Article pubs.acs.org/cm
Band-Gap Reduction and Dopant Interaction in Epitaxial La,Cr Codoped SrTiO3 Thin Films Ryan B. Comes,† Peter V. Sushko,† Steve M. Heald,‡ Robert J. Colby,§ Mark E. Bowden,§ and Scott A. Chambers*,† †
Fundamental and Computational Sciences Directorate, Pacific Northwest National Laboratory, Richland, Washington 99352, United States ‡ Advanced Photon Source, Argonne National Laboratory, Argonne, Illinois 60439, United States § Environmental Molecular Sciences Laboratory, Pacific Northwest National Laboratory, Richland, Washington 99352, United States S Supporting Information *
ABSTRACT: We show that by co-doping SrTiO3 (STO) epitaxial thin films with equal amounts of La and Cr, it is possible to produce films with an optical band gap ∼0.9 eV lower than that of undoped STO. Sr1−xLaxTi1−xCrxO3 thin films were deposited by molecular beam epitaxy and characterized using X-ray photoelectron spectroscopy and X-ray absorption near-edge spectroscopy to show that the Cr dopants are almost exclusively in the Cr3+ oxidation state. Extended X-ray absorption fine structure measurements and theoretical modeling suggest that it is thermodynamically preferred for La and Cr dopants to occupy nearest-neighbor A- and B-sites in the lattice. Transport measurements show that the material exhibits variable-range hopping conductivity with high resistivity. These results create new opportunities for the use of doped STO films in photovoltaic and photocatalytic applications.
I. INTRODUCTION The unique electronic properties of SrTiO3 (STO) make this material a particularly intriguing candidate for a wide range of technological applications. Epitaxial SrTiO3 films have been shown to exhibit the highest electron mobility of all perovskite oxides with values as large as 30 000 cm2/(V s) at low temperatures.1 The material also exhibits ferroelectric behavior when epitaxial tensile strain is applied through growth on DyScO3 substrates.2 The observed ferroelectricity and large mobility could also make the material ideal for photovoltaic applications, which have also been examined in another ferroelectric perovskite, BiFeO3 (BFO).3,4 STO is also an excellent catalyst for solar hydrolysisthe splitting of H2O into H2 and O2 via visible lightbecause of the favorable alignment of its conduction band edge with the electrochemical potential of the H+/H2 reaction.5,6 However, the large optical band gap of STO (3.2 eV) makes this oxide an inefficient absorber of visible light and limits its potential usefulness.7 With these applications in mind, a variety of approaches have been employed to reduce the band gap of STO in both epitaxial films and powders. The application of 2% biaxial tensile strain in epitaxial STO films has been predicted to reduce the gap by ∼0.2 eV.8 The bands have been mapped using angle-resolved photoelectron spectroscopy to show that uniaxial strains can enhance electron mobilities.9 The effect of dopants has also been explored, particularly Cr substitution for the Ti4+ ion on the B site of the perovskite lattice. The fully occupied Cr3+ 3d © XXXX American Chemical Society
t2g majority spin band lies immediately above the O 2p valence band maximum of STO, raising the valence band maximum of the system so that the band gap is reduced. However, charge conservation dictates that Cr be in the 4+ oxidation state, which acts as an acceptor trap state and is therefore undesirable for most optical applications.10 In SrCrxTi1−xO3−δ films, Cr typically exhibits a mix of 3+ and 4+ oxidation states after growth, with Cr3+ present due to oxygen vacancies that occur because Cr4+ is metastable in octahedral coordination.11 Mixedvalent substitutional Cr and O vacancies have been useful as ionic conductors in memristive devices.12 However, trap states induced by these defects make the material ill-suited for solar hydrolysis and photovoltaic applications. Instead, a material free of O vacancies with Cr in the 3+ state is highly desirable. Recent work in powder samples has suggested a possible solution: doping STO with equal amounts of Cr on the B-site and La on the A-site.13,14 Modeling based on density functional theory (DFT) showed that substituting La3+ for Sr2+ (LaSr) would compensate the Cr4+ acceptors to produce stable Cr3+ more efficiently than other co-dopant ions, including YSr, NbTi, and SbTi.13,15 This approach has been reported to reduce the optical band gap by ∼1.1 eV in powders with doping levels of ∼5% Cr and La substituting for Ti and Sr, respectively, thereby Received: September 25, 2014 Revised: November 6, 2014
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These results create new opportunities for the use of titanate thin films for photovoltaic and photocatalytic applications. The co-doping process demonstrated here also offers new pathways for engineering other functionalities in perovskite films using a variety of transition-metal dopants.
improving the performance of the material for solar hydrolysis.16 Recent theoretical models have predicted that higher doping concentrations, up to 50%, may further reduce the band gap.17 However, dopant concentrations at these levels have not been explored experimentally, to the best of our knowledge. Many of the technological advances are contingent on the availability of thin films with specific properties. However, making compositionally controlled co-doped complex oxide epitaxial films is far more demanding than the analogous powder synthesis. Yet, co-doping STO epitaxial thin films with La and Cr can serve two purposes: (i) enable the engineering of a material with a reduced band gap that could lead to new applications for STO, and (ii) generate fundamental insight into the behavior of co-doped systems prepared one atomic layer at a time. This approach has been employed for La and Fe co-doped STO grown via pulsed laser deposition (PLD), but there was evidence for oxygen vacancies and mixed-valent Fe.18 A similar co-doping scheme with La and Co has also been employed in Bi4Ti3O12 films in the Aurivillius phase grown by PLD, resulting in a reduced band gap.19 To our knowledge, there are no examples of co-doped perovskite films grown using molecular beam epitaxy (MBE).The ability to control the oxidation state of dopants in the perovskite lattice is critical to enhancing the functionality of these films, making MBE growth particularly worthwhile. By changing the co-dopant ratio from 1 A-site to 1 B-site dopant to 2 A-sites for every B-site, it should be possible to stabilize other functional oxidation states. These might include Ni2+, which has been shown to exhibit photoluminescence in BaxSr1−xTiO3,20 and Co2+, which exhibits ferromagnetism when doped in TiO2.21 Such an approach would be readily achievable using the MBE scheme applied in this work. The behavior of compensating dopants on the two cation sites of the perovskite lattice and the tendency to occupy nearest neighbor sites has been the subject of significant theoretical work. The interaction energies of transition-metal dopants, including Ni, Fe, and Cr, along with compensating Asite ions, including Y3+, La3+, Pr3+, Gd3+, Tb3+, and Er3+ in BaTiO3 was explored through atomistic simulations.22 The interaction between Cr3+ and the A3+ compensating dopants was predicted to be strong, with a binding energy of 2.35 eV. Other works predicted a substantially smaller binding energy of 0.15 eV, which was not believed to be sufficient to promote ordering, given the thermal energy present during hightemperature (1100 °C) annealing of powder samples.13 However, there was no experimental verification of this prediction. Extended X-ray absorption fine structure (EXAFS) experiments have been used to confirm the presence of nearestneighbor oxygen vacancies in oxygen-deficient Cr3+-doped SrTiO3,23 but explorations of the A-site and B-site coupling have not been reported. Here, we describe the properties of Sr1−xLaxTi1−xCrxO3 (SLTCO) epitaxial thin films grown via MBE. We show that it is possible to precisely control La and Cr dopant concentrations and reduce other compensating defects to levels that are within the uncertainty of the growth process. Through theoretical modeling and EXAFS measurements, we show that the growth conditions employed in this work lead to the incorporation of La3+ and Cr3+ dopants on nearest-neighbor A- and B-sites in the STO lattice. Finally, we demonstrate that the films exhibit semiconducting behavior with a reduction in the optical band gap of ∼0.9 eV, relative to that of pure STO.
II. METHODS A. Experiments. Oxygen-assisted molecular beam epitaxy (OAMBE) was employed to grow a series of SLTCO films with Cr(La) A(B) site doping levels (x) of 0.03, 0.10, and 0.20. Details of the growth process can be found in the Supporting Information. The doped films were grown in molecular O2 at a pressure of 3 × 10−6 Torr and at a substrate temperature of 700 °C. The Sr and Ti fluxes were calibrated via homoepitaxial growth of STO on an STO(001) substrate, with effusion cell temperatures adjusted to produce consistent reflection high-energy electron diffraction (RHEED) intensity oscillations for the (11) Bragg rods.24 This technique has been shown by others to produce films with Sr:Ti ratios within 1 at. % for STO films, as confirmed by in situ X-ray photoelectron spectroscopy (XPS) through comparison of the Ti, Sr, and O corelevel peaks areas with those measured on a STO(001) reference crystal. The SLTCO growth was subsequently carried out on a new STO substrate, with an initial 5−10 nm thick homoepitaxial film deposited first to remove any surface defects on the substrate and confirm that no significant drift in the atom beam fluxes had occurred. Molecular O2 was employed during SLTCO growth, rather than activated oxygen supplied by an electron cyclotron plasma or ozone source, because of the propensity of Cr to overoxidize to Cr6+ during growth in such an atmosphere.25 To co-dope the films, a shuttered growth scheme was employed beginning with a SrO deposition sufficient to produce 1 − x unit cells of the desired material. LaCrO3 (LCO) was then deposited using the Cr and La effusion cells to nucleate x unit cells, followed by a 1 − x unit-cell deposition of TiO2. We have previously shown that it is possible to control the La:Cr ratio to within ∼5 at. % in pure LCO films.25 Films were capped with 2 unit cells of pure STO to prevent overoxidation of Cr during the cooling process, which was done in 3 × 10−6 Torr O2. Figure 1a shows a representative plot of the time dependence of the RHEED intensity near the end of a 50-nm growth with annotations and an inset showing the opening times for each shutter. Instantaneous transients are seen when shutters open and close due to interference of the magnetic shutter actuators with the RHEED gun. Figure 1b shows the RHEED pattern along the [110] azimuth following the growth, an out-of-plane X-ray diffraction (XRD) ω-2θ scan, a reciprocal space map about the STO(103) peak, and an atomic force micrograph of the surface topography for a film with x = 0.10. These data clearly indicate that the film grew in a layer-by-layer fashion, coherent to the substrate with a well-ordered and atomically flat surface at the end. After growth, the samples were transferred under ultrahigh vacuum to an appended XPS system with a monochromatic Al Kα1 source (hν = 1486.7 eV) and a VG Scienta R3000 electron energy analyzer. Details of the XPS measurements and data analysis can be found in the Supporting Information. Selected samples were analyzed using a variety of ex situ techniques including optical absorptioņ resistivity, Xray diffraction (XRD), scanning-transmission electron microscopy (STEM), X-ray absorption near-edge spectroscopy (XANES), and extended X-ray absorption fine structure (EXAFS). XANES and EXAFS measurements were performed at the Advanced Photon Source on Beamline 20-BM. For additional details on the various characterization techniques, see the Supporting Information. B. Modeling. Ab initio simulations were carried out using DFT. Most of the calculations were performed using the Perdew−Burke− Ernzerhof generalized gradient approximation functional (PBE)26 and the projected augmented waves method27 implemented in the Vienna Ab Initio Simulation Package.28,29 For the most stable configurations, we also used the PBE modified for solidsPBEsol30and investigated the effects of the Hubbard U correction, plane-wave basis set cutoff, and the lattice parameters. Details of these calculations can be found in Table S2 in the Supporting Information section. The B
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the vacuum gap is over 17 Å. Since one of the slab surfaces is an artifact of the model and, in fact, corresponds to a part of the system that is buried inside the film, the relaxation and rumpling associated with this surface of the slab were suppressed by fixing atoms in two outermost atomic planes at the ideal bulk lattice sites. The calculations were carried out for the Γ point of the Brillouin zone only.
III. RESULTS AND DISCUSSION A. Oxidation States and Stoichiometry. The heteroepitaxial growth of SLTCO films with controlled stoichiometry presents unique challenges, because of the range of oxidation states exhibited by Cr25 and the impact that oxygen vacancies31 and La dopants1 have on the transport properties of SrTiO3 films. Films were grown in a molecular oxygen environment, which might be expected to lead to oxygen deficiency in the films, leading to unintentional carrier doping. Similarly, slight variations in the flux ratio for La and Cr could unintentionally dope the material. To examine these effects, XPS and electron energy loss spectroscopy (EELS) via scanning transmission electron microscopy (STEM) measurements were made on the films and compared to reference spectra for undoped STO. Figure 2 shows XPS and EELS spectra for 10%-doped SLTCO, with the O 1s XPS peaks (Figure 2a), the Ti 2p (Figure 2b) XPS peaks, and the O K-edge EELS overlaid on a STEM image of the sample (Figure 2c). The O 1s XPS peak areas match to better than 1% for all films, with no systematic variation between the undoped films and doped films. This result shows that not using activated oxygen during film growth does not lead to an increase in oxygen vacancy concentrations, as has also been seen for STO films grown in O2.32 STEM EELS O K-edge spectra acquired for the doped film and the substrate are very similar below 570 eV as well, with almostidentical peak positions for all features and no additional peaks in the doped film that might suggest the presence of holes in the valence band.10,33 The Cr L-edge dopant peak is also visible in the film spectrum at 575−590 eV. The area-normalized Ti 2p spectra shown in Figure 2b for a 10%-doped sample and a STO reference film show no evidence of Ti3+. Ti3+ appears when La donors in La-doped STO films provide electrons that occupy Ti 3d states at the bottom of the conduction band. The Ti3+ 2p3/2 peak falls at slightly lower binding energy (∼1.5−2 eV) than the Ti4+ peak and has been observed for La doping levels as low as 0.7%.34,35 The absence of such a peak reveals that Ti3+ comprises, at most, ∼0.5% of the total Ti in the lattice, and this matches the experimental uncertainty with which we can measure our fluxes. This result is consistent with La donors being compensated by the Cr acceptors and not contributing electrons to the conduction band, as expected for equal concentrations of La and Cr. The same is true for all other films. Cr 2p XPS for 3%-, 10%-, and 20%-doped SLTCO films are shown in Figure 3a, along with spectra for LaCrO336 and SrCrO310 reference samples. The Cr 2p spectra shown in Figure 3a allow us to tentatively assign the oxidation states of the Cr dopants. The binding energies and the multiplet splitting of the 2p3/2 peak at ∼575−576 eV in the 10%- and 20%-doped samples suggests that the majority of Cr ions are in the 3+ state. However, this peak is broader than its equivalent in pure LCO, leaving uncertainty as to whether there may be some Cr ions in a different valence state. The 3%-doped film shows intensity in the binding energy range characteristic of higher oxidation states (576−578 eV). The presence of the STO capping layer significantly reduces the signal intensity,
Figure 1. (a) RHEED (11) diffracted beam intensity oscillations at the end of a 50-nm 10%-doped film growth with a 2-unit-cell STO capping layer, with inset showing sequential shuttering process. (b) Out-ofplane XRD scan of a 10%-doped film overlaid on a representative AFM surface topography measurement. (Inset) Reciprocal space map of the same 10%-doped sample about STO(103); RHEED pattern in the (110) azimuth for the same sample shown in panel (a). plane-wave basis set cutoff was set to 400 eV. The bulk STO was modeled using the 320-atom cubic supercell constructed of 4 × 4 × 4 STO crystallographic unit cells. In addition, two models of STO slab were used in order to simulate the interaction between the La and Cr impurities at different stages of the film growth. The slab containing eight STO atomic planes corresponds to a stage, at which a full TiO2 plane of STO is already formed and the next SrO plane has not started growing yet. In other words, it corresponds to the completed SrO (97%)−LCO (3%)−TiO2 (97%) growth cycle (see inset in Figure 1a). In contrast, the slab containing seven atomic planes is used to model the La−Cr configurations that are formed after the SrO plane and the LCO are deposited but the growth of the next TiO2 plane has not started yet. In other words, it corresponds to the half-completed growth cycle (see Figure 1a). Thus, the former model helps to evaluate stability of the La−Cr configurations after the film is grown, while the latter model is need to evaluate the La−Cr interaction and association in the process of growth. For both slab model the lateral cell was 4a0 × 4a0 (a0 is the STO bulk lattice constant). In all slab model calculations, C
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Figure 3. Cr K-edge X-ray absorption near-edge spectroscopy (XANES) data for SLTCO films, with reference spectra from LaCrO3,36 CrO2, and Cr3+-doped SrTiO3.11
XANES results, along with reference spectra for LCO36 and Cr3+-doped STO11 are shown in Figure 3b. The line shape closely resembles that of a published spectrum for Cr3+-doped STO near the K-edge and LaCrO3 at higher energies, suggesting that the dopants are almost exclusively in the 3+ state. There is no evidence of a pre-edge peak at ∼5990 eV that would be associated with Cr5+ or Cr6+ dopants. In addition, the spectrum for Cr4+ from a CrO2 sample (also shown in Figure 3) shows a higher absorption onset energy due to the larger binding energy of 1s electrons on Cr4+, compared to Cr3+.36 To summarize, spectroscopic analysis of the samples has confirmed that Cr dopants are stabilized in the 3+ oxidation state through the La co-doping process. O 1s XPS spectra and K-edge EELS show no evidence of oxygen vacancies in the films, indicating that the La donors compensate the Cr dopants to produce a film free of the intrinsic electrons of oxygen vacancies that would be expected for Cr3+-doped STO. Ti 2p XPS spectra show that the Ti valence is unchanged from the ideal Ti4+ state in STO, indicating that the excess electrons donated from La lie on the Cr acceptors. Collectively, these observations confirm that the co-doping scheme produces an intrinsically doped system with no compensating defects.
Figure 2. (a) O 1s XPS spectra for a 10%-doped film and a reference SrTiO3 (STO) substrate, showing equal oxygen peak intensities; (b) Ti 2p XPS spectra for 10%-doped film and STO reference; (c) HAADF-STEM image of 10%-doped film and substrate, with interface denoted by horizontal white bars; and (d) O K-edge EELS spectra for the substrate and film.
making precise exclusion of other oxidation states difficult. It is also possible that some Cr dopants at the surface overoxidize during the cooling process, even with the STO cap. To better determine the oxidation state in the bulk of the films, as well as the Cr coordination environment, Cr K-edge XANES measurements were performed on 3%- and 10%-doped samples. The D
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which is close to the Cr−O bond length in LaCrO3.25 This value is slightly smaller than bond lengths used in previous EXAFS models of oxygen-deficient Cr-doped STO of 1.95− 1.97 Å.23 The lattice parameter of oxygen-deficient Cr-doped STO would be expected to expand, because of ionic repulsion of the cations near oxygen vacancies, indicating that the bond length in the model is reasonable and that the films have ideal oxygen stoichiometry. In modeling the EXAFS, a mixture of Sr and La ions on the A-site was assumed and the relative fraction of Sr and La was varied to determine if La preferentially occupies B-sites near the Cr dopant ions. At a dopant concentration of 3%, there should be 7.76 Sr ions (NSr) and 0.24 La ions (NLa) at nearestneighbor A-sites, if there is no preferential La occupancy near the Cr dopant. A model assuming this condition produced an R-factor of 0.0091, indicating good statistical agreement with the data. However, by increasing the occupancy of La ions (NLa) on the nearest neighbor A-site, the fit could be improved. The best fits occurred for NLa between 1.5 and 2.5, with Sr ions on the remaining sites. In this range, the R-factor was within 10% of the minimum value, which was achieved for NLa equal to 2.0. The EXAFS data for χ(R) and fits with NLa = 0.24, 1.5, and 2.0, and R-factor dependence on NLa for the fits are shown in Figure 4a and 4b, respectively. Fits based on larger values of NLa resulted in much-higher R-factors, ruling out the possibility that clusters of pure LaCrO3 formed within an STO lattice. In addition, the EXAFS data for a LaCrO3 reference sample are a very poor match to the observed data for the doped films for R = 2−4 Å. Instead, we observe that the positions of La and Cr, on A- and B-sites, respectively, are correlated, possibly driven by the stabilization of Cr3+ at B-sites via electron donation from La3+ at adjacent A-sites. Details of the 3% fit are shown in Table S1 and Figure S2 in the Supporting Information, along with the reference LaCrO3 EXAFS data. To corroborate this result, we constructed a series of computational models for La and Cr dopants in STO to estimate the interaction energies between the dopants. First, we considered the La and Cr impurities in the bulk STO, modeled using a Sr63La1Ti63Cr1O192 supercell, and calculated the total energies for the systems with the smallest distance between the La and Cr impurities being 3.4, 6.6, 8.6, and 10.3 Å. We found that, although configurations with smaller La−Cr distances are thermodynamically preferred, the La3+−Cr3+ association energy is small. In particular, for the system with all atoms occupying unrelaxed, bulk-like lattice sites, the La−Cr association energy is ∼0.3 eV. However, if the total energy of the system is minimized, with respect to the fractional coordinates and the lattice parameters, this energy decreases to 0.02 eV due to atomic displacements in the lattice that effectively screen the electrostatic interaction of the La3+ and Cr3+ impurities. This suggests that there would not be any appreciable binding of La and Cr at the growth temperatures and, therefore, no appreciable deviation from the random spatial distribution of these impurities in the bulk. This result is in agreement with the theoretical work of Reunchan et al., which showed that ordering would not be expected in bulk synthesized powders.13 It should be noted that the effect of the atomic displacements on the screening of the electrostatic interaction between the La3+ and Cr3+ impurities in the STO is expected to be different in the vicinity of the surface, i.e., at the growth front, than in the bulk. In addition, this screening is also affected by substrateinduced constraints on the in-plane lattice parameters. Hence, the association energy of the La3+ and Cr3+ ions calculated
B. Dopant Configuration. XRD measurements confirmed that all films were coherently strained to the STO substrates, as expected. The unit-cell volume for representative 3%- and 10%doped samples was estimated by assuming in-plane lattice parameters of 3.905 Å and measuring the out-of-plane lattice parameter from the (002) diffracted peak. For both samples, the unit cell volume is reduced slightly from that of undoped STO and the total volume agrees within 0.1% of the value that would be expected from applying Vegard’s Law37 to alloys of STO and LCO (apc = 3.885 Å) at the corresponding concentrations.25 These results suggest that the Cr−O bond length in the doped films should be nearly equal to that in LCO. EXAFS measurements at the Cr K-edge were also acquired for 3%- and 10%-doped samples and showed very similar results for the two samples. The Fourier transform of the EXAFS data for the 3%-doped sample is shown in Figure 4a, along with fits to the data obtained by modeling all electron scattering events that occur within a one-unit-cell effective radius (∼3.9 Å) from the Cr dopant ion. The Cr dopant was assumed to be at the B-site with a Cr−O bond length of 1.94 Å,
Figure 4. (a) Cr K-edge EXAFS data and fits assuming different La nearest-neighbor occupancy; (b) R-factor of fits for different La occupancies. Circles represent R-factors within 10% of the minimum value. The star indicates what is expected for a completely random arrangement of La and Cr dopants on the A- and B-sites, respectively, and colors match plots from panel (a). E
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plane. This configuration replicates the expected adatom configuration based on the actual film growth protocol shown in Figure 1a, with La and Cr ions are deposited on an almostcompleted SrO plane. A Cr atom was positioned at various distances from the LaSr site and both La and Cr were coordinated with oxygen in either atomic or molecular form, with the number of the oxygen species varying between one and four. In this scenario, since the La3+ dopant has an effective charge of +1, with respect to the lattice, it may attract the compensating negatively charged CrOx species. Alternatively, an oxygen atom or an O2 molecule adsorbed at the La3+ site form O−/La3+ or O2−/La3+ complexes. These complexes have the same charge as lattice Sr2+ ions, i.e., neutral with respect to the lattice, which makes the electrostatic interaction between the La and Cr complexes negligible. Several configurations of La−Cr−O defects are shown in Figure S4 in the Supporting Information. The dependence of the free energies for several of the most stable La−Cr−O configurations on the oxygen chemical potential (ΔμO) is shown in Figure 5. If the La and Cr atoms are ∼10 Å apart and ΔμO is low, the most stable configurations are formed by separated La3+ and [CrO]− (ΔμO < −4.5 eV, not shown), La3+ and [CrO2]− (−4.5 eV < ΔμO < −4.1 eV), and La3+ and [CrO3]− (ΔμO > −4.1 eV). Essentially, the oxidation state of the Cr surface ion increases with increasing oxygen chemical potential, as would be expected. As the value of ΔμO increases above ΔμO ≈ −1 eV, the most stable configuration becomes a O2−/La3+ complex, which is neutral with respect to the STO lattice, and the [CrO3]0 complex (denoted in blue in Figure 5c). A value of ΔμO ≈ −1 eV corresponds to the bulk synthesis regime, where dopant pairing is not expected. In contrast, if the La and Cr are located at the neighboring sites, the most stable configuration at low ΔμO (ΔμO < −3.9 eV) is formed by [CrO2]− adsorbed near the La3+ site, whereas at ΔμO > −3.9 eV, it becomes the [CrO3]−/La3+ configuration (denoted in black in Figure 5c). The ions occupy nearest-neighbor sites and form LaCrO2 (ΔμO < −3.9 eV) and LaCrO3 (ΔμO > −3.9 eV) neutral species (denoted in green in Figure 5c). We also note that our experimental conditions (3 × 10−6 Torr O2 at 973 K) correspond to ΔμO ≈ −2 eV.38 Judging from the tetrahedral local environment of Cr in the [CrO3]− complexes, as well as from the spin-density distribution, we conclude that the oxidation state of Cr at this chemical potential is 4+ for both separated and neighboring La and Cr dopants. Finally, comparison of the free energies calculated for these configurations shows that the association energy of La3+, buried in the outermost SrO plane, and [CrO3]− complex deposited on this plane prior to the growth of the complete BO2 plane, is between 0.1 eV and 0.2 eV. The lower limit value was obtained using the PBE and the experimental value of the lattice constant for the lateral cell of the slab (3.905 Å). Using the lattice parameters obtained for the bulk STO via PBE (3.945 Å) increases the La−Cr binding energy by ∼0.05 eV. These results are consistent with the La−Cr association energy calculated for the slab model above. Applying the +U correction, where U = 3 eV, increases this energy by another ∼0.01 eV. Finally, using PBEsol+U (U = 3 eV) raises this energy further to 0.2 eV. To estimate the value of NLa on the basis of these calculations, we constructed a model that takes into account the layer-by-layer growth of the La,Cr co-doped SrTiO3, the dependence of the thermodynamic stability of the La−Cr pairs on the distance between them, as well as the entropic
using the bulk model described above may not be applicable to our experimental conditions. Therefore, we performed a similar series of calculations using the STO slab model, in which the Cr and La dopants were located in the outermost and secondoutermost STO atomic planes, respectively (see Figure 5a) and
Figure 5. (a) Energies of La−Cr configurations at the film surface in the slab model. (b) View of adatom configuration for third-nearest neighbor La and Cr. (c) View of adatom configuration for nearest neighbor La and Cr. (d) Energy for various La dopant configurations as a function of oxygen chemical potential. In these configurations, La ions are shown in orange (largest size), Cr in yellow (tetrahedron), Sr in green (medium size), O in red/pink (smallest), and Ti in blue (octahedra).
found the calculated association energy of ∼ 0.1 eV. Given that the growth temperature (700 °C) corresponds to a thermal energy of 84 meV, the La−Cr interaction energy is of the same order of magnitude as the thermal energy of the material and a small amount of ordering might be expected. However, these results are not sufficiently conclusive to categorically predict the dopant ordering in SLTCO that was observed in EXAFS. To better assess La−Cr interaction during film growth, we explored surface structures that are likely to form immediately after deposition of the fractional monolayer coverage of LaCrO3 but prior to the TiO2 monolayer deposition. To this end, we considered the STO slab terminated with an SrO plane and containing a single LaSr impurity in the outermost SrO F
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contributions due to the configurational space and temperature (see Figure S3 in the Supporting Information for details). According to this model, in the layer-by-layer growth of the film, La species tend to associate with Cr deposited at the previous growth step and vice versa. As a result, La and Cr assemble into quasi-one-dimensional structures with NLa of 1.25 for the La−Cr association energy of 0.2 eV and the temperature of 1000 K. For comparison, in the bulk model, where the La−Cr pair assumed to not interact, NLa is estimated to be ∼0.82 and can never exceed 1. Thus, we propose that the electrostatic interaction between the La3+ and [CrO3]− favors their association. Since the LaCrO3 deposition step occurs over a period of ∼2 s prior to the opening of the Ti shutter, the association of La3+ and [CrO3]− proceeds unimpeded by the presence of other species. Once the TiO2 deposition step begins and the Ti portion of the BO2 atomic layer is formed, Cr4+ ions in tetrahedral oxygen coordination are reduced to Cr 3+ ions in octahedral coordination. At this stage of the process, the Ti and additional oxygen species in the BO2 plane would confine the Cr dopants and, thus, suppress the rate of dissociation of Cr3+ from La3+, thereby freezing their spatial arrangement. The observed dopant coupling is expected to be unique to the layer-bylayer growth process employed here, where films grow in a kinetically limited regime with lower oxygen chemical potential than that used in bulk synthesis. C. Band Gap. Valence-band (VB) XPS measurements were performed to examine the effects of Cr3+ doping level on the overall VB density of states. VB spectra for undoped STO and 3%- and 10%-doped films are shown in Figure 6a. The spectra are aligned so that the O 2p band edges match, with the edge closest to the Fermi level falling at 3.1 eV on the scale. Precise determination of the Fermi level is not possible due to the use of the low-energy electron flood gun compensating surface charging from photoemission. The practice typically adds a constant offset to the binding energy scale dictated by the flood gun settings. Thus, the binding energy scale is arbitrary. For both the 3%- and 10%-doped samples, a small Cr 3d t2g derived feature appears above the edge of the O 2p band. The same result occurs for 20%-doped samples, but these are not included in the figure. The observed band matches theoretical predictions for the Cr-derived partial density of states when Cr3+ is doped on the A-site of STO performed in our group and elsewhere.13,14 The projected density of states for our model is shown as Figure S5 in the Supporting Information. Fitting to a Fermi function yields a VB maximum of 1.9 ± 0.2 eV for both samples. The O 2p bands exhibit identical shapes and widths in the doped films and pure STO, indicating that co-doping has a negligible effect on the more tightly bound portions of the VB. Thus, the primary effect of the Cr3+ dopants is to produce additional electronic states above the O band edge and reduce the band gap of the material. The band edge extrapolations suggest that the gap may be reduced by as much as 1.2 eV, which is in good agreement with what was observed by others in powder samples.13 To measure the optical band gap of the doped films, optical absorption measurements were performed on representative 3%-, 10%-, and 20%-doped samples, along with a bare STO reference substrate, to directly determine the experimental gaps. Details of the experimental methods and data analysis techniques can be found in the Supporting Information. Absorption versus photon energy plots for both films are shown in Figure 6b. Photographs of representative samples are
Figure 6. (a) Valence-band (VB) XPS spectra for doped films and SrTiO3 reference. (b) Optical absorption spectra for the three films, with insets showing the actual color of samples and percent absorption for all three samples and a reference SrTiO3 substrate.
shown in the inset of Figure 6b, along with the absorption percentage for all three samples and the STO reference. The 20%-doped film was 25 nm thick, while the 3%- and 10%doped films were 50 nm thick, producing a similar absorption percentage in the 20%-doped film, relative to the other two samples, despite the greater absorption coefficient (α). The photographs show that the doped films are yellow and that the intensity of the yellow color increases with doping level. Negligible absorption is observed for photon energies below 2.3 eV, with increasing absorption above that energy up to an energy of ∼3.2 eV, where absorption from to the substrate occurs. Linear extrapolation of the data between 2.4 and 2.6 eV produces an intersection with the background signal of ∼2.3 eV for all three samples. The absorption intensity is greater for more highly doped samples, because of the larger number of occupied Cr t2g valence states available for electron−hole pair excitation. Computational modeling predicts that, for dopant levels from 3.7% to 25%, the leading edge of the Cr 3d t2g VB G
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maximum would decrease in binding energy by ∼0.3 eV, decreasing the band gap accordingly.17 We do not observe such a decrease, although the uncertainty in the measurements is comparable to the predicted change. The band gap is not dependent on the doping level, in agreement with the observations from the XPS VB measurements. The combination of VB XPS and optical absorption measurements confirms that the doping scheme employed here reduces the band gap of the films by 0.9−1.1 eV, relative to the 3.2 eV gap for undoped STO. This reduction indicates that the films absorb at wavelengths below ∼550 nm. D. Electronic Transport. Electronic resistivity measurements were performed on 3%- and 10%-doped samples to determine the effects of co-doping on conductivity. Ohmic contacts of Cr were sputter-deposited onto the samples39 using a shadow mask in the van der Pauw configuration. Resistivity measurements were performed over a temperature range of 50−300 K using a Quantum Design Physical Property Measurement System. The temperature-dependent resistivity is shown in Figure 7. The 3%-doped sample exhibits
variations in dopant concentration on either the A-site or the Bsite at these levels are beyond the detection limits for XPS and RBS and are likely to be beyond the capabilities of secondary ion mass spectrometry as well. However, 0.5% excess La would be expected to produce an electron-doped sample at room temperature. Included in Figure 7 is a point indicating the resistivity of 0.5% La-doped SrTiO3 at 300 K.43 The temperature dependence for doping at this level indicates that the material is metallic. The room-temperature resistivity is very close to the measured value at 300 K for the 10%-doped sample, suggesting that a slight excess of La may be the origin of the relatively low resistance of that sample. The resistivity of the 3%-doped film is roughly 1 order of magnitude larger at room temperature than the 10%-doped sample, despite having only about three times fewer dopant ions. This suggests that the dopant concentrations in this film match with significantly better than 5% uncertainty, producing what might be considered a more ideal film.
IV. CONCLUSIONS We have prepared and investigated epitaxial Sr1−xLaxTi1−xCrxO3 films with doping levels of 3%, 10%, and 20%. XPS measurements confirm that the oxygen content of the films matches that of pure STO, indicating that the growth technique produces high-quality films with extremely low oxygen vacancy concentrations. Through Cr K-edge XANES measurements, we have shown that the Cr dopants are in the 3+ oxidation state with no evidence to suggest overoxidation to Cr4+ or Cr6+. Fitting of Cr K-edge EXAFS data suggests that the La concentration on the nearest-neighbor A-site is larger than would be expected for a random distribution. These results indicate that there is a driving force for La and Cr dopants in the STO lattice to occupy nearest-neighbor sites and thus exhibit a high degree of spatial correlation. Computational modeling of the film surface structure at different growth stages is in agreement with these results. Strong Cr−La interaction is likely to be specific to the film growth process when compared to other forms of bulk synthesis due to the lower oxygen chemical potential during growth. Future experimental and theoretical studies on judiciously selected dopant combinations could further elucidate the effect of the oxygen chemical potential on the electrostatic interaction and binding of the dopants. Using optical absorption and valence band XPS measurements, we have shown that the band gap of the SLTCO films is ∼0.9 eV lower than that of undoped STO. Resistivity measurements show that the material exhibits semiconducting three-dimensional variable-range hopping behavior. This reduction in the optical band gap offers intriguing possibilities for the use of SLTCO films in photocatalytic and photovoltaic applications. The co-doping technique can also be readily employed to other B-site dopants to control their oxidation states and engineer new functionalities in STO and other perovskite films.
Figure 7. Resistivity measurements for 3%- and 10%-doped films with 3D variable range hopping fits to data shown in the inset.
significantly higher resistivity than the 10%-doped one, although both exhibit similar semiconducting behavior. Fits to the data, which are shown in the inset to Figure 7, indicate that the transport in both films can be well-modeled via a Mott three-dimensional variable range hopping (3D-VRH) model: log(ρ) ∝ (T0/T)1/4.40 Fits to other models, such as the Efros− Shklovskii hopping model,41 log(ρ) ∝ (T0/T)1/2, were less satisfactory. The fitting parameter T0 for the Mott 3D-VRH model decreases with increasing doping density, most likely due to the decreased hopping distance between dopant ions. While the films exhibit semiconducting behavior, the resistivity at room temperature is somewhat lower than what might be expected, given that nominally equal doping concentrations of La and Cr should not produce free carriers in the system. As noted in the discussion of film stoichiometry, however, the precision of the La and Cr flux ratios is limited to roughly 5% from sample to sample, because of the intrinsic uncertainty in Rutherford back scattering (RBS) fitting and quartz crystal rate monitors.42 Thus, for a 10%-doped film, the uncertainty in the fraction of La occupying the A-site of the lattice is ∼0.5%, and likewise for Cr on the B-site. Small
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ASSOCIATED CONTENT
S Supporting Information *
Additional details and figures of the modeling procedures, film growth process, XPS and EXAFS data analysis, and other characterization techniques. This material is available free of charge via the Internet at http://pubs.acs.org. H
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AUTHOR INFORMATION
Corresponding Author
*E-mail:
[email protected]. Author Contributions
R.B.C. performed film growth, X-ray photoelectron spectroscopy (XPS), and optical absorption measurements. P.V.S. performed computational modeling. S.M.H. performed X-ray absorption measurements, and S.M.H. and R.B.C. analyzed the data. R.J.C. and R.B.C. performed electron microscopy measurements and analyzed the data. M.E.B. performed X-ray diffraction (XRD) measurements. R.B.C., P.V.S., and S.A.C. wrote the manuscript. S.A.C. and R.B.C. conceived of the project. All authors participated in discussions of the work and have approved of the final version of the manuscript. Notes
The authors declare no competing financial interest.
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ACKNOWLEDGMENTS R.B.C. was supported by the Linus Pauling Distinguished Postdoctoral Fellowship at Pacific Northwest National Laboratory (PNNL LDRD PN13100/2581). S.A.C. and M.E.B. were supported at PNNL by the U.S. Department of Energy, Office of Science, Division of Materials Sciences and Engineering, under Award No. 10122. PVS was supported by the Laboratory Directed Research and Development Program at PNNL. The PNNL work was performed in the Environmental Molecular Sciences Laboratory (EMSL), which is a national science user facility sponsored by the Department of Energy’s Office of Biological and Environmental Research and located at Pacific Northwest National Laboratory. R.J.C. was supported by the EMSL William Wiley Postdoctoral Fellowship program. PNC/XSD facilities at the Advanced Photon Source, and research at these facilities, are supported by the U.S. Department of Energy−Basic Energy Sciences, the Canadian Light Source and its funding partners, the University of Washington, and the Advanced Photon Source. Use of the Advanced Photon Source, an Office of Science User Facility operated for the U.S. Department of Energy (DOE) Office of Science by Argonne National Laboratory, was supported by the U.S. DOE, under Contract No. DE-AC02-06CH11357.
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