Home
Search
Collections
Journals
About
Contact us
My IOPscience
Morphology dependent magnetic properties of -Fe2O3 nanostructures
This content has been downloaded from IOPscience. Please scroll down to see the full text. 2014 Mater. Res. Express 1 046104 (http://iopscience.iop.org/2053-1591/1/4/046104) View the table of contents for this issue, or go to the journal homepage for more
Download details: IP Address: 116.203.205.246 This content was downloaded on 11/10/2014 at 17:12
Please note that terms and conditions apply.
Morphology dependent magnetic properties of αFe2O3 nanostructures S Chakrabarty1, T K Jana1, K De2, S Das3, K Dey4 and K Chatterjee1 1
Dept. of Physics and Technophysics, Vidyasagar University, Midnapore—721102, India NITMAS, Jhinga, D. H. Road, 24 Pgs (S)-743368, India 3 Dept. of Electronics and Communication Engineering, Guru Ghasidas Vishwavidyalaya, Bilaspur, India 4 Department of Solid State Physics, Indian Association for the Cultivation of Science, Jadavpur, Kolkata 700 032, India E-mail:
[email protected] 2
Received 17 June 2014, revised 14 August 2014 Accepted for publication 11 September 2014 Published 10 October 2014 Materials Research Express 1 (2014) 046104 doi:10.1088/2053-1591/1/4/046104
Abstract
Well crystalline α-Fe2O3 nanomaterials with a wide range of morphology variation have been successfully synthesized by solvothermal route. The synthesized products have been characterized for structural and morphological details by employing x-ray diffraction patterns, transmission electron microscopy, field emission scanning electron microscopy and energy dispersive x-ray spectroscopy. Various unique shapes of α-Fe2O3 nanocrystal have been modelled on the basis of their growth evolution. The effect of morphology of α-Fe2O3 nanocrystals on their magnetic behaviour has been studied by investigating temperature and field dependence of magnetization. The results are analyzed considering all the possible surface anisotropy and lattice strain evolved due to their surface structure. This comprehensive study of morphology dependent magnetic behaviour of α-Fe2O3 nanomaterials offers a better opportunity to tune the materials in the desired technological applications. Keywords: iron oxide, nanostructures, magnetic properties
1. Introduction
Research on synthesis of nanostructured materials has gained increasing attention due to their morphology dependent functional properties such as optical, electrical, magnetic, catalytic, mechanical and chemical [1–5]. Therefore, scientists are actively engaged in the issue of morphology controllable synthesis of low dimensional structures. In the past decade iron (III) oxide, especially hematite (α-Fe2O3), has been at the focus of research interest due to its huge Materials Research Express 1 (2014) 046104 2053-1591/14/046104+17$33.00
© 2014 IOP Publishing Ltd
S Chakrabarty et al
Mater. Res. Express 1 (2014) 046104
potential for tuning the properties depending upon its morphology [4–10]. α-Fe2O3 is thermodynamically the most stable iron oxide with a band gap of 2.2 eV [11] and the material is being investigated extensively for its various technological aspects such as photoinduced water splitting [12], catalysis [13], gas sensing [14], magnetic recording [5], drug delivery [15], tissue repair engineering [16], lithium-ion batteries [5], spin electronic devices [17] and pigments [18]. Understanding the correlation, between the morphology and the magnetic properties, is the prerequisite for the successful and efficient applications of α-Fe2O3 nanomagnetism in future technology. Appreciable attention has been paid to synthesizing different morphology of α-Fe2O3 such as plate-like, stalactite-like, coral-like, hexagonal, nanotube, nanorods, urchin-like, ring-like, nanospheres, flower-like, nanorhombohedral, nanospindle [8, 19–29]. Therefore, it is evident that there is a strong thrust on the synthesis of morphologically varied α-Fe2O3 nanostructure. Complex 3D architectures, expecting interesting magnetic properties, have also been attempted with great enthusiasm such as dendrite and snowflake-like [30], airplane-like [31], cantaloupelike [32], shuttle-like [33], nanocages [34], dendritic micro-pines [35], branched topology [36] and so on. However, the studies are in general, bearing limited morphology variations in each individual approach and that, in most of the cases, employs complex synthesis mechanism. Integrated version of wide morphology variation is somehow less attempted. Say for example, from nano network systems to descrete nanostructures of α-Fe2O3, in a common pathway, are of enormous interest owing to their high complexity and high anisotropic surface texture. In our previous article [37], we have reported such variation in α-Fe2O3 nanostructures with simple, surfactant free, hydrothermal approach and presented the growth evolution mechanism in details. Here it opens up a huge possibility of study addressing morphology dependent magnetic properties of α-Fe2O3 nanostructures in a much wider length scale. Actually, the main dragging force behind the morphology variation of α-Fe2O3 is having to tune the inherent nanomagnetism for the desired technological applications. Magnetic behaviour of α-Fe2O3 is very much sensitive to its morphology and a decent amount of articles are available reporting morphology dependent magnetic properties of hematite nanostructures [2, 4, 10, 19, 26– 28, 30, 35, 36]. Recently, Bharathi and co-workers [30] have reported controlled growth of dendrites, single-and double-layered snowflakes of α-Fe2O3 to show shape dependent magnetic properties with variation in coercivity values. Dendritic micropines α-Fe2O3 structure [35], owing to their shape anisotropy and lattice strain, results in lowering of Morin transition temperature TM at 216 K and coercive force as high as 1510 Oe at 300 K. Mitra et al [10] prepared nanospindle, nanorhombohedron and nanocube structured α-Fe2O3 and reported strong dependence of magnetic behaviour on their morphological aspects. Jagadeesan et al [12] introduced complex shape anisotropy in the form of nanocups in α-Fe2O3 structure to get drastic change in magnetic results compared to their hollow spherical counterpart. Porous α-Fe2O3 nanostructures with branched topology [36] display two TM, one at 195 K and the other at 243 K. Bo Tang and his team [38] have shown α-Fe2O3 nanorods exhibiting weakly ferromagmetic behaviour at low temperature and superparamagnetic property at room temperature. Mandal et al also observed lowering of TM in the structured α-Fe2O3 samples depending upon their size, shape and lattice parameters [39]. Hexapods of α-Fe2O3 with arm diameter of 60–80 nm and length of 400–900 nm exhibit TM = 233 and 245 K under field cooled (FC) and zero field cooled (ZFC) conditions, respectively [40]. Absence of TM is also reported in the literature as Mika Sillanpaa et al have synthesized chainlike hematite [41] having no
2
S Chakrabarty et al
Mater. Res. Express 1 (2014) 046104
apparent Morin transition and Zhao et al reported absence of Morin transition in hematite nanorods [42]. Nevertheless, as the synthesis of hematite structure varied in wide range, to address its morphology-controlled magnetic behaviour remains a challange. Although several attempts are reported addressing unique structure of α-Fe2O3 and structure dependent magnetic study, an integrated version of extensive morphology variation to understand the tuning of magnetic properties is still lacking. In particular, network type morphology in α-Fe2O3, to the best of our knowledge, has not yet been reported. Herein we report the variation of α-Fe2O3 in morphological aspect, and the wide range of magnetic response originating from different samples is presented. Network type systems, in general, show high ferromagnetic response with suppressed TM, whereas discrete morphological systems reveal strong Morin transition with the presence of blocking temperature. The samples produced by modified synthesis route show the signature of either the Morin transition or the blocking temperature. The details of magnetic behaviour of each individual sample is reported and analyzed on the basis of their surface related properties such as surface anisotropy, lattice strain etc. 2. Experimental section
All the reactants are Merck made of analytical grade and used without further purification. In our synthesis procedure we have adopted simple solvothermal route followed by proper annealing treatment and the details of the synthesis technique are described in our previous article [37]. To tune the morphology FeCl3 precursor was taken in different organic solvents such as ethanol amine [EA], ethylene diamine [ED], ethylene glycol [EG], acetic acid [AA], ethanol [EtOH], acetaldehyde [AH] and in inorganic solvent water [H2O] also. Here it is worth mentioning that the solvents are the key factor for modulating the morphology. To expand the dimension of morphology variation we have also altered the reaction time in furnace for the solvothermal reaction. As was reported earlier [37] the initial reaction time was 18 h and here another set of samples have been prepared with 12 h reaction time keeping the other things unaltered. Henceforth, we will refer to the samples, treated with 18 h reaction time, by only their respective solvent code name such as EA, EG etc, and the samples perpared in 12 h reaction time as EG12, EA12 etc. Structural analysis of all the powdered samples was carried out by Rigaku Mini-Flex x-ray diffractometer using Cu Kα radiation (λ = 1.541 78 Å) source. Morphological analysis was done by both JEM 2100 transmission electron microscope (TEM) at an accelerating voltage of 200 keV and NEON 40 (CARL ZEISS) scanning electron microscope (SEM). Energy dispersive x-ray (EDX) spectrocopy was carried out in an S-4200, Hitachi. DC Magnetization was measured in a vibrating sample magnetometer (VSM) with a field range from 0 to 10 T. In the case of zero-field-cooled (ZFC) mode the sample was cooled down to the desired temperature at zero magnetic fields, while for the field-cooled (FC) mode the sample was cooled in a static magnetic field. 3. Results and discussions
Figure 1 shows powder diffraction pattern for nine different samples. It is obvious from the figure that the yield materials are well crystalline and well matched with hexagonal α-Fe2O3 3
S Chakrabarty et al
Mater. Res. Express 1 (2014) 046104
Figure 1. X-ray diffraction pattern of as prepared iron oxide samples synthesized in different reaction environment. The respective sample name in the figure refers the respective reaction condition. Inset shows the magnified version of the two major peaks originated from (104) and (110) planes.
phase, matching with ICDD card number 33-0664. The unit cell is of a = b = 5.0356 Å and c = 13.7489 Å. As there is no unidentified peak in the XRD data it confirms the chemical purity of the product specimens. In the inset XRD peaks for the plane (104) and (110) reveal that there is slight peak shifting in the main two crystalline planes compared to the bulk counterpart. Those peak positions as well as the peak broadenings are different for different samples. It implies that the produced α-Fe2O3 samples are bearing lattice strain with different shape and size. It is already established that lattice symmetry for α-Fe2O3 in nanodimension are not rigidly valid [43] and the lattice strain in α-Fe2O3 affects the XRD peak positions appreciably [44]. As the lattice strain is associated with the nanodimensional morphology and our samples are expected to vary in morphology, so the result is quite expected. The result also confirms that no other crystalline phase of iron oxide, except α-Fe2O3, has evolved in the synthesis procedure with different solvents. Figure 2 shows the photogallery of FESEM images of different as-prepared samples. It can be seen from figures 2(a)–(c) that in the case of EA, ED and EG the grown Fe2O3 appears in interconnected network type formation. EA sample looks like branched network while ED resembles nano-ginger morphology. EG develops nano-coral like structures bearing high aspect ratio in its length-to-width dimension. From the SEM image of sample AA, figure 2(d), the structure is bead-like, i.e., one-dimensional structure having fluctuant diameter along the length and the shape of this α-Fe2O3, is nano-bead type. Here each segment was found to be seperated. The sample treated with EtOH, shown in figure 2(e), has hexagonal thin disc-like arrangement. The structure is almost uniform, discrete and homogeneous. Sample AH, figure 2(f), gives almost spherical nanoparticles that are also uniform and homogeneous. The sample treated with only water, instead of mixing water with organic solvent, produces flat nanorods-like structure of α-Fe2O3 and the morphology is shown in figure 2(g). Sample EA12 presents uniform spindle-like crystalline Fe2O3, whereas EG12 produces highly asymmetric multipodal-like 4
S Chakrabarty et al
Mater. Res. Express 1 (2014) 046104
Figure 2. FESEM images of the prepared α-Fe2O3 nanocrystals prepared in different
reaction medium and reaction time. The sample name and the scale are indicated in each figure. Inset in each figure shows the optical microscopic image of corresponding sample.
particles, and these are shown in figures 2(h) and (i), respectively. Both EA12 and EG12 samples are distinctly different in morphology from their 18 h reacted sample counterpart. All, these nine samples are of unique structure and are bearing different morphology. The insets of figures 2(a)–(i) show the corresponding appearance of the sample in optical microscope and it is interesting to note the difference in color of different samples. All the dimensional parameters of different samples are presented in table 1 in this article. Figure 3 shows the photogallery of TEM images of the prepared iron oxide samples. In agreement with the SEM images TEM images of EA, ED and EG samples, figures 3(a)–(c), respectively, show nonlinear interconnected structure of Fe2O3 bearing a nano-network-like formation. These nano-network colonies of Fe2O3 samples are inevitably highly anisotropic in surface nature. Most importantly, this network type morpholgy is expected to provide strong interparticle interaction, which should be quite different from any discrete morphology to influence various physical properties. To study the magnetic property of these structures would be of immense interest. In the TEM image of AA sample a bead-like structure is quite prominent and some seperated/dissociated segments, in almost spherical nature, are also present. Well faceted, flat, almost hexagonal disc-like structure of α-Fe2O3 is shown in 5
C (Å)
Volume of unit cell (Å3)
Δc c−1 (%)
Average particle dimension
5.034 62
13.743 44
301.689
0.039 73
ED
5.019 14
13.678 70
298.425
0.513 21
EA
5.018 74
13.689 13
298.604
0.436 62
L: 500–700 nm, W: 56 nm (network) L: 275 nm, W: 75nm (network) L: 130 nm, W: 35 nm (network)
EtOH
5.014 88
13.665 07
297.621
0.613 46
AA
5.026 13
13.706 93
299.874
0.3062
AH
5.017 45
13.655 06
297.708
0.687 22
H2O
5.016 72
13.666 95
297.881
0.599 62
EA12
5.026 97
13.706 99
299.975
0.305 76
EG12
5.020 49
13.702 77
299.110
0.336 65
Name of the sample
A or b (Å)
EG
6
Flat side ∼2–3 μm, H: 250 nm L: 300 nm, W: 30 nm 140 nm L: 1 μm, W:200 nm H: 20 nm L: 1 μm W: 350 nm 63 nm × 40 nm
MR (emu g−1) (300 K, 5 K)
MS (emu g−1) (300 K, 5 K) 54.78, 32.4
Shape
TM (K)
TB(K)
HC (T) (300 K, 5 K)
Nano-coral
234 (feeble)
—
0.0248, 0.0640
4.41, 7.38
Nano-ginger
—
0.0190, 0.0725
0.374, 3.6
5.51, 20.22
Branched network
229 (feeble) 225 (feeble)
—
0.031, 0.080
0.788, 1.982
7.34, 11.32
Hexagonal disc
246 (sharp)
50 (broad)
0.075, 0.075
0.484, 0.062
6.747, 2.30
Nano-bead
235 (sharp) 211 (sharp)
62 (sharp)
0.290, 0.067
0.37, 0.013
5.524, 1.65
62 (supressed)
0.0282, 0.0892
0.138, 0.025
2.55, 2.385
0.278, —
0.0495, —
Nano-sphere Flat Nano-road
236 (sharp)
—
Nano spindle
—
140
—, 0.048
—, 1.165
Asymmetric particle
228 (sharp)
—
0.0054, 0.0874
0.138, 0.025
Mater. Res. Express 1 (2014) 046104
Table 1. Structural, morphological and magnetic variation of as prepared α-Fe2O3 samples.
—
0.37, 8.856 0.66, — S Chakrabarty et al
S Chakrabarty et al
Mater. Res. Express 1 (2014) 046104
Figure 3. TEM images obtained for different Fe2O3 samples (a)–(i) and the sample
names are indexed accordingly.
figure 3(e) for the EtOH sample. Dispersed, spherical α-Fe2O3 nanoparticles in the AH sample are presented in figure 3(f). In the case of H2O, the product is porous flat rod-like and the porosity is evident from figure 3(g). One single spindle-like Fe2O3 is shown in figure 3(h) for the EA12 sample and asymmetric particle nature for EG12 sample is evident from figure 3(i). To reveal the chemical composition and crystal structure, EDX and high resolution TEM (HRTEM) and selected area electron diffraction (SAED) have been carried out, and typical results of a few samples are shown in figure 4. EDX taken from the EtOH sample in figure 4(a) has shown the presence of Fe and O. Here the Si peak is generated from the Si substrate onto which the powder sample was loaded. A high-resolution TEM image shown in figure 4(b) shows the typical crystal nature of Fe2O3 sample prepared in EA12. Two other typical HRTEM images (c) and (d) show the atomic spacing of Fe2O3 crystals originated from ED and AH samples, respectively. The SAED pattern obtained from H2O sample is shown in figure 4(e). The corresponding crystal planes of α-Fe2O3 in HRTEM and SAED have been indexed and they are in good agreement with the planes revealed from the XRD study of the samples. So far from our study, we have seen that several nanostructures of α-Fe2O3 have been grown using different solvents and different times in reaction. To demonstrate the evolution of different shape we propose a model. It is interesting to categorically divide all the samples into
7
S Chakrabarty et al
Mater. Res. Express 1 (2014) 046104
Figure 4. EDX (a), HRTEM (b)–(d) and SAED (e) taken from EtOH, EA12, ED, AH and water sample, respectively. The elements (a) and crystal planes (b)–(e) are indexed within the respective figure.
some groups depending on their morphology and synthesis technique. Before we start the systemetic magnetic measurement, if we can group the samples methodically it will be beneficial for the analysis of further results, and the commonality will bring the opportunity to make the discussion simple. In figure 5, we have demonstrated the schematic diagram of morphology evolution of different α-Fe2O3 nanostructures starting from various nucleating points and guided differently by different solvents used in the reaction. It can be assumed that spherical units of Fe(OH)3/Fe(O)(OH) nuclei, resulting from the reaction of FeCl3 and NaOH, come into the picture first to initiate the formation of Fe2O3. It is shown in the square box of the centre of the model as first phase of evolution. Those nuclei are now guided and oriented by reaction medium and are markedly influenced by the microenvironment produced within the respective solvent used in the reaction. Three big arrows in upper, side and lower directions are indicating three groups of the product materials, namely ‘network system’, ‘discrete system’ and ‘synthesis modified system’. Network system represents the morphology of the α-Fe2O3 nanostructures as a whole in a network form. Similarly, discrete system represents α-Fe2O3 nanostructures in discrete form and synthesis modified system shows the variation of α-Fe2O3 when the synthesis route has been altered by reducing the reaction time and by excluding any organic solvent. In each division ‘a’, ‘b’ and ‘c’ represents the oriented attachment of Fe(OH)3 nuclei guided by the reaction environment, formation of final Fe2O3 after annealing at 600 °C and the corresponding SEM image of the sample, respectively. It is interesting to note the transformation of morphology in EG12 to EG sample and in EA12 to EA sample. Here EG12 and EA12 signify samples prepared in a 12 h reaction, whereas EG and EA signify samples produced in an 18 h reaction keeping other parameters unaltered. It is quite evident that in the first case asymmetric, multipodal-like nanoparticles, produced in EG medium in a 12 h reaction, get associated with each other with the passage of time and consequently develop the nanocoral-like structure in the EG sample. On the other hand, in the case of EA medium 12 h reaction produces spindle-like structure α-Fe2O3 but when it is allowed to react further the particles firstly get dissociated into several segments and then the segments are attached to form branched network-like system in the EA sample.
8
S Chakrabarty et al
Mater. Res. Express 1 (2014) 046104
Figure 5. Schematic representation of the evolution of morphology for different αFe2O3 nanocrystals grouped in three directions. Central box, ‘a’ and ‘b’ represent the three phases of reaction respectively and ‘c’ presents real time HRSEM images of the corresponding samples. ΔT signifies the heat treatment of each sample by 600 °C.
For the detailed description of magnetic behaviour of each sample now, we will proceed groupwise. The magnetic properties of the studied samples were characterized from temperature and field dependence of magnetization curves. Both the zero field cooled magnetization (MZFC) and field cooled magnetization (MFC) curves at an applied field (H) of 200 Oe were measured in heating cycle for all the samples. The field dependence of magnetization (M–H) has been taken at temperatures 300 K and 5 K only. The magnetic measurement of the network system is shown in figure 6. Figures 6(a)–(c) shows the curves for the temperature dependence of MZFC and MFC under an applied field of 200 Oe for EG, ED and EA systems, respectively. For all these samples, we observe the suppression of spin flip Morin transition and dominating ferromagnetic behaviour irrespective of their variation in lattice strain (table 1). This observation is further supported by their microstructure analysis giving a clear demonstration of the connected network system for all of them. Great surface anisotropy can lead to tuning of the alignment of spins of sublattices in favour of ferromagnetic ordering, and can be a probable explanation of strong ferromagnetism in this series of samples EA, ED and EG, respectively. A large field-cooled effect is observed in the studied temperature range of 5–300 K for all these three systems with a weak signature of Morin transition observed at 225 K, 229 K and 234 K for EA, ED and EG, respectively. The suppression of Morin transition from that of the bulk hematite is more prominent and seems to be due to variation of particle size distribution with varying diameters for these connected network systems. For all three cases, figures 6(a)–(c), MZFC and MFC curves split significantly; 9
S Chakrabarty et al
Mater. Res. Express 1 (2014) 046104
Figure 6. (a)–(c) Temperature (T) dependence of MZFC and MFC curves measured at
H = 200 Oe for EG, ED and EA, respectively. The inset in (a) shows T-dependence of MFC curve for a selective temperature range. The inset in (b) exhibits T-dependence of MFC–MZFC. The inset (c) shows T-dependence of corresponding differential of MZFC curve. (d)–(e) Field-dependent M–H curves at 300 and 5 K for EG, ED and EA, respectively. The insets in (d)–(e) selectively reveal corresponding low field regions of the M–H curves measured at 300 and 5 K.
the MFC rises regularly with decreasing temperature, while the MZFC curve decreases slowly indicating strong ferromagnetic behaviour and magnetization sensitive to the size evolution from submicron particles to the 500–700 nm nano-coral. The divergence between the magnetization in ZFC and FC process, as obtained in figures 6(a) to (c), does not exist in bulk hematite and is characteristic of small particles [45–47]. The MZFC and MFC curves clearly indicate that the particles are quite thermally stable without blocking below TM. The consistent increase of MFC with decreasing temperature indicates the presence of spontaneous magnetization and long-range magnetic ordering in these systems [4, 48, 49]. The dominating ferromagnetism of all these three samples suppresses the spin flip Morin transition temperature. Figures 6(d)–(f) presents the field dependent magnetization at 300 K and 5 K for EG, ED and EA, respectively. The hysteresis loops measured at 5 K and 300 K are indicative of soft ferromagnet. The values of MR, HC and MS at 300 K and 5 K are listed in table 1. Larger values of HC and MR indicate strong ferromagnetic behaviour at 5 K, which gets relatively weak at 300 K. The results of the M(T) and M(H) measurements ensure the presence of long-range magnetic ordering in all these systems that suppresses the Morin transition. Figure 7 shows the magnetic response of as prepared discrete system consisting of EtOH, AA and AH samples. The temperature (T) dependence of magnetization measured at 200 Oe is shown in figures 7(a) to (c) for the samples EtOH, AA and AH, respectively. As obtained, for all these samples, we get a sharp decrease in magnetization in both ZFC and FC curves with decrease in temperature, which can be assigned to Morin temperature. The Morin temperature
10
S Chakrabarty et al
Mater. Res. Express 1 (2014) 046104
Figure 7. (a)–(c) Temperature (T) dependence of MZFC and MFC curves measured at
H = 200 Oe for EtOH, AA and AH, respectively. The insets in (a) and (b) show Tdependence of corresponding differential of MZFC curve. The inset in (c) shows Tdependence of the MZFC and MFC curve for a selective temperature range. (d)–(e) M–H curves at 300 and 5 K for EtOH, AA and AH, respectively. The insets in (d)–(e) selectively reveal corresponding low field regions of the M–H curves measured at 300 and 5 K.
(TM) derived from the sharp peak of the corresponding differential ZFC curves, gives its value of 246 K for EtOH, 235 K for AA and 211 K for AH, respectively. In order to explain the variation of magnetic behaviour in the entire temperature range from 300 K to 5 K for the system of EtOH, AA and AH, we use table 1 which shows the interfering effect of lattice strain along with the shape and size of the discrete particles of these systems. These lower values of Morin temperature (TM) of these samples compared to that of the bulk hematite (TM = 263 K) can be explained as a consequence of their submicron form giving high lattice strain in the system (table 1). The samples EtOH, AA and AH are all in the form of individual particles in the shape of hexagonal discs (EtOH), nano-beads (AA) and nano-spheres (AH) having therefore large surface area. The sizes of the particles range from 150 nm to a maximum of 300 nm (table 1). This enhanced surface effect is also visible in executing the blocking temperature (TB) by all these three samples in their differential ZFC curves, which corresponds to the blocking of the surface spins of the average-small size particles [50, 51]. The values of TB obtained are 50 K, 62 K and 62 K for EtOH, AA and AH, respectively. Wang et al also reported quite similar TB for mesoporous α-Fe2O3 nanostructures towards their application in gas sensor and lithium ion battaries [52]. Thermomagnetic irreversibility is depicted for EtOH in the lower inset of figure 7(a) from the distinct difference between MZFC (T) and MFC (T). The split of the MZFC (T) and MFC (T) curve starts below TM and significantly increases below 100 K reflecting the existence of a large size distribution of magnetic units, whose moments block progressively with decreasing temperature [53]. The progressive increase of MFC below 220 K reflects the existence of weak ferromagnetism, while MZFC (T) curve decreases rapidly below 100 K for 11
S Chakrabarty et al
Mater. Res. Express 1 (2014) 046104
EtOH. For the sample AA with bead like morphology and having moderate strain (0.3%), a consistent decrease of MZFC and MFC curves down to 100 K in figure 1(b), indicates predominant antiferromagnetic (AFM) ordering in the sublattices of the studied compound. Below the temperature TB the MZFC and MFC curves decrease down to 29 K and then decrease down to 5 K. The decrease of MFC curve in 62 K ⩽ T ⩽ 29 K is attributed to the AFM like spin interaction between Fe ions. The increase of MFC curve below 29 K implies presence ferromagnetic (FM) interaction at T ⩽ 29 K. For the sample, AH with sphere-like morphology, we get similar behaviour to that of AA down to blocking temperature TB, below which a slow increase of MZFC and MFC is observed, which is ascribed due to the presence of weak FM interaction below TB in this sample. The field (H) dependence of magnetization measured at temperatures 300 K and 5 K are shown in figures 7(d) to (f) for the samples EtOH, AA and AH, respectively. The existence of a hysteresis loop at 300 K, as well as the absence of magnetization saturation at 5 K for all the samples should be noted. The hysteresis behaviours of M–H curves evidence a weak ferromagnetic state of the studied samples at room temperature. The M–H curve at higher field shows almost linear dependence of magnetization on the applied magnetic field. The remanent magnetization (MR), coercivity (HC) and saturation magnetization (MS) values (determined by extrapolating 1/H to zero-field in the M versus 1/H plot based on the high field data) of all the above samples are listed in table 1. For the sample EtOH, the M−H curves show a steep linear increase with the field at the low-field region (H ⩽ 0.35 T) then a downward curvature followed by almost a linear behaviour up to the highest applied field without saturation. Hence there are two main contributions to magnetization in this case as given by M = MNC(H) + χH. The initial rapid increase in magnetization is attributed to the weak ferromagnetic behaviour, MNC in the samples, which tries to saturate the moment. The non-saturation of moments at high field may probably be coming due to residual antiferromagnetism in the samples, and expressed by the term χH, where χ is the magnetic susceptibility and H, the magnetic field. Hence, from M (H) isotherms it is clear that the contribution of ferromagnetic phase at room temperature is stronger than that at 5 K. For sample AA, the magnetization curve at low field region (inset of figure 7(e)) reflects further evidence of a delicate ferromagnetic behaviour at 5 K. The shape of the loop and the large values of HC, MR and MS indicate strong ferromagnetic behaviour for sample AA at 300 K, which becomes relatively weak at 5 K due to progressive influence of antiferromagnetic ordering in the sublattices. It is understood that enhanced surface anisotropy of the studied sample induced larger value of HC, MR and MS. It is worth noting that high value of HC (≈0.29 T) at 300 K can be explained due to the presence of large surface anisotropy of bead-shaped particle with aspect ratio more than 1:10 [19]. For sample AH, similar to sample EtOH, a hysteretic behaviour of M–H curve evidences a weak ferromagnetic state of the present sample at room temperature. A different shaped hysteresis loop is observed at 5 K; showing the coexistence of the AFM and FM phases in the system which bears the maximum strain in lattice (0.68%) in a network of discrete particles of average size 140 nm. Figure 8 represents the magnetic behaviour of as-prepared samples within the synthesis modified system. The temperature (T) dependence of magnetization measured at 200 Oe is shown in figures 8(a) to (c) for the samples water, EA12 and EG12, respectively. A similarity of magnetic behaviour is obtained for the samples of water and EG12 showing sharp decrease in magnetization below 260 K and 240 K, respectively, and distinct Morin transition is obtained at 236 K and 228 K, respectively, while the sample EA12 suppressed the Morin transition completely. This behaviour can be correlated with the structural analysis giving maximum 12
S Chakrabarty et al
Mater. Res. Express 1 (2014) 046104
Figure 8. (a)–(c) Temperature (T) dependence of MZFC and MFC curves measured at
H = 200 Oe for water, EA12 and EG12, respectively. The insets (a) and (c) show Tdependence of corresponding differential of MZFC curves. (d)–(e) Field-dependent M–H curves at 300 and 5 K for water, EA12 and EG12, respectively. The insets in (d)–(e) selectively reveal corresponding low field regions of the M–H curves measured at 300 and 5 K.
strain of 0.59% for water sample and minimum of 0.3% for EA12. The nanoparticles for all these three samples are of different shape with particle size maximum for EA12. Large particle size along with lesser strain in EA12 helps in dominating ferromagnetic behaviour suppressing Morin transition in this system. For the systems of water, the low value of TM may be influenced due to the reduced size effect of Fe2O3 nanorods along with the porosity of the nanorod. At temperature below 300 K, MZFC and MFC split up, which can be attributed to the presence of the competition between the shape and magnetocrystalline anisotropies [39]. For the system of EA12, the deviation of MZFC and MFC accompanied by broad maxima in MZFC at 140 K corresponding to blocking temperature (TB) is obtained. The thermo-magnetic irreversibility starts at 200 K, somewhat higher than TB along with broader ZFC peak in the system. A consistent increase in MFC with decreasing temperature is observed implying predominant FM interaction. It is worth mentioning that Hartley et al observed quite similar magnetic response for cubic iron oxide nanoparticles [54]. Figure 8(d) shows the field dependent magnetization behaviour at 300 K and 5 K of water sample. The parasitic FM behaviour is observed at 300 K, higher than the value of TM, with HC and MR being 2777 Oe and 0.0495 emu gm−1 (table 1), respectively. The high value of HC is attributed to the enhanced shape anisotropy, which is induced by the porous structure of the nanorods. This porous Fe2O3 nanorod with high value of HC is a potent material to be used in digital magnetic storage devices in the future. The M−H behaviour of EA12 in figure 8(e) shows hysteresis loops (inset of figure 8(e)) measured at 5 and 300 K which is indicative of a soft ferromagnet. The larger values of HC and MR indicate strong ferromagnetic behaviour at 5 K, which is relatively weak at 300 K. The
13
S Chakrabarty et al
Mater. Res. Express 1 (2014) 046104
results of the M(T) and M(H) measurements indicate that long-range magnetic ordering occurs, which suppresses the Morin transition. The magnetic hysteresis of EG12 at 300 K and 5 K is shown in figure 8(f). A clear hysteresis loop is observed at 300 K, which is indicative of the presence of ferromagnetic components, however, no saturation in magnetization is observed up to the maximum applied magnetic field at 300 K showing the influence of other magnetic ordering present in the system. At 5 K, a typical superparamagnetic behaviour with small hysteresis loop at low field region is observed. However, the magnetic behaviour of EG12 needs further investigation to explain its field dependence at 5 K. After analysing the magnetic response of three different groups of samples, it is clearly seen that the network type systems, in general, show high ferromagnetic behaviour where Morin transiton is almost suppressed. The blocking temperature is also absent, in general, for the network like α-Fe2O3 nanocrystals. Strong interparticle interaction, in these network-like samples, is mainly attributed to their dominating ferromagnetic behaviour. Synthesized systems with discrete morphology and homogeneous surface structure experience strong Morin transition with the presence of blocking temperature. High aspect ratio (in the case of AA and EtOH) and high lattice strain (in the case of EtOH and AH) are believed to be the main factors in determining their magnetic behaviour. The samples produced by ‘modified synthesis route’ show signature of either the Morin transition (in water and EG12) or the blocking temperature (in EA12). In this group, size and shape inhomogeneity of samples play a crucial role for their magnetic response. Finally, all the synthesised α-Fe2O3 nanocrystals with their shape, size and lattice parameters are presented in table 1 to show their major signatures in magnetic behaviour. It is seen from the table that coercivity of AA and water sample at room temperature are even higher than that measured for highly anisotropic Fe2O3 ‘nanocup’ system [2]. From the coercivity value recorded for the network type system it is evident that they are typically soft magnet, which was also seen in branched topology of porous α-Fe2O3 nanostructures [36]. The Morin transition temperature varies in a wide range of temeprature starting from 211 K for spherical AH sample to 246 K for hexagonal disc like EtOH sample. It is also noteworthy that the nanocoral shaped EG sample provides excellent magnetic moment at both 300 K and 5 K.
4. Conclusion
Here we have synthesized α-Fe2O3 nanostructure with different unique morphology by simple solvothermal technique followed by proper heat treatment. The prepared samples are characterized structurally and morphologically to understand their growth evolution. The article reports the effects of morphology in the magnetic behaviour of α-Fe2O3 nanostructures. Magnetic response varies in a wide range for the as-obtained α-Fe2O3 nanocrystals. The results have been discussed on the basis of their surface anisotropy and lattice strain. The report presents an integrated version of morphology dependent magnetic behaviour of α-Fe2O3 nanocrystals. It allows the creation of numerous recipes of α-Fe2O3 for optimizing and scaling up production towards technological applications.
14
S Chakrabarty et al
Mater. Res. Express 1 (2014) 046104
Acknowledgment
This work was financially supported by the UGC, India [42/1069/2013(SR) & 42-908/2013 (SR)] and Special Assistance Program-UGC, India. S Chakrabarty also gratefully acknowledges her fellowship from UGC-BSR Scheme, India. The authors gratefully acknowledge Professor Jason Chang, IOP, Academia Sinica, Taiwan for his thoughtful discussions and for providing SEM/EDX facilities. References [1] Lian J, Duan X, Ma J, Peng P, Kim T and Zheng W 2009 Hematite (α-Fe2O3) with various morphologies: ionic liquid-assisted synthesis, formation mechanism, and properties ACS Nano 3 3749–61 [2] Jagadeesan D, Mansoori U, Mandal P, Sundaresan A and Eswaramoorthy M 2008 Hollow spheres to nanocups: tuning the morphology and magnetic properties of single-crystalline α-Fe2O3 nanostructures Angew. Chem. 120 7799–802 [3] Cui Y and Lieber C M 2001 Funtional nanoscale electronic devices assembled using silicon nanowire building blocks Science 291 851–3 [4] Lui L, Kou H-Z, Mo W, Liu H and Wang Y 2006 Surfactant-assisted synthesis of α-Fe2O3 nanotubes and nanorods with shape-dependent magnetic properties J. Phys. Chem. B 110 15218–23 [5] Wu C, Yin P, Zhu X, OuYang C and Xie Y 2006 Synthesis of hematite (α-Fe2O3) nanorods: diameter-size and shape effects on their applications in magnetism, lithium ion battery, and gas sensors J. Phys. Chem. B 110 17806–12 [6] Garcia-Labato M A, Martinez A I, Castro-Roman M, Falcony C and Escobar-Alarcon L 2011 Correlation between structural and magnetic properties of sprayed iron oxide thin films Physica B 406 1496–500 [7] Wang D, Wang Q and Wang T 2011 Controlled synthesis of mesoporous hematite nanostructures and their application as electrochemical capacitor electrodes Nanotechnology 22 135604–15 [8] Wang L and Gao L 2011 Controlled synthesis and tunable properties of hematite hierarchical structures in a dual-surfactant system CrystEngComm. 13 1998–2005 [9] Elias V R, Oliva M I, Vaschetto E G, Urreta S E, Eimer G A and Silvetti S P 2010 Magnetic properties of iron loaded MCM-48 molecular sieves J. Magn. Magn. Mater. 322 3438–42 [10] Mitra S, Das S, Mandal K and Chaudhuri S 2007 Synthesis of a α-Fe2O3 nanocrystal in its different morphological attributes: growth mechanism, optical and magnetic properties Nanotechnology 18 275608–16 [11] Dieckmann R 1993 Point defects and transport in haematite (Fe2O3−ε) Philos. Mag. A 68 725–45 [12] Cesar I, Kay A, Martinez J A G and Gratzel M 2006 Translucent thin film Fe2O3 photoanodes for efficient water splitting by sunlight: nanostructure-directing effect of Si-doping J. Am. Chem. Soc. 128 4582–3 [13] Ohmori T, Takahashi H, Mametsuka H and Suzuki E 2000 Photocatalytic oxygen evolution on α-Fe2O3 films using Fe3+ ion as a sacrificial oxidizing agent Phys. Chem. Chem. Phys. 2 3519–22 [14] Gou X, Wang G, Park J, Liu H and Yang J 2008 Monodisperse hematite porous nanospheres: synthesis, characterization, and applications for gas sensors Nanotechnology 19 125606–11 [15] Widder K J, Senyei A E and Scarpelli D G 1978 Magnetic microspheres: a model system of site specific drug delivery in vivo Proc. Soc. Exp. Biol. Med. 158 141–6 [16] Garcon G, Garry S, Gosset P, Zerimech F, Martin A, Hannothiaux M-H and Shirali P 2001 Benzo(a)pyrenecoated onto Fe2O3 particles-induced lung tissue injury: role of free radicals Cancer Lett. 167 7–15 [17] Busch M, Gruyters M and Winter H 2006 Spin polarization and structure of thin iron oxide layers prepared by oxidation of Fe(110) Surf. Sci. 600 4166–9 [18] Walter D 2006 Characterization of synthetic hydrous hematite pigments Thermochim. Acta 445 195–9
15
S Chakrabarty et al
Mater. Res. Express 1 (2014) 046104
[19] Tadić M, Čitaković N, Panjan M, Stojanović Z, Marković D and Spasojević V 2011 Synthesis, morphology, microstructure and magnetic properties of hematite submicron particles J. Alloys Compd. 509 7639–44 [20] Teja A S and Koh P-Y 2009 Synthesis, properties, and applications of magnetic iron oxide nanoparticles Prog. Cryst. Growth Charact. Mater. 55 22–45 [21] Liu Z and Zheng Y 2011 Effect of Fe(II) on the formation of iron oxide synthesized from pyrite cinders by hydrothermal process Powder Technol. 209 119–23 [22] Su C, Wang H and Liu X 2011 Controllable fabrication and growth mechanism of hematite cubes Cryst. Res. Technol. 46 209–14 [23] Wang H, Geng W and Wang Y 2011 Preparation of nanoparticles and hollow spheres of α-Fe2O3 and their properties Res. Chem. Intermed. 37 389–95 [24] Wu W, Xiao X, Zhang S, Zhou J, Fan L, Ren F and Jiang C 2010 Large-scale and controlled synthesis of iron oxide magnetic short nanotubes: shape evolution, growth mechanism, and magnetic properties J. Phys. Chem. C 114 16092–103 [25] Tadić M, Kusigerski V, Marković D, Čitaković N, Remškar M and Spasojević V 2009 Morphological, structural and magnetic properties of α-Fe2O3 nanoparticles in an amorphous alumina matrix obtained by aqueous combustion method J. Alloys Compd. 486 839–43 [26] Mandal S and Müller A H E 2008 Facile route to the synthesis of porous α-Fe2O3 nanorods Mater. Chem. Phys. 111 438–43 [27] Liang H-F and Wang Z-C 2010 Template-free synthesis and characterization of snowflake-like α-Fe2O3 microstructures Mater. Lett. 64 2410–2 [28] Song F, Guan J, Fan X and Yan G 2009 Single-crystal star-like arrayed particles of hematite: synthesis, formation mechanism and magnetic properties J. Alloys Compd. 485 753–8 [29] Xu X, Cao R, Jeong S and Cho J 2012 Spindle-like mesoporous α-Fe2O3 anode material prepared from MOF template for high-rate lithium batteries Nano Lett. 12 4988–91 [30] Bharathi S, Nataraj D, Seetha M, Mangalaraj D, Ponpandian N, Masuda Y, Senthil K and Yong K 2010 Controlled growth of single-crystalline, nanostructured dendrites and snowflakes of α-Fe2O3: influence of the surfactant on the morphology and investigation of morphology dependent magnetic properties CrystEngComm. 12 373–82 [31] Li S, Zhang H, Wu J, Ma X and Yang D 2006 Shape-control fabrication and characterization of the airplanelike FeO(OH) and Fe2O3 nanostructures Cryst. Growth Des. 6 351–3 [32] Zhu L-P, Xiao H-M and Fu S-Y 2007 Template-free synthesis of monodispersed and single-crystalline cantaloupe-like Fe2O3 superstructures Cryst. Growth Des. 7 177–82 [33] Zhang Y, Wang S, Li X, Chen L, Qian Y and Zhang Z 2006 CuO shuttle-like nanocrystals synthesized by oriented attachment J. Cryst. Growth 291 196–201 [34] Wang X, Chen X, Ma X, Zheng H, Ji M and Zhang Z 2004 Low-temperature synthesis of α-Fe2O3 nanoparticles with a closed cage structure Chem. Phys. Lett. 384 391–3 [35] Cao M, Liu T, Gao S, Sun G, Wu X, Hu C and Wang Z L 2005 Single-crystal dendritic micro-pines of magnetic α-Fe2O3: large-scale synthesis, formation mechanism, and properties Angew. Chem. Int. Ed. 44 4197–201 [36] Yang H et al 2010 Porous α-Fe2O3 nanostructures with branched topology: growth, formation mechanism, and properties CrystEngComm. 12 1842–9 [37] Chakrabarty S and Chatterjee K 2013 Oriented growth of α-Fe2O3 nanocrystals with different morphology and their optical behaviour J. Cryst. Growth 381 107–13 [38] Tang B, Wang G, Zhuo L, Ge J and Cui L 2006 Facile route to α-FeOOH and α-Fe2O3 nanorods and magnetic property of α-Fe2O3 nanorods Inorg. Chem. 45 5196–200 [39] Mitra S, Das S, Basu S, Sahu P and Mandal K 2009 Shape-and field-dependent Morin transitions in structured α-Fe2O3 J. Magn. Magn. Mater. 321 2925–31
16
S Chakrabarty et al
Mater. Res. Express 1 (2014) 046104
[40] Gou X, Wang G, Kong X, Wexler D, Horvat J, Yang J and Park J 2008 Flutelike porous hematite nanorods and branched nanostructures: synthesis, characterisation and application for gas-sensing Chem. Eur. J. 14 5996–6002 [41] Muruganandham M, Amutha R, Sathish M, Singh T S, Suri R P S and Sillanpää M 2011 Facile fabrication of hierarchical α-Fe2O3: self-assembly and its magnetic and electrochemical properties J. Phys. Chem. C 115 18164–73 [42] Zhao Y, Dunnill C W, Zhu Y, Gregory D H, Kockenberger W, Li Y, Hu W, Ahmad I and McCartney D G 2007 Low-temperature magnetic properties of hematite nanorods Chem. Mater. 19 916–21 [43] Lu L, Li L, Wang X and Li G 2005 Understanding of the finite size effects on lattice vibrations and electronic transitions of nano α-Fe2O3 J. Phys. Chem. B 109 17151–6 [44] Kleiman-Shwarsctein A, Huda M N, Walsh A, Yan Y, Stucky G D, Hu Y-S, Al-Jassim M M and McFarland E W 2010 Electrodeposited aluminum-doped α-Fe2O3 photoelectrodes: experiment and theory Chem. Mater. 22 510–7 [45] Li Z, Lai X, Wang H, Mao D, Xing C and Wang D 2009 Direct hydrothermal synthesis of single-crystalline hematite nanorods assisted by 1,2-propanediamine Nanotechnology 20 245603–11 [46] Gupta R K, Ghosh K, Dong L and Kahol P K 2011 Structural and magnetic properties of phase controlled iron oxide rods Mater. Lett. 65 225–8 [47] Xia C, Hu C, Xiong Y and Wang N 2009 Synthesis of α-Fe2O3 hexagons and their magnetic properties J. Alloys Compd. 480 970–3 [48] Jiao F, Harrison A, Jumas J-C, Chadwick A V, Kockelmann W and Bruce P G 2006 Ordered mesoporous Fe2O3 with crystalline walls J. Am. Chem. Soc. 128 5468–74 [49] Li G, Liu M and Kou H-Z 2011 Mesoporous α-Fe2O3 nanospheres: structural evolution and investigation of magnetic properties Chem. Eur. J. 17 4323–9 [50] Cangussu D, Nunes W C, Corrêa H L S, Macedo W A A, Knobel M, Alves O L, Filho A G S and Mazali I O 2009 γ-Fe2O3 nanoparticles dispersed in porous Vycor glass: a magnetically diluted integrated system J. Appl. Phys. 105 013901–7 [51] Zysler R D, Fiorani D, Testa A M, Suber L, Agostinelli E and Godinho M 2003 Size dependence of the spinflop transition in hematite nanoparticles Phys. Rev. B 68 212408–11 [52] Sun B, Horvat J, Kim H S, Kim W-S, Ahn J and Wang G 2010 Synthesis of mesoporous α-Fe2O3 nanostructures for highly sensitive gas sensors and high capacity anode materials in lithium ion batteries J. Phys. Chem. C 114 18753–61 [53] Mansilla M V, Zysler R, Fiorani D and Suber L 2002 Annealing effects on magnetic properties of acicular hematite nanoparticles Physica B 320 206–9 [54] Zhen G et al 2011 Comparative study of the magnetic behavior of spherical and cubic superparamagnetic Iron oxide nanoparticles J. Phys. Chem. C 115 327–34
17
