Room-Temperature Ferromagnetic Enhancement and Crossover of

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Room Temperature Ferromagnetic Enhancement and Crossover of Negative Magnetoresistance to Positive Magnetoresistance in N-Doped In2O3 Films Luhang Shen, Yukai An, Dandan Cao, Zhonghua Wu, and Jiwen Liu J. Phys. Chem. C, Just Accepted Manuscript • DOI: 10.1021/acs.jpcc.7b08732 • Publication Date (Web): 07 Nov 2017 Downloaded from http://pubs.acs.org on November 14, 2017

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The Journal of Physical Chemistry

Room Temperature Ferromagnetic Enhancement and Crossover of Negative Magnetoresistance to Positive Magnetoresistance in N-Doped In2O3 Films Luhang Shen†, Yukai An*†, Dandan Cao†, Zhonghua Wu‡, Jiwen Liu† †

Key Laboratory of Display Materials and Photoelectric Devices, Ministry of

Education; Tianjin Key Laboratory for Photoelectric Materials and Devices; National Demonstration Center for Experimental Function Materials Education; School of Material Science and Engineering, Tianjin University of Technology, Tianjin 300384, China ‡

Beijing Synchrotron Radiation Facility (BSRF), Institute of High Energy Physics,

Chinese Academy of Sciences, Beijing 100049, China

Abstract The effects of N-induced acceptor defects on tuning optical, transport and magnetic properties of the In2O3 films fabricated by magnetron sputtering technique were investigated systematically by X-ray diffraction (XRD), X-ray photoelectron spectroscopy (XPS), UV-Vis, Hall effect, ρ-T, MR and magnetic measurements. The detailed structural analyses reveal that the N-doped In2O3 films have cubic bixbyite structure with the substitutional N defect at the O sites of In2O3 lattice. The N-doped In2O3 films display clear room-temperature (RT) ferromagnetic behavior and Mott variable range hopping (VRH) transport behavior. With the increase of N doping concentration, the saturated magnetization (Ms) of the films monotonically increases and the conductivity transforms into p-type. Crossover of negative magnetoresistance to positive magnetoresistance and the red shift of optical band gap Eg are also 1

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observed with N doping. The first principles calculations show that the localized holes induced by N doping can mediate the magnetic interaction by short-ranged N1:pIn:d/p-N2:p hybridization in N-doped In2O3 system. Therefore, the intrinsic ferromagnetic ordering in the N-doped In2O3 films can be attributed to the p-p interaction between N 2p orbital, which causes large Zeeman-split effect to suppress the carrier’s hopping path, leading to form the positive MR. Key word: In2O3; N doping; Ferromagnetism; Magnetoresistence;

1. INTRODUCTION Diluted magnetic semiconductors (DMSs), a class of magnetic semiconductors, have attracted intensive scientific interest due to their unique physical properties and potential applications in the emerging field of next-generation spin-based magnetic devices.1-3 Although some magnetic devices based giant magnetoresistance (GMR) effects have been successfully realized, the global development of DMS materials is still rather important for spintronices. The ideal DMSs should exhibit ferromagnetic ordering at or above room temperature (RT) for practical applications and have a homogeneous distribution of magnetic dopants. Since Dietl et al. theoretically predicted

that

RT

ferromagnetism

might

exist

in

wide-band-gap

oxide

semiconductors,4 In2O3 doped with transition metal (TM) elements has received considerable attention due to its integration of optical, electronic and magnetic properties into one single material.5-9 However, many studies show that the RT ferromagnetism in TM-doped In2O3 may arise from the precipitation of magnetic clusters or oxide secondary magnetic phases. For examples, Jiang et al. reported RT ferromagnetism in Fe/Sn-codoped In2O3 powders, but is directly related to the precipitated Fe3O4 impurity.10 Ohno et al. reported that the nanoscale Fe clusters result in the observed RT ferromagnetism for Fe-doped ITO films.11 Kohiki et al. 2


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reported that the ferromagnetism in Fe-doped In2O3 bulk ceramics originates from a formation of ferromagnetic Fe2O3 clusters.12 These extrinsic magnetic behaviors are undesirable for practical applications. Since Pan et al. observed RT ferromagnetism in C-doped ZnO,13 the nonmagnetic element doping (such as C and N) has triggered extensive study, which can effectively avoid the extrinsic ferromagnetism due to the magnetic precipitations associated with the magnetic dopants. The magnetic properties induced by nonmagnetic elements are also different from the traditional 3d or 4f elements induced ferromagnetism. The magnetic interaction comes from the p-p orbital coupling rather than s-p,d exchange coupling. Therefore, many researchers have considered utilizing these nonmagnetic elements in engineering the magnetic interaction or enhancing RT ferromagnetism in oxide based DMSs. Recently, doping a nonmagnetic N element in some oxide semiconductors has been recognized as an effective way to achieve RT ferromagnetic ordering. However, the origin of ferromagnetism still remains controversial and uncertain. Yu et al. deposited the RT ferromagnetic N-doped ZnO films, and the ferromagnetism is attributed to the bound magnetic polaron (BMP) mechanism based on oxygen vacancies.14 However, Wu et al. reported that RT ferromagnetism is directly associated with the intrinsic defects: Zn vacancies in N-doped ZnO films.15 Sun et al. suggested that the RT ferromagnetism in N doped In2O3 films is associated with the N incorporation and can be mediated by the indirect ferromagnetic coupling between the N 2‫ ݌‬states.16 Bao et al. reported RT ferromagnetism in N-doped TiO2 films, which is ascribed to the O substitution with N.17 Theoretically, a localized hole-mediated ferromagnetism through a ‫݌‬-‫ ݌‬interaction between the 2p states of O and N was consistently suggested in N-doped ZnO, TiO2 and In2O3 systems.18-20 The similar hole-mediated magnetic mechanism was also found in nonmagnetic C-doped ZnO and In2O3.21-23 To the best 3



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of our knowledge, up to now, limited reports are available in the N-doped In2O3 systems by experimental insights. In this letter, we introduce N in In2O3 lattice to fabricate nonmagnetic element doped DMSs by magnetron sputtering technique. Systematical experimental and theoretical studies on the structural, optical, magnetic and transport properties of N-doped In2O3 films were preformed. It was found that the 5% N-doped In2O3 film has strongest RT ferromagnetism and p-type conductivity. The variation of N doping concentration leads to the crossover of negative magnetoresistance to positive magnetoresistance. The ferromagnetism can be attributed to the p-p exchange coupling by unpaired electrons on substitutional dopant N atoms. These results give new sights for understanding the mechanism of magnetic interactions in N-doped In2O3. They also prove that the nonmagnetic element N can be employed to induce RT ferromagnetism and remarkably tune the magnetic and transport properties by the p-p exchange interaction.

2. EXPERIMENTAL DETAILS The pure In2O3 and N-doped In2O3 films with a thickness of about 400 nm were fabricated on (001) Si and ultra-white glass substrates by RF-magnetron sputtering technique. The sputter chamber was evacuated to a base pressure of 8×10-5 Pa before the deposition. All the films were sputtered in 0.8 Pa of pure Ar, O2 and N2 mixed gases (purity 99.999%) at a substrate temperature of 400℃. By adjusting the flow rate of N2, the 2 at% and 5 at% N-doped In2O3 films were prepared. The N concentration in the films was determined by inductively coupled-plasma (ICP) emission spectra. The crystal structures of the films were examined by θ/2θ X-ray diffraction (XRD) with Cu Κα radiation (λ=0.15406nm). The surface morphology of films was determined by Scanning electron microscopy (SEM) of JOEL 6700F. The valence state of the films was confirmed by X-ray photoelectron spectroscopy (XPS), 4


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which was carried out on a PHI-1600 photoelectric spectrometer with a focused monochromatic Mg Kα (hν=1253.6eV). The optical transmission properties were carried out using TU-1901 double beam UV-visible spectrophotometer with scanning wavelength from 200nm to 800nm. The film resistivity (ρ), magnetoresistence (MR) and Hall effect measurements were performed by conventional the van der Pauw fourprobe configuration with a Closed Cycle Cryo-cooler with the applied magnetic fields up to 11 kOe and temperature range of 10-300K. The magnetic measurements were carried out by a Quantum Design superconducting quantum interference device (SQUID) magnetometer as a function of magnetic field (0 to ±15 kOe). The first principles calculations on the total energy and electronic structure are carried out in a supercell of In2O3 with cubic bixbyite structure using the MedeAVASP code. The exchange and correlation functionals are treated under the generalized gradient approximation (GGA) of Perdew-Burke-Ernzerhof (PBE). The lattice constants are optimized and obtained as a=b=c=10.29Å, which are in good agreement with experimental values. Brillouin zone integrations are performed with the 2×2×2 centered K-point grid to make sure the total energy is relatively stable. The cutoff energy of plane waves and the maximum force are set to be 500 eV and 0.02V/Å, respectively. Meanwhile, self-consistency is considered not to be achieved until the difference in energy of a single atom between succeeding iterations is less than 1.0×10-5 eV/ atom.

3. RESULTS AND DISCUSSION Figure 1 shows the typical XRD patterns of the un-doped, 2 at% and 5 at% Ndoped In2O3 films. All the films show the same cubic bixbyite structure as pure In2O3 with the preferential orientation of (222). The N-doping does not lead to the appearance of any extra peaks from N-related secondary phases or other impurities 5



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within the XRD detection limit. These suggest that the structure of the N-doped In2O3 films still retains the cubic bixbyite phase belong to the space group I a 3 . In order to investigate the effects of N doping, a careful analysis of diffraction peaks is carried out, it is obvious that the prominent (222) diffraction peak shifts to lower 2θ angles with the increase of N concentration, as show in the inset of Fig.1, suggesting a dspacing expansion, namely a gradual increase in the average lattice parameter. It is possible that N ion (RN3- =1.32Å) has a larger atomic radius than that of O ion (RO2=1.24Å), such the expansion of lattice parameter can be caused by the incorporation of N ions into O sites of the In2O3 lattice.24,25 Moreover, it is also observed that the intensity of (222) diffraction peaks is decreased and the linewidth becomes broaden with N doping. The mean grain size, D, of the N-doped In2O3 films is estimated using the Debye-Scherrer equation: D = 0 . 9 λ , where λ is the wavelength of X-ray β cos θ

radiation (λ=0.15406nm), θ is the Bragg angle and β is the full width at half maximum. The corresponding mean grain size is calculated to be 37.8nm, 11.2nm and 10.9nm for the un-doped, 2 at% and 5 at% N-doped In2O3 films, respectively, suggesting that the mean grain size gradually decreases with the increase of N concentration in the films. This is also verified by the SEM mages as shown in Fig.2(a)-(c). One can see that all the films show uniform and dense morphology without visible voids and defects overall the surface. The average grain size and roughness of surface obviously decrease with N doping. It may due to that the N impurities can act as nucleation centers and result in decreasing the grain size. On the other hand, it is clear from Fig.1 that many diffraction peaks disappear with N-doping. It maybe due to that the N doping decreases the grain size and intensity of peaks, resulting in the disappearance of many weak diffraction peaks. 6


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In order to determine the valence state of the N-doped In2O3 film, XPS measurements were carried out and the binding energy scale was calibrated using the C1s binding energy at 284.6 eV. The core-level XPS spectra of In 3d, O 1s and N 1s for the 5 at% N-doped In2O3 film are shown in Figs. 3(a)-(c). The XPS spectrum of In 3d core level is symmetric, ruling out the existence of multiple components of In in the film. The observed energies for In 3d5/2 and 3d3/2 states are located at 444.5 eV and 452 eV, respectively, which match closely with the binding energy of In3+ in In2O3. The asymmetric O 1s spectrum can be deconvoluted into two characteristic peaks, I and II, located at 530.2 eV and 531.6 eV, respectively. The peak I on the lower binding energy side may be attributed to O surrounded by In atoms in the bixbyite structure of In2O3, namely lattice oxygen (In-O bond). The peak II on the higher binding energy side is usually associated with O2- ions existing in oxygen deficient regions in the In2O3 matrix.26 The N 1s XPS spectrum in the Fig.3(c) can be also well deconvoluted by two characteristic peaks. The N 1s speak located at 397eV is assigned to the substitution of N atom for O sublattice (NO) in In2O3. The other N 1s peak at about 399.1eV is usually related to a (NC)O and/or (N2)O complexes in In2O3.27 The carrier concentration and resistivity ρ were measured at room temperature by Hall effect and the corresponding experimental data were shown in Table 1. One can see that the semiconductivity of the films is strongly dependent on the N doping concentration. For the pure In2O3 film, the type of carriers is confirmed to be n-type, as expected for the film growth in an oxygen-deficient environment. It is clearly that the resistivity ρ increases and electron concentration nc decreases remarkably with N doping (2 at%). This can be explained due to the recombination of electrons and hole carriers. Further increasing the N concentration to 5 at%, the resistivity ρ becomes 7



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larger. However the conductivity transforms into p-type with a lower hole concentration of 1.51×1017 cm-3. N is a useful element for achieving p-type conductivity in the In2O3 films. For the low N doping concentration, the holes produced by the substitution of N3- ions for the O2- sites of In2O3 lattices can compensate the electrons arising from the oxygen vacancies, resulting in a decrease in the carrier concentration nc, then the scattering from more N ions may also contribute to the increase of the resistivity ρ. For the higher N doping concentration (5 at%), there will be net holes, which result in p-type conductivity. Wu et al. also observed ptype conductivity in Cu-doped SnO2 films, and the p-type conductivity is due to the acceptor states of Cu dopants which substitute Sn atoms in the SnO2 lattice.28 The electric parameters of the films are characterized by kFl, which is estimated using the formula: k F l = h (3π 2 )

2

3

1

(e 2 ρnc 3 ) , where h is the Planck constant, e is the electron

charge, ρ is the resistivity and nc is the carrier concentration.29 From table 1, it is obviously that the calculated values of kFl for the N-doped In2O3 films are much smaller than 1, suggesting that the semiconductivity of the films is in the strongly localized regime.30 Figure 4 shows the magnetization as a function of applied magnetic field at 300K using a SQUID magnetometer for the 2 at% and 5 at% N-doped In2O3 films. The diamagnetic contribution from the Si substrates has been subtracted from the raw data. The un-doped In2O3 film exhibits only diamagnetic behavior (not shown here). However, the N-doped In2O3 films show clear hysteresis loops, suggesting the existence of RT ferromagnetism. The saturation magnetization versus N doping concentration is ploted in the inset of Fig.4. The saturation magnetic moment (Ms) of the 5 at% N-doped In2O3 film is 0.6 emu/cm3, much larger than that (0.15 emu/cm3) of the 2 at% N-doped In2O3 film, suggesting that the further N doping can largely 8


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enhance RT ferromagnetism of the In2O3 films. As is well known, neither nitrogen nor any nitrides exhibited ferromagnetic signals. The absence of the secondary magnetic phases has ruled out the possibility of RT ferromagnetism due to the extrinsic origin, which is also consistent with the XRD analyses. However, the magnetic mechanism need be further clarified. In order to understand the magnetic mechanism and spin polarization of N doped In2O3, we further employed first principle calculations based density functional theory to investigate the electronic structures and magnetic interactions between N ions. The N-doped In2O3 system was modeled with a cubic bixbyite In2O3 supercell consisting of 32 In atoms and 48 O atoms. The supercell size is large enough to allow us to investigate different N doping configurations. In the supercell, O atoms were substituted by N atoms, and were considered by different doping configurations: (1) two N atoms substitute for the nearest neighboring O atoms in the In2O3 lattice, namely 2N-close; (2) two N atoms substitute for the O atoms in a far distance, namely 2N-far; (3) two N atoms substitute for the nearest neighboring O atoms and an oxygen vacancy (VO) is introduced by removing one oxygen atom adjacent to neighboring N atom in In2O3 lattice, namely VO-2N-close. In order to compare the relative stability of the N-doped In2O3 systems, we calculated the formation energy Ef of three doping configurations. For the (1) and (2) configurations Ef=E1(doped)-E(pure)-2µN + 2µO

(1)

for the (3) configuration Ef=E2(doped)-E(pure)-2µN + 3µO

(2)

in which E1(doped) and E2(doped) correspond to the total energy of In2O3 supercell containing N impurities, N impurities and oxygen vacancy, respectively. E(pure) is the total energy for pure In2O3 supercell, whereas µN and µO represent the chemical 9



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potential of the N and O atoms, respectively. The chemical potential depends on growth conditions, which may be either In- or O-rich. µIn and µO obey the equation 2µIn + 3µO =µ(In2O3). Under In-rich growth conditions, the chemical potential of In is assumed to be the energy of one atom in bulk In (µIn=µInmetal) and µO can be calculated by the above formula. Under O-rich growth conditions, the chemical potential of O is related to molecular oxygen (µO=µ(O2)/2). The chemical potential µN was estimated by consideration of one N2 molecule (µN=µ(N2)/2). The calculated formation energy Ef for three doping configurations is summarized in Table 2. It can be seen that under both In-rich and O-rich growth conditions, the (1) configuration, namely 2N-close, possesses the lowest formation energy, indicating that two N atoms substitute the nearest neighbor oxygen atoms is energetically favorable in the In2O3 lattice. Both FM and AFM couplings between two N atoms for the (1)-(3) configurations have been also calculated. The magnetization energy difference △EAFM-FM between the FM and AFM spin orderings for the (1)-(3) configurations is 63.5meV, 0.81meV and 1.3meV, respectively, indicating that the FM state is the ground magnetic ordering. The magnitude of △EAFM-FM (63.5meV) is larger than some other systems exhibiting RT ferromagnetism, such as N doped ZnO and C doped ZnO,13,31 suggesting that the observed RT ferromagnetic ordering in N-doped In2O3 system is expected. Figs.5(a)-(c) show the density of states (DOSs) of different N doping configurations. The DOS curves of spin-polarized N, O and In atoms for the (1) “closed” configuration, namely two N atoms substitute for the nearest neighboring O atoms, indicate that the N 2p states significantly overlap with the O 2p and In 5p and 4d states around the Fermi level (EF), forming two obvious N 2p hole acceptor levels about 0.16eV and 0.55eV above EF, respectively, as shown in Fig.5(a). The hybridization between N dopant and its neighboring atoms leads to the splitting of the 10


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energy levels near the EF and forms a couple chain of N1:p-In:d/p-N2:p, which contribute to an indirect FM coupling between N dopants mediated by the holes introduced. Our calculated results are similar with the theory predicted by Peng et al. and Pan et al. that holes mediate the ferromagnetic ordering by the p-p interaction in C-doped ZnO.13,14 The hybridization introduces a large total magnetic moment of 2.0µB, which is mainly contributed by the N 2p orbital, as can be observed by the spin density distribution in Fig.5(e). This is also consistent with the results of N doped ZnO and C doped TiO2.31,32 Fig.5(b) shows the DOS curves for the (2) configuration, namely the two N atoms are larger separations. The N 2p states are still significantly overlap with the O 2p and In 5p and 4d states, while the hole states are obvious less than that of the (1) “closed” configuration, as show in Fig.5(b). These indicate that larger N-N separations do not favor to form a coupling chain of N1:p-In:d/p-N2:p and the indirect FM coupling between N dopants is also rather weak. The magnetic moment is also mainly contributed by the two N 2p orbitals, as indicated by the spin density distribution from Fig.5(f). Therefore, it can be proposed that the localized holes induced by N doping mediate the magnetic interaction by short-ranged N1:pIn:d/p-N2:p hybridization in N-doped In2O3 systems. However, for the (3) configuration with one oxygen vacancy, it is obvious that the N 2p hole acceptor levels almost completely disappear, but the DOS result still indicates a p-type character near the EF. The hybridization of N 2p states with the O 2p and In 5p and 4d states only leads to the weak splitting of the energy levels and introduces a very low magnetic moment of 0.25µB, as shown in Fig.5(g). To further check the effect of oxygen vacancy, the N doped system with two oxygen vacancies was also calculated, as shown in Fig.5(d). It was found that the configuration (2VO-2N-close) does not

11



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show a magnetic moment, and the p-type character also completely disappears, which is consistent with the n-type conductivity for the 2 at.% N doped film. Fig. 6(a) shows the UV-Visible transmission spectra of the un-doped, 2 at% and 5 at% N-doped In2O3 films deposited on ultra-white glass substrates. The transmittance spectra for the various N-doped In2O3 films show clear wave forms, which is the characteristic of the interference of light. All the films show the optical transparency about 75-90% in the visible light range. The N doping changes the optical properties of the films. It increases the absorption edge towards higher wavelength, indicating an obvious red shift for the absorption band. The optical transmission intensity slightly decreases with N doping, which may be due to the Ninduced acceptor defects enhancing the light scattering on the transparency of the films. The optical band gap Eg can be determined by measuring absorption coefficient (α) given by the following equation: α=ln(1/T)/d, where T is the transmittance and d is the film thickness.33 The optical band gap draw on the absorption coefficient (α) which is defined by the equation: (αhν)=A(hν-Eg)1/2, where A and Eg are constant and optical band gap, respectively. It can be calculated by plotting (αhν)2 versus hν and extrapolating the linear portion of the plot to (αhν)2=0, as shown in Fig. 6(b). The variation of Eg for the N-doped In2O3 films is shown in the inset of Fig. 6(b).One can see that the un-doped In2O3 film has an optical band gap of about 3.66 eV, which is slightly lower than the usually cited value of 3.75 eV. With increasing N concentration, the Eg value gradually decreases from 3.66 eV to 3.35 eV, implying that the incorporation of N ions in the In2O3 lattice plays an important role in tuning the band gap of In2O3. Generally, a passivated impurity band will occur in the bandgap due to the existence of compensated donor-acceptor complexes, leading to a red shift of optical band gap.34 In our N-doped In2O3, the similar band engineering can 12


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be obtained by introducing oxygen vacancies and N impurity in the In2O3 matrix. Obviously, after N doping, a narrowing of optical band gap Eg was observed with respect to the un-doped In2O3 film, indicating the presence of the N impurity band in the N-doped In2O3 films. The resistivity ρ was also measured as a function of temperature to understand the electrical conduction mechanism for the N-doped In2O3 films. Figs.7 (a)-(c) show the temperature dependence of the resistivity ρ(T) for the un-doped, 2 at% and 5 at% N-doped In2O3 films. All the films show the semiconducting transport behaviors characterized by the negative temperature coefficient in the whole temperature range. The conduction mechanism at low temperature cannot be described by thermal excitation of carriers from localized states to the conduction band. Therefore, the Mott variable-range-hopping (VRH) mechanism, namely ρ ～ exp[(T0/T)1/4], was proposed.35 Fig. 7(d) shows the plots of lnρ versus T−1/4 for the un-doped, 2 at% and 5 at% N-doped In2O3 films. Despite the varying carrier concentration nc, the almost linear relationship can be obtained, indicating that the Mott VRH mechanism plays an important role in the low temperature conductivity and the carriers can hop from one site to another with only a specific hopping probability. Magnetoresistance (MR) effect provides a new sight for understanding the mechanism of ferromagnetic order and for determining the contribution of magnetic exchange interaction in DMSs. Fig. 8 shows the normalized field dependence of MR behavior for the un-doped, 2 at% and 5 at% N-doped In2O3 films at 10K. The MR curves were measured with the applied magnetic field perpendicular to the films, defined as MR= ρ H ρ 0 . It is obvious that the un-doped In2O3 film shows the negative MR. Generally, the suppression effect of magnetic field on the interference between scattered electrons in the weak localization regime can cause the negative MR 13



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behavior.36 With the N doping (2 at%), the negative MR still dominates in the film, but the magnitude slightly decreases. Further increasing the N doping concentration to 5 at%, the MR becomes positive at the whole magnetic field range. As shown in Fig.4, the 5 at% N-doped In2O3 film shows the strongest Ms. The strong spin-split of band states caused by s-p exchange interactions can result in an obvious positive MR in DMSs.37,38

As aforementioned,

the

short-ranged

N1:p-In:d/p-N2:p

magnetic

interaction is the origin of RT ferromagnetic ordering, which is medicated by the localized holes induced by N doping. Therefore, the magnetic coupling between local hole carriers and N dopants induced p-p exchange can result in the Zeeman-split (spin-split) of band states. With the increase of external magnetic field, the Zeemansplit effect becomes remarkable and suppresses the hopping path of carriers. These lead to the increase of resistivity and form the positive MR behavior. For the 2 at % N-doped In2O3 film with n-type conductivity, the holes produced by N doping are compensated by the electrons arising from the oxygen vacancies. Therefore, the ferromagnetism meditated by the holes is weakened, leading to decrease the p-p exchange interaction. In this situation, the suppression effect of the applied magnetic field becomes dominate and the negative MR behavior still appears. It is obvious that the N doping has profound effects on the positive and negative contributions of MR. The crossover phenomenon from the negative MR to positive MR were also observed in Fe doped In2O3 films and Mn doped ZnO films, et al..38,39 These results suggest that the N doping induced RT ferromagnetism has similar behaviors with that induced by TM-doped oxide semiconductors. The variation of N2 flow rate during deposition effectively changes the N doped concentration in the In2O3 films. The detailed structural analyses suggest that a single phase with cubic bixbyite structure is observed and the doped N ions are 14


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substitutional for the O sites of In2O3 lattices. Extrinsic ferromagnetism from N precipitations and other secondary phases is safety excluded. Therefore, the observed ferromagnetic ordering in the N-doped In2O3 films should arise from the intrinsic magnetic interaction between N ions. Contacting the monotonous increases in Ms with N doping as well as the strong localization effect of carriers, the carrier-mediated RKKY magnetic interaction can be ruled out. The first principles calculations further confirm that the ferromagnetism in the N doped In2O3 systems has a direct correlation with the substituted N atoms. The N doping introduces spinpolarized hole states in the band gap and generates a total magnetic moment of 2.0µB, which arises from the short-ranged N1:p-In:d/p-N2:p magnetic interaction mediated by localized holes. As is well known, In2O3 is a natural n-type semiconductor due its intrinsic oxygen vacancy defects. The N element doping can induce holes and achieve p-type conductivity in In2O3. For the 2 at% N-doped In2O3 film, the conductivity of the film still exhibits an n-type character, implying that the electrons arising from the oxygen vacancies are not completely compensated by the holes induced by N doping. Therefore, the observed weak ferromagnetism is not meditated by the holes and is possible related to the oxygen vacancies. Some experimental results have proved that oxygen vacancies can introduce spin split impurity bands and cause local moments with longer range coupling.40,41 On the other hand, the trapped electron by oxygen vacancies can occupy a hydrogeniclike orbit which overlaps the p orbit of the N ions and form the magnetic polarons. When the oxygen vacancies surrounding the localized N dopants are enough, some magnetic polarons will overlap with each other and lead to a local ferromagnetic ordering. As the concentration of N doping decreases to 5 at%, the conductivity of the film transforms into p-type and obvious hole states of N 2‫ ݌‬orbitals in bridging In 5‫݌‬ and 4݀ orbitals will appear, which significantly hybridizate with the O 2p, In 5p and 15



1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

4d states. Therefore, an indirect ferromagnetic coupling chain of N1:p-In:d/p-N2:p can be established, which leads to an obvious increase in MS and the crossover of negative MR to positive MR.

4. CONCLUSIONS In summary, the systematic studies on the structural, optical, magnetic and transport properties of the N-doped In2O3 films grown by RF-magnetron sputtering technique were performed. The XRD results show that N dopants are incorporated substitutionally into In2O3 lattice at O sites and no N-related secondary phases are detected. The N-doped In2O3 films are ferromagnetic at 300K and the Ms increases monotonically with the increase of N doping concentration. The increased p-p coupling induced by N doping results in the strongest MS for the 5 at% N-doped In2O3 film. It can be concluded that the magnetic interaction between N dopants is activated through a short-ranged N1:p-In:d/p-N2:p coupling chain, which plays a key role in achieving RT ferromagnetism in the N-doped In2O3 films. The p-p interaction between N 2p orbital also causes large Zeeman-split effect, which can suppress the hopping path of the holes and lead the crossover from the negative MR to positive MR. AUTHOR INFORMATION Corresponding Author *

E-mail: [email protected]

Notes The authors declare no competing financial interest. ACKNOWLEDGMENTS

16


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This work was supported by National Natural Science Foundation of China (Grant 11604242, 11174217) and Tianjin Natural Science Foundation of china (Grant No. 17JCYBJC17300). REFERENCES (1) Saadaoui, H.; Luo, X.; Salman, Z.; Cui, X. Y.; Bao, N. N.; Bao, P.; Zheng, R. K.; Tseng, L. T.; Du, Y. H.; Prokscha, T. Intrinsic Ferromagnetism in the Diluted Magnetic Semiconductor Co:TiO2. Phys. Rev. Lett. 2016, 117, 227202. (2) Quan, Z. Y.; Zhang, X. P.; Liu, W.; Li, X. L.; Addison, K.; Gehring, G. A.; and Xu, X. H. Enhanced Room Temperature Magnetoresistance and Spin Injection from Metallic Cobalt in Co/ZnO and Co/ZnAlO Films. ACS Appl. Mater. Interfaces 2013, 5, 3607-3613. (3) Pellicer, E.; Menendez, E.; Fornell, J.; Nogues, J.; Vantomme, A.; Temst, K.; Sort, J. Mesoporous Oxide-Diluted Magnetic Semiconductors Prepared by Co Implantation in Nanocast 3D-Ordered In2O3-y Materials. J. Phys. Chem. C 2013, 117, 1708417091. (4) Dietl, T.; Ohno, H.; Matsukura, F.; Cibert, J.; Ferrand, D. Zener Model Description of Ferromagnetism in Zinc-Blende Magnetic Semiconductors. Science 2000, 287, 1019-1022. (5) Farvid, S. S.; Sabergharesou, T.; Hutfluss, L. N.; Hegde, M.; Prouzet, E.; Radovanovic, P. V. Evidence of Charge-Transfer Ferromagnetism in Transparent Diluted Magnetic Oxide Nanocrystals: Switching the Mechanism of Magnetic Interactions. J. Am. Chem. Soc. 2014, 136, 7669-7679. (6) Green, R. J.; Regier, T. Z.; Leedahl, B.; McLeod, J. A.; Xu, X. H.; Chang, G. S.; Kurmaev, E. Z.; Moewes, A. Adjacent Fe-Viacancy Interactions as the Origin of

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(41) Coey, J. Weak d0 Magnetism in C and N Doped ZnO. J. Solid State Sci. 2005, 7, 660-667.

Table 1. Parameters of the un-doped, 2 at% and 5 at% N-doped In2O3 films obtained from Hall effect measurements. samples

ρ (Ω·cm)

carrier concentration (cm-3)

type

kFl

0.12

3.74×1018

n

0.212

17

n p

0.007 0.003

un-doped 2 at% 5 at%

7.00 22.58

6.03×10 1.51×1017

Table 2. Formation Energy Ef (eV) for different N doping configurations of In2O3. Configuration

In-rich

O-rich

2N-close

14.15

10.84

2N-far

14.71

11.40

VO-2N-close

17.01

12.05

22


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Figure captions Fig.1. XRD patterns of the un-doped, 2 at% and 5 at% N-doped In2O3 films. The inset shows the enlarge view of (222) diffraction peaks. Fig.2. SEM images of the un-doped (a), 2 at% (b) and 5 at% (c) N-doped In2O3 films. Fig.3. (a) In 3d, (b) O1s and (c) N 1s core-level XPS spectra for the 5 at% N-doped In2O3 film. Fig.4. Magnetic hysteresis loops of the 2 at% and 5 at% N-doped In2O3 films recorded at 300K. Fig.5. The DOS curves (a)-(d) for different N doping configurations and the corresponding spin density distribution (e)-(g) in the N-doped In2O3 systems. The vertical dotted line denotes the position of the Fermi level. The yellow, red and blue balls represent the In, O and N atoms, respectively. Fig.6. (a) The UV-visible transmission spectra of the un-doped, 2 at% and 5 at% Ndoped In2O3 films. (b) (αhν)2 versus hν curves for the films to estimate the optical band gap Eg. The inset shows the variation of optical band gap Eg values as a function of N concentration. Fig.7. (a) Resistivity-temperature (ρ-T) curves and (b) plots of lnρ versus T−1/4 for the un-doped, 2 at% and 5 at% N doped In2O3 films. Fig.8. MR curves measured at 10K for the un-doped, 2 at% and 5 at% N-doped In2O3 films.

23


Intensity (a.u.)

31

(541)

(611)

(440)

(431) SiO2(853)

(400)

32

(622)

30

2θ (degree)

(411) (420) (332)

(321)

Page 24 of 32

(222)

29

(211)

Intensity (a.u.)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

(222)



20

30

40

50

2θ (degree) Figure. 1

24


60

Page 25 of 32

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


(a)

(b)

(c)

Figures. 2(a)-(c)

25



In 3d

In 3d5/2

(b)

O 1s

530.2 eV

444.5eV In 3d3/2 452.0eV

I

Intensity (a.u.)

(a) Intensity (a.u.)

531.6 eV

II 440

444

448

452

456

527

528

Binding energy (eV)

529

530

531

(c)

N 1s

397.1 ev

399.1 ev

394

532

Binding energy (eV)

Intensity (a.u.)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 26 of 32

396

398

400

Binding energy (ev)

Figures. 3(a)-(c)

26


402

533

534

Page 27 of 32

0.8 3

MS (emu/cm )

0.6

5 at%

0.4

M (emu/cm3)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


0.4

0.2 0.0

2 at%

0.0 0.1 0.2 0.3 0.4 0.5

0.0

N concentration (at%)

-0.4

-0.8

-12000 -8000

-4000

0

4000

8000

H (Oe) Figure. 4

27


12000


N-2p O-2p In-5p In-4d

1 (a) 0 -1

DOS (state/ev)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

2N-close

2 (b) 1 0 -1 -2 2 (c)


(f)

2N-far N-2p O-2p In-5p In-4d

0

(g)

2N-close-Vo

-2


2 (d) 0

2N-close-2VO

-2 -4

(e)

-3

-2

-1

0

Energy (ev)

1

2

Figures. 5(a)-(g)

28


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Page 29 of 32

100

0.8

(αhυ)²(ev²)

80

0.6

0.4

0 at% 2 at% 5 at%

0.2

500

600

700

800

3.50 3.43

0.00

0.02

0.04

0.06

N concentration (x)

40

0

400

3.57

3.36

60

0 at% 2 at% 5 at%

20

0.0 300

3.64

Eg (eV)

(b)

(a) Transmittance (%)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


2.7

3.0

3.3

hυ (eV)

Wavelength (nm)

Figures. 6(a)(b)

29


3.6

3.9

4.2


(b)

(a)

0.18 0.16

2000 0.14 1000

un-doped

0.12

ρ (Ω⋅cm)

ρ (Ω⋅cm)

3000

2 at% 0

0.10 10 8 21000

2 at%

6 4

14000

2 7000

un-doped

5 at%

0 -2

0

0

50 100 150 200 250 300

T(K)

0.3

0.4

-1/4

-1/4

T (K )

Figures. 7(a)-(d)

30


0.5

lnρ (Ω⋅cm)

(d)

5 at%

(c)

28000

ρ (Ω⋅cm)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 30 of 32

Page 31 of 32

1.020

10K 1.015

MR (ρH/ρ0)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


1.010


1.005 1.000 0.995 0.990 0.985 -1.0

-0.5

0.0

0.5

H (T)

Figure. 8

31


1.0


Table of Contents Graphic

0.8

5 at%

0.4

2 at% 1.020

0.0

10K

1.015

MR ( ρH/ρ0)

M (emu/cm3)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 32 of 32

-0.4

1.010


1.005 1.000 0.995 0.990

-0.8

0.985 -1.0

-0.5

0.0

0.5

1.0

H (T)

-12000 -8000

-4000

0

4000

8000

H (Oe)

32


12000

Room-Temperature Ferromagnetic Enhancement and Crossover of

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