Enhanced ferromagnetism in nano-sized Zn0.95Mn0.05O grains

ARTICLE IN PRESS Journal of Magnetism and Magnetic Materials 322 (2010) 2340–2349

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Enhanced ferromagnetism in nano-sized Zn0.95Mn0.05O grains R.N. Bhowmik a,n, Asok Poddar b, A. Saravanan a a b

Department of Physics, Pondicherry University, R. Venkataraman Nagar, Kalapet, Pondicherry-605014, India Experimental Condensed Matter Physics, Saha Institute of Nuclear Physics, 1/AF Bidhannagar, Kolkata-700064, India

a r t i c l e in f o

a b s t r a c t

Article history: Received 20 January 2010 Received in revised form 2 February 2010 Available online 18 February 2010

The present work reports ferromagnetism by doping magnetic Mn atoms in the diamagnetic ZnO matrix and the ferromagnetism has been extended up to 640 K in nano-grained Zn0.95Mn0.05O samples. The bulk and nano-grained samples were stabilized in hexagonal crystal structure with space group p63mc. The grain size and lattice strain of the samples were estimated from room temperature XRD spectrum. Surface morphology of the samples was examined at room temperature using SEM picture and EDX spectrum. The ferromagnetism of the bulk material shows enhancement in nano-grained samples, which was mainly due to the solution of Mn atoms into the lattice sites of ZnO by mechanical milling. The enhancement of magnetic moment and ferromagnetic ordering temperature with reduction in grain size has been understood in terms of the core–shell structure and existing theoretical models. The present work also demonstrated the role of surface spin disorder on the enhancement of ferromagnetism in Zn0.95Mn0.05O nanograins. & 2010 Elsevier B.V. All rights reserved.

Keywords: Ferromagnetism in semiconductor Mechanical milled nanomaterial Magnetic nanoparticle Core–shell structure.

1. Introduction Dilute magnetic semiconductors (DMS) are the solid solution of magnetic elements (Mn, Fe, Co and Ni) in non-magnetic semiconductor (e.g., ZnO) matrix. Recently, DMS proved to be a promising candidate for potential applications in spintronic devices [1,2]. Dietl et al. [3] predicted that 5% Mn doping in ZnO can exhibit ferromagnetism with Curie temperature (TC) above room temperature, although ZnO is a typical diamagnetic semiconductor. The multifunctional properties in ZnO based DMS can be applicable in many technological devices [4–7]. The interesting aspect of studying such materials is to understand whether the room temperature ferromagnetism (RTF) is intrinsic or extrinsic to the doping effect. There is no clear mechanism for the origin of ferromagnetism in Mn doped ZnO (i.e., Mn–ZnO). Different types of magnetic ordering (ferromagnetic, antiferromagnetic, paramagnetic and spin glass) at room temperature (300 K) were observed in previous reports [8–14]. Different views were presented (e.g., carrier mediated ferromagnetism) to realize the magnetism of ZnO based DMS. Monte Carlo simulation [11] assumed the competition between two indirect exchange interactions, i.e., carrier mediated ferromagnetic interactions and antiferromagnetic superexchange interactions. The simulation suggested

n

Corresponding author. Tel.: + 91 9944064547; fax: + 91 413 2655734. E-mail addresses: [email protected], [email protected] (R.N. Bhowmik). 0304-8853/$ - see front matter & 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.jmmm.2010.02.035

that p-type Mn–ZnO (Mn up to 20%) can exhibit paramagnetism for hole concentration from 1  1016 to 1  1018 cm  3 and ferromagnetism for hole concentrations 1  1019 cm  3 and above. Similar magnetic nature was also predicted by Dietl et al. [3] and Spaldin [15]. Experimentally, the carrier mediated ferromagnetism was verified by Ivill et al. [10] from the co-doping effect in ZnO matrix. On the other hand, antiferromagnetism [12], spin glass [16], and even no phase transition [11] were suggested for n-type Mn doped ZnO samples. Many other mechanisms, e.g., vacancy of cation (Zn) [17–19] and coexisting impurity phases [13,20,21], also came into discussion as the source of RTF in Mn–ZnO. Otherwise, the material was suggested to be paramagnetic down to 5 K [22]. But the role of secondary phase in showing RTF in Mn doped ZnO is not consistent in the literature. In fact, the preparation technique for most of the samples, where the impurity phase seems to play a significant role, is not straightforward. Jayakumar et al. [9] ruled out the role of secondary phase. They showed that RTF in Mn–ZnO nanomaterial arises due to the doping of Mn2 + ions in the matrix of Zn2 + ions. It is well established that magnetism in Mn–ZnO strongly depends on the routes and conditions of material synthesis, irrespective of micron size (bulk) or nano-size particles [17,21–23]. Hence, proper understanding of the structure and size-dependent properties is essential before applying DMS as the excellent material for next generation spintronic devices. In this case the route of material synthesis is not much important, but the reproducibility and stabilization of reported RTF are essential. Most of the DMS were synthesized following a bottom-up approach, i.e., chemical synthesis followed by thermal heating [24]. The properties

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of such DMS are not consistent with each other and also not free from structural defects, impurity phases and substrate effects, especially for thin films. They need special atmospheric environment, which is not convenient for accurate control during material synthesis. On the other hand, top-down approaches such as mechanical milling and alloying have been recognized as a powerful tool for producing a variety of ZnO based DMS nanomaterials [25,26]. This technique has effectively shown RTF in the milled powder of Mn– ZnO, which needs further experimental works for its reproducibility and to understand their structural changes and physical properties. This paper deals with the synthesis of Zn0.95Mn0.05O samples using mechanical milling and presents their structural and ferromagnetic aspects. Noting the room temperature ferromagnetism in bulk and nano-grained samples, an attempt was made to record the ferromagnetic to paramagnetic ordering temperature (TC) that is less explored in the literature. Mechanical alloying, a different technique compared to mechanical milling, was also applied to crosscheck the reproducibility of the basic ferromagnetic properties of the mechanical routed Zn0.95Mn0.05O samples. The experimental results will be compared with reported works and features will be understood in the framework of existing models.

2. Experimental 2.1. Sample preparation Bulk Zn0.95Mn0.05O (Mn–ZnO) sample was prepared using the high temperature solid-state reaction technique in argon atmosphere. Stoichiometric amounts of high purity ZnO (99.99%, Alpha Aesar) and MnO (99.99% Alpha Aesar) powders were mixed

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thoroughly and fired around 900–970 1C for 24 h. The fired material was slowly cooled to room temperature and subsequently ground for 2 h. The powder was made into pellets. Finally, the pellets were sintered at 970 1C for 36 h and furnace cooled. Polycrystalline phase of the final product was checked using an X-ray diffraction spectrum. Being sure about the structural phase purity of bulk Zn0.95Mn0.05O sample, Fritsch Planetary Mono Mill ’’Pulverisette 6’’ was used to mechanical mill the powder form of the bulk sample. The milling was carried out in argon atmosphere in an 80 ml agate vial with 10 mm agate balls. The ball to sample mass ratio was maintained at 7:1. For proper mixing, the milling process was stopped every 4 h. The milling was continued up to 60 h. During this process, the vial was opened at 10, 20 and 35 h to remove a small quantity of material for checking the structural phase and grain size reduction. The samples were denoted as MH10, MH20, MH35 and MH60 for milling times of 10, 20, 35 and 60 h, respectively. We also synthesized the Zn0.95Mn0.05O sample using mechanical alloying technique. Initially, stoichiometric amounts of high purity ZnO and MnO powders were mixed using mortar and pestle for 2 h. The mixed powder was mechanical milled up to 72 h with intermediate stopping. Similar conditions, e.g., argon atmospheres, balls and bowl, ball to sample mass ratio, milling speed, etc., as applied for mechanical milling of bulk sample were also maintained during mechanical alloying of MnO and ZnO to form the compound Zn0.95Mn0.05O. The alloyed sample with a milling time of 72 h was denoted as MA72. 2.2. Sample characterization and measurements The structural study of Zn0.95Mn0.05O samples was carried out by recording the X-ray diffraction (XRD) spectrum at room

Fig. 1. (a) EDX spectra and (b) XRD spectra for selected samples.

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5.229

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temperature using an X-ray diffractometer (model: X pert PANalytical). The XRD spectrum was recorded in the 2y range 10–901 with step size 0.011 using Cu-Ka radiation (wavelength ˚ The surface morphology of the samples was investi1.54056 A). gated by scanning electron microscopy (model: Hitachi S-3400N, Japan). The elemental composition of the samples was carried out using Energy Dispersive analysis of X-ray (EDX) spectrum (Nortan System Six, Thermo Electron Corporation Instrument Super DRY II, USA). The elemental distribution of the samples was checked from the point and shoot microanalysis at ten different points and also using elemental mapping over a selected zone. The EDX spectrum of each sample was recorded at 10 various points. DC magnetization of the samples was studied as a function of magnetic field and temperature using a vibrating sample magnetometer (Model: Lakeshore 7400), attached with a low temperature cryostat and a high temperature oven. The temperature dependence of magnetization was carried out at 2 kOe magnetic field by increasing the temperature from 100 to 780 K and field dependence of magnetization was investigated within 715 kOe at room temperature. It is worth mentioning that most of the room temperature measurements, e.g., X-ray diffraction study and field dependence of dc magnetization, were carried out in atmospheric condition. There is no structural and magnetic transformation in the samples in air.

V 47.6 3. Experimental results

47.4 0

20

40 Milling time (hour)

60

Fig. 2. (a) Variation of lattice parameters (a and c) and (b) unit cell volume (V) with milling hours for Zn0.95Mn0.05O samples. The dotted lines guide the general trend of a and V.

3.1. Structural study Elemental composition (Zn, Mn, O) of the samples is confirmed from the EDX spectrum (Fig. 1(a)). The EDX spectra of both milled and alloyed samples are identical with respect to the bulk Mn–ZnO sample. There is no significant contamination (Si atom) from the

Bulk Mn-ZnO

MH35

MH60

MA72

Fig. 3. SEM pictures of selected (bulk (a), MH35 (b), MH60 (c) and MA72 (d)) samples.

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agate balls and container. The elemental composition in bulk and milled samples is close to that expected for the Zn0.95Mn0.05O compound. The atomic ratio of Zn, Mn and O is 47.53:2.47:50; 47.77:2.23:50; 47.64:2.36:50; 47.60:2.40:50; and 47.46:2.54:50 for bulk, MH20, MH35, MH60 and MA72 samples, respectively. Fig.1(b) shows the XRD spectrum of selected Zn0.95Mn0.05O (bulk, MH20, MH35, MH60 and MA72) samples. The XRD patterns of both bulk and mechanical (milled and alloyed) samples are identical (i.e., position, number and shape symmetry of the peaks). Based on the log scale plot (not shown in the figure) of XRD intensity, we conclude there is no significant additional phase in the bulk and mechanical milled samples. We also did not observe any additional phase for the MA72 (mechanical alloyed) sample, except for a minor signature of starting MnO phase at about 2y  43.421. The presence of a non-reactive starting material (MnO) in the MA72 sample is very small. The crystalline cell parameters of the samples were obtained by matching the XRD peaks to the single-phased hexagonal crystal structure with space group p63mc. The lattice parameters and unit cell volume of the Zn0.95Mn0.05O samples are shown in Fig. 2. The lattice parameters of the bulk sample are ˚ The general trend of mechanical a=b=3.247 A˚ and c=5.2083 A. milled samples is that lattice parameter(a and b) clearly increases with milling time whereas c shows a non-monotonic increase. Overall, the unit cell volume increases with milling time. The increasing effect of cell parameters is significant for the MA72 sample. The increase of lattice parameters and cell volume may be small, but for the 5% Mn doped sample this increase is significant and definitely suggests the incorporation of Mn atoms ˚ into the Zn (Zn2 + =0.74 A) ˚ sites of the ZnO crystal (Mn2 + =0.80 A) structure [14,27,28]. The cell expansion process is enhanced by the milling induced effect (grain size reduction and lattice strain) [29] and gives an indirect evidence of better solution of Mn atoms in ZnO matrix [28]. The simultaneous effects of Mn doping and grain size reduction on the lattice expansion are further supported by the observation of significantly large cell parameters in the mechanical alloyed (MA72) sample. We employed the Williamson–Hall method [30] to estimate accurate grain size for the mechanical milled samples where an appreciable amount of mechanical strain induced effect increases the broadening of XRD peaks [31]. We used the forms bL =0.89l/(od4cos yC) and b2G =8p(tan2 yC)(erms)2 to calculate the grain size od4 and root mean square lattice strain erms, respectively. Here, l is wavelength of X-ray radiation ˚ bL and bG are the Lorentzian and Gaussian (1.54056 A). components of integral width bhkl (defined as the peak area divided by peak height) of (h k l) peak (in degrees) with the centre at 2yC (in degrees). bhkl was obtained for 6 to 7 prominent XRD peaks by fitting each peak profile to a Pseudo-Voigt function, consisting of Gaussian and Lorentzian components. To extract meaningful information on XRD peak broadening due to simultaneous size and strain effects, it is assumed in the Williamson–Hall method that the Lorentzian (L) component can be attributed solely to grain size (crystallite size) effect and Gaussian (G) component can be attributed to lattice strain effect (the lattice strain is assumed to be non-uniform in the mechanical milled material). The effective Lorentzian (bL) and Gaussian (bG) components of bhkl were obtained from bL = bhkl–b0 and  1=2 bG ¼ b2hkl b0 2 , respectively. b0 is the integral width of standard silicon powder (0.1231) taken as the measurement of instrumental broadening. We found that grain size calculation is not completely isolated from the strain induced effect, because bL cos yC was not totally independent of cos yC. However, lattice strain was calculated in a more accurate manner. From the best values of grain size (od4) and lattice strain (erms), we observed that grain size of the material continued to decrease (od4  75, 55, 40, 31 nm for MH10, MH20, MH35 and MH60, respectively) and

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lattice strain (erms  0.53, 0.57, 0.62, 0.65 for MH10, MH20, MH35 and MH60, respectively) continued to increase with the increase in milling time. The od4 and erms for the MA72 sample are nearly 35 nm and 0.62, respectively, and consistent with the range of milled samples. Furthermore, structural aspects (particle size, surface morphology and spatial distribution of elements) were studied using SEM pictures of the samples. The SEM pictures (using 200 nm–5 mm scale) of the selected samples are shown in Fig. 3(a–d). The SEM picture (Fig.3(a)) confirms the hexagonal crystal structure for bulk sample with grain sizes  1.84–4.37 mm. The hexagonal shape of the crystallites transforms into a more or less spherical one for milled samples. From SEM pictures (Fig. 3(b–d)) the particle size for the MH35 sample is 200–435 nm, for the MH60 sample it is  100–130 nm and for the MA72 sample it is  50–80 nm. It is true that the estimated particle size (from SEM pictures) is larger than the grain size from XRD data, but the pattern of reducing the grain (single crystallite domain) size with milling time is also confirmed from the SEM pictures. At the same time, SEM pictures suggest the size of multi-grained particles or agglomeration of many single domain particles. The grain sizes estimated from XRD data will be used in future discussion. In the absence of transmission electron microscopic (TEM) data, we captured a typical particle of the MH35 sample using the SEM picture. The SEM picture (Fig. 4) indicates a clear core (300 nm)—–shell ( 40 nm) type morphology in multi-grained particle of milled samples. The line scan of EDX spectrum over a length of 45 mm (Figs. 5(a–c)) suggests a reasonably good homogeneity of elements (Zn, Mn, O) distribution (within the experimental error due to varying surface roughnesses) in the milled samples. More interestingly, the mapping of elements, as shown in Fig. 5(d) for the MH60 sample, indicated no appreciable clustering or segregation of Mn atoms in the ZnO matrix. This suggests the doping of Mn into Zn sites of milled samples and also rules out any significant role of Mn clustering on the observed RTF [8]. After obtaining a sufficient amount of structural information on the samples, we try to understand the doping effect of Mn atoms into ZnO matrix through magnetic properties of the bulk and nano-grained milled samples. 3.2. Magnetic study First, we identified the magnetic properties of the starting materials, i.e., ZnO, MnO, mixture of 5% MnO and 95% ZnO, and also Mn powder for reference purpose, using field (H) dependence

Fig. 4. Core–shell type structure of a particle obtained from SEM picture of MH35 sample.

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Fig. 5. (a–c) Elemental scan [O(red), Mn (green), Zn (blue-K line and magenta-L line) as indicated in (c) for MA72 sample] over the length 0–45 mm. (d) Elemental mapping for MH60 sample. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)

of magnetization (M) measurement at room temperature (300 K). The experimental data are shown in Fig.6. For better representation, the M(H) data of ZnO, a mixture of 5% MnO and 95% ZnO and bulk Mn–ZnO samples are shown in low field range (Fig. 6a). The negative slope in M(H) confirmed the diamagnetic properties of the bulk ZnO sample (means no ferromagnetic contribution present). The mixture of 5% MnO and 95% ZnO showed only paramagnetic character. This means the magnetic properties of the mixture are mainly dominated by the paramagnetism of 5% MnO over the diamagnetism of 95% ZnO and there is no signature ferromagnetic contribution. The striking feature is that the solid state routed bulk Zn0.95Mn0.05O (Mn–ZnO) sample exhibited a clear ferromagnetic hysteresis loop in the low field range 71 kOe (Fig. 6a) and an unusual (up-curvature with field) increase of magnetization without any loop at the higher

magnetic field (inset of Fig. 6b). Fig. 6b shows an expected paramagnetic behaviour for MnO (TN  120 K [23,32]) and ferromagnetic loop for Mn powder. It is interesting to note that the ferromagnetism of Mn–ZnO bulk sample neither matches the field response of magnetization of any starting component nor is consistent with the observed ferromagnetic feature of Mn powder (Fig. 6b). Hence, we exclude the possibility of segregated ferromagnetic Mn atoms in our Mn doped ZnO sample. The spontaneous (ferromagnetic) magnetization (MS  0.006 emu/g) of Mn–ZnO bulk sample is obtained from the polynomial fit of M(H) data at Ho7 kOe. On the other hand, the extrapolation of (positive) high field M(H) data intercept to the negative M axis and negative high field M(H) data to positive M axis (inset of Fig. 6b). The M(H) data at HZ + 7 kOe nearly fit to the linear equation ML(H)= M0 + w0H with M0 =  0.006 emu/g and

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0.2 5% MnO+ 95% ZnO mixture

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Mn pow der

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Fig. 7. Field dependence of magnetization for different samples (main panel) and inset shows the M(H) data (0–15 kOe range) after subtracting linear component from high field (47 kOe) data.

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Fig. 6. M(H) behaviour at 300 K (a) for ZnO, mixture of ZnO and MnO, bulk Mn– ZnO samples within low field range and (b) for MnO and Mn powder up to 15 kOe. Inset shows the field induced magnetism for bulk Mn–ZnO sample above 7 kOe.

w0 = 2.46  10  6 emu/g/Oe for positive H values. Considering a very unusual increase of M(H) curve for bulk Mn–ZnO sample with the larger slope at higher field than at smaller fields, the negative value of M0 is justified. The negative magnetization (M0) should not be treated as the negative spontaneous magnetization. Rather, it represents some kind of high field induced paramagnetic component whose contribution per gram of bulk Mn–ZnO sample increases with field. More precisely, a high field induced paramagnetic component is superimposed on the low field ferromagnetism of bulk Mn–ZnO sample. The room temperature M(H) of bulk sample resembles the modulated magnetic order that appeared due to the solid solution of high magnetic Fe moments (from Fe2O3 canted ferromagnet) into the sites of low magnetic Cr moments (from Cr2O3 paramagnet) [30]. In the present sample, the ferromagnetic phase of Mn atoms mixes into the diamagnetic phase of Zn sites. The unusual increase (up-curvature) of the M(H) curve of Mn–ZnO bulk sample at higher magnetic fields suggests that the magnetic solution between ferromagnetic Mn atoms and diamagnetic Zn atoms is not fully completed in the solid state routed bulk sample. Rather, a significant fraction of Mn spins were pinned in the diamagnetic matrix and by de-pinning them at higher field magnetic contribution increases in the material. Although structurally we could see a good solid solution of Mn atoms in ZnO matrix without any significant impurity phase, there is a scope of improving the ferromagnetic order by increasing the

magnetic solubility of Mn atoms into Zn sites and reducing the de-pinning effect of Mn spins. Now, we understand the room temperature ferromagnetism of nano-grained mechanical milled samples. Fig. 7 shows that field dependence of magnetization [M(H)] of milled samples is drastically different from the Mn–ZnO bulk sample, as well as from the ferromagnetic Mn powder. The M(H) isotherm at relatively higher fields [HZ7 kOe] linearly increases with field (ML(H) =M0 + w0H) for milled samples. Now, the intercept of the straight line ML(H) on magnetization axis gives a positive value for M0 that is equal to the spontaneous magnetization (MS) of the sample and w0H is the paramagnetic contribution superimposed on the ferromagnetic order. The estimated values of paramagnetic susceptibility (w0) are  1.79, 4.22, 4.21, 2.67 and 1.09 (in 10  6 emu/g/Oe unit) for MH10, MH20, MH35, MH60 and MA72 samples, respectively. We estimated an effective ferromagnetic Meff(H) curve (inset of Fig. 7) by subtracting the paramagnetic w0H (H= 0–15 kOe) term from the measured M(H) curve. The striking feature is that spontaneous magnetization (i.e., ferromagnetic order parameter) has been significantly enhanced by reducing the grain size of the Mn doped ZnO bulk sample. The obtained values of MS and other magnetic parameters from the M(H) loops, such as coercivity (HC), remanent magnetization (MR) and loop area (S), are shown in Fig. 8. It is interesting to note that all the magnetic parameters (MS, HC, MR, S) show an increasing trend at the initial stage of mechanical milling whereas the parameters decrease at the higher milling time. The reported values of coercivity [33,34] are also reproduced in the present samples. The range of coercivity and hysteresis loop area (representing magnetic energy loss due to domain wall motion/domain rotation) suggest the soft ferromagnetic nature of the synthesized material. It may be noted that the magnetic parameters of MA72 sample are significantly lower in comparison with the MH60 sample, suggesting

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Fig. 8. Variation of magnetic parameters, i.e., (a) spontaneous magnetization (MS), (b) coercivity (HC), (c) remanent magnetization(MR), and (d) hysteris loop area, with milling time and calculated M(H) data at 300 K.

more soft ferromagnetic properties of the MA72 sample. The magnetic parameters of the MA72 sample are within the range of mechanical milled samples, although the mechanical alloying technique used to produce the MA72 sample is significantly different from the mechanical milled samples. We would like to mention that time dependence of magnetization for MH60 sample (data not shown) showed a gradual increase of magnetization with time and was nearly time independent after 30 min. This confirmed a typical ferromagnetic character at room temperature with finite disorder in the material. Now, we examine the ferromagnetic order of Zn0.95Mn0.05O samples below and above room temperature (100–800 K). The temperature dependence of dc magnetization curves of the samples at 2 kOe (Fig. 9) decreased continuously when the temperature increased up to the ferromagnetic to paramagnetic transition temperature (TC), as indicated by arrows for each sample. One could see a significant transformation in the temperature dependent magnetization of milled samples compared with the character of the Mn–ZnO bulk sample. A slight up-curvature in the Mn–ZnO bulk sample on lowering the

temperature below 300 K indicates the coexistence of a paramagnetic contribution with ferromagnetism and similar coexisting magnetic phases as also previously indicated [35]. The shape of the M(T) curve is completely changed to downcurvature for the mechanical milled samples and shows a dominated ferromagnetic order in the nano-grained samples. We noted that TC (  500 K) of the Mn–ZnO bulk sample is comparable to the predicted value above 425 K [8]. It is highly remarkable that TC of the material increases (  595, 605, 620 and 640 K for MH10, MH20, MH35 and MH60 samples, respectively) by reducing the grain size using mechanical milling. It is interesting to note that the ferromagnetic order above room temperature is also reproduced in the MA72 sample with TC  615 K and is comparable in the range of mechanical milled samples. The increase of TC with the decrease of grain size in the present mechanical samples is different from the decreasing trend in chemical routed samples [36]. Some reports [20] even attempted to compare the magnetism of mechanical milled samples to that of chemical routed samples. Such comparison may not be appropriate, because surface magnetism plays a

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Fig. 9. Temperature dependence of zero field cooled magnetization at 2 kOe for different Zn0.95 Mn0.0 O samples. The arrows mark TC of the respective sample.

significant role in controlling the nano-grained magnetism in mechanical activated samples [37].

4. Discussion Understanding ferromagnetism at and above room temperature in Mn doped ZnO is a matter of recent interest. ZnxMn3  xO4 type additional impurity phase was clearly visible in the XRD spectrum of some works [13,20,21,38], which attributed ferromagnetism in Mn doped ZnO due to impurity phases. It is highly relevant to mention that the ZnxMn3  xO4 spinel phase is paramagnetic at T Z300 K [39]. It is very unlikely that paramagnetic impurity phase can be the cause for ferromagnetism above room temperature in Mn doped ZnO. Interestingly, those works used MnO2 and ZnO as the starting materials for synthesizing Mn doped ZnO compound and various magnetic features were noted at room temperature [33,34,40,41]. In our experiments, we used ZnO and MnO as the starting materials and the XRD spectrum of bulk and milled samples of Mn doped ZnO does not show any significant additional phase. The additional phases, if they really exist at all, are well below the detection limit from XRD spectrum and such minute effect of an additional phase cannot be the origin of enhanced ferromagnetism of the present

samples. For the sake of argument if we assume that some ferromagnetic impurity phase (which is below the XRD detection limit) increases, one must expect a monotonic increase of magnetic parameters (MR and MS) with the increase of milling time. However, experimental results are not consistent with this assumption of increasing ferromagnetic impurity with milling time. The mechanical alloyed sample (starting with ZnO and MnO powders) also reproduces the structural phase purity and room temperature ferromagnetism, although a small signature of MnO (starting material) was detected from the XRD spectrum of the MA72 sample. The incomplete solution of MnO in the ZnO matrix most probably is the reason as to why magnetic parameters of the MA72 sample are lower than those of the MH60 sample. However, one thing that is clear is that the present work effectively shows that mechanical activated (milling and alloying) technique is useful in producing nanomaterials of DMS. It excludes the complexity in material synthesis and phase destabilization related to the high temperature annealing effects when using other routes [13,20,21]. The non-magnetic contamination (Si) is minimized below 1 at% and it has no role in producing RTF. The role of different atmospheres on the generation of RTF in Mn–ZnO samples is also inconclusive in the literature [23,33,34]. We believe that the basic origin of RTF in transition metal doped ZnO is something else, which is not clear till date and cannot be

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attributed to the impurity phase alone. We proceed to search for an alternative mechanism of room temperature ferromagnetism in Mn doped ZnO. In this work, we observed an unconventional field dependent magnetic order in Mn–ZnO bulk sample, consisting of a ferromagnetic order saturated within 1 kOe field and an induced paramagnetic order at higher field ( Z7 kOe). The coexistence of ferromagnetism and paramagnetism is consistent with the theoretical prediction [3,8,11,18] for Mn doped ZnO. The doping of Mn atoms into Zn sites of ZnO was checked from XRD data. For the sake of argument, if there was no Mn doping into Zn sites before and after milling of the Mn–ZnO bulk sample, only a paramagnetic phase of the mixture of 95% ZnO and 5% MnO was expected. However, room temperature M(H) measurement of the Mn–ZnO bulk sample showed ferromagnetic spontaneous moment  0.006 emu/g–0.0017 mB/Mn atom and the magnetic moment at 15 kOe including paramagnetic component is  0.032 emu/g–9.0  10  3 mB/Mn atom. These magnetic moment values are smaller than the 5 mB/Mn atom at 0 K for high spin Mn2 + states [42]. One reason for the reduced magnetic moment in bulk Mn–ZnO could be the higher measurement temperature (300 K). The second reason is the quenching of Mn moment in the nonfavorable (diamagnetic) environment of Zn. The alloying of ferromagnetic a-Fe2O3 and paramagnetic Cr2O3 [30], and the formation of the ferromagnetic Co core surrounded by antiferromagnetic CoO shell or paramagnetic Al2O3 [43] may be good examples that showed a strong environmental affect on the magnetic ordering of high magnetic moments. Undoubtedly, the core–shell model can describe the magnetism of nano-sized grains or particles in a better way. The concept of core–shell model was used [8] to explain the ferromagnetism in Mn doped ZnO particles by assuming the fact that the Mn rich (segregated) ferromagnetic layer surrounds the non-magnetic ZnO core. We argue that an Mn rich additional phase sitting on the grain boundary of nano-grained samples might not be seen from the XRD spectrum [28], but the Mn rich secondary phase, if it really exists, could be expected in the XRD spectrum of Mn–ZnO bulk samples. This is due to the fact that grain boundary is relatively small in bulk samples (coherent crystalline zone  4 mm from SEM picture) and a part of the crystalline zone must have a Mn rich layer. In our bulk Mn–ZnO sample, neither the secondary phase was prominent nor was a strong ferromagnetic moment noted. The experimental observations of enhanced ferromagnetic moment and ordering temperature in nano-grained samples may be related to the increasing solubility of Mn atoms in the ZnO matrix by decreasing the grain size of the material [28] or by increasing the disorder in DMS particles [44]. Although magnetic mechanism in thin films is related to many concurrent effects, a recent report [45] suggested that ferromagnetism in Mn doped ZnO can be better understood by assuming the formation of a shell type ferromagnetic grain boundary over non-magnetic grain. In the present work, we propose an alternate explanation based on the core–shell model and taking into account the disorder effects of nanograins for the enhancement of ferromagnetism in mechanical milled Mn–ZnO samples. The application of core–shell gives a realistic explanation of the experimental observations. We explain the lower magnetic moment in bulk sample considering the fact that Mn atoms may be dissolved structurally into the ZnO matrix, but magnetically it is an inhomogeneous mixture of weak ferromagnetism and large paramagnetism throughout the grain and negligible grain boundary. This means magnetically the bulk material is not the perfect solid solution. Some Mn moments are pinned in diamagnetic ZnO matrix. For nano-grained samples, the shell or grain boundary part increases, as well as disorder effect also that may act as the de-pinning field to Mn moments in ZnO matrix. The core is

essentially less ferromagnetic Zn0.95Mn0.05O (i.e., a mixture of weak ferromagnetic and large paramagnetic components as in bulk) surrounded by the shell of more ferromagnetic Zn0.95Mn0.05O (i.e., ferromagnetic part is enhanced by the disorder activated grain boundary and better magnetic solution at the interface of shell and core). Therefore, the decrease of grain size by mechanical milling induces more ferromagnetic ordering at the shell and also at the interface of core and shell. These combined effects result in the increase of magnetic parameter at the initial stage of milling. The shell disorder increases at lower grain size (i.e., higher milling time), which reduces effective ferromagnetic moment in the shell on the one hand and continued ferromagnetic enhancement at the interface of core and shell on the other hand. This resulted in the decrease of room temperature magnetic parameters (MS, HC, MR and loop area) for samples with higher milling time, but the magnetic parameters were definitely enhanced in comparison with un-milled bulk samples. The magnetic parameters of the MA72 sample were relatively small due to increased shell disorder [29]. The experimental results suggest that increasing shell disorder is favorable for enhanced ferromagnetic interactions among Mn atoms [18,33,46–48]. We suggest that disorder (cation vacancy and other lattice defects as suggested in Refs. [17–19]) in shell structure can promote the required number of hole carriers, as predicted for carrier mediated ferromagnetism in Mn doped ZnO [3,8,11]. However, the facts that (95%) Zn atoms diffuse into the surface of (5%) MnO2 particles [13] and double exchange type mechanism is responsible for RTF in Mn doped ZnO [34,38,40] are not consistent with our experimental results. This is due to the fact that double exchange mechanism showed the maximum TC in perovskite compound up to 370 K [46]; whereas the TC in our samples is up to 640 K, it is up to 980 K as reported by another group [13]. On the other hand, ferromagnetism with TC up to 620 K was noted in 0 0 double perovskite (Sr2BB O6; B= Fe, Cr; B =Mo, Re) materials [49,50], where super exchange type interactions between cations 0 among B and B sublattices via O anion determined the long range ferromagnetic ordering temperature, inspite of increasing disorder in the lattice structure and decreasing magnetic moment in the material. Monte Carlo simulation [11] also considered the competition between antiferromagnetic superexchange interactions and oscillating carrier mediated interactions to explain the high temperature ferromagnetism in ZnO based DMS materials. Further experimental and theoretical studies are needed to confirm the effects of superexchange interactions in DMS.

5. Conclusion Mechanical milling and alloying are capable of producing nano-grained Zn0.95Mn0.05O (DMS) without showing any significant additional phase. Ferromagnetism at and above room temperature is confirmed in Mn doped ZnO. We believe that secondary phase, if really present below the detection limit, does not play any significant role in the exhibition of ferromagnetism in Mn doped ZnO samples. The coexisting paramagnetic contributions account for the quenching of magnetic moment per Mn atom in micron size bulk Zn0.95Mn0.05O. The magnetic solution of Mn atoms into the Zn sites of ZnO lattice structure is enhanced by reducing the grain size using mechanical milling. Although a challenging problem, the generalization of the mechanisms of high temperature ferromagnetism in different magnetic oxides (DMS, double perovskites) would be interesting. It needs many theoretical models and further experimental works in future. The core–shell model, as introduced in the present work, suggests that the (core–shell) interfacial disorder of nanograins affects the solution of Mn atoms into the Zn sites of the ZnO structure.

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Acknowledgments The authors thank CIF, Pondicherry University, for providing the experimental facilities. RNB and AS also thank UGC for financial support [F.NO. 33-5/2007 (SR)]. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13]

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