Lithium Niobate-Type Oxides as Visible Light Photovoltaic Materials

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Lithium Niobate-type Oxides as Visible Light Photovoltaic Materials Jiangnag He, Cesare Franchini, and James M. Rondinelli Chem. Mater., Just Accepted Manuscript • DOI: 10.1021/acs.chemmater.5b03356 • Publication Date (Web): 08 Dec 2015 Downloaded from http://pubs.acs.org on December 11, 2015

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Chemistry of Materials

Lithium Niobate-type Oxides as Visible Light Photovoltaic Materials Jiangang He,†,‡ Cesare Franchini,‡ and James M. Rondinelli∗,†,¶ †Department of Materials Science and Engineering, Northwestern University, Evanston, IL 60208, USA ‡University of Vienna, Faculty of Physics and Center for Computational Materials Science, Vienna, Austria ¶Materials Science Division, Argonne National Laboratory, Argonne, Illinois 60439, USA E-mail: [email protected]

Abstract Although transition metal oxides are mechanically, chemically, and thermally stable, most of the compounds possess a large band gap (≥ 3 eV) and low carrier mobilities (owing to the small d orbital band dispersions), which significantly impedes the absorption of visible solar light (≤ 2 eV) and the transfer of photo-excited carriers. By means of advanced first-principles calculations combined with crystal-chemistry and electronic structure theory, we design a new and simple polar oxide ZnPbO3 . It adopts the LiNbO3 -type structure in the polar R3c space group and is predicted to have a direct band gap of 1.45 eV with an extremely small electron effective mass (0.21 m0 ). We propose it and related compositions as promising single phase photovoltaic materials that operate in the visible light spectrum.

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Photovoltaic (PV) materials, which directly convert solar radiation into electrical energy, are one of the most important renewable energy materials platforms. The power conversion efficiency of many candidate PV compounds, however, is often too low for broad commercial use. In general, there are three main features to optimize for efficient power conversion: (i) absorption of solar light, (ii) separation of photon-excited electron and hole pair, and (iii) transportation of electron and hole carriers from bulk to surface. First, the band gap should be smaller than 2 eV to match the energy distribution of solar spectrum and large enough to generate sufficient open circle voltages, which makes 1 to 2 eV the optimal band gap range. Second, a built-in electric field can improve charge separation and is usually achieved via a p-n junction. Finally, the charge carrier diffusion length should exceed the absorption depth, and this can be supported by selecting chemistries that give large band dispersions (small electron and hole effective masses) near the Fermi level. For these reasons, small band gap III-V and IV semiconductors with extended s and p orbitals that are facilely heterostructured are found industrially; however, their low thermal and chemical stabilities under solar irradiation and in atmospheric environments lead to long-term power generation reliability concerns. 1 In contrast, all oxide PV compounds are attractive owing to their environmental stability, low-cost processing, 2 and potential to support various ferroic orders. 3 Ferroelectric crystals allow a single phase compound to exhibit a PV response—the so-called bulk photovoltaic effect (BPVE) in noncentrosymmetric materials 4,5 that is completely different from the p-n junction mechanism—with above band gap open circuit voltages and tunable photocurrent behavior owing to the electric field switchable intrinsic dipole moment and domain structure. 6 Note that the effective built-in electric field in ferroelectric oxides is usually one order magnitude larger than that in a p-n junction. 6 Nonetheless, common ferroelectric oxides are unsuitable for PV applications because they possess band gaps beyond the visible region (BaTiO3 , 3.2 eV; PbTiO3 , 3.5 eV) or have narrow bands near the Fermi level (3d states of Fe3+ in BiFeO3 ), which notably increases electron-hole recombination events. These features pose a materials chemistry


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challenge for improving the efficiency of oxide PV compounds: 7 How to achieve a direct gap semiconducting ferroelectric oxide with broad band dispersions? The electronic band gaps in ABO3 perovskite oxides, which exhibit large chemical and structure flexibility, 8 can be remarkably reduced by judicious chemical substitution; for example, replacing alkaline earth A2+ metal cations with post-transition metal cations exhibiting nd10 electronic configurations (Zn2+ , Cd2+ , and Hg2+ ) will red shift the band gap (Table 1). Furthermore, upon substitution of second-order Jahn-Teller active d0 transition metal B-site cations 9 with d10 cations (In3+ , Sn4+ , Pb4+ , Sb5+ ) allows for stronger overlap between (n + 1)s states of the A cation and the antibonding states of the B (n + 1)s–O 2p states. Together this effect widens the conduction band and shifts the valence band towards higher energy, narrowing the band gap. 10 At the same time, the A cation radius also influences the distortion of the BO6 octahedra (B-O-B angles and B-O bond length) and therefore the dpσ hybridization, which alters the gap size: 10,11 In general the gap decreases with decreasing √ distortion of the octahedra as the tolerance factor, t = (rA−O )/( 2rB−O ), where rA−O and rB−O are the bond lengths of A-O and B-O, respectively, approaches unity (Table 1), which provides a another crystal-chemistry knob to tailor the band gap. Recent research has established routes to induce ferroelectric distortions in low-tolerance factor ABO3 oxides 24 in centric R¯3 ilmenite or polar LiNbO3 -type R3c (LN in what follows) polymorphs with edge- or corner-shared BO6 octahedra, respectively. Chemistries that will give low t < 0.85 structures include small radii A-site cations (e.g., Li+ , Mg2+ , and Zn2+ ) combined with large radii B cations (e.g., Sn4+ , Pb4+ , Bi5+ ). Note that the LN phase can be obtained by hydrostatic pressure, 25 and LN-type PbNiO3 (t=0.84), ZnSnO3 (t=0.814), LiOsO3 (t=0.82) have been synthesized by high-pressure methods. 12,26,27 For the APbO3 (A=Zn, Cd, and Hg) series, LN-type CdPbO3 was obtained at high-pressure and high-temperature conditions. 17 Interestingly, although hexagonal HgPbO3 was synthesized as early as 1973, 28 an accurate crystal structure, i.e., whether it is polar (R3c) or nonpolar (R¯3c), remains ambiguous. To the best of our knowledge, there is also no reported high-pressure or



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Table 1: Evolution in the crystal and electronic structure of select perovskite oxides.a ABO3

t

SG

b

ZnSnO3 CdSnO3 HgSnO3 CaSnO3 SrSnO3 BaSnO3

0.814 0.883 0.892 0.905 0.957 1.016

R3c 12 P nma 10 R¯3c 14 P nma 15 P nma 15 P m¯3m 16

142 -47 259 -423 -781 -1114

3.02 2.48 1.18 4.12 3.50 2.43

3.72 3.12 1.51 5.20 4.47 3.23

ZnPbO3 CdPbO3 HgPbO3 CaPbO3 SrPbO3 BaPbO3

0.764 0.828 0.836 0.848 0.897 0.953

R3c R3c 17 R3c 18 P nma 19 P nma 20 Imma 22

32 -156 84 -372 -764 -1088

1.13 0.91 f M 1.73 1.43 0.59

1.45 0.78

∆Em

c

EgHSE06

d

EgGW

1.97 0.85

e

Egexp

3.9 13 3.0 10 1.6 14 4.4 10 4.1 10 3.1 13

1.78 21 g SM 23

a

For each compound the space group (SG), crystallographic tolerance factor (t), band gap (Eg , eV), and mixing energy (∆Em , meV/f.u.) are specified. b Mixing energies defined as ∆Em = E(ABO3 ) − E(AO) − E(BO2 ) were calculated at the PBEsol level. c Theoretical band gaps obtained using HSE06 in this work. d Theoretical band gaps obtained using GW 0T C−T C@HSE in this work. e Experimental band gaps. f Metallic g Semimetallic

thin film synthesis of ZnPbO3 . Here we predict ZnPbO3 as a new semiconducting polar PV compound exhibiting the main features required for high efficiency power conversion in the LN-type structure using ab initio electronic structure methods. We first perform a rigorous crystal structure search based on lattice dynamical density functional theory calculations and representation theory. We find the ground state polymorph for ZnPbO3 to be the polar R3c LN-type phase, whereas isostructural CdPbO3 and HgPbO3 exhibit metastable R3c structures consistent with highpressure experiments. 17,28 Lastly, the electronic structures and optical properties of LN-type APbO3 (A=Zn, Cd, and Hg) are obtained using the many-body GW method and solutions to the Bethe-Salpeter equation, respectively, to show that ZnPbO3 has a direct band gap of 1.45 eV and strong absorption at ∼ 750 nm and ∼ 660 nm with an extremely small electron effective mass (0.21 m0 ).


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− Figure 1: Illustration of the unstable phonon modes described by irreps (a) Γ− 5 , (b) Γ4 , (c) + + M3 and (d) R4 . Blue arrows indicated relative size and direction for atomic displacements.

By building on the described crystal-chemistry understanding, we first examine the dynamical stability of the APbO3 compounds by calculating the phonons for the fully-relaxed high-symmetry P m¯3m structure at the DFT-PBEsol level (see Supporting Information for theoretical methods). The phonon dispersion curves for ZnPbO3 reveal numerous unstable branches (Figure S1): The softest modes, transforming as irreducible representations (irreps), − + + −1 −1 −1 −1 are Γ− 4 (ν = 232i cm ), Γ5 (ν = 323i cm ), M3 (ν = 388i cm ), and R4 (ν = 392i cm ).

Each mode’s real-space displacements are sketched in Figure 1. The phonon dispersion curves for CdPbO3 and HgPbO3 are very similar to those of ZnPbO3 (see Figure S1); therefore, we only discuss the latter. The main zone-center instability Γ− 5 is a nonpolar mode that consists of anti-polar displacements of in-plane oxide ions (Figure 1a). In contrast, the Γ− 4 mode is polar; it involves the cooperative displacements of the oxide ions forming the PbO6 octahedra along a direction opposite to the Pb and Zn cations (Figure 1b). Interestingly, the largest and smallest eigen-displacements of this polar mode are derived from Zn2+ and Pb4+ , respectively, which is remarkably different from large and moderate tolerance factors oxides, such as BaTiO3 and SrPdO3 , 11,29 where the contribution to the mode from the A cation is much smaller than from the B cation. This behavior can be understood as resulting from a highly uncoordinated A cation in small tolerance factor ABO3 oxides. The largest unstable modes appearing at the M and R points in the cubic Brillouin zone describe the in-phase 5 ACS Paragon Plus Environment

Chemistry of Materials CdPbO3 HgPbO3

Mode amplitude (Å)

−

2.1

Γ4

−

+

Γ5

ZnPbO3

R4

+

M3

CdPbO3

1.8

HgPbO3

1.5 1.2 0.9 0.6 -1

∆E (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

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-2 -3 -4 c a c bm mm 3 mm mm m2 3m Pm Cm 4m2 mm2 R32 Pm C2 /mcmmm R3 C2/m C2/ P1 /m /m Im Im I P A P4 Am R 4 I P4 I4

Figure 2: Mode amplitude and total energy gain with respect to the cubic P m¯3m perovskite structure (∆E, eV/f.u.) obtained by relaxing the unstable phonon modes transforming − + + as irreps Γ− 4 , Γ5 , R4 , and M3 along different directions, giving the specified subgroup symmetries, for APbO3 compounds. (M3+ ) and out-phase (R4+ ) PbO6 octahedral rotations common for perovskites with t < 1. 11,30 We now use these main unstable phonon modes to systematically search for the ground state structure. We consider first the phases obtained by condensing only as each gives us clues to the global structure minimum in the potential energy surface spanned by these modes. Figure 2 shows that the relative energetic stability of the fully relaxed subgroup structures depends both on the nature of the mode (irreps) and also the direction of the atomic displacements relative to the crystallographic axes of the cubic parent phase; for example, along [001] versus [110], etc., which gives rise to the different low-symmetry space groups for each irrep. 31 The largest energy gain obtained by condensing a single unstable mode is for R4+ along the [111] direction, giving rhombohedral or monoclinic structures for − each compound. The other three modes, M3+ , Γ− 4 , and Γ5 , also have considerable energy gain

as well, and therefore, when the modes combine a larger energy gain is expected provided the interaction among modes is cooperative. The large mode direction-dependent energy gain, especially for the Γ− 4 mode (Figure 2), requires more careful consideration of the interaction among multiple modes than usually garnered in most polymorphic studies. All possible combinations of modes condensing with


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different directions and the resulting subgroups require structural relaxation within the given symmetry to give confidence to the identified global structure minimum. All the possible and meaningful combinations of two and three of the unstable modes were considered, which generated 36 and 79 subgroup structures, respectively, see Tables S1-S3. From this analysis we find the lowest energy structure for ZnPbO3 is the LN-type phase (R3c), which is generated by combining R4+ along [111] direction and Γ− 4 along [111] direction. It is very surprising that R3c ZnPbO3 is 129 meV/f.u. lower in energy than the ilmenite phase, whereas the ground state structures for CdPbO3 and HgPbO3 are the centrosymmetric ilmenite phase, which are 65 and 27 meV/f.u. lower in energy than the R3c phase (see Table S3), respectively. Finally, the dynamic stabilities of the R3c APbO3 (A=Zn, Cd, and Hg) phases were checked by phonon calculations. As shown in Fig. S2, these three compounds are dynamically stable in the polar R3c structure. As expected from tolerance factors considerations, the orthorhombic P nma structure common for perovskite oxides with moderate tolerance factors is disfavored by the APbO3 (A=Zn, Cd, and Hg) as well. To shed light on this observation, we briefly review the crystal chemistry of these structures (see Figure 3 for schematic illustrations). The LN-type structure is a very distorted perovskite derivative with a much smaller coordinate number (6 versus 12 for the ideal perovskite) for the A cation, 32 and can be derived from the cubic phase via two irreps (R4+ ⊕ Γ− 4 ) as noted above. In constrast, ilmenite is significantly different from perovskite in that both the A and B cations are octahedrally coordinated with oxygen and the BO6 octahedra are linked together by sharing octahedral edges rather than corners. Although both the A and B cations are octahedrally coordinated with oxygen in the ilmenite and LN-type (with more distorted AO6 octahedra) structures, the ordering of these polyhedra in sheets along the c-axis are different (Figure 3). The structure that is favored mainly depends on the relative size of A and B cations, with the ilmenite structure being adopted if the tolerance factor is smaller than 0.85. 33 However, for compounds with much smaller tolerance (larger A and B-site cation mismatch), the LN-type structure is favored since the A and B



Energy (eV/f.u.)

-51

ZnPbO3

CdPbO3

-50

-45

HgPbO3

LN-type Pb Zn

R3c (LN-type) Pnma (perovskite) R3 (ilmenite)

-52

-51

-46

-53 (a) -54

∆H (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

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104 112 120 3 Volume (Å )

(b) -52

120 128 136 3 Volume (Å )

Perovskite

(c) -47 120 128 136 144 3 Volume (Å )

0.3 0.1 0.04 0.2 0 0 0.1 (d) (e) (f) -0.1 -0.04 0 0 10 20 30 0 10 20 30 0 10 20 30 Pressure (GPa) Pressure (GPa) Pressure (GPa)

Ilmenite

Figure 3: Thermodynamic stability of APbO3 phases indicated by the evolution in total energy with unit cell volume and enthalpy changes (∆H) with hydrostatic pressure for ZnPbO3 (a) and (d), CdPbO3 (b) and (e), and HgPbO3 (c) and (f) in the main structural polymorphs shown to the right. In the LN-type structure, both the Zn and Pb cations are octahedrally coordinated; the polyehedra are linked together in the same layer. In contrast, perovskite only exhibits octahedrally coordinated Pb cations which are corner connected in three-dimensions (Zn cations fill voids formed by the PbO6 units). Ilmenite is similar to the LN-type structure with the exception that the PbO6 and ZnO6 units order on separate layers.

cations stacking pattern (ordered in a 50:50 ratio within each layer) produces smaller in-plane strain than does the alternating stacking of A and B cations (fully ordered planes). As the result of these differences, displacive phase transitions can happen between perovskite and the LN-type structure with varying pressure (see below), whereas the ilmenite to LN-type transition requires the cation order to be rearranged, which is more difficult to achieve. 33 Next, we examine the pressure dependent phase stability of the polar APbO3 compounds as an experimental guide to synthesis and materials realization in the laboratory. Highpressure methods have been widely used to stabilize LN-type small tolerance factor ABO3 compounds. 17 At high pressure, however, the LiNbO3 phase, similar to ilmenite, usually transforms to a perovskite structure, typically P nma symmetry. Owing to the large structural difference between perovskite and ilmenite phases, however, when the P nma phase is stabilized


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at high pressure, then it usually does not return to the equilibrium ilmenite structure after the pressure is released. Rather, it transforms to a metastable LN phase. 33 LN-type ZnTiO3 , ZnSnO3 , and CdPbO3 , e.g., have been synthesized with this process. 12,17,33 Figure 3 shows the total energy and enthalpy change as a function of the unit cell volume and pressure, respectively for the APbO3 compounds at 0 K. The polar LN-type structure is always the most stable phase for ZnPbO3 over the pressure range explored (0-30 GPa), likely owing to it having the smallest tolerance factor (t=0.764) and perovskite-like PbO6 rotations that permit proper Zn bond coordination even at high pressure. Therefore, we propose it should be experimentally accessible using either soft-chemistry or high pressure synthesis routes. The behavior of CdPbO3 and HgPbO3 , however, is quite different. Although the centric ilmenite phases of CdPbO3 and HgPbO3 are stable at ambient and in lowpressure regions, we find at moderate pressures, the ferroelectric R3c phase is predicted to be inaccessible for CdPbO3 . In contrast, it is anticipated to exist for HgPbO3 . Note, that although the R3c phase is not predicted to be stable at zero Kelvin, the calculated critical pressure for the R3c to P nma transformation in CdPbO3 may be estimated as ∼8.7 GPa, and this value is in close agreement with previous (experimental) studies. 17,34 Thus, we are able to capture the correct sequence of transitions, albeit the stability of the finite temperature phase likely requires the addition of the vibrational entropy contribution. Finally at elevated pressure (∼10 GPa), the centric perovskite (P nma) is the thermodynamic equilibrium structure for both compounds. This behavior is consistent with the trend observed in other small tolerance factors ABO3 compounds. 35,36 Next we calculated the electronic structures of LN-type R3c APbO3 (A=Zn and Cd) using a state-of-the-art many-body GW method 37 based on the converged HSE06 wavefunctions. To minimize the effects of the starting single-particle wavefunctions on the GW calculation, we updated the Green’s functional (G) and fixed the screened Coulomb potential (W ) until the quasi-particle energy converged. Excitonic effects are also included in the calculation of W via the T C − T C kernel. 38 This entire approach we refer to as GW0T C−T C @HSE06 (see SI).



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This method can significantly improve the agreement with experiment, as shown in Table 1 for this series of compounds. Our calculations at the GW0T C−T C @HSE06 level show that both ZnPbO3 and CdPbO3 are direct band gap semiconductors with calculated band gaps of 1.45 eV and 0.78 eV (Table 1), respectively, whereas HgPbO3 is a metal with bands at the Fermi level with nearly linear E(k) dispersion relationships (Figure S4). Although HgPbO3 is isostructural to the recently synthesized 27 polar metal LiOsO3 , the origin of the metallic state in HgPbO3 is different. 39 LiOsO3 is metallic owing to fractional band filling of the Os 5d manifold with a nominal t32g electronic configuration, whereas the electron counts for Hg2+ and Pb4+ in HgPbO3 would indicate insulating behavior. However, the electron-lattice interactions are more important in HgPbO3 , specifically those supported by cooperative displacements of the oxide ligands surrounding Pb, i.e., displacement modes R4+ and Γ− 4: We find a clear band gap enlargement with increasing mode distortion amplitude (Q) found ˚ in the equilibrium LN-type structure for HgPbO3 , Q(R4+ ) = 1.26 ˚ A and Q(Γ− 4 ) = 0.38 A, ˚ to ZnPbO3 , Q(R4+ ) = 1.65 ˚ A and Q(Γ− 4 ) = 0.83 A. An increase in these mode amplitudes reduces the hybridization between the B cations and oxide anions, consistent with known effects of octahedral rotations in orthorhombic perovskite oxides. 10,11 Figure 4a shows the quasiparticle band structure for ZnPbO3 . Similar to most oxides, the top of valence band is dominated by O 2p orbitals (Figure 4b). Furthermore, the shallow and fully occupied 3d orbitals in ABO3 compounds with nd10 A cation electronic configurations (Zn2+ , Cd2+ , and Hg2+ ) strongly hybridize with these oxygen states; this behavior pushes up the valence band maximum and reduces the band gap. 40 It also explains why nd10 cation containing stannates have smaller band gaps than the A nd0 cation compounds with Ca2+ , Sr2+ , and Ba2+ presented in Table 1. Another remarkable character of these band structures is the large band dispersions, derived mainly from the Pb 6s orbitals, at the bottom of conduction band located at the Γ point, indicating light electron band masses and potentially high electron mobility. Indeed, we estimate the electron effective mass for ZnPbO3 to be m∗ = 0.21m0 within the parabolic approximation, which is lighter than that reported for


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Figure 4: Quasiparticle (a) band structure and (b) density of states for LN-type R3c ZnPbO3 calculated at the GW0T C−T C @HSE06 level. (c) Optical absorption coefficient for LN-type ZnPbO3 and CdPbO3 calculated from the BSE dielectric function on the top of GW 0T C−T C @HSE06 band structure; the experimental absorption coefficients of In0.27 Ga0.73 N (InGaN, Ref. 42 ) and CH3 NH3 PbI3 (MAPbI3 , Ref. 43 ) are reproduced along side the solar radiation energy distribution (gray shading) for comparison. CH3 NH3 PbI3 . 41 As a result, the electron carriers in ZnPbO3 should have a higher mobility and longer diffusion length, which would significantly enhance the electron-hole pair separation and carrier transport from the bulk of the sample to surface. Next, the frequency-dependent dielectric functions for ZnPbO3 and CdPbO3 were calculated by solving the Bethe-Salpeter equation (BSE) based on the GW0T C−T C @HSE06 quasiparticle band structure. Figure 4c shows the wavelength dependent absorption coefficients for ZnPbO3 and CdPbO3 compared to the experimental absorption spectra of In0.27 Ga0.73 N (InGaN) and CH3 NH3 PbI3 (MAPbI3 ). 42,43 The first strong absorption peak for ZnPbO3 is at ∼750 nm, which is longer than that of the widely used PV material InGaN, and comparable to the most intensely studied organometallic perovskite PV material MAPbI3 . Moreover, ZnPbO3 shows much better absorption than InGaN and MAPbI3 over most of the visible portion of the electromagnetic spectrum. Although CdPbO3 has stronger absorption and a lower absorption edge than ZnPbO3 its small band gap (0.78 eV) impedes its application as a PV material. Having established ZnPbO3 has electron carriers with low effective masses and an optimal 11 ACS Paragon Plus Environment


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band gap, we now examine the electric polarization in the noncentrosymmetric material, which is known in related oxides 6 to produce an above band gap open-circuit voltage owing to the BPVE and higher PV efficiency. 44 The HSE06 calculated polarization using the Born effective charges (see Table S4), which includes the electronic contribution to the electric dipole per unit cell, is 94 µC cm−2 , whereas the polarization obtained by using the nominal ionic charges is only 80 µC cm−2 . This result indicates that although the principal contribution to the polarization is from ionic displacements, there is a non-negligible electronic contribution owing to dynamical charge transfer. Remarkably the calculated polarization is larger than that of BiFeO3 , 45 which also exhibits the BPVE and the corresponding large built-in electric field that significantly enhances the electron-hole separation and PV efficiency. In summary, we used electronic structure calculations combined with a representationtheory based polymorph search to predict that the lithium niobate-like nd10 compound ZnPbO3 should be a promising photovoltaic material with light electron masses. The manybody GW method and solutions to the Bethe-Salpeter equation allow us to conclude that it should be a direct band gap semiconductor with a 1.45 eV gap exhibiting strong optical absorption at ∼ 750 nm and ∼ 660 nm. The polar structure for ZnPbO3 makes it useful as a PV material without the need for heterostructuing owing to the BPVE and its spontaneous polarization of ∼ 94 µC cm−2 . Similar studies on CdPbO3 and HgPbO3 reveal that the metastable polar phase (R3c), which should be accessible by high-pressure synthesis, exhibit semiconducting (0.78 eV band gap) and metallic behavior, respectively. The former can be used as photoelectric material working in the infrared, whereas the latter is a realization of the elusive class of ‘ferroelectric metals’ proposed more than 50 years ago by Anderson. 46,47

Acknowledgement Work at the University of Vienna was sponsored by the FWF-SFB project VICOM (Grant No. F41). J.M.R. was supported by the U.S. DOE, Office of Basic Energy Sciences, under


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grant no. DE-AC02-06CH11357. All calculations were performed on the Vienna Scientific Cluster (VSC).

Supporting Information Available Supplementary information is available in the online version of the paper: Description of the theoretical methods, calculation of the absorption coefficient, complete structure search results, and additional tables and figures. This material is available free of charge via the Internet at http://pubs.acs.org. This material is available free of charge via the Internet at http://pubs.acs.org/.

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Lithium Niobate-Type Oxides as Visible Light Photovoltaic Materials

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