Room temperature ferromagnetism of Mn-doped SnO2 thin films fabricated by sol–gel method

Applied Surface Science 254 (2008) 7459–7463

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Room temperature ferromagnetism of Mn-doped SnO2 thin films fabricated by sol–gel method Yuhua Xiao a, Shihui Ge a,*, Li Xi a, Yalu Zuo a, Xueyun Zhou a, Bangmin Zhang a, Li Zhang a, Chengxian Li, Xiufeng Han b, Zhenchao Wen b a b

Key Laboratory for Magnetism and Magnetic Materials of Ministry of Education, Lanzhou University, Lanzhou 730000, PR China State Key Laboratory of Magnetism, Institute of Physics, Chinese Academy of Sciences, Beijing 100080, PR China

A R T I C L E I N F O

A B S T R A C T

Article history: Received 6 May 2008 Received in revised form 31 May 2008 Accepted 2 June 2008 Available online 17 June 2008

Sn1 xMnxO2 (x  0.11) thin films spin-coated on Si (1 1 1) substrate were fabricated by sol–gel method. X-ray diffraction revealed that single-phase rutile polycrystalline structure was obtained for x up to about 0.078. Evolution of the lattice parameters and X-ray photoelectron spectroscopy studies confirmed the incorporation of Mn3+ cations into rutile SnO2 lattice. Magnetic measurements revealed that all Sn1 xMnxO2 thin films exhibit ferromagnetism at room temperature, which is identified as an intrinsic characteristic. Magnetization data showed that the average magnetic moment per Mn atom decreased and the coercivity increases with increasing Mn content. The origin of room temperature ferromagnetism can be understood in terms of the percolation of the bound magnetic polaron. Our experimental results prove that the sol–gel method is an effective method for fabrication of transition metal doped SnO2 nanostructures with room temperature ferromagnetism by chemical synthesis. ß 2008 Elsevier B.V. All rights reserved.

Keywords: Room temperature ferromagnetic semiconductor Sol–gel method Mn-doped SnO2 film The percolation of bound magnetic polaron

1. Introduction Diluted magnetic semiconductors (DMSs) have been attracting intense interests because of their potential applications in spinbased information-processing technologies [1–3]. A key work for realizing spintronic devices is to develop DMSs with Curie temperature higher than room temperature. Since Dietl et al. predicted the existence of RTFM in transition metal (TM) doped semiconductors based on mean-field Zener model [4], there have been many reports on ferromagnetism above 300 K in thin films of wide-gap semiconductors doped with a few percent of 3d transition-metal ions. The host materials are oxides – TiO2 [5,6], ZnO [7–11,16], SnO2 [12–16], or nitrides – GaN [17], AlN [18], and the 3d dopants are Sc, Ti, V, Cr, Mn, Fe, Co, Ni, Cu or Zn. Recently, there is a review of results in ZnO and GaN by Liu et al. [19]. SnO2 is an attractive material for solar cells and gas-sensing applications due to its high optical transparency and electrical conductivity [20]. As a wide-band-gap (3.6 eV) oxide semiconductor, SnO2 has a rutile structure similar to that of TiO2, in which RTFM was first discovered [5]. Because of the native oxygen vacancies and high carrier density, SnO2 may be an attractive host semiconductor for the fabrication of DMS.

* Corresponding author. E-mail address: [email protected] (Y. Xiao). 0169-4332/$ – see front matter ß 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.apsusc.2008.06.026

Although there have been relatively few reports on the SnO2based DMSs, compared to those on other oxide-based DMSs [19], some interesting properties have been observed in Co-doped SnO2 systems. For example, transparent Co-doped SnO2 films with a giant magnetic moment (7.5  0.5mB/Co) and low coercivity (50 Oe) were reported [12], while a relatively low magnetic moment (0.133mB/Co) as well as a high coercivity (630 Oe) were observed in Sn1 xCoxO2 powders [21]. Most Sn1 xMxO2 (M = V, Cr, Mn, Fe, Co, or Ni) samples as mentioned above were fabricated by pulsed laser deposition (PLD) and show single crystal or close to single crystal structure [12– 16,22]. Only a few Sn1 xMxO2 systems made by other methods, such as ceramic synthetic method [23], spray pyrolysis [24], chemical synthetic method [25], were reported. In the Mn-doped SnO2 systems studied by different groups, the experimental results are quite different even controversial. Kimura et al. investigated pulse laser deposited thin films of Mn-doped SnO2 and found that the magnetic moments per Mn site decrease with increasing Mn content x, indicating an antiferromagnetic (AFM) nature of Mn–O–Mn interaction reminiscent of AFM in rutile MnO2 [20]. On the other hand, Fitzgerald et al. reported that bulk ceramic Mn-doped SnO2 showed room temperature ferromagnetism and suggested that the exchange interaction mechanism of the observed RTFM involve an electron trapped at an oxygen vacancy adjacent to the transition metal ions (F-center exchange) [23]. They also stated: In spite of the fact that some minor amounts

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of impurity compounds were detected, the magnetization could not be attributed to their presence, but to the transition metals substituting tin in SnO2. Chemically synthesized Sn1 xMnxO2 powders have been studied by Duan et al. [25]. The authors suggested magnetic properties in Mn-doped SnO2 nanocrystalline powders depend on the competition between the free carriers mediated Ruderman–Kittel–Kasuya–Yosida (RKKY) interaction and the antiferromagnetic superexchange interaction. Obviously, it seems that the magnetism is very sensitive to the sample preparation methods. In this work, we fabricated Sn1 xMnxO2 thin films on Si (1 1 1) substrate by a simple sol–gel method and investigated the structure and magnetic properties of Mn-doped SnO2 thin films.

To validate the oxidation states of the Mn ions in the samples, Fig. 1 presents XPS spectra of the Sn 3d (a), O 1s (b), and Mn 2p (c) electrons for the Sn0.948Mn0.052O2 film (all the samples have similar XPS spectra), in which the adventitious C 1s peak at 284.6 eV binding energy has been removed by Ar sputtering. As shown in Fig. 1(a), the binding energies of the Sn 3d5/2 and Sn 3d3/2

2. Experiment Tin tetrachloride [SnCl4] and the required amount of manganese nitrate hydrate [Mn(NO3)2
6H2O] were dissolved in distilled water and anhydrous ethanol respectively, following by stirring and circumfluencing at 80 8C for 4 h and aging for a week, then the solution was acquired. For the film preparation, the solution was spin-coated on Si (1 1 1) substrate following by being heated at 120 8C for 10 min. After multilayer-coating, the film precursors were obtained. Finally, the precursors of films were calcined in air at 600 8C for 1 h to obtain Sn1 xMnxO2 thin films. An ellipsometer was used to measure the thicknesses of the films which were determined to be around 180–220 nm. The content of manganese element was determined by inductive coupled plasma emission spectrometer (ICP) and energy dispersive X-ray analysis (EDAX). Xray photoelectron spectroscopy (XPS) was employed to test the chemical valence of the films on a VG ESCALAB MK II spectrometer at room temperature using an Al Ka X-ray source (hn = 1486.6 eV). In order to remove the surface contamination, the sample surface was sputter-etched for 10 min using an Ar+ sputter gun (E = 1 kV, emission current of 18 mA) before testing. X-ray diffraction (XRD) patterns were collected on a Philips X’ pert Pro diffractometer with Cu Ka1 radiation (l = 1.5406 A˚). Magnetization studies were carried out using a superconducting quantum interference device (SQUID) magnetometer (MPMS, Quantum Design) in the temperature range 5–350 K. Room temperature magnetization was recorded using a vibrating sample magnetometer (VSM). The electrical resistivity was measured using the standard four-point technique with Keithley-2000 Multimeter. 3. Results and discussions Table 1 lists the Mn concentrations of the Sn1 xMnxO2 films determined by both EDAX and ICP analyses. It is shown that the actual Mn contents of the films are very different from those of the film precursors. The calcining does not maintain the stoichiometry of the precursor in face of the volatility of SnO2. In fact the films contain roughly 1.7 times as much of the Mn element as the precursors. A lesser enhancement of 3d dopant concentration has been reported in Co or other transition metals doped ZnO and SnO2 films made by PLD [26,27]. Furthermore, the Mn concentrations of the samples determined by ICP are well consistent with those of the EDAX. Table 1 Nominal Mn concentrations in the film precursors, the actual Mn concentrations determined by both ICP and EDAX analyses Mn% nominal

0.0

0.7

1.6

3.1

4.8

6.3

Mn% EDAX analysis Mn% ICP analysis

0.0 0.0

1.3 1.8

2.8 3.2

5.2 5.1

7.8 7.9

11.0 10.8

Fig. 1. (a–c) XPS spectrum of the Sn0.948Mn0.052O2 film.

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Fig. 2. XRD patterns of Sn1 xMnxO2 films on Si (1 1 1) substrates. The inset shows the FWHM value of (1 1 0) peaks according to the XRD data.

are 486.7 and 495.1 eV, respectively. The binding energy peak of O 1s in Fig. 1(b) is 530.6 eV. The binding energy values of Sn 3d and O 1s evidence that the valence of Sn and O in Sn1 xMnxO2 films is +4 and 2, respectively. The Mn 2p core level spectra, as seen in Fig. 1(c), are split into 2p3/2 and 2p1/2 components due to the spin– orbit coupling. From the smoothed spectra (thick solid lines), the core level binding energies of Mn 2p3/2 and Mn 2p1/2 for Sn0.948Mn0.052O2 film are estimated to be about 641.8 and 653.4 eV, respectively, and the difference in binding energy between Mn 2p3/2 and Mn 2p1/2 is 11.6 eV. The satellite peak appearing at about 646.9 eV also is the characteristic of Mn3+. Based on above analysis, we confirm that Mn ions have a chemical valence of +3 in all samples. Fig. 2 shows the XRD patterns of the Sn1 xMnxO2 films spincoated on Si (1 1 1) substrate with the actual Mn contents x of the samples determined by ICP method. It is evident that only the peaks corresponding to the rutile-type phase SnO2 (space group P42/mnm) are detected with x up to 0.078, indicating that the samples have pure rutile polycrystalline structure. The Mn doping does not impact the crystal structure of SnO2. With increasing Mn

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doping to x  0.11, the diffraction pattern displays a weak undesired peak (labeled by *) at 33.038, which corresponds to the (2 2 2) diffraction peak of Mn2O3 (space group Pcab), implying the occurrence of Mn2O3 phase. In addition, the inset shows that the full width at half maximum (FWHM) of (1 1 0) peaks increase with Mn concentration, indicating the decrease of the crystallite size. The evolution of lattice parameters a and c with Mn content x is shown in Fig. 3. The a and c decrease with x from 0 to 0.078, which could be attributed to the difference between the effective ionic radius of high-spin Mn3+ (0.645 A˚, coordination number CN = 6) while it is smaller than that of Sn4+ (0.69 A˚, CN = 6). For impurityfree Sn1 xMnxO2 samples, therefore, it is reasonable to consider that Mn3+ cations incorporate into rutile SnO2 matrix, which is consistent with the results of XPS. As shown in Fig. 3, the lattice parameters are practically unchanged for x  0.078, so the solubility limit of Mn in SnO2 films by the sol–gel method is estimated to be x  0.078, which is larger than those reported on Sn1 xMnxO2 bulk or powders [25,28]. To avoid any possible magnetic contamination, samples and the materials contacting them were dealt with very carefully during the preparation processes. Fig. 4 shows the room temperature VSM hysteresis loops for Sn1 xMnxO2 thin films on Si (1 1 1) substrates, where the magnetic field was applied in the film planes and the diamagnetism of substrates were subtracted linearly. The undoped SnO2 films on Si (1 1 1) substrates under the same preparation conditions were examined and shown to be of diamagnetic behavior. This result confirms that there is no magnetic contamination from the substrates, pure SnO2, and their interface. The hysteresis loops clearly indicate that the polycrystalline thin films are ferromagnetic with TC above room temperature. The inset of Fig. 4(a) shows the coercivity increases from 51 to 77 Oe with increasing Mn content, which can be ascribed to the decreasing of crystalline size. The relationship between the magnetic moment per Mn ion of films corresponding to saturation magnetization (MS) and Mn content is shown in the inset of Fig. 4(b). It is obvious that the value of magnetic moment per Mn drops rapidly with the increase of Mn contents. Similar results have been observed in 3d elements doped oxide semiconductor films [8,29], and the origin can be understood by antiferromagnetic superexchange interaction between neighbouring Mn3+ ions through O2 ions. With an increase in Mn doping, the average distance between the Mn3+ ions

Fig. 3. The evolution of cell parameters a and c of Sn1 xMnxO2 thin films as a function of Mn content x.

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Fig. 4. The magnetization curves of Sn1 xMnxO2 films on at room temperature. The inset (a) and (b) shows Mn content dependence of magnetization and coercivity of the films on Si (1 1 1) substrates, respectively.

decreases leads to enhancement in the antiferromagnetic contribution. In fact, for the Sn0.89Mn0.11O2 sample, no hysteresis behavior was observed in the M–H curves (not shown here). Since the doping limit has been achieved and the surplus Mn cannot occupy the Sn position any more, this induces the formation of an antiferromagnetic impurity phase (Mn2O3). Additionally, the dominant magnetic interaction between the incorporation Mn3+ ions into SnO2 lattice also becomes nearest-neighbour antiferromagnetic. Hence, it is reasonable to see a disappearance of RTFM for the sample. For the Sn0.987Mn0.013O2 film, the moment as large as 3.26mB/Mn is observed which is less than spin-only moment of Mn3+ ions (4.90mB/Mn). The smaller values of magnetic moment per Mn atom indicate that not all the Mn3+ ions contribute to the ferromagnetism. It is a key requirement for understanding RTFM in transitionmetal-doped semiconductor oxides to find out whether the magnetism originates from an inherent feature of DMSs or from the formation of a secondary phase. Some small ferromagnetic clusters may also induce RTFM. Fig. 5 presents zero-field-cooled (ZFC) and field-cooled (FC) magnetization curves for the Sn0.948Mn0.052O2 thin films on Si (1 1 1) substrates. As can be

Fig. 5. ZFC and FC magnetization variation as a function of temperature under 200 Oe for Sn0.948Co0.052O2 film on Si (1 1 1) substrate.

seen from the figure, there are no discernible peaks corresponding to small particles or clusters in either curve, revealing that our sample has no spin glass or cluster glass characteristics [30]. The difference between ZFC and FC magnetization indicates the hysteretic behavior [31]. In fact, it is impossible to form some Mn clusters under our experimental conditions (600 8C calcined 1 h in air). In addition, bulk Mn is antiferromagnetic and cannot account for the observed RTFM. In Mn-related oxides, Mn3O4 or Mn2SnO4 is the only possible impurity having FM phase because other Mn oxides (including Mn2O3) are antiferromagnetic. But this phase cannot contribute to the RTFM because the TC is 43 or 53 K, respectively; they may only be responsible for the increased ferromagnetism at low temperature. Even if some unidentified secondary phases were present, they could not account for the experimental results: (1) the value of magnetic moment per Mn decreases with the increase of Mn contents; (2) no ferromagnetic character is observed in Sn0.89Mn0.11O2. Based on the experiments, we can rule out possible magnetic contamination and the secondary phase of metallic Mn clusters and Mn-related oxides as the origins of the RTFM of Sn1 xMnxO2 (x  0.078) films. The observed RTFM in Mn-doped SnO2, therefore, must be intrinsic and does not originate from the secondary phase. The question now arises: what is the nature of the exchangemechanism leading to RTFM in the Sn1 xMnxO2 system? It is to be emphasized here that the VSM only looks at the global response of the sample and hence cannot alone ascertain or refute any claim for a possible exchange mechanism. Since the maximum measurable resistance with the resistivity option is 10 MV, a Keithley2000 Multimeter was used in the resistivity measurements (the maximum measurable value is 100 MV in the four probe resistance option). However, the resistances of our samples remain too large to be measured. From the distance between the two voltage electrodes and the area of cross-section for our sample, we can deduce that the resistivities for the samples are at least greater than 107 V cm at room temperature, which means that Mn doped the samples are ferromagnetism in the insulated state. Therefore, the free-carrier mediated RKKY interaction may not be responsible for the FM behaviors in our samples. The double-exchange mechanism between d states of transition metal elements is another possible candidate to induce ferromagnetism in magnetic oxides. However, XPS data (see Fig. 1) indicate the single valence state of Mn (i.e., Mn3+) in SnO2, which excludes the doubleexchange mechanism. Considering the high resistivity (close to insulator) of our samples, we prefer to interpret the origin of RTFM in terms of the percolation of bound magnetic polarons (BMP) [32]. In this model, the ferromagnetism may arise from the interaction of bound magnetic polarons, which basically consist of a localized hole/electron and a large number of magnetic impurities around the localized charge. Though the direct exchange between localized holes is antiferromagnetic, the interaction between BMP may be FM-like at a large enough concentration of magnetic impurities. This model is applicable when the carrier concentration is much less than the concentration of magnetic impurities. In the present case, the conditions are favorable since the samples contain >1021 Mn ions cm 3, much greater than concentrations of the carriers. Thus, the overlapping polarons (or the so-called BMP) might correlate with and give rise to the FM contribution as is experimentally observed (Fig. 4). However, at the same time, the isolated polarons can contribute to the PM fraction, i.e. those noninteracting Mn ions, not taking part in forming the BMPs, account for the PM/AFM fraction. It has also been reported by X-ray magnetic circular dichroism measurements on Zn0.95Co0.05O thin films that Co ions have a similar electronic structure in the PM and FM phases, and only a part of the doped Co ions are ferromagnetic [33]. Although, for both, FM as well as PM/AFM, Mn3+ are in the

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that the Sn1 xMnxO2 thin films are ferromagnetic at room temperature, and the ferromagnetism shows intrinsic characteristic. Magnetization data shows that the average magnetic moment per Mn atom decreases and the coercivity increases with increasing Mn content. The percolation of BMP is suggested to be the origin of the RTFM. Although it remains challenging to utilize ferromagnetic Sn1 xMnxO2 thin films for practical applications, our experimental results prove that the sol–gel method is an effective method for fabrication of TM-doped SnO2 nanostructures with RTFM by chemical synthesis. Acknowledgement This work is supported by Grant Nos. 50371034 from the National Natural Science Foundation of China. References Fig. 6. The magnetization curves of Sn0.948Mn0.052O2 film at 5 and 300 K.

octahedral crystal field, subtle differences such as the kind of nearest neighbour (Sn or Mn), neighbour defects and local lattice distortions may lead to different magnetic behavior as suggested by Kobayashi et al. [33]. Further studies based on microscopic investigations are required to validate any such claims of the ferromagnetic ordering in Sn1 xMnxO2. Fig. 6 shows the temperature dependent hysteresis loops of Sn0.948Mn0.052O2 thin films at T = 300 and 5 K in the SQUID magnetometer. It is seen that the significant paramagnetic contribution from magnetic ions is evidenced by the portion of nonsaturation magnetization at 5 K. The percolation theory of ferromagnetism based on BMP predicts that a small fraction of Mn spins participate in percolation volume to settle in a ferromagnetic state [34]. The rest do not participate in ferromagnetism. These free spins align either at very high fields or at very low temperatures. Thus, one expects a rise in magnetization both with the application of higher fields and at very low temperature, as has been confirmed through simulations [34]. Comparing with other preparation methods, such as PLD, molecular beam epitaxy (MBE), sputtering and so on, the sol–gel method is lower-cost and easier to control the composition of DMS. In addition, the solution made by chemical technique could help a homogeneous distribution of Mn2+ ions throughout the films. For all samples, the calcining at 600 8C in air provided a rich-oxygen atmosphere. Under this condition, we could exclude the existence of the metallic clusters in a large part. At this point, our experimental results prove that the sol–gel method is an effective method for fabrication the TM-doped SnO2 nanostructures with RTFM. The observation of the FM in the Sn1 xMnxO2 thin films by the preparation route is of scientific significance and technological value. This is an important progress in sample fabrication.

[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]

4. Conclusions

[28]

We successfully fabricated Mn-doped SnO2 films on Si (1 1 1) substrates with the rutile structure by sol–gel method (x  0.078) and investigated the crystal structure and magnetic properties of the films by means of different measurements. Evolution of cell parameters and XPS studies confirmed the incorporation of Mn3+ cations into rutile SnO2 lattice. Magnetic measurements revealed

[29] [30] [31] [32] [33] [34]

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