Designing a Lower Bandgap Bulk Ferroelectric Material with a Sizable

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Designing a Lower Bandgap Bulk Ferroelectric Material with a Sizable Polarization at the Room Temperature Shyamashis Das, Somnath Ghara, Priya Mahadevan, Sundaresan Athinarayanan, Jagannatha Gopalakrishnan, and D. D. Sarma ACS Energy Lett., Just Accepted Manuscript • DOI: 10.1021/acsenergylett.8b00492 • Publication Date (Web): 20 Apr 2018 Downloaded from http://pubs.acs.org on April 22, 2018

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Designing a Lower Bandgap Bulk Ferroelectric Material with a Sizable Polarization at the Room Temperature Shyamashis Das,† Somnath Ghara,# Priya Mahadevan, ⊥ A. Sundaresan, # J. Gopalakrishnan, † and D. D. Sarma†,* †

Solid State and Structural Chemistry Unit, Indian Institute of Science, Bengaluru 560012

#

Chemistry and Physics of Materials Unit, Jawaharlal Nehru Centre for Advanced Scientific

Research, Jakkur, Bengaluru 560064 ⊥

Department of Condensed Matter Physics and Material Sciences, S.N. Bose National Centre

for Basic Sciences, Block JD, Sector-III, Saltlake, Kolkata 700106

AUTHOR INFORMATION Corresponding Author * D. D. Sarma (E-mail : [email protected])

1 ACS Paragon Plus Environment

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ABSTRACT: A low band gap ferroelectric material with a sizable polarization at ambient conditions would constitute an ideal photovoltaic material to harvest solar energy owing to their efficient polarization driven charge carrier separation. We achieve this elusive goal by co-doping a Jahn-Teller Mn3+ and Nb5+ pair for two Ti4+ ions in ferroelectric BaTiO3. Representing a charge-neutral dipole doping, this approach achieves for the first time a bulk ferroelectric oxide with the lowest bandgap of 1.66 eV with a sizable polarization of nearly 70% of BaTiO3. We contrast with the analogous system with Mn3+ replaced by the non-Jahn-Teller Fe3+ (3d5) ion, that, even at a much lower level of doping reduces the polarization to 25% without reducing the bandgap significantly, establishing the efficacy of the present strategy. TOC GRAPHICS


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If any ferroelectric material would have a small bandgap, thereby effectively absorbing a large part of solar spectrum, it would present itself as a promising material candidate for photovoltaic (PV) applications, owing to their efficient polarization driven charge carrier separation.1-7 However, one of the major deterrent in adopting the large number of known ferroelectric materials for this purpose is the large band gap associated with these. This leads to harvesting only a small part of the solar spectrum. In general, common ferroelectric oxides with the perovskite structure and a generic chemical formula, ABO3, exhibit wide band gaps. This arises from the large difference in electronegativities of oxygen and the transition metal atom at the B site. At the same time, in most perovskite ferroelectric materials, the transition metal ion at the B site also plays the crucial role to drive ferroelectricity.8 Any attempt to reduce the band gap by substitution of the B ion leads to a deterioration of ferroelectric and dielectric properties.9 Many of the well-known wide band gap ferroelectric oxides like KNbO3,10-12 PbTiO3,13-14 Pb(Zr,Ti)O3,15 BiFeO3,16-19 Bi4Ti3O1220, Bi5Ti3FeO1521 have been investigated both of theoretically and experimentally. Unfortunately, however, almost all such substituted systems end up with either a large bandgap or a small or zero polarization (or both). On the other hand, many of the low band gap ferroelectric materials22-28 fail to show siginificant polarization suitable for a PV device. This has prompted efforts to an alternate route of making thin-film heterostructures29 or multilayers30-32 using sophisticated growth technologies to achieve desired properties. This approach has led to at least two examples so far in the literature29-30 with sizable polarization and a reduced bandgap. The standard chemical route of substitution however has not yielded any comparable result. In the present study, we report a new class of potential materials for photovoltaic applications. These materials are based on band-gap tuning of classical ferroelectric material BaTiO3 (BTO)



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via chemical substitutions. Pure BTO has been extensively studied for its PV related properties.33-36 The single-most significant drawback of this material is that like many other perovskite ferroelectric materials, BTO also possesses a wide band gap (3.2 eV)37. This large band gap is a consequence of the 3d0 configuration of Ti4+ ions. On the other hand, the ferroelectricity in BTO has been related to this “d0”-ness.38 Thus, it is not surprising to find that an effort to reduce the band gap by substitution of Ti4+ with non-d0 transition metal ions usually destroys the ferroelectricity. For example, isovalent substitutions of Ti4+ with other tetravalent metal ions (eg. Mn4+) is known to stabilize the high temperature hexagonal non-ferroelectric phase of BTO at ambient conditions.39 So, to synthesize BTO based reduced band gap ferroelectric materials we exploit the strategy to substitute two Ti4+ ions with two different transition metal ions, an acceptor (TM1) and a donor (TM2), balancing the dopant level to achieve a net zero charge. This pair however has a finite dipole associated with it and the general composition is given by BaTi1-x(TM11/2TM21/2)xO3. The introduction of the dipole is intended to help retain the polarization. We choose TM1 = Mn3+ as non-d0 transition metal ions (3d4) to reduce the band gap and TM2 = Nb5+ which is known to support ferroelectric states40 with its 4d0 configuration. Our strategy involves choosing the Mn3+ 3d4 Jahn-Teller ions for TM1. This choice, in addition to leading to an effective reduction in the bandgap, also has lattice distortions which lift the degeneracy of the 3d orbitals and should result in an insulating state. It is well-known that the stability of a ferroelectric ground state is significantly contributed by changes in the lattice energy arising from local atomic displacements and not only by dipole interactions. Therefore, the Mn3+ substitution, with its intrinsic property for local atomic distortions, may help in retaining the ferroelectric ground state better than using a non-JahnTeller ion for TM1. This anticipation is supported by investigating the analogous series of


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substituted compounds with Mn3+ 3d4 ions replaced by non-Jahn-Teller Fe3+ 3d5 ions. We provide additional support to our experimental findings by carrying out extensive theoretical calculation. We refer to the two separate series of co-doped BTO compositions, namely Mn-Nb co-doped and Fe-Nb co-doped as BTMNO and BTFNO, respectively. In pure BTO, cooperative displacements of Ti-atoms along the crystallographic c-direction below the transition temperature (393 K) lead to the tetragonal distortion and, consequently, give rise to the ferroelectricity. Substitution of Ti with other transition metal ions has been reported to decrease this tetragonal distortion.39 However, the extent of this decrease depends on the specific substituent transition metal ion. Here we compare the effect of Fe-Nb co-doping and Mn-Nb codoping in BTO on its tetragonal distortion. The tetragonal distortion is reflected in the powder Xray diffraction pattern (PXRD) of pure BTO as each peak is split in a tetragonal environment except the (111) peak. Figure 1a shows PXRD patterns of these two series of materials BTMNO (in red) and BTFNO (in blue) along with pure BTO (in black). We have expanded the plot around the (200) reflection for better clarity in Figure 1b to follow the tetragonal distortion, as reflected by the splitting of this peak. It is evident from this figure that the splitting of the (200) peak due to the tetragonal distortion decreases upon Ti substitution with both Mn-Nb and Fe-Nb pairs. In addition, Figure 1b also establishes that the splitting of the (200) peak is much smaller for the Fe-Nb substituted composition compared to the Mn-Nb one. This is evident from the results of x = 0.05 and 0.075 compositions of both series shown in Figure 1b. This implies that the tetragonal distortion, associated with ferroelectric property of BTO, decreases much faster for the non-Jahn-Teller Fe3+ substitution compared to the Jahn-Teller Mn3+ substitution. We have extracted lattice constants of all samples of BTMNO series from Rietveld refinement of


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corresponding X-ray diffraction patterns and parameters extracted from refinement are listed in Table S1 of Supporting Information (SI). A representative diffraction pattern for BTMNO,

b)

BaTi1-x(TM1/2Nb1/2)xO3 (200)

(200)

c) 3x10 Intensity (counts)

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c 4.02

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44.8 45.2 45.6

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a 0.00

0.04

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0.16

Composition (x)

Figure 1. Evolution of structural change with doping in BTO. (a) Powder X-ray diffraction pattern of samples of BaTi1-x(Mn1/2Nb1/2)xO3 series (in red lines) with 0.025 ≤ x ≤ 0.1 and of BaTi1-x(Fe1/2Nb1/2)xO3 series (in blue line) with 0.02 ≤ x ≤ 0.075 along with pure BTO (in black line). (b) Magnified (020) family of reflection indicating decrease of tetragonal distortion with Ti-substitution. (c) Diffraction pattern from Rietveld refinement of the X-ray ray diffraction data of the sample of composition BaTi1-x(Mn1/2Nb1/2)xO3, x = 0.025. (d) Evolution of lattice parameters with increasing concentration of Mn-Nb dopant. x = 0.025 composition obtained from Rietveld refinement is shown in Figure 1c in comparison with experimental data. The values of lattice constants of BTMNO series extracted from the refinement are plotted as a function of the doping concentration which is shown in Figure 1d. We find that the room temperature lattice constants ‘c’ and ‘a’ are indistinguishable beyond 7.5% Mn-Nb co-doping which indicates cubic symmetry and a paraelectric state for higher doped compositions of this material.


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A systematic decrease of the tetragonal distortion with an increasing doping concentration is suggestive of a systematic reduction in the ferroelectric properties in both series of compounds. It is known41 that with decreasing temperature, BTO shows three distinct phase transitions corresponding to (i) cubic paraelectric to tetragonal ferroelectric phase at 393 K, (ii) tetragonal ferroelectric to orthorhombic ferroelectric phase at 285 K and (iii) orthorhombic to rhombohedral at 196 K. These transitions are reflected in the temperature dependent dielectric constant of BTO with a prominent sharp peak in the dielectric constant at a paraelectric to ferroelectric transition at 393 K.42-43 Therefore, we have measured the temperature dependent dielectric constant to determine the paraelectric to ferroelectric transition temperature (TC) in both series. Results of dielectric constant as a function of temperature in the applied electric field frequency range of 1 kHz-1MHz are shown in Figure 2 for the samples of BTMNO and BTFNO series. In the BTMNO series the first three compositions x = 0.025, 0.05 and 0.075 (Figure 2a-c), clearly reflect the cubic paraelectric to tetragonal ferroelectric and tetragonal ferroelelctric to

2000

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BTMNO x = 0.05

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BTMNO x = 0.025 1 kHz 10 kHz 4000 100 kHz 1 MHz


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100

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Figure 2. Characterization of ferroelectric to paraelectric transition temperature. Dielectric constant as a function of temperature a)-d) of four different samples of BaTi1-x(Mn1/2Nb1/2)xO3 7 ACS Paragon Plus Environment


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series with 0.025 ≤ x ≤ 0.1 and e)-g) of three different samples of BaTi1-x(Fe1/2Nb1/2)xO3 series with 0.02 ≤ x ≤ 0.075. For each composition dielectric constant is shown for ac field frequency range 1 kHz-1 MHz. Red arrows in the plots indicate different phase transitions. Ferroelectric to paraelectric phase transition temperature is shown as a function of composition is shown in the plot h). Green horizontal line in h), corresponds to room temperature. orthorhombic ferroelectric phase transitions in temperature dependent dielectric constants, as indicated by red arrows in the respective plots. For these three compositions, phase transition temperatures are also independent of the applied electric field frequency; this indicates the absence of any glassy behavior, indicating a long range order state. For the higher doped composition, x = 0.1, only the cubic paraelectric to tetragonal ferroelectric phase transition is distinctly visible below the room temperature (Figure 2d); this is consistent with the observation of the cubic phase in XRD of this composition collected at the room temperature (Figure 1). The peak in the dielectric constant as a function of the temperature is also broad for this sample; this is indicative of presence of some disordered dipoles within the ferroelectric phase of this material. The evolution of dielectric constant with temperature for the samples of BTFNO series are shown in Figure 2e-g. Here we find that, except for the lowest doped composition (x = 0.02, Figure 2e) the peak in the dielectric constant as a function of temperature is broad for the two other compositions, indicating presence of highly disordered dipoles in these samples. A comparison of the temperature dependent dielectric responses of these two series of samples brings out the fact that each of the Mn-Nb co-doped compositions shows a sharper ferroelectric transition with a higher transition temperature than the corresponding Fe-Nb co-doped composition. This is in accordance with powder X-ray diffraction data which also indicate reduced tetragonal distortion in the samples of the BTFNO series compared to those of the BTMNO series at room temperature. In Figure 2h, we have plotted the transition temperature


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(TC) between the paraelectric and the ferroelectric phases as a function of the dopant concentration for the samples of both BTMNO and BTFNO series. It is clear from this figure that the transition temperature decreases with an increasing dopant concentration and falls below the room temperature (marked by green horizontal line in Figure 2h) for the higher doped compositions. A comparison of TC of the two series of samples indicate that transition temperature falls more rapidly for the samples of BTFNO series than BTMNO series. A closer inspection of this data suggests that the three compositions of BTMNO series with lower doping concentrations (x = 0.025, 0.05 and 0.075) are suitable for photovoltaic applications, from the point of view of the internal ferroelectric polarization helping to separate the photoexcited electron-hole pairs. In order to compare the extent of the ferroelectric polarization for samples in these two series, we have measured electric polarization vs applied electric field (P-E) hysterisis loops. P-E loops measured at liquid nitrogen temperature for x = 0.075 compositions of BTMNO and BTFNO are shown in Figure 3a. Frequency dependent P-E loops of these two samples are shown in Figure S1 of SI. We notice for BTMNO sample, the P-E loop (Figure S1a of SI) is relatively invariant with the change in the frequency of the electric field, whereas for BTFNO (Figure S1b of SI), the shape of the P-E loop has relatively more dependence on the frequency of the electric field. We also find that the magnitude of the saturation polarization (15.1 µC/cm2) is much higher for the x = 0.075 sample of BTMNO series compared to that (9.5 µC/cm2) in the corresponding composition of BTFNO series (Figure 3a). It is to be noted that saturation polarizations, Psat, for both these samples are less than that (24.1 µC/cm2) of pure BTO. Thus, these data together with the structural (Figure 1b) and dielectric (Figure 2) data suggest a decreasing stability of the ferroelectric state with an increasing level of dopants; this suppression of the ferroelectric state 9 ACS Paragon Plus Environment

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is clearly more pronounced in the BTFNO series with larger reduction in TC and Psat compared to those for the BTMNO sample. This reduced adverse impact on the ferroelectric state due to Mn3+

5 0 -5 -10

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5 0 -5

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Figure 3. Comparison of extent of ferroelectricity in different compositions of doped BTO samples. a) Polarization vs electric field hysteresis loops of compositions BaTi1-x(TM1/2Nb1/2)xO3 with TM = Mn, x = 0.075 (in black) and with TM = Fe, x = 0.075 (in red) at applied electric field frequency (f) = 100 Hz and at temperature (T) = 77 K. b) Polarization vs electric field hysteresis loops for two compositions of BaTi1-x(Mn1/2Nb1/2)xO3 series (with x = 0.05 and 0.075) measured at room temperature. c) Invertible spontaneous polarization as a function of temperature as obtained from the measurement of the composition BaTi1-x(Mn1/2Nb1/2)xO3 with x = 0.075. 3d4 doping in BTMNO is critical to achieve a reduced bandgap ferroelectric material, as shown in this paper. Figure 3b shows P-E loops measured at the room temperature for x = 0.05 and 0.075 compositions of BTMNO series. Both these materials show a reasonably large saturation polarization of about 15 µC/cm2; this value favorably compares with the saturation polarization value of the other ferroelectric materials employed for PV applications.20 A little higher value of saturation polarization of x = 0.075 composition in comparison to x = 0.05 composition is within the experiemntal uncertainties and this may be contributed by difference in sample desnity or stray field of the surrounding medium. P-E loop measurements of samples with composition x ≥ 0.10 did not yield any ferroelectric loop as TC of these samples are below the room temparature. Samples of BTFNO series, on the other hand, shows very small saturation polarizations at room temperature with a narrow P-E loop and hence suitable for photovoltaic application. 10 ACS Paragon Plus Environment

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Measurement of pyroelectric current provides a firm evidence for any ferroelectric phase transition. Spontaneous polarization (Ps) is derived by integrating pyroelectric current over the measurement duration. At paraelectric to ferroelectric phase transition temperature (TC), Ps appears without apparent discontinuity and increases steeply below this temperature. As dielectric measurement and ferroelectric hysteresis loop indicate presence of inherent polarization in Mn-Nb co-doped BTO samples, we accordingly measured pyroelectric current (Ip) as a function of temperature for BTMNO sample with x = 0.075, as shown in Figure 3c, in terms of Ps calculated from Ip. We find Ps starts changing continuously across the two phase transitions, one at T∼350 K, corresponding to the cubic paraelectric to the tetragonal ferroelectric phase transition and another at T∼270 K corresponding to the tetragonal to the orthorhombic ferroelectric phase transition. It is important to note here that these two phase transitions have also been observed from the measurement of dielectric constant as a function of temperature for this particular composition (Figure 2c). The inversion of polarization upon applying an opposite poling field of the same magnitude further confirms the ferroelectricity in this sample. In order to understand the microscopic origin of the enhanced ability of Mn-Nb substitutions in retaining the ferroelectric state, we have analyzed the local geometry around the dopant sites and compared ferroelectric stabilization energy of the samples of both series. The various M-O bond lengths in three crystallographic directions as obtained from density functional theory (DFT) calculations are listed in Table S2 and S3 of SI, for x ≈ 0.11 compositions of BTMNO and BTFNO samples, respectively. Clearly, the most prominent structural difference between two series is that there is a marked Jahn-Teller distortion at the MnO6 octahedra, which is entirely absent in the FeO6 octahedra as seen from respective M-O bond lengths of ground state 11 ACS Paragon Plus Environment


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configuration (configuration III, Figure S2 of SI). This is obviously a consequence of Mn3+ dopant used in one case being a Jahn-Teller ion, whereas Fe3+ with 3d5 configuration used in the other series is not. Interestingly, the distortion around the Mn site forces a compensating distortion at the co-dopant, Nb site, as is most easily seen for the ground state configuration (configuration III, Figure S2 of SI) in Table S2 of SI. Thus, a strong electron-lattice coupling is ensured by the substitution of Mn-Nb pair that has the ability to couple coherently to the lattice distortions associated with the off-centering of Ti atoms in the tetragonal structure. On the other hand, insertion of undistorted octahedral pairs around Fe and Nb sites effectively decouple lattice distortions around Ti sites in the tetragonal structure by inserting undistorted octahedra in between, thereby reducing ferroelectric stability and the TC. Examining the other dopant configurations which are at slightly higher energies, we find that the presence of the Mn-Nb pair results in a significant polarization being induced in the Ti sublattice as well as an off-centering of the Nb5+ ion. In contrast, the induced polarizations as evidenced by the Ti-O bond lengths are smaller in the case of Fe-Nb co-doping, supporting our observations. The ferroelectric stabilization energies, defined as energy stability of the ferroelectric configuration compared to the corresponding paraelectric state and shown in Table S4 of SI, predict the experimentally observed trends very well. We find that for both x = 0.06 and 0.11 compositions of BTMNO series possess higher ferroelectric stabilization energy than that of BTFNO series and thereby helping us to understand microscopic origin of our experimental findings. We determined optical band gaps of these samples using UV-visible optical reflectance spectroscopy. In Figure 4a, Kubelka-Munk function44 calculated from optical reflectance spectra of BTMNO samples of compositions x = 0-0.15 and one composition of BTFNO (x = 0.05) is 12 ACS Paragon Plus Environment

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plotted against energy. According to Kubelka-Munk model, the band gap is determined from the point of intersection between tangents drawn from the inflection point of the curves shown in Figure 4a and horizontal energy axis. The band gap obtained in this method is listed out in Table 1 for all the compositions. The band gap decreases continuously from 3.18 eV for pure BTO which is white in color, to 1.3 eV for the composition of BTMNO, x = 0.15 making the material dark yellow in color. We have also shown Kubelka-Munk function of composition BTFNO, x = 0.05 to compare its band gap with equivalent Mn composition BTMNO, x = 0.05. We find the bandgap (2.64 eV) for BTFNO composition is much higher than that (2.17 eV) of the corresponding BTMNO composition, thus, making BTMNO series a better photovoltaic material than corresponding compositions of BTFNO both in terms of superior ferroelectric properties (higher ordering temperature and higher polarization) and lower optical bandgaps (see Table 1). 10

Spectral Irradiance (Wm-2 nm-1)

1.6

1.2

0.8

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3 2 1

3.5

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4 3 2 1

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Figure 4. Change in optical band gap with increase in concentration of dopants. a) KubelkaMunk function plotted as a function of incident energy of BTMNO, x = 0-0.15 and BTFNO, x = 0.05. Band gap of a few standard materials which were used in PV device has also been shown. Total and orbital projected electronic density of state of b) BTO and c) BTMNO, x = 0.11 composition. Top of the oxygen 2p state is set as zero in energy scale in both the plots. Table 1. Comparison of relevant parameters for PV device of present materials.

BTO

Room Temperature Saturation Polarization (µC/cm2) 24.1

Curie Temperature Tc (K) 393

Band Gap Eg (eV) 3.18

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314.2

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A comparison of optical band gap values of present material with those reported in the literature so far, establish that the optimal composition of x = 0.075, BTMNO has the best combination of a sufficiently low bandgap (1.66 eV) and substantially high polarization (70% of the pure BTO) at the room temperature among all bulk ferroelectric materials; it also compares favorably against the best results29-31 among all multilayers and heterostructure devices reported so far. To understand the microscopic origin of the band gap lowering due to substitution of Ti4+ ions by transition metal ions with finite number of d-electrons, we examined the electronic structure obtained from DFT calculations. We note that there are many symmetry inequivalent realizations of dopant distributions, as already illustrated in Figure S2 of SI; we also note that most of these structures are almost degenerate in total energies (see Table S2 of SI) indicating that in real


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samples, we are likely to achieve random combinations of different dopant distributions without any significant preference of one structure over another. Thus, we have performed calculations for the electronic structures of each of these nearly degenerate configurations and averaged over, to arrive at the representation of the overall electronic structure of the experimentally obtained sample. The results are shown in terms of electronic density of states for pure BTO and BTMNO, x ≈ 0.11 in Figure 4b and 4c, respectively. We find that, in presence of dopants, midgap states arise with predominant Mn 3d character. Presence of these Mn 3d states both at the top of the valence band and at the bottom of the conduction band decreases the band gap of this doped material significantly. We find that the band gap decreases from 3.1 eV in pure BTO to 1.8 eV in the ~11% Mn-Nb substituted composition. These calculated band gap values are in good agreement with experimentally obtained values of equivalent compositions (see Table 1). The agreement between calculated and experimental band gap values is a consequence of the use of HSE06 hybrid functional which is known for its success in describing such properties of oxide materials, establishing the presence of mid gap states as the main cause of the rapid decrease in the bandgap with the composition, x. It is to be noted that Fe-Nb substituted samples do not show mid-gap states like Mn-Nb substituted samples. The Jahn-Teller distortion of Mn3+ ions controls the splitting of the mid-gap states that lead to the reduction of the band gap (Figure 4c). The splitting between the highest occupied states and the lowest unoccupied states on Fe is governed by the exchange splitting. This is the reason why we do not find a substantial reduction of the band gap for a non-Jahn-Teller ion like Fe3+. The present work has shown a new direction to design compositions to tune relevant properties of the classical ferroelectric material BTO, making it suitable for PV applications. Specifically, we show a way to reduce its large band gap drastically without significantly compromising its 15 ACS Paragon Plus Environment


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ferroelectric polarization properties via a doping strategy involving a compensating co-doping with Nb5+ 4d0 ions and Mn3+ 3d4 Jahn-Teller ions in place of Ti4+ 3d0 ions in BTO. This ensures (1) maximizing the presence of d0 ions that are known to promote ferroelectricity, (2) a supporting spontaneous lattice distortion of the Jahn-Teller non-d0 ion that is not as deleterious to the ferroelectric state as non-Jahn-Teller ions; and (3) a major reduction of the bandgap due to the presence of non-d0 Mn3+ ions. We also show marked reductions of all desirable properties when Mn3+ ions are replaced by non-Jahn-Teller Fe3+ 3d5 ions, establishing the essential role played by Mn3+ ions. Past attempts in the literarure to decrease the band gap of ferroelectric compounds by substitution have resulted in a rapid loss of polarization even before a substantial bandgap tuning could be achieved. This makes these Mn-Nb co-doped samples particularly attractive with a substantial reduction in the bandgap without comprosing the polarization significantly. At the optimal composition of BaTi0.925(Mn1/2Nb1/2)0.075O3, we achieve a bulk ferroelectric material with the smallest bandgap (1.66 eV) reported so far, while retaining a sizeabale polarization. It makes this bulk compound comparable to the best materials achieved via artificial lattice structuring such as, multilayers and heterostructured thin films. ASSOCIATED CONTENT Supporting Information. Experimental section with details of sample synthesis and measurement methods, details of Rietveld refinement and list of parameters extracted from refinement of PXRD patterns, computational methods and results of ab initio calculation on geometry optimized structure two different compositions, additional polarization vs electric field hysteresis loops (PDF) AUTHOR INFORMATION


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ACS Energy Letters

D. D. Sarma (E-mail : [email protected]) Notes The authors declare no competing financial interest. ACKNOWLEDGMENT We thank IISc, JNCASR and the Department of Science and Technology, Government of India for financial support. SD thanks CSIR for a research Fellowship. A.S. acknowledges Sheikh Saqr Laboratory at Jawaharlal Nehru Centre for Advanced Scientific Research for providing experimental facilities. J.G. thanks the National Academy of Sciences Allahabad, India (NASI) for the award of a Senior Scientist Fellowship. D.D.S. thanks Jamsetji Tata Trust for support. We also thank Prof. Olof Karis for helpful discussion. REFERENCES (1) Fridkin, V. M.; Grekov, A. A.; Rodin, A. I.; Savchenko, E. A.; Volk, T. R. Photoconductivity in Ferroelectrics. Ferroelectrics 1973, 6, 71-82. (2) Butler, K. T.; Frost, J. M.; Walsh, A. Ferroelectric Materials for Solar Energy Conversion: Photoferroics Revisited. Energy Environ. Sci. 2015, 8, 838-848. (3) Rühle, S.; Anderson, A. Y.; Barad, H.-N.; Kupfer, B.; Bouhadana, Y.; Rosh-Hodesh, E.; Zaban, A. All-Oxide Photovoltaics. J. Phys. Chem. Lett. 2012, 3, 3755-3764. (4) Kreisel, J.; Alexe, M.; Thomas, P. A. A Photoferroelectric Material is More Than the Sum of its Parts. Nat. Mater. 2012, 11, 260. (5) Paillard, C.; Bai, X.; Infante, I. C.; Guennou, M.; Geneste, G.; Alexe, M.; Kreisel, J.; Dkhil, B. Photovoltaics with Ferroelectrics: Current Status and Beyond. Adv. Mater. 2016, 28, 51535168. (6) Martin, L. W.; Rappe, A. M. Thin-film Ferroelectric Materials and Their Applications. Nat. Rev. Mater. 2016, 2, 16087. (7) Huang, H. Ferroelectric Photovoltaics. Nat. Photonics 2010, 4, 134. (8) Cohen, R. E. Origin of Ferroelectricity in Perovskite Oxides. Nature 1992, 358, 136. (9) Benedek, N. A.; Fennie, C. J. Why Are There So Few Perovskite Ferroelectrics? J. Phys. Chem. C 2013, 117, 13339-13349. (10) Grinberg, I.; West, D. V.; Torres, M.; Gou, G.; Stein, D. M.; Wu, L.; Chen, G.; Gallo, E. M.; Akbashev, A. R.; Davies, P. K.; Spanier, J. E.; Rappe, A. M. Perovskite Oxides for Visiblelight-absorbing Ferroelectric and Photovoltaic Materials. Nature 2013, 503, 509-512.



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(11) Yang, Y.; Jung, J. H.; Yun, B. K.; Zhang, F.; Pradel, K. C.; Guo, W.; Wang, Z. L. Flexible Pyroelectric Nanogenerators using a Composite Structure of Lead-Free KNbO3 Nanowires. Adv. Mater. 2012, 24, 5357-5362. (12) Sun, Y.; Guo, F.; Chen, J.; Zhao, S. Improved Ferroelectric and Photovoltaic Properties of BiMnO3 Modified Lead-free K0.5Na0.5NbO3 Solid-solution Films. Appl. Phys. Lett. 2017, 111, 253901. (13) Bennett, J. W.; Grinberg, I.; Rappe, A. M. New Highly Polar Semiconductor Ferroelectrics through d8 Cation-O Vacancy Substitution into PbTiO3: A Theoretical Study. J. Am. Chem. Soc. 2008, 130, 17409-17412. (14) Gou, G. Y.; Bennett, J. W.; Takenaka, H.; Rappe, A. M. Post Density Functional Theoretical Studies of Highly Polar Semiconductive PbTi1-xNixO3-x Solid Solutions: Effects of Cation Arrangement on Band Gap. Phys. Rev. B 2011, 83, 205115. (15) Zheng, F.; Xin, Y.; Huang, W.; Zhang, J.; Wang, X.; Shen, M.; Dong, W.; Fang, L.; Bai, Y.; Shen, X.; Hao, J. Above 1% Efficiency of a Ferroelectric Solar Cell Based on the Pb(Zr,Ti)O3 Film. J. Mater. Chem. A 2014, 2, 1363-1368. (16) Choi, T.; Lee, S.; Choi, Y. J.; Kiryukhin, V.; Cheong, S.-W. Switchable Ferroelectric Diode and Photovoltaic Effect in BiFeO3. Science 2009, 324, 63-66. (17) Yi, H. T.; Choi, T.; Choi, S. G.; Oh, Y. S.; Cheong, S. W. Mechanism of the Switchable Photovoltaic Effect in Ferroelectric BiFeO3. Adv. Mater. 2011, 23, 3403-3407. (18) Bhatnagar, A.; Roy Chaudhuri, A.; Heon Kim, Y.; Hesse, D.; Alexe, M. Role of Domain Walls in the Abnormal Photovoltaic Effect in BiFeO3. Nat. Commun. 2013, 4, 2835. (19) Yang, S. Y.; Seidel, J.; Byrnes, S. J.; Shafer, P.; Yang, C. H.; Rossell, M. D.; Yu, P.; Chu, Y. H.; Scott, J. F.; Ager Iii, J. W.; Martin, L. W.; Ramesh, R. Above-bandgap Voltages from Ferroelectric Photovoltaic Devices. Nat. Nanotechnol. 2010, 5, 143. (20) Choi, W. S.; Chisholm, M. F.; Singh, D. J.; Choi, T.; Jellison Jr, G. E.; Lee, H. N. Wide Bandgap Tunability in Complex Transition Metal Oxides by Site-specific Substitution. Nat. Commun. 2012, 3, 689. (21) Bai, Y.; Chen, J.; Wu, X.; Zhao, S. Photovoltaic Behaviors Regulated by Band-Gap and Bipolar Electrical Cycling in Holmium-Doped Bi5Ti3FeO15 Ferroelectric Films. J. Phys. Chem. C 2016, 120, 24637-24645. (22) Han, H.; Song, S.; Lee, J. H.; Kim, K. J.; Kim, G.-W.; Park, T.; Jang, H. M. Switchable Photovoltaic Effects in Hexagonal Manganite Thin Films Having Narrow Band Gaps. Chem. Mater. 2015, 27, 7425-7432. (23) He, J.; Franchini, C.; Rondinelli, J. M. Lithium Niobate-Type Oxides as Visible Light Photovoltaic Materials. Chem. Mater. 2016, 28, 25-29. (24) Kolb, B.; Kolpak, A. M. First-Principles Design and Analysis of an Efficient, Pb-Free Ferroelectric Photovoltaic Absorber Derived from ZnSnO3. Chem. Mater. 2015, 27, 5899-5906. (25) Inaguma, Y.; Yoshida, M.; Katsumata, T. A Polar Oxide ZnSnO3 with a LiNbO3-Type Structure. J. Am. Chem. Soc. 2008, 130, 6704-6705. (26) Son, J. Y.; Lee, G.; Jo, M.-H.; Kim, H.; Jang, H. M.; Shin, Y.-H. Heteroepitaxial Ferroelectric ZnSnO3 Thin Film. J. Am. Chem. Soc. 2009, 131, 8386-8387. (27) Omata, T.; Nagatani, H.; Suzuki, I.; Kita, M.; Yanagi, H.; Ohashi, N. Wurtzite CuGaO2: A New Direct and Narrow Band Gap Oxide Semiconductor Applicable as a Solar Cell Absorber. J. Am. Chem. Soc. 2014, 136, 3378-3381.


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Designing a Lower Bandgap Bulk Ferroelectric Material with a Sizable

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