Intrinsic or Interface Clustering-Induced Ferromagnetism in Fe-Doped

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Functional Inorganic Materials and Devices

Intrinsic or interface clustering induced ferromagnetism in Fe-doped InO diluted magnetic semiconductors 2

3

Xi Luo, Li-ting Tseng, Yiren Wang, Nina Bao, Zunming Lu, Xiang Ding, Rongkun Zheng, Yonghua Du, Kevin Huang, Lei Shu, Andreas Suter, Wai Tung Lee, Rong Liu, Jun Ding, Kiyonori Suzuki, Thomas Prokscha, Elvezio Morenzoni, and Jia Bao Yi ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.8b04046 • Publication Date (Web): 12 Jun 2018 Downloaded from http://pubs.acs.org on June 12, 2018

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ACS Applied Materials & Interfaces

Intrinsic

or

interface

clustering

induced

ferromagnetism in Fe-doped In2O3 diluted magnetic semiconductors Xi Luo, † Li-Ting Tseng, † Yiren Wang, † Nina Bao, ‡ Zunming Lu, § Xiang Ding†, Rongkun Zheng, § Yonghua Du, ∥ Kevin Huang, ⊥ Lei Shu, ⊥ Andreas Suter, # Wai Tung Lee,¶ Rong Liu, ‡ Jun Ding, ‡ Kiyonori Suzuki, ※ Thomas Prokscha, # Elvezio Morenzoni, # and Jia Bao Yi‡* † School of Materials Science and Engineering, UNSW, Kensington, NSW 2052, Australia ‡ Department of Materials Science and Engineering, National University of Singapore, 119260,

Singapore § School of Physics, The University of Sydney, NSW 2006, Australia ∥ Institute of Chemical and Engineering Science, Agency for Science, Technology and Research (A*STAR), 1 Pesek Road, Jurong Island, 627833, Singapore ⊥Department of Physics, Fudan University, 200438, China #Laboratory for Muon Spin Spectroscopy, Paul Scherrer Institute, 5232 Villigen, Switzerland

¶Bragg institute, ANSTO, New Illawarra Road, Lucas Heithers, NSW, 2234, Australia ‡ SIMS Facility, Office of the Deputy-Vice Chancellor (Research and Development), Western Sydney University, Locked Bag 1797, Penrith, New South Wales, 2751, Australia ※ Department of Materials Science and Engineering, Monash University, 3800, Victoria, Australia ‡Global Innovative Centre for Advanced Nanomaterials, School of Engineering, The University of Newcastle, Callaghan 2308, New South Wales, Australia Abstract 5% Fe-doped In2O3 films were deposited using a pulsed laser deposition system. X-ray diffraction and transmission electron microscopy analysis show that the films deposited under oxygen partial pressures of 10-3 and 10-5 torr are uniform without clusters or secondary phases. However, the film deposited under 10-7 torr has a Fe rich phase at the interface. Magnetic measurements demonstrate that the magnetization of the films increases with decreasing oxygen partial pressure. Muon spin relaxation (µSR) analysis indicates that the volume fraction of the ferromagnetic phases in PO2=10-3, 10-5 and 10-7 torr deposited samples are 23%, 49% and 68% respectively, suggesting clusters or secondary phases may not be the origin of the

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ferromagnetism, and that the ferromagnetism is not carrier-mediated. We propose that the formation of magnetic bound polarons is the origin of the ferromagnetism. In addition, both µSR and polarized neutron scattering demonstrate that the Fe rich phase at the interface has a lower magnetization compared to the uniformly distributed phases.

Keywords: Diluted magnetic semiconductor; Ferromagnetism; Clustering; Intrinsic ferromagnetism; In2O3, and Muon spin relaxation (µSR).

Introduction Diluted magnetic semiconductors (DMSs) have drawn extensive attention due to their potential applications in spintronics devices, which are based on the concepts simultaneously utilizing charge and spin degrees of freedom of electrons 1. Since the room temperature ferromagnetism of Mn doped ZnO has been theoretically predicted by the mean field theory 2, the wide-bandgap oxide semiconductors have motivated extensive research interest. Ferromagnetism based on a variety of promising oxide semiconductors, such as ZnO 3-9, TiO2 10-13 SnO2 14-16 and In2O3 17-21

, as the semiconductor host, has been extensively investigated and reported to exhibit

ferromagnetism above room temperature. Mn doped GaAs has been considered as one of the best candidates for diluted magnetic semiconductors, whereas, its low Curie temperature makes it non-eligible for practical applications. Hence, oxide based diluted magnetic semiconductors have attracted continuous research interest due to their possible Curie temperatures higher than room temperature. Up to now, Co doped TiO2 may be one of the best candidates for diluted magnetic semiconductor based on oxide semiconductors for its magnetism uniformity, carrier mediated ferromagnetism and relatively large saturation magnetization as well as spin manipulation by electric field 10-12.

Being one of those candidate hosts, In2O3 is another important transparent oxide semiconductor with a direct band gap of 3.75 eV and can be fabricated as an n-type semiconductor with high electrical conductivity by introducing oxygen deficiencies 22-23. Similar to other wide gap oxide semiconductors, indium oxide has also been widely investigated as the semiconductor host for achieving an intrinsic DMS with a Curie temperature higher than room temperature. In addition, Fe doped In2O3 is one of the few systems which has shown room temperature ferromagnetism in bulk state 24. However, the results reported from different research groups varied, and were even controversial. Room temperature ferromagnetism has been observed in Fe-doped In2O3 thin films 17-25, powders 18-20 and bulk materials 24. The Curie temperature can reach as high as


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900 K 25. A number of groups reported that room temperature ferromagnetism was observed by co-doping Cu or Mn element into In2O3, while Fe-only samples showed paramagnetic behaviour, suggesting carrier-meditated magnetism

23, 26-28

. Because ferromagnetism was

mostly observed in the samples prepared under oxygen poor environment, some researchers attributed the ferromagnetism to the presence of oxygen vacancies instead of charge carriers 22, 24, 29

. On the other hand, some other researchers claimed that the ferromagnetism of Fe-doped

In2O3 is mainly originating from iron oxide nano-clusters 30,31. Currently, there is no evidence for spatially uniform ferromagnetism in the whole volume of the samples of Fe doped In2O3. So far, researchers only used transmission electron microscopy (TEM), energy dispersive spectroscopy (EDS), or electron energy loss spectroscopy (EELS) mapping to confirm the structure uniformity or clusters. However, it is well known that these techniques are very localized 18. Hence, whether the ferromagnetism in Fe doped In2O3 is intrinsic or not has been a debate for a long time. This debate also exists in the whole family of diluted magnetic semiconductors based on wide gap oxide semiconductors, and it has strongly impeded the progress of diluted magnetic semiconductors.

In this work, we used both polarized neutron reflectometry (PNR) and low energy muon spin relaxation (LE-µSR), the only available equipment in the world for the thin film µSR measurements located at PSI, Switzerland, for the study of spatial uniformity of magnetic moments in Fe doped In2O3 systems, thus to elucidate the origin of ferromagnetism in Fe doped In2O3. This work is the first time to prove that both extrinsic and intrinsic ferromagnetism can co-exist in Fe doped In2O3 system. In addition, intrinsic ferromagnetism plays the dominant role in the magnetization contribution. This intrinsic ferromagnetism is due to the formation of bound magnetic polarons mediated by the defects, such as oxygen vacancies, but it is not carrier-mediated.

Materials and methods Fe doped In2O3 films were deposited on 2 cm×2 cm MgO (001) substrates under a variety of oxygen partial pressures using a pulsed laser deposition (PLD) system. In detail, α-Fe2O3 and In2O3 (both from Sigma-Aldrich, 99.99%) were first mixed in a clean mortar, followed by sintering in a furnace at 1173 K for 3 hrs. The sintered powders were examined by X-ray diffraction (XRD). No iron oxide phase could be observed. The sintered powders were then pressed into a pellet with a diameter around 1.25 inches and sintered in a furnace at 1520 K for 10 hrs. The sintered target was examined by XRD again to confirm a pure In2O3 phase without



any impurities or secondary phases. The films were then deposited using a PLD system with a KrF excimer laser operating at 248 nm and a fluence of 1.0-1.8 J/cm2 (corresponding to a deposition rate of approximately 3-10 nm/min) at 873 K. The base pressure of the system is 10-8 torr. The film thickness was controlled to be approximately 40 nm, determined by the deposition time and measured by a depth profilometer. XRD (PANalytical Xpert Multipurpose X-ray diffraction system) using Cu Kα radiation was used for the structure and phase characterization. Energy dispersive X-ray spectroscopy (EDS), X-ray photoelectron spectroscopy (XPS, Kratos AXIS Ultra DLD) and X-ray absorption near edge structure spectroscopy (XANES, Singapore synchrotron light source (SSLS)) was used for the composition and valence state analysis. The magnetic and transport properties were measured using a superconducting quantum interference device (SQUID, Quantum Design, MPMS, XL5, USA) and a physical property measurement system (PPMS, Quantum Design, 14T, USA). Substrate MgO paramagnetic signal has been carefully subtracted by measuring the magnetic signal of MgO at different temperatures. The polarized neutron reflectometry (PNR) measurement was carried at the Bragg Institute, Australia Nuclear Science Organization (ANSTO). Low energy µSR was carried out at PSI, Switzerland. The analysis of the data is described in detail in Ref. 12.

Figure 1: TEM image and EDS mapping of 5%Fe doped In2O3 films. TEM image of (a) PO2=10-3 torr. (b) PO2=10-5 torr. (c) PO2=10-7 torr. EDS mapping of (d) PO2=10-3 torr. (e) PO2=10-5 torr. (f) PO2=10-7 torr.


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1.2 1.0 0.8 Fe-In2O3 Fe2O3 Fe3O4 FeO Fe foil

0.6 0.4 0.2 0.0 -0.2

7100

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Energy (eV) Figure 2: Near edge X-ray absorption spectroscopy of 5%Fe doped In2O3 films deposited under an oxygen partial pressure of 10-7 torr.

Results and Discussion The deposited films were first characterized by XRD analysis. It turned out that the films grew with two directions, namely (222) and (004), indicating the highly textured growth of the films, which is consistent with other literature data

32

(Supporting information Figure S1). The

thicknesses of all the films are around 40 nm respectively. Figure 1 shows the TEM images of Fe- In2O3 films deposited under an oxygen partial pressure of 10-3, 10-5 and 10-7 torr respectively. Due to the textured growth of the films, no polycrystalline grains could be observed. In Fig.1 (c), there is a transitional area in the interface. Dislocations are found as labelled in Fig. 1(c). This secondary phase has a thickness of 3-4 nm. However, there is no secondary phase observed in the interface for the films deposited under PO2=10-3 and 10-5 torr (Fig. 1a (a) and (b)), though the interface between the film and the substrate is not so clear. For an intrinsic diluted magnetic semiconductor, the basic requirement for the dopant is to be



uniformly distributed in the semiconductor host. Figures 1 (d)-(f) show the Fe dopant mapping by EDS, which is attached to the TEM system. All the films show thicknesses around 38 nm. No apparent clustering of Fe ions can be observed in the samples deposited under PO2=10-3 and PO2=10-5 torr. However, in the PO2=10-7 torr sample, it can be seen that there is an accumulation of Fe dopants in the interface between the film and the substrate. The accumulated thickness is around 4 nm, consistent with the high-resolution TEM analysis. XPS analysis indicates that dopant Fe is in 3+ on the surface, suggesting that Fe dopant in the film may be in a substitutional state, though XPS is surface sensitive. In order to understand the valence state inside the film, depth profile of XPS analysis for PO2=10-7 torr sample was carried out. The results indicate that inside the film, the valence state begins to change to 2+ until to the interface, where the valence state of Fe is with a mixture of Fe2+/Fe3+. From XPS analysis, we also can obtain the composition of Fe is 5.6% on the surface, and it gradually decreases in the middle of the film. Close to the interface, the composition of Fe increases slightly. In order to confirm the Fe dopant distribution in In2O3 film, secondary ion mass spectrometry (SIMS) was performed. The spectra are shown in supporting information (Figure S5). It can be seen that elements, In and O, are uniformly distributed. Fe concentration is slightly higher on the surface and decreases gradually in the middle, similar to that obtained from XPS analysis, whereas, Fe has a quite concentration in the interface. XPS also obtained relatively high concentration of Fe in the interface, but not so evidently. It may be due to the step for XPS analysis is too large to achieve interface information. From EDS analysis, the overall doping concentration of Fe is close to 5 %. Furthermore, from Fig. 1(c), the high-resolution TEM image of the interface structure is different from that of the film in other areas. Dislocations can be clearly seen, suggesting a mismatch during film growth.


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Figure 3: M-H loops of 5% Fe doped In2O3 deposited under (a) PO2=10-7 torr. (b) PO2=10-5 torr. (c) PO2=10-3 torr. (d) Saturation magnetization dependence on temperature of the three samples.

In order to further understand the status of Fe in the interface, we carried out X-ray absorption near edge spectroscopy (XANES) measurement, as shown in Fig. 2

33

. Compared with

reference samples Fe2O3, Fe3O4, FeO and Fe foil, the spectrum of Fe in Fe doped In2O3 in the near edge is different from both Fe2O3 and Fe3O4, but very close to Fe3O4, suggesting the Fe in In2O3 has a mixing valence state of Fe2+/Fe3+ 34. From XRD analysis, we cannot detect any Fe3O4 phases. It may be due to its detection limitation of XRD system. Hence, the Fe rich phase may be composed of Fe3O4 phase. The high vacuum (10-7 torr) may be attributed to the formation of Fe3O4 clusters.



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Figure 4: (a) R-T curve of the sample deposited under oxygen partial pressures of 10-3, 10-5 and 10-7 torr, respectively. The inset is a zoom of the Y axis of the PO2=10-7 torr sample, showing a metal-semiconductor transition around 150 K. (b) Plots of ln(ρ) versus T-1/4 from data of (a). The inset is a zoom of the Y axis of the PO2=10-7 torr sample.

Fe doped In2O3 has been one of the most important oxide based diluted magnetic semiconductors. Several works have shown that the magnetic properties of Fe doped In2O3 are strongly dependent on the oxygen partial pressure during film growth

34,35

. In this work, the

magnetic properties of Fe doped In2O3 films deposited under an oxygen partial pressure of 103

, 10-5 and 10-7 torr, respectively, are shown in Fig. 3. MgO shows strong paramagnetism at

low temperature. We have carefully subtracted the MgO signal from every sample. However, it will affect the shape of the hysteresis loop for the samples with very small saturation magnetization. Room temperature ferromagnetism is observed in all the samples, evidenced by the small coercivity in the hysteresis loops taken at room temperature (Figs. 3 (a)-(c)). The saturation magnetization increases with decreasing oxygen partial pressure. The saturation magnetization of PO2=10-7 torr sample is close to 19 emu/cm3 at room temperature, corresponding to 1µB/Fe, and 40 emu/cm3, corresponding to 2.1 µB/Fe, whereas, PO2=10-3 and PO2=10-5 samples exhibit very small saturation magnetization, 3.8 and 1.3 emu/cm3, corresponding 0.22 and 0.08 µB/Fe. It may be due to that at PO2=10-7 torr, more oxygen vacancies have been produced, which play critical role in the formation of ferromagnetic ordering 34. It may also be possible due to the formation of Fe3O4 clusters, which contribute mainly to the overall saturation magnetization. In addition, for the PO2=10-3 and PO2=10-5 samples, it is observed that the saturation magnetization at 5 K is much higher than that at 300 K, indicating that there might be a large amount of paramagnetic phase at room temperature. It should be noted that the coercivity at 5 K is very small in Fig. 3(b) and 3(c) compared to that


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at 100 K. The reason is that the magnetization of the both samples is very small, very close to the limitation of SQUID sensitivity. At 5 K, the possible paramagnetic phase of Fe-In2O3 and MgO substrate both show ferromagnetic/paramagnetic signal at 5 K, which may smear the signal of coercivity, though the magnetization has been carefully corrected by measuring pure MgO substrate. Figure 3(d) shows the temperature dependence of saturation magnetization of the three samples. The saturation magnetization does not change too much from low temperature (5 K) to room temperature (300 K) for the PO2=10-7 torr sample, confirming the ferromagnetism. It should be noted that there is no much magnetic signal for the pure In2O3 deposited under the oxygen partial pressure of 10-7 torr.

Figure 5: Normalized muon stopping profiles for various muon implantation energies.

Figure 4 shows the temperature dependence of electrical resistivity for the samples deposited under PO2=10-3 torr, PO2=10-5 torr, and PO2=10-7 torr, respectively. The decreasing resistivity with decreasing oxygen partial pressure can be attributed to an increasing concentration of oxygen vacancies. The sample deposited under PO2=10-3 torr has a relatively high resistivity. Both PO2=10-3 and PO2=10-5 samples show semiconductor behaviour and the resistance increases with decreasing temperature. However, the PO2=10-7 sample shows a metallic behaviour at relatively high temperature, whereas, at low temperature, it shows semiconductor behaviour (Inset of Fig. 4(a)). This transition has been observed before and it is associated with itinerant carriers induced by dopant doping

21, 36

. For the R-T curves, we have used several

ways to do the fitting, such as ln(ρ) versus temperature T-1/4 (from Mott variable range hopping (VRH)) and ln(ρ) versus temperature T-1/2 (from Efros VRH). We found that best linear fits are



obtained for ln(ρ) versus T-1/4 in the low temperature range (T≤65 K). This implies that the electrical transport of the film at this temperature range may exhibit Mott VRH behaviour 35, 37, 38

, which indicates that the majority of the carriers is localized, suggesting that the

ferromagnetism in the films may not be carrier-mediated. Bound magnetic polarons, forming in the presence of oxygen vacancies, may be the origin for the emergence of ferromagnetism 39

. For the PO2=10-3 torr sample, the linear fit can be only applied down to 25 K. Below 25 K,

the large resistivity makes its measurement more unreliable, probably causing systematic errors which may lead to a deviation from the linear dependence.

The three films all show magnetoresistance (MR) behaviour. Due to the relatively high resistivity and weak ferromagnetism, the MR of both PO2=10-3 and PO2=10-5 samples are very weak and negative. However, there is a positive MR for the PO2=10-7 sample at a temperature higher than 150 K (see supporting information), which coincides with the metal-semiconductor transition at around 150 K (Fig. 4 (a)). Hence, we attribute the positive MR to itinerant electrons, which mediate the magnetic moment, resulting in a strong s-d interaction. At low temperature, the electrons are localized, where we attribute the negative MR to reduced spin scattering from variable range hopping of electrons. For the PO2=10-3 and PO2=10-5 torr samples, there is no metallic behaviour. Hence, no positive MR can be observed.

(a) 1.0 Ref. 0.5 0.0 -0.5 -1.0 1.0 0 10-3 torr2 0.5 0.0 -0.5 -1.0 2 1.0 0 -5 0.5 10 torr 0.0 -0.5 -1.0 2 1.0 0 -7 0.5 10 torr 0.0 -0.5 -1.0 0 2

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Figure 6: (a)Typical normalized LE-µSR spectra measured by applying a transverse magnetic field of 30 Oe at 5 K. The reference sample is a non-doped In2O3 sample grown in 10-7 torr, and the other samples are doped with 5% of Fe and grown in an oxygen partial pressure of PO2=10-3 , PO2=10-5 and PO2=10-7 torr, respectively.(b) LE-µSR spectra of the samples in (a), measured without applying a magnetic field (zero-field µSR).


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Low energy muon spin relaxation (LE-µSR) is one of the methods to identify the spin environment from surface to the interface of thin films due to its tuneable muon beam energy (from 1 to 30 keV). This technique has been successfully applied to a variety of systems such as diluted magnetic semiconductors

12, 40

or oxide heterostructures

41-44

. Magnetic uniformity

has been reported in Mn doped GaAs magnetic semiconductors, confirming the intrinsic and 40

carrier-mediated ferromagnetism

. Figure 6 shows calculated muon stopping distributions.

Low-energy muons with an energy of E = 5 keV can reach a depth of 40 nm at the interface to the MgO substrate, with a mean stopping depth of about 25 nm near the center of the Fe-doped In2O3 film. Normalized LE-µSR fitted spectra taken at 5 K and 5 keV implantation energy at an applied transverse field of 30 Oe are shown in Figure 6. 0.18 0.17

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The TF data is fitted to a cosine function multiplied by an exponential taking into account the field inhomogeneities with a damping rate λS: 𝑃(𝑡) = 𝐴𝑐𝑜𝑠(𝛾𝑢 𝐵𝑡 + 𝜑)𝑒 λ𝑠 𝑡 .


The decay


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asymmetry A, or amplitude of the precession signal, is a measure of the non-magnetic volume fraction, 𝛾𝑢 is the muon gyromagnetic ratio, B is the magnetic field at the muon stopping site, and 𝜑 is the phase of the precession signal in the positron detector. By comparing this asymmetry to the reference asymmetry in a non-magnetic sample of similar geometry, we can deduce the non-magnetic volume fraction of the magnetic films, which is simply given by (1 normalized asymmetry).

The oscillation of the reference sample (pure In2O3) and the relatively slow damping with the rate λ indicate the non-ferromagnetic nature of the film. Here, the damping is caused by nuclear dipolar fields from the In nuclei. When the film was deposited under PO2=10-3 torr, an evident reduction of the oscillation amplitude can be observed, suggesting some ferromagnetic phases in the film. When the film was deposited under PO2=10-5 torr, a larger reduction of the oscillation amplitude is observed. Further decreasing the oxygen partial pressure to PO2=10-7 torr leads to an almost flat spectrum (disappearing oscillation). From the reduction of the decay asymmetry, we can calculate the volume fractions of the ferromagnetic phase, which are 23%, 49% and 68%, respectively, indicating that none of the samples has a uniform magnetic phase (i.e. 100% volume fraction).

Typical zero-magnetic-field (ZF) asymmetry spectra taken at 5 K with 5 keV muons are shown in Fig. 7 (b). The ZF spectra in strongly magnetic samples with dilute magnetic moments can be fitted by an exponential Kubo-Toyabe (KT) function multiplied by an exponential: 1

2

A(t)=A3 + 3 (1 − ∆𝑡)𝑒 −∆𝑡 ]𝑒 −λ𝑒𝑡 . Here ∆ is related to the width of the local field distribution ∆B of the dilute static magnetic moments by ∆= = 𝛾𝑢 ∆𝐵, and λ𝑒 is the damping rate due to some slow electronic dynamics. The signal in pure In2O3 and weakly magnetic samples grown at PO2=10-3 torr is better accounted for using a simple fast-depolarization exponential plus a slow-depolarization

exponential

function: 𝑃(𝑡) = 𝐴𝑓 𝑒 −λ𝑓 𝑡 + 𝐴𝑠 𝑒 −λ𝑠 𝑡 ,

with λ𝑠 ≈

0.2(1)𝜇𝑠 −1 . The slow component represents a combination of the paramagnetic and background signal, and the non-damping component of the small magnetic signal corresponding to the first term of the KT function. The depolarization rate of the fast 4

component λ𝑓 is related to ∆ by ∆λ𝑓 ≈ 3 ∆ 40. As in the TF measurements the signal in the undoped reference sample is slowly damped, typical of a non-magnetic sample. The PO2=10-3 torr sample exhibits a similar behaviour to that of the undoped samples, suggesting a nonuniform and weak ferromagnetic ordering, while the PO2=10-5 and PO2=10-7 torr samples


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exhibit a fast damping rate, suggesting a relatively strong ferromagnetic ordering causing a large width of the magnetic field distribution measured by the muons.

Figure 7 shows the asymmetry and damping rate dependence on temperature and implantation energy of muons in TF mode. The asymmetry and damping rate as a function of temperature were measured at a muon implantation energy of 5 keV. The temperature dependence of the asymmetry (Fig. 7 (a)) of the reference sample shows the full asymmetry for a sample with a size of 2×2 cm2, evidencing nonmagnetic behaviour. The increase in the asymmetry on increasing temperature for all the doped samples indicates the weakening of ferromagnetic ordering with increasing temperature due to the thermal vibrational energy since increasing temperature can randomize the spin alignment. Similarly, the damping rate has an opposite trend to that of the asymmetry (Fig. 7 (b)). Note, that the damping rate of ~ 0.15 𝜇𝑠 −1 of the reference sample is due to the nuclear moments of In. The increased damping rate of the magnetic samples is due to muons stopping in the non-magnetic regions of the sample, but experiencing magnetic stray fields from the ferromagnetic regions of the sample.

Figure 7 (c) and (d) show the muon energy dependence of the asymmetry and damping rate at 5 K. Whereas the damping rate of the reference sample is nearly unchanged, in the three doped samples, the damping rate increases with increasing implantation energy until it reaches a maximum value at an implantation energy of 5 keV, when most of the muons are stopping close to the center of the film. At higher energies, when muons reach the interface and begin to penetrate into the substrate, the stray magnetic fields in the interface region are getting weaker, followed by a further reduction in the non-magnetic MgO, causing a decrease of depolarization rate. Below 5 keV, on approaching the surface, the damping rate also decreases, suggesting a weakening of ferromagnetic ordering at the surface. The relatively low ferromagnetic ordering at the surface is attributed to a surface ``deadlayer''. From the TEM analysis, there is a Fe rich area in the PO2=10-7 torr sample at the interface. Therefore, the LEµSR results indicate that the Fe rich area at the interface actually has a lower magnetization than the center of the film, which means that the Fe rich area is not the major origin of the observed ferromagnetism. The non-uniformity of the magnetic ordering is supported by the asymmetry dependence on muon implantation energy. For the reference and PO2=10-3 torr sample, the asymmetry increases with increasing muon energy at first. The energy dependence of the reference sample up to 8 keV can be explained by a well-known apparatus effect, where the energy dependent muon backscattering causes a reduction of observable asymmetry at low



implantation energies. Above 8 keV muons are stopping in MgO where muonium formation (bound state of positive muon and an electron) 45, 46 leads to a reduction of the fraction of muons precession at the muon's Larmor frequency, i.e. a reduction of observable asymmetry. For the PO2=10-3 sample the lower asymmetry while stopping in the Fe-In2O3 film indicates the presence of ferromagnetic regions. For the PO2=10-5 and PO2=10-7 samples, if the energy is smaller than 5 keV, the asymmetry is already significantly reduced compared to the other two samples due to the larger magnetic volume fractions. In contrast to the reference and the PO2=10-3 sample the asymmetry decreases with increasing muon energy, where a simultaneous increase of the damping rate is observed. This can be attributed to an increasing magnetic volume fraction as a function of depth, suggesting increasing ferromagnetic ordering. The peak of the depolarization rate around 5 keV further indicates that the local magnetic moment is not uniform, reaching a maximum value in the center of the film. When the energy is larger than 5 keV, there is a sudden decrease of the asymmetry, which has the same trend as in the reference and the PO2=10-3 torr sample, and is again attributed to muonium formation in the substrate. In addition, we deposited PO2=10-7 torr samples with thicknesses of 20 and 100 nm. Magnetic measurements indicate that the magnetization does not vary significantly, suggesting that neither the interface nor the surface plays a critical role in the magnetization, supporting the µSR results.


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Figure 8: Polarized neutron reflection curves taken at room temperature of PO2=10-7 torr sample. (a) In the positive neutron spin state. (b) In the negative neutron spin state. (c) NSLD and MSLD curves versus film thickness. (d) Asymmetry of the curves measured at two spin states (blue, hollow sphere) and the corresponding fitting curve (red, solid square).

In order to further confirm the low magnetic moment of the Fe rich area in the interface, polarized neutron reflectometry (PNR) is employed to identify the structure and magnetic profile of the three samples. Due to the low magnetization of the PO2=10-3 and PO2=10-5 torr samples, the PNR measurement could not be carried out for these two samples. Figure 8 shows PNR measurement results of the PO2=10-7 torr sample at room temperature. Figure 8 (a) and (b) show the neutron reflection data for spin states R+ and R- with fitted results. The spin polarized reflectivity spectra were fitted with the MOTOFIT software packages

47

. The

reflectivity profile for a particular model is calculated using the Abeles matrix method 48. The fitted results indicate that the film is composed of a two-layer structure. The low scattering length density (SLD) is corresponding to Fe doped In2O3. The higher SLD in the interface should be corresponding to the Fe rich area since Fe3O4 has a neutron scattering length density of 6.935 (10-6/Å), whereas, In2O3 has a neutron scattering length density of 3.9578 (10-6/Å).



Page 16 of 22

The thickness of this area is approximately 4 nm, which is consistent with the TEM and EDS mapping analysis. From the calculation, the uniformly distributed area has a magnetic moment of 40 emu/cm3, which agrees well with the SQUID measurement. The Fe rich area has a saturation magnetization of approximately 20 emu/cm3. This lower magnetization in the Ferich interface region is consistent with the LE-µSR measurements. Hence, at room temperature, the magnetization is mainly contributed by areas other than the Fe rich area. It should be noted that bulk Fe3O4 has a saturation magnetization over 475 emu/cm-3 at room temperature. Nanostructured Fe3O4 has shown much lower saturation magnetization 49. At 5 K, the magnetization will significantly increase, which may attribute to nanostructured clusters having the enhanced magnetization at 5 K for PO2=10-7 torr sample, as shown in Fig. 3(a). In addition, the study of different film thicknesses showed that the saturation magnetization did not increase with decreasing thickness, indicating that the interface does not play a critical role in the magnetization. Furthermore, we suggest that - as indicated by the R-T measurements in Fig. 4, the dependence of the magnetic signal on oxygen vacancy concentration and the inhomogeneous magnetism shown by µSR - that the ferromagnetism is mainly due to the formation of bound magnetic polarons, in analogy to Co-doped ZnO

50

. Carrier mediated

ferromagnetism may exist at a temperature higher than 150 K for the PO2=10-7 torr sample due to its metallic behaviour at this temperature range, though µSR measurement has shown that the asymmetry at 200 K is similar to that of 150 K, as shown in Fig. 7.

In summary, we have deposited 5 % Fe doped In2O3 films using a pulsed laser deposition system. The doped films all show room temperature ferromagnetism and the magnetization increases with decreasing oxygen partial pressure. EDS mapping shows that the films deposited under an oxygen partial pressure of 10-3 and 10-5 torr have a uniform distribution of Fe dopants within the detection limit. However, the film deposited under an oxygen partial pressure of PO2=10-7 torr has an accumulation of Fe dopants at the interface. LE-µSR shows that the ferromagnetic phases in the PO2=10-3, 10-5 and 10-7 samples occupy volume fractions of 23%, 49% and 68%, respectively, which confirms that neither clusters nor secondary phases are the major contribution to the ferromagnetism in In2O3 based diluted magnetic semiconductor. Instead, we propose that bound magnetic polarons, forming in the presence of oxygen vacancies, are the main origin of ferromagnetism in Fe doped In2O3 films. In addition, from LE-µSR and PNR measurement, the Fe rich area at the interface to the substrate has a lower magnetization compared to other areas. XANES shows that the accumulated Fe dopants are in


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the 2+ state and that the Fe rich region exists as Fe3O4 as substitutional defects. The lower magnetization in this region is attributed to the antiferromagnetic coupling of dopants.

ASSOCIATED CONTENT Supporting information XRD spectra of the 5%Fe doped In2O3 samples under the oxygen partial pressure of 10-3, 10-5 and 10-7 torr. EDS mapping of PO2=10-3 and 10-5 torr samples. XPS spectra of 5% Fe-In2O3 surface and depth profile and magnetoresistance curves of PO2=10-7 torr sample at different temperatures.

AUTHOR INFORMATION *Tel: 61(0)249261625; Email address: [email protected] The authors declare no competing financial interest.

ACKNOWLEGEMENT J.Yi acknowledges the support of the Australia Research Council discovery project grants DP140103041 and Future Fellowship FT160100205. The muon experiments were performed at the Swiss Muon Source, Paul Scherrer Institute, Villigen, Switzerland.

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Intrinsic or Interface Clustering-Induced Ferromagnetism in Fe-Doped

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