Electron Paramagnetic Resonance Study on Oxygen Vacancies and Site Occupations in Mg-Doped BaTiO3 Ceramics

Nominal (Ba1−xMgx)TiO3 (x = 0.015) (BM1T) and (Ba1−xMgx)TiO3 (x = 0.03–0.20) (BMT) ceramics were prepared by the mixed-oxide route at sintering temperatures (Ts) of 1200−1400 °C and 1200 °C, respectively. The solubility limit of Mg2+ in BMT was determined by XRD to be x = 0.05, and evidence was found for occupation of the A site by Mg2+. Electron paramagnetic resonance (EPR) was employed as a key technique to investigate the effect of Ts on oxygen vacancies in BM1T. The structure of BM1T changed from pseudocubic at Ts = 1200 °C to tetragonal at 1300 °C to mixed phases of hexagonal and tetragonal at 1400 °C. When Ts ≥ 1300 °C, a g = 1.956 EPR signal was observed at T = −188 °C and assigned as ionized oxygen vacancies. Mg2+ exhibited amphoteric behavior of substituting for the double cation sites. When Ts = 1400 °C, B-site Mg2+ and oxygen vacancies mainly existed in the hexagonal phase and A-site Mg2+ was dominant in the tetragonal phase. The higher tan δ was attributed to the higher concentrations of oxygen vacancies and Ti3+ in the hexagonal phase.

Many transition metal ions with lower valence states occupy the B site when doped in BaTiO 3 , for example, Mn 2+ [7,8]. Similar to these dopants, Mg 2+ was also considered to be substituted for the B site as an acceptor because 6-CN Mg 2+ is closer to Ti 4+ in ionic size, and the defect notation was written as Mg Ti according to the defect notation proposed by Kröger and Vink [9]. Considering the electroneutrality, Mg Ti was usually compensated by one oxygen vacancy (V •• O ) and Mg Ti − V •• O pairs were supposed to exist in BaTiO 3 [10][11][12][13].
At present, two scientific problems remain unsettled and need further investigation: (1) direct evidence for observing V •• O in Mg-doped BaTiO 3 is still lacking; and (2) the possibility of occupying the A site for Mg 2+ has not been determined. Our previous study confirmed the amphoteric behavior of Dy 3+ in BaTiO 3 , which can occupy both A and B sites [14][15][16]. The ionic radius of 12-coordinate Dy 3+ at the A site is 1.19 Å [14], which is little smaller than Mg 2+ (1.23 Å) with the same coordinate number (CN). Ionic radii with different CN are given in Table 1 [17]. The amphoteric behavior of Dy 3+ and the similar ionic radii between Dy 3+ and Mg 2+ indicate that although the structures and properties of B-site Mg-doped BaTiO 3 have been studied, the possibility of Mg 2+ occupying the A site should not be excluded. Although MgTiO 3 has a distorted rhombohedral structure [18], which is completely different from the perovskite structure of BaTiO 3 , the difference in crystalline structure between MgTiO 3 and BaTiO 3 may not be a key factor for Mg 2+ to enter the A site. Thus, the amphoteric nature of Mg 2+ in BaTiO 3 is still a scientific problem to be clarified.
In this work, BaTiO 3 −MgTiO 3 (BMT) solid solutions were prepared at different sintering temperatures. At a lower sintering temperature (T s ) such as 1150 • C, Mg 2+ was considered to segregate to the surfaces of the grains and play an important role in the formation of the core-shell structure [6]. Therefore, a longer sintering time in this work was used to ensure the incorporation of Mg 2+ into the BaTiO 3 lattice. The site occupation and amphoteric behavior of Mg 2+ and the dependence of V •• O on T s were investigated. The electron paramagnetic resonance (EPR) technique was employed to detect the existence of V •• O in the low-temperature range.
Powder X-ray diffraction (XRD) data were collected using a DX-2700 X-ray diffractometer (Dandong Haoyuan, Dandong, China). The lattice parameters were calculated by MS Modeling (Accelrys, Inc., San Diego, CA, USA) using Rietveld refinement in the Reflex Package and Cu Kα1 radiation (λ = 1.540562 Å). Scanning electron microscope (SEM) investigations were performed using an EVOMA 10 SEM (Zeiss, Oberkochen, Germany) operated at 15 kV. The sample surfaces were first polished and then thermally etched at the same sintering temperatures for a few minutes before SEM measurement. The dielectric properties were investigated at 1 kHz, from −75 to 200 • C, at a heating rate of 2 • C/min using a Concept 41 dielectric/impedance spectrometer (Novocontrol) with an applied voltage of 1 V. Temperature-dependent electron paramagnetic resonance (EPR) measurements were performed using an A300-10/12 X-band spectrometer (Bruker, Rheinstetten, Germany) operating at 9.43 GHz. The EPR cavity of the spectrometer was replaced with an ER 4102ST cavity.

Results
Powder XRD patterns of nominal (Ba 1−x Mg x )TiO 3 (x = 0.015) (BM1T) ceramics prepared at T s = 1200-1400 • C are shown in Figure 1. BM1T sintered at T s = 1200 • C exhibited a pseudocubic perovskite structure (space group: Pm3m) marked by a symmetric and broad characteristic (200) peak at~45 • (Figure 1a, inset). As T s was increased to 1300 • C, this peak evolved into slight the tetragonal and hexagonal phases coexisted in BM1T. It was inferred from the main (110) peak at 31 • that the amount of the hexagonal phase was approximately 30% of the tetragonal phase for BM1T sintered at T s = 1400 • C.   [20] and the rhombohedral MgTiO 3 appeared in BMT when x ≥ 0.07. Thus, the solubility limit of Mg 2+ in BMT sintered at T s = 1200 • C was determined by XRD to be x = 0.05. The variation in unit cell volume (V 0 ) as a function of x for BMT is shown in the inset in Figure 3. In the monophasic region of x ≤ 0.05, V 0 decreased linearly with increasing x. In the multiphasic region of x > 0.05, V 0 increased. Temperature dependencies of the dielectric permittivity (ε') and dielectric loss (tan δ) for BM1T are shown in Figure 4. The ε'-T curve of BM1T sintered at T s = 1200 • C was smooth and even, showing a rounded hill at around T m = 110 • C. The Curie peak of BaTiO 3 was dramatically suppressed due to Mg doping, and this ceramic satisfied the X8S specification (|(ε'−ε' RT )/ε' RT | ≤ 22% in a temperature range from −55 to 125 • C) with ε' RT = 1200. BM1T exhibited a very low tan δ (0.0176) at room temperature and lower tan δ (<0.05) in a T range of −55 to 110 • C. Subsequently, tan δ increased with increasing T.
When T s = 1300 • C, the ε'-T curve of BMT exhibited a bimodal structure, corresponding to a tetragonal-cubic (t-c) transition at dielectric peak temperature T m = 96 • C and an orthorhombic-tetragonal (o-t) transition at T 2 = 12 • C.
As T s was increased to 1400 • C, the bimodal feature in the ε'-T curve became more distinct and t-c and o-t transitions occurred at T m = 106 and 14 • C, respectively. The ε' RT decreased and tan δ increased rapidly above T = 50 • C. Temperature-dependent EPR spectra for BM1T are shown in Figure 5. For BM1T sintered at T s = 1200 • C, only the g = 2.004 signal existed over the measuring temperature (T) range of −188 to 150 • C (Figure 5a). This signal was assigned as ionized Ti vacancies [21][22][23]. The g = 2.004 signal was activated in the cubic phase above T m and in the rhombohedral phase below T = −100 • C. This activation confirmed the nature of Ti vacancies [23]. The pair of weak lines denoted as g 1 = 1.944 and g 3 = 2.060 appeared at T = −188 • C, forming a centrosymmetric pattern around g 2 = 2.004. This phenomenon is similar to the low-temperature EPR spectrum observed for (Ba 0.85 Sr 0.15 )TiO 3 [24], which may relate to the occupation of Mg 2+ on the A site. When T s = 1300 • C, except for the g = 2.004 signal, two additional signals at g = 1.974 and 1.956 observed at T = −188 • C (Figure 5b) were assigned as ionized Ba [14,22] and oxygen ( [19] vacancies, respectively. BM1T sintered at T s = 1400 • C existed in mixed forms of the hexagonal and tetragonal phases. Five EPR signals appeared below T = −100 • C and their intensity increased with decreasing T (Figure 5c). The presence of three signals at g = 2.004, 1.974, and 1.957 implies the coexistence of V Ba , V Ti , and We attributed two additional signals at g = 1.934 and 1.942 to a hexagonally distorted d 1 ion from Ti 3+ (Ti Ti ) because low temperatures can effectively prolong the spin-lattice relaxation time (τ) [19,22]. This indicates that during high-temperature sintering of T s = 1400 • C, the electrons in BM1T can be trapped by Ti 4+ ions to cause a reduction from Ti 4+ to Ti 3+ . It has been reported that the (Ba 1−x Ca x )TiO 3 (x = 0.03) ceramic sintered at T s = 1500 • C showed a more ordered tetragonal structure, and only a Ti 3+ -related signal at g = 1.932 was observed at T = −188 • C [19,22]. However, this signal did not appear in the tetragonal BMT sintered at T s = 1300 • C (Figure 5b). In the mixed hexagonal and tetragonal phases of BMT sintered at T s = 1400 • C, the Ti 3+ -related signal split into two signals at g = 1.934 and 1.942. It is obvious that these two signals originated from the hexagonal phase in BM1T.

Site Occupation of Mg 2+ in BM1T at Different Sintering Temperatures
On the basis of a simple comparison of 12-CN ionic size between Ba 2+ (1.61 Å) and Mg 2+ (1.23 Å) and 6-CN ionic size between Ti 4+ (0.605 Å) and Mg 2+ (0.72 Å), a continuous decrease in V 0 with x (≤0.05) for BM1T sintered at T s = 1200 • C (Figure 3, inset) provides sufficient evidence for occupation of the A site by Mg 2+ . When x is higher than the solubility limit of 0.05, Mg 2+ cannot continuously enter the A site, accompanied by separation of Mg-rich phases (  [10][11][12][13]. BO 6 octahedrons are characteristic of the perovskite lattice. Hence, higher energy is required to incorporate doping ions into the B site. It is inferred that the sintering temperature of T s = 1200 • C is too low to incorporate Mg 2+ into the B site because the V •• O -related EPR signal was not observed (Figure 5a). On the other hand, BM1T has a pseudocubic structure and its V 0 (= 64.40 Å 3 ) is equal to the tetragonal BaTiO 3 (V 0 = 64.41 Å 3 , JCPDS Card No. 6-526). This implies that Mg 2+ tends to remain close to the surfaces of the grains and plays an important role in the temperature-stable X8S behavior in BM1T, as suggested by Chang et al. [5]. At this time, Mg 2+ exists only at the A site as Mg × Ba . El Ghadraoui et al. indicated that the solubility limit of Mg 2+ in (Ba 1−x Mg x )TiO 3 was 0.15. They neglected a small amount of the secondary phases of BaMg 6 Ti 6 O 19 and MgTiO 3 , which also appeared in their samples with x ≥ 0.05 [25]. Their report undoubtedly supports that Mg 2+ may exist at the A site.
When T s was increased to 1300  [26,27]. The Jahn-Teller distortion encased by the Mn Ti ions is proposed to be the driving force of the phase transition from tetragonal to hexagonal [28]. This implies that Mg Ti and Mn Ti acceptors on the Ti sites in BaTiO 3 play the same role in the formation of the hexagonal phase. Thus, Mg Ti  when T s is lower than 1200 • C and tan δ at T s = 1300 • C is astonishingly low over the T range of −55 to 150 • C (tan δ ≤ 0.03).
The increase in T s can create V •• O and Ti Ti , giving rise to phase splitting into hexagonal and tetragonal at T s = 1400 • C. The high value of tan δ is attributed to the high concentrations of V •• O and Ti Ti in the hexagonal phase in BM1T (Figure 4).

Conclusions
The solubility of Mg 2+ in (Ba 1−x Mg x )TiO 3 ceramics sintered at 1200 • C was 0.05, and rhombohedral MgTiO 3 and hexagonal BaMg 6 Ti 6 O 19 phases were observed with higher doping content. The evolution of unit cell volume provided sufficient evidence for the A-site occupation of Mg 2+ . The x = 0.015 ceramic had a pseudocubic crystal structure when the sintered temperature was 1200 • C and exhibited a temperature-stable X8S dielectric specification with ε RT = 1200. The structure transformed into a tetragonal phase when sintered at 1300 • C, and tetragonal and hexagonal phases coexisted when sintered at 1400 • C.