Highly conductive Nb doped BaSnO3 thin films on MgO substrates by pulsed laser deposition

Journal of Alloys and Compounds 680 (2016) 343e349

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Highly conductive Nb doped BaSnO3 thin films on MgO substrates by pulsed laser deposition Bing Li, Qinzhuang Liu*, Yongxing Zhang, Zhongliang Liu, Lei Geng Department of Physics and Electronic Information, Huaibei Normal University, Huaibei 235000, People’s Republic of China

a r t i c l e i n f o

a b s t r a c t

Article history: Received 21 January 2016 Received in revised form 30 March 2016 Accepted 16 April 2016 Available online 19 April 2016

Ba(NbxSn1-x)O3 (BNSO) films were epitaxially grown on (001) oriented MgO substrates with Nb doping content from 0.00 to 0.15 by pulsed laser deposition. The structural, electrical, and optical properties of the films were investigated in detail by x-ray diffraction, x-ray photoelectron spectroscopy, Hall effect and transmittance spectroscopy. The lowest room temperature resistivity value of 4.81
104 U cm with carrier concentration and mobility of 6.59
1020/cm3 and 19.65 cm2/V was observed in the film at x ¼ 0.05, indicating an excellent electrical conductivity. All the BNSO films exhibit a high transmittance of more than 80% in the visible range. The variation of band gap Eg with Nb doping content was interpreted by the Burstein-Moss effect and the occurrence of octahedral tilt distorting. Excellent optical and electrical properties suggest that BNSO films have potential applications in the optoelectronic devices as transparent and conducting oxide material. © 2016 Elsevier B.V. All rights reserved.

Keywords: Transparent conducting oxides Transport property Optical property Alkaline earth stannate Pulsed laser deposition

1. Introduction Transparent conducting oxides (TCOs) films are unique materials combining the seemingly conflicting properties of high electrical conductivity in a wide temperature range and optical transparency at visible wavelengths within a single material. They have many applications in optoelectronic devices such as flat panel displays, organic light-emitting diodes, solar cells, and transparent electronics [1e6]. As one of the TCOs, Sn doped In2O3 (ITO) is widely used because of its excellent optical transparency and electrical transport property, which exhibits a high transmittance of more than 90% and very low resistivity about 104 U cm [7]. However, limited by the rarity and high cost of indium with the increasing demand for high-performance TCOs, many efforts have been made to develop new TCOs materials, such as Al, Sb, Ta and Nb doped SnO2 [8e11], Al and Ga doped ZnO [12,13], Nb doped TiO2 [14], etc. Besides these binary oxides, perovskite-type ternary oxides have been employed to exploit new type TCOs and attracted much attention. For example, La doped SrSnO3 [15], Sb and Nd doped SrSnO3 [16], Sb doped ZnSnO3 [17], La and Sb doped SrTiO3 [18,19], Sr doped LaCrO3 [20], and Nb doped CaTiO3 [21] have been reported. Compared with the conventional TCOs, they have the

* Corresponding author. E-mail address: [email protected] (Q. Liu). http://dx.doi.org/10.1016/j.jallcom.2016.04.157 0925-8388/© 2016 Elsevier B.V. All rights reserved.

advantage of a simple and flexible structure that benefits the ionic substitution on both the A and B sites, which enables the formation of a large set of important materials for a wide variety of academic and industrial research [22,23]. Significantly, the structural compatibility and composition range of perovskite oxides allow one to combine thin layers of very different materials to design all perovskite-type heteroepitaxial structures to improve performances and further explore functions [24,25]. As we all know, undoped BaSnO3 has an ideal cubic perovskite structure and corresponds to a transparent wide band gap n-type semiconductor with the band gap of 3.4 eV [26,27], which is separated by the top of the valence band composed of the O 2p orbitals and the bottom of the conduction band dominated by the contributions of nonbonding Sn 5s orbitals [28]. Previously, some research groups achieved the conducting property in La and Sb doped BaSnO3 ceramic [29,30]. Afterwards, Woodward and Mryasov calculated its electronic band structure and pointed out that BaSnO3 could be a good transparent conducting oxide [31,32]. Recently, BaSnO3 films doped with some elements have gained considerable attention owing to their excellent optical and transport properties. Our group first prepared La and Sb doped BaSnO3 films on MgO substrates by pulsed laser deposition with room temperature resistivity 3.94 U cm and the transmittance of more than 90% [33,34]. Then, we prepared (Ba, Gd)SnO3 films on MgO substrates with the lowest resistivity of 6.18 mU cm and the Hall mobility 11.35 cm2/V [35]. Anoop et al. investigated the correlation

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between the structural, electrical, and luminescence behaviors of vacuum-annealed (Ba, La)SnO3 epitaxial films on SrTiO3 substrates and indicated that the electrical properties of epitaxial films could be controlled by the Sn2þ defects generated with oxygen vacancies during the vacuum-annealing of the films [36]. Kim et al. prepared high quality (Ba, La)SnO3 films on SrTiO3 substrates and realized a high electrical mobility (70 cm2/V) at room temperature [37]. Mun et al. confirmed that the electron mobility in BaSnO3 films was reduced by almost 7 times when the dopant was changed from A site La to B site Sb despite little change in the effective mass of the carriers, which indicates that the scattering rate of conduction electrons in the BaSnO3 system is strongly affected by the site at which the dopants are located [38]. In order to further investigate the influence of B site dopants species, we choose pentavalent Nb to substitute for tetravalent Sn. In fact, Nb doping can generate carriers with high density, which has been confirmed both by the firstprinciple calculations and experimental studies [14,39,40]. Considering the similar ionic radius of Nb5þ (0.64 Å) and Sn4þ (0.69 Å), a high electrical conductivity might be obtained in Nb doped BaSnO3 thin film. In this paper, we prepared Nb doped BaSnO3 (BNSO) epitaxial films by pulsed laser deposition with Nb doping content x ¼ 0.00, 0.02, 0.05, 0.07, 0.10 and 0.15. In order to avoid affecting the transmittance measurement of the films, the MgO single crystal substrates with a wide band gap of 7.8 eV were used to deposit BNSO films. The structural, electrical and optical properties of the BNSO films as a function of Nb doping content were studied in detail. The optimal Nb doping content was observed in 5% Nb doped BaSnO3 films with the room temperature resistivity, carrier density, and Hall mobility of 4.81
104 U cm, 6.59
1020 cm3, and 19.65 cm2/V, respectively. 2. Experimental The Ba(NbxSn1-x)O3 ceramic targets with x ¼ 0.00, 0.02, 0.05, 0.07, 0.10, 0.15 were prepared by conventional high-temperature solid reactions using BaCO3 (99.5%), SnO2 (99.5%), and Nb2O5 (99.99%). Stoichiometric mixtures of the ingredients were grinded in an agate mortar for 2 h and presintered in the air at 1300  C for 12 h. Then the products were grinded again and sintered at 1400  C for 12 h. After grinding, the powders were pressed to disk with a 30 MPa pressure and then sintered at 1460  C for 24 h. The BNSO thin films were prepared on (001) oriented MgO substrates by pulsed laser deposition (PLD) technique using a 248 nm KrF excimer laser. During deposition, the substrates were kept at 800  C with a flowing oxygen pressure of 30 Pa. The laser irradiating repetition rate was 3 Hz with a 360 mJ energy per pulse, and targetsubstrate distance was set at 5.5 cm. After a 20 min deposition time, the films were in situ annealed for 20 min and then cooled down to room temperature in the same oxygen ambient. The nominal composition of the films was assumed to be the same as that of the targets. The film thicknesses were determined to be about 300 nm by the film deposition rate and time. The structures of the films were studied by x-ray diffraction (XRD, PANalytical Empyrean/X’pert MRD, Resolution 0.0001 ) using Cu Ka radiation, including qe2q line scan, rocking curves(RCs), and reciprocal space mapping (RSM). The composition and binding energy of the films were characterized by x-ray photoelectron spectroscopy (XPS, VG Scientific, ESCALAB-250, Energy resolution 0.45 eV/(Ag 3d5/2)) with Al-anode x-ray source (Al Ka hn ¼ 1486.6 eV). The binding energies were calibrated relative to the C1s peak (284.8 eV) from hydrocarbons absorbed on the surface of the films. The temperature dependent resistivity curves of the films were tested using the conventional four-probe method. The room temperature carrier concentration, carrier mobility and

electrical resistivity were determined using van der Pauw geometry using a Hall effect measurement (HMS-3000/0.56T). The transmittances of the films were measured via a UVevis spectrometer (Lambda 950, Perkin-Elmer, USA, Wavelength repeatability UV/VIS: ±0.08 nm) with a wavelength range of 200e2000 nm. 3. Results and discussion Fig. 1(a) shows the XRD qe2q line scan data for BNSO films with various Nb content grown on (001) oriented MgO substrates by PLD technique. Only well-defined (00l) diffraction peaks appeared and no other phases or randomly oriented grains appeared in the scans, indicating that all the films exhibit preferred orientation along the c-axis. It was calculated that the out-of -plane lattice constants of all the films have nearly the same value of 4.107 Å. There are no obvious changes with the Nb doping content, which can be ascribed to the similar ionic radius of Nb5þ (0.64 Å) and Sn4þ (0.69 Å). This value is slightly smaller than the 4.118 Å of the bulk BNSO. The u scan rocking curve of the BNSO (002) peak is also shown in the inset to Fig. 1(a). The full-width at half-maximum (FWHM) value for the 0.05 doped film is 0.46 .The peak broadening of this film may be related to defects such as grain boundaries and dislocations. In order to check the epitaxial properties, the F scans were measured on BNSO (202) and MgO (202) reflections by setting the reflection plane and substrate surface plane at J ¼ 45 as shown in Fig. 1(b). A set of four distinct peaks with 90 of separation indicates the films were epitaxially grown on MgO(001) substrates with a cube-on-cube relationship. In order to check the strain state in the films, asymmetrical reciprocal space mappings (RSM) were measured around the (113) reflections. Fig. 1(c) shows the asymmetrical RSM of BNSO film with Nb doping content at x ¼ 0.05, in which it can be seen that only reflections spots from BNSO(113) and MgO(113) were observed and the BNSO films are fully strainrelaxed [41]. Considering the large lattice mismatch (2.3%) between the film and the substrate, the strain might be relaxed by the generation of misfit dislocations [42]. The broader (113) reflection is consistent with the XRD rocking curve result. Fig. 2(a)e(d) displays the XPS spectra of BNSO films with x ¼ 0.00, 0.05, 0.10 covering Ba 3d, Nb 3d, Sn 3d and O 1s core-level emissions. Fig. 2(a) shows the binding energies of Ba 3d5/2 and Ba 3d3/2, located at 779.6 eV and 794.9 eV respectively, indicating the Ba2þ state in the BNSO films [43]. Fig. 2(b) shows the Nb 3d spectrum for the doped film exhibiting two peaks at 207.4 eV and 209.9 eV. The peaks represent the 3d5/2 and 3d3/2 components respectively, with a spin-orbit splitting of 2.5 eV, and correspond to that of Nb5þ oxidation state [44]. In other word, we do not find the characteristic of tetravalent Nb4þ, because its binding energy is lower than that of Nb5þ about 2 eV. The similar phenomenon was also observed in ceramic Nb doped BaSnO3 [45]. Therefore, Nb acts as a donor impurity and provides electron carriers in the BNSO film. It is also seen that the intensity of the Nb5þ peaks increases with the increasing of Nb content in the film, which further confirmed that Nb substitution occurred in the doping range. The Sn 3d photoemission peaks are shown in Fig. 2(c). It can be seen that the binding energy of Sn 3d3/2 and 3d5/2 is located at 495.2 eV and 486.6 eV respectively, corresponding to Sn4þ [46]. When scrutinizing the Sn 3d XPS spectra, we find that the binding energy shifts very slightly towards lower energy with increasing Nb doping content, indicating the partial change of Sn valence state from Sn4þ to Sn2þ. Considering the change is very slight and there is no obvious asymmetry of the peaks resulting from Sn2þ, Sn exists mainly in the valency of Sn4þ. This is different from the La doped BaSnO3 system, in which Sn doping results in obvious charge disproportionation of Sn4þ and Sn2þ [47]. The charge disproportionation blocks the increase of carrier density and influences the

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Fig. 1. (a) XRD linear scans for BNSO/MgO films at different Nb concentration (from x ¼ 0 to 0.15). The inset shows the u scan for BNSO (002) reflection. (b) The F scans patterns of the BNSO/MgO around (202) reflections. (c) The reciprocal space mapping around the (113) asymmetric Bragg diffraction point for BNSO/MgO film at x ¼ 0.05.

Fig. 2. The XPS spectra of (a) Ba 3d, (b) Nb 3d, (c) Sn 3d, and (d) O 1s from BNSO films.

conductivity. In fact, the existence of Sn2þ (1.12 Å) would cause the

lattice constant increase with the increasing Nb doping content,

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which is not observed in the XRD data. Fig. 2(d) presents the measured O1s signal. The spectrum is consisted of an asymmetric peak dominated by one major peak positioned at 530.1 eV and another peak at 532 ± 0.3 eV, which is related to the lattice oxygen and absorbed oxygen respectively [48]. The electrical resistivity (r), carrier concentration (n) and mobility (m) (see Table 1) of the BNSO films as a function of Nb doping content are shown in Fig. 3. Hall effect measurement suggests that the majority carriers in the films are electrons within the doping range. The resistivity decreases firstly and reaches the lowest value of 4.81
104 U cm at x ¼ 0.05, which is very close to the widely used ITO [7]. Then the resistivity increases monotonously with the Nb content. Considering the relationship of r ¼ 1/ nem, the changes of resistivity are related to the variation of the carrier concentration and mobility. It can be seen that the carrier concentration is 2.64
1020/cm3 and 6.59
1020/cm3 at x ¼ 0.02, 0.05, respectively, then reaches the highest value of 8.01
1020/ cm3 at x ¼ 0.07. The initial carrier concentration increase with the Nb doping content increasing can be explained by the increasing substitutional incorporation of Nb5þ at Sn4þ sites and forming more free carriers in the films. If each Nb5þ dopant produces one free electron, the expected carrier concentration in the BNSO film for x ¼ 0.02, 0.05 would be 2.86
1020/cm3 and 7.51
1020/cm3. It means that the ionization efficiency of Nb can reach above 90%. However, as the Nb doping content more than 0.07, the carrier concentration decreases gradually. At x ¼ 0.15, it decreases nearly by three orders of magnitude to the lowest value of 1.53
1017/cm3. This is explained by the fact that, as the Nb doping content further increasing, lots of electrically inactive Nb that cannot generate free carriers will appear in the form of interstitial atoms or dopant cluster. It will cause the disorders and trap the electrons, and then lead to the decrease in carrier concentration [49,10]. Furthermore, these disorders and trapped electrons will become an additional scattering center to block the electron transport and cause a decrease of mobility [49,10], as shown in Fig. 3. Also, the increasing small angle grain boundaries and dislocation in the epitaxial films play an important role for the decrease of mobility [37]. In fact, the mobility decreases dramatically after x ¼ 0.05, indicating the formation of scattering center. The initial increase of mobility may be due to the reduced barriers caused by the higher carrier concentration when the Nb element doping content is less than x ¼ 0.05 [50]. Moreover, we should pay attention to the high mobility m of 19.65 cm2/V, which is larger than that of 11.35 cm2/V in the Gd doped BaSnO3 films on MgO substrates [35]. Considering the same substrates, the main reason resulting the difference of m values can be ascribed to the doped element species. The Nb5þ has a similar radius with Sn4þ and reduces the degree of lattice distortion because of doping, which makes the carriers suffer a smaller scattering than that of Gd doped BaSnO3. The higher mobility and carrier concentration make the BNSO film come into the optimal doping at x ¼ 0.05. Fig. 4(a)e(e) show the effect of Nb doping on the temperature dependent resistivity of the BNSO films with Nb contents from 0.02 to 0.15. It can be seen that the resistivity of the films is tuned by the Nb doping level. The BNSO films with x ¼ 0.02, 0.05, 0.07 present a

Fig. 3. Room temperature resistivity (r), carrier concentration (n), and mobility (m) of the BNSO films as a function of Nb doping content.

metal-semiconductor transition (MST) at 19 K, 40 K and 69 K with metallic conductivity at higher temperatures and semiconducting behavior at lower temperatures. The similar MST behavior was also observed in our La and Sb doped BaSnO3 [33,34], as well as many other heavily-doped films [10,13,14,51]. These MST behaviors at lower temperatures have also been observed in the spin disordered system and called Kondo-like scattering [52], which is impossible for the BNSO film because of its nonmagnetic nature. In our case, the MST implies that at least two conducting mechanisms coexist in the films. The semiconducting behavior observed below the transition temperature (TC) may be due to the localization of electrons by disorder at low temperature [53]. The conductivity behavior above TC can be attributed to the formation of a degenerate band in the heavily doped semiconductor. As a large concentration of carriers is introduced into the system, the electrons start to occupy the bottom of the conduction band and consequently the Fermi level lies in the conduction band, which induces the metallic behavior of the films [54]. The room temperature resistance of 0.05 doped BNSO film measured by four-probe method is only 10.4 U, showing a high electrical conductivity. With Nb doping content increasing to x ¼ 0.10 and 0.15, the resistivity value of the BNSO films increases and the semiconducting behavior dominates in the measurement temperature range. In order to well understand the conducting mechanism, a detailed analysis of temperature dependent resistivity was carried out. For the BNSO films with x ¼ 0.10, a well linear relations can be obtained by r and lnT, as shown in the inset of Fig. 4(d), suggesting that conducting mechanism obeys two dimensional weak localization model [55]. However, for BNSO films with x ¼ 0.15, the well linear relation of lnr versus T1/4 can be obtained to describe the film resistivity, indicating that the electrical conductivity was predominated by variable range hopping [56]. It can be explained that, with the Nb doping content increasing, many scattering centers such as clusters or defects will form and enhance the scattering potential, resulting in Anderson

Table 1 Summary of the carrier concentration, mobility, the resistivity and hall coefficient at 300 K by Hall effect measurement. Doping content x x x x x

¼ ¼ ¼ ¼ ¼

0.02 0.05 0.07 0.10 0.15

Carrier concentration (/cm3) 2.64 6.59 8.01 1.95 1.53








1020 1020 1020 1020 1017

Mobility (cm2/Vs)

Resistivity (mU cm)

Hall coefficient cm3/C

10.43 19.65 5.01 2.83 2.59

2.27 0.48 1.55 11.33 15770

2.37 9.48 7.79 3.20 4.09








102 103 103 102 101

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347

Fig. 4. Temperature dependent resistivity of the BNSO films with Nb doping content at (a) x ¼ 0.02, (b) x ¼ 0.05, (c) x ¼ 0.07, (d) x ¼ 0.10 and (e) x ¼ 0.15. Inset to figure (d) shows the linear fit of r vs lnT, and inset to figure (e) is the linear fit of lnr vs T1/4.

localization of electronic states in the system. The change of electric conduction mechanism from weak localization to strong localization regime with increasing Nb concentration can be further substantiated by the increased resistivity. The resistivity value of x ¼ 0.15 is larger than that of x ¼ 0.10 by three orders of magnitude at room temperature. Fig. 5(a) shows the optical transmission of the BNSO/MgO thin films as a function of Nb doped comtent in the wavelength region 200e2000 nm, and that of pure MgO substrate was also given for comparison. A high transmittance of more than 80% was achieved in the visible region for all the films. The inset to Fig. 4(a) shows the BNSO/MgO film with x ¼ 0.05, illustrating a high transparency. However, in the near-infrared region, the transparency decreases greatly at first and then increases slightly again with increasing Nb doping content. An important optical characteristic for TCOs is that they have a transparency window. In the near-infrared region, reflection occurs and light cannot be transmitted because of the plasma edge. The plasma frequency is given by the relation [57],




.

1=2

up ¼ ne2 ε0 ε∞ m *

(1)

where n is the carrier concentration, e is the electronic charge, ε0 is the permittivity of free space, ε∞ is the high-frequency permittivity and m* is the conductivity effective mass. It is found from the

relation that the plasma frequency is highly dependent on the carrier concentration and its shift, which can explain the phenomena of our BNSO films in the near-infrared region. This is consistent with the change of carrier concentration as shown in Fig. 3. In the near-ultraviolet region, the light is absorbed due to the fundamental band gap and the absorption-edge shifts towards higher energy. The optical band gaps Eg of the BNSO films were determined from the absorption spectra using the equation [58],

  ðhnaÞ2 ¼ A hn  Eg

(2)

where a is the absorption coefficients, A is a constant corresponding to the electronehole mobility, h is the Planck constant, and n is photon frequency. Here, the absorption coefficients were calculated using the well-known relation, a¼(1/d)ln[(1-R)/T] [59], where d is the film thickness, T is the transmittance, and R is the reflectance (we ignored it in the calculation). The (hna)2 versus hn plot was shown in Fig. 5(b), and Eg could be obtained by extrapolating the linear portion of (hna)2 to abcissa axis. As shown in the inset to Fig. 5(b), the band gap energy Eg increases from 3.48 eV to 4.11 eV with the Nb doping content. The error bars means the uncertainty during extrapolating the linear portion of (hna)2 to abcissa axis. The optical band gap Eg 3.48 eV of the undoped BaSnO3 films equals to

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BaSnO3 system. Hall effect measurement suggests that the majority carriers in the films are electrons within the doping range. With the increasing Nb doping content, the resistivity decreases firstly and reaches the lowest value of 4.81
104 U cm at x ¼ 0.05, which is very close to the widely used ITO. The excellent electrical conductivity can be ascribed to the higher mobility and carrier concentration. The transport properties of BNSO films with x ¼ 0.02, 0.05, 0.07 are dominated mainly by the metallic conductivity at a wide temperature region, which can be attributed to the formation of a degenerate band in the heavily doped semiconductor. A high transmittance of more than 80% was achieved in the visible region for all the films. The band gap variation with the Nb doping content was interpreted by the Burstein-Moss effect and the occurrence of octahedral tilt distorting. Acknowledgments This work was supported by the National Natural Science Foundation of China (Grant Nos. 11504120, 11004071, 51302102, 51402120), Natural Science Foundation of Anhui Province (Grant Nos. 1608085QE90, 1408085QA19), Natural Science Foundation of Anhui Higher Education Institutions of China (Grant No. KJ2015A095), Talent Support Program of Anhui Province (Grant No. gxyqZD2016110), and Collaborative Innovation Center of Advanced Functional Materials (Grant No.XTZX103732015012). References

Fig. 5. (a) Optical transmission of BNSO/MgO heterostructures at different Nb doping content, and pure MgO substrate was given for comparison. (b) (hna)2 vs hn plots of BNSO/MgO films, and the inset shows the band gaps of BNSO films.

the energy separation between the conduction and valence band edges. The value is slight larger than that (3.4 eV) of undoped BaSnO3 reported by Refs. [26,27] and can be attributed to the oxygen deficiency during deposition [60]. The increase of Eg with Nb doping content from 0.02 to 0.07 can be explained by the BursteinMoss effect [61,62]. With the increasing Nb doping content, more carriers generate in the film and fill the bottom of the conduction band. As a result, larger energy will need to excite the electrons to the high energy level of conduction band. However, Eg still increases at x ¼ 0.10 and 0.15 even though the carrier density was reduced. This could be ascribed to the occurrence of the octahedral tilt distorting [63], as the distortion of the linear SneOeSn bonds narrows the conduction band. 4. Conclusion BNSO films have been grown on (001) oriented MgO substrates with Nb doping content from 0.00 to 0.15 by pulsed laser deposition. The XRD data reveal that all the films are epitaxially grown on the substrates with a cube-on-cube relationship. The RSM around the (113) reflections indicates that the films are almost strain relaxed. The XPS spectra show the presence of only Nb5þ irons and give no evidence of charge disproportionation for Sn4þ to Sn2þ and Sn4þ, illustrating that Nb is an excellent donor element in the

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