Topotactic Reduction toward a Noncentrosymmetric Deficient

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Topotactic reduction towards a non-centrosymmetric 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 Chem. Mater., Just Accepted Manuscript • DOI: 10.1021/acs.chemmater.7b04115 • Publication Date (Web): 30 Oct 2017 Downloaded from http://pubs.acs.org on November 2, 2017

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Chemistry of Materials

Topotactic reduction towards a non-centrosymmetric 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, 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, China

§

Beijing National Laboratory for Condensed Matter Physics, Institute of Physics, Chinese Academy of Sciences, Beijing 100190, P. R. China. E-mail: [email protected], [email protected]

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 I 2 3 (a ~ 15.27 Å) based on a 4ap x 4ap x 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 quasi-static 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 spincanting below 40 K. This work provides a new way towards synthesizing unconventional acentric materials, in the absence of second-order Jahn-Teller active “distortion centers”.

Introduction Searching for new crystal structured solid state materials has attracted much interest since new crystal structure could intrinsically leads to undiscovered physical and electronic properties. These new materials, if crystallizing in non-centrosymmetric (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 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 towards non-centrosymmetry by high temperature fabrication is to introduce d0 transitional metal cations (Ti4+, Zr4+, Nb5+, W6+) which would undergo off-centering displacements in their coordination polyhedra, as illustrated by second-order Jahn-Teller

(SOJT) distortion.10-16 However, the drawback is that these symmetry-breaking cations are not compatible with dspin behaviors, which are the basis of magnetic or spintronic devices because d-spin behaviors require metal ions to have unpaired d electrons coupled with each other through exchange interaction. In order to achieve the coexistence of NCS and d-spin related properties, many approaches have been made, such as introducing “Longpair” ns2 cations (Pb2+, Bi3+),17-22 creating charge order,23,24 manipulating spiral spin arrangements,25,26 or more recently utilizing cooperative tilting distortion.27,28,29 Compared with various approaches, low-temperature topotactic reactions,30 which specifically target at preparing new structural compounds, is another promising way towards new NCS compounds especially with d-spin behavior. In topotactic reaction, the oxygen ions could be extracted from 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 metal-anion 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. Besides, topotactic reactions would further tune the d-

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spin/state of transitional metal centers which interact with each other through extended metal-anion coordination frameworks to give rise to new electronic structure and related properties.

spin-canting behavior. This work proves that low temperature topotactic reaction can be a powerful technique in searching unconventional NCS materials with new crystal structure and d-spin related properties.

Many excellent works have been made in topotactic field,31-38 but only few of them achieved noncentrosymmetry in the lattice. For reduced Ba4CaFe3O9.5,39 its complex cation-ordered structure has a low-symmetry arrangement of CaO6, FeO5 and FeO4 polyhedra on the Bcation 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 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 radius. One reason that most topochemical phases fail to achieve non-centrosymmetry may be the fact that only alkaline metal (Ca2+, Sr2+, Ba2+) and transitional metals are primarily concerned during materials preparation, yet Lanthanide, especially post Lanthanide were 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 A-site of the lattice, would considerably affect the whole metalanion frameworks in reduced structure since they require higher coordination number than alkaline metal cations. One example is the reduced La1-xCaxMnO2+δ,41 in which its lattice has to adopt chain-like 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.

Experimental Section

In this report, a new structured non-centrosymmetric 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 4apx4apx4ap lattice expansion on the basis of a simple cubic perovskite unit cell, crystallizing in a non-centrosymmetic space group I 2 3, a ~ 15.27 Å. The metal-anion coordination frameworks in the structure evolve in a very unusual way that the original corner-shared 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 ordered Mn vacancy in deficient perovskite systems. Room temperature piezoelectricity was successfully measured, with quasi-static d33 value of ~0.32 pC·N-1 and inverse d33 value of ~10.5 pm·V-1. When making into multilayered pellets, the whole pellets 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

Preparation of Tb0.5Ca0.5MnO3 perovskites as Precursor. Perovskite Tb0.5Ca0.5MnO3 was synthesized by citric method. Appropriate stoichiometric ratios of Tb4O7 (A.R.), CaCO3 (A.R. 99.5%) and MnCO3 (A.R. 99.5%) were firstly dissolved separately in nitric acid and then mixed together. About two mole equivalents of citric acid were then added. The solution was dried while stirring to form gel, in which all the ions were well distributed. Later, these prepared gel was calcined under 1300 ºC for 48 hours with 3 times 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 5 g of Tb0.5Ca0.5MnO3 phase was thoroughly ground with two mole equivalents of NaH in agate mortar in an argonfilled 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 hours. The sample was then taken out, reground in glovebox and resealed in tubes before being heated for 7 further periods of 48 h at 210 ˚C. Finally, sample was washed with methanol for several times under a nitrogen atmosphere to remove sodiumcontaining 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: 6s/step) at 50 kV and 40 mA. TOF neutron powder diffraction data were collected on WISH diffractometer (ISIS neutron source, U.K.). Rietveld refinements of XRD and Neutron diffraction data were performed by Fullprof software.42 Selected area electron diffractions (SAED) were carried out on a JEM 2100F 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 of 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 three-dimensional 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) operated at 200 keV and equipped with double spherical aberration (Cs) correctors. The at-


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tainable resolution of the probe defined by the objective pre-field is 78 picometers. The ABF (angular bright-field) images were also collected on this TEM. In order to facilitate the observation we inversed 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 NETZSCH STA449C. Magnetic properties were measured on SQUID MPMS. For electric properties, since the reduced sample cannot be calcined under high temperature, a little quantity of epoxy resin was added into the powder (about 1 drop epoxy for ~0.2 g 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 hours for solidification. After that, these harden 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 quasistatic d33 meter (ZJ-3D; Institute of Acoustics, Beijing, China.) and TF analyzer 2000. Dielectric and pyroelectric measurements were carried out on SQUID PPMS with WK 6500B and KEITHLEY 6517A devices, respectively. 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 Supporting Information SI 1. As the reaction goes, the reflection peaks of the precursor are gradually decreasing while a new set of reflection peaks are getting 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 on reflection peaks were observed in the XRD pattern. The XRD pattern of our final product is provided in SI 2. With the help of TEM techniques (discussed later in the next section), all reflection peaks can be well indexed by cubic I lattice (a ~ 15.27 Å) and no impurities are found. The product has 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 strongest peaks (044)/(404)/(440) (d ~ 2.7 Å) of the product, demonstrating that the new lattice adopts higher symmetry and may still remain perovskite topological frameworks. The stoichiometry of elements Tb, Ca, 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. The detailed information of these measurements was provided in SI 3. 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 Tb0.50Ca0.50Mn0.962.33+O2.37. Our Tb0.50Ca0.50Mn0.96O2.37 is a unique and final phase that can be obtained under current reaction condition. This is because: first, further reduction at 210 °C will cause no difference in XRD pattern, proving that the current phase is stable and no more structural changes would happen under 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 Tb1xCaxMnO2+δ, which exhibits similar structure with reported La1-xCaxMnO2+δ41 (see SI 4); Third, although the precursor Tb1-xCaxMnO3 is actually a solid solution in the whole 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 precursor (such as Tb0.3Ca0.7MnO3) will yield impurity CaMnO2 (see SI 5) 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 relationship with its parent phase. As shown in SI 6, the Tb0.50Ca0.50Mn0.96O2.37 is stable at room temperature, but will begin re-oxidation above 150 °C. At 480 °C, the sample will be finally re-oxidized back to its parent phase Tb0.50Ca0.50Mn0.96O2.95 with perovskite structure, though with Mn vacancy. This facile re-oxidation 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 SI 7) under Ar atmosphere reveal that the sample can sustain up to 600 °C. But at 820 °C, the sample will totally decompose to CaMnO244,45 and Tb2O3. Crystal Structure determination. At the beginning, the HRXRD pattern indexing suggests I-, P-, or even Runit cell with various possible space group and cell length. Therefore, TEM RED (rotation electron diffraction) techniques are specially used to identify its unit cell and space group (Figure 1 a-c.), with the help of TEM SAED analysis (Figure 1 d, e.). Several possible types of lattice could be derived as cubic I m -3 m, I m -3 and I 2 3; tetragonal I 4/m m m; orthorhombic I 2 2 2 and I m m m, with cell length in all three dimension 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 to be 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, cubic I lattice is preferred and tried for later structural determination.



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It is clear in SAED data (Figure 1 d, e.) that a series of reflection (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 these imply that during reduction, the Tb0.50Ca0.50Mn0.96O2.37 undergoes 4-fold expansion in all crystallographic a, b, c-axis on the basis of cubic perovskite unit cell. Based on the above, we initially built a 4-fold expanded structural model with I 2 3, a ~ 15.27 Å (Later we will discuss other space group.) that contains 64 cubic perovskite units, but only heavy ions were set, as shown in Figure 2 a, b. The Tb/Ca ions are placed at 8c, 24f sites that are equal to A-site in perovskite, and the Mn ions are placed at 2a, 6b, 12d, 12e, 24f sites that are equal to B-site. The structure model could also be seen as a 3a x 3b x 2c expansion on the basis of distorted orthorhombic perovskite unit cell of precursor Tb0.5Ca0.5MnO3-x.46 (Pbnm, a = 5.334 Å, b = 5.458 Å, c = 7.467 Å).

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es projecting along both [100] and [111]. The refined cation lattice (In Figure 3 a) is considerably different from perovskite Tb0.5Ca0.5MnO3-x, indicating novel metal-anion coordination geometries. If using I m -3 m, the refined position of Tb/Ca would not fit the images. In Figure 3 b and 3 c of [111] zone, helical arrangements were clearly observed around each apical and body-centered sites (marked by green circles) of the lattice. This helical arrangement could result from 3 or -3 -rotation axis. Carefully scrutinizing the atom information on these images, we found that the symmetry operation m and 2-rotation could not be differentiated. But the diagonal mirror plane m (like in I m -3 m) and 4-rotation axis (like in I 4/m m m) could be precluded because if they exist, the [111] zone helical arrangements would not be allowed. The space group I m -3 is centric while its subgroup I 2 3 is acentric. Yet our diffraction data did not provide a clear choice between them. But since room temperature piezoelectric responses were successfully detected, (discussed in next section) the non-centrosymmetric space group I 2 3 were 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 Difference-Fourier procedure were repeated. Every time after the procedure, at least one or two oxygen ions could be identified, with residual density of 0.2~0.3. These oxygen ions were then added into the structure and further refined. The Difference-Fourier procedure was repeated for 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 be still readily reached and all coordinate parameters were stable. All obtained bond lengths in the structure are among reasonable range of 1.78~2.35 Å . The inversed ABF images as shown in Figure 3 d could provide information about the oxygen location (marked by red circles). We can see that the projection of oxygen arrangements is in good agreements with the image.

Figure 1. TEM data, (a), (b), (c) rotation electron diffraction (RED) data and (d), (e) selected area electron diffraction (SAED) data of the Tb0.50Ca0.50Mn0.96O2.37. The strategy for structure determination (the final structure is shown in Figure 2 c.) is introduced here in details. The first step is to locate all cations in the lattice. XRD data was firstly used for the cation lattice refinements and 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 I m -3 or I 2 3. However, when the space group was I m -3 m, the convergence could not be reached and the cations could not be settled. One possible reason is that the symmetry I m -3 m is too high to describe the structure. HAADF-STEM images were collected to further confirm the cation lattice. As shown in Figure 3 a-c, the refined cation structure in I 2 3 and I m -3 could agree well with the HAADF-STEM imag-

However, the strange thing is that no bonding oxygen ions could be found around the apical (0, 0, 0) and also body-centered (1/2, 1/2, 1/2) Mn cations, as shown in Figure 2 b and Figure 3 c, marked by green sphere. At first we thought that the problem is about the symmetry and we tried again with lower symmetry I 3 and even I 2 2 2, 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 body centered sites could also be unreasonably high (more than 60) if refined. This anomaly indicates that the reduction


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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, the heavy cations’ arrangements of Tb0.50Ca0.50Mn0.96O2.37; (c) The final crystal structure of Tb0.50Ca0.50Mn0.96O2.37. on its Mn and O stoichiometric components. This provides more evidence for the existence of ordered Mn vacancies.

Figure 3. (a), (b), (c) HAADF-STEM images of obtained Tb0.50Ca0.50Mn0.96O2.37 projected along [100] and [111]; (d) Inversed ABF images revealing possible oxygen sites (red circle). Helical cation arrangements (white circles) can be seen from (b)(c), with Mn vacancies (green circles). 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 firstly reset the Mn ions occupancy to 0.14 and refined the coordination 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 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 seems to be empty channels around the apical sites (marked by green circles) penetrating the whole lattice along [111] direction, as viewed by [111] STEM images in Figure 3 b, 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

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



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Figure 5. Structure description of Tb0.50Ca0.50Mn0.96O2.37, with space group I 2 3, 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), 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. Pattern A, B and C are extracted to describe coordination frameworks for each layer. Local geometries of around each Mn cation with bond length and displacements are also plot here. 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, 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. In order to precisely describe the structure, anisotropic atom displacements Uaniso were used for these ions and highly anisotropic displacements were obtained. After this, anisotropic Uaniso 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 2 c, and will be further discussed in Figure 5. More detailed information about refinement and structure parameters is summarized in SI 8. All bond lengths are among 1.75~2.35 Å, as shown in Table 1, with 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


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polyhedra. Totally, there are five types of coordination polyhedra, and their central Mn cations locate at 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).

Mn(3)O4 tetrahedron would rotate around its axial axis due to O(7) equatorial displacements, and the Mn(1)O5 pyramids would undergo large distortion due to relatively high displacements of these ions.

In layer 1 (Figure 5 a), the green circle represents apical Mn vacancy. Part 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 are surrounded by twelve 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 corner-shared 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) would move up and another two O(2), (4) down from the pyramid base.

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 chain-like geometries. The distance between Mn(2) and Mn(4) is 3.048 Å, and they link in edge-shared manner. The bond angular is ~92.1°, and 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 co-edged linking manner has driven the O(1) away from it, which breaks its 3-rotation symmetry in the equatorial plane, with different bond length and larger angular ~ 133.95°. The two axial O(6) anions also deviate slightly from the axial axis towards O(3).

Table 1. Bond distance and BVS calculation. bond length and BVS calculation

Distance Å

BVS calculation

Mn(1) 24f

O(2)

1.966(1)

0.602

O(4)

1.909(2)

0.725

O(5)

2.100(2)

0.432

O(6)

2.245(1)

0.276

O(7)

2.040(3)

0.511

O(1) x 4

2.234(3)

0.301

O(3) x 2

2.298(1)

0.253

O(5) x 2

2.016(2)

0.543

O(7) x 2

2.136(4)

0.393

O(1) x 2

2.129(2)

0.400

O(3)

2.292(3)

0.258

O(6) x 2

2.117(1)

0.390

O(2) x 3

2.072(2)

0.453

O(4) x 3

1.989(1)

0.584

Mn(2) 6b

Mn(3) 12d

Mn(4) 12e

Mn(5)

8c

Total valence 2.546

1.710

1.872

1.838

3.111

Calculated average valence of Mn: 2.2705+.

The Uaniso of Mn(1), O(7) and O(2) are relatively high. The Mn(1) cation displaces mainly towards O(7) or O(6), while the O(7) vibrates primarily in Mn(3)O4 equatorial plane. In this condition, there is strong possibility that the

The anisotropic Uaniso of Mn(2) is unconventionally high in one dimension. Yet 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 edge-shared manner. This competition render the Mn(2) to displace towards 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 5 b), 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. Totally, 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 I 2 3, 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 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. 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 of Mn(2), Mn(3) and Mn(4) are calculated 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), Mn(4) ions, 2.5+ for Mn(1) ions, and 3+ for Mn(5) ions, the final averaged oxidation state will be Mn2.323+ that is almost equal to 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 SI 9, after polarized under an electric field of 300 V, one single pellet (thickness ~0.3 mm, area ~0.2 cm2) only



(SI 11, Figure S12 b), no charging peaks can be seen, indicating no electric polarization reversed during the test. Based on these the sample should be piezoelectric but not ferroelectric. More measurements under higher voltage also demonstrate this point (SI 11, Figure S12 c, d).

(a) 35

Pyroelectric measurements were also performed (SI 12.) but no pyro-current peak could be observed which represents pyro- to para-electric phase transition. Above 200 K, the pyro-current curve evolves and reaches its maximum around 300 K. This is caused by thermal release of electrons trapped at grain boundaries, as mentioned above.

(b) 0.10

30

2

Polarization (uC/cm )

Displacement (nm)

exhibits a d33 of 0.1 pC · N-1 that is to the limit of instrument detection. In order to confirm the piezoelectricity, multi-layered pellets were specially prepared to amplify the piezoelectric signal, as illustrated in SI 9 and SI 10. 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 three times higher than the d33 ~ 0.1 pC · N-1 obtained from one single pellet measurements. The reason is that the stacked pellets are harder for better stress exertion during test and correspondingly they would release more electrons.

25 20 15 10 5

0.05 0.00 -0.05

0 -3

-1 0 1 E (kV/mm)

2

Frequency 1 KHz 10 KHz 100 KHz 1 MHz

19 18

-0.10 -1.5

3

(d)

17 16 0

(e)

50

100 150 Temperature (K)

(f)

Frequency 1 KHz 10 KHz 100 KHz 1 MHz

0.10

0.05

-1.0

-0.5 0.0 0.5 E (kV/mm)

1.0

1.5

300 250 200 150 100 50 0 200

200

250 300 350 Temperature (K)

400

0.7 0.6

Dielectric Loss

Dielectric Constant ε

(c) 20

-2

Dielectric Constant ε

-5

Dielectric Loss

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 8 of 13

0.5 0.4 0.3 0.2 0.1

0.00 0

50


200

0.0 200

250


400

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 K to 380 K with different AC frequency. 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 10.5pm · V-1. For ferroelectric measurements, bananatyped47 hysteresis loops (Figure 6 b, also SI 11) were observed during 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

Dielectric measurements (Figure 6.) were also carried out but unfortunately no dielectric anomaly peaks could be observed, indicating no observable phase transition happened from 2 K to 380 K. As the frequency increased, both the dielectric constants and loss curves would shift towards high temperature bands. This is related to the Maxwell-Wagner relaxation where charges trapped at interfaces (grain 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.

However, since our sample can be stable up to 600 ºC (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. But 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 field cooled (FC) DC magnetization data were collected in an applied field of 100 Oe for Tb0.50Ca0.50Mn0.96O2.37 in the temperature range of 5 < T/K< 300, as shown in Figure 7. The data collected 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 agreements 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 firstly 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 not saturated 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-shape below 70 K, which should be 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 Tesla and remnant magnetic moments of ~0.39 μB per formula. The


Page 9 of 13

magnetic moments can reach ~ 2.12 μB per formula unit under the magnetic fields of 6 Tesla. This hysteresis loop is typical of spin-canting ferromagnetic features. Based all above, we propose that the obtained phase is primarily antiferromagnetic below TN ~ 70 K, with mixed spincanting ferromagnetic responses below 40 K. Discussion Reduction of perovskite Tb0.5Ca0.5MnO3-x with sodium hydrate yields novel deficient oxide Tb0.50Ca0.50Mn0.96O2.37, which unexpectedly exhibits a 4apx4apx4ap lattice expansion on the basis of cubic perovskite unit with 2/64 Mn ordered vacancies and oxygen rearrangements. The symmetry has been lifted from orthorhombic P b n m (of precursor) to non-centrosymmetric cubic I 2 3. Such expansion in lattice with breaking of inversion symmetry was not observed in other reduced phases.

χ (emu⋅mol-1⋅Oe-1)

6 5 4 3

ture 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 get 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 corner-shared or edge-shared manners. The anion extraction in this phase is not limited to the oxygen on MnO6 equatorial plane, and the rearrangement of oxygen anion lattice and coordination linkage adopt complex ordering schemes.

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

T'~42 K

2 T"~70 K

1 0 0

50

15000 10000

Μ (emu⋅mol-1)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


5000

100 150 200 Temperature (K)

250

3.5 K 42 K 60 K 68 K

0 500

-5000 0

-10000 -500

-15000 -6

-0.1

-4

-2

0 2 H (Tesla)

0.0

0.1

4

6

Figure 7. M-T and M-H magnetic properties measurements of Tb0.50Ca0.50Mn0.96O2.37. 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 de-intercalation 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 struc-

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 rocksalttyped MnO sheets, as in CaMnO244, which is more thermodynamically stable because it can be directly synthesized under high temperature45. 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 sheet-like 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 Å and 2.569 Å that are longer than conventional 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 Ca2Fe2O5,50 La141 and La1−xSrxMnO3−(0.5+x)/251, nor forms edgexCaxMnO2+δ shared chains or groups as in Sr4Mn3O6.5Cl252 and BaMnO2+x53. 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 selfconnected pyramids in Ca2MnO3.554 and Sr2MnO455.



1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

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Figure 9. In asymmetric I 2 3 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), cyan for Mn(5). The BVS calculation clearly demonstrate 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 bi-pyramidal Mn(4), and ~3+ for octagonal Mn(5). Yet the Mn(1) exhibits an averaged valence of ~2.5+ and we fails 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 for making multiferroics23,24, but here in our structure, only piezoelectricity was detected. Although this charge ordering may not create detectable electric polarization, they 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 I m -3, this non-centrosymmetric (NCS) nonpolar space group I 2 3 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 O5, 0.136(17) Å for O6, and 0.173(0.017) Å for O7 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 bonds of ~1.989 Å and longer 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 cancelled. Here, the piezoelectric properties are originating from the microscopic deformation in the asymmetric unit cell. Several reasons are proposed to explain the symmetry breaking. First, from synthetic perspective, the compound Tb0.50Ca0.50Mn0.96O2.37 is a not one-step calcined thermostable phase. Instead, it is a metastable phase obtained by post-treatment that needs not necessarily to be energyfavored or close-packing, which is often the case for centrosymmetric phases. Second, from 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 breaks the limitation of original topological frameworks and thus allow 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 Ca2+ and Y3+ do in Ba4CaFe3O9.539 and Ba2YFeO5.540. In conclusion, low temperature topotactic reduction of perovskite Tb0.5Ca0.5MnO3-δ yielded novel deficient phase


Page 11 of 13

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


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 non-centrosymmetry of the phase. Low temperature antiferromagnetic ordering with mixed ferromagnetic responses were also observed. By this low temperature approach, we successfully achieved the breaking of local inversion symmetry of d spin/state transitional metal centers, offering a new route towards multifunctional solid states materials.

ASSOCIATED CONTENT Supporting Information. XRD, ICP, TG, DSC, Piezoelectric and pyroelectric measurements. This material is available free of charge via the Internet at http://pubs.acs.org.

AUTHOR INFORMATION Corresponding Author * Ruqiang Zou Beijing Key Laboratory for Theory and Technology of Advanced Battery Materials, Department of Materials Science and Engineering, College of Engineering, Peking University, Beijing 100871, China. Email: [email protected] * Guobao Li College of Chemistry and Molecular Engineering, Peking University, Beijing, China

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

Funding Sources 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.

ACKNOWLEDGMENT 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.

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Topotactic Reduction toward a Noncentrosymmetric Deficient

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