Metal‐to‐insulator transition in oxide semimetals by anion doping

Oxide semimetals exhibiting both nontrivial topological characteristics stand as exemplary parent compounds and multiple degrees of freedom, offering a promise for the realization of novel electronic states. In this work, we report the structural and transport phase transition in an oxide semimetal, SrNbO3, achieved through effective anion doping. Notably, the resistivity increased by more than three orders of magnitude at room temperature upon nitrogen‐doping. The extent of electronic modulation in SrNbO3 is strongly correlated with misfit strain, underscoring its phase instability to both chemical doping and crystallographic symmetry variations. Using first‐principles calculations, we discern that elevating the level of nitrogen doping induces an upward shift in the conductive bands of SrNbO3−δNδ. Consequently, a transition from a metallic state to an insulating state becomes apparent as the nitrogen concentration reaches a threshold of 1/3. This investigation shows effective anion engineering in oxide semimetals, offering pathways for manipulating their physical properties.


Main text
Dirac semimetals emerge as remarkable quantum materials, brimming with enigmatic topological peculiarities.In the Dirac semimetal, the conduction and valence bands contact only at discrete (Dirac) points in the Brillouin zone and disperse linearly along all directions around these critical points [1][2][3][4], exhibiting giant diamagnetism, linear quantum magnetoresistance, and quantum spin Hall effect.Researchers have employed many strategies to realize precise control over the structure and physical properties of Dirac semimetals, encompassing the incorporation of 3d magnetic elements exhibiting strong correlations through doping schemes [5] and the fabrication of heterostructures incorporating magnetic and/or superconducting materials [6].However, the investigation pertaining to in situ regulation methods remains relatively unexplored, and there is a dearth of comprehensive research in this domain.The study of the physical properties of transition metal oxides reveals its paramount significance in scientific research.Recently, the 3D Dirac semimetals have been experimental explored in oxides, such as -cristobalite BiO2 [1] and SrNbO3.Ok et al. [4] reported that SrNbO3 thin film exhibit an emerging topological band structure.By manipulating the octahedral rotations in transition metal oxides, a novel semimetal was created in strained SrNbO3, leading to the realization of a novel correlated topological quantum state.Although this method exhibits commendable efficacy, it is not devoid of limitations, encompassing challenges such as the large degree of stress regulation.Consequently, to facilitate the comprehensive exploration of both in situ control and the physical properties of oxide semimetals, novel methodologies need to be adopted.
The burgeoning field of anion engineering in transition metal oxides holds immense promise in fine-tuning their physical properties through the introduction of anions with different size, electronegativity, and charge [7][8][9][10].Heteroanionic materials, exemplified by oxyhydrides [11,12], oxynitrides [13,14], oxyfluorides [15,16], and others [17,18], present a fertile ground for the exploration of novel or enhanced responses, including superconductivity [19], giant magnetoresistance effect [20,21], and visible photocatalytic activity [22,23].In particular, a technology of in-situ nitrogen doping in oxide films is realized through the combination between pulsed laser deposition and a radio frequency nitrogen atom source [24,25].This approach exhibits simplicity, efficiency, and environmental compatibility, and also enables realtime adjustment of dopant concentration.Consequently, the method is poised to take on a great role in the precise control and regulation of material properties.
In this paper, we report the successful synthesis of highly epitaxial SrNbO3 (SNO) thin films and SrNbO3-N (SNON) oxynitride thin films using in situ nitrogen (N) doping.This deliberate substitution of N drives consequential alterations in the valence state of Nb ions and subsequently modifies the electrical transport properties of the materials under investigation.Furthermore, we discuss the strain dependency exemplified by both SNO and SNON films.
These observations are in excellent agreement with our first-principles calculations.In this manner, utilizing the paradigm of anion engineering, we unveil a meaningful pathway toward the regulation and manipulation of the distinctive properties characterizing the semimetal phase in oxide thin films.

Results
Structural transition induced by nitrogen doping.The SNO and SNON thin films with thickness of ~20 nm were grown on (001)-oriented SrTiO3 (STO) substrates by pulsed laser epitaxy.A radio frequency plasma source was used to generate highly active nitrogen atoms for doping N 3-ions into the SNO films during the synthesis (see Supplemental Material, experimental section).Figure 1a shows a comparison of XRD -2 scans from SNO and SNON films.The clear thickness fringes and sharp diffraction peaks indicate the high-quality epitaxial growth for both films.We calculated the out-of-plane lattice constants of SNO and SNON films to be (4.06 ± 0.02) Å and (4.14 ± 0.02) Å, respectively (inset of Figure 1a).The films elongate along the [001] orientation by 2% after N doping.We performed the optical second harmonic generation (SHG) polarimetry measurements on both samples to determine the symmetry variation after N doping.Figures 1b and 1c show  − 2 and  − 2 SHG signals for SNO (blue curves) and SNON (red curves) thin films, respectively.The experimental results were fitted by different models.The best fits to the data were given by the point group symmetry of mm2 for the undoped SNO films and 4mm for the SNON films, respectively.These results indicate that N incorporation improves the spatial symmetry of the lattice.Furthermore, we performed scanning transmission electron microscopy (STEM) measurements on a SNON film to illustrate the microscopic structure and chemical distribution.A cross-sectional high-angle annular darkfield (HAADF) STEM image indicates a chemically sharp and coherent interface between SNON film and STO substrate (Figure 1d and Supplementary Figure S1, [26]).The selected area electron diffraction (SAED) pattern from a SNON film along the pseudocubic [100] zone axis was shown in the inset of Figure 1d.Spatially resolved electron energy loss spectroscopy (EELS) mapping was performed at Nb L-, Ti K-, N K-, and O K-edges, respectively, as shown in Figures 1e-1h.The chemical analysis suggests that the N element is uniformly and randomly distributed within the SNON film.To assess the N content and its distribution in the SNON films, we further performed Time-of-Flight Secondary Ion Mass Spectrometry (ToF-SIMS) measurements (Supplementary Figure S2).Using the NO -signal as a surrogate for N, the depth profiles reveal both a N content enrichment and uniform distribution within the SNON film.
Electronic states of SNO and SNON films.The electronic states of SNO and SNON films were examined by X-ray absorption spectroscopy (XAS) and X-ray photoelectron spectroscopy (XPS) at room temperature.Figures 2a and 2b present XAS spectra at N K-edges and XPS spectra of Nb 3p from SNO and SNON films, respectively.The XAS spectra are normalized to their pre-and post-edges for direct comparison.As shown in Figure 2a, the N Kedge XAS of SNON films exhibit three main features positioned at 389.1, 403.0, and 409.5 eV, respectively.Similarly, XPS results in Figure 2b reveal a distinct N 1s peak located at ≈ 396.5 eV for SNON films.Based on previous studies on N-doped perovskite oxides [27,28], this N 1s peak is attributed to the N chemical state associated with N-Nb bonds.The formation of these N-Nb bonds arises from the substitution of nitrogen atoms in lattice oxygen sites.
Additionally, we found that the peak shapes of both Nb 3p (Figure 2b) and 3d (Supplementary Figure S3c) become broadener after N doping, indicating a significant change in the valence state of Nb ions.To quantitatively assess the change in Nb valence resulting from N doping, we performed a fitting analysis on the Nb 3d XPS spectra with a Shirley background and three sets of spin-orbit doublets, as shown in Figures 2c and 2d.Specifically, the area ratio between the 3d3/2 and 3d5/2 peak pairs was held constant at 2:3, while the spin-orbit splitting energy difference was fixed at 2.75 eV.Based on previous studies [29][30][31][32], the intense doublet (purple) at higher binding energy, ~206.8 eV, is assigned to Nb 5+ , while the other two doublets located at lower binding energies, ~206.0 eV and ~204.4 eV, correspond to the Nb 4+ and Nb 2+ states, respectively.In the case of stoichiometric SNO, the Nb valence state should be Nb 4+ .The emergence of the Nb 5+ state can likely be attributed to surface over-oxidation, while the Nb 2+ state may be associated with the presence of a small amount of oxygen vacancies.Notably, the Nb 5+ content significantly drops while the Nb 4+ content increases after N doping, suggesting that N doping could prevent the surface over-oxidation and make SNO more stable.

Electrical transport properties of SNO and SNON films and their strain dependency.
Transport properties of as-grown films were further examined using standard van der Paw geometry.Figure 3a shows the temperature dependent resistivity () of SNO and SNON films.
For thin films grown on STO substrates, the SNO films undergo a clear transition from metallic to semiconducting behavior after N incorporation.The  of SNO films increases by three orders of magnitude at room temperature and further enhances to seven orders of magnitude at 10 K after N doping.In addition, Hall measurements confirmed that the majority carriers for both samples are electrons.We plotted the carrier concentration (n) as a function of temperature for both films.At room temperature, the n of SNO films is ~10 22 cm -3 , consistent with the calculated value based on a d 1 electron configuration of n theory ≈ 1.53×10 22 cm -3 [4].After N doping, the n of SNON films decreases by an order of magnitude compared to that of SNO films.This reduction in n becomes even more obvious, reaching two orders of magnitude at low temperatures.A similar trend was observed in carrier mobility (μ) (Supplementary Figure S5), suggesting the effective modulation of electronic behavior by N doping.
We further investigated the influence of N doping on the electrical transport properties of SNO films at different strain states.The SNO and SNON films were epitaxially grown on (001)oriented KTaO3 (KTO,  ≈ -0.85%) and LaAlO3 (LAO,  ≈ -5.79%) substrates, together with STO ( ≈ -2.93%).Here,  represents the in-plane misfit strain at the interface and is defined by  = (as-af)/as, where as and af are the bulk lattice parameters of the substrate and SNO film.
The structure characterizations of these strained films were shown in Supplementary Figure S4.
All samples show a significant elongation along the out-of-plane direction after N doping.
Although the c/a ratio of SNO films reduces dramatically as increasing the compressive strain, the c/a ratio of SNON films exhibits a slight unexpected increase.This abnormal behavior is attributed to the strain relaxation at a high level of compressive strain.The SNO films grown on all these three substrates exhibit a metallic phase across all temperatures we measured.
However, when subjected to either a heavy compressive strain or a tensile strain, the roomtemperature resistivity increases by an order of magnitude.Similarly, the SNON films on STO and LAO become an insulator at room temperature and show a reduced n, whereas the SNON films on KTO undergo an insulator-to-metal-to-insulator transition as decreasing temperature.
Band structure and first principles calculations.To gain deeper insights into the impact of N doping on the electronic structure of SNO films, we show the occupied (Figure 4a) and unoccupied (Figure 4b) electronic states around Fermi level (EF) for both SNO and SNON films, which were taken from XPS valence band (VB) spectra and O K-edge XAS, respectively.As shown in Figure 4a, a clear finite density of states (DOS) is observed across EF for SNO, confirming its metallic behavior (Figure 3).The VB spectrum of SNO, consistent with previous report [31], consists of three features (labeled as A, B, and C) around EF.According to previous studies [31,32], the occupied states close to EF (feature A) are assigned to Nb 4d states, while to hybridized O 2p-Sr 4d (Sr-O bond), refer to peaks in a green background [35].Due to the crystal-field interactions, the four shape-resonances reflect the splitting of main transitions (i) (531.5 and 532.6 eV) and (ii) (536.0 and 537.7 eV) into t2g and eg levels [36].First of all, the peak intensity of (i) (Nb-O bond) in the SNON film is lower than that in the SNO film, indicating a weaker hybridization between O 2p and Nb 4d states after the N substitution [37], in consistent with XAS O K-edge results.Secondly, the increment of eg peak intensity at 532.6 eV and simultaneous the reduction of t2g peak intensity at 531.5 eV imply the increased number of Nb 4+ (4d 1 ) ions in SNON films.
To corroborate the origin of the features observed in XPS VB as well as the changes in O K pre-edge XAS and the corresponding evolution of electronic structure in SNO, we turn to density functional theory (DFT) simulations.Prior to simulating the scenarios that involves Ndoping, we first examined the fundamental properties of the undoped bulk SNO.This includes assessing lattice constants and the magnetic moment of the Nb ion, with a specific focus on their relationship to the effective Hubbard parameter Ueff (= U -J), where U is the Hubbard parameter and J is the exchange interaction.As depicted in Figure S8a, the lattice constants are optimally adjusted at Ueff = 0 eV, aligning closely with experimental values [38].The magnetic moment at Ueff = 0 eV is effectively quenched, in agreement with its non-magnetic fact [39].
Thus, Ueff = 0 eV is chosen as the default value for the subsequent calculations.Moving forward, we explore the influence of N doping.Initially, we consider a low N concentration of 1/12 (i. e.  = 1/4), using both the virtual crystal approximation (VCA) method and the substitution of one O atom by one N atom in one unit cell.In the latter case, we examine both equatorial and apical O-sites.For comparison, the DOS of these two cases are present in Figures 4c and 4d, revealing remarkably similar DOS profile although different methods were employed during the simulations.Consequently, it is reasonable to use VAC method in other doping cases, mitigating the need to consider an excessive number of N atom doping configurations.In the VCA method, the SNON system can be effectively simulated with varying N concentration by adjusting the relative weights of N and O atoms.In fact, our experiments did not find any specific ordered arrangement of doped N ions in SNON.Thus, the VCA method is even better in the DFT calculations, than the N-ordered state.The DOS for varying N concentrations are present in Figures 4e-4i.As the N doping concentration increases, there is a gradual upward shift in the conduction bands originating from Nb 4d orbitals.In Figure 4j, we illustrate the band diagram schematics for SNO, SNON1 (VCA, 1/6) and SNON2 (VCA, 1/3), with increasing the N doping level.When the N concentration reaches 1/3 (i.e.  = 1), a complete insulating state emerges, resulting in full Nb 5+ (4d 0 ).Further evidence of the metal-insulator transition can be gleaned from the band structures displayed in Figure S9.

Discussions and conclusions
In summary, we have unveiled striking structural and transport phase transitions within an semimetallic oxide via effective N-doping.These changes, characterized by substantial alterations in electronic states and carrier density, culminate in a compelling metal-to-insulator transition.This experimental result closely aligns with the predictions made in our theoretical models.Consequently, our study not only offers profound insights into the practical achievement of distinct physical ground states in oxide semimetals through anionic manipulation but also opens up a promising avenue for further exploration and innovation in the realm of advanced materials and condensed matter physics.This work not only deepens our understanding of the intricate interplay between structure and electronic properties but also paves the way for transformative applications in emerging technologies.

Synthesis of nitrogen doped thin films
The SrNbO3 (SNO) and SrNbO3-δNδ (SNON) films were grown on (001)-oriented KTO, STO and LAO substrates (Hefei Kejing Materials Technology Co. Ltd) using pulsed laser deposition (PLD).The ceramic target was synthesized by sintering mixtures of stoichiometric amounts of SrCO3 and Nb2O5 powder.The heating process was performed at 20 MPa and 1100 °C for 10 hours.Then, the recovered powder was sintered as a target at 20 MPa and 1000 °C for 12 hours.
The undoped SNO films were fabricated in vacuum at the substrate's temperature of 750 ℃.
The laser furnace was 0.52-0.65J cm -2 , and the laser repetition was 2 Hz.For the N-doped SNON films, the highly active nitrogen atoms were in-situ doped into the SNO films during the deposition using RF plasma generated atomic nitrogen under the same experimental conditions.
The input power was 250 W and the nitrogen flow was maintained 1.5 sccm/min, keeping the partial pressure of N2 gas at 10 -4 Torr.The plasma source was equipped with a parallel plate capacitor to remove ionic specie.The films were cooled down to room-temperature under the irradiation of nitrogen plasma in order to compensate the nitrogen vacancies (NVs).The film thickness was controlled by counting the number of laser pulses and further confirmed by Xray reflectivity (XRR) measurements.

Structural characterization and elemental mapping
Synchrotron X-ray diffraction (sXRD) θ-2θ scans were conducted at the beamline 1W1A of the Beijing Synchrotron Radiation Facility (BSRF).Reciprocal space mapping (RSM) and XRR measurements were carried out using a PANalytical X'Pert 3

Second harmonic generation (SHG) measurements
Optical SHG measurements of SNO and SNON films on STO substrates were performed in the reflection geometry at room temperature.The pumping beam is a Ti: sapphire femtosecond laser (Tsunami 3941-X1BB, Spectra-Physics, =800 nm).The linearly polarized light incidents on the sample at the angle of 45°.The polarization direction () of the incident field (E) was rotated through a half-wave (/2) plate.Based on different polarization combinations of incident and reflected light, two configurations were used to conduct the SHG measurement.

P-out (𝐼 𝑝−𝑜𝑢𝑡 2𝜔
) represents the analyzer polarization parallel to the plane of incidence and incident light polarization being rotated, and s-out (  −

2𝜔
) as the analyzer polarization perpendicular to the plane of incidence and incident light polarization being rotated [40,41].
The optical signals were detected by a photon multiplier tube.Theoretical fittings of the SHG polarimetry data were performed with analytical models using standard point group symmetries.

Electronic state characterization and elemental content
X-ray absorption spectroscopy (XAS) measurements were conducted for both N Kand O Kedges at beamline 4B9B of the BSRF.The incident direction of polarized light is parallel to the surface normal.Spectra were collected in the total electron yield mode at ambient temperature.
Room-temperature X-ray photoelectron spectroscopy (XPS) measurements were performed at Pacific Northwest National Laboratory (PNNL).Spectra were measured using an electron flood gun to compensate the positive photoemission charge because SNON films were not sufficiently conductive, and conductive SNO films were not grounded.In order to easily visualize the change in Nb 3p and 3d line shapes with N-doping, we align all XPS spectra to place the corresponding O 1s peaks at 530.0 eV.A small polycrystalline Au foil was affixed to the corner of each film surface using Cu tape.For VB spectra, the Au 4f7/2 peak was used to calibrate the binding-energy scale.The distribution of element valence states was obtained by fitting Nb 3d XPS using Casa XPS software.ToF-SIMS measurements were carried out using a ToF-SIMS V (ION-TOF GmbH, Münster, Germany) at PNNL.The mass spectrometer was equipped with a reflection type time-of-flight analyzer.A dual-beam depth profiling strategy was employed, in which a 1.0 keV Cs + beam (~40 nA, 300 µm × 300 µm scanning area) was used for sputtering and a 25 keV Bi + beam (~1.1 pA, 100 µm × 100 µm scanning area within the center of the Cs+ crater) was used for negative spectra data collection.Additionally, a flood gun (~1 µA) was used for charge compensation.The film/substrate interface was determined via the secondary ion signals of NbO -and 49 TiO2 -.

Electrical Transport Measurements
The transport properties were performed by using a 9T-PPMS.All samples were measured in a standard van der Pauw geometry with a four-probe method (to eliminate contact resistance).
The contacts were prepared using wire-bonding to ensure that all interfaces were electrically well connected.

First-Principles Calculations
First-principles density functional theory (DFT) calculations are performed using the Vienna ab initio Simulation Package (VASP) [42,43] based on the projected augmented wave pseudopotentials.For the exchange-correlation functional, the PBEsol (Perdew-Burke-Ernzerhof revised for solids) [44] parametrization of the generalized gradient approximation (GGA) [45][46][47] method is used .The plane-wave cutoff is set to 500 eV and the Dudarev approach [48] is adopted when Hubbard U is imposed on Nb ion.The atomic positions and lattice constants are fully optimized until the Hellman-Feynman forces converged to less than 0.01 eV/Å.The virtual crystal approximation (VCA) method is used in our calculations which tunes the weight of N and O atoms in one site to simulate the uniform doping [49].The N signal in SNON is more than two orders of magnitude larger than that in SNO, implying that the N ions were sufficiently doped.The N content in SNO exceeds the natural abundance, possibly attributed to the N introduced during the preparation of the SNO PLD target.The dash lines marked with a0, b0, and c0 represent the experimental values [37].

Figures and figure captions
feature B is mostly formed by O 2p bands and featured C is assigned to Nb 4d/O 2p and Nb 4p/ O 2p hybridized bands, respectively.After N doping, the vanishing DOS across EF matches well with the insulating behavior of SNON films.Moreover, a broad shoulder (labeled as feature D) located at the top of feature B appears.Additionally, O K-edge XAS of SNO and SNON films show a consistent change in the spectral line shape and peak intensity, as shown in Figure 4b.The characteristic features at the low-energy region reflect the bonding strength between transition metal cations and ions [34].The shape resonances in the low-energy region come from two main transitions: (i) from O 1s to hybridized O 2p-Nb 4d (Nb-O bond), see peaks in a purple background, and (ii) from O 1s MRD diffractometer with Cu Kα1radiation equipped with a 3D pixel detector.The thicknesses of films and X-ray scattering length densities (SLD) were obtained by fitting XRR curves using GenX software.Crosssectional TEM specimen of the SNON/STO film was prepared using the standard focused ion beam (FIB) lift-off process.High-angle annular dark-field (HAADF) images and the selected area electron diffraction (SAED) patterns were taken along a pseudo-cubic [100]pc zone axis using the JEM ARM 200CF microscopy at the Institute of Physics, Chinese Academy of Sciences.Elemental-specific electron-energy-loss-spectroscopy (EELS) and energy dispersive x-ray (EDX) spectroscopy mappings were obtained by integrating the Nb L-, Ti K-, N Kand O K-edges signals from selected regions after background subtracting.All data were analyzed using Gatan DigitalMicrograph software.

Figure 2 .
Figure 2. Electronic states of SNO and SNON films.(a) N K-edges XAS spectra for the SNO and SNON films.(b) N 1s and Nb 3p XPS spectra for SNO and SNON films.(c) and (d) Measured (open symbols) and fitting (solid lines) results of Nd 3d XPS spectra for SNO and SNON films, respectively.The purple, green, and yellow curves are Lorentzian functions fit to the raw spectrum, black curve is the Shirley background, and the blue and red curve are the sum of the individual Lorentzian functions.The area ratio for Nb 5+ , Nb 4+ , and Nb 2+ is 83.0%

Figure 3 .
Figure 3. Strain-dependent transport properties of SNO and SNON films.(a) ρ-T curves of SNO and SNON films grown on KTO, STO and LAO substrates.The temperature dependent carrier density (n) of SNO and SNON films on three substrates were shown in (b)-(d).n decreased by 1-2 orders of magnitude after N doping.Strain dependent (e) c/a ratio and (f) activation energy (ΔE) of SNO and SNON films.The impact of the compressive strain on the c/a ratio and ΔE increases significantly.

Figure 4 .
Figure 4.The evolution of electronic structures upon N doping.(a) Valence band spectra for SNO (blue) and SNON (red) films.The spectra were shifted to reflect the direct comparison.(b) O K-edges XAS spectra for the SNO and SNON films.(c-d) Comparison of DFT DOS for the 1/12 doping case with different methods, which are very similar.(e-i) DFT DOS's around the Fermi level obtained with VCA method for SNO and SNON with N dopant from 0 to 1/3.(j) Band diagram schematics of SNO, SNON1 (VCA, 1/6) and SNON2 (VCA, 1/3), with increasing N doping level.

Figure S1 .
Figure S1.STEM image and EDX maps of a 23 nm-thick SNON film.(a) Low-magnified HADDF-STEM image of a SNON film grown on a STO substrate.The STEM results indicate a sharp interface between SNON films and STO substrates.The SNON films show a high crystallinity within the observed region.The colored panels show the integrated EDX intensities of (b) Sr K-, (c) Nb L-, (d) Ti K-, (e) O K-, and (f) N K-edges.EDS results suggest the elemental distribution is uniform and the SNON/STO interface does not show apparent chemicalintermixing.The O content in SNON films is less than that in the STO substrates.We notice that the N content in SNON films is significantly higher than the noise level that present in the STO substrates and Pt coatings.

Figure
Figure S2.X-ray reflectometry (XRR) and SIMS measurements on SNO and SNON films.
(a) and (b) XRR curves of SNO and SNON films, respectively.The solid lines are the best fittings to the experimental data (open symbols).(c) and (d) X-ray scattering length densities (SLDs) of SNO and SNON as a function of distance, respectively.At the oxide interfaces, there are unavoidable structural transition and dislocations due to the misfit strain, resulting in a reduction of SLDs at the interfaces (shadow area in both panels).(e) and (f) SIMS results of SNO and SNON films on STO substrates.The interface locations (marked by grey dashed lines) of SNO/STO and SNON/STO were confirmed by the secondary ion signals of NbO -and 49 TiO2 -.

Figure
Figure S3.N-doping increases the Nb valence state.(a) Survey, (b) O 1s, (c) Nb 3d XPS spectra for the SNO (blue) and SNON (red) films.Apparently, a tremendous N 1s signal has been detected after N doping.Compared with SNO, the C 1s signal (a) of SNON displays reduced intensity, resulting in a diminished intensity of the CO3 feature in the O 1s spectrum (b).The existence of Nb 5+ in the SNO film may potentially arise from surface oxidationprocesses.We find that the peak intensities of Nb 4+ increased as well as Nb 3+ decreased after N inserting, while the Nb 5+ contribution remains relatively unchanged, which is in consistent with the hole doping.

Figure S4 .
Figure S4.Structure characterizations of strained SNO and SNON films.The SNO and SNON films were epitaxially grown on (001)-oriented KTO, STO and LAO substrates.(a) XRD θ-2θ scans for SNO and SNON films around the substrates' 002 reflection peaks.The film's peaks of strained SNON films shifted to the low angle after N doping, indicating that the c-axis lattice constant increases after replacing partial O atoms into N atoms.(b)-(g) Reciprocal space maps (RSMs) around the substrates' pseudocubic 103 reflections for SNO and SNON films grown on different substrates.From RSMs, we could obtain both a and c lattice constants.As increasing the lattice mismatch, the SNO and SNON films relax their in-plane compression, resulting in a reduction of lattice constants.
(a)-(c)Temperature dependent µ for SNO and SNON films grown on KTO, STO, and LAO substrates, respectively.In most cases, µ reduces with N doping.With an exception, µ of SNON/KTO is slightly larger than that of SNO/KTO at low temperatures.

Figure S6 .
Figure S6.Fitting curves for transport properties.(a)-(c) Fitted curves of ρ-(1000/T) for SNO and SNON films grown on KTO, STO and LAO substrates, respectively.All curves fit to

Figure S7 .
Figure S7.Transport properties of SNO and SNON films grown on (001)-, (110)-, and (111)-oriented KTO substrates.(a) ρ-T curves of SNO and SNON films.Films grown on different oriented substrates show the same trend: The film gradually changes from the conductivity of SNO films to the insulation of SNON films.(b) and (c) Room-temperature n and µ for SNO and SNON films, respectively.

Figure S8 .
Figure S8.The tests of Ueff value on the lattice constants and magnetic moment of Nb ion.

Figure S9 .
Figure S9.Band structures of SNON with different N doping levels.The band structure of SNON changes systematically with N concentrations.Apparently, the conduction bands gradually move upwards across the Fermi level (EF).When the doping level increases to 1/3, the SNON becomes an insulator.