For La-doped BaSnO3 thin films grown by pulsed laser deposition, we combine chemical surface characterization and electronic transport studies to probe the evolution of electronic states in the band structure for different La-doping contents. Systematic analyses of spectroscopic data based on fitting the core electron line shapes help to unravel the composition of the surface as well as the dynamics associated with increasing doping. These dynamics are observed with a more pronounced signature in the Sn 3d core level, which exhibits an increasing asymmetry to the high binding energy side of the peak with increasing electron density. The present results expand the current understanding of the interplay between the doping concentration, electronic band structure, and transport properties of epitaxial La:BaSnO3 films.

The perovskite La-doped BaSnO3 (La:BaSnO3) is a novel transparent oxide semiconductor that exhibits outstanding room temperature (RT) electron mobility (μe) with high carrier density together with high optical transmittance.1–3 Owing to its unique electronic and optical properties, La:BaSnO3 has the potential for applications in transparent electronics,4–7 photovoltaics,8–11 as well as in thermoelectric12–15 and multifunctional perovskite-based optoelectronic devices.10,16,17 Furthermore, its low-power consumption combined with its ability to be heavily doped and its good stability at high temperatures make La:BaSnO3 a suitable material for integration in thermally stable capacitors, field effect transistors, and power electronic devices.3,4,17–19

The discovery of a RT μe of 320 cm2 V−1 s−1 (with corresponding carrier density, n = 8 × 1019 cm−3) in La:BaSnO3 single crystals1–3 stimulated intense investigation into this material.4 Particularly, the potential of La:BaSnO3 for device applications and heterostructures triggered considerable interest in thin films grown from this compound.5–7,15,17,19–38 However, the reported μe in La:BaSnO3 thin films have only reached a maximum value of 183 cm2 V−1 s−1 (n ≃ 1.2 × 1020 cm−3) for epitaxial films grown by molecular beam epitaxy (MBE).33 Other growth techniques resulted in the following electron mobilities: 140 cm2 V−1 s−1 (n ≃ 5.2 × 1020 cm−3) for pulsed laser deposition (PLD),27 121 cm2 V−1 s−1 (n ≃ 4.0 × 1020 cm−3) for high-pressure magnetron sputtering,36 and 53 cm2 V−1 s−1 (n ≃ 2.0 × 1020 cm−3) for chemical solution deposition.39 Various strategies to improve the mobility in La:BaSnO3 epitaxial films have been explored. Such efforts include, for example, the incorporation of undoped BaSnO3 buffer layers to compensate for the lattice mismatch between the substrate and the active La:BaSnO3 top layers,7,19,26,33 adsorption-controlled MBE for improved stoichiometry control,31–33,40 a very high-temperature grown insulating buffer layer to reduce the density of threading dislocations,27 and post growth annealing processes.22,24,41 Besides the ongoing efforts for RT μe improvement, in order to gain a better understanding of the conduction mechanisms in La:BaSnO3 films, it is important to establish a proper correlation between the transport characteristics and the behavior of the electronic states in the conduction band. This is crucial because the high ambient μe in La:BaSnO3 has been proposed to originate from both the small effective mass of the electrons at the conduction band minimum (CBM),25,42 which is associated with the largely dispersive Sn 5s conduction band and the low optical phonon scattering rate.19,43

Although several studies used photoemission spectroscopy techniques to investigate the electronic structure of La:BaSnO3 films,32,43–46 only a few reports have combined electronic transport and spectroscopic studies to explore the evolution of electronic states in La:BaSnO3 films and heterostructures at different La-doping levels.32,43 In particular, recent ex-situ hard x-ray photoemission spectroscopy (HAXPES) experiments on La:BaSnO3 films demonstrated that both the CBM and the valence band maximum (VBM), as well as the core electrons, are effectively modified with increasing carrier density.32 Thus, this result calls for additional combined spectroscopic and electrical characterizations to facilitate a more quantitative exploration of the evolution of the intrinsic properties of La:BaSnO3 films and heterostructures at different doping levels.

In this paper, we combine chemical surface analysis as a function of La doping using x-ray photoelectron spectroscopy (XPS) and electronic transport studies to explore the evolution of the electronic states in La:BaSnO3 films and heterostructures. From the transport measurements, we extract the transport characteristics, as well as n and μe, of La:BaSnO3 samples. The surface properties of these samples are subsequently investigated using spectroscopic techniques. A direct connection between the electronic transport characteristics and the spectroscopic data is demonstrated. We used XPS as a probing tool to measure the changes in the film spectra associated with the increasing amount of La3+ dopant. Through the analysis and systematic fits of the core XPS spectra, we are able to extract the binding energy values of the constituent elements along with the associated oxidation states. These data are consistent with the electron energy loss spectroscopy (EELS) data as well as with the literature. By increasing the doping concentration, we observe shifts of the valence band leading edges toward higher binding energies as well as increases in the states in the conduction band. More importantly, we provide a quantitative understanding of the effect of conduction band filling in La:BaSnO3 films and heterostructures. This effect is manifested by an increasing asymmetry in the line shape of the Sn 3d core spectra and leads to considering an additional plasmon satellite peak in the analysis of Sn 3d spectra.

Epitaxial La:BaSnO3 films and heterostructures (samples labeled A to E in Table I) were prepared by PLD (λ = 248 nm). Prior to deposition, the (100) oriented SrTiO3, (110) oriented DyScO3. and TbScO3 crystalline substrates (5 × 5 × 1 mm3) were terminated in situ at 1300 °C using a CO2 laser substrate heating system.47Figure 1(a) depicts a schematic view of the sample types investigated. The films were grown from La:BaSnO3 targets of 2%, 4%, and 6% La doping contents. Details about the growth and systematic characterization of the films are provided in Ref. 27 (see details on the PLD growth conditions in Table S1 of the supplementary material.

TABLE I.

Electronic transport characteristics (carrier density and mobility) of the samples discussed in this study.

Sample nameSample layoutCarrier density (×1020 cm−3)Carrier mobility (cm2 V−1 s−1)
6% La:BaSnO3 (25 nm)/BaSnO3 (100 nm)/DyScO3 1.24 ± 0.02 71 ± 2 
6% La:BaSnO3 (25 nm)/BaSnO3 (100 nm)/TbScO3 4.92 ± 0.05 20 ± 1 
6% La:BaSnO3 (25 nm)/SrTiO3 4.05 ± 0.05 91 ± 2 
2% La:BaSnO3 (25 nm)/TbScO3 1.35 ± 0.02 75 ± 2 
4% La:BaSnO3 (100 nm)/BaSnO3 (25 nm)/SrTiO3 0.80 ± 0.05 18 ± 1 
Sample nameSample layoutCarrier density (×1020 cm−3)Carrier mobility (cm2 V−1 s−1)
6% La:BaSnO3 (25 nm)/BaSnO3 (100 nm)/DyScO3 1.24 ± 0.02 71 ± 2 
6% La:BaSnO3 (25 nm)/BaSnO3 (100 nm)/TbScO3 4.92 ± 0.05 20 ± 1 
6% La:BaSnO3 (25 nm)/SrTiO3 4.05 ± 0.05 91 ± 2 
2% La:BaSnO3 (25 nm)/TbScO3 1.35 ± 0.02 75 ± 2 
4% La:BaSnO3 (100 nm)/BaSnO3 (25 nm)/SrTiO3 0.80 ± 0.05 18 ± 1 
FIG. 1.

(a) Schematic layout of the different thin film samples investigated in this study. (b) High-resolution scanning transmission electron microscopy (HRSTEM) image of a representative La:BaSnO3/BaSnO3 heterostructure (sample A). Misfit dislocations indicated by the ⊥ symbols characterize the relaxed interface between the film and the (110) oriented DyScO3 substrate. The inset represents a high magnification around one of the misfits delimited by the white rectangle, where a lack of closure of the Burgers circuit is seen (orange rectangular-like contour). (c) Electron energy-loss spectroscopy (EELS) elemental mapping for sample A showing the distribution of the elements in the sample. The signal is divided into regions as indicated by the dashed orange lines. The residual color in the La plot is noise. (d) A representative low energy electron diffraction (LEED) image taken at 48 eV displaying a clean La:BaSnO3 (001) surface. The diffraction spots form square lattices (white rectangles drawn on the image) in reciprocal space.

FIG. 1.

(a) Schematic layout of the different thin film samples investigated in this study. (b) High-resolution scanning transmission electron microscopy (HRSTEM) image of a representative La:BaSnO3/BaSnO3 heterostructure (sample A). Misfit dislocations indicated by the ⊥ symbols characterize the relaxed interface between the film and the (110) oriented DyScO3 substrate. The inset represents a high magnification around one of the misfits delimited by the white rectangle, where a lack of closure of the Burgers circuit is seen (orange rectangular-like contour). (c) Electron energy-loss spectroscopy (EELS) elemental mapping for sample A showing the distribution of the elements in the sample. The signal is divided into regions as indicated by the dashed orange lines. The residual color in the La plot is noise. (d) A representative low energy electron diffraction (LEED) image taken at 48 eV displaying a clean La:BaSnO3 (001) surface. The diffraction spots form square lattices (white rectangles drawn on the image) in reciprocal space.

Close modal

Electrical transport properties were measured in a physical property measurement system (PPMS) in a van der Pauw geometry obtained by wire bonding aluminum wires to the samples’ corners [see Fig. S1(a) of the supplementary material]. The carrier concentration, n, and the electron mobility, μe, were determined following the procedure discussed elsewhere.27,48–50Table I gives the carrier density and electron mobility of the different sample types investigated in this study. The La:BaSnO3 samples analyzed here in both spectroscopy and transport experiments have a RT carrier concentration ranging from n = 0.80 × 1020 to 4.92 × 1020 cm−3; and carrier mobility from μe = 18 to 91 cm2 V−1 s1. The measured carrier densities reported are smaller than the expected values associated with the nominal concentration of the La dopant. In fact, the doping levels corresponding to 2%, 4%, and 6% La are 2.72 × 1020, 5.44 × 1020, and 8.16 × 1020 cm−3, respectively. The observed discrepancy can be ascribed to the high density of dislocation defects present in the films, as seen in the weak beam dark field scanning transmission electron microscopy micrographs [Figs. S1(b)–S1(d)]. These defects arise because the films are grown on substrates to which they are poorly lattice matched, and this often results in the reduction of the carrier density and electron mobility in epitaxial La:BaSnO3 films2,27,31,33 (see more details in Sec. S1 of the supplementary material).

Following the procedure described in Ref. 51, the samples were systematically cleaned in ultra-high vacuum (UHV) before photoemission experiments. The surface structure of the samples was characterized using low-energy electron diffraction (LEED) [see Figs. S2(a)-S2(b) in the supplementary material]. The surface of clean La:BaSnO3 (001) surface showed a stable 1 × 1 surface structure [Fig. 1(d)]. The diffraction spots form square lattices in reciprocal space corresponding to the cubic lattice structure of BaSnO3 in real space, thus indicating the high crystallinity of the films.52 The cleanliness of the samples was checked by monitoring the LEED patterns directly after an annealing cycle and also by comparing the XPS survey scans before and after cleaning (see Fig. S2 in the supplementary material).

Figure 1(b) shows a representative scanning transmission electron microscopy (STEM) image of the investigated La:BaSnO3 heterostructure (sample A). To investigate stoichiometry and evaluate the composition of the epitaxial layers, EELS measurements were performed on the cross section of the samples during STEM characterization. Figure 1(c) displays a representative EELS elemental mapping (atomic layer distribution) of sample A for a total layer thickness of 120 nm, of which the La-doped layer is 25 nm thick. The black color suggests zero intensity, whereas the other colors indicate the presence of different elements (La, Ba, Sn, Dy, Sc, and O) that are resolved in the scanned region. The elemental composition of the different layers of the La:BaSnO3/BaSnO3 heterostructure is also evidenced in the EELS spectra for the same sample presented in Fig. 2. Looking at all the ionized edges, it can be seen that the spectra of the BaSnO3 (red curve) and La:BaSnO3 (green curve) are very similar, apart from the transfer of spectral weight to the La peaks in the latter. This spectral weight transfer is illustrated by the reduction of the intensity at the Ba-M edge. Intensity reduction is also noticeable at the O–K edge and could well be due to electronic (charge) modulation associated with the substitution of the bivalent Ba2+ atoms with the trivalent La3+ ones.53 Furthermore, the line shape of the spectra around the O–K feature (see inset Fig. 2) indicates good stoichiometry (no oxygen deficiency) in bulk layers, since these are in good agreement with the O–K EELS spectra reported for stoichiometric BaSnO3 and La:BaSnO3 films.33,54,55

FIG. 2.

Electron energy loss spectra of a representative La:BaSnO3/BaSnO3 heterostructure. Two regions in the specimen were investigated: an area with only BaSnO3 (red curve) and another with only La:BaSnO3 (green curve). The ionized edges of La-M4,5, Ba-M4,5, and Sn-M4,5 (corresponding to the excitation of 3d states), Sn-M2,3 (corresponding to the excitation of 3p states), and O–K (corresponding to the excitation of the 1s state) are resolved. The inset shows the enlargement of the O–K edge delimited by the blue rectangle.

FIG. 2.

Electron energy loss spectra of a representative La:BaSnO3/BaSnO3 heterostructure. Two regions in the specimen were investigated: an area with only BaSnO3 (red curve) and another with only La:BaSnO3 (green curve). The ionized edges of La-M4,5, Ba-M4,5, and Sn-M4,5 (corresponding to the excitation of 3d states), Sn-M2,3 (corresponding to the excitation of 3p states), and O–K (corresponding to the excitation of the 1s state) are resolved. The inset shows the enlargement of the O–K edge delimited by the blue rectangle.

Close modal

Table II presents the elemental composition of the surface of the films obtained from the fit of the XPS spectra. The composition is consistent throughout the surface of the various samples, with a larger proportion of Ba compared to Sn. Most importantly, the oxygen proportion in these high temperature annealed samples demonstrates the stability of the oxygen atoms in La:BaSnO31,3 and highlights the robustness of the sample cleaning procedure.

TABLE II.

Surface atomic percentages for the samples measured in XPS.

Sample nameBa (%)Sn (%)O (%)
20 16 64 
25 15 60 
22 16 62 
23 13 64 
Sample nameBa (%)Sn (%)O (%)
20 16 64 
25 15 60 
22 16 62 
23 13 64 

Figure 3 represents the XPS spectra of the Ba 3d core electrons together with the Voigt function fits. The Ba 3d spectra of all the samples showed an asymmetric line shape, suggesting the presence of multiple components at the core level. Two symmetric Voigt doublets (Ba I and Ba II) were used to fit the spectra. The main doublet, Ba I, located at the binding energy of 780.00 eV, is assigned to lattice barium in the Ba2+ state, consistent with previous spectroscopic results on powder and epitaxial thin films of BaSnO3.56,57 The second doublet, Ba II, situated at higher binding energy (781.13 eV), has been attributed to a surface character in several reports on epitaxial BaTiO3 films.58–60 This component was suggested to originate either from under-coordinated barium at a BaO terminated surface or from lattice relaxation.59,60

FIG. 3.

Angle- and energy-dependent XPS spectra of the Ba 3d core level for sample B. After subtracting a Shirley background (black dashed lineshape), the Ba 3d core lines are fitted (red lineshape) with two Voigt doublet peaks. The spectra in (a) and (b) were acquired with an Al Kα anode (Al Kα source) at 35° and 90° take-off angles, respectively. The spectrum in (c) was measured with an Ag Lα anode (Ag Lα source) at 90° take-off angle. Two satellite features (cyan and orange peaks) are resolved in the spectra. The peak heights were normalized for clarity.

FIG. 3.

Angle- and energy-dependent XPS spectra of the Ba 3d core level for sample B. After subtracting a Shirley background (black dashed lineshape), the Ba 3d core lines are fitted (red lineshape) with two Voigt doublet peaks. The spectra in (a) and (b) were acquired with an Al Kα anode (Al Kα source) at 35° and 90° take-off angles, respectively. The spectrum in (c) was measured with an Ag Lα anode (Ag Lα source) at 90° take-off angle. Two satellite features (cyan and orange peaks) are resolved in the spectra. The peak heights were normalized for clarity.

Close modal

In comparing the XPS spectra for cleaned surfaces with those of surfaces measured as-inserted (see Figs. S3 and S4 of the supplementary material), the relative intensity of Ba II was observed to decrease after the treatment of the surfaces, while that of Ba I increases. This trend suggests that Ba II could be a surface component, which is amplified with contamination. The nature of the Ba II peak was carefully investigated by carrying out a systematic analysis of its fraction with respect to the probing depth. This was achieved by performing angle-dependent XPS measurements as depicted in Fig. 3. The measurements were first performed using the Al Kα anode (photon energy of 1486.71 eV) at electron take-off angles of 35° [Fig. 3(a)] and 90° (normal emission) [Fig. 3(b)], and later, the excitation source was changed to Ag Lα anode (photon energy of 2984.31 eV) for acquisition at normal emission [Fig. 3(c)]. For the take-off angle of 35°, the photoelectrons emitted originate from a region nearer the surface, whereas at the take-off angle of 90°, the emitted photoelectrons are from a deeper depth within the sample. Hence, the measurement gets more bulk sensitive as the photoelectron take-off angle increases from 35° to 90° and as the excitation source is changed from Al to Ag. The parameters of the fits pertaining to the angle-dependent analysis are given in Table III. The ratio of the Ba II feature was observed to decrease considerably with bulk sensitivity measurements, thus confirming its surface character.

TABLE III.

Peak ratios for angle-dependent XPS of the Ba 3d spectra.

Peak assignmentPeak position (±0.05 eV)Relative intensity (±1%)FWHM (±0.05 eV)
35° Ba I 780.08 82 1.58 
Al anode Ba II 781.21 18 1.58 
90° Ba I 780.00 85 1.53 
Al anode Ba II 781.13 15 1.53 
90° Ba I 780.00 90 1.97 
Ag anode Ba II 781.13 10 1.97 
Peak assignmentPeak position (±0.05 eV)Relative intensity (±1%)FWHM (±0.05 eV)
35° Ba I 780.08 82 1.58 
Al anode Ba II 781.21 18 1.58 
90° Ba I 780.00 85 1.53 
Al anode Ba II 781.13 15 1.53 
90° Ba I 780.00 90 1.97 
Ag anode Ba II 781.13 10 1.97 

Two satellite features labeled Sat I and Sat II were also detected in the XPS spectra around the Ba 3d peaks (Fig. 3). These satellites result from shake-up processes involving Ba 3d photoelectrons and valence electrons.61,62 These broad satellites are located about 10 eV on the high binding energy side of the associated Ba 3d52 and Ba 3d32 peaks.

Figure 4 depicts the Sn 3d core level XPS spectra together with the Voigt function fits after subtraction of a Shirley background, for three representative samples of different total carrier densities n. The Sn 3d spectral lineshapes display an asymmetry to the high binding energy side of the peak, which increases with increasing carrier density. In metallic systems, the asymmetry in core photoemission spectra arises from intrinsic plasmon excitations associated with the creation of the core hole, which results in an additional component satellite to the main core line.63,64 It is known that the Coulomb potential of the core hole creates a localized trap state by capturing a conduction electron.65–67 In La:BaSnO3 systems, the conduction band is derived from highly dispersive Sn 5s bands,32,43 and the observed doping effect in the Sn 3d core level lineshape is most probably due to screening responses of the conduction electrons introduced by doping.63 Therefore, two doublet components were used to fit the Sn 3d spectra, assuming that the Koopmans’ state (i.e., the excited state after the removal of a core electron from the atom) is projected into screened and unscreened final eigenstates.63 

FIG. 4.

XPS spectra around the Sn 3d regions for three samples of different total carrier densities: (a) sample D with n3D=1.35×1020 cm−3, (b) sample C with n3D=4.05×1020 cm−3, and (c) sample B with n3D=4.92×1020 cm−3. After subtracting a Shirley background (black dashed lineshape), the Sn 3d spectra are fitted (red lineshape) with two Voigt doublet components. The data were acquired at normal emission with the Al Kα excitation source.

FIG. 4.

XPS spectra around the Sn 3d regions for three samples of different total carrier densities: (a) sample D with n3D=1.35×1020 cm−3, (b) sample C with n3D=4.05×1020 cm−3, and (c) sample B with n3D=4.92×1020 cm−3. After subtracting a Shirley background (black dashed lineshape), the Sn 3d spectra are fitted (red lineshape) with two Voigt doublet components. The data were acquired at normal emission with the Al Kα excitation source.

Close modal

To understand the effect of increasing carrier density in the Sn 3d core level, the spectral lineshape of four samples (B, C, D, and E) of different n values were investigated, and the analysis results are summarized in Table IV. The core lines were fitted to two Voigt components, which give an excellent description of the overall line shape of the spectra. In each spectrum, the main component is the peak at low binding energy labeled “screened.” This peak has a dominant Gaussian line shape. The component at high binding energy labeled “unscreened” is dominantly Lorentzian, which is a satellite associated with intrinsic plasmon excitations.67 This peak is broader than the screened component, as evidenced by their FWHM values. Similar satellite structures were reported in the Sn 3d and In 3d core photoemission spectra of binary transparent conducting oxides (Sb-doped SnO2,63,64,68 In2O3–ZnO,69 and Sn-doped In2O367,70). To better visualize the connection between carrier density and the screened/unscreened intensity and energy, the evolution of both the satellite energy and the intensity of the peaks with n−1/3 is plotted (Fig. 5). As can be seen from Table IV, the binding energy value of the main component suggests a valence state of 4+ for Sn,71,72 and the energy separation (satellite energy) between the main and satellite components increases with n [Fig. 5(a)]. Furthermore, the relative intensity of the screened component increases with increasing n, while that of the unscreened peak decreases. This is indicated in Fig. 5(b) by the increase of the intensity ratio of the peaks, conveying good agreement with previous reports.64,65,67,69,73 The discrepancy at the lowest n (sample E) could be ascribed to the fact that the thickness of this sample, 100 nm, surpasses a critical thickness above which additional structural defects are induced in the film.74,75 We speculate that defect scattering dominates, as only 15% of the carriers are activated compared to more than 50% activation for the other samples.

TABLE IV.

Fitted parameters of the Sn 3d regions along with the calculated carrier density in each sample.

Sample namePeak assignmentPeak position (±0.05 eV)FWHM (±0.05)Relative intensity (±1%)Satellite energy (eV)Carrier density (×1020 cm−3)
Screened 486.76 1.07 54 0.30 0.8 ± 0.05 
Unscreened 487.06 1.66 46 
Screened 486.51 1.15 52 0.52 1.35 ± 0.02 
Unscreened 487.03 1.34 48 
Screened 486.80 1.06 57 0.70 4.05 ± 0.05 
Unscreened 487.50 1.34 43 
Screened 486.75 1.12 60 0.80 4.92 ± 0.05 
Unscreened 487.55 1.40 40 
Sample namePeak assignmentPeak position (±0.05 eV)FWHM (±0.05)Relative intensity (±1%)Satellite energy (eV)Carrier density (×1020 cm−3)
Screened 486.76 1.07 54 0.30 0.8 ± 0.05 
Unscreened 487.06 1.66 46 
Screened 486.51 1.15 52 0.52 1.35 ± 0.02 
Unscreened 487.03 1.34 48 
Screened 486.80 1.06 57 0.70 4.05 ± 0.05 
Unscreened 487.50 1.34 43 
Screened 486.75 1.12 60 0.80 4.92 ± 0.05 
Unscreened 487.55 1.40 40 
FIG. 5.

Variation of the (a) satellite energy and (b) intensity ratio of the screened and unscreened components as a function of n−1/3. The error bars in the carrier density are too small to be seen in this plot. The dashed lines are guides to the eye. The black arrow indicates increasing n (plotted data are from Table IV).

FIG. 5.

Variation of the (a) satellite energy and (b) intensity ratio of the screened and unscreened components as a function of n−1/3. The error bars in the carrier density are too small to be seen in this plot. The dashed lines are guides to the eye. The black arrow indicates increasing n (plotted data are from Table IV).

Close modal

For fitting the spectra, the Gauss–Lorentz ratio was allowed to vary freely. Boundaries were set for the satellite energy. These constraints were applied to the lower limit of the position of the unscreened peak with the consideration that the satellite energy corresponds to the plasmon energy (i.e., the surface plasmon energy) and increases with the carrier density.63,64,67,76,77 Since the surface plasmon energy is proportional to the carrier density,64,70 the constraints were such that the satellite energy would be in the range of values reported for photoemission spectra of the 3d orbitals in binary transparent conducting oxides63,64,67,69 with comparable n values as in samples B, C, D, and E. Therefore, the consistent observation of a narrower low binding energy peak and a broader high binding energy peak supports the applicability of the plasmon model to the analysis of the Sn 3d core XPS spectra in these films, which is consistent with previously reported Sn 3d core level spectra in Sb-doped SnO2 samples.63,64,68

Next, the effect of increasing carrier density on the valence and conduction band spectra was explored. Same La:BaSnO3 samples characterized for the core level spectra were used. The XPS spectra of the valence and conduction bands are depicted in Figs. 6(a) and 6(b), respectively.

FIG. 6.

(a) XPS valence band spectra of three samples of different total carrier densities excited with the Al Kα anode: sample D (green curves), sample C (blue curves), and sample B (red and red dashed curves). To investigate how surface absorbed carbonate and hydroxide layers affect the states in the valence band region, after initial measurements of the core levels and valence band spectra, sample B (red dashed curve, surface doped) was intentionally exposed to contamination in the loadlock chamber operated at 1 × 10−8 mbar (not UHV). The carrier densities of all the samples are indicated. The blue, green, and red dashed curves are normalized to the maximum intensity of the red curve. The magenta arrow indicates the shift of the valence band with increasing carrier density. (b) Magnified view of the low energy part of the spectra on (a). The bottom left inset is a schematic illustration of the Moss–Burstein shift, which shows how the doping process affects the electronic band structure of the material. The top right inset is enlarged spectra for the region around the Fermi level.

FIG. 6.

(a) XPS valence band spectra of three samples of different total carrier densities excited with the Al Kα anode: sample D (green curves), sample C (blue curves), and sample B (red and red dashed curves). To investigate how surface absorbed carbonate and hydroxide layers affect the states in the valence band region, after initial measurements of the core levels and valence band spectra, sample B (red dashed curve, surface doped) was intentionally exposed to contamination in the loadlock chamber operated at 1 × 10−8 mbar (not UHV). The carrier densities of all the samples are indicated. The blue, green, and red dashed curves are normalized to the maximum intensity of the red curve. The magenta arrow indicates the shift of the valence band with increasing carrier density. (b) Magnified view of the low energy part of the spectra on (a). The bottom left inset is a schematic illustration of the Moss–Burstein shift, which shows how the doping process affects the electronic band structure of the material. The top right inset is enlarged spectra for the region around the Fermi level.

Close modal

Three main features are observed in the valence band spectra: (i) a mixture of Sn 5s and bonding O 2p orbitals located at 10.6 eV, (ii) the states at 8.3 eV originating from hybridized Sn 5p and O 2p orbitals, and (iii) the bands at binding energies between 4 and 6 eV associated with O 2p bonding or anti-bonding character.78–80 Additionally, shifts of the valence band leading edge to high binding energies upon increasing carrier density can be observed. This indicates an increase in the optical band gap as proposed previously in ellipsometry and HAXPES results.25,32,81 These shifts are correlated with the shifts observed in the core levels (see Sec. S2 of the supplementary material) as well as with the increasing asymmetry in the Sn 3d core lines. Similar trends were observed in other degenerate doped transparent conducting oxides, which were attributed to the increasing occupation of the states in the conduction band.63,67,69 It is noteworthy that an opposite trend (i.e., the shift of the valence band spectra toward lower binding energies with increases in La doping) was reported in recent angle resolved photoemission spectroscopy (ARPES) experiments on La:BaSnO3 films, and it was suggested to originate from the opposite evolution of surface and bulk chemical potentials.46 

To explore further spectral features arising from the occupation of the conduction bands, high resolution scans around the valence band leading edges were acquired [Fig. 6(b)]. To achieve an adequate signal to noise ratio and resolve the fine features in the region close to the Fermi energy (EF), each spectrum was acquired over a period of about 20 hours. A peak at ∼4 eV deriving from O 2p orbitals is observed in all spectra as indicated by their first derivatives (see Fig. S5 of the supplementary material).45,46 For the contaminated surface (sample with the highest n intentionally exposed to contamination in the load lock), a shift to the higher binding energy of ∼0.23 eV is clearly visible in the leading edge valence band spectrum [see red dashed curve in Fig. 6(b)]. Furthermore, a bump is detectable in all spectra in the region between 2 eV and EF [see top right inset in Fig. 6(b)]. The spectra exhibit a weak structure close to EF, which terminates in a sharp Fermi edge. This structure is associated with occupied states in the conduction band (Sn 5s orbital character with a small contribution from O 2p orbitals).32,43 Moreover, the intensity of this CBM peak is observed to increase on the contaminated surface as evidenced by the red dashed curve [see top right inset in Fig. 6(b)]. This suggests that exposure of the surface to contamination results in increasing occupied states in the conduction band. We attribute this behavior to the Moss–Burstein effect, i.e., the apparent optical band gap of the material is increased as the absorption edge is pushed to higher energies as a result of some states close to the conduction band being populated [see, bottom left inset in Fig. 6(b)].82,83 Indeed, the valence band shifts associated with the increasing density of electrons occupying the conduction band were reported in several transparent conducting oxides, resulting in an increase in the intensity of the conduction band feature.25,32,67,69,84

To date, few ARPES studies of the electronic band structure of La:BaSnO3 and BaSnO3 films have been reported.44–46 In Fig. 7, we present the ARPES data of the band structure of a representative La:BaSnO3 film (Sample D) exposed to air for days before cleaning in vacuum.85 These data were collected at room temperature. Although ARPES is a very surface sensitive technique and the samples were exposed to ambient conditions, valence band dispersion is observed from about 3.10 to 10.62 eV [Fig. 7(a)]. The fact that these bands are not clearly resolved is understood in terms of the need for a very particular surface treatment associated with ex-situ ARPES measurements.44Figure 7(b) depicts a high-resolution 2D ARPES map in the binding energies ranging from 6.8 to ∼9.8 eV around the hybridized Sn 5s and O 2p states. Some highly dispersive bands are resolved: the black markers overlaying the dispersing bands are extracted band dispersions obtained from peak fitting of the momentum distribution curves (MDCs), whereas the blue markers are for bands fitted to peaks in the energy distribution curves (EDCs). Figure 7(c) represents the density of states (DOS) integrated from the ARPES map (i.e., EDC obtained from the ARPES map over the entire momentum space). The DOS spectrum exhibits well resolved band features that are similar to the XPS valence band spectra [Fig. 6(a)]. The Fermi–Dirac edge straddling 0 eV is visible [Fig. 7(d)], and a linear extrapolation of the valence band leading edge reveals that the VBM is situated at 3.1 eV [see, black dashed line in Fig. 7(c)]. The extracted VBM value is in close agreement with previous theoretical and experimental values.43–46 For this sample investigated in both XPS and ARPES, the same value for the VBM is obtained from both techniques, and the CBM is not well developed as evidenced by the data for both techniques shown in the insets in Figs. 6(b) and 7(d).

FIG. 7.

(a) Representative 2D ARPES map for sample D. The map was acquired in the Γ̄X̄ direction. (b) ARPES map of the region indicated by the pink dashed rectangle. The inset shows the first Brillouin zone together with the Γ̄,X̄ and M̄ high symmetry points. (c) A representative density of states (DOS) is integrated over the entire momentum space from the ARPES map in (a). The black dashed line is a linear extrapolation of the valence band leading edge, revealing that the valence band maximum is located at 3.1 eV. (d) Enlargement of the region around the Fermi level in (c).

FIG. 7.

(a) Representative 2D ARPES map for sample D. The map was acquired in the Γ̄X̄ direction. (b) ARPES map of the region indicated by the pink dashed rectangle. The inset shows the first Brillouin zone together with the Γ̄,X̄ and M̄ high symmetry points. (c) A representative density of states (DOS) is integrated over the entire momentum space from the ARPES map in (a). The black dashed line is a linear extrapolation of the valence band leading edge, revealing that the valence band maximum is located at 3.1 eV. (d) Enlargement of the region around the Fermi level in (c).

Close modal

In summary, we have systematically investigated the evolution of electronic states in the band structure of La:BaSnO3 films at different La doping levels. A close connection between the transport and the spectroscopic characteristics is demonstrated. In particular, increasing the carrier concentration in the conduction band by doping is observed to significantly affect the core and valence band spectra. The Sn 3d core line shape presents a pronounced asymmetry variation with the carrier density and is fitted following the plasmon model applicable to metallic systems. Scans around the valence band spectra allowed the detection of the occupied states in the conduction bands. It is determined that surface contamination could potentially induce surface carrier accumulation, supported by the increase in the intensity of the CBM detected on the surface exposed to contamination. This study presents a detailed characterization of the chemical composition of the near-surface region of La:BaSnO3, and it provides a better picture of the interplay between the doping concentration, electronic band structure, and transport properties of epitaxial La:BaSnO3 films. The ARPES data presented in this study highlight the challenge of surface preparation for ex-situ ARPES measurements of epitaxial La:BaSnO3 films and heterostructures. Preferably, a portable vacuum suitcase is ideal for long-distance transport of epitaxial La:BaSnO3 films in UHV conditions to perform further in situ ARPES analysis at different locations86 and/or development of protective capping methods for in situ capping to avoid any possible surface contamination when films are exposed to ambient conditions.

See the supplementary material for details on electronic transport measurements, and additional microstructural and photoemission spectroscopy characterizations of La:BaSnO3 films and heterostructures.

The authors acknowledge fruitful discussions and technical support from Kathrin Küster. B. P. Doyle, A. P. Nono Tchiomo, A. R. E. Prinsloo, and E. Carleschi acknowledge funding support from the National Research Foundation (NRF) of South Africa under Grant Nos. 93205, 90698, 99030, and 111985. W. Sigle and P. van Aken acknowledge funding from the European Union’s Horizon 2020 research and innovation program under Grant Agreement No. 823717 -ESTEEM3.

The authors have no conflicts to disclose.

Arnaud Pastel Nono Tchiomo: Formal analysis (equal); Investigation (equal); Writing – original draft (equal); Writing – review & editing (equal). Emanuela Carleschi: Data curation (equal); Formal analysis (equal); Investigation (equal); Supervision (equal); Writing – review & editing (equal). Aletta R Prinsloo: Supervision (equal); Validation (equal); Writing – review & editing (equal). Wilfried Sigle: Investigation (equal); Writing – review & editing (equal). Peter A. van Aken: Investigation (equal); Writing – review & editing (equal). Jochen Mannhart: Supervision (equal); Writing – review & editing (equal). Prosper Ngabonziza: Conceptualization (equal); Formal analysis (equal); Supervision (equal); Writing – original draft (equal); Writing – review & editing (equal). Bryan Doyle: Formal analysis (equal); Investigation (equal); Supervision (equal); Writing – original draft (equal); Writing – review & editing (equal).

The data that support the findings of this study are available from the corresponding authors upon reasonable request.

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The La:BaSnO3 films were prepared at the Max Planck Institute for Solid State Research, Stuttgart, Germany and transported at ambient conditions for further spectroscopic measurements (XPS and ARPES) to the University of Johannesburg, South Africa. Details on the cleaning procedure of the samples and spectroscopic measurements are presented in the supplementary material.

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Supplementary Material