Barium Titanium Oxynitride from Ammonia-Free Nitridation of Reduced BaTiO 3

: We investigated the nitridation of reduced BaTiO 3 , BaTiO 2.60 H 0.08 , corresponding to an oxyhydride with a large concentration of O defects (>10%). The material is readily nitrided under ﬂowing N 2 gas at temperatures between 400 and 450 ◦ C to yield oxynitrides BaTiO 2.6 N x ( x = 0.2 − 0.22) with a slightly tetragonally distorted perovskite structure, a ≈ 4.01 and c ≈ 4.02 Å, and Ti partially remaining in the oxidation state III. The tetragonal structure was conﬁrmed from Raman spectroscopy. 14 N MAS NMR spectroscopy shows a single resonance at 270 ppm, which is typical for perovskite transition metal oxynitrides. However, largely different signal intensity for materials with very similar N content suggests N/O/vacancy ordering when prolonging nitridation times to hours. Diffuse reﬂectance UV-VIS spectroscopy shows a reduction of the intrinsic band gap to 2.4–2.45 eV compared to BaTiO 3 (~3.2 eV). Mott-Schottky measurements conﬁrm n-type conductivity and reveal a slight negative shift of the conduction band edge from –0.59 V (BaTiO 3 ) to ~–0.65 eV.


Introduction
In 2012 Kobayashi et al. reported that the simple perovskite BaTiO 3 can be converted to BaTiO 3-x H x (with H-contents up to x ≈ 0.6) at comparatively mild temperatures, around 600 • C, by hydride reduction [1]. The paramagnetic oxyhydride BaTiO 3-x H x attains a cubic structure in which O 2− and H − ions commonly-and as in a solid-solution-form the octahedral environment around Ti that is now in a mixed IV/III oxidation state. Shortly after, it was shown that due to the labile nature of the hydride species BaTiO 3-x H x represents a versatile precursor toward new mixed-anion compounds [2]. This includes low temperature conversion to heavily nitridized BaTiO 3 by either ammonolysis (nitridation under a stream of NH 3 ) but, remarkably, also nitridation with inert elemental N 2 [2,3]. Because of the ease of activation of N 2 , BaTiO 3-x H x has been investigated as a potential catalyst for ammonia synthesis [4].
Nitridized BaTiO 3 has been previously obtained from the ammonolysis of mixtures of BaO (BaCO 3 ) and TiO 2 at 950 • C, which produced BaTiO 2.85 N 0.1 (0.57 wt% N) [5]. By applying an excess of BaO nitrogen contents may be increased to 0.8 wt% [6]. The low-temperature conversion from reduced BaTiO 3 precursors will give even higher N contents (up to 2.3 wt% as in BaTiO 2.4 N 0.4 [3]) and provide at the same time a broader range of compositions because of their variable O content, which can be-at least partiallycontrolled through the reaction conditions in the hydride reduction. Yet there have been only few investigations into the nitridation of reduced BaTiO 3 and, in particular, little is known with respect to the properties of nitridized BaTiO 3 . According to Yajima et al. [3] and Masuda et al. [2] barium titanium oxynitride materials attain, depending on the N content, a cubic or a slightly distorted tetragonal simple perovskite structure. The tetragonal Inorganics 2021, 9,62 2 of 14 phases are non-centrosymmetric and ferroelectric. Temperature induced tetragonal-tocubic transition occurs at around 120 • C, which is close to that of BaTiO 3 .
Here we re-examine the nitridation of reduced BaTiO 3 with N 2 . We point out that reduced BaTiO 3 range from oxyhydrides where O 2− in the anion substructure is stoichiometrically replaced by H − (as in refs [1][2][3]) to materials with rather low contents of H and simultaneously large concentrations of O vacancies, BaTiO 3-x H y (x-y) . Hydride reduction of BaTiO 3 -involving the sintering of pelletized mixtures of BaTiO 3 and metal hydride (e.g., CaH 2 ) over prolonged periods of time-is not well understood, but it is most likely that O-deficient forms BaTiO 3-x represent intermediates in the formation of oxyhydrides. Reactions conditions in hydride reductions can be varied by the choice and concentration/activity of metal hydride, reaction temperature, time and also the simultaneous presence of a (pressurized) H 2 atmosphere [7][8][9][10]. Depending on conditions, highly O-deficient disordered cubic phases with x up to 0.6 and y in a range 0.04-0.25 may be obtained [10,11]. We show that such O-deficient variants show equally high reactivity toward oxynitride formation as originally investigated BaTiO 3-x H x . The sequence of materials conversions is sketched in Figure 1.
Inorganics 2021, 9, x FOR PEER REVIEW 2 of 14 with respect to the properties of nitridized BaTiO3. According to Yajima et al. [3] and Masuda et al. [2] barium titanium oxynitride materials attain, depending on the N content, a cubic or a slightly distorted tetragonal simple perovskite structure. The tetragonal phases are non-centrosymmetric and ferroelectric. Temperature induced tetragonal-to-cubic transition occurs at around 120 °C, which is close to that of BaTiO3.
Here we re-examine the nitridation of reduced BaTiO3 with N2. We point out that reduced BaTiO3 range from oxyhydrides where O 2− in the anion substructure is stoichiometrically replaced by H − (as in refs [1][2][3]) to materials with rather low contents of H and simultaneously large concentrations of O vacancies, BaTiO3-xHy◻(x-y). Hydride reduction of BaTiO3-involving the sintering of pelletized mixtures of BaTiO3 and metal hydride (e.g., CaH2) over prolonged periods of time-is not well understood, but it is most likely that O-deficient forms BaTiO3-x represent intermediates in the formation of oxyhydrides. Reactions conditions in hydride reductions can be varied by the choice and concentration/activity of metal hydride, reaction temperature, time and also the simultaneous presence of a (pressurized) H2 atmosphere [7][8][9][10]. Depending on conditions, highly O-deficient disordered cubic phases with x up to 0.6 and y in a range 0.04-0.25 may be obtained [10,11]. We show that such O-deficient variants show equally high reactivity toward oxynitride formation as originally investigated BaTiO3-xHx. The sequence of materials conversions is sketched in Figure 1.

BTON Synthesis and Characterization
The starting material for nitridation was obtained from reacting BaTiO3 with CaH2 at 550 °C for 48 h and corresponded to BaTiO2.60H0.08◻0.32, in the following abbreviated as BTOH. BTOH was characterized by powder X-ray diffraction (PXRD), thermogravimetric (TG) analysis, and 1 H Magic Angle Spinning (MAS) Nuclear Magnetic Resonance (NMR) spectroscopy. The TG curve is shown in Figure 2 whereas the PXRD and 1 H-MAS NMR analyses are detailed as supplementary material, Figures S1 and S2. The characterization results are summarized in Table 1. BTOH has a cubic lattice parameter of a = 4.022 Å resulting in a unit cell volume slightly higher than BaTiO3. From TG analysis the H content and/or O defect concentration of reduced BaTiO3 can be assessed [1,11,12]. An experiment under flowing air monitors the reactions BaTiO3-xHx + 0.75x O2 → BaTiO3 + 0.5x H2O (1) and BaTiO3-x + 0.5x O2 → BaTiO3 (2). Typically there is a small initial weight loss (between 0.1 and 0.15%) in the temperature range up to 300 °C which stems from the loss of surface hydroxyl and secondary water [13,14]. The onset of oxidation is seen at around 400 °C. Oxidation appears to be completed above 650 °C and leads to a weight increase of 2.79%.
The associated x values according to Equations (1) and (2) are contained in Table 1. 1 H-MAS-NMR then revealed a rather small H content in BTOH, x ≈ 0.08, and accordingly BTOH possesses a high concentration of O defects.

BTON Synthesis and Characterization
The starting material for nitridation was obtained from reacting BaTiO 3 with CaH 2 at 550 • C for 48 h and corresponded to BaTiO 2.60 H 0.08 0.32 , in the following abbreviated as BTOH. BTOH was characterized by powder X-ray diffraction (PXRD), thermogravimetric (TG) analysis, and 1 H Magic Angle Spinning (MAS) Nuclear Magnetic Resonance (NMR) spectroscopy. The TG curve is shown in Figure 2 whereas the PXRD and 1 H-MAS NMR analyses are detailed as Supplementary Materials, Figures S1 and S2. The characterization results are summarized in Table 1. BTOH has a cubic lattice parameter of a = 4.022 Å resulting in a unit cell volume slightly higher than BaTiO 3 . From TG analysis the H content and/or O defect concentration of reduced BaTiO 3 can be assessed [1,11,12]. An experiment under flowing air monitors the reactions BaTiO 3-x H x + 0.75x O 2 → BaTiO 3 + 0.5x H 2 O (1) and BaTiO 3-x + 0.5x O 2 → BaTiO 3 (2). Typically there is a small initial weight loss (between 0.1 and 0.15%) in the temperature range up to 300 • C which stems from the loss of surface hydroxyl and secondary water [13,14]. The onset of oxidation is seen at around 400 • C. Oxidation appears to be completed above 650 • C and leads to a weight increase of 2.79%. The associated x values according to Equations (1) and (2) are contained in Table 1. 1 H-MAS-NMR then revealed a rather small H content in BTOH, x ≈ 0.08, and accordingly BTOH possesses a high concentration of O defects.    Figure 3 shows combined TG and differential scanning calorimetry (DSC) traces of BTOH when heated in a stream of N2 to 500 °C (a) and 600 °C (b) at a rate of 10 °C/min. The exothermic event at around 435 °C, concomitant with a weight increase of 0.99% and 1.38% respectively, is attributed to the formation of a nitridated phase, in the following abbreviated as BTON. The onset of nitridation occurs at a similar temperature as oxidation, perhaps relating to the activation of labile H − , but appears to be a sudden process compared to more gradual oxidation. Figure 3c compares TG traces of BTOH when heating for a prolonged period of time after reaching the target temperatures 500 and 600 °C, respectively. There was no further weight increase with time in the 600 °C experiment, which implies that the nitridation reaction had already been completed when reaching this temperature. In the 500 °C experiment, continued weight increase was seen for about additionally 15 min. At the same time the TG trace plateaued at a lower weight gain level, which suggests that the 600 °C experiment resulted in a product with a slightly higher N content (by 15%). The N content corresponding to the TG weight increase is x = 0.19 and 0.22 for the 500 and 600 °C experiment, respectively. The N content was also determined by chemical analysis for samples obtained at 600 °C without annealing (BTON-600C-0 h) and after 6 h prolonged nitridation (BTON-600C-6 h) and found to be in good agreement with the TG results. The characterization results for BTON samples are summarized in Table 2 and as supplementary material. It is important to emphasize that nitridation experiments require strictly air-free conditions, otherwise (partial) oxidation will occur. In air BTON samples are stable up to about 400-500 °C.   Figure 3 shows combined TG and differential scanning calorimetry (DSC) traces of BTOH when heated in a stream of N 2 to 500 • C (a) and 600 • C (b) at a rate of 10 • C/min. The exothermic event at around 435 • C, concomitant with a weight increase of 0.99% and 1.38% respectively, is attributed to the formation of a nitridated phase, in the following abbreviated as BTON. The onset of nitridation occurs at a similar temperature as oxidation, perhaps relating to the activation of labile H − , but appears to be a sudden process compared to more gradual oxidation. Figure 3c compares TG traces of BTOH when heating for a prolonged period of time after reaching the target temperatures 500 and 600 • C, respectively. There was no further weight increase with time in the 600 • C experiment, which implies that the nitridation reaction had already been completed when reaching this temperature. In the 500 • C experiment, continued weight increase was seen for about additionally 15 min. At the same time the TG trace plateaued at a lower weight gain level, which suggests that the 600 • C experiment resulted in a product with a slightly higher N content (by 15%). The N content corresponding to the TG weight increase is x = 0.19 and 0.22 for the 500 and 600 • C experiment, respectively. The N content was also determined by chemical analysis for samples obtained at 600 • C without annealing (BTON-600C-0 h) and after 6 h prolonged nitridation (BTON-600C-6 h) and found to be in good agreement with the TG results. The characterization results for BTON samples are summarized in Table 2 and as Supplementary Materials. It is important to emphasize that nitridation experiments require strictly air-free conditions, otherwise (partial) oxidation will occur. In air BTON samples are stable up to about 400-500 • C.  The composition of nitridized BaTiO3 can be expressed as BaTiO3-xN2x/3. Accordingly, the inherent O defect concentration is x/3 [3]. The upper limit for x for reduced BaTiO3 appears to be x ≈ 0.6 [1]. Thus the composition BaTiO2.4N0.4, as reported by Yajima et al. [3], should represent the highest achievable N content for Ba oxynitrides. With our BTOH starting material, the composition after nitridation was expected to be BaTiO2.6N0.267. This  The composition of nitridized BaTiO 3 can be expressed as BaTiO 3-x N 2x/3 . Accordingly, the inherent O defect concentration is x/3 [3]. The upper limit for x for reduced BaTiO 3 appears to be x ≈ 0.6 [1]. Thus the composition BaTiO 2.4 N 0.4 , as reported by Yajima et al. [3], should represent the highest achievable N content for Ba oxynitrides. With our BTOH starting material, the composition after nitridation was expected to be BaTiO 2.6 N 0.267 . This N content is clearly not achieved, which implies that Ti partially remains as Ti 3+ (about 5%). Thus, our BTON materials should be paramagnetic and possibly metallic. To accommodate for this situation, the formula for nitridized BaTiO 3 is rewritten as BaTiO 3-x N (2x/3)-y (x/3)+y where y = 0 refers to the stoichiometric, diamagnetic case. The application of higher nitridation temperatures and/or changed heating rates did not result in fully nitridized BTON samples (cf. Supplementary Materials, Figure S3). Figure 4 shows scanning electron microscopy (SEM) images of the materials constituting the sequence of conversions (cf. Figure 1). The average particle size of tetragonal BaTiO 3 used for hydride reduction was specified by the manufacturer as 0.5 micron. This is confirmed in Figure 4a, although some particle size distribution is noticeable. The particle morphology of BaTiO 3 is retained in BTOH (Figure 4b). However, upon nitridation the surface topography of particles becomes noticeable changed (Figure 4c,d). The step-like features suggest recrystallization, involving ion transport during the solid-gas reaction. Thus the nitridation process of reduced BaTiO 3 may not be truly topochemical.
Inorganics 2021, 9, x FOR PEER REVIEW N content is clearly not achieved, which implies that Ti partially remains as Ti 3 5%). Thus, our BTON materials should be paramagnetic and possibly meta accommodate for this situation, the formula for nitridized BaTiO3 is rewritten as xN(2x/3)-y◻(x/3)+y where y = 0 refers to the stoichiometric, diamagnetic case. The applic higher nitridation temperatures and/or changed heating rates did not result nitridized BTON samples (cf. supplementary material, Figure S3). Figure 4 shows scanning electron microscopy (SEM) images of the materials tuting the sequence of conversions (cf. Figure 1). The average particle size of tet BaTiO3 used for hydride reduction was specified by the manufacturer as 0.5 micro is confirmed in Figure 4a, although some particle size distribution is noticeable. T ticle morphology of BaTiO3 is retained in BTOH (Figure 4b). However, upon nitr the surface topography of particles becomes noticeable changed (Figure 4c,d). Th like features suggest recrystallization, involving ion transport during the solid-g tion. Thus the nitridation process of reduced BaTiO3 may not be truly topochemic   Figures S4-S13, Tables S1 and S2). The tetragonal structure of is clearly noticeable in the laboratory PXRD pattern, a = 3.996, c = 4.031 Å (c/a = whereas the patterns for BTOH and BTON appear cubic, with the unit cell vo BTON slightly smaller. Compared to BTOH one could infer a slight broadening reflections (200), (201), (211) (in the 2θ range 20-40°), which most clearly would sh tetragonal split. However, the refinement of laboratory data does not allow for d nation between cubic and tetragonal phase (c/a < 1.003), Figure 5b. The tetragona ture of BTON is then clearly revealed from Raman spectroscopy, as shown in F Thus, we chose to present the result of the tetragonal refinement (c/a ≈ 1.002(2)) i 2, whereas the result of the cubic refinements is given in Table S2.   Figures S4-S13, Tables S1 and S2). The tetragonal structure of BaTiO 3 is clearly noticeable in the laboratory PXRD pattern, a = 3.996, c = 4.031 Å (c/a = 1.009), whereas the patterns for BTOH and BTON appear cubic, with the unit cell volume of BTON slightly smaller. Compared to BTOH one could infer a slight broadening of the reflections (200), (201), (211) (in the 2θ range 20-40 • ), which most clearly would show the tetragonal split. However, the refinement of laboratory data does not allow for discrimination between cubic and tetragonal phase (c/a < 1.003), Figure 5b. The tetragonal structure of BTON is then clearly revealed from Raman spectroscopy, as shown in Figure 6. Thus, we chose to present the result of the tetragonal refinement (c/a ≈ 1.002(2)) in Table 2, whereas the result of the cubic refinements is given in Table S2.  The Raman spectrum of tetragonal BaTiO3 is well understood [15][16][17]. All twelve optical modes 3×A1 + B1 + 4×E are Raman active. Poly/nanocrystalline samples are characterized by sharp bands at about 170 and 305 cm −1 and asymmetric broader bands at about 260, 515, and 720 cm −1 [18,19]. Furthermore, a pronounced dip at around 180 cm −1 is also characteristic, which has been attributed to interference from anharmonic coupling between three A1 modes [15]. Lastly, the E1 softmode (x,y displacement of Ti) expresses itself with a sharp rise of Raman scattering intensity at lowest wavenumbers [19]. The BTOH spectrum is featureless, in agreement with a cubic perovskite structure where all modes are Raman inactive. The Raman spectra of the BTON samples show broad bands at 347, 506 and 716 cm −1 and sharp bands at 100, 180 and 302 cm −1 . The latter (assigned to the nonpolar B1 mode) is characteristic for a tetragonal phase [20,21]. The same holds true for the band at 720 cm −1 which would be extremely broad and week in the cubic para-electric phase [20]. Because of the considerably lower tetragonality of the BTON phases compared  The Raman spectrum of tetragonal BaTiO3 is well understood [15][16][17]. All twelve optical modes 3×A1 + B1 + 4×E are Raman active. Poly/nanocrystalline samples are characterized by sharp bands at about 170 and 305 cm −1 and asymmetric broader bands at about 260, 515, and 720 cm −1 [18,19]. Furthermore, a pronounced dip at around 180 cm −1 is also characteristic, which has been attributed to interference from anharmonic coupling between three A1 modes [15]. Lastly, the E1 softmode (x,y displacement of Ti) expresses itself with a sharp rise of Raman scattering intensity at lowest wavenumbers [19]. The BTOH spectrum is featureless, in agreement with a cubic perovskite structure where all modes are Raman inactive. The Raman spectra of the BTON samples show broad bands at 347, 506 and 716 cm −1 and sharp bands at 100, 180 and 302 cm −1 . The latter (assigned to the nonpolar B1 mode) is characteristic for a tetragonal phase [20,21]. The same holds true for the band at 720 cm −1 which would be extremely broad and week in the cubic para-electric phase [20]. Because of the considerably lower tetragonality of the BTON phases compared The Raman spectrum of tetragonal BaTiO 3 is well understood [15][16][17]. All twelve optical modes 3×A 1 + B 1 + 4×E are Raman active. Poly/nanocrystalline samples are characterized by sharp bands at about 170 and 305 cm −1 and asymmetric broader bands at about 260, 515, and 720 cm −1 [18,19]. Furthermore, a pronounced dip at around 180 cm −1 is also characteristic, which has been attributed to interference from anharmonic coupling between three A 1 modes [15]. Lastly, the E 1 softmode (x,y displacement of Ti) expresses itself with a sharp rise of Raman scattering intensity at lowest wavenumbers [19]. The BTOH spectrum is featureless, in agreement with a cubic perovskite structure where all modes are Raman inactive. The Raman spectra of the BTON samples show broad bands at 347, 506 and 716 cm −1 and sharp bands at 100, 180 and 302 cm −1 . The latter (assigned to the non-polar B 1 mode) is characteristic for a tetragonal phase [20,21]. The same holds true for the band at 720 cm −1 which would be extremely broad and week in the cubic para-electric phase [20]. Because of the considerably lower tetragonality of the BTON phases compared to BaTiO 3 , bands are generally broader and weaker. However, there are two clear major differences. The softmode in BaTiO 3 appears stiffened and the corresponding band is at around 100 cm −1 in the spectra of the BTON phases. Further, the band at around 260 cm −1 in BaTiO 3 (A 1 (2) ) is not apparent in BTON and instead a band at 347 cm −1 is visible. The A 1 (2) mode is the mutual displacement of Ba atoms and the TiO 6/2 octahedron framework in the z direction [19]. It is not clear why this mode has either a very low intensity for BTON or is possibly blue-shifted by almost 90 cm −1 .
The Raman spectra of the four different BTON samples are very much alike (possibly that the B 1 peak at~302 cm −1 is somewhat sharper for the annealed samples, especially for BTON-600C-6 h). This is in contrast with 14 N-MAS-NMR spectra, shown in Figure 7. The nuclear quadrupole moment of the 14 N isotope implies drastic broadening of the resonance lines unless N is in a highly symmetric environment for which the quadrupolar coupling constant is relatively small. This is the case for perovskite oxynitrides and 14 N-MAS-NMR has been used previously to probe N [5,[22][23][24]. Somewhat surprising and still not explained, 14 N expresses as a single resonance at 270 ppm, regardless the kind of transition metal (Ti, Nb, Ta) or A component (Ca, Sr, Ba, La, Ce) in the perovskite oxynitride [25]. In our spectra the 270 ppm resonance is clearly there, but the signal intensity is considerably less for BTON-600-0 h compared to BTON-600C-6 h although the two materials possess virtually the same N content. We speculate that quadrupolar coupling is rather large for BTON-600C-0 h and that prolonged nitridation at 600 • C resulted in a more ordered N/O/vacancy structure for BTON-600C-6 h (cf. Figure 1c). 1 H-MAS-NMR spectra show clearly the absence of hydridic H in BTON-600-0h and BTON-600-6h. As earlier shown, hydridic H as part of the bulk structure can be discriminated from protic H (from (surface) hydroxyl and secondary water) as negative and positive chemical shift contributions, respectively [11,12]. Hydric H expresses itself as a single broad resonance peak at negative chemical shift whereas protic resonances appear at~1 ppm (surface OH species) and in the region 6-7 ppm (secondary water). Interestingly one can recognize two H − environments for BTOH, one centered at −4 ppm and one at −55 pm. The occurrence of two hydridic environments in reduced BaTiO 3 is highly unusual and perhaps relates to the large O defect concentration of BTOH (see Supplementary Materials for more details). Although the phenomenon is interesting and deserves more detailed investigations, it is not considered important for the outcome of the nitridation reactions.

Spectroscopic and Mott-Schottky Analysis
It is well known that BaTiO 3 is a wide gap semiconductor with a band gap energy of 3.2-3.4 eV [26,27]. According to first principles calculations, the band gap is indirect for the tetragonal form and direct for the cubic (high temperature) polymorph [28]. However, experimentally it appears to be difficult to assign the nature of the band gap and frequently a direct band gap is also assumed for the tetragonal form. Mixed anion oxynitrides show reduced band gaps, by ca 1 eV, since the top of the valence band constitutes N-2p based states, which are at higher energy than O-2p [29]. Perovskite oxynitrides involving Ti are typically based on lanthanoid metals as the perovskite A component and ATiO 2 N (A = La, Nd, Ce, Pr, Eu) have been extensively investigated because of their photocatalytic activity and dielectric properties [30][31][32]. As outlined above, when A corresponds to an alkaline earth metal, the reduced valence has to be compensated with defects in the anion substructure, i.e., BaTiO 3-x N (2x/3)-y (x/3)+y where y = 0 refers to the fully nitridized, diamagnetic case. For these materials properties are not well investigated.

Spectroscopic and Mott-Schottky Analysis
It is well known that BaTiO3 is a wide gap semiconductor with a band gap energy of 3.2-3.4 eV [26,27]. According to first principles calculations, the band gap is indirect for the tetragonal form and direct for the cubic (high temperature) polymorph [28]. However, experimentally it appears to be difficult to assign the nature of the band gap and frequently a direct band gap is also assumed for the tetragonal form. Mixed anion oxynitrides show reduced band gaps, by ca 1 eV, since the top of the valence band constitutes N-2p based states, which are at higher energy than O-2p [29]. Perovskite oxynitrides involving Ti are typically based on lanthanoid metals as the perovskite A component and ATiO2N (A = La, Nd, Ce, Pr, Eu) have been extensively investigated because of their photocatalytic activity and dielectric properties [30][31][32]. As outlined above, when A corresponds to an alkaline earth metal, the reduced valence has to be compensated with defects in the anion substructure, i.e., BaTiO3-xN(2x/3)-y◻(x/3)+y where y = 0 refers to the fully nitridized, diamagnetic case. For these materials properties are not well investigated. Figure 8 compares the UV-VIS diffuse reflectance spectra of the BTON-500C and BTON-600C samples with BTOH and pristine BaTiO3. As earlier reported, reduced BaTiO3 materials exhibit a broad absorption in the visible and near-IR region (Figure 8a) [12]. The actual absorption edge, corresponding to interband transitions, is around 400 nm, which corresponds well to the band gap energy of BaTiO3. The near IR-VIS absorption enters a minimum upon approaching the upturn of the absorption edge, which is typical of (intraband) free carrier absorption in heavily doped semiconductors [33]. In metallic BTOH states at the bottom of the conduction band are occupied. The absorption edge of BTON is significantly red-shifted. In addition, the exponential onset of absorption indicates that   (Figure 8a) [12]. The actual absorption edge, corresponding to interband transitions, is around 400 nm, which corresponds well to the band gap energy of BaTiO 3 . The near IR-VIS absorption enters a minimum upon approaching the upturn of the absorption edge, which is typical of (intraband) free carrier absorption in heavily doped semiconductors [33]. In metallic BTOH states at the bottom of the conduction band are occupied. The absorption edge of BTON is significantly red-shifted. In addition, the exponential onset of absorption indicates that charge carriers are localized in traps below the conduction band [34]. So in contrast with reduced BaTiO 3 where electrons form delocalized band states [35], they appear localized in polarons in not fully nitridized BTON. charge carriers are localized in traps below the conduction band [34]. So in contrast with reduced BaTiO3 where electrons form delocalized band states [35], they appear localized in polarons in not fully nitridized BTON.  Figure 8b shows the Kubelka-Munk (K-M) transformed spectra, which-for comparability across the samples-was normalized. The K-M transformed spectrum of metallic BTOH is different from the BTON materials because of the free carrier absorption. Its absorption edge is slightly (by 0.1 eV) shifted upwards compared to BaTiO3, which is in line with earlier observations [12]. The absorption edge of BTON is in a range 500-550 nm. Tauc plots (presented as Supplementary Material, Figure S14) suggest band gaps in a range 2.4-2.5 eV as summarized in Table 3, which agrees with the green-yellow color of the material.
Mott-Schottky measurements were carried out for BTON-600-0 h and BTON-600-6 h as well as the BTOH precursor and the results compared with pristine BaTiO3 (Figure 9). All materials show a positive slope, confirming the semiconducting n-type behavior. The flat band potential of BaTiO3 is −0.59 V vs. Reversible Hydrogen Electrode (RHE). In case of a n-type semiconductor, the flat band potential is close to the bottom of the conduction band [36]. BTOH exhibits a flat band potential at −0.54 V vs. RHE. A slight positive shift of the flat potential, which may reach up to −0.36 V compared to the pristine BaTiO3, has also been observed for other reduced materials BaTiO3-xHy◻(x-y) with different x,y [12]. In contrast, BTON displays a negative shift −0.65 V, without any significant difference between the differently annealed samples (cf. Table 3). The slopes of BaTiO3 and BTOH are very similar, although the charge carrier density of the latter is considerably higher. Since the slope is also depending on the dielectric constant, this was explained with a much smaller dielectric constant of cubic BTOH [12]. The dielectric constant of tetragonal BTON  Figure 8b shows the Kubelka-Munk (K-M) transformed spectra, which-for comparability across the samples-was normalized. The K-M transformed spectrum of metallic BTOH is different from the BTON materials because of the free carrier absorption. Its absorption edge is slightly (by 0.1 eV) shifted upwards compared to BaTiO 3 , which is in line with earlier observations [12]. The absorption edge of BTON is in a range 500-550 nm. Tauc plots (presented as Supplementary Materials, Figure S14) suggest band gaps in a range 2.4-2.5 eV as summarized in Table 3, which agrees with the green-yellow color of the material.
Mott-Schottky measurements were carried out for BTON-600-0 h and BTON-600-6 h as well as the BTOH precursor and the results compared with pristine BaTiO 3 (Figure 9). All materials show a positive slope, confirming the semiconducting n-type behavior. The flat band potential of BaTiO 3 is −0.59 V vs. Reversible Hydrogen Electrode (RHE). In case of a n-type semiconductor, the flat band potential is close to the bottom of the conduction band [36]. BTOH exhibits a flat band potential at −0.54 V vs. RHE. A slight positive shift of the flat potential, which may reach up to −0.36 V compared to the pristine BaTiO 3 , has also been observed for other reduced materials BaTiO 3-x H y (x-y) with different x,y [12]. In contrast, BTON displays a negative shift −0.65 V, without any significant difference between the differently annealed samples (cf. Table 3). The slopes of BaTiO 3 and BTOH are very similar, although the charge carrier density of the latter is considerably higher. Since the slope is also depending on the dielectric constant, this was explained with a much smaller dielectric constant of cubic BTOH [12]. The dielectric constant of tetragonal BTON will relate more to that of tetragonal BaTiO 3 and, thus, the smaller slope compared to BaTiO 3 reflects its higher charge carrier density.
will relate more to that of tetragonal BaTiO3 and, thus, the smaller slope compared to Ba-TiO3 reflects its higher charge carrier density.

Synthesis of the Materials
As starting materials we used BaTiO3 (500 nm particle size, 99.9% purity, ABCR GmbH, Karlsruhe, Germany) and CaH2 powder (95%, Sigma Aldrich, Darmstadt, Germany). Prior to use, BaTiO3 was dried overnight in an oven at 200 °C. All steps of sample preparation were performed in an Ar filled glove box. For a typical synthesis ~0.25 g (1.05 mmol) of BaTiO3 was intimately mixed with 0.053 g (2.36 mmol) CaH2, i.e., the BaTiO3:H ratio was about 4.5. The BaTiO3/CaH2 mixture was subsequently pressed into a pellet with a diameter of 9 mm. The pellet was then sealed inside a stainless steel ampule (with dimensions 10 mm ID and 50 mm length), which-after removal from the glove box-was placed in a box furnace. A K-type thermocouple was inserted and located in close proximity to the metal ampule. Ampules were heated for 48 h at 550 °C. After cooling to room temperature, ampules were opened and the products washed with 0.1 M acetic acid (HAc) solution to remove excess CaH2 and metal oxide/hydroxide formed during hydride reduction. For washing, the pellets were crushed and sonicated with 50 mL HAc and then centrifuged. The procedure was repeated 3 times. As a last step, products were treated with 20 mL 95% ethanol and then dried at 120 °C under dynamic vacuum (<10 −5 bar). The dried product corresponded to dark blue powder. Portions of about 20 mg of the so obtained oxyhydride material were then heated in a 100 mL/min nitrogen stream (99.9999% (6N) purity) at a rate of 10 °C/min to various temperatures using a Netzsch ® STA 449 F1 Jupiter thermal analysis apparatus (Netzsch, Selb, Germany). The reactions were protected by an oxygen getter to avoid oxidation of oxyhydride materials. The color or the products after nitridation was changed to yellow-green (cf. Figure 5a).

Synthesis of the Materials
As starting materials we used BaTiO 3 (500 nm particle size, 99.9% purity, ABCR GmbH, Karlsruhe, Germany) and CaH 2 powder (95%, Sigma Aldrich, Darmstadt, Germany). Prior to use, BaTiO 3 was dried overnight in an oven at 200 • C. All steps of sample preparation were performed in an Ar filled glove box. For a typical synthesis~0.25 g (1.05 mmol) of BaTiO 3 was intimately mixed with 0.053 g (2.36 mmol) CaH 2 , i.e., the BaTiO 3 :H ratio was about 4.5. The BaTiO 3 /CaH 2 mixture was subsequently pressed into a pellet with a diameter of 9 mm. The pellet was then sealed inside a stainless steel ampule (with dimensions 10 mm ID and 50 mm length), which-after removal from the glove boxwas placed in a box furnace. A K-type thermocouple was inserted and located in close proximity to the metal ampule. Ampules were heated for 48 h at 550 • C. After cooling to room temperature, ampules were opened and the products washed with 0.1 M acetic acid (HAc) solution to remove excess CaH 2 and metal oxide/hydroxide formed during hydride reduction. For washing, the pellets were crushed and sonicated with 50 mL HAc and then centrifuged. The procedure was repeated 3 times. As a last step, products were treated with 20 mL 95% ethanol and then dried at 120 • C under dynamic vacuum (<10 −5 bar). The dried product corresponded to dark blue powder. Portions of about 20 mg of the so obtained oxyhydride material were then heated in a 100 mL/min nitrogen stream (99.9999% (6N) purity) at a rate of 10 • C/min to various temperatures using a Netzsch ® STA 449 F1 Jupiter thermal analysis apparatus (Netzsch, Selb, Germany). The reactions were protected by an oxygen getter to avoid oxidation of oxyhydride materials. The color or the products after nitridation was changed to yellow-green (cf. Figure 5a).

Powder X-ray Diffraction (PXRD) Analysis
PXRD patterns were collected on a Panalytical X'Pert PRO diffractometer (Malvern Panalytical, Malvern, UK) operated with Cu Kα radiation and in θ-2θ diffraction geometry. Powder samples were mounted on a Si wafer zero-background holder and diffraction patterns measured in a 2θ range of 20-120 • with 0.016 • step size. The contribution of Kα 2 radiation to the PXRD patterns was removed using the Panalytical X'Pert HighScore Plus software (Malvern Panalytical, Malvern, UK). The Rietveld method as implemented in the FullProf program was used for structure and phase analysis [37]. A six-coefficient polynomial function was applied for the background. The peak shape was described by a Pseudo-Voigt function. Site occupancies for the O atoms could not be refined reliably and were constrained to 1.

Scanning Electron Microscopy (SEM) Investigations
SEM investigations were carried out using a JEOL JSM-7000F microscope (Jeol, Tokyo, Japan) equipped with a Schottky-type field emission gun (Oxford Instruments, Oxfordshire, UK). Images were recorded in second-electron mode with an accelerating voltage of 15 kV. Samples were prepared by first producing a homogeneous suspension of particles in ethanol by sonication. Then droplets of the suspension were applied onto surface-polished Aluminum pin stubs and left to dry.

Raman Spectroscopy
Raman spectra were measured using a Labram HR 800 spectrometer (Horiba, Kyoto, Japan) equipped with an 800 mm focal length spectrograph and an air cooled, back thinned CCD detector. Powdered sample was placed on a glass slide and excited with a double frequency Nd:YAG laser (532 nm). To avoid decomposition/oxidation the lowest possible laser power (ND filter 1%, corresponding to a power density of 5.5 × 10 −6 mW·µm −2 ) was applied for the oxyhydride and oxynitride samples. Spectra were collected at room temperature in a range 0-1800 cm −1 and using a grating of 600 lines/mm. For oxynitride samples large exposure times (of up to 25 min) were employed.

Thermogravimetric Analysis (TGA)
TGA experiments were carried out using a TA instruments Discovery system (TA Instruments, New Castle, DE, USA). The samples (~15 mg powders) were heated in a platinum crucible from room temperature to 900 • C with a heating rate of 10 • C/min. A dry air gas flow of 20 mL/min was applied.

UV-VIS Diffuse Reflectance Spectroscopy
UV-VIS diffuse reflectance measurements were performed at room temperature on finely ground samples. Spectra were recorded in the range from 200 to 800 nm with an Agilent Cary 5000 UV-VIS-NIR spectrometer (Agilent, Santa Clara, CA, USA) equipped with a diffuse reflectance accessory from Harrick (Harrick Scientific Products Inc, New York, NY, USA). A polytetrafluoroethylene (PTFE) pellet (100% reflectance) was used as the reference.

Magic Angle Spinning (MAS) NMR Spectroscopy
The 1 H and 14 N MAS NMR experiments were performed at a magnetic field of 14.1 T (600.1 and 43.4 MHz Larmor frequencies for 1 H and 14 N, respectively) and MAS frequency of 60.00 kHz on a Bruker Avance-III spectrometer (Bruker, Billerica, MA, USA) equipped with a 1.3 mm MAS HX probe. Proton spectra were acquired using a rotor-synchronized, double-adiabatic spin-echo sequence with a 90 • excitation pulse of 1.1 µs followed by a pair of 50 µs tanh/tan short high-power adiabatic pulses (SHAPs) with 5 MHz frequency sweep [38,39]. All pulses were applied at a nutation frequency of 210 kHz. In total, 256 signal transients with a 5 s recycle delay were accumulated for each sample. The 14 N MAS spectra were collected using a Hahn-echo sequence with 3.0/6.0 µs 90 • /180 • rf pulses and 131,072 scans collected per sample using a 1 s relaxation delay. 1 H shifts were referenced with respect to tetramethylsilane (TMS) at 0 ppm, whereas 14 N shifts to solid NH 4 Cl at 0 ppm (−342.4 ppm with respect to nitromethane).

Mott-Schottky Measurements
Powder samples were assembled into particle-base thin film electrodes on fluorinedoped tin oxide (FTO) glass substrate via electrophoretic deposition. The homogeneous suspension was prepared by mixing 30 mg powder samples and 10 mg iodine in 30 mL of acetone via sonication. Two pre-cleaned FTO slides were immersed in the above suspension with 1 cm distance, and a DC potential of 30 V was applied between them for 3 min. After Electrophoretic Deposition (EPD) process, the FTO slides with thin film were washed with water and dried naturally. 0.1 M NaOH (pH 13), in which oxyhydrides are stable, was used as electrolyte to perform Mott-Schottky measurements. The measurements were conducted in the conventional three-electrode system where a platinum wire, a 1 M Ag/AgCl electrode and the thin film on FTO were used as counter electrode, reference electrode and working electrode, respectively. The potentials were recorded versus 1 M Ag/AgCl and converted versus reversible hydrogen electrode (RHE) according to E RHE (V) = E Ag/AgCl + (0.059 × pH) + E θ Ag/AgCl . The measurements were carried out by using the Gamry INTERFACE 1010T Potentiostat/Galvanostat/ZRA workstation (Gamry Instruments, Warminster, PA, USA) in darkness at ac amplitude of 5 mV and a frequency of 10 Hz.

Conclusions
It was shown that reduced BaTiO 3 containing simultaneously hydride and defects in the anion substructure, BaTiO 3-x H y (x-y) , can be readily nitrided with elemental N 2 at temperatures around 450 • C to yield tetragonal phases BaTiO 3-x N (2x/3)-y (x/3)+y with c/a ≈ 1.002. In contrast with the nitridation of vacancy free oxyhydride BaTiO 3-x H x , oxynitrides from BaTiO 3-x H y (x-y) appear not fully nitrided and titanium remains partially as Ti 3+ , i.e., y = 0. Electrons are localized in traps below the conduction band. The band gap energy of BaTiO 3-x N (2x/3)-y (x/3)+y materials is in a range 2.4-2.5 eV. Since the conduction band edge lies above the reduction potential for water and optical absorption is in the visible light region, it will be interesting to investigate photocatalytic properties.
Supplementary Materials: The following are available online at https://www.mdpi.com/article/ 10.3390/inorganics9080062/s1, Figure S1: Rietveld fit to the PXRD data of the oxyhydride starting material BaTiO 2.60 H 0.08 0.32 (BTOH), Figure S2: 1 H MAS NMR spectrum of BTOH showing the division into protic and hydridic signal intensity and explaining the quantification of hydridic H content, Figure S3: Simultaneously measured TG and DSC traces for BTOH upon heating under a dry N 2 stream up to 700 • C with a heating rate 10 • C/min and to 600 • C with a heating rate 5 • C/min, Figure S4-S13: Rietveld fits to PXRD data of BTON samples, Figure S14: Tauc plots for band gap evaluation, Table S1: Refinement result summary for BTON samples using the tetragonal structure model, Table S2: Refinement result summary for BTON samples using the cubic structure model.  Data Availability Statement: All data presented is available in this manuscript.