PolarâNonpolar Phase Transition Accompanied by Negative Thermal

Subscriber access provided by BUFFALO STATE

Article

Polar-nonpolar phase transition accompanied by negative thermal expansion in perovskite-type Bi PbNiO 1-x

x

3

Yuki Sakai, Takumi Nishikubo, Takahiro Ogata, Hayato Ishizaki, Takashi Imai, Masaichiro Mizumaki, Takashi Mizokawa, Akihiko Machida, Tetsu Watanuki, Keisuke Yokoyama, Yoichi Okimoto, Shin-ya Koshihara, Hena Das, and Masaki Azuma Chem. Mater., Just Accepted Manuscript • DOI: 10.1021/acs.chemmater.9b00929 • Publication Date (Web): 29 May 2019 Downloaded from http://pubs.acs.org on May 29, 2019

Just Accepted “Just Accepted” manuscripts have been peer-reviewed and accepted for publication. They are posted online prior to technical editing, formatting for publication and author proofing. The American Chemical Society provides “Just Accepted” as a service to the research community to expedite the dissemination of scientific material as soon as possible after acceptance. “Just Accepted” manuscripts appear in full in PDF format accompanied by an HTML abstract. “Just Accepted” manuscripts have been fully peer reviewed, but should not be considered the official version of record. They are citable by the Digital Object Identifier (DOI®). “Just Accepted” is an optional service offered to authors. Therefore, the “Just Accepted” Web site may not include all articles that will be published in the journal. After a manuscript is technically edited and formatted, it will be removed from the “Just Accepted” Web site and published as an ASAP article. Note that technical editing may introduce minor changes to the manuscript text and/or graphics which could affect content, and all legal disclaimers and ethical guidelines that apply to the journal pertain. ACS cannot be held responsible for errors or consequences arising from the use of information contained in these “Just Accepted” manuscripts.

is published by the American Chemical Society. 1155 Sixteenth Street N.W., Washington, DC 20036 Published by American Chemical Society. Copyright © American Chemical Society. However, no copyright claim is made to original U.S. Government works, or works produced by employees of any Commonwealth realm Crown government in the course of their duties.

Page 1 of 12 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

Chemistry of Materials

Polar-‐nonpolar phase transition accompanied by negative thermal expansion in perovskite-‐type Bi1−xPbxNiO3 Yuki Sakai,*,†,‡ Takumi Nishikubo,‡ Takahiro Ogata,‡ Hayato Ishizaki,‡ Takashi Imai,‡ Ma-‐ saichiro Mizumaki,§ Takashi Mizokawa,¶ Akihiko Machida,# Tetsu Watanuki,# Keisuke Yokoya-‐ ma,| Yoichi Okimoto,| Shin-‐ya Koshihara,| Hena Das‡,\ and Masaki Azuma*,†,‡ †

Kanagawa Institute of Industrial Science and Technology, 705-‐1 Shimoimaizumi, Ebina, 243-‐0435, Japan Laboratory for Materials and Structures, Tokyo Institute of Technology, 4259 Nagatsuta, Midori, Yokohama, 226-‐ 8503, Japan § Japan Synchrotron Radiation Research Institute, SPring-‐8, Sayo-‐gun, Hyogo 679-‐5198, Japan ¶ Department of Applied Physics, School of Advanced Science and Engineering, Waseda University, 3-‐4-‐1 Okubo, Shinjuku-‐ku, Tokyo 169-‐8555, Japan # Synchrotron Radiation Research Center, National Institutes for Quantum and Radiological Science and Technolo-‐ gy, Sayo, Hyogo 679-‐5148, Japan | Department of Chemistry, Tokyo Institute of Technology, Meguro, Tokyo 152-‐8551, Japan ‡

\

World Research Hub Initiative, Institute of Innovative Research (IIR), Tokyo Institute of Technology, 4259 Na-‐ gatsuta, Midori, Yokohama, 226-‐8503, Japan

ABSTRACT: Perovskite-‐oxide Bi1−xPbxNiO3 for 0.60 ≤ x ≤ 0.80 was found to show a polar-‐orthorhombic to non-‐polar-‐ orthorhombic phase transition accompanied by negative thermal expansion (NTE). Bi1−xPbxNiO3 showed successive crys-‐ tal structure changes depending on the amount of Pb. As the amount of Pb increased, the crystal structure changed from a triclinic one with Bi3+/Bi5+ long-‐range ordering to an orthorhombic one with Bi3+/Bi5+ short-‐range ordering; then it changed into a polar orthorhombic structure without Bi3+/Bi5+ ordering and finally to a polar LiNbO3-‐type one. The key to the inversion symmetry breaking in PbNiO3, where both 6s2 lone-‐pair and Jahn-‐Teller (JT) active cations are absent, is the high valency state of Pb4+. Our results suggest that the polar orthorhombic phase can be realized by using high valence A-‐ site cations in addition to controlling the tolerance factor.

21

1. INTRODUCTION

xBixVO3.

Nanoscale manufacture of electronic devices and optical communications require precise positioning, and thus, even small amounts of thermal expansion can be a prob-‐ lem. For this reason, materials showing negative thermal expansion (NTE) have attracted much attention because composites made from them are expected to be able to compensate for the thermal expansion.1-‐7 Compounds with flexible frameworks in their crystal structures such as β-‐LiAlSiO4,1,2 ZrW2O88 and Cd(CN)29 can be catego-‐ rized into the first generation of NTE materials. The last decade has seen remarkable development in materials with NTE resulting from phase transitions. In particular, a large NTE of αL = −30 ppm K−1 coupled with a magnetic transition was discovered in anti-‐perovskite-‐type manga-‐ nese nitride.10-‐16 PbTiO3-‐based perovskite-‐type oxides were found to show NTE originating from a ferroelectric-‐ paraelectric transition.17-‐21 Colossal volume shrinkage, as large as 7.8 %, induced by essentially the same tetragonal-‐ to-‐cubic phase transition was recently found in Pb1-‐

BiNiO3-‐based perovskite-‐type oxides show large nega-‐ tive coefficients of thermal expansion.24-‐29 The parent ma-‐ terial, BiNiO3, has an unusual valence distribution of Bi3+0.5Bi5+0.5Ni2+O3.30-‐32 Bi3+ and Bi5+ occupy distinct crystal-‐ lographic sites in a √2ap × √2ap × 2ap triclinic unit cell with space group P-‐1, where ap is the lattice constant of a simple cubic perovskite-‐type structure. Since Bi 6s and Ni 3d levels are close to each other in energy, an intersite charge transfer between Bi5+ and Ni2+ takes place under pressure, resulting in a Bi3+Ni3+O3 high-‐pressure (HP) phase that crystallizes into a √2ap × √2ap × 2ap ortho-‐ rhombic GdFeO3-‐type structure with a Pbnm space group and unique Bi site. The change in the Ni valence state from 2+ to 3+ leads to shrinkage of the Ni-‐O bond; the unit-‐cell volume decreases by 2.9%. Lanthanide (Ln = La, Nd, Eu, Dy) substitution for Bi or Fe substitution for Ni destabilizes the Bi3+/Bi5+ disproportionated state, and the

An intersite charge transfer transition was shown to cause volume shrinkage in A-‐site ordered per-‐ ovskite-‐type oxides upon heating.22,23

ACS Paragon Plus Environment

Chemistry of Materials 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 2 of 12

intersite charge transfer occurs upon heating.24-‐27 This transition is of first order, but since low temperature (LT) triclinic and high temperature (HT) orthorhombic phases coexist in varying fraction over a wide temperature range, the weighted average unit-‐cell volume decreases smooth-‐ ly.

at a weight ratio of 5:1 and sealed in Pt capsules. The sam-‐ ples were treated at 8 GPa and 1473 K for 30 min in a cu-‐ bic anvil-‐type HP apparatus. The remaining KCl was re-‐ moved by washing with distilled water. Sintered pellets for the HAXPES measurements were prepared by treating the sample powder at 6 GPa and 973 K for 10 min.

Bi1−xPbxNiO3 shows a triclinic-‐to-‐orthorhombic phase transition at x = 0.20.33,34 Recently, we reported a curious NTE behavior in Bi1−xPbxNiO3 with x ≤ 0.25.34 Samples with x = 0.05, 0.10 and 0.15 showed a triclinic-‐to-‐ orthorhombic phase transition accompanied by Bi5+ and Ni2+ charge transfer on heating, similar to what is obser-‐ ved in Bi1−xLnxNiO3 and BiNi1−xFexO3. On the other hand, samples with x = 0.20 and 0.25 showed an orthorhombic-‐ to-‐orthorhombic phase transition with a 0.6–0.7% volume shrinkage. A comprehensive study involving a Rietveld analysis of synchrotron X-‐ray diffraction (SXRD) data, Ni L-‐edge X-‐ray absorption spectroscopy (XAS), hard X-‐ray photoemission spectroscopy (HAXPES), and pair distribu-‐ tion function (PDF) analysis of the synchrotron X-‐ray total scattering data revealed that the LT orthorhombic phase had a short range ordering of Bi3+/Bi5+. Therefore, the LT phases of (Bi3+0.5Bi5+0.5)1−xPb4+xNi2+O3, triclinic for 0.05 ≤ x ≤ 0.15 and orthorhombic for 0.20 ≤ x, had the same valence distribution.

SXRD patterns were collected using a large Debye-‐ Scherrer camera installed at the BL02B239,40 and BL19B2 beamlines of SPring-‐8 and were analyzed using RIETAN-‐ FP programs.41 A wavelength of ~0.42 Å was used. The total scattering data for PDF analysis were collected using the diffractometer at the BL22XU beamline of SPring-‐8.42 Powder samples were loaded into cylinder polyimide ca-‐ pillaries that had a diameter of 1 mm. Monochromatized synchrotron X-‐rays of λ = 0.177926 Å were irradiated in the sample at RT. Scattering data were obtained up to the maximum momentum transfer Qmax of 25 Å−1 (2θmax = 45°). The scattering signal from the polyimide capillary was removed by subtracting the empty capillary data. Various other corrections were made, as shown in the reference.43 The PDF information was obtained by using the PDFgetX2 program.44 Local structural analyses were per-‐ formed by using the PDFgui program.45 The Qdamp and Qbroad values were calibrated by refining a CeO2 standard.

The perovskite phase of PbNiO3 with a Pb4+N2+O3 charge distribution was reported to crystalize into a Ge-‐ FeO3-‐type structure.35,36 However, the isotropic displace-‐ ment parameter of the oxygens was fixed at an unreason-‐ ably large value, i.e., B = 2.0 Å2 (Uiso = 0.0253 Å2), during the Rietveld refinement; the refined value for nickel, B = 1.50(3) Å2 (Uiso = 0.0190(4) Å2), was also unreasonably large. This suggests that the actual structure has a lower symmetry. In the present work, we investigated the evolu-‐ tion of crystal structure, valence distribution, thermal expansion and magnetic properties of Bi1−xPbxNiO3 with 0.30 ≤ x ≤ 1.00. All the synthesized samples had √2ap × √2ap × 2ap orthorhombic structures, but the samples with x ≥ 0.50, including PbNiO3, had polar Pbn21 symmetry. The samples with x ≥ 0.50 did not have the short-‐ranged ordering of disproportionated Bi3+/Bi5+ that has been ob-‐ served in samples with 0.25 ≤ x ≤ 0.40. The samples with 0.60 ≤ x ≤ 0.80 showed NTE, but it originated from a po-‐ lar-‐nonpolar transition, as in PbTiO3, instead of intersite charge transfer.

2. EXPERIMENTAL SECTION Precursors for Bi1−xPbxNiO3 (x = 0.25, 0.30, 0.35, 0.40, 0.50, 0.60, 0.70, 0.80 and 1.00) were prepared by using the pol-‐ ymerized complex method.23,37,38 Stoichiometric mixtures of Bi(NO3)3·∙5H2O, Pb(NO3)2 and Ni(NO3)2·∙6H2O were dis-‐ solved in aqueous solutions of citric acid and ethylene glycol. The molar ratios of citric acid and ethylene glycol to the mixtures were 10:1 and 30:1, respectively. The gels were obtained by stirring the solution for 12 h at 473 K in air. The gels were pyrolyzed for 2 h at 723 K in air by using a mantle heater. The obtained powders were annealed at 973 K for 12 h in air using a furnace after grinding with a mortar. The obtained precursors were mixed with KClO4

Pb-‐4f and Bi-‐4f HAXPES measurements were made at 300 K with E = 7940 eV at the BL09XU and BL47XU un-‐ dulater beamlines of SPring-‐8. A hemispherical photoe-‐ lectron analyzer (VG Scienta, R4000) was used in these measurements. The polycrystalline samples used in the HAXPES measurements were fractured in situ. The binding energies were calibrated using the Au 4f7/2 peak (84.0 eV) and the Fermi edge of gold reference samples. Total energy resolution was set to 280 meV at E = 7940 eV. Second harmonic generation (SHG) measurements of the powder sample were performed in the transparent geometry. The light source was a mode-‐locked Ti: sap-‐ phire regenerative amplified laser (pulse width about 120 fs, repetition rate of 1 kHz, photon energy of 1.58 eV). The sample was irradiated with the pulse, and the frequency doubled light (3.16 eV), i.e., the SHG signal, was detected by a photomultiplier after passing through high-‐pass fil-‐ ters and a grating-‐type monochromator to cut out the fundamental 1.58 eV pulses. The temperature dependence of the magnetic suscepti-‐ bility was measured with a SQUID magnetometer (Quan-‐ tum Design, MPMS 5) in an external magnetic field of 1 kOe. We used the first-‐principles density functional theory (DFT) + U method46 with the Perdew-‐Burke-‐Ernzerhof (PBE) form of the generalized gradient approximation (GGA) of the exchange-‐correlation functional47 to study the structural and magnetic phase stability of PbNiO3. The calculations were performed using the plane-‐wave pseudopotential method and projector augmented wave (PAW) potentials as implemented in the Vienna Ab initio Simulation Package (VASP).48,49 The rotationally invariant scheme of Dudarev et al.50 was used for the Hubbard U correction to GGA. We used a kinetic energy cut-‐off of


Page 3 of 12 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


Figure 1. Temperature dependences of unit cell volume calculated from the SXRD patterns confirming the orthorhombic-‐to-‐ orthorhombic phase transition accompanied by NTE. (a) Temperature variation of the magnified SXRD patterns around the prominent peaks of Bi1−xPbxNiO3 (x = 0.10, 0.25, 0.50, 0.60 and 0.80) upon heating. The data of x = 0.10 and 0.25 are reprinted 34 with from the previous report. λ = 0.42002 Å for x = 0.10, λ = 0.41996 Å for x = 0.25, λ = 0.41842 Å for x = 0.50, λ = 0.42091 Å for x = 0.60 and λ = 0.42011 Å for x = 0.80. LN represents the reflection of the LiNbO3-‐type phases. (b) Composition-‐temperature phase diagram of Bi1−xPbxNiO3 obtained by the variable temperature SXRD studies at ambient pressure (AP). (c) Temperature dependence of the weighted average unit cell volume of Bi1−xPbxNiO3.

520 eV to expand the wave functions for all calculations and optimized the k-‐point mesh following the crystal symmetry. The Hellman–Feynman forces converged to 0.01 eV Å−1. We performed calculations considering an effective U from 2.0 eV to 8.0 eV for the Ni 3d state. We calculated the electric polarization for the polar phases by using the modern theory of polarization.51-‐53

3. RESULTS AND DISCUSSION Structural transformation. Figure 1a shows the temper-‐ ature variation upon heating of the SXRD patterns of Bi1− xPbxNiO3 with x = 0.10, 0.25, 0.50, 0.60 and 0.80. Tempera-‐ ture-‐induced triclinic-‐to-‐orthorhombic phase transitions were observed for x = 0.10 on heating at around 500 K. The x = 0.25 sample showed an unusual shift in the peak to a lower angle at 550 K, indicating an orthorhombic-‐to-‐ orthorhombic transition. These were NTEs induced by charge transfer, as previously reported. The SXRD pat-‐ terns of Bi1−xPbxNiO3 with x > 0.25 are indexed with a √2ap × √2ap × 2ap orthorhombic unit cell, the same as the x = 0.25 sample. However, the sample with x ≥ 0.50 showed a new temperature-‐induced orthorhombic-‐to-‐

orthorhombic phase transition, where the transition tem-‐ perature increased as x increased. A large shift in the peak for the 200 reflection of the x = 0.6 sample between 400 and 500 K and the patterns at 420-‐480 K showing that LT and HT phases coexist indicate that this phase transition is of first order. Therefore, the crystal structure at RT changes from a triclinic one to an orthorhombic one, and finally to another orthorhombic one. Figure 1b shows the composition-‐temperature phase diagram of Bi1−xPbxNiO3. It should be noted that this is not a true thermodynamic phase diagram because the samples synthesized at HP and HT condition and recovered to ambient condition are metastable at AP. We name the LT and HT phases of Bi0.75Pb0.25NiO3 orthorhombic phases I and II, and the new orthorhombic phase observed in Bi0.4Pb0.6NiO3 at LT or-‐ thorhombic phase III. In addition to the orthorhombic-‐ to-‐orthorhombic phase transitions, the x = 0.8 sample showed a temperature-‐induced irreversible phase transi-‐ tion from perovskite-‐type to LiNbO3-‐type structure, as has been observed in PbNiO3.36 On the other hand, the orthorhombic I-‐II and II-‐I transitions are both reversible. Figure 1c shows the temperature dependence of the



Figure 2. Temperature dependence of the lattice parameters of Bi1−xPbxNiO3. The orthorhombic lattice parameters are normalized to a pseudo-‐cubic subcell with ap = a/√2, bp = b/√2 and cp = c/2. In the orthorhombic I-‐II transition, shrinkage of the b-‐ and c-‐axes and expansion of a-‐axes occur upon heating. On the other hand, the a-‐ and b-‐axes shrink and the c-‐axis expand upon heating in the orthorhombic III-‐I transition.

weighted average unit-‐cell volumes, fLTVLT + fHTVHT (where fLT(HT) represents the phase fraction of the LT(HT) phase and VLT(HT) represents the unit-‐cell volume of the LT(HT) phase) for Bi1−xPbxNiO3 with 0.20 ≤ x ≤ 0.80. NTE occurred in the two composition ranges of 0.20 ≤ x ≤ 0.25 and 0.60 ≤ x ≤ 0.80, while the x = 0.30 and 0.50 samples showed almost zero thermal expansion. The difference between the orthorhombic phases I and II was in the va-‐ lence distribution and local structure, namely, phase I had a (Bi3+0.5Bi5+0.5)1−xPb4+xNi2+O3 valence distribution with short-‐ranged Bi3+/Bi5+ ordering, whereas phase II had a (Bi, Pb)3+Ni3+O3 valence distribution without such a charge ordering. NTE occurs owing to the intersite charge trans-‐ fer between Bi5+ and Ni2+. The changes in the lattice pa-‐ rameters in the III-‐I transition show a different anisotropy from those in the I-‐II transition. Figure 2 shows the tem-‐ perature dependence of the lattice parameters of Bi1−xPbxNiO3 with 0.20 ≤ x ≤ 0.80. The orthorhombic lat-‐ tice parameters are normalized to pseudo-‐cubic subcell. In the I-‐II transition, shrinkage of the b-‐ and c-‐axes and expansion of the a-‐axis occurred upon heating. On the other hand, in the III-‐I transition, the a-‐ and b-‐axes shrunk and the c-‐axis expanded upon heating. If this transition is accompanied by an intersite charge transfer, orthorhombic phase III should have a Bi5+1−xPb4+xNi(1+x)+O3 valence distribution. Structure refinement. To elucidate the valence distri-‐ bution of orthorhombic phase III and the mechanism of NTE in the III-‐I transition, a Rietveld analysis was per-‐ formed on the SXRD data. The reflection conditions of 0kl (k = 2n), h0l (h + l = 2n), h00 (h = 2n), 0k0 (k = 2n) and 00l (l = 2n) in the orthorhombic phase III afford two possible space groups, nonpolar Pbnm and polar Pbn21. Figure 3a shows the laser power dependence of the SH intensity for Bi0.2Pb0.8NiO3. (In this measurement, we used the x = 0.8

Page 4 of 12

sample instead of PbNiO3 to investigate the polar sym-‐ metry in the orthorhombic phase III, because perovskite-‐ type PbNiO3 gradually changes into the polar LiNbO3-‐ type one even under ambient conditions, whereas the x = 0.8 sample retains a perovskite-‐type structure at RT.) The SH intensity was found to be proportional to the square of the incident laser power, clearly indicating breaking of the inversion symmetry of the x = 0.8 sample. The results of the Rietveld refinement assuming the Pbn21 structure model for Bi1−xPbxNiO3 with x = 0.60, 0.80 and 1.00 at 300 K are shown in Figure 3b and the crystallographic param-‐ eters are summarized in Table 1. Satisfactory small R fac-‐ tors and reasonable isotropic displacement parameters (Uiso) were obtained. It should be noted that the refine-‐ ment assuming nonpolar Pbnm structure model resulted in unusually large Uiso values for Ni and O2 sites (see Ta-‐ ble 1). Figure 3c shows a schematic representation of the refined Pbn21 structure and that of Pbnm for comparison. Viewing from the c-‐axis direction, the NiO6 octahedron in adjacent planes overlap in the Pbnm structure, whereas they don’t in the Pbn21 structure. In other words, the O2 site of the Pbnm structure splits into O2 and O3 sites in the Pbn21 structure. The unusually large Uiso value for the O2 site observed in the refinement with the Pbnm model should be attributed to this crystallographic site splitting. P-‐E measurement was attempted on Bi0.4Pb0.6NiO3, but ferroelectric loop was not observed because of relatively large leakage current. Valence distribution. The results of bond valence sum (BVS) calculations at RT indicates the formal valence state of (Bi, Pb)4+Ni2+O3 in all the compositions, despite the unstable nature of the Bi4+ state (see Figure 4a). HAX-‐ PES measurements were performed to investigate the possible disproportionation into Bi3+ and Bi5+ without long-‐range ordering. Figure 4b and 4c compare the HAX-‐ PES spectra for Bi1−xPbxNiO3 with data for Bi and Pb standard samples. The intense peaks at 163.2 and 157.9 eV in the Bi 4f data for BiNiO3 that are absent from the data for Bi3+Fe3+O3 indicate the presence of Bi5+ in addition to Bi3+. The Bi 4f spectrum for Bi1−xPbxNiO3 is similar to that of BiNiO3, indicating coexistence of Bi3+ and Bi5+ valence states. It should be noted that the 6s0 electronic configu-‐ ration (Bi5+ or Pb4+) leads to lower binding energy than the 6s2 configuration (Bi3+ or Pb2+) because of a strong screening effect.34,54 On the other hand, the Pb 4f spec-‐ trum consists of a single peak, and the peak position is rather closer to that of Pb4+Ni2+O3 than that of Pb2+Ti4+O3, indicating that the valence state of Pb is Pb4+. Figure 4d shows the composition dependence of the Bi3+ ratio esti-‐ mated from the area ratio of the HAXPES peaks. The Bi3+ ratio remains one half for all compositions, indicating the formal valence state of (Bi3+0.5Bi5+0.5)1−xPb4+xNi2+O3, the same as in the x = 0.20 and 0.25 samples. Focusing on the width of Bi3+ and Bi5+ peaks, the Bi3+ ones become nar-‐ rower as the amount of Pb increases, whereas Bi5+ ones become broader. This trend may be thought of as follows. The inhomogeneous potential energy of the Bi3+ site stemming from the lone-‐pair effect makes the peak broad, but the effect becomes weak when the amount of Bi is


Page 5 of 12 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


Figure 3. SHG measurement and Rietveld refinement indicating polar symmetry in the orthorhombic phase III. (a) Laser power dependence of the SH intensity for Bi0.2Pb0.8NiO3. (b) Observed (red points), calculated (green lines), and difference (blue lines) patterns from the Rietveld analysis of the SXRD data of the Bi1−xPbxNiO3 with x = 0.60 and 0.80 and 1.00 at 300 K. The green tick marks in the x = 0.60 and 0.80 samples and the top tick marks in the x = 1.00 sample correspond to the positions of Bragg reflec-‐ tions of the Pbn21 perovskite-‐type phase, and the bottom tick marks in the x = 1.00 sample correspond to that of the R3c LiNbO3-‐ type phase. (c) Schematic representation of the polar Pbn21 structure and the Pbnm structure models viewed from the c-‐axis direction.

small. On the other hand, the high valence Bi5+ site, which is considered to share oxygen holes with Ni site,55,56 has inhomogeneous potential energy when the oxygen holes are biased by the Ni off-‐center displacement in the polar symmetry. The broadening of the Pb4+ peak with increas-‐ ing the amount of Pb can be explained by the same mech-‐ anism as that of the Bi5+ peak. Local structure characterization. We examined the local structures in the orthorhombic phase III by perform-‐ ing a PDF analysis of the synchrotron X-‐ray total scatter-‐ ing data in order to investigate the possible short-‐ranged Bi3+/Bi5+ charge ordering in 0.25 ≤ x ≤ 0.30 samples. Figure 5 shows the fitting results of the PDF analysis for Bi0.2Pb0.8NiO3 at low r range (1-‐20 Å). We employed a nonpolar P-‐1 triclinic structural model with two Bi/Pb sites, which is the local structure of orthorhombic phase I, in addition to the Pbnm and Pbn21 ones. The isotropic displacement parameter of each element was fixed at Ui-‐ 2 2 so(Bi) = Uiso(Pb) = Uiso(Ni) = 0.00633 Å (B = 0.5 Å ) and 2 2 Uiso(O) = 0.01267 Å (B = 1.0 Å ) for all models. The initial analysis assuming the Pbnm structure model gave a poor reliability factor, RWP = 12.47%. Employing the polar Pbn21 model significantly decreased RWP to 7.52%. On the other hand, the nonpolar P-‐1 triclinic model provided only a minor improvement (RWP = 9.64%). These results indicate the presence of a polar local structure and the absence of Bi3+/Bi5+ short-‐range ordering in the orthorhombic phase III. The absence of the Bi3+/Bi5+ ordering is reasonable

because the Bi occupancy of the Bi/Pb site is only 20%, as will be discussed later. Therefore, considering the presen-‐ ce of polar symmetry in the orthorhombic phase III struc-‐ ture and the absence of intersite charge transfer upon heating, NTE occurring at the III-‐I transition should be due to the polar-‐nonpolar phase transition, as in PbTiO3. Such a temperature-‐induced polar-‐nonpolar phase transi-‐ tion from Pbn21 to Pbnm has been reported in perovskite-‐ type CdTiO3 at 77 K.57,58 These results strongly suggest that PbNiO3 also has a polar Pbn21 structure. However, NTE was absent in PbNiO3 because the perovskite-‐type structure changed into the LiNbO3-‐type one upon heating before a transition to the Pbnm structure. DFT Calculations. A striking feature of our present study is the stabilization of the polar Pbn21 phase for one of the end members, high pressure stabilized PbNiO3. Interestingly, previous studies have identified the high-‐ pressure phase as a non-‐polar Pbnm structure.36,59 In or-‐ der to understand the stability of this polar phase, we first determined the ground state structure by calculating the energies of the experimentally reported LiNbO3-‐type R3c, Pbnm and Pbn21 phases, and various other structures identified on the basis of unstable phonon distortions at the zone center (q = 0) and at the zone boundary (q ≠ 0) symmetry points of the Brillouin zone for the undistorted Pm-‐3m structure. We carried out full structural optimiza-‐ tions considering various collinear ferromagnetic (FM) and antiferromagnetic (AFM) arrangements between Ni



Table 1. Structural Parameters for Perovskite-‐type Bi1−xPbxNiO3 at 300 K atoms

site

g

x

y

2

z

100×Uiso(Å )

−0.7627(3)

0.86(1)

0.5005(12) −0.0116(8) 0 0.6275(13) −0.0675(15) −0.7618(52) 0.3353(20) 0.3565(20) −0.5483(17) 0.8008(20) 0.2182(21) −0.4210(14) b x = 0.8, Space Group Pbn21

0.78(4) 1.84(20) 1.35(30) 1.00(29)

a

x = 0.6, Space Group Pbn21 Bi Pb Ni O1 O2 O3 Bi Pb Ni O1 O2 O3

4a 4a 4a 4a 4a 4a 4a 4a 4a 4a

0.4 0.6 1 1 1 1 0.2 0.8 1 1 1 1

0.0111(1)

0.0476(1)

0.0085(1)

0.0445(1)

−0.7663(3)

0.56(1)

0.4999(9) 0.6260(9) 0.3470(14) 0.7968(14)

−0.0169(5) −0.0721(12) 0.3719(14) 0.2216(17)

0 −0.7514(23) −0.5447(12) −0.4190(10)

0.58(3) 1.52(21) 1.04(23) 0.76(21)

c

Pb Ni O1 O2 O3

4a 4a 4a 4a 4a

1 1 1 1 1

x = 1.0, Space Group Pbn21 0.0079(1) 0.0417(1) −0.7691(2) 0.5014(6) −0.0203(4) 0 0.6259(10) −0.0793(10) −0.7436(12) 0.3395(14) 0.3683(14) −0.5439(11) 0.7951(14) 0.2232(16) −0.4145(10)

0.51(1) 0.65(3) 0.26(16) 0.84(19) 1.06(19)

d

x = 0.6, Space Group Pbnm Bi Pb Ni O1 O2

4c 4a 4c 8d

0.4 0.6 1 1 1

Bi Pb Ni O1 O2

4a 4c 8d

0.2 0.8 1 1 1

Pb Ni O1 O2

4c 4a 4c 8d

1 1 1 1

4c

0.0112(1)

0.4524(1)

0.25

0.84(1)

0 0 0 0.6280(13) 0.5704(15) 0.25 0.3164(13) 0.1789(13) 0.0654(8) e x = 0.8, Space Group Pbnm

1.33(3) 1.62(20) 4.90(24)

0.0087(1)

0.4554(1)

0.25

0.54(1)

0 0.6240(11) 0.3212(12)

0 0.5780(13) 0.1673(13)

0 0.25 0.0649(8)

1.46(3) 1.38(20) 6.56(27)

f

x = 1.0, Space Group Pbnm 0.0080(1) 0.4583(1) 0.25 0 0 0 0.6257(12) 0.5881(14) 0.25 0.3213(14) 0.1640(15) 0.0616(9)

0.47(1) 1.93(3) 1.35(18) 6.92(28)

a

Z = 4; a = 5.34660(1) Å, b = 5.52299(2) Å, c = 7.72237(2) Å; b Rwp = 6.97%, Rp = 4.93%, RB = 1.57%, RF = 0.97%, S = 2.24. Z = 4; a = 5.35698(1) Å, b = 5.49382(1) Å, c = 7.71333(1) Å; Rwp = c 3.57%, Rp = 2.65%, RB = 1.28%, RF = 0.77%, S = 1.74. Z = 4; a = 5.35800(1) Å, b = 5.46365(1) Å, c = 7.70772(1) Å; Rwp = 6.72%, Rp = 4.70%, RB = 1.24% for perovskite-‐type and 1.15% for LiNbO3-‐type, RF = 0.59% for perovskite-‐type and 0.43% for d LiNbO3-‐type, S = 2.87. Z = 4; a = 5.34639(2) Å, b = 5.52280(2) Å, c = 7.72208(2) Å; Rwp = 7.37%, Rp = 5.25%, RB = 2.02%, RF = e 1.24%, S = 2.37. Z = 4; a = 5.35651(1) Å, b = 5.49332(1) Å, c = 7.771263(2) Å; Rwp = 4.45%, Rp = 3.32%, RB = 2.66%, RF = 1.76%, f S = 2.17. Z = 4; a = 5.35805(1) Å, b = 5.46372(1) Å, c = 7.70778(1) Å; Rwp = 8.85%, Rp = 6.43%, RB = 2.24% for perov-‐ skite-‐type and 1.89% for LiNbO3-‐type, RF = 1.06% for perov-‐ skite-‐type and 0.72% for LiNbO3-‐type, S = 3.77.

spins. The relaxed energetics for the PbNiO3 system for the lowest energy spin structure is depicted in Figure 6a. We found that the ground state structure crystallizes with LiNbO3-‐type polar R3c symmetry and the Ni spins are ordered in an AFM G-‐type pattern, in agreement with the previous report.59 The R3c phase was found to be an insu-‐

Page 6 of 12

lator with an energy gap between O 2p states and Pb 6s states hybridized with O 2p orbitals (see Figure S1). The Ni 3d states form conduction bands in the energy range ~ −6.0 to −2.0 eV in one spin channel, while in the other spin channel, the t2g states are located between ~ −4.0 eV and the Fermi level and the eg states form conduction bands (see Figure S1). This indicates a +2 nominal oxida-‐ tion state for Ni (d8: t2g6eg2), where the computed mag-‐ netic moment is ~1.8 µB. On the other hand, the Pb 6s states hybridize strongly with the O 2p states with stabili-‐ zation of the hole on the oxygen, as was observed in β-‐ PbO257 and pentavalent Bi compounds.31,61 The basic na-‐ ture of the electronic structure remains the same for all the phases under the present conditions. The polar R3c structure is composed of an a−a−a− octahedral tilt distor-‐ tion that has the R5− symmetry and a polar distortion that has the Γ4− symmetry of the Pm-‐3m structure. The esti-‐ mated electric polarization is ~35 µC/cm2 directed along the cubic [111] axis. The Pbn21 and Pbnm structures are respectively 55 meV/f.u. and 85 meV/f.u. higher in energy compared to the lowest energy polar R3c phase (see Fig-‐ ure 6a), with a smaller cell volume. The details of the op-‐ timized structures are provided in Table S1. Our calcula-‐ tion results explain why PbNiO3 tends to crystallize in the R3c structure under ambient conditions. Next, we studied the changes in the structural phase stability induced by the volume reduction, as is expected to occur under hydrostatic pressure. To this end, we cal-‐ culated the energies of R3c and the low-‐energy Pbn21 and Pbnm phases for a range of uniformly varied volumes tak-‐ ing into consideration the c/a ratio of the fully optimized structures. The atomic positions were optimized for each volume. The results are summarized in Figure 6b. We observed a phase transition from the polar R3c to the po-‐ lar Pbn21 structure at around a 4.9% reduction in volume and at around 3.2 GPa applied pressure (see Figure S3). Here, although the energy difference between Pbn21 and Pbnm decreased as the volume decreased, no further transformation of phase was observed even when the cell volume was reduced by 12%. This indicates that under high pressure, the Pbn21 structure should be stabilized, as was observed in the high-‐pressure synthesis. This howev-‐ er, has not been reported in the previous studies.33,56 We cross validated our results by taking into account a range of U values at the Ni 3d states (see Figure S2). The prima-‐ ry phonon distortions that contribute to the Pbn21 struc-‐ ture are as follows: (i) an a−a−c0 octahedral tilt distortion that follows R5−, (ii) an a0a0c+ octahedral rotation that follows M2+ and (iii) a polar distortion that follows Γ4− of the Pm-‐3m structure. Note that the Pbnm phase is primar-‐ ily composed of the octahedral a−a−c0 tilt and a0a0c+ rota-‐ tional distortions. The Pbn21 structure is expected to be stabilized by the Γ4− polar distortion, which is unstable in the Pm-‐3m structure. The estimated electric polarization for the fully optimized Pbn21 structure is ~44 µC/cm2 along the cubic [001] axis, in good agreement with the value calculated from the experimentally determined structural parameters (37.8 µC/cm2) in Table 1.


Page 7 of 12 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


3+

5+

4+

2+

Figure 4. HAXPES measurements indicating the formal valence state of (Bi 0.5Bi 0.5)1−xPb xNi O3. (a) Composition depend-‐ ence of BVSs for Bi1−xPbxNiO3 at RT. (b) Bi 4f HAXPES results for Bi1−xPbxNiO3 at RT together with those for BiFeO3 and BiNiO3 3+ 3+/5+ as standard materials for Bi and Bi . (c) Pb 4f HAXPES results for Bi1−xPbxNiO3 at RT together with those for PbTiO3 and 2+ 4+ 3+ PbNiO3 as standard materials for Pb and Pb . (d) Composition dependence of the fraction of Bi calculated from the area ratio of HAXPES peaks.

Magnetic property. We performed magnetic meas-‐ urements to evaluate the spin and valence states of Ni. Figure 7a and S4 shows the temperature dependence of magnetic susceptibility and the magnetization curves at 100 K for Bi1−xPbxNiO3. Weak ferromagnetism, most prob-‐ ably due to the canted spins induced by a Dzyalonshin-‐ sky-‐Moriya interaction, was observed in the samples with 0.25 ≤ x ≤ 0.80. The data between 370 and 400 K, well above the Néel temperatures (TN), can be fitted to the Curie-‐Weiss law with a temperature independent term χ0, χ = C/(T − θ) + χ0. Here, C is the Curie constant, and θ is the Weiss temperature. Figure 7b shows the composition dependence of TN and the effective magnetic moments (peff) calculated from the Curie constant. The peff values of 2.89 -‐ 3.13 µB are close to 2.83 µB for Ni2+ with a high spin configuration with S = 1 and g = 2, indicating that Ni re-‐ mains divalent in all the compositions. On the other hand, TN is almost independent of x until x = 0.40, at which point it begins to decrease with increasing x. These find-‐ ings are consistent with the results of structural analysis showing that Ni remains divalent and that the polar-‐ nonpolar phase transition occurs around x = 0.40 at RT. TN decreases for x ≥ 0.60, suggesting a decrease in the antiferromagnetic interaction owing to the reduction of the Ni-‐O-‐Ni angles. Indeed, the average Ni-‐O-‐Ni angles (θ) calculated from the Rietveld refinement for x ≥ 0.60 at 100 K (see Table S2), well below TNs, systematically de-‐ crease with increasing x: θ = 137.0° (x = 0.6), 135.9° (x = 0.8), and 135.5° (x = 1.0). In the polar structure, the magni-‐ tude of the antiferromagnetic interaction changes be-‐

cause of the off-‐center displacement of Ni in the NiO6 octahedron, which is absent in the nonpolar structure. Phase stability. Finally, we discuss the experimentally and theoretically determined stability of the polar struc-‐ ture. Figure 8 is a schematic illustration of the Bi/Pb layer showing the successive phase transitions of Bi1−xPbxNiO3 depending on the amount of Bi. As the amount of Bi de-‐ creases, the crystal structure changes from a triclinic one with Bi3+/Bi5+ long-‐range ordering to an orthorhombic one with Bi3+/Bi5+ short-‐range ordering; then it changes into a polar orthorhombic one without Bi3+/Bi5+ ordering and finally to a polar LiNbO3-‐type one, which is closely related to the perovskite-‐type one.36 Under the HP synthesis con-‐ dition, Bi is all trivalent and partial ordering of Bi3+ and Pb4+ occurs because of the differences in the valence and the ionic radius. When the sample is cooled and the pres-‐ sure is released, the valence state changes to (Bi3+0.5Bi5+0.5)1−xPb4+xNi2+O3, where Bi3+/Bi5+ disproportiona-‐ tion occurs. However, when x exceeds the critical concen-‐ tration of 0.2, the development of Bi3+/Bi5+ long-‐range ordering is hindered by the presence of small Pb4+ at the large Bi3+ site.34 The disappearance of Bi3+/Bi5+ short-‐range ordering can be similarly understood considering the de-‐ crease in the number of Bi ions. Moreover, the increased stability of the LiNbO3-‐type structure can be understood in terms of Goldschmidt’s tolerance factor (t), which has been used as an indicator for the stability of perovskite-‐ type structures. The ionic radii of Shannon (r) for six-‐fold coordination are rPb4+ = 0.775 Å, rBi = (rBi5+ + rBi3+)/2 = 0.895 Å, rNi2+ = 0.69 Å and rO2− = 1.40 Å, leading to t of 0.78



Figure 5. Observed (red points), calculated (blue line), and difference (green line) PDFs for Bi0.2Pb0.8NiO3 with nonpolar Pbnm orthorhombic, nonpolar P-‐1 triclinic and polar Pbn21 orthorhombic structural models at RT. ~ 4.9%

(b)

(a) 0.8 0.6

P1

P1 R3c

0.4 Ima2

0.2 0 -0.2

Im3

Pbn21

Imma Pbnm !3c

Energy (eV/f.u.)

I4/mmm

Energy (eV/f.u.)

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 12

R3c Pbn21 Pbnm

0.2

2+

0.15 0.1 0.05 0 54

56

58

60

62

64

66

3

Volume (Å /f.u.)

Figure 6. DFT calculations indicating the ground states of LiNbO3-‐type R3c and perovskite-‐type Pbn21 structures under AP or HP conditions. (a) Computed energies of various dis-‐ torted PbNiO3 structures with respect to lowest energy R3c phase. The phases with energy within 1 eV are shown. (b) Computed total energy versus volume per formula unit for the R3c, Pbn21 and Pbnm phases. The results are for the lowest energy AFM G-‐type spin structure and for the effec-‐ tive U = 6.0 eV at the Ni site. The results for the other U val-‐ ues are provided in Figure S2. (x = 0), 0.76 (x = 0.5) and 0.74 (x = 1.0) for Bi1−xPbxNiO3. The perovskite-‐type structure is expected to be stable within the 62,63 whereas the LiNbO3-‐type structure is limits 0.75 < t < 1.1, stable for t < 0.75. The calculated t values are near the boundary of the perovskite-‐type and LiNbO3-‐type structures. The inversion symmetry breaking in perovskite-‐type PbNiO3 2 seems peculiar because there are neither 6s lone pairs nor second-‐order Jahn-‐Teller (JT) active cations. However, polar materials without these cations have been found in the 64 65 LiNbO3-‐type InFeO3 (t = 0.76) and ScFeO3 (t = 0.74), and the driving force for the inversion symmetry breaking is con-‐ sidered to be the increased Coulomb interactions between 64,66 the high valency A-‐site cation and oxygen.

Figure 7. Magnetic measurements indicating Ni high spin state and Ni off-‐center displacement in the NiO6 octahedron. (a) Temperature dependence of molar magnetic susceptibil-‐ ity for Bi1−xPbxNiO3. (b) Composition dependence of the Néel temperature (TN) and effective magnetic moment (peff). The peff were obtained from the fittings to the Curie-‐Weiss law.

As the A-‐O Coulomb interactions increase, the R-‐3c struc-‐ ture with a nine-‐coordinated A-‐site cation becomes un-‐ stable and changes into an R3c one with six-‐coordinations through polar displacement of A-‐site cation. The large A-‐ O Coulomb interactions can also be expected to stabilize the polar symmetry in the orthorhombic perovskite-‐type materials, from analogy with LiNbO3-‐type case. The Γ4− polar distortion is indeed found to be unstable for the Pm-‐3m structure of PbNiO3. Therefore, the occurrence of the polar orthorhombic perovskite-‐type structure should be reasonable considering the high valency state of Pb4+ and the small t value close to those of LiNbO3-‐type mate-‐ rials and the polar orthorhombic perovskite-‐type BiInO3 (t = 0.78).67

4. CONCLUSION The structural analysis using SXRD revealed the presence of the three different orthorhombic phases in Bi1−xPbxNiO3 and temperature-‐induced orthorhombic-‐to-‐orthorhombic phase transitions accompanied by NTE in two composi-‐ tion ranges, 0.20 ≤ x ≤ 0.25 and 0.60 ≤ x ≤ 0.80. The for-‐ mer NTE is induced by intersite charge transfer between Bi5+ and Ni2+, as has already been clarified in a previous


Page 9 of 12 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


Figure 8. Illustration of the Bi/Pb layer showing systematic phase transition of Bi1−xPbxNiO3 according to the amount of Bi. As 3+ 5+ the amount of Bi decreases, the crystal structure changes from a triclinic one with Bi /Bi long-‐range ordering to an ortho-‐ 3+ 5+ 3+ 5+ rhombic one with Bi /Bi short-‐range ordering; then it changes into polar orthorhombic one without Bi /Bi ordering and a polar LiNbO3-‐type one.

study. The results of the Rietveld refinements, PDF anal-‐ yses, and SHG measurements revealed that the LT ortho-‐ rhombic phase with 0.40 ≤ x ≤ 1.00 has a polar structure with space group Pbn21. The latter NTE originates from a polar-‐nonpolar transition. The change in the magnetic properties observed at x = 0.50 is consistent with off-‐ center displacement of the Ni in NiO6 octahedron due to the polar nature. Perovskite-‐related materials with small tolerance factors have attracted much attention because they tend to have a polar structure. However, most polar materials found in the region of small tolerance factor have LiNbO3-‐type structures; only a few polar ortho-‐ rhombic perovskite-‐type structures have been found, which is probably due to the high stability of LiNbO3-‐type structure in this region. Our results suggest that the polar orthorhombic phase can be realized by using high valence A-‐site cations in addition to controlling the tolerance fac-‐ tor. This approach will be new strategy for searching for functional materials having a polar nature, for example, ferroelectric and large NTE materials.

ASSOCIATED CONTENT Supporting Information The Supporting Information is available free of charge on the ACS Publications website at DOI: Computed G-‐type antiferromagnetic density of states (DOS) for the R3c phase; Computed total en-‐ ergy versus cell volume per formula unit for R3c, Pbn21 and Pbnm phases; Calculated enthalpy as a function of applied pressure for R3c, Pbn21 and Pbnm phases; The magnetization curves at 100 K for Bi1−xPbxNiO3; Optimized structural parameters for PbNiO3 for the ground state R3c structure and low

energy Pbn21 and Pbnm structures; Structural Pa-‐ rameters for Perovskite-‐type Bi1-‐xPbxNiO3 at 100 K (PDF)

AUTHOR INFORMATION Corresponding Author * [email protected] * [email protected]

Notes

The authors declare no competing financial interest.

ACKNOWLEDGMENT This work was partially supported by the Grant-‐in-‐Aid for Scientific Research 16H02393 and 18H05208 from the Japan Society for the Promotion of Science (JSPS) and by the Pho-‐ ton and Quantum Basic Research Coordinated Development Program of the Ministry of Education, Culture, Sports, Scien-‐ ce and Technology (MEXT), Japan. The synchrotron-‐ radiation experiments were performed at SPring-‐8 with the approval of the Japan Synchrotron Radiation Research Insti-‐ tute (2017A1242, 2017A1388, 2017B3751, 2017B1721, 2018A1642, 2018A3751, 2018A1667, 2018B3751, 2018B1797, 2018B1672 and 2018B1860).

REFERENCES (1) Chu C. N.; Saka, N.; Suh, N. P. Negative thermal expansion ceramics: A review. Mater. Sci. Eng. 1987, 95, 303-‐308. (2) Sleight, A. W., Compounds That Contract on Heating. In-‐ org. Chem. 1998, 37, 2854-‐2860. (3) Barrera, G. D.; Bruno, J. A. O.; Barron, T. H. K.; Allan, N. L. Negative thermal expansion. J. Phys. Condens. Matter 2005, 17, R217-‐R252.



(4) Takenaka, K. Negative thermal expansion materials: tech-‐ nological key for control of thermal expansion. Sci. Technol. Adv. Mater. 2012, 13, 013001. (5) Chen, J.; Hu, L.; Deng, J.; Xing, X. Negative thermal expan-‐ sion in functional materials: controllable thermal expansion by chemical modifications. Chem. Soc. Rev. 2015, 44, 3522-‐3567. (6) Mittal, R.; Gupta, M. K.; Chaplot, S. L. Phonons and anom-‐ alous thermal expansion behaviour in crystalline solids. Prog. Mater. Sci. 2018, 92, 360-‐445. (7) Attfield, J. P. Mechanisms and Materials for NTE. Front. Chem. 2018, 6, 371. (8) Mary, T. A.; Evans, J. S. O.; Vogt, T.; Sleight, A. W. Nega-‐ tive Thermal Expansion from 0.3 to 1050 Kelvin in ZrW2O8. Sci-‐ ence 1996, 272, 90-‐92. (9) Phillips, A. E.; Goodwin, A. L.; Halder, G. J.; Southon, P. D.; Kepert, C. J. Nanoporosity and Exceptional Negative Thermal Expansion in Single-‐Network Cadmium Cyanide. Angew. Chem. Int. Ed. 2008, 47, 1396-‐1399. (10) Takenaka, K.; Takagi, H. Giant negative thermal expan-‐ sion in Ge-‐doped anti-‐perovskite manganese nitrides. Appl. Phys. Lett. 2005, 87, 261902. (11) Takenaka, K.; Asano, K.; Misawa, M.; Takagi, H. Negative thermal expansion in Ge-‐free antiperovskite manganese nitrides: Tin-‐doping effect. Appl. Phys. Lett. 2008, 92, 011927. (12) Takenaka, K.; Inagaki, T.; Takagi, H. Conversion of mag-‐ netic structure by slight dopants in geometrically frustrated an-‐ tiperovskite Mn3GaN. Appl. Phys. Lett. 2009, 95, 132508. (13) Matsuno, J.; Takenaka, K.; Takagi, H.; Matsumura, D.; Nishihata, Y.; Mizuki, J. Local structure anomaly around Ge do-‐ pants in Mn3Cu0.7Ge0.3N with negative thermal expansion. Appl. Phys. Lett. 2009, 94, 181904. (14) Hamada, T.; Takenaka, K. Giant negative thermal expan-‐ sion in antiperovskite manganese nitrides. J. Appl. Phys. 2011, 109, 07E309. (15) Tong, P.; Louca, D.; King, G.; Llobet, A.; Lin, J. C.; Sun, Y. P. Magnetic transition broadening and local lattice distortion in the negative thermal expansion antiperovskite Cu1 − xSnxNMn3. Appl. Phys. Lett. 2013, 102, 041908. (16) Takenaka, K.; Hamada, T.; Kasugai, D.; Sugimoto, N. Tai-‐ loring thermal expansion in metal matrix composites blended by antiperovskite manganese nitrides exhibiting giant negative thermal expansion. J. Appl. Phys. 2012, 112, 083517. (17) Chen, J.; Xing, X.; Sun, C.; Hu, P.; Yu, R.; Wang, X.; Li, L. Zero Thermal Expansion in PbTiO3-‐Based Perovskites. J. Am. Chem. Soc. 2008, 130, 1144-‐1145. (18) Hu, P.; Chen, J.; Sun, X.; Deng, J.; Chen, X.; Yu, R.; Qiao, L.; Xing, X. Zero thermal expansion in (1−x)PbTiO3– xBi(Mg,Ti)1/2O3 piezoceramics. J. Mater. Chem. 2009, 19, 1648-‐ 1652. (19) Chen, J.; Nittala, K.; Forrester, J. S.; Jones, J. L.; Deng, J.; Yu, R.; Xing, X. The Role of Spontaneous Polarization in the Negative Thermal Expansion of Tetragonal PbTiO3-‐Based Com-‐ pounds. J. Am. Chem. Soc. 2011, 133, 11114-‐11117. (20) Chen, J.; Fan, L.; Ren, Y.; Pan, Z.; Deng, J.; Yu, R.; Xing, X. Unusual Transformation from Strong Negative to Positive Ther-‐ mal Expansion in PbTiO3-‐BiFeO3 Perovskite. Phys. Rev. Lett. 2013, 110, 115901. (21) Yamamoto, H.; Imai, T.; Sakai, Y.; Azuma, M. Colossal Negative Thermal Expansion in Electron-‐Doped PbVO3 Perovskites. Angew. Chem. Int. Ed. 2018, 57, 8170-‐8173. (22) Long, Y. W.; Hayashi, N.; Saito, T.; Azuma, M.; Muranaka, S.; Shimakawa, Y. Temperature-‐induced A-‐B intersite charge transfer in an A-‐site-‐ordered LaCu3Fe4O12 perovskite. Nature 2009, 458, 60-‐63. (23) Yamada, I.; Tsuchida, K.; Ohgushi, K.; Hayashi, N.; Kim, J.; Tsuji, N.; Takahashi, R.; Matsushita, M.; Nishiyama, N.; Inoue, T.; Irifune, T.; Kato, K.; Takata, M.; Takano, M. Giant negative

Page 10 of 12

thermal expansion in the iron perovskite SrCu3Fe4O12. Angew. Chem. Int. Ed. 2011, 50, 6579-‐6582. (24) Ishiwata, S.; Azuma, M.; Hanawa, M.; Moritomo, Y.; Ohishi, Y.; Kato, K.; Takata, M.; Nishibori, E.; Sakata, M.; Terasa-‐ ki, T.; Takano, M. Pressure/temperature/substitution-‐induced melting of A-‐site charge disproportionation in Bi1 − xLaxNiO3 (0≤x≤0.5). Phys. Rev. B 2005, 72, 045104. (25) Azuma, M.; Chen, W.; Seki, H.; Czapski, M.; Olga, S.; Oka, K.; Mizumaki, M.; Watanuki, T.; Ishimatsu, N.; Kawamura, N.; Ishiwata, S.; Tucker, M. G.; Shimakawa, Y.; Attfield, J. P. Colossal negative thermal expansion in BiNiO3 induced by intermetallic charge transfer. Nat. Commun. 2011, 2, 347. (26) Oka, K.; Sakaguchi, C.; Sinclair, A.; Ritter, C.; Mizumaki, M.; Attfield, J. P.; Azuma, M. Intermetallic charge-‐transfer tran-‐ sition in Bi1−xLaxNiO3 as the origin of the colossal negative ther-‐ mal expansion. Phys. Rev. B 2013, 88, 014112. (27) Nabetani, K.; Muramatsu, Y.; Oka, K.; Nakano, K.; Hojo, H.; Mizumaki, M.; Agui, A.; Higo, Y.; Hayashi, N.; Takano, M.; Azuma, M. Suppression of temperature hysteresis in negative thermal expansion compound BiNi1 − xFexO3 and zero-‐thermal expansion composite. Appl. Phys. Lett. 2015, 106, 061912. (28) Naka, M.; Seo, H.; Motome, Y. Theory of Valence Transi-‐ tion in BiNiO3. Phys. Rev. Lett. 2016, 116, 056402. (29) Nishikubo, T.; Sakai, Y.; Oka, K.; Mizumaki, M.; Wa-‐ tanuki, T.; Machida, A.; Maejima, N.; Ueda, S.; Mizokawa, T.; M. Azuma, M. Optimized negative thermal expansion induced by gradual intermetallic charge transfer in Bi1−xSbxNiO3. Appl. Phys. Express 2018, 11, 061102. (30) Ishiwata, S.; Azuma, M.; Takano, M.; Nishibori, E.; Takata, M.; Sakata, M.; Kato, K. High pressure synthesis, crystal struc-‐ ture and physical properties of a new Ni (II) perovskite BiNiO3. J. Mater. Chem. 2002, 12, 3733-‐3737. (31) Azuma, M.; Carlsson, S.; Rodgers, J.; Tucker, M. G.; Tsu-‐ jimoto, M.; Ishiwata, S.; Isoda, S.; Shimakawa, Y.; Takano, M.; Attfield, J. P. Pressure-‐induced intermetallic valence transition in BiNiO3. J. Am. Chem. Soc. 2007, 129, 14433-‐14436. (32) Mizumaki, M.; Ishimatsu, N.; Kawamura, N.; Azuma, M.; Shimakawa, Y.; Takano, M.; Uozumi, T. Direct observation of the pressure-‐induced charge redistribution in BiNiO3 by x-‐ray ab-‐ sorption spectroscopy. Phys. Rev. B 2009, 80, 233104. (33) Ishiwata, S.; Azuma, M.; Takano, M. Structure and Physi-‐ cal Properties of Perovskite Bi0.8Pb0.2NiO3 in Unusual Valence State A4+B2+O3. Chem. Mater. 2006, 19, 1964-‐1967. (34) Nakano, K.; Oka, K.; Watanuki, T.; Mizumaki, M.; Machida, A.; Agui, A.; Kim, H.; Komiyama, J.; Mizokawa, T.; Nishikubo, T.; Hattori, Y.; Ueda, S.; Sakai, Y.; Azuma, M. Glassy Distribution of Bi3+/Bi5+ in Bi1 – xPbxNiO3 and Negative Thermal Expansion Induced by Intermetallic Charge Transfer. Chem. Mater. 2016, 28, 6062-‐6067. (35) Inaguma, Y.; Yoshida, M.; Tsuchiya, T.; Aimi, A.; Tanaka, K.; Katsumata, T.; Mori, D. High-‐pressure synthesis of novel lithium niobate-‐type oxides. J. Phys.: Conf. Ser. 2010, 215, 012131. (36) Inaguma, Y.; Tanaka, K.; Tsuchiya, T.; Mori, D.; Katsuma-‐ ta, T.; Ohba, T.; Hiraki, K.; Takahashi, T.; Saitoh, H. Synthesis, structural transformation, thermal stability, valence state, and magnetic and electronic properties of PbNiO3 with perovskite-‐ and LiNbO3-‐type structures. J. Am. Chem. Soc. 2011, 133, 16920-‐ 16929. (37) Yasukawa, M.; Kono, T.; Ueda, K.; Yanagi, H.; Kim, S. W.; Hosono, H. Thermoelectric properties and figure of merit of perovskite-‐type Ba1-‐xLaxSnO3 with x=0.002-‐0.008. Solid State Commun. 2013, 172, 49-‐53. (38) Kakihana, M. Invited review “sol-‐gel” preparation of high temperature superconducting oxides. J. Sol-‐Gel Sci. Technol. 1996, 6, 7-‐55.


Page 11 of 12 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


(39) Kawaguchi, S.; Takemoto, M.; Osaka, K.; Nishibori, E.; Moriyoshi, C.; Kubota, Y.; Kuroiwa, Y.; Sugimoto, K. High-‐ throughput powder diffraction measurement system consisting of multiple MYTHEN detectors at beamline BL02B2 of SPring-‐8. Rev. Sci. Instrum. 2017, 88, 085111. (40) Nishibori, E.; Takata, M.; Kato, K.; Sakata, M.; Kubota, Y.; Aoyagi, S.; Kuroiwa, Y.; Yamakata, M.; Ikeda, N. The large De-‐ bye–Scherrer camera installed at SPring-‐8 BL02B2 for charge density studies. Nucl. Instrum. Methods Phys. Res. A 2001, 467-‐ 468, 1045-‐1048. (41) Izumi, F.; Momma, K. Three-‐dimensional visualization in powder diffraction. Solid State Phenom. 2007, 130, 15-‐20. (42) Watanuki, T.; Machida, A.; Ikeda, T.; Ohmura, A.; Kaneko, H.; Aoki, K.; Sato, T. J.; Tsai, A. P. Development of a single-‐ crystal X-‐ray diffraction system for hydrostatic-‐pressure and low-‐ temperature structural measurement and its application to the phase study of quasicrystals. Philos. Mag. 2007, 87, 2905-‐2911. (43) Egami, T.; Billinge, S. J. L. Underneath the Bragg Peaks: Structural Analysis of Complex Materials; Pergamon Press: Ox-‐ ford, U.K., 2003. (44) Qiu, X.; Thompson, J. W.; Billinge, S. J. L. PDFgetX2: a GUI-‐driven program to obtain the pair distribution function from X-‐ray powder diffraction data. J. Appl. Cryst. 2004, 37, 678. (45) Farrow, C. L.; Juhas, P.; Liu, J. W.; Bryndin, D.; Božin, E. S.; Bloch, J.; Proffen, T.; Billinge, S. J. L. PDFfit2 and PDFgui: computer programs for studying nanostructure in crystals. J. Phys.: Condens. Matter 2007, 19, 335219. (46) Anisimov, V. I.; Aryasetiawan, F.; Lichtenstein, A. I. First-‐ principles calculations of the electronic structure and spectra of strongly correlated systems: the LDA+U method. J. Phys.: Con-‐ dens. Matter. 1997, 9, 767-‐808. (47) Perdew, J. P.; Burke, K.; Ernzerhof, M. Generalized gradi-‐ ent approximation made simple. Phys. Rev. Lett. 1996, 77, 3865-‐ 3868. (48) Kresse, G.; Hafner, J. Ab initio molecular dynamics for liquid metals. Phys. Rev. B 1993, 47, 558–561. (49) Kresse, G.; Furthmüller, J. Efficient iterative schemes for ab initio total-‐energy calculations using a plane-‐wave basis set. Phys. Rev. B 1996, 54, 11169-‐11186. (50) Dudarev, S. L.; Botton, G. A.; Savrasov, S. Y.; Humphreys, C. J.; Sutton, A. P. Electron-‐Energy-‐Loss Spectra and the Struc-‐ tural Stability of Nickel Oxide: An LSDA+U Study. Phys. Rev. B 1998, 57, 1505-‐1509. (51) King-‐Smith, R. D.; Vanderbilt D. Theory of polarization of crystalline solids. Phys. Rev. B 1993, 47, 1651-‐1654. (52) Resta, R. Macroscopic polarization in crystalline dielec-‐ trics: the geometric phase approach. Rev. Mod. Phys. 1994, 66, 899-‐915. (53) Vanderbilt, D.; King-‐Smith, R. D. Electric polarization as a bulk quantity and its relation to surface charge. Phys. Rev. B 1993, 48, 4442-‐4455. (54) Sakai, Y.; Yang, J.; Yu, R.; Hojo, H.; Yamada, I.; Miao, P.; Lee, S.; Torii, S.; Kamiyama, T.; Ležaić, M.; Bihlmayer, G.; Mi-‐ zumaki, M.; Komiyama, J.; Mizokawa, T.; Yamamoto, H.; Nishikubo, T.; Hattori, Y.; Oka, K.; Yin, Y.; Dai, J.; Li, W.; Ueda, S.; Aimi, A.; Mori, D.; Inaguma, Y.; Hu, Z.; Uozumi, T.; Jin, C.; Long, Y.; Azuma, M. A-‐Site and B-‐Site Charge Orderings in an s-‐ d Level Controlled Perovskite Oxide PbCoO3. J. Am. Chem. Soc. 2017, 139, 4574-‐4581.

(55) Wadati, H.; Takizawa, M.; Tran, T. T.; Tanaka, K.; Mizo-‐ kawa, T.; Fujimori, A.; Chikamatsu, A.; Kumigashira, H.; Oshima, M.; Ishiwata, S.; Azuma, M. Takano, M. Valence changes associ-‐ ated with the metal-‐insulator transition in Bi1-‐xLaxNiO3. Phys. Rev. B 2005, 72, 155103. (56) Foyevtsova, K.; Khazraie, A.; Elfimov, I.; Sawatzky, G. A. Hybridization effects and bond disproportionation in the bis-‐ muth perovskites. Phys. Rev. B 2015, 91, 121114. (57) Shan, Y. J.; Mori, H.; Imoto, H.; Itoh, M. Ferroelectric Phase Transition in Perovskite Oxide CdTiO3. Ferroelectrics 2002, 270, 381-‐386. (58) Shan, Y. J.; Mori, H.; Tezuka, K.; Imoto, H.; Itoh, M. Fer-‐ roelectric Phase Transition in CdTiO3 Single Crystal. Ferroelec-‐ trics 2003, 284, 107-‐112. (59) Hao, X. F.; Stroppa, A.; Barone, P.; Filippetti, A.; Franchini, C.; Picozzi, S. Structural and ferroelectric transi-‐ tions in magnetic nickelate PbNiO3. New J. Phys. 2014, 16 015030. (60) Payne, D. J.; Egdell, R. G.; Paolicelli, G.; Offi, F.; Panac-‐ cione, G.; Lacovig, P.; Monaco, G.; Vanko, G.; Walsh, A.; Watson, G. W.; Guo, J.; Beamson, G.; Glans, P. -‐A.; Learmonth, T.; Smith, K. E. Nature of electronic states at the Fermi level of metallic β-‐ PbO2 revealed by hard x-‐ray photoemission spectroscopy. Phys. Rev. B 2007, 75, 15310. (61) Saiduzzaman, M.; Takei, T.; Yanagida, S.; Kumada, N.; Das, H.; Kyokane, H.; Wakazaki, S.; Azuma, M.; Moriyoshi, C.; Kuro-‐ iwa, Y. Hydrothermal Synthesis of Pyrochlore-‐Type Pentavalent Bismuthates Ca2Bi2O7 and Sr2Bi2O7. Inorg. Chem. 2019, 58, 1759-‐ 1763. (62) Reaney, I. M.; Colla, E. L.; Setter, N. Dielectric and Struc-‐ tural Characteristics of Ba-‐ and Sr-‐based Complex Perovskites as a Function of Tolerance Factor. Jpn. J. Appl. Phys. 1994, 33, 3984-‐ 3990. (63) Li, C.; Lu, X.; Ding, W.; Feng, L.; Gao, Y.; Guo, Z. Forma-‐ bility of ABX3 (X = F, Cl, Br, I) halide perovskites. Acta Cryst. B. 2008, 64, 702-‐707. (64) Fujita, K.; Kawamoto, T.; Yamada, I.; Hernandez, O.; Hayashi, N.; Akamatsu, H.; Lafargue-‐Dit-‐Hauret, W.; Rocque-‐ felte, X.; Fukuzumi, M.; Manuel, P.; Studer, A. J.; Knee, C. S.; Tanaka, K. LiNbO3-‐Type InFeO3: Room-‐Temperature Polar Mag-‐ net without Second-‐Order Jahn–Teller Active Ions. Chem. Mater. 2016, 28 (18), 6644-‐6655. (65) Kawamoto, T.; Fujita, K.; Yamada, I.; Matoba, T.; Kim, S. J.; Gao, P.; Pan, X.; Findlay, S. D.; Tassel, C.; Kageyama, H.; Stu-‐ der, A. J.; Hester, J.; Irifune, T.; Akamatsu, H.; Tanaka, K. Room-‐ temperature polar ferromagnet ScFeO3 transformed from a high-‐ pressure orthorhombic perovskite phase. J. Am. Chem. Soc. 2014, 136, 15291-‐15299. (66) Xiang, H. J. Origin of polar distortion in LiNbO3-‐type “ferroelectric” metals: Role of A-‐site instability and short-‐range interactions. Phys. Rev. B 2014, 90, 094108. (67) Belik, A. A.; Yu, S.; Lazoryak, B. I.; Takayama-‐Muromachi, E. BiInO3: A Polar Oxide with GdFeO3-‐Type Perovskite Structure. Chem. Mater. 2006, 18, 1964-‐1968.



Page 12 of 12


12

PolarâNonpolar Phase Transition Accompanied by Negative Thermal

Recommend Documents

PolarâNonpolar Phase Transition Accompanied by Negative Thermal