ bond lengths (1.70−2.35 Å) for bonding. Second, the O(1) and its related four Mn ions all deviate from the original sites in MnO square plane. The Mn(1) and Mn(3) move a little bit away from Mn(2) and Mn(4), while the Mn(2) and Mn(4) move closer to each other, dragging O(1) away from Mn(1)/Mn(3). Third, the Mn(1), Mn(3), and Mn(4) polyhedra are not in ideal octahedral geometry. Therefore, there should be no extra bonding between O(1) and Mn(1)/Mn(3). The linkage pattern of Mn(3)O4 tetrahedron is very exceptional. It neither forms twisted tetrahedra chains between layers as in brownmillerite phases Ca 2 F e 2 O 5 , 5 0 La1−xCaxMnO2+δ,41 and La1−xSrxMnO3−(0.5+x)/2,51 nor forms edge-shared chains or groups as in Sr4Mn3O6.5Cl252 and BaMnO2+x.53 Instead, these tetrahedra are only connected by Mn(1)O5 distorted pyramids. Also, the Mn(1)O5 pyramids in the structure are isolated by other-typed polyhedral. These Mn(1)O5 pyramids are different from those self-connected pyramids in Ca2MnO3.554 and Sr2MnO4.55 The BVS calculation clearly demonstrates charge ordering for Mn(3), Mn(2), Mn(4), and Mn(5) in expended lattice, that is, ∼2+ for tetragonal Mn(3), edge-shared octagonal Mn(2), and bipyramidal Mn(4); and ∼3+ for octagonal Mn(5). However, the Mn(1) exhibits an averaged valence of ∼2.5+, and we fail to figure out its valence distribution. The resulting averaged oxidation state is consistent with iodometry results and M−T Curie−Weiss fitting. Charging ordering has been utilized to introduce local polarization to make multiferroics,23,24 but here in our structure, only piezoelectricity was detected. Although

this charge ordering may not create detectable electric polarization, it could be helpful to stabilize the acentric lattice. The breaking of inversion symmetry is responsible for the piezoelectricity. As shown in Figure 9, compared to centrosymmetric Im3̅, this noncentrosymmetric (NCS) nonpolar space group I23 allows the oxygen O(1), O(6), O(5), and O(7) to deviate slightly from the centrosymmetric special sites, leading to polyhedral distortion and symmetry breaking. The refined structure yields a deviation of 0.035(18) Å for O(1), 0.118(17) Å for O(5), 0.136(17) Å for O(6), and 0.173(0.017) Å for O(7) from their original equatorial or axial sites. For cations, the Mn(5) ions also exhibit a displacement of 0.029(27) Å from specific 8c sites along the local C3 direction and bond with two unequal oxygen O(2) and O(4), with shorter ⟨Mn−O(4)⟩ bonds of ∼1.989 Å and longer ⟨Mn− O(2)⟩ bonds of ∼2.072 Å. The Mn(1) ions could also deviate ∼0.013(18) Å from the plane in an asymmetric way. Large anisotropic displacements would also contribute to the symmetry breaking. For instance, in Mn(2)O6 octahedron, the O(1) and Mn(2) may displace from centric position, and in Mn(5)O6, its two inequivalent O(2) and O(4) share different displacements that may also lead to local symmetry breaking. Generally, there should be local polarization for each asymmetric polyhedron, but for the whole lattice their polarization moments are canceled. Here, the piezoelectric properties originate from the microscopic deformation in the asymmetric unit cell. Several reasons are proposed to explain the symmetry breaking. First, from a synthetic perspective, the compound Tb0.50Ca0.50Mn0.96O2.37 is not a one-step calcined thermostable phase. Instead, it is a metastable phase obtained by posttreatment that does not necessarily need to be energy-favored or close-packing, which is often the case for centrosymmetric phases. Second, from a structural perspective, compared with other topotactic reduced phases, the existence of Tb components forces the lattice to expel its apical Mn cations during the reduction. The ordered extraction of Mn cations 9848

DOI: 10.1021/acs.chemmater.7b04115 Chem. Mater. 2017, 29, 9840−9850

Article

Chemistry of Materials

(2) Scott, J. F.; Paz de Araujo, C. A. Ferroelectric Memories. Science 1989, 246, 1400−1405. (3) Lines, M. E.; Glass, A. M. Principles and Applications of Ferroelectrics and Related Materials; Oxford University Press: Oxford, 1991. (4) Halasyamani, P. S.; Poeppelmeier, K. R. Noncentrosymmetric Oxides. Chem. Mater. 1998, 10, 2753−2769. (5) Haertling, G. H. Ferroelectric Ceramics: History and Technology. J. Am. Ceram. Soc. 1999, 82, 797−818. (6) Saito, Y.; Takao, H.; Tani, T.; Nonoyama, T.; Takatori, K.; Homma, T.; Nagaya, T.; Nakamura, M. Lead-Free Piezoceramics. Nature 2004, 432, 84−87. (7) Wu, H.; Pan, S.; Poeppelmeier, K. R.; Li, H.; Jia, D.; Chen, Z.; Fan, X.; Yang, Y.; Rondinelli, J. M.; Luo, H. K3B6O10Cl: A New Structure Analogous to Perovskite with a Large Second Harmonic Generation Response and Deep UV Absorption Edge. J. Am. Chem. Soc. 2011, 133, 7786−7790. (8) Horiuchi, S.; Tokura, Y. Organic Ferroelectrics. Nat. Mater. 2008, 7, 357−366. (9) Zhang, W.; Xiong, R. G. Ferroelectric Metal-Organic Frameworks. Chem. Rev. 2012, 112, 1163−1195. (10) Burdett, J. K. Use of the Jahn-Teller Theorem in Inorganic Chemistry. Inorg. Chem. 1981, 20, 1959−1962. (11) Cohen, R. E. Origin of Ferroelectricity in Perovskite Oxides. Nature 1992, 358, 136−138. (12) Kang, S. K.; Tang, H.; Albright, T. A. Structures for d0 ML6 and ML5 Complexes. J. Am. Chem. Soc. 1993, 115, 1971−1981. (13) Kunz, M.; Brown, I. D. Out-of-Center Distortions around Octahedrally Coordinated d0 Transition Metals. J. Solid State Chem. 1995, 115, 395−406. (14) Pearson, R. G. The Second-Order Jahn-Teller Effect. J. Mol. Struct.: THEOCHEM 1983, 103, 25−34. (15) Wheeler, R. A.; Whangbo, M. H.; Hughbanks, T.; Hoffmann, R.; Burdett, J. K.; Albright, T. A. Symmetric vs. Asymmetric Linear MXM Linkages in Molecules, Polymers, and Extended Networks. J. Am. Chem. Soc. 1986, 108, 2222−2236. (16) Bersuker, I. B. Modern Aspects of the Jahn-Teller Effect Theory and Applications to Molecular Problems. Chem. Rev. 2001, 101, 1067− 1114. (17) Lefebvre, I.; Lannoo, M.; Allan, G.; Ibanez, A.; Fourcade, J.; Jumas, J. C.; Beaurepaire, E. Electronic Properties of Antimony Chalcogenides. Phys. Rev. Lett. 1987, 59, 2471−2474. (18) Lefebvre, I.; Szymanski, M. A.; Olivier-Fourcade, J.; Jumas, J. C. Electronic Structure of Tin Monochalcogenides from SnO to SnTe. Phys. Rev. B: Condens. Matter Mater. Phys. 1998, 58, 1896−1906. (19) Seshadri, R.; Hill, N. A. Visualizing the Role of Bi 6s “Lone Pairs” in the Off-Center Distortion in Ferromagnetic BiMnO3. Chem. Mater. 2001, 13, 2892−2899. (20) Stoltzfus, M. W.; Woodward, P. M.; Seshadri, R.; Klepeis, J. H.; Bursten, B. Structure and Bonding in SnWO4, PbWO4, and BiVO4: Lone Pairs vs Inert Pairs. Inorg. Chem. 2007, 46, 3839−3850. (21) Watson, G. W.; Parker, S. C. Origin of the Lone Pair of α- PbO from Density Functional Theory Calculations. J. Phys. Chem. B 1999, 103, 1258−1262. (22) Watson, G. W.; Parker, S. C.; Kresse, G. Ab Initio Calculation of the Origin of the Distortion of α-PbO. Phys. Rev. B: Condens. Matter Mater. Phys. 1999, 59, 8481−8486. (23) Choi, Y. J.; Yi, H. T.; Lee, S.; Huang, Q.; Kiryukhin, V.; Cheong, S.-W. Ferroelectricity in an Ising Chain Magnet. Phys. Rev. Lett. 2008, 100, 047601. (24) Ikeda, N.; Ohsumi, H.; Ohwada, K.; Ishii, K.; Inami, T.; Kakurai, K.; Murakami, Y.; Yoshii, K.; Mori, S.; Horibe, Y.; Kito, H. Ferroelectricity from Iron Valence Ordering in the Charge-Frustrated System LuFe2O4. Nature 2005, 436, 1136−1138. (25) Kimura, T.; Goto, T.; Shintani, H.; Ishizaka, K.; Arima, T.; Tokura, Y. Magnetic Control of Ferroelectric Polarization. Nature 2003, 426, 55−58. (26) Wang, J.; Neaton, J. B.; Zheng, H.; Nagarajan, V.; Ogale, S. B.; Liu, B.; Viehland, D.; Vaithyanathan, V.; Schlom, D. G.; Waghmare, U.

breaks the limitation of original topological frameworks and thus allows all cations and oxygen ions to rearrange in more unusual ways, bringing more opportunity for breaking the local inversion symmetry. Third, ordered B-site vacancies would exert lattice strain to break the lattice symmetry. They may play similar roles as Ca 2+ and Y3+ in Ba4CaFe3O9.539 and Ba2YFeO5.5.40 In conclusion, low-temperature topotactic reduction of perovskite Tb0.5Ca0.5MnO3−δ yielded novel deficient phase Tb0.496(2)3+Ca0.504(10)2+Mn0.963(14)2.334(3)+O2.372(1) with unexpected 4-fold lattice expansion in all dimensions resulting in new metal-anion coordination geometries. Ordered Mn cation vacancy was observed as apex and body center of the I lattice and played an important role in maintaining the lattice structure. Room-temperature piezoelectric behavior was detected, confirming the noncentrosymmetry of the phase. Low-temperature antiferromagnetic ordering with mixed ferromagnetic responses was also observed. By this lowtemperature approach, we successfully achieved the breaking of local inversion symmetry of d-spin/state transitional metal centers, offering a new route toward multifunctional solid-state materials.

■

ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.chemmater.7b04115. XRD, ICP, XPS, TG, DSC, piezoelectric, and pyroelectric measurements (PDF)

■

AUTHOR INFORMATION

Corresponding Authors

*E-mail: [email protected]. *E-mail: [email protected]. ORCID

Lin Gu: 0000-0002-7504-031X Junliang Sun: 0000-0003-4074-0962 Ruqiang Zou: 0000-0003-0456-4615 Author Contributions

The manuscript was written through contributions of all authors. All authors have given approval to the final version of the manuscript. Funding

This work is supported by the National Natural Science Foundation of China (Grants 21271014, 51772008, 11275012), National Program for Support of Top-notch Young Professionals, and Changjiang Scholar Program. Notes

The authors declare no competing financial interest.

■

ACKNOWLEDGMENTS The authors thank Prof. Sihai Yang for his help on neutron analysis, ISIS neutron beam center for data collection, Peking University for ICP and XPS measurements, and Tsinghua University for inverse piezoelectric measurements.

■

REFERENCES

(1) Nye, F. J. Physical Properties of Crystals; Oxford University Press: Oxford, U.K., 1957. 9849

DOI: 10.1021/acs.chemmater.7b04115 Chem. Mater. 2017, 29, 9840−9850

Article

Chemistry of Materials

Rock-salt Related Nanoclusters in CaMnO2. J. Am. Chem. Soc. 2009, 131, 8660−8668. (45) Sagdeo, P. R.; Anwar, S.; Lalla, N. P. Powder X-ray Diffraction and Rietveld Analysis of La1‑xCaxMnO3 (0< X< 1). Powder Diffr. 2006, 21, 40−44. (46) Machida, A.; Moritomo, Y.; Ohoyama, K.; Nakamura, A. Neutron Investigation of Tb1‑xCaxMnO3 (x ≥ 0.5). J. Phys. Soc. Jpn. 2001, 70, 3739−3746. (47) Scott, J. F. Ferroelectrics Go Bananas. J. Phys.: Condens. Matter 2008, 20, 021001. (48) Lin, Y. Q.; Chen, X. M.; Liu, X. Q. Relaxor-like Dielectric Behavior in La2NiMnO6 Double Perovskite Ceramics. Solid State Commun. 2009, 149, 784−787. (49) Lunkenheimer, P.; Bobnar, V.; Pronin, A. V.; Ritus, A. I.; Volkov, A. A.; Loidl, A. Origin of Apparent Colossal Dielectric Constants. Phys. Rev. B: Condens. Matter Mater. Phys. 2002, 66, 340− 351. (50) Berggren, J. Refinement of the Crystal Structure of Dicalcium Ferrite, Ca2Fe2O5. Acta Chem. Scand. 1971, 25, 3616−3624. (51) Dixon, E.; Hadermann, J.; Hayward, M. A. Structures and Magnetism of La1‑xSrxMnO3‑(0.5+x)/2 (0.67 ≤ x ≤ 1) Phases. Chem. Mater. 2012, 24, 1486−1495. (52) Romero, F. D.; Hayward, M. A. Structure and Magnetism of the Topotactically Reduced Oxychloride Sr4Mn3O6.5Cl2. Inorg. Chem. 2012, 51, 5325−5331. (53) Hadermann, J.; Abakumov, A. M.; Adkin, J. J.; Hayward, M. A. Topotactic Reduction as a Route to New Close-Packed Anion Deficient Perovskites: Structure and Magnetism of 4H-BaMnO2+x. J. Am. Chem. Soc. 2009, 131, 10598−10640. (54) Leonowicz, M. E.; Poeppelmeier, K. R.; Longo, J. M. Structure Determination of Ca2MnO4 and Ca2MnO3.5 by X-ray and Neutron Methods. J. Solid State Chem. 1985, 59, 71−80. (55) Hernden, B. C.; Lussier, J. A.; Bieringer, M. Topotactic SolidState Metal Hydride Reductions of Sr2MnO4. Inorg. Chem. 2015, 54, 4249−4256. (56) Sarma, D. D.; Rao, C. N. R. XPES Studies of Oxides of 2nd-Row and 3rd-Row Transition-Metals Including Rare-Earths. J. Electron Spectrosc. Relat. Phenom. 1980, 20, 25. (57) Pemba-Mabiala, J. M.; Lenzi, M.; Lenzi, J.; Lebugle, A. XPS Study of Mixed Cerium Terbium Orthophosphate Catalysts. Surf. Interface Anal. 1990, 15, 663−667.

V.; Spaldin, N. A.; Rabe, K. M.; Wuttig, M.; Ramesh, R. Epitaxial BiFeO3 Multiferroic Thin Film Heterostructures. Science 2003, 299, 1719−1722. (27) Pitcher, M. J.; Mandal, P.; Dyer, M. S.; Alaria, J.; Borisov, P.; Niu, H.; Claridge, J. B.; Rosseinsky, M. J. Tilt Engineering of Spontaneous Polarization and Magnetization above 300 K in a Bulk Layered Perovskite. Science 2015, 347, 420−424. (28) Zhang, R.; Senn, M. S.; Hayward, M. A. Directed Lifting of Inversion Symmetry in Ruddlesden–Popper Oxide–Fluorides: Toward Ferroelectric and Multiferroic Behavior. Chem. Mater. 2016, 28, 8399− 8406. (29) Zhang, R.; Abbett, B. M.; Read, G.; Lang, F.; Lancaster, T.; Tran, T. T.; Halasyamani, P. S.; Blundell, S. J.; Benedek, N. A.; Hayward, M. A. La2SrCr2O7: Controlling the Tilting Distortions of n = 2 Ruddlesden–Popper Phases through A-Site Cation Order. Inorg. Chem. 2016, 55, 8951−8960. (30) Hayward, M. A. Topochemical Reactions of Layered TransitionMetal Oxides. Semicond. Sci. Technol. 2014, 29, 064010. (31) Hayward, M. A.; Green, M. A.; Rosseinsky, M. J.; Sloan, J. Sodium Hydride as a Powerful Reducing Agent for Topotactic Oxide Deintercalation: Synthesis and Characterization of the Nickel (I) Oxide LaNiO2. J. Am. Chem. Soc. 1999, 121, 8843−8854. (32) Hayward, M. A.; Cussen, E. J.; Claridge, J. B.; Bieringer, M.; Rosseinsky, M. J.; Kiely, C. J.; Blundell, S. J.; Marshall, I. M.; Pratt, F. L. The Hydride Anion in an Extended Transition Metal Oxide Array: LaSrCoO3H0.7. Science 2002, 295, 1882−1884. (33) Tsujimoto, Y.; Tassel, C.; Hayashi, N.; Watanabe, T.; Kageyama, H.; Yoshimura, K.; Takano, M.; Ceretti, M.; Ritter, C.; Paulus, W. Infinite-Layer Iron Oxide with a Square-Planar Coordination. Nature 2007, 450, 1062−1065. (34) Dixon, E.; Hayward, M. A. The Topotactic Reduction of Sr3Fe2O5Cl2 Square Planar Fe (II) in an Extended Oxyhalide. Inorg. Chem. 2010, 49, 9649−9654. (35) Romero, F. D.; Coyle, L.; Hayward, M. A. Structure and Magnetism of Sr3Co2O4Cl2: An Electronically Driven Lattice Distortion in an Oxychloride Containing Square Planar Co (II) Centers. J. Am. Chem. Soc. 2012, 134, 15946−15952. (36) Yamamoto, T.; Kobayashi, Y.; Hayashi, N.; Tassel, C.; Saito, T.; Yamanaka, S.; Takano, M.; Ohoyama, K.; Shimakawa, Y.; Yoshimura, K.; Kageyama, H. Sr1‑xBaxFeO2 (0.4 < x < 1): A New OxygenDeficient Perovskite Structure. J. Am. Chem. Soc. 2012, 134, 11444− 11454. (37) Denis Romero, F.; Leach, A.; Moller, J. S.; Foronda, F.; Blundell, S. J.; Hayward, M. A. Strontium Vanadium Oxide–Hydrides: “SquarePlanar” Two-Electron Phases. Angew. Chem., Int. Ed. 2014, 53, 7556− 7559. (38) Patino, M. A.; Zeng, D.; Bower, R.; McGrady, J. E.; Hayward, M. A. Coupled Electronic and Magnetic Phase Transition in the InfiniteLayer Phase LaSrNiRuO4. Inorg. Chem. 2016, 55, 9012−9016. (39) Turp, S. A.; Hargreaves, J.; Baek, J.; Halasyamani, P. S.; Hayward, M. A. Noncentrosymmetric Cation Order in the Cubic Perovskite Ba4CaFe3O9.5. Chem. Mater. 2010, 22, 5580−5587. (40) Luo, K.; Johnson, R. D.; Tran, T. T.; Halasyamani, P. S.; Radaelli, P. G.; Hayward, M. A. Ba2YFeO5.5: A Ferromagnetic Pyroelectric Phase Prepared by Topochemical Oxidation. Chem. Mater. 2013, 25, 1800−1808. (41) Dixon, E.; Hadermann, J.; Ramos, S.; Goodwin, A. L.; Hayward, M. A. Mn (I) in an Extended Oxide: The Synthesis and Characterization of La1‑xCaxMnO2+δ (0.6 ≤ x ≤ 1). J. Am. Chem. Soc. 2011, 133, 18397−18405. (42) Rodriguez-Carvajal, J. Recent Advances in Magnetic Structure Determination by Neutron Powder Diffraction. Phys. B 1993, 192, 55−69. (43) Wan, W.; Sun, J.; Su, J.; Hovmöller, S.; Zou, X. ThreeDimensional Rotation Electron Diffraction: Software RED for Automated Data Collection and Data Processing. J. Appl. Crystallogr. 2013, 46, 1863−1873. (44) Varela, A.; Dios, S. d.; Parras, M.; Hernando, M.; FernandezDiaz, M. T.; Landa-Canovas, A. R.; Gonzalez-Calbet, J. M. Ordered 9850

DOI: 10.1021/acs.chemmater.7b04115 Chem. Mater. 2017, 29, 9840−9850