Open Access

Magnetic modulation in mechanical alloyed Cr1.4Fe0.6O3oxide

PMC Physics B20081:20

DOI: 10.1186/1754-0429-1-20

Received: 22 October 2008

Accepted: 29 December 2008

Published: 29 December 2008

Abstract

The compound Cr1.4Fe0.6O3 has been synthesized through mechanical alloying of Cr2O3 and α-Fe2O3 powders and subsequent thermal annealing. The XRD spectrum, SEM picture and microanalysis of EDAX spectrum have been used to understand the structural evolution during alloy formation. The alloyed samples have been matched to rhombohedral structure with R3C space group. The observation of a modulated magnetic order confirmed the systematic diffusion of Fe atoms into the Cr sites of lattice structure. A field induced magnetic behaviour has been noted in the field dependence of magnetization data of the annealed samples. This feature is significantly different from that of the mechanical alloyed samples. The experimental results also provided the indications of considering the present material as a potential candidate for opto-electronic applications.

PACS Codes: 75.30.-m, 75.20.En, 81.40.Rs

1. Introduction

In recent years extensive research efforts have been given to the formation of new magnetic oxides, considering their potential applications in the field of micro-electronics and multifunctional devices [1, 2]. In order to search for the new class of materials, it has been found that solid solution of mixed metal oxides might be a potential candidate, having advantage of easy alloying due to ionic radius of same order [3, 4]. Fe2-xCrxO3 is one such compound series, which can be formed through the alloying of α-Fe2O3 and Cr2O3 oxides [5, 6]. The interesting points are that both α-Fe2O3 and Cr2O3 stabilized into rhombohedral crystal structure with space group R3C [4, 7] and both are antiferromagnetic insulator [8]. Cr2O3 oxide is a well known compound for the prediction and experimental observation of large magneto-electric (ME) effect [9, 10]. The ME effect is generally small in symmetric non-magnetic insulators [11]. The low magnetic moment and low antiferromagnetic ordering temperature (TN) of single phase Cr2O3 (~310 K) is also not suitable for obtaining large magneto-electric effect, as well as for applications [12, 13]. Recently, the exhibition of large magnet-electric effect in Cr2O3 was attributed [14] to a possible change of magnetic space group symmetry. This triggered the alloying of Fe2-xCrxO3series by mixing a suitable amount of Fe2O3 (TN ~ 950 K) [4, 8] in Cr2O3 oxide. The study of this series also remained attractive due to its promising applications in the field of opto-electronic materials. The immediate effect is that Fe2-xCrxO3 series has shown the enhancement in photo-conductivity with the increase of Cr content [2, 15]. In order to understand the modified photo-conductivity, as well as magneto-electric effect, one needs to realize the correlation between crystal structure and magnetic ordering in the compound. The correlation between its crystal structure and magnetism would also be relevant to gain the properties of Cr-Fe interface [2].

In literature, the antiferromagnetic ordering of spins in both α-Fe2O3 and Cr2O3 has been explained in terms of super exchange interactions (Cr-O-Cr, Fe-O-Fe) [4, 8]. Neutron diffraction study explained the drastic variation of TN (~310 K for Cr2O3 and ~950 K for Fe2O3) due to a different kind of magnetic structure, although α-Fe2O3 and Cr2O3 have shown identical crystal structure. Earlier reports [4, 7, 8] suggested that the spin moments of Cr3+ (3d3) ions in Cr2O3 are arranged in + - + - (+ ≡ up spin, - ≡ down spin) sequence along the [111] axis, whereas the spins of Fe3+ (3d5) ions in α-Fe2O3 are ordered in the + - - + sequence. Hence, it is expected that the substitution of Fe3+ (3d5) by Cr3+ (3d3) in Fe2-xCrxO3 solid solution would modify the sequence of spins ordering, as well as the nature of superexchange interactions. Form the phase diagram of Fe2-xCrxO3 [4] one could estimate an equal probability of Fe-O-Fe, Fe-O-Cr, Cr-O-Cr superexchange interactions near to x = 1 and no significant change in the ordering of spins. This resulted in a slow variation of TN with Cr content in the region 0.90 < x ≤ 1.52 [4]. However, the nature of spin ordering may not follow the conventional structure due to the change of magnetic space group symmetry. The modulated (perturbed) local magnetic order could be expected during the diffusion of Fe atoms into Cr sites. To our knowledge, the magnetic ordering of Fe2-xCrxO3 compound is not clear till date. Attempts were made to understand the structural and magnetic properties of Fe2-xCrxO3 compound by reducing the particle size in nanometer range using various chemical routes [16, 17]. However, enough attention was not given in the reported works to realize the effect of modulated spin structure on the properties of Fe2-xCrxO3compound.

In this work, Cr1.4Fe0.6O3 compound has been synthesized using the novel technique of Mechanical alloying [18]. Attempts have been made to investigate the associated structural and magnetic evolution in different states of mechanical alloying and annealing of the samples. The nature of magnetic order, i.e., antiferromagnet or ferromagnet or the mixed properties of ferromagnet and antiferromagnet, was understood for Cr1.4Fe0.6O3 compound.

2. Experimental

2.1. Sample preparation

The stoichiometric amounts of high purity α-Fe2O3 and Cr2O3 were mixed for the preparation of Cr1.4Fe0.6O3compound. The initial colours of α-Fe2O3 and Cr2O3 were red and green, respectively. The mixture of Fe2O3 and Cr2O3 was ground using agate mortar and pestle for nearly two hours in atmospheric conditions. The ground powder was mechanical alloyed using Fritsch planetary mono mill (pulverisette 7). The material and balls (combination of 10 mm Silicon Nitride and 5 mm Tungsten Carbide) mass ratio was maintained to 1:7. The mechanical alloying was carried out in atmospheric condition up to 84 hours in a silicon nitride (Si3N4) bowl with rotational speed 300 rpm. The non-magnetic balls and bowl (Silicon Nitride and Tungsten Carbide) were selected to avoid the magnetic contamination during the milling process. The milling process was intermediately stopped to monitor the uniform alloying of the mixture and to minimize the local heat generation that might occur during continuous milling. A small quantity of alloyed sample after 24 hours and 48 hours was taken out to check the structural phase evolution. The alloyed samples with different milling hours were made into pellets. The pellets of 84 hours milled sample were placed in alumina crucibles and annealed at 700°C in atmospheric conditions. After annealing for 1 hour, 3 hours and 17 hours, individual pellet was directly air quenched to room temperature. The mechanical alloyed samples was denoted as MAh, where h = 0, 24, 48 and 84 for alloying time 0, 24 hours, 48 hours and 84 hours, respectively. The samples annealed at 700°C were denoted as SNt, where t = 1, 3 and 17 for annealing time 1 hour, 3 hours and 17 hours, respectively.

2.2. Sample characterization and measurements

The crystalline phase of alloyed and annealed samples was characterized by recording the X-ray Diffraction spectra in the 2θ range 10–90° with step size 0.01°. The Cu-Kα radiation from the X-ray Diffractometer (model: X pert Panalytical) were employed to record the room temperature spectrum of each sample. The scanning electron microscopic (SEM) picture of the samples was taken using HITACHI S-3400N model. Elemental composition of the samples was obtained from the energy dispersive analysis of x-ray (EDX) spectrum. The magnetic properties of the samples were studied by the measurement of magnetization as a function of temperature and magnetic field using vibrating sample magnetometer (Model: Lakeshore 7400). The temperature (300 K–900 K) dependence of magnetization was measured by attaching a high temperature oven to the vibrating sample magnetometer. The temperature dependence of magnetization was carried out at 1 kOe magnetic field by increasing the temperature from 300 K to 900 K (ZFC mode) and reversing back the temperature to 300 K in the presence of same applied field 1 kOe (FC mode). It should be noted that the ZFC mode followed here is slightly different from the conventional zero field cooling (ZFC) measurement, where the sample is first cooled without applying magnetic field from the temperature greater than TC to the temperature lower than TC and magnetization measurement in the presence of magnetic field starts with the increase of temperature. The field dependence of magnetization at 300 K was measured within ± 15 kOe.

3. Results and discussion

3.1. Structural properties

Fig. 1 shows the XRD spectrum of mechanical alloyed and subsequent annealed samples. The MA0 sample (before milling) shows a mixed pattern of Fe2O3 and Cr2O3 phases. The peaks of Fe2O3 and Cr2O3 phases are well separated before mechanical alloying. As the alloying process continued by increasing the milling time, the individual peaks of Fe2O3 and Cr2O3 phases are merging to each other. The intensity of peaks is decreasing as the milling time increases up to 84 hours. At the same time broadness and symmetry in the shape of the peaks are increasing. Subsequent annealing of MA84 sample at 700°C, again, increases the peak intensity, along with increasing sharpness, with the increase of annealing time up to 17 hours. The undergoing phase evolution during the mechanical alloying and annealing process is clearly demonstrated in Fig. 2 for (104) and (110) peak positions. In the (MA0) mixed sample, the peaks position of Cr2O3 sample are at higher 2θ value with respect to Fe2O3 sample. The notable feature is that peaks of Fe2O3 sample disappear with the increase of milling time and the XRD pattern of alloyed compound is tending to achieve the character of Cr2O3. This indicates the diffusion of Fe atoms into the lattice sites of Cr atoms to form a solid solution of Cr1.4Fe0.6O3.
Figure 1

The evolution of XRD spectrum in alloyed sample with milling time and annealing time.

Figure 2

The evolution of spectrum of the alloyed samples for (104) and (110) XRD peaks.

Now, we understand the grain size refinement process. We have determined the grain size of alloyed samples from the prominent (012), (104), (110) and (116) XRD peaks using Debye-Scherrer formula: d = 0.089 * 180 * λ 3.14 * β * cos θ c n m MathType@MTEF@5@5@+=feaagaart1ev2aaatCvAUfKttLearuWrP9MDH5MBPbIqV92AaeXatLxBI9gBaebbnrfifHhDYfgasaacH8bkY=wiFfYlOipiY=Hhbbf9v8qqaqFr0xc9vqpe0di9q8qqpG0dHiVcFbIOFHK8Feei0lXdar=Jb9qqFfeaYRXxe9vr0=vr0=LqpWqaaeaabiGaciaacaqabeaabeqacmaaaOqaamaaamaabaGaemizaqgacaGLPmIaayPkJaGaeyypa0tcfa4aaSaaaeaacqaIWaamcqGGUaGlcqaIWaamcqaI4aaocqaI5aqocqGGQaGkcqaIXaqmcqaI4aaocqaIWaamcqGGQaGkcqaH7oaBaeaacqaIZaWmcqGGUaGlcqaIXaqmcqaI0aancqGGQaGkcqaHYoGycqGGQaGkcyGGJbWycqGGVbWBcqGGZbWCcqaH4oqCdaWgaaqaaiabdogaJbqabaaaaOGaemOBa4MaemyBa0gaaa@4BF5@ , where 2θC is the position of peak center, λ is the wavelength of X-ray radiation (1.54056 Å), β is the full width at the half maximum of peak height (in degrees). The 2θC and β values were calculated by fitting the XRD peak profile to Lorentzian shape. The average grain size of the alloyed compound is found to be in nanometer range, which suggested the nanocrystalline structure of the material. The grain size of the alloyed compound slowly decreases with the increase of milling time (i.e., <d> ~ 26 nm for MA24, 22 nm for MA48, and 15 nm for MA84 sample). The thermal annealing of MA84 (lowest grain size) sample at 700°C activates the increase of grain size (i.e., <d> ~ 19 nm, 21 nm and 22 nm for annealing time 1 hour, 3 hours and 17 hours, respectively). However, a finite amount of instrumental broadening and lattice strain may affect the XRD peak broadening in mechanical alloyed samples and hence, the grain size estimation based on the Debye-Scherrer formula may be far from good accuracy. To check these two factors, Williamson-Hall plot [19] was used. The Williamson-Hall equation is βeff cosθC = Kλ/<d> + 2ε sinθC, where K is Scherrer constant (0.89), ε is the lattice strain and βeff2 = β2 - β02. In this case, the 2θC and β values were calculated by fitting the XRD peak profile to Voigt shape and β0 is the full width at half maximum of standard Silicon powder. Finally, seven prominent XRD peaks were used for the βeff cosθC vs. 2sinθC plot. The linear extrapolation of this plot to βeff cosθC axis gives Kλ/<d> (and hence, grain size <d>) and the slope gives the lattice strain (ε). The calculated values of grain size and lattice strain are shown in Table 1. It is observed that the grain size calculated using Willimson-Hall plot is larger than that obtained from simple Debye-Scherrer formula, although the nature of the variation of grain size with milling time and annealing time is identical in both the cases. It is also observed that lattice strain of the alloyed compound increases with milling time and thermal annealing at 700°C reduces the lattice strain of the alloy. Similar effects were observed in other mechanical alloyed compound [20]. The lattice parameter of the samples (shown in Table 1) was calculated by matching the XRD peaks into rhombohedral structure with R3C space group. The lattice parameters (aand c) and cell volume of the alloyed compound were increasing with the milling time. But, the lattice parameters of the annealed samples showed decreasing values with the increase of annealing time at 700°C. The finite amount of mechanical milling induced lattice strain and its release by thermal annealing may have some effect on the variation of lattice parameters of the samples. However, the understanding of the variation of cell parameters is more appropriate from the fact that lattice parameters of bulk Cr2O3 sample (a= 4.9510 Ǻ, c= 13.5996Ǻ, cell volume = 288.69Ǻ3) is smaller than that for bulk α-Fe2O3 sample (a= 5.0386Ǻ, c= 13.7498Ǻ, cell volume = 302.3Ǻ3). The calculated lattice parameters for both bulk samples are consistent with literature values [4, 5]. The continuous increase of lattice parameters and cell volume with milling time indicates the kinetic diffusion of larger size (0.67 Ǻ) Fe3+ ions into the boundary of smaller size (0.65 Ǻ) Cr3+ ions [2, 4, 21]. On the other hand, the decrease of cell parameters with the increase of annealing time indicated that the alloying process is, still, continued to attain the crystal structure of Cr2O3 dominated phase, because atomic ratio of Cr and Fe is 7:3. Here, it is pointed out that thermal heating plays a major role for obtaining the Cr2O3 dominated phase. It is believed that the variation of cell parameters is not significantly affected by the nano-sized grains in comparison with the undergoing change during alloying and annealing process. The cell parameters and grain size of the MA0 sample were not calculated, as the pattern is not due to single phase XRD spectrum.
Table 1

Grain size (<d>), lattice strain (ε) based on Willimson-Hall method and cell parameters of the samples are calculated using XRD data.

Sample

Grain size (nm)

Latice-strain

Lattice parameter a(Ǻ)

Lattice parameter c(Ǻ)

Volume (Ǻ)3

α-Fe2O3

--

--

5.0386 ± 0.0014

13.7498 ± 0.0002

302.30 ± 0.0840

MA0

--

--

--

--

--

MA24

63

0.4817 ± 0.0756

4.9521 ± 0.0021

13.6063 ± 0.0004

288.96 ± 0.1221

MA48

46

0.5121 ± 0.0365

4.9535 ± 0.0016

13.6041 ± 0.0002

289.08 ± 0.0923

MA84

35

0.5771 ± 0.0604

4.9546 ± 0.0013

13.6082 ± 0.0003

289.29 ± 0.0781

SN1

55

0.4495 ± 0.0416

4.9526 ± 0.0019

13.6142 ± 0.0002

289.19 ± 0.1090

SN3

74

0.3709 ± 0.0118

4.9520 ± 0.0018

13.6084 ± 0.0002

288.99 ± 0.1099

SN17

86

0.3202 ± 0.0121

4.9512 ± 0.0012

13.6070 ± 0.0002

288.87 ± 0.0997

Cr2O3

--

--

4.9510 ± 0.0011

13.5996 ± 0.0001

288.69 ± 0.0667

The parameters are not calculated for MA0 sample, due to the coexistence of two XRD patterns.

The surface morphology (SEM picture) and chemical composition (EDX spectrum) can be studied for selected (MA0, MA48, MA84 and SN17) samples. The selected zone (~80 μm × 80 μm) of the samples, along with SEM picture and EDX spectrum, are shown in Fig. 3. The SEM picture of MA0 sample (Fig. 3a) suggests the heterogeneous mixed character of α-Fe2O3 and Cr2O3 particles. The mechanical alloying between α-Fe2O3 and Cr2O3 particles can be understood from the SEM picture of MA84 sample (Fig. 3b), where nanoparticles are distributed with a little bit agglomeration. A better uniformity in the particle size distribution is observed from the SEM picture (Fig. 3c) of thermal activated (SN17) sample. The elemental composition of the samples is determined from the point and shoots microanalysis at 10 selected points over a selected zone. The elements are found to be Cr, Fe and O. The atomic percentage of the elements is estimated from the EDX spectrum of the samples and Fig. 3d represents the spectrum for MA84 sample. The composition of Cr and Fe varies from 0.33 to 1.34 and 1.64 to 0.57, respectively, with respect to O composition 3 in MA0 sample. This suggests a random mixing of Fe2O3 and Cr2O3 oxides (i.e., an inhomogeneous distribution of elements) in MA0 sample. We obtained the atomic % of Fe (12.12), Cr (27.88) and O (60) (i.e., elemental composition Cr1.38Fe0.6O2.97) for MA84 sample; atomic % of Fe (11.96), Cr (27.06), O (60.17) (i.e., elemental composition Cr1.36Fe0.6O3.01) for SN17 sample. Therefore, elemental composition of the alloyed as well as annealed samples is close to the expected composition (Cr1.4Fe0.6O3). The distribution of Cr, Fe and O atoms in the alloyed samples are cross checked from elemental mapping over a selected zone, as shown in Fig. 4 for MA48 sample. The mapping (Fig. 4c–e) suggests a uniform distribution of Cr, Fe, O atoms over the zone. The particle size (~500 nm) of MA48 sample estimated from the SEM picture (Fig. 4b) is much larger than that obtained from XRD data (~46 nm using Williamson-Hall method). This means the SEM picture indicates the size of polycrystalline particles. Achieving some basic knowledge of the structural properties (i.e., size, shape, composition, and crystal structure) of the samples, attempt has made below to understand the magnetic properties of the samples.
Figure 3

SEM picture of selected samples MA0 (a), MA84 (b), SN17 (c) and EDX spectrum of samples MA0 (d), MA84 (e), SN17 (f).

Figure 4

(Colour online) Selected zone (a) is used to estimate the particle size (b) and mapping of Cr (c), Fe (d) and O (e) atoms in MA48 sample.

3.2. Magnetic properties

Fig. 5 shows the temperature dependence of ZFC and FC magnetization curve of different samples. The zero field cooled magnetization (MZFC), measured at 1 kOe field, does not change significantly (plateau like behaviour) in the temperature range 300 K to 600 K. The MZFC increases rapidly above 650 K to show a peak at about Tm ~ 810 K and then, MZFC rapidly decreases to indicate weak temperature dependence above 870 K. On the other hand, field cooled magnetization (MFC) at 1 kOe separates out from MZFC (i.e., onset of magnetic irreversibility) below the irreversibility temperature Tirr ~ 860 K. The separation between MFC and MZFC continued to increase with the decrease of temperature down to 300 K. The above features of MZFC and MFC are seen in the bulk α-Fe2O3 sample, as well as in the alloyed Cr1.4Fe0.6O3 samples, except some typical changes in the shape of the curves. The continuous increase of MFC with down curvature is observed in both bulk α-Fe2O3 and alloyed samples (e.g., MA48 and MA84). Interestingly, the shape of MFC curves of the annealed samples (e.g., SN3) has transformed into up curvature as an effect of thermal heating. The magnetization of bulk α-Fe2O3 sample decreases in the alloyed samples. The decrease is, further, accelerated by subsequent annealing of MA84 sample at 700°C. The decrease of magnetization is also reflected in the suppression of peak magnetization at Tm, when the sample changes from bulk α-Fe2O3 to SN17. We noted that the position of MZFC maximum at Tm is highly sensitive to the magnitude of measurement field. This is confirmed from the decrease of Tm ~ 810 K (at 1 kOe) to 750 K (at 2 kOe) in MA48 sample (Fig. 4b). The paramagnetic to canted ferromagnetic ordering temperature TN ~ 950 K of bulk α-Fe2O3 (hematite) sample [reported in Ref. [4]] is higher than Tm ~ 810 K at 1 kOe (in the present work). The field dependence of Tm (~940 K at 100 Oe and ~845 K at 200 Oe) has already reported in bulk α-Fe2O3 (hematite) sample [22, 23]. The present work showed that Tirr (~860 K) is higher than Tm (~810 K) and Tm decreases with the increase of measurement field. Based on these facts, we suggest that Tm does not represent a paramagnetic to canted ferromagnetic ordering temperature of the samples. Rather, Tm can be attributed as the temperature below which magnetic domains are blocked along the local anisotropy axes, where anisotropy energy might be higher than the measurement field [2426]. The interesting point is that magnetic phase of α-Fe2O3 is, still, dominating in the alloyed compound, of course with diminished magnitude with milling time. This can be attributed to the presence of larger atomic moment of Fe3+ ions (~5 μB) in comparison with Cr3+ ions (~3 μB). The effect of α-Fe2O3 phase is becoming weak in the annealed samples, as seen from the weak magnetic irreversibility and upward curvature in the MFC curve of SN3 sample. The results clearly showed that Cr2O3 dominated typical antiferromagnetic phase in the alloyed compound is yet to achieve at 700°C. The diffusion of Fe atoms into the Cr lattices either still continues, although the XRD data indicated the stabilization of alloyed Cr1.4Fe0.6 O3 close to the pattern of Cr2O3 sample. The appearance of a different kind of magnetic behaviour in the annealed samples may also be attributed to a modulated magnetic ordering at the interface of nano-sized grains.
Figure 5

Temperature dependence of zero field cooled (solid symbol) and field cooled (open symbol) magnetization at 1 kOe for the selected (Fe 2 O 3 bulk, MA48, MA84, SN3) samples. Tm represents the peak temperature of zero field cooled magnetization maximum. Magnetization data at 2 kOe are also shown to show the field dependence of Tm.

The existence of possible modulated magnetic phase in the samples is investigated from the field dependence of dc magnetization. The magnetization (M) with the variation of field (H) at room temperature is shown in Fig 6. The M(H) data of bulk Fe2O3 and Cr2O3 are shown in Fig. 6a to compare the magnetic change in alloyed compound. Bulk Cr2O3 sample does not show any noticeable hysteresis loop, whereas Fe2O3 sample has shown a well defined loop. The mixture of Fe2O3 and Cr2O3, i.e., MA0 sample (Fig. 6b), shows a clear loop, but the magnitude of loop of MA0 sample is less in comparison with Fe2O3. This is due to the effect of Cr2O3 coexisting with Fe2O3. As the alloying process continued by increasing the milling time, the loop area gradually decreases, but noticeable, as shown for MA48 sample in the inset of Fig. 6c. It is interesting to note that the nature of M(H) curves of annealed samples (Fig. 6e–f) drastically differ in comparison with the mechanical alloyed samples. The extrapolation of high field M (positive H side) data (shown by dotted line in Fig. 6f) intersects the M axis to negative value and the negative magnetization confirmed the field induced magnetic behaviour in the annealed samples. We would like to mention that neither bulk Fe2O3 nor bulk Cr2O3 sample exhibited such induced magnetization. Recently, similar field induced magnetization has been observed in the Cr2O3 nanoparticles [27], originating from the competitive spin ordering at the core-shell interface [28]. Viewing the absence of such field induced magnetic ordering in mechanical alloyed (with out heat treatment) samples, although grain size is in nanometer range, we believe that the field induced magnetic behaviour in our annealed samples is not solely due to particle size effect, but can be attributed to a modulated magnetic ordering due to diffusion of two different magnetic elements [29]. The induced magnetic behaviour in annealed samples is also reflected in the field dependence of differential magnetization (dM/dH) curves (Fig. 7). The differential susceptibility (dM/dH) increases to exhibit a maximum for mechanical alloyed samples when the applied field approaches to zero value, whereas dM/dH exhibits a minimum near to the zero field value for the annealed samples. The observation of induced magnetism confirms an experimental evidence of modulated spin structure in this compound, which was suggested from Neutron diffraction study [8]. To get the insight of modulated magnetic ordering, we analyze the M(H) data in terms of a general power series of H, i.e., the combination of liner and non-linear components in the equation
Figure 6

Variation of magnetization (M) with applied field (H) for different alloyed samples, along with bulk Fe 2 O 3 and Cr 2 O 3 samples. The existence of a small hysteresis loop for MA48 sample is shown in the inset of (c).

Figure 7

(Colour online) The field dependence of differential susceptibility (dM/dH) is shown for selected (MA0, MA48, MA84, SN1, SN3, SN17) samples.

M(H) = χ1H + χ2H2 + χ3H3 + (rest of the terms)

In the above equation, we have assumed that magnetic susceptibility is not a scalar quantity, rather a tensor with directional property in the material. The coefficients χ1 and χ3 are assumed to be symmetric with field (H), where as χ2 is antisymmetry to preserve the vector nature of magnetization (M). The validity of equation (1) is tested from Fig. 8. The M vs. H2 data (in Fig. 8a) indicate that magnetization of all the samples are symmetric about the H2 axis, irrespective of alloyed or annealed samples, and the change of magnetic direction is taken care by χ2, i.e., χ2(-H) = -χ2(H). It is really interesting to note that M vs. H3 plot (Fig. 8b) represents a ferromagnetic like appearance, irrespective of the samples and confirms the symmetric nature of χ3, where the change of magnetization direction is taken care by H. The investigation of non-linear components (specially the second-harmonic generation) in antiferromagnetic material may be one of the important technological aspects of the present material.
Figure 8

(Colour online) (a) M vs. H 2 plot and (b) M vs. H 3 plot for selected samples.

Now, look at the magnetic parameters. The remanent magnetization (MR) and coercive field (HC) of the samples were calculated using M(H) data. The variation of MR and HC of the alloyed compound with milling time and annealing time is shown in Fig. 9. We noted that both HC (~1620 Oe) and MR (~172 memu/g) of α-Fe2O3 sample are reduced by nearly 2.5 times in MA0 sample (HC ~ 700 Oe and MR ~ 66.3 memu/g), although Cr2O3 is not showing any hysteresis loop. This feature suggests that properties of magnetic particles, irrespective of bulk (present work) or nanosize (reported work [30]), are strongly influenced by the surrounding matrix. The HC of MA0 sample is continued to decrease with milling time (e.g., HC ~ 120 Oe for MA24). On the other hand, HC has shown a significant increase by annealing MA84 sample at 700°C up to 3 hours, followed by a slight decrease at 17 hours annealed sample (SN17). The MR of MA0 sample also decreases to attain an almost constant value ~3.4 memu/g for MA24 and MA48 samples, followed by a slow increase (~4.4 memu/g) for MA84 sample and continued for (~5.3 memu/g) SN1 sample. The MR value, again, starts to decrease at higher annealing time, e.g. ~3.9 memu/g for SN17 sample. The decrease of both remanent magnetization and coercivity at higher annealing time indicates that the alloyed compound is approaching to the magnetic state dominated by non-hysteretic Cr2O3 sample or a different kind of magnetic ordering that is already indicated earlier. The fact is that mechanical alloying has introduced a certain amount of lattice strain (and defects) in the material. However, we did not find the increase of coercivity with increasing milling time (and associated increasing strain), rather the coercivity decreases. Hence, the effect of strain and associated defects is minor on the magnetic properties of the present alloyed compound and such effect can be included in the combined contribution of grain boundary spins.
Figure 9

The variation of coercivity (H C ) and remanent magnetization (M R ), calculated from hysteresis loop, are shown with the variation of milling time and annealing time.

4. Conclusion

The magnetic properties of mechanical alloyed Cr1.4Fe0.6O3 compound strongly depends on the structural change associated with the variation of milling time and annealing temperature. The variation of cell parameters is attributed to the effect of diffusion of Fe3+ ions into the boundary of Cr3+ ions. The structural analysis, using XRD and SEM with EDX spectrum, may suggest the completion of alloy formation during milling process, but the magnetic behaviour confirmed that the alloying process is, still, continued at 700°C to approach the structure dominated by Cr2O3 phase. Experimental results suggested that magnetic ordering of mechanical alloyed Cr1.4Fe0.6O3 compound, annealed at 700°C, belongs to neither a canted ferromagnetic ordering of α-Fe2O3 nor a typical antiferromagnetic ordering of Cr2O3. Rather, a field induced magnetic order appeared in the annealed samples, which we attribute to the appearance of a modulated spin structure in the compound. This experimental work may instigate subsequent investigations on the non-linear magneto-opto-electronic properties of similar materials.

Declarations

Acknowledgements

We thank to CIF, Pondicherry University for providing experimental facilities. We also thank to FIST Program in the department of Physics for providing XRD measurement facilities. RNB also thanks UGC for providing financial support (F.No. 33-5/2007(SR)).

Authors’ Affiliations

(1)
Department of Physics, Pondicherry University

References

  1. Cox PA: Transition Metal Oxides: An Introduction to Their Electronic Structure and Properties (Clarendon, Oxford, 1992).
  2. Jaffe JE, Dupuis M, Gutowski M: Phys Rev B. 2004, 69: 205106-10.1103/PhysRevB.69.205106.View ArticleADSGoogle Scholar
  3. Levistein HJ, Robbins M, Capio C: Mater Res Bull. 1972, 7: 27-10.1016/0025-5408(72)90066-9.View ArticleGoogle Scholar
  4. Yin-Yuan Li: Phys Rev. 1956, 101: 1450-10.1103/PhysRev.101.1450.View ArticleADSGoogle Scholar
  5. McCarty KF, Boehme DR: Journal of Solid State Chemistry. 1989, 79: 19-10.1016/0022-4596(89)90245-4.View ArticleADSGoogle Scholar
  6. Music S, Lenglet M, Popovic S, Hannoyer B, Czako-Nagy I, Ristc M, et al: J Mater Sci. 1996, 31: 4067-10.1007/BF00352669.View ArticleADSGoogle Scholar
  7. Osmond WP: Proc Phys Soc. 1962, 79: 394-10.1088/0370-1328/79/2/319.View ArticleADSGoogle Scholar
  8. Cox DE, Takei WJ, Shirane G: J Phys Chem Solids. 1963, 24: 405-10.1016/0022-3697(63)90199-9.View ArticleADSGoogle Scholar
  9. Dzyaloshinskii IE: Zh Eksp Teor Fiz. 1959, 37: 881-[Sov. Phys. JETP 10, 628 (1960)]Google Scholar
  10. Astrov DN: Zh Eksp Teor Fiz. 1960, 38: 984-[Sov. Phys. JETP 11, 708 (1960)]Google Scholar
  11. New D, Lee NK, Tan HS, Castner TG: Phys Rev Lett. 1982, 48: 1208-10.1103/PhysRevLett.48.1208.View ArticleADSGoogle Scholar
  12. Schmid H: Ferroelectrics. 1994, 161: 1-View ArticleGoogle Scholar
  13. Hill NA: J Phys Chem B. 2000, 104: 6694-10.1021/jp000114x.View ArticleGoogle Scholar
  14. Di Matteo S, Natoli CR: Phys Rev B. 2002, 66: 212413-10.1103/PhysRevB.66.212413.View ArticleADSGoogle Scholar
  15. Merchant P, Collins R, Kershaw R, Dwight K, Wold AJ: Solid State Chemistry. 1979, 27: 307-10.1016/0022-4596(79)90173-7.View ArticleADSGoogle Scholar
  16. Bhattarcharya AK, Hartridge A, Mallick KK, Majumdar CK, Das D, Chintalapudi SN: Journal of Material Science. 1997, 32: 557-10.1023/A:1018506912078.View ArticleADSGoogle Scholar
  17. Manjanna J, Venkateswaran G: The Canadian Journal of Chem Eng. 2002, 80: 882-View ArticleGoogle Scholar
  18. Chicinas IJ: Optoelectronics and Advanced Materials. 2006, 8: 439-Google Scholar
  19. Williamson GK, Hall WH: Acta Metall. 1953, 1: 22-10.1016/0001-6160(53)90006-6.View ArticleGoogle Scholar
  20. Hamzaoui R, Guessasma S, ElKedim OJ: Alloys and Comp. 2008, 462: 29-10.1016/j.jallcom.2007.07.111.View ArticleGoogle Scholar
  21. Chambers SA, Droubay T: Phys Rev B. 2001, 64: 075410-10.1103/PhysRevB.64.075410.View ArticleADSGoogle Scholar
  22. Zysler RD, Vasquez-Mansilla M, Arciprete C, Dimitrijewits M, Rodriguez-Sierra D, Saragovi C: J Magn Magn Mater. 2001, 224: 39-10.1016/S0304-8853(00)01365-2.View ArticleADSGoogle Scholar
  23. Vasquez Mansilla M, Zyslera R, Fiorani D, Suber L: Physica B. 2002, 320: 206-10.1016/S0921-4526(02)00683-X.View ArticleADSGoogle Scholar
  24. Bhowmik RN, Ranganathan R, Nagarajan R: J Magn Magn Mater. 2006, 299: 327-10.1016/j.jmmm.2005.04.021.View ArticleADSGoogle Scholar
  25. Aharoni A, Frei EH, Schieber M: Phys Rev. 1962, 127: 439-10.1103/PhysRev.127.439.View ArticleADSGoogle Scholar
  26. Lin ST: Phys Rev. 1959, 116: 1447-10.1103/PhysRev.116.1447.View ArticleADSGoogle Scholar
  27. Tobia D, Winkler E, Zysler RD, Granada M, Troiani HE: Phys Rev B. 2008, 78: 104412-10.1103/PhysRevB.78.104412.View ArticleADSGoogle Scholar
  28. Bhowmik RN, Nagarajan R, Ranganathan R: Phys Rev B. 2004, 69: 054430-10.1103/PhysRevB.69.054430.View ArticleADSGoogle Scholar
  29. Abe R, Tsunoda Y, Nishi M, Kakurai K: J Phys Condens Mater. 1998, 10: L79-10.1088/0953-8984/10/4/001.View ArticleADSGoogle Scholar
  30. Frandsen C, Morup S: J Magn Magn Mater. 2003, 266: 36-10.1016/S0304-8853(03)00453-0.View ArticleADSGoogle Scholar

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