Site-Selective Oxygen Vacancy Formation Derived from the Characteristic Crystal Structures of in Sn-Nb complex Oxides

Divalent tin oxides have attracted considerable attention as novel p-type oxide semiconductors, which are essential for realizing future oxide electronic devices. Recently, p-type Sn2Nb2O7 and SnNb2O6 were developed; however, enhanced hole mobility by reducing defect concentrations is required for practical use. In this work, we investigate the correlation between the formation of oxygen vacancy which may reduce the hole-generation efficiency and hole mobility, and the crystal structure in Sn-Nb complex oxides. Extended X-ray absorption fine structure spectroscopy and Rietveld analysis of x-ray diffraction revealed the preferential formation of oxygen vacancy at the O site bonded to the Sn ions in both the tin niobates. Moreover, a large amount of oxygen vacancy around the Sn ions were found in the p-type Sn2Nb2O7, thereby indicating the effect of oxygen vacancy to the low hole-generation efficiency. The dependence of the formation of oxygen vacancy on the crystal structure can be elucidated from the Sn-O bond strength that is evaluated based on the bond valence sum and Debye temperature. The differences in the bond strengths of the two Sn-Nb complex oxides are correlated through the steric hindrance of Sn2+ with asymmetric electron density distribution. This suggests the importance of the material design with a focus on the local structure around the Sn ions to prevent the formation of oxygen vacancy in p-type Sn2+ oxides.


INTRODUCTION
Oxide semiconductors with high electrical conductivity and transparency have attracted significant attention owing to their application in several technologies such as thin-film solar cells and/or touch screens. However, most of the existing oxide semiconductors are n-type oxides, such as Sndoped In 2 O 3 (ITO) and InGaO 3 (ZnO) 5 (IGZO) [1], and the lack of practical p-type oxide semiconductors limits their applications to unipolar devices. Fabrications of the p-n junctions, as same as technologies for Si semiconductors, will lead to more energy saving, and more complex transparent devices, and thus it is important to develop p-type oxides for the future "oxide electronics" [2].
One of the intrinsic problems in the development of p-type oxide semiconductors is the low hole mobility due to a flat valence band maximum composed of an O 2p orbital [3,4]. To reduce strong localization of holes, modifying the valence band structure by hybridization of the O 2p orbital with a metallic s or d orbital is theoretically suggested [3]. Because the s orbital has an electron density with a large spatial spread and strong interactions with the neighboring atoms, the degree of valence band dispersion is particularly large for the s orbital-based materials. Thus, oxides including cations with ns 2 electronic configuration (such as Sn 2+ , Pb 2+ , and Bi 3+ ) have attracted attention as new candidates for p-type oxides [3][4][5][6][7]. This approach was demonstrated through the fabrication of SnO films with a hole mobility of ~21 cm 2 V -1 s -1 [8][9][10][11]. However, because SnO has an indirect band gap of 0.7 eV [9], it is not suitable for future transparent device applications. For realizing both high hole mobility and transparency, p-type tin oxides with wide band gaps are required. is provided in previous reports [13,14] [13,14]. Although their valence band is composed of Sn 5s orbital, they showed low hole mobility of 0.03-0.4 cm 2 V -1 s -1 so far [12][13][14]. In general, since structural defects act as the scattering centers for electrons (holes), the mobility declines with the increase of the defect density, and hence, the reduction of these defects is a common strategy for improving hole mobility [18]. By the same time, the carriers are generated by the defect formations, therefore, a high carrier generation efficiency of the donor (accepter) is required to improve the hole mobility for practical applications. Thus, it is important to focus on the underlying mechanism of the defect formations.

Structural Characterization
The crystal structure was confirmed by X-ray diffraction (XRD), which was conducted using the Bragg-Brentano configuration with Cu Ka radiation (PANalytical, X'Pert Pro MPD). The structure parameters of Sn 2 Nb 2 O 7 were refined by the Rietveld analysis using the RIETAN-FP software [25]. The crystal structure was obtained by the VESTA software [26].
The local crystal structures were investigated by EXAFS measurements. The EXAFS spectra of all the samples were measured by the transmission mode at 40 K. The Sn K-edge and Nb K-edge spectra were observed at the beamline AR-NW10A of the Photon Factory Advanced Ring (PF-AR) and BL-9C of Photon Factory, KEK [27]. The EXAFS spectra were Fourier-transformed using the Hanning window function within the k range of 3.0-14.0 Å -1 for the Sn K-edge, Nb K-edge of SnNb 2 O 6 , and Nb K-edge of Sn 2 Nb 2 O 7 , and 3.0-13.0 Å -1 for the Sn K-edge of Sn 2 Nb 2 O 7 . The data processing of the EXAFS spectra was performed using Athena [28].
The 119 Sn Mössbauer spectroscopy was conducted by the conventional transmission method under the same experimental geometry at 78 and 300 K. The powder samples were mixed with silicone grease and plastered on pure Al foil to eliminate any orientation effect. The range of the Doppler velocity of the Ca 119m SnO 3 source was set to ±8 mm s -1 . The peak position was calibrated using a CaSnO 3 reference.

Comparing the formation of oxygen vacancies through crystal structure analysis
Before discussing the possible V ! •• formation and its crystal structure dependency, we provide an  Fig. 2(a), the peak between 1.5 and 1.7 Å in the FT of the Sn K-edge EXAFS spectra corresponds to the first-neighbor Sn-O bond. The peak intensity of the Sn-O bond in the n-type sample is distinctly weaker than that in the p-type sample. Here, the spectral intensities of the FT feature mainly reflect the coordination number (N j ) of the absorbing atom and the Debye Waller factor (s j ) of the scattering atom, which corresponds to the magnitude of thermal vibrations depending on the crystalline quality. Considering the higher annealing temperature of the n-type sample than the p-type one (|DT| = 450 K), s j of the n-type sample is naturally expected to be lower than that of the p-type one. Namely, the FT magnitude of the n-type sample could be stronger than that of the p-type one; however, this contradicts the experimental findings. Therefore, the significant decrease in the Sn-O peak intensity can be assumed to be a decrease in N j , that is, formation of V ! •• . Simultaneously, we notice that the EXAFS oscillation becomes unclear for the n-type sample (inset of Fig. 2(a)), suggesting the formation of structural disorder around the Sn ions. In contrast, in Fig. 2 (b), two peaks are observed in the FT of the Nb K-edge EXAFS spectra for radial distance (r) < 2 Å. However, a slightly complex behavior is observed in the FT spectra of the Nb K-edge compared to that of the Sn K-edge; the intensity of the first shell (r = 1.1 Å) for the n-type sample seems to be stronger than that for the p-type one; however, the second shell (r = 1.7 Å) shows an opposite behavior. As mentioned above, the FT magnitude of the n-type sample will be higher than that of p-type one, when we assume a low s j values with higher annealing temperatures for the n-type samples. Thus, this complex feature could reflect the decrease in N j . It is noteworthy that the decrease in the intensity is significant for the Sn K-edge as compared to Nb K-edge. From the spectral calculations using FEFF program [29,30], the first metal-oxide scattering peaks in the Sn K-edge (r ~1.5 Å) and Nb K-edge (r ~1.8 Å) EXAFS spectra mainly originate from the Sn-O3 bonds and Nb-O1 and Nb-O2 bonds, respectively (see supplementary S2).
Given the large decrease in the intensity of the FT EXAFS spectra for the Sn K-edge, we can consider that V ! •• is preferentially formed at O3 sites rather than at O1 and/or O2 sites. The siteselective V ! •• formation is consistent with the findings in the Rietveld analysis of the XRD patterns (see supplementary S3).

Figures 2 (c) and (d) show the FT of the Sn K-edge and Nb K-edge EXAFS spectra of Sn
respectively. The intensity of the n-type sample for both the FT spectra is slightly weaker than that of the p-type one, suggesting larger formation of V ! •• for the n-type sample. Initially, negligible differences are observed between the Sn K-edge and Nb K-edge; the preferential formation of V !

••
for n-type is not evident for Sn 2 Nb 2 O 7 as it is for SnNb 2 O 6 . However, upon closer look at the EXAFS oscillations at wavenumber (k) > 6 Å, Sn K-edge EXAFS oscillation is unclear for both p-type and n-type Sn 2 Nb 2 O 7 (inset of Fig. 2(c)), which is similar to the behavior observed for n-

Strength of Sn-O bond in Sn-Nb complex oxides
To understand the different V ! •• formations around the Sn ions between SnNb 2 O 6 and Sn 2 Nb 2 O 7 , we note the Sn-O bond strength in each structure. Bond valence sum (BVS) is a suitable indicator of the bond strength [31]. The bond valence S ij is directly related to the strength of the Sn-O bond and inversely with the bond length [32,33]. It can be approximated as where R 0 is the bond valence parameter, 1.984 Å for Sn-O bond and 1.916 Å for Nb-O bond, R ij and B are the actual bond length and 0.37 which is a constant independent of an element, respectively [34,35]. The BVS is the sum of S ij , which is expressed as follows [36].
The values of the BVS for the Sn and O of SnNb 2 O 6 and Sn 2 Nb 2 O 7 are summarized in Table 2.
The R ij (< 3 Å) values, which were evaluated from the crystal structure data of the p-type samples refined by the Rietveld analysis, are summarized in  Another indicator for the Sn-O bond strength is Debye temperature (q D ), which is based on the amplitude of thermal vibration [37]. We estimated q D of Sn 2+ for p-type SnNb 2 O 6 and Sn 2 Nb 2 O 7 by means of 119 Sn Mössbauer spectroscopy. Based on the high-temperature approximation of the Debye model, the Debye temperature is expressed with recoil-free fractions f(T) as follows [38], where A, q D , T, k B , and E R are the integral absorption intensity of Mössbauer peak, Debye temperature, measurement temperature, Boltzmann constant, and recoil energy, respectively. The temperature dependence of ln A and the estimated values of q D are shown in Fig. 4 (details are provided in supplementary section S6). The slope of Sn 2 Nb 2 O 7 was larger than that of SnNb 2 O 6 ; Therefore, the q D value of SnNb 2 O 6 (237 K) was higher than that of Sn 2 Nb 2 O 7 (174 K). This is qualitatively in good agreement with the results obtained through the BVS estimations.

Comparing local structures around the Sn ions
Finally, we will discuss the origin of the differences in the Sn-O bond strengths of SnNb 2 O 6 and Sn 2 Nb 2 O 7 by focusing on the local structures around the Sn ions. It is known that Sn 5s 2 electrons have an asymmetric electron density and a large volume as same as that of oxygen ions [23].
Moreover, they act like the lone electron pair in the Sn 2+ oxides [39,40]. The polarization of the 5s 2 electrons occurs due to the hybridization of Sn 5s with O 2p and Sn 5p orbitals [39].

Supporting Information
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