Functional Magnetic Materials

3 Functional Magnetic Materials: Fundamental and Technological Aspects S.M. Yusuf Solid State Physics Division, Bhabha Atomic Research Centre, Trombay, Mumbai, Maharashtra, India

3.1

Introduction

Magnetic materials play an important role in the advancement of industrial and scientific growth. They are invariably used in power generation and transmission, electronic appliances, analogue and digital data storage, medical appliances like magnetic resonance imaging (MRI), magnetic therapy and drug delivery, sensors and actuators, scientific instruments, etc. Functional magnetic materials are a group of materials having important and interesting physical properties, which can be influenced by application of an external perturbation such as applied magnetic field. These materials are also called the smart magnetic materials of the future. Some important results of our studies, performed on various functional magnetic materials such as colossal magnetoresistance (CMR) manganites, high magnetocaloric materials, hexacyanide-based molecular materials and magnetic nanoparticles, are presented in this chapter. These materials have potential for use in information storage and processing, spintronics, drug delivery, cooling technology, etc. A large change of magnetic entropy across a magnetic ordering temperature of a material can have an application in magnetic refrigerators. This functionality of a magnetic material has enormous possibility for use as an alternative cooling technology. The basis of this is the magnetocaloric effect (MCE), defined as the cooling or heating of a magnetic material upon application of a varying magnetic field. This functionality offers the prospect of a compact, highly efficient, less noisy and environment-friendly alternative to the most commonly used vapour-compressionbased refrigeration system. The main challenges in the realization of a magnetic refrigerator are the availability of high magnetocaloric materials in large quantities that exhibit large MCE at room temperature in a moderate magnetic field as well as low hysteretic losses. The design of the magnetic field is also crucial as properly designed permanent magnet arrays can enhance the efficiency of these refrigerators to a great extent. The structural and magnetic properties of some of the interesting Functional Materials. DOI: 10.1016/B978-0-12-385142-0.00003-9 © 2012 Elsevier Inc. All rights reserved.

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high magnetocaloric materials, namely TbCo22xFex and La0.67Ca0.33MnO3, will be presented in view of their usefulness in magnetic cooling at or near room temperature and in low temperature regimes. Hexacyanometallate-based molecular magnetic materials have been studied extensively in past two decades because of their multifunctional properties like light-, pressure- and humidity-induced magnetization and its reversal, mixed ferro-ferrimagnetism, solubility and transparency. These materials can be grown easily at an ambient temperature using synthetic chemistry methods. The nature of short- or long-ranged structural correlations in these compounds plays an important role in controlling these functionalities. The role of various short- and long-ranged structural correlations leading to various defects and interesting magnetic properties will be presented for the (CoxNi1 2x)1.5[Fe(CN)6]
zH2O compounds. The correlations among various magnetic properties have been studied as a function of Co ion substitution. The control of magnetization as well as its polarity in (CuxMn12x)1.5[Fe(CN)6]
zH2O hexacyanide compounds as a function of magnetic field and temperature have been studied. Possible applications of the magnetic pole reversal phenomenon in magnetoelectronic and magnetocaloric devices such as magnetic memory and magnetic cooling/heatingbased constant temperature baths will also be revealed. The thickness- and stoichiometry-dependent magnetic properties of electrochemically prepared crystalline thin films of Prussian blue analogues (PBAs) KjFek[Cr(CN)6]l
mH2O will also be discussed. Magnetic nanoparticles have been the focus of research because of their attractive properties, which potentially could see use in data storage and processing, spintronics, catalysis, drug delivery, MRI, environmental studies, etc. These materials show unusual magnetic behaviour compared with bulk materials, mainly because of their surface/interface effects, electronic charge transfer and magnetic interactions. The typical phenomena related to nanoscale structures are the increased relevance of surface effects, defects and the existence of new or metastable phases. Hence these phenomena can be exploited to develop new magnetic nanoparticles. The role of structural and magnetic properties of the nanoparticles in designing the magnetic nanoparticle systems for their use in spintronics, high-density magnetic recording, biological applications, radionuclide separation, etc. has been investigated by us. Rare-earth-based CMR manganites exhibit a range of extraordinary magnetic, electronic and structural properties including CMR effect, charge ordering, magnetic-field-induced changes in structure and transport properties. However, the microscopic understanding of various magnetic and electronic phase transitions due to ionic size effects is needed to explore their use in various technological applications. This aspect will be addressed by studying the structural and magnetic properties of the Dy-substituted La12xCaxMnO3 CMR manganites. The results of our detailed studies, performed for compounds of the functional magnetic materials mentioned above, are given in the forthcoming sections of this chapter.

Functional Magnetic Materials: Fundamental and Technological Aspects

3.2

113

Magnetocaloric Effect

The MCE is among the most fundamental physical properties of magnetic solids. It describes thermal behaviour of a material in the presence of magnetic field. It is the change in magnetic entropy of a magnetic material when an external magnetic field is applied. In the case of a ferromagnet, near its magnetic ordering temperature TC, the isothermal application of a magnetic field reduces the magnetic entropy. The magnetic material in turn is heated up owing to the increase in its lattice entropy. In a reversible process, upon adiabatic removal of the magnetic field, the magnetic material is cooled as the magnetic entropy increases at the expense of the lattice entropy. This warming and cooling of a magnetic material in response to a changing magnetic field is similar to the warming and the cooling of a gaseous medium in response to an adiabatic compression and expansion. Therefore, by magnetizing and demagnetizing, a functional magnetic material that shows a large MCE can operate as a magnetic refrigerant. The temperature of the final and the initial states of the functional material depends on various intrinsic and extrinsic factors. The important intrinsic material factors that determine the MCE of a material are the chemical composition, crystal structure and the magnetic state of the compound. The extrinsic factors include the temperature, surrounding pressure and the change in magnetic field. Over the past few years MCE has been intensively investigated because of its potential application in magnetic refrigeration near room temperature. The development of magnetic refrigeration technology, based upon the MCE [1], can be thought of as an alternative to the conventional gas compression technique. By magnetic refrigeration technology, a cooling efficiency of up to 60% of a Carnot cycle can be reached, whereas it is only up to 10% for refrigeration by conventional gas compression. Therefore, magnetic refrigeration is more efficient than conventional gas compression. Other advantages of magnetic refrigeration technology are that it is an environmental friendly cooling technology as it does not use any ozone-depleting or global-warming greenhouse-effect gases and hazardous chemicals. MCE was discovered in 1881 [1], but the major advance occurred in the late 1920s when Debye [2] and Giauque [3] independently proposed cooling by adiabatic demagnetization. A few years later, in 1933 [4], the procedure was demonstrated for the first time using paramagnetic salts (e.g. Gd2(SO4)3
8H2O). The temperature achieved was 0.25 K. As a proof-of-principle, in 1997, a magnetic refrigerator demonstrated that it is a practical and competitive cooling technology of the future around room temperature [5] with potential energy savings. Since then there has been an increase in research in this area.

3.2.1

Fundamentals of Magnetic Cooling and Heating

When a magnetic material is subjected to a magnetic field, the magnetic moments of the atoms get re-oriented. On applying a varying magnetic field, two processes

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may occur in the magnetic material. The first is the isothermal process that occurs when the magnetic field is applied but the material remains connected to the surroundings (heat reservoir) and, therefore, remains at constant temperature. The entropy of the magnetic material is then decreased. The change in the magnetic entropy is denoted by ΔSM. The lattice entropy of the material, on the other hand, increases in the form of heat, which is released to the surroundings. The second is an adiabatic process that occurs when the magnetic field is removed but the material is isolated from the surroundings and, therefore, the total entropy of the material remains constant. In the latter process, the temperature of a magnetic material is decreased by ΔTad and is called an adiabatic temperature change. This warming and cooling of a functional material in response to the application and removal of an external magnetic field is called the MCE. The magnetic free energy of a system can be written as FðT; HÞ 5 U 2 TS 2 HM

ð3:1Þ

Here U 5 internal energy T 5 temperature S 5 entropy H 5 magnetic field M 5 magnetization.

Taking the partial derivative, we have dF 5 2 SdT 2 MdH

ð3:2Þ

@F


52S @T H

ð3:3Þ

@F


52M @H T

ð3:4Þ

By applying the following Maxwell’s relation 

      @ @y @ @y 5 @z @x z x @x @z x z

ð3:5Þ

we have 

      @ @F @ @F 5 @H @T H T @T @H T H

ð3:6Þ

If both the magnetization and entropy are continuous functions of the temperature and magnetic field, then the infinitesimal isobaricisothermal magnetic entropy

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115

change (S) can be related to the magnetization (M), the magnetic field strength (H) and the absolute temperature (T) by using Eqs (3.3) and (3.4) as     @S @M 5 ð3:7Þ @H T @T H The magnetic entropy change, ΔSM (T, H), is calculated by ΔSM ðT; HÞ 5 SM ðT; HÞ 2 SM ðT; 0Þ  ðH  @MðT; HÞ dH 5 @T 0 H

ð3:8Þ

The magnetic entropy change is related to the heat capacity of the material as ΔSM ðT; HÞ 5

ðT 0

CðT; HÞ 2 CðT; 0Þ dT T

ð3:9Þ

where C(T, H) and C(T, 0) are the values of the heat capacity measured in a field H and in zero field, respectively. Therefore the adiabatic temperature change can be evaluated by integrating Eq. (3.9) over the magnetic field as ΔTad 5 2

ðH 0

  T @M dH CP;H @T H

ð3:10Þ

From Eqs (3.8) and (3.10), it can  be seen  that for large MCE, i.e. large ΔSM and ΔTad, one needs to have large @M=@T H and small C(T, H).

3.2.2

Magnetic Transition and Magnetocaloric Effect

The magnitude of the magnetic entropy change and its dependence on temperature and magnetic field depend strongly on the nature of the magnetic phase transition. Most ferromagnetic materials show a second-order magnetic phase transition. In a first-order transition, the MCE is concentrated over a narrow temperature range, whereas second-order transitions are usually spread over a broad temperature range. This is important for an active magnetic refrigeration. Even though the change in magnetic entropy is large in materials showing a first-order magnetic transition, they exhibit a large thermal and field hysteresis on variation of magnetization with temperature and magnetic field, respectively, which, for practical magneticrefrigeration (MR) applications, should be as small as possible.

3.2.3

Relative Cooling Power

The relative cooling power (RCP) or refrigeration capacity is an important parameter in determining the cooling efficiency of a magnetic refrigerant material for its

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Functional Materials

use in magnetic refrigeration. It is the heat transfer between hot and cold reservoirs during an ideal refrigeration cycle. This represents numerically the area under the 2ΔSM versus T curve. The RCP is defined as [6] RCP 5

ð THot

ΔSM dT

ð3:11Þ

TCold

Here, THot and TCold are higher and lower temperatures at half maximum of the ΔSM peak and can be considered as the temperatures of the hot and cold reservoirs, respectively.

3.2.4

Magnetocaloric Materials

Current research aims to achieve new functional materials that show a giant magnetocaloric effect (GMCE) with a small change in magnetic field near room temperature. For example, the rare-earth element gadolinium (Gd) has been widely investigated for its use as an active magnetic refrigerant near room temperature [7]. Later it was shown that Gd5Si2Ge2 alloy exhibits GMCE [7]. Some other compounds that show a large MCE are MnFe(P12xAsx) [8], La(Fe132xSix) [9], NiMnGa Heusler alloys [10,11], RCo2 intermetallic compounds [12] and perovskites [1315]. For practical applications, materials with a large MCE over a wide range near room temperature are desirable. Hence it is challenging to increase the operating temperature range of the magnetic refrigerant materials. We have investigated the effect on MCE due to substitution of Fe at the Co site in TbCo2 [16]. We have shown that an appropriate (Fe) chemical substitution not only tunes the transition temperature to around room temperature but also increases the operating temperature range of magnetic refrigeration without significant reduction in the ΔSM value. In TbCo22xFex samples, TC increases with Fe substitution (231, 275, 290 and 303 K for x50, 0.06, 0.08 and 0.1, respectively). The observed increase in TC on substituting Fe at the Co site is due to the enhanced exchange interaction between 3d transition metal ions [17]. Figure 3.1 shows the variation of 2ΔSM with temperature for the samples with x 5 0, 0.06 and 0.1. The magnetic entropy change ΔSM was calculated by using Eq. (3.8). The maximum value of ΔSM is found to be around TC and it increases with the increase in the applied magnetic field. At a field variation of 3 T, 2ΔSM values are found to be 5.0, 2.6 and 2.5 J kg21 K21 for x 5 0, 0.06 and 0.1, respectively. We also observe that on increasing the Fe content from x 5 0.06 to x 5 0.1, TC is successfully tuned to a desired value (room temperature) without any significant change in the value of 2ΔSM. The values of RCP were calculated by using Eq. (3.11). The values of RCP at a field variation of 5 T are found to be 357, 299 and 271 J kg21, for x 5 0, 0.06 and 0.1, respectively. For x 5 0.06, the operating temperature range (defined as the difference between THot and TCold) is B89 K at Δμ0H 5 5 T and, for the x 5 0.1 sample, the operating temperature range is B95 K, whereas for the sample x 5 0, it is only 50 K [16]. Such a broad operating

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117

–ΔSM (J kg–1 K–1)

temperature range is important for practical applications. Thus, by substituting Fe in TbCo2, we not only tune the TC towards room temperature but also increase the operating temperature range. To increase the operating temperature range further, composites of these alloys with different Fe concentrations can be synthesized. A series of such compositions can be used in a cascaded way for cooling over a broad temperature range. In another study, we have investigated the MCE in the Fe-doped La0.67Ca0.33MnO3 CMR perovskite [18]. The parent compound La0.67Ca0.33MnO3 [13,15] shows a first-order magnetic phase transition. For the Fe-substituted La0.67Ca0.33Mn0.9Fe0.1O3 compound, the magnetic transition was found to be quite broad (indicating a secondorder-type magnetic phase transition) compared with that for the parent compound. Here, the presence of the MCE over a broad temperature region with a reduction in both thermal and field hysteresis has been shown. For Δμ0H 5 2 T, it reaches the maximum value of 0.83 J kg21 K21 at 113 K [18]. The rounding of the first-order magnetic phase transition after Fe doping in the parent compound La0.67Ca0.33MnO3 possibly arises owing to an introduction of the quenched disorder. The value of RCP shown in Figure 3.2 has been derived to be B54 J kg21 for Δμ0H 5 2 T.

(A)

5

x=0

4

1T 3T

(B)

1T 3T

x = 0.06

(C)

1T 3T

x = 0.1

3 2 1 0

180 200 220 240 260 200

240

280

320 240

260

280

300

320

Temperature (K)

Figure 3.1 The magnetic entropy change 2ΔSM with temperature at various fields for x 5 0, 0.06 and 0.1 compositions of TbCo22xFex compounds.

Figure 3.2 The ΔSM versus T curve of La0.67Ca0.33Mn0.9Fe0.1O3 compound for Δμ0H 5 2 T. The shaded area is the RCP.

–ΔSM (J kg–1 K–1)

0.9 2T 0.6

0.3 Thot

Tcold 0.0

50

100 150 Temperature (K)

200

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Functional Materials

χDC (emu g–1 Oe–1)

0.03 Warming cycle Cooling cycle

Figure 3.3 χDC versus T curves of La0.67Ca0.33Mn0.9Fe0.1O3 in the field-cooled warming and cooling cycles.

0.02

0.01

0.00 50

100

150

Temperature (K)

Thus, by substituting Fe at the Mn site, we are able to increase the operative temperature range considerably and to retain a large value of the RCP. There is no observable hysteresis over the studied temperature range. The DC susceptibility (χDC) versus T curves of La0.67Ca0.33Mn0.9Fe0.1O3 compound in the field-cooled cooling and warming cycles are shown in Figure 3.3. The thermal hysteresis is found to be very small. For the parent compound La0.67Ca0.33MnO3, a thermal hysteresis of B5 K was reported [19]. By substituting Fe at the Mn site in the parent compound La0.67Ca0.33MnO3, we successfully introduced the quenched disorder which has modified the nature of the phase transition from the first order to the second order. Consequently, the thermal hysteresis losses have been reduced. The efficiency of RCP is less for materials showing the hysteresis in ΔSM in the magnetic field or thermal cycle. The negligible thermal or magnetic field hysteresis in magnetization for La0.67Ca0.33Mn0.9Fe0.1O3 perovskite over a wide temperature range around the transition temperature makes this material very useful as a magnetic refrigerant.

3.2.5

Challenges in Using GMCE Materials in Magnetic Refrigerators

Designing and testing of large-scale materials for commercial cooling devices is an area where an intensive research is required. The major problems are related to the production of the magnetic refrigerants on a large scale and fabrication of these materials at a low cost for regenerators (spheres, foils, etc.) without sacrificing their MCE. Most of the GMCE materials are intermetallic compounds and are brittle in nature. If we consider an average lifetime of 1015 years for a commercial magnetic cooling device, and that the refrigerant material will undergo 50500 million cycles, it is likely that most of these brittle materials will undergo some fracture. Another fact is the delay for the ΔTad rise to achieve its maximum value in a cycle. For most of the GMCE materials, the measured temperature changes may be about 40% less than the equilibrium value because of the kinetics of the phase transformation. This is a problem as the magnetic refrigerators will operate between 0.5 and 10 Hz and most of the GMCE may not be working properly during the rapid change of magnetic field. Other problems are the thermal and magnetic field

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119

hysteresis, which limit the use of MCE materials. Hence, there is a need for new and better functional magnetic materials, and to improve the critical properties of existing materials for their technological applications.

3.3

Molecular Magnetic Materials

Traditional magnetic materials are composed of inorganic atoms or ions of transition metals, lanthanides and/or actinides containing spin units. A new area of research and development is in molecular magnetism [2023], where magnets are constructed from molecular building blocks bearing magnetic moments. Molecular magnets [2427], as the name suggests, are composed of molecular subunits. Molecular magnetism essentially involves the study of organic molecules having magnetic moment. The advantage of using organic molecules as building blocks in magnets comes in terms of additional functionalities of these organic molecules. For example, one can take advantage of various molecular properties of the organic molecules, such as self-assembly, tunability, solubility, transparency, chirality, low density, biocompatibility and mechanical flexibility, and attach them with magnetic properties to get multifunctional or smart materials. Molecular magnets can be classified into the following four categories.

3.3.1

Purely Organic Molecular Magnets

In these materials, the unpaired electrons reside in the s or p orbitals of the organic ligands. The organic ligands, in general, are mostly diamagnetic with closed shell structure. Any unpaired electron in s or p orbitals gives an instability to the ligand, which tends to make a covalent bond with another ligand. Even if we make an organic radical (with an odd number of electrons) stable by attaching a bulky aromatic ring to it, the stabilization of a triplet state (parallel alignment of spins) is difficult, so they prefer to get antiferromagnetically coupled with no spontaneous magnetic moment. This is why only a few purely organic compounds show bulk magnetization at very low temperature. The very first purely organic molecular magnet was the β-phase crystal of the p-nitrophenyl nitronyl nitroxide (p-NPNN) radical [28] with a magnetic ordering temperature (TC) of 0.6 K.

3.3.2

OrganicInorganic Molecular Magnets

In these compounds, the unpaired spins reside not only in the s and p orbitals of organic radicals but also in the d and f orbitals of inorganic metallic ions. The magnetic properties of these molecules arise because of the magnetic interactions between organic radicals and metallic ions. The area of organicinorganic (mixed) molecular magnets was pioneered by initial work on charge transfer salts of tetracyanoethylene (TCNE) by Miller, Epstein and co-workers [2931]. The first charge transfer compound that was discovered, [Fe(C5Me5)2][TCNE], where C5Me5 5 pentamethylcyclopentadienide,

120

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orders ferromagnetically at 4.8 K [31] with a large value of the coercive field (B1 kOe) at 2 K. A similar compound, [Mn(C5Me5)2][TCNE], shows magnetic ordering at 8.8 K [32] because of a larger moment at the Mn (S 5 1) site compared with the Fe (S 5 1/2) site. Subsequently, the reaction of [VI(C6H6)2]1, which is iso-electronic to [Fe(C5Me5)2], with TCNE in a dichloromethane solution yielded V[TCNE]2
1/2(CH2Cl2) [33] with a magnetic ordering temperature of 400 K. The high TC results from the ferrimagnetic ordering between VII (S53/2) and [TCNE]2 (S 5 1/2) moments. Another interesting radical with an open shell (p-electron) structure is tetrathiafulvalene (TTF). The derivatives of TTF with BEDT (bisethylenedithio) give a variety of functional molecular systems including organic metals and superconductors. For example, (BEDTTTF)3CuCl4
H2O is the first organic compound with metallic properties down to 200 mK [34]. However, because of the localized d-electron spins, this system is paramagnetic down to 200 mK. The compound (BEDTTTF)3[MnCr (C2O4)3], a layered structure of alternating oxalate-bridged hexagonal network and β-packed BEDTTTF molecules, is the very first ferromagnetic organic metal with a TC of 5.5 K [35]. Another very interesting compound, (BEDTTTF)2H2OFe (C2O4)3C6H5CN [36], which has a superconducting transition temperature of 7 K, is the first molecular paramagnetic superconductor.

3.3.3

Inorganic Molecular Magnets

The unpaired spins in these systems reside mainly in the d or f orbitals of metal ions. The exchange pathways are provided by the organic radicals, which do not contain spins. The metal hexacyanides or hexacyanometallates, An[B(CN)6]m
zH2O, where A and B are the 3d and 4d transition metal ions, are the most studied compounds in this category. Prussian blue, Fe4[Fe(CN)6]3
14H2O [37], is the parent compound of the hexacyanide family and, therefore, such compounds are also called PBAs or derivatives of Prussian blue. A wide variety of functional magnetic compounds with very interesting properties are obtained by substituting the two Fe ions in Prussian blue with other suitable magnetic ions. For example, V[Cr(CN)6]x
yH2O compounds [38,39] show magnetic ordering temperatures up to 376 K, which is well above room temperature. Some PBAs are photo-active, i.e. their magnetization changes with an exposure to light. The cobaltiron-based PBAs, AxCoy[Fe(CN)6]
zH2O [40,41], where A is alkali metal cation, switch between a paramagnetic and a ferromagnetic state with an exposure to a visible light. The mixed ferro-ferrimagnetic compounds, {NixMn12x}1.5[Cr(CN)6]
zH2O, show the novel phenomenon of magnetic pole reversal because of incorporation of ferromagnetic (between NiII and CrIII ions moments) as well as antiferromagnetic (between MnII and CrIII ions moments) interactions [42]. The phenomenon of magnetic pole reversal is very important from a technological point of view as it can be used in the next generation of magnetic memories. The control of the magnetic pole reversal by light, (Fe0.40Mn0.60)1.5[Cr(CN)6]
7.5H2O [43], by pressure, Rb0.64Ni0.31Mn0.87[Fe(CN)6]
2.8H2O [44], and by humidity, (CoxMn12x)[Cr(CN)6]2/3
zH2O [45], has also been reported for PBA compounds.

Functional Magnetic Materials: Fundamental and Technological Aspects

3.3.4

121

Molecular Magnetic Clusters

These compounds are small clusters of molecules containing a finite number of magnetic centres. These materials are also called single molecular magnets because one molecular unit behaves like a tiny magnet below a certain (blocking) temperature. One of the most studied and firstly discovered molecules in this group is the cluster of Mn12-acetate [Mn12O12(CH3COO)16(H2O)4]
2CH3COOH
4H2O [46,47]. The unit cell of the crystal has a ferromagnetic core consisting of Mn12O12 clusters separated by the non-magnetic CH3COOH radical and coordinated water molecules. The inner four Mn atoms (S 5 3/2) of an Mn12O12 cluster are antiferromagnetically coupled to the outer eight Mn atoms (S 5 2), yielding a net spin S 5 10 [46]. These compounds are known to exhibit the phenomenon of macroscopic quantum tunnelling of magnetization below the blocking temperature. The hysteresis curve for the Mn12-acetate cluster depicts many steps in its magnetization value before saturation. The observed magnetization behaviour can be attributed to thermally assisted resonant tunnelling between (2S 1 1) quantum spin states in the Mn12-acetate cluster [47]. High-spin molecules with large anisotropy can have a very stable orientation of the molecular magnetic moment. Therefore, such systems can provide the ultimate limit of high-density magnetic memory if single-molecule magnets at room temperature can be synthesized. These compounds are also potential candidates for applications in quantum computers.

3.3.5 Functionalities in Molecular Magnets As is evident from the above paragraphs, molecular magnetic materials have the clear advantage of low density, excellent structural integrity and enhanced functionalities, which make them suitable for new and advanced magnetoelectronic device applications. These magnetic materials can be grown easily at room temperature using organic, organometallic and/or coordination metal chemistry methods. The biggest advantage of the ambient temperature flexible synthesis is an easy incorporation of various molecular functionalities, for example photo-activity, electrical conductivity, polarizability and transparency into these magnetic compounds. They are the right class of magnetic materials where magnetic properties can be combined with their mechanical, electrical and/or optical properties. So, they can be called the ‘smart materials’ of the future. In this chapter we will give a few examples of Prussian blue-based molecular magnetic compounds showing interesting magnetic properties.

3.3.6

Controlling the Magnetic Hardness by Co Substitution in the (CoxNi12x)1.5[Fe(CN)6]
zH2O (x 5 0, 0.25, 0.5, 0.75 and 1) PBAs

Prussian blue-based molecular magnetic materials are generally soft as they have very small values of coercivity (HC). For example, the observed values of HC for V[Cr(CN)6]0.86
2.8H2O [39], Cr3[Cr(CN)6]2
10H2O [48] and CsNi[Cr(CN)6]
2H2O [49] are below 100 Oe. A high value (of the order of 1 kOe) of coercivity is very

122

Functional Materials

1800

297 K

(A)

1200

x=0 x = 0.25 x = 0.5 x = 0.75 x=1

Lattice parameter (Å)

X-ray intensity (a.u.)

much essential for their use in permanent magnets. In the present example, we have studied the control of magnetic hardness in Co-substituted (CoxNi12x)1.5[Fe(CN)6]
zH2O (x 5 0, 0.25, 0.5, 0.75 and 1) [5052] PBAs. The correlations of the coercivity with other parameters like structural disorder, ordering temperature, saturation magnetization and lattice parameters have been brought out. Polycrystalline compounds of the series were prepared by the co-precipitation method [52]. X-ray powder diffraction patterns were recorded at room temperature using a CuKα radiation. The infrared spectra were recorded by incorporating the samples in a KBr pellet on a Bomen MB-102 FTIR spectrometer. DC magnetization measurements were performed using a commercial vibrating sample magnetometer. Figure 3.4A shows the X-ray diffraction (XRD) patterns for the compounds of the series. The structural analysis, using Rietveld refinement [53], reveals a face-centred cubic structure (space group 5 Fm3m) for all the compounds. However, the lattice constant (a) does not vary linearly (Figure 3.4B) with Co doping but has a dip at the x 5 0.75 composition. The infrared spectra of the compounds (Figure 3.5) show two CN stretching frequencies around 2165 and 2101 cm21 corresponding to Fe31CNCo21 /Ni21 and Fe31NCCo21 /Ni21 -type linear chains, respectively [50,54]. However, for

600 0 15

30

45

60

75

Scattering angle (°)

10.32

(B)

10.28 10.24 10.20 10.16 0.0

0.2

0.4

0.6

0.8

1.0

Composition (x)

Figure 3.4 (A) Room temperature X-ray diffraction patterns of (CoxNi12x)1.5[Fe(CN)6]
zH2O compounds. (B) Variation of lattice parameter with composition [52].

Transmission (%)

95

Figure 3.5 Infrared spectra of (CoxNi12x)1.5[Fe(CN)6]
zH2O compounds at room temperature. Positions of arrows show the absorption corresponding to CN stretching frequencies [52].

x=0

76

x = 0.25 x = 0.5

57

x = 0.75

38 1900

x=1 2000 2100 2200 Wave number (cm–1)

2300

Functional Magnetic Materials: Fundamental and Technological Aspects

123

10

12 x = 0.5

8 0.0

x=1

0.3

0.6

0.9

x

5 x = 0.75

0

Oe )

–3

16

8

(10 emu g

20

12

–1

x = 0.25

(C)

0.2

χ

15

(B)

24

–1 –1

0.4

(A)

χ.T (emu K Oe–1 g–1)

x=0

TC (K)

Magnetization (emu g–1)

20

–1

the x 5 0.75 composition, one more absorption at 2055 cm21 is observed, which corresponds to diamagnetic Fe21CNCo31 /Ni31 -type linear chains in these compounds [50,52,54]. The presence of {Fe21CNCo31 /Ni31}-type chains causes a sharp dip in the ˚ lattice parameter as the ferrocyanides have a lattice constant smaller by 0.150.4 A 31 21 21 compared with corresponding ferricyanides {Fe CNCo /Ni } [55]. The presence of the last two types of chains in these compounds can be viewed as structural defects. Here, the partial occupancies of the Fe, C, N and the non-coordinated water molecule sites, and the random distribution of the Co and Ni ions at the (0, 0, 0) site, constitute the main structural defects [52]. Figure 3.6A and B show the field-cooled (FC) magnetization curves and the variation of magnetic ordering temperature (TC), respectively, with the Co doping. Interestingly, the observed TCs for the compounds follow a behaviour similar to lattice parameter variation. This finding confirms the presence of diamagnetic chains (Fe21CNCo31 /Ni31 ) in the x 5 0.75 composition. Figure 3.6C depicts the plots of χ
T versus T for the compounds of the series. For all compositions (except x 5 1), the χ
T versus T curve rises continuously until TC is reached, with lowering of temperature, a behaviour typical for a ferromagnet [56]. However for x 5 1, the χ
T versus T curve shows a minimum around 31 K before rising to its maximum around TC, as expected for a ferrimagnet [56]. This finding is also consistent with the derived paramagnetic curie temperatures (Θ) from the fitting of χ 21 versus T curves (Figure 3.6D) where (for x51) Θ is negative whereas a positive Θ is obtained for all other compounds of the series. TCs for the ferromagnetic (except x 5 1) compounds, therefore, have some correlation with the amount of structural disorder. The analysis of high field magnetization data (Figure 3.7A) reveals a ferromagnetic ordering among low-spin Co21 (S 5 1/2), low-spin Fe31 (S 5 1/2) and high-spin Ni21 (S 5 1) for Co-doped compounds, and ferrimagnetic ordering between low-spin Fe31 (S 5 1/2) and high-spin Co21 (S 5 3/2) for the x51 composition. The low-spin state of Co21 ions in all the Co-doped compounds may arise

(D)

4 0 42

x = 1.125 x = 0.5 x = 0.25

63

84

x=0 x=1

0.0 0

15

30

Temperature (K)

45

60

0

50

100

Temperature (K)

Figure 3.6 (A) FC magnetization curves of (CoxNi12x)1.5[Fe(CN)6]
zH2O compounds at 200 Oe. (B) Variation of TC with the Co doping. (C) χ
T versus T curves. (D) χ21 versus T curves with linear fitting between 30 and 85 K [52].

150

Functional Materials 55

x=0 x = 0.25 x=1

44

x = 0.50 x = 0.75

33 55

22

(B)

50 45

11 1.5 K

40 35

0 0

20

40

60

Coercive field

32

Remanence

3

24

2 16 1.5 K

1 8

0.0 0.2 0.4 0.6 0.8 1.0 x

Magnetic field (kOe)

(C)

4 Coercive field (kOe)

Magnetization (emu g–1)

(A)

Remanence (emu g–1)

124

0 80

0.00

0.25

0.50

0.75

1.00

Composition (x)

Figure 3.7 (A) M versus H curves of (CoxNi12x)1.5[Fe(CN)6]
zH2O compounds at 1.5 K. Variation of (B) max magnetization and (C) HC and MR with Co substitution [52].

because of the modification of the crystal field around the Co2 1 ion in the presence of the Ni21 ion. Consequently, the magnetic interaction changes from ferrimagnetic (x 5 1) to ferromagnetic for all Co-doped compounds owing to the change of orbital symmetry [50,52]. The variation of maximum magnetization (Figure 3.7B) with Co doping depicts a behaviour similar to lattice parameter and TC variation. The observed values of the coercive field (HC) and remanence (MR) obtained from the 1.5 K hysteresis curves are shown in Figure 3.7C. The values of HC are of the order of a few kOe, which are one or two orders of magnitude higher than that for other compounds of the hexacyanide family. The HC and MR values decrease linearly with Co substitution up to x 5 0.75, and thereafter an upturn is seen. It is therefore very much evident from these observations that various parameters like lattice constant, TC, maximum magnetization, HC, and MR are highly correlated and have minima around x 5 0.75 composition. This minimum is caused by the maximum amount of structural disorder [52] in these compounds.

3.3.7

Implications of the Magnetic Pole Reversal Phenomenon in the Cu0.73Mn0.77[Fe(CN)6]
zH2O Molecular Magnetic Compound

In 1948, Neel [57] envisaged the existence of negative magnetization in ferrimagnets. However, a microscopic understanding of it remains unclear. We observed the peculiar phenomenon of magnetic pole reversal (negative magnetization) in the Prussian blue-based molecular magnetic compound Cu0.73Mn0.77[Fe(CN)6]
zH2O. The first microscopic (experimental) explanation of this unusual phenomenon was given by us using reverse Monte Carlo simulation [58] and Rietveld refinement techniques on the neutron scattering data [59]. Our study revealed an antiferromagnetic ordering of Mn ionic moments with respect to Cu as well as Fe ionic moments. The different temperature dependences of the Cu, Mn and Fe sub-lattice magnetizations cause

Magnetization (emu g–1)

Functional Magnetic Materials: Fundamental and Technological Aspects

125

Figure 3.8 FC magnetization curves of Cu0.73Mn0.77[Fe(CN)6]
zH2O compound at various fields.

2 1 0 1.5 kOe 1 kOe 500 Oe 200 Oe

–1 6

12

18

24

30

Temperature (K)

a magnetization cross-over for this compound [59]. Here we present the control of the magnetic pole reversal by temperature and magnetic field. We also show the physics principles for its use in magnetoelectronics as well as magnetocaloric device applications. A polycrystalline sample of Cu0.73Mn0.77[Fe(CN)6]
zH2O was prepared by the precipitation method. The sample was characterized thoroughly by neutron diffraction, X-ray fluorescence and Mo¨ssbauer spectroscopic techniques. DC magnetization measurements were performed using a commercial (Oxford Instruments) vibrating sample magnetometer as a function of temperature and magnetic field. FC DC magnetization curves at various fields are shown in Figure 3.8. For fields less than 1 kOe, the FC magnetization curves increase sharply below 20 K, followed by a peak around 13.5 K and become negative below 10 K. However, it is quite evident that the magnitude of negative magnetization changes with applied magnetic field and it is no more negative at any temperature for fields above 1 kOe. The reversal of Mn spins in the direction of applied field could be a possible reason for the observed magnetization behaviour. It is also noticeable that the observed magnetization cross-over is not sharp. This is probably due to the presence of various structural defects [59] in this compound. The consequence of the magnetic-field-dependent magnetization behaviour is shown in Figure 3.9. The magnetization of the compound can be switched between negative and positive values without any noticeable decay in its magnitude by changing the applied fields (100 and 700 Oe) in a cyclical manner. The observed field-actuated reproducible magnetization reversal can lead to new device applications such as magnetic data storage (memory) and switches. The basic requirement of a data storage material is the presence of two noticeably different values (bistable states) of a physical property, e.g. electrical resistance, magnetization, electrical polarization and reflectivity. These two states of the material should be accessible with the help of an external actuator like heat, light and electric and magnetic fields. In our experiment, the realization of the two (positive and negative) magnetization states was achieved by externally applied magnetic fields (700 and 100 Oe, respectively). However, once the magnetic field is switched off,

126

Functional Materials

10 K

700

0.03 0.00

400

–0.03

Field (Oe)

Magnetization (emu g–1)

0.06

Figure 3.9 Switching of magnetization of Cu0.73Mn0.77[Fe(CN)6]
zH2O compound between positive and negative values under the application of 700 and 100 Oe fields, respectively, in a cyclical manner.

100

–0.06 0

500

1000

1500

2000

Time (s)

Figure 3.10 Normal and inverse magnetocaloric effects present in Cu0.73Mn0.77[Fe(CN)6]
zH2O compound at low fields.

–ΔSm (ergs gauss–1 K–1)

600 1 kOe 2 kOe

400

200

0

–200 0

9

18

27

36

45

Temperature (K)

the corresponding magnetization state ceases to exist, indicating that this kind of material may be suitable for volatile magnetic memories [60]. Another interesting aspect of the magnetic pole reversal phenomenon is seen in the unusual MCE [6163]. A positive value of 2ΔSm results in a decrease of temperature of the material [62]. The MCE curves for the present compound at 1 and 2 kOe fields are shown in Figure 3.10. A polarity reversal of 2ΔSm is quite evident. This polarity reversal of 2ΔSm can be used in a novel device such as a magneticentropy-driven constant temperature bath [60]. The positive value of 2ΔSm below 20 K will result in cooling and it will be maximum at 15 K where a peak in the 2ΔSm value is obtained. However, the cooling starts decreasing below 15 K and becomes zero at 13 K. On further cooling below 13 K, owing to inverse MCE [63], an increase in the temperature of the material will take place. Therefore, the temperature of the MCE material will always tend to the equilibrium temperature of 13 K. However, this mechanism will work only when the strength of the external heat perturbation is within the maximum and minimum limits of 2ΔSm. Furthermore, the

Functional Magnetic Materials: Fundamental and Technological Aspects

127

stability of the equilibrium temperature around 13 K depends on the sharpness of the ΔSm versus T curve. A steep rise and fall in ΔSm across the ΔSm 5 0 point would lead to a better stability of the equilibrium temperature [60].

3.3.8

Thickness- and Stoichiometry-Dependent Magnetic Properties of Electrochemically Prepared Crystalline Thin Films of PBAs KjFeIIk[CrIII(CN)6]l
mH2O

PBAs constitute a linear ACNB bridge in three dimensions, where the CN unit promotes the strong magnetic interactions between adjacent spin centres. The magnetic properties can be tailored by substituting a wide range of metal ions with different spin states and oxidation states at the A/B sites, as this gives considerable control over the nature and the magnitude of the local magnetic exchange interactions. Apart from substituting a wide range of metal ions at A/B sites, one can also modify the magnetic properties by varying the A/B ratio with an introduction of alkali metal ions into the structure of PBAs. One example is CsINiII[CrIII(CN)6]
2H2O, where TC90 K was reported compared with TC60 K for Ni3[Cr(CN)6]2
12H2O [49]. Alkali metal ions are usually introduced in the structure by taking an additional compound of alkali metals as a starting material during the process of synthesis. However, we report for the first time that the alkali metal ions can be introduced into the structure of the thin film of PBAs without taking any additional compound of alkali metals but by using a suitable electrode voltage during the electrochemical deposition process. Most of the PBAs studied so far, including those with a magnetic ordering temperature above or close to room temperature, mainly exist in bulk powder form. However, many potential applications of these exciting materials will require fabrication of thin films. There are very few reports in the literature of thin films of PBA owing to limitations in fabrication methods. A thin film of Prussian blue was first deposited using the electrochemical method by Neff [64]. Later, epitaxial Prussian blue films onto a single-crystal Au (110) substrate were prepared by Nakanishi et al. [65] using the same method. Alternatively, Frye et al. [66] prepared films of FeCo-based PBAs with different alkali metals by using a sequential adsorption method. However, among all fabrication methods available at present for a thin film of PBAs, electrochemical deposition is the most efficient and widely used. This is because (i) film with mixed valence transition metals can be prepared, which provides a way of controlling the structure of the film by means of the electrode potential and thereby potentially raising TC, and (ii) the oxidation states of the metal ions in films can be electrochemically controlled by making use of their zeolitic properties, allowing modification of their magnetic properties [67]. Therefore, in the present study, thin films of PBAs, KjFeIIk[CrIII(CN)6]l
mH2O, have been prepared electrochemically with variation in (i) film thickness and (ii) composition, and their structural and magnetic properties have been investigated in detail. The ability of enhancing the Curie temperature just by changing electrode voltage has been demonstrated.

128

Functional Materials

S1

(B)

(111)

80 60 40 20 0.0

S0.7

0.3 0.6 0.9 K+ ion content

S0.2

420

400

220

200

111

S1

Intensity (a.u.)

(A)

Intensity (a.u.)

Thin films of KjFeIIk[CrIII(CN)6]l
mH2O were deposited on indium tin oxide (ITO)-coated glass substrates at room temperature. Fresh aqueous solutions of 0.10 M K3[Cr(CN)6] and 0.15 M FeCl3 (10 mL each) were mixed and then electrochemically reduced to prepare KjFeIIk[CrIII(CN)6]l
mH2O thin films. We prepared two set of samples for the present study. In the first set of samples, the electrochemical reduction voltage of 20.5 V was kept constant, whereas deposition time was varied (10, 30 and 60 min), resulting in insoluble polynuclear metal cyanide films of orange colour with thicknesses of B1, 3, and 5 μm, respectively. In the second set of samples, the electrode voltage was varied from 20.5 to 20.9 V with respect to the reference electrode, whereas deposition time (25 min) was kept constant. Hereafter, the thin-film samples deposited at electrode voltages of 20.5, 20.6, 20.8 and 20.9 V, are referred to as S0, S0.2, S0.7 and S1, respectively. Atomic force microscopy (AFM) images and XRD patterns of the second set of thin-film samples (with varying stoichiometry) are shown in Figure 3.11. The XRD patterns for these films are also consistent with the face-centred cubic (fcc) structure with an Fm3m space group similar to that for the bulk powder sample. ˚ for sample S0 However, the lattice constant increases successively from 10.55 A ˚ ˚ to 10.66 A for sample S1. The crystallite size of B750 A was found to be same for all the samples with various compositions. The changes in intensities of Bragg peaks are also observed with varying compositions. The integrated intensity of the Bragg peak (111) as a function of K1 ion content is plotted in the inset of Figure 3.11B. It is clear that the Bragg peak (111) for the film prepared with higher voltage (20.9 V) is intense in comparison with that for the film prepared with a

S0.7

S0.2

S0 Experimental data Fitted Difference Bulk sample

S0

10

20

30

40 2θ (°)

50

60

Figure 3.11 (A) AFM images and (B) room temperature XRD patterns of thin-film samples of KjFeIIk[CrIII(CN)6]l
mH2O with various compositions. An XRD pattern of bulk sample is also shown for comparison. Rietveld refinement of the XRD pattern of the bulk sample is done by using the FULLPROF program. The inset shows the variation of integrated intensity of the strongest Bragg peak (111) with K1 ion content [70].

Functional Magnetic Materials: Fundamental and Technological Aspects

129

lower voltage (20.5 V). This could be due to a preferred orientation of the film in (111) direction. However, the relative intensities of the diffraction peaks change depending on the concentration of alkali cations [68,69]. The presence of a higher amount of alkali cations (K1), with increasing electrode voltage, is evident from the elemental analysis. The sample morphology (Figure 3.11A) shows a similar kind of growth behaviour for samples with varying thicknesses, confirming that all films are uniformly deposited on the substrates. Figure 3.12 shows the normalized FC magnetization versus temperature curves measured in 500 Oe field for all thin-film samples. The paramagnetic- to ferromagnetic-like transition has been observed for all the samples. For the films with varying thickness, a small increase in the transition temperature has been observed with increasing film thickness (Figure 3.12A). The value of TC for each film is derived from the corresponding curve of dM/dT versus T. The variation of TC as a function of thickness is depicted in the inset of Figure 3.12A. For the films with thicknesses of 1, 3 and 5 μm, TC values of B11, 13 and 21 K have been observed. It is seen from the XRD and AFM study that for the film with thickness of 1 μm, the crystallites are small and, with increasing film thickness, the size of crystallites increases, which results in an increase in TC. For thin films with varying composition, a drastic increase in the TC, compared with films with thickness variation, has been observed with increasing K1 ion.

1.0

70 Tc (K)

(B)

MT /M5K

0.8

10

0.6

0.0

0.3 0.6 0.9 + K ion content

0.4 0.2

S0

0.0 1.0

S0.2

S0.7

(A)

0.6

S1

21 Tc (K)

0.8 MT /M5K

50 30

18 15 12 1

0.4

2 3 4 Thickness (μm)

5 5 μm 3 μm 1 μm

0.2 0.0 5

25

45 Temperature (K)

65

Figure 3.12 Normalized field-cooled magnetization versus temperature data in 500 Oe field for (A) Fe1.5[Cr(CN)6]
7.5H2O thin-film samples with varying film thickness and (B) KjFeIIk[CrIII(CN)6]l
mH2O thin films with composition variation. The inset shows a variation in TC with thickness and composition for the corresponding films [70].

130

Functional Materials

The inset of Figure 3.12B shows variation in TC as a function of K1 ion content in the film. For sample S0, a TC of B11 K was derived from the minima of the dM/dT versus T curve. However, for the thin film with K1 ion content 5 1, a TC of B65 K is observed. This is the first experimental result where an enhancement in TC for PBAs has been observed just by a using suitable electrode voltage in the electrochemical deposition process. The enhancement in TC is mainly attributed to the change in the FeII/CrIII ratio because of incorporation of K1 ions with increasing electrode voltage [70]. The ability of tuning TC just by changing the electrode voltage could be useful in designing thin films of molecule-based magnets.

3.4

Magnetic Nanoparticles

Magnetic nanoparticles have attracted a lot of attention because of their various technological applications. These include high-density storage, spintronics, biological applications such as therapeutic drugs, gene and radionuclide delivery, radionuclide separation and contrast-enhancement agents for MRI. The size range of 1100 nm makes them suitable for biomedical applications as their size is less than or comparable to various biological entities such as cells, viruses, genes and proteins. Magnetic nanoparticles form the bridge between bulk systems and atomic-level structures or molecules. The magnetic properties of nanoparticles can be significantly different from their corresponding bulk materials. We discuss below the role of structural and magnetic properties of the nanoparticles in designing magnetic nanoparticle systems for their use in spintronics, high-density magnetic recording, biological applications, radionuclide separation, etc.

3.4.1

Spintronics Materials

Spintronics (or spin-based electronics) is currently an active area of research because spin-based multifunctional electronic devices have several advantages over conventional charge-based devices in data-processing speed, non-volatility, higher integration densities, etc. Spintronics uses individual magnetic moments to integrate logic function and non-volatile information storage. Giant magnetoresistance (GMR) and diluted magnetic semiconductors (DMS) are such two examples of spintronics. DMS or semiconductors with dilute concentration of uniformly distributed magnetic dopants are promising materials for spin-based devices. In DMS, apart from an electron’s charge degree of freedom, one uses the spin degree of freedom, which can lead to new classes of devices and circuits. Semiconductors (p-type or n-type) with high Curie temperature (more than room temperature) and high magnetic moment are required for practical applications. Searching for new DMS materials with intrinsic room temperature ferromagnetism as well as understanding the origin of the ferromagnetism is an active area of research in the field of spintronics.

Functional Magnetic Materials: Fundamental and Technological Aspects

131

Transition-metal-doped semiconducting materials such as GaN and ZnO with 5% Mn and large hole concentration are predicted by Dietl et al. [71] to show room temperature ferromagnetism. However, in the earlier work of Sato and Katayama-Yoshida, [72] from ab initio calculations based on local density approximation, it was predicted that incorporating 525% V, Cr, Fe, Co or Ni in ZnO can give ferromagnetic behaviour without any additional hole doping. After this theoretical prediction, there was extensive study on transition-metal-doped semiconductors. There are various reports in the literature showing that doping with transition metal does not show any ferromagnetism at room temperature. However, several other reports showed the presence of ferromagnetism in the DMS materials above room temperature. So, it is still controversial whether the ferromagnetism in these systems is intrinsic to the DMS materials or not, leaving scope for probing the matter further. The magnetic properties of colloidal 5% Fe-doped ZnO nanocrystal (Figure 3.13) are presented here. For this purpose, colloidal Fe-doped ZnO nanocrystals were prepared by thermal decomposition of zinc(II) acetylacetonate and iron(II) acetylacetonate at 200 C in hexadecylamine, which acts as a solvent as well as a capping agent. The sample was prepared at relatively low temperature to avoid the formation of secondary ferromagnetic phases. No annealing of the samples was done. The XRD pattern confirmed the exclusive formation of the host ZnO wurtzite structure and did not show any reflections attributable to any of the iron oxides (namely Fe2O3, Fe3O4 and FeO) or ZnFeO (namely ZnFe2O4 and ZnxFe32xO4) compounds [73]. The oxidation state of iron in the nanocrystals has been determined from X-ray photoelectron spectroscopy. Figure 3.14 shows the spectra of Zn 2p and the Fe 2p regions. The peaks shown in the figure are in agreement with the binding energies of Zn 2p3/2 and Zn 2p1/2. The Fe 2p spectrum shows a very weak and broad peak centred around 709.6 eV. The weak signal may have arisen owing to low Figure 3.13 TEM pattern of Zn0.95Fe0.05O nanocrystals [72].

132

Functional Materials

(A) Zn, 2p

(B) Fe, 2p

1021.1

Intensity (a.u.)

Intensity (a.u.)

1043.4

1010

1020

1030 1040 BE (ev)

1050

709.7

705

710

715

720

BE (ev)

Figure 3.14 XPS spectra for Zn0.95Fe0.05O nanocrystals [72].

concentration of Fe in the present sample. However, owing to a very weak and broad signal, it is not possible to assign the oxidation state of Fe unambiguously. Figure 3.15A depicts the zero-field-cooled (ZFC) and FC magnetization as a function of temperature over 55315 K under the applied magnetic field of 500 Oe. From the observed magnetization it is very clear that the sample is not paramagnetic at room temperature. In fact the ordering temperature is higher than the highest measurement temperature of 315 K. The ferromagnetism at room temperature is further confirmed from the hysteresis loop measurement as shown in Figure 3.15B. A well-defined hysteresis loop at room temperature with a coercive field of 58 Oe and a remanent magnetization of 0.006 emu g21 has been observed. The corresponding values at 150 K are 60 Oe and 0.016 emu g21 and at 5 K are 110 Oe and 0.020 emu g21. From the hysteresis plot at 5 K, the maximum moment (5 T applied field) per Fe ion is found to be B0.894 μB. The corresponding values at 150 and 300 K are 0.093 and 0.0472 μB, respectively. It is evident that the observed maximum magnetization is significantly lower than the theoretically expected spin-only value of saturation magnetization for a ferromagnetically ordered system. These observations are somewhat similar to those of other ferromagnetic ZnO-based DMS.

3.4.2 Nanoparticles for High-Density Magnetic Recording The data storage density of hard disks is growing very rapidly. The demand for higher storage density in information storage media constitutes an important motivation for the fabrication of smaller and smaller nanoparticles. The main objective of advanced research in this area is to reduce the size of the magnetic particles and increase the magnetic stability (magnetic anisotropy). There has been an unabated tendency for a long time to prepare nanoparticles of various sizes and shapes. The geometry of a nanomagnet has a great impact on its magnetic properties resulting from the interplay among different types of magnetic energy. The elongated magnetic nanoparticles

Functional Magnetic Materials: Fundamental and Technological Aspects

Figure 3.15 (A) Temperature dependence of ZFC and FC magnetization for Zn0.95Fe0.05O under the applied magnetic field of 500 Oe. M versus H curves at (B) 5 K and (C) 150 and 300 K. Insets expand the low-field region [72].

FC

0.04

ZFC M (emu g–1)

133

0.02

500 Oe (A) 0.00 60

120

180

240

300

Temperature (K) 3.0

(B) 5K

1.5 0.0 M (emu g–1)

0.03

–1.5

0.00 –0.03

–3.0

–0.04 –0.02 0.00 0.02 0.04

(C)

150 K

0.25

300 K

0.00 0.03 0.00

–0.25

–0.03 –0.04 –0.02 0.00 0.02 0.04

–4

–2

0

2

4

H (T)

are of particular interest for data storage applications because they have two stable magnetization states. The ability to prepare air-stable ferromagnetic bimetallic particles is therefore a challenge. We have studied some magnetic nanoparticle systems to understand their magnetic properties at nanolength scale. Here we give NiPt nanoalloy and γ-Fe2O3 nanoparticles, as examples. A room-temperature wet chemical method of synthesis for rod-shaped magnetic ‘solid solution’-type NiPt nanoparticles in cetyltrimethylammonium bromide (CTAB) surfactant has been used [74]. At higher concentrations, it acts as a template in such a fashion that the number of rod-shaped particles becomes more prevalent. The advent of new methods for preparing highly soluble and processable colloidal metallic, semiconductor and magnetic nanocrystals with very tight control over size and shape creates significant opportunities in the study of inorganic liquid

134

Functional Materials

Figure 3.16 TEM of NiPt bimetallic nanoparticles. Preparation condition [NiSO4] 5 0.014 M, [H2PtCl6] 5 4.8 3 1024 M, [KOH] 5 0.36 M and [CTAB] 5 0.25 M [73].

crystals. The transmission electron microscopy (TEM) image (Figure 3.16) shows a large proportion of nanorods of length B20 nm (aspect ratio B7) along with spherical particles having an average diameter B10 nm. The temperature-dependent magnetization of NiPt nanoparticles was studied by a superconducting quantum interference device (SQUID) between 5 and 300 K following the ZFC and FC procedures in an applied field of 500 Oe. This study indicates that particles at room temperature are not ferromagnetic, but rather show superparamagnetic behaviour. Here it is observed that the ‘blocking temperature’ (Tb) for the samples increases (from 65 to 150 K) with a decrease in CTAB concentrations (0.251024 M). It is well known that Tb of different nanoparticles varies with the variation of particle size and shape. Two representative ZFC and FC curves for two different samples keeping the Ni/Pt proportion fixed but varying the CTAB concentrations are shown in Figure 3.17. Here the variation in blocking temperature is due to size and shape effects of the particles. As the CTAB concentration decreases, the particle size increases, hence blocking temperature also increases. Again Tb (65 K) for the alloy having maximum rod-shaped particles was compared with that of pure Ni(0) of comparable size. This difference in Tb is due to the alloying of the nanorods. In this case, change in blocking temperature of the alloyed particles is due to change in magnetic anisotropy from pure to alloy state. The particle-size-dependent magnetic properties of γ-Fe2O3 nanoparticles have been studied [75,76]. These nanoparticles were prepared using a reverse micelle technique. Figure 3.18 shows the TEM images and the corresponding number distribution of particle size (diameter) graphs for the nanocrystalline samples with average diameters B8.6 and 10.2 nm, respectively. Figure 3.19 illustrates the ZFC and FC magnetization curves, recorded under an applied magnetic field of 100 Oe for the samples annealed at 200 C, 250 C and 350 C. The ZFC and FC curves show irreversibility, which is typical of the blocking process for an assembly of superparamagnetic nanoparticles. The values of Tirr (the temperature below which irreversibility in ZFC and FC magnetization occurs) are found to be 150, 175 and B315 K for the samples annealed at 200 C, 250 C and 350 C, respectively. With increasing applied field, Tirr decreases for all samples. A maximum field of 10 kOe is found to be enough to suppress the bifurcation between ZFC and FC magnetization curves for all the three samples. Peaks in the

Functional Magnetic Materials: Fundamental and Technological Aspects

135

Figure 3.17 FC and ZFC curves for rod-shaped and spherical NiPt magnetic particles of sets a and b. Condition: for both the sets [NiSO4] 5 1.4 3 1022 M, [KOH] 5 0.36 M, [H2PtCl6] 5 4.8 3 1024 M; but [CTAB] 5 0.25 M for (A) and [CTAB] 5 1024 M for (B) [73].

ZFC magnetization curve occur at a temperature Tpeak B147, 166 and 303 K (below Tirr) for the samples annealed at 200 C, 250 C and 350 C, respectively. With increasing particle size, Tpeak shifts towards higher temperatures. However, for any given sample, Tpeak is found to shift towards lower temperatures with increasing applied field. The observed field dependence of Tpeak and a weak temperature dependence of FC magnetization indicate a cluster spin-glass-like behaviour, with strong interparticle interactions. From the observed magnetic behaviour, it is, therefore, evident that such nanoparticles are useful in information storage media.

3.4.3

Possible Application in Radionuclide Separation

Various methods are used for the separation of radionuclides from non-radionuclides in a mixture of nuclear waste. One of the most common techniques is solvent

136

Functional Materials

10 nm

20 nm

No. of particles

100 75 50 25 0 2 (A)

4 6 Particle radius (nm)

8 2 (B)

6

10 14 Particle radius (nm)

Figure 3.18 TEM images and the corresponding number distribution of particle size (diameter) graphs for the (A) 200 C and (B) 250 C annealed γ-Fe2O3 samples [75].

Figure 3.19 ZFC (solid symbols) and FC (open symbols) magnetization curves of γ-Fe2O3 samples measured under 100 Oe in the heating cycles [74].

extraction. Extractants such as crown ether find extensive applications in the removal of radionuclides from liquid wastes. In the column-based separation method, resins loaded with extractants reduce the quantity of the required volume of the extractants. However, the magnetic separation method can be more efficient than column-based

Functional Magnetic Materials: Fundamental and Technological Aspects

137

separation. For magnetic separation, the superparamagnetic nanoparticles can be coated with suitable solvent extractant. It is quite important to understand the magnetic properties of such superparamagnetic nanoparticles before and after loading solvent extractant. Here we present the magnetic properties of the crown-ether-loaded magnetitebased polymer beads [77]. The superparamagnetic beads can be used in magnetically assisted chemical separation of potassium and strontium from aqueous medium. A sample of polyester polystyrene polyvinyl alcohol blended magnetite (magnetic beads) was prepared by using a chemical method. Dibenzo-18-crown-6 (1 mM) was dissolved in a mixture of tributyl phosphate and toluene (1:1.8 v/v) to obtain solvent extractant for potassium. A 2 mL mixture of extractant, acetone and methanol (1:2:1 v/v) was added to 1 g magnetic beads and allowed to stand for 24 h. Air-dried samples were used for extraction studies. Figure 3.20A and B shows the bright-field TEM image of the base magnetite particles and corresponding particle size histogram, respectively. Figure 3.20C shows the bright-field TEM image of the polymer-coated magnetite beads indicating a non-uniform agglomeration. The clustered nature of the particles suggests a random precipitation of the magnetic particles within a polymer having random coil structure. Polymer coating on magnetite nanoparticles is evident from the observed diffused nature of the diffraction pattern (not shown here). The average particle size of the base magnetite nanparticles has been found from TEM studies to be B19 nm. The field (H) dependence of DC magnetization (M) at 300 and 5 K is shown in Figure 3.21A. From data of M versus H, the average size of the magnetic particles was calculated to be about 16 nm. Figure 3.21B depicts the temperature dependence of ZFC and FC magnetization under the applied field of 100 Oe. A thermomagnetic irreversibility is observed below about 120 K. A broad peak in the ZFC magnetization is found at about 56 K. Such behaviour is usually observed in superparamagnetic systems consisting of single-domain particles. These room-temperature superparamagnetic particles have filter regeneration capability and with suitable extractants may form potential systems for radionuclide removal from aqueous wastes [77].

3.4.4

Scope in Biomedical Science

Magnetic nanoparticles have various biological applications: therapeutic drugs, gene and radionuclide delivery, radionuclide separation and contrast-enhancement agents for MRI to name a few. A contrast enhancement using magnetic nanoparticles is of particular importance as it is possible to detect tumours that contain the magnetic nanoparticles at the cellular level. Among these applications, the use of magnetic nanoparticles in MRI as contrast enhancement media has recently drawn considerable interest. This means that if the nanoparticles can be selectively inserted into tumour cells, it is possible to detect tumours at their initial growth stage. Application of nanoparticles in MRI is still at its infancy. Problems of particle agglomeration, uncertainty in selective uptake of

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Functional Materials

Figure 3.20 (A) TEM image of the base magnetite nanoparticles, (B) corresponding particle size histogram and (C) TEM image of the magnetic polymer beads [76].

(A)

100 nm

(B)

Percentage

30

20

10

0 10

15

20 25 Size (nm)

30

30

(C)

500 nm

particles by target cells and low efficiency of particle internalization by cells after the particles are attached onto their surfaces have yet to be solved. For magnetic nanoparticles used in the human body, physicians need particles that are free from aggregation and highly stable. Another critical property is the size. Magnetic nanoparticles for in vivo biomedical use must be small enough to avoid detection by the immune system or to be circulated through the bloodstream and should be stable enough to remain in the body for a sufficient time.

Functional Magnetic Materials: Fundamental and Technological Aspects

1.6 5K

6

300 K

3 0 –3 –6

Magnetization (emu g–1)

Magnetization (emu g–1)

9

–9 –60 –40 –20 (A)

139

0

20

40

Magnetic field (kOe)

FC ZFC

1.4 1.2 1.0 0.8 0.6 0

60 (B)

100 200 Temperature (K)

300

Figure 3.21 (A) M versus H recorded at 300 and 5 K and (B) FC and ZFC M versus T curves at 100 Oe for the magnetic polymer beads [76].

The magnetic nanoparticles need to be biocompatible for their use in biological applications. There are various techniques to achieve this; coating the magnetic nanoparticles with some biocompatible entity is one of the simplest methods. The magnetic properties of iron oxide nanoparticles coated with gold and silver [78] have been presented here. The Fe3O4 nanoparticles with size B13 nm were prepared in aqueous micellar medium at about 80 C. To make Fe3O4 nanoparticles resistant to surface poisoning, a new route is developed for coating Fe3O4 nanoparticles with noble metals such as gold or silver as a shell. The shell thickness of the coreshell particles becomes tunable through the adjustment of the ratio of the constituents. Thus, the route yields well-defined coreshell structures of size 1830 nm with varying proportions of Fe3O4 to the noble metal precursor salts [78]. Figure 3.22A shows the XRD patterns for the magnetic nanoparticles of pure Fe3O4 (M1) and the particles with gold coating (M2, M3, M4 with Fe3O4:Au of 1:0.1, 1:0.5, 1:1, respectively). The experimentally obtained patterns were identified through comparison with standard Fe3O4 and Au patterns. For thin coating of the Au metal, the peaks corresponding to both Fe3O4 and Au patterns appeared in the XRD pattern of the M2 sample. For the M3 and M4 samples, XRD patterns showed peaks mainly due to Au as X-rays were hardly able to penetrate the Au shell around the Fe3O4 nanoparticle core. TEM images for the corresponding samples are shown in Figure 3.22B. From TEM images, it is evident that the particles are well separated and almost monodisperse. The size of pure Fe3O4 particles is 1360.5 nm, whereas after gold coating, the particle size increases to 1861, 2560.5 and 3060.2 nm for the M2, M3 and M4 samples, respectively. Figure 3.23A shows the FC and ZFC magnetization curves for the M1, M2, M3 and M4 samples. With increasing temperature, the ZFC curve comes to a maximum, which corresponds to the case when most of the particles behave superparamagnetically. This temperature is the blocking temperature (Tb). At temperatures

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Functional Materials

Intensity

(A) 1.2 0.8 0.4 0.0

111 200

M4 220 311222

111 200

0.4 0.2 0.0

M3 220 311 222

0.3 0.2 0.1 0.0

M2

0.3 0.2 0.1 0.0

311 111 220 0

20

440 400 511 533

40

60 2θ

80

M1 731 100

120

(B) 50 nm

100 nm

M2

M1 M3

50 nm

M4

50 nm

Figure 3.22 (A) XRD and (B) TEM of the bare (M1) and Au-coated (M2M4) Fe3O4 nanoparticle samples [77].

higher than Tb, the magnetization decreases. In the case of FC magnetization, both FC and ZFC curves coincide until they come to Tb (equilibrium magnetization). Below Tb, the FC curve splits from the ZFC curve because the ZFC does not correspond to the equilibrium condition. The values of Tb for M1, M2, M3 and M4 are 150, 143, 138 and 135 K, respectively. Hence Tb decreases with the increase of the thickness of the gold shell because of reduced dipoledipole interaction between highly separated magnetic nanoparticles. The plots of magnetization versus magnetic field at room temperature and 10 K for the typical magnetic nanoparticles without (M1) and with (M3) a coating of gold are shown in Figure 3.23B. At room

Functional Magnetic Materials: Fundamental and Technological Aspects

141

10

12

FC

Magnetization (emu g–1)

Magnetization (emu g–1)

(A)

ZFC

10

M1 H = 200 Oe

8 6 4 0

50

100 150 200 250 Temperature (K)

FC ZFC

8 6

M2 H = 200 Oe

4 2

300

0

50

100 150 200 Temperature (K)

250 300

Magnetization (emu g–1)

Magnetization (emu g–1)

3.0 FC

8

ZFC

6 M3 4

H = 200 Oe

6 0 0

50

100 150 200 Temperature (K)

250

FC

2.5

ZFC

2.0 1.5

M4

1.0

H = 200 Oe

0.5 0.0 0

300

50

100

150

200

250

300

Temperature (K)

40

M1

0

–40

10 K 300 K –5000 0 5000 Magnetic field (Oe)

Magnetization (emu g–1)

Magnetization (emu g–1)

(B) 40

M3

0 10 K 300 K –40 –5000 0 5000 Magnetic field (Oe)

Figure 3.23 (A) ZFC and FC magnetization plots of samples M1, M2, M3, M4. (B) Hysteresis curves for two samples, M1 and M3, at 10 K and room temperature [77].

temperature no hysteresis loop was obtained for the samples, but at 10 K a clear hysteresis revealed the resultant magnetic nanoparticles to be superparamagnetic in nature, and the particles were so small that they may be considered to have a single magnetic domain. Saturation magnetization (Ms) of bulk Fe3O4 is 92 emu g21, but here for these Fe3O4 nanoparticles it is 38 emu g21 at room temperature. So it is reduced by 57.6% from the bulk. These stable particles with strong magnetic

142

Functional Materials

properties might find applications in MRI techniques and DNA sensing methods after selective surface functionalization. Similarly, our DC magnetization study on γ-Fe2O3 nanoparticles [76] reveals that these nanoparticles are superparamagnetic at room temperature with a particle moment of B28,000 μB without any magnetic interaction between the nanoparticles. The observed high value of magnetization makes the studied material more attractive for possible practical applications in drug delivery. Recently we have built a prototype magnetic-nanoparticle-loaded membrane device using such nanoparticles for use in support of artificial heart pumps [79].

3.5

CMR Manganites

The phenomenon of huge change (B100%) in the electrical resistance of a material under the application of a magnetic field is called the CMR effect. The basic understanding of the CMR effect offers tremendous opportunities for the development of new technologies such as data-storage devices with increased data density and reduced power requirements. The oxide compounds of ABO3 (Figure 3.24) type (A: rare earths, Ca, Sr, Ba, etc. and B: transition metal ions) exhibit a variety of structural, electronic and magnetic properties. The rare-earth ion site plays an important role in stabilizing the crystal structure, whereas the valence state of the transitionmetal ions decides the magnetic and the transport properties in these systems. Among these compounds, the manganites (AMnO3) have been of great interest because of the CMR effect, charge ordering and metal-insulator transition properties observed in such materials. We have studied the structural and magnetic properties of many such mixed perovskite manganites [8087]. We have controlled the double-exchange interaction as well as the underlying superexchange and Coulomb interactions among Mn ions in several manganites such as La0.67Ca0.33Mn0.9Fe0.1O3, Figure 3.24 Crystal structure of ABO3-type cubic perovskites.

O

O

O A

B

O

O A

A

Functional Magnetic Materials: Fundamental and Technological Aspects

143

La0.67Ca0.33Mn12xGaxO3, (La12xDyx)0.7Ca0.3MnO3, (Nd12xTbx)0.55Sr0.45MnO3 and Sm0.52Sr0.48MnO3 [8087]. Our study essentially brings out the tunability as well as the tailoring aspects of the CMR effect in these manganites.

3.5.1

Study of Ionic Size Effect in Dy-Substituted La0.7Ca0.3MnO3 CMR Perovskite

For the (La12xDyx)0.7Ca0.3MnO3 system, by tuning the size mismatch of the A site (La/Dy/Ca site) ions, we have successfully controlled the structural parameters such as MnaOaMn angles and MnaO bond lengths [83,8890]. In our study, the starting parent compound La0.7Ca0.3MnO3 shows a paramagnetic to ferromagnetic phase transition at B250 K. The substitution of La31 by Dy31 in La0.7Ca0.3MnO3 ˚ and 0.91 A ˚ , respectively, keeps the ratio of Mn31 /Mn41 with ionic radii of 1.06 A unaltered from its original value of 7/3 as in the La0.7Ca0.3MnO3 parent compound. We present the effect on the magnetic properties due to substitution of La31 by Dy31 in La0.7Ca0.3MnO3 using AC susceptibility, DC magnetization and low-temperature neutron diffraction techniques. The results of the neutron diffraction structural study have been used to discuss the microscopic origin of the observed changes in magnetic properties. Figure 3.25 shows the temperature dependence of the real part of AC susceptibility χAC for both La0.7Ca0.3MnO3 and (La0.757Dy0.243)0.7Ca0.3MnO3. A paramagnetic to ferromagnetic type transition at 235 K for the x 5 0 sample is found. With the Dy substitution, the magnetic ordering temperature decreases and the peak in the χAC curve is found to occur at 60 K for the x 5 0.243 sample. The behaviour of the χAC curve for the x 5 0.243 sample is very different from that of the ferromagnetic x 5 0 sample. The peak in the χAC curve found at 60 K for the Dy-substituted sample is mimic to a canted spin system or spin-glass-type system. Figure 3.26 shows the

AC susceptibility (a.u.)

1.00

0.75 x=0 x = 0.243

0.50

0.25

0.00 0

60

120

180

240

300

T (K)

Figure 3.25 Real part of AC susceptibility as a function of temperature for (La1.02xDyx)0.7Ca0.3MnO3 (x 5 0 and 0.243) perovskites [82].

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Functional Materials

2450

x=0

Paramagnetic

1550

297 K

650

Neutron counts (a.u.)

–250 2000 1250

Ferromagnetic

x=0

3.4 (6) μB

15 K

500 –250 1200

x = 0.243

800

Paramagnetic 297 K

400 Counts (arb. units)

0 1400 1000 600

x = 0.243

500 400 300

50

100 T (K)

150

0.0 (6) μB

200

15 K

200 –200 11

17

23

29

35

41

2θ (°)

˚ ) for Figure 3.26 Observed (open circles) neutron diffraction patterns (λ51.094 A (La12xDyx)0.7Ca0.3MnO3 of the x50 and 0.243 samples at 297 and 15 K recorded over the lower angular range where magnetic Bragg reflection intensities, if any, are expected predominantly. The solid lines represent the Rietveld refined patterns. The difference between the observed and calculated patterns is also shown at the bottom of each curve by solid lines. The insets show the combined peak intensity of the (110) and (002) inner Bragg peaks (2θ16.6 ) of the x50.243 sample as a function of temperature [82].

observed and Rietveld refined neutron powder diffraction patterns at 15 and 297 K for the x 5 0 and 0.243 samples. For both samples, the diffraction patterns at 297 K could be fitted with only nuclear intensities, confirming the paramagnetic nature of the samples at room temperature. The Rietveld refinement of data recorded at 15 K shows a ferromagnetic ordering for the x 5 0 sample with a net Mn-site ordered moment of 3.42 6 0.06 μB along the crystallographic c-direction. However, for the x 5 0.243 sample, the diffraction pattern at 15 K could be fitted with only nuclear intensities, confirming the absence of any observable long-range magnetic ordering. The absence of any magnetic contribution to the (110) and (002) nuclear Bragg peaks over 15200 K (shown in the inset of Figure 3.26) confirms the absence of a longrange magnetic order at any intermediate temperatures as well. It may be stressed that no additional Bragg peaks are found either, indicating the absence of any other ordered magnetic phase (antiferromagnetic, spiral, etc.) in the Dy-substituted sample.

Functional Magnetic Materials: Fundamental and Technological Aspects

No. of scattered neutrons

3000

145

297 K

2000 x = 0.0

1000 0 2500 297 K 1500

x = 0.243 500 –500

15

35

55

75

95

2θ (°)

Figure 3.27 Rietveld refinement of the neutron diffraction data for the La0.7Ca0.3MnO3 (top) and (La0.757Dy0.243)0.7Ca0.3MnO3 (bottom) samples measured at room temperature. Observed (open circles) and calculated (solid lines) patterns are shown above. The difference between measured and calculated data is plotted below. The vertical lines indicate the position of allowed Bragg peaks [87].

The effect of ionic size (owing to Dy substitution) on the structure of the crystalline perovskite compound has been studied by neutron diffraction experiments at room temperature. Figure 3.27 shows the Rietveld refined (using the FULLPROF program [53]) neutron powder diffraction patterns at room temperature for both La0.7Ca0.3MnO3 and (La0.757Dy0.243)0.7Ca0.3MnO3. The refinement shows an orthorhombic perovskite structure (space group: Pbnm). The stoichiometric nature of both samples is confirmed from the analysis. The refinement shows that the La/ Dy/Ca atoms are at 4c(x, y, 1/4) position, Mn at 4b(1/2, 0, 0), O(1) at 4c(x, y, 1/4) and O(2) at 8d(x, y, z) position. The lattice parameters for the x50.243 sample are ˚ , b 5 5.460(2) A ˚ and c 5 7.678(3) A ˚ compared with found to be a 5 5.423(2) A ˚ ˚ ˚ a 5 5.479(4) A, b 5 5.482(3) A and c 5 7.750(4) A for the parent compound. The ˚ 3 for ˚ 3 and 232.78 A unit cell volume decreases with the Dy substitution (227.373 A the substituted and parent perovskites, respectively). The MnaO(2) distances decrease with the Dy substitution. The values of MnaO(2) are found to be 1.96(1) ˚ 3 2, 1.98(1) A ˚ 3 2 for the x50 sample and 1.958(7) A ˚ 3 2, 1.962(7) A ˚ 3 2 for A ˚ the x50.243 sample. The MnaO(1) distance (51.960(4) A 3 2) is found to remain unchanged with the Dy substitution. For the x50.243 sample, the MnaO1aMn and MnaO2aMn bond angles are found to be 156.5(9) and 157.9(4) , respectively, compared with 162(1) and 159.9(6) , respectively, for the parent La0.7Ca0.3MnO3 compound. It is therefore evident that the distortion of the MnO6 octahedra (Figure 3.24) on going from the x50 sample to the x50.243 sample increases significantly. The geometrical characteristics indicate that the distortion of the perovskite structure is in good agreement with the size effect of ions. The observed changes

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Functional Materials

9

10

0 KOe

8

ρ (Ω cm)

% MR = 100 × (PH – Po)/PH

(A) 2000 1500 14.5 KOe 7 KOe 5 KOe 4 KOe 2 KOe

1000

500

30

PH / Po

(B)

40

1.1 1.0 1.9 1.8 1.7 1.6 1.5 1.4 1.3 1.2 0.1 0.0

10 107 106 105

14.5 KOe

104 103 40

60 80 100 120 Temperature (K)

50 60 70 Temperature (K)

80

95 K

74 K

Figure 3.28 (A) Temperature dependence of the magnetoresistance under several magnetic fields. Magnetoresistance is calculated from the ρ(H,T) curves as 100 3 [ρ(H,T)2ρ(0,T)]/ρ(H,T). The inset shows the resistivity iso-field curves under 0 and 14.5 kOe. The maximum measurable resistance limits the 90 temperature range in which ρ(0,T) can be tracked to T .35 K; consequently, magnetoresistance curves cannot be calculated below this temperature. (B) Magnetoresistance isotherms at selected temperatures [83].

40 K 45.5 K 0

15 5 10 Applied field (kOe)

20

in the magnetic properties due to substitution of Dy (ionic size effect), therefore, arise from the structural buckling of the MnO6 octahedra, which reduces the strength of the ferromagnetic double-exchange interaction in the Dy-substituted compound. Figure 3.28A shows the observed magnetoresistance dependence on the temperature under several external magnetic fields for the x 5 0.243 sample. All the ρ(H,T) curves were measured while increasing the temperature at a constant rate of 3 K min21. As displayed in the inset of Figure 3.28A, the resistivity under both 0 and 14.5 kOe exhibits insulating behaviour in whole temperature range, i.e. above and below Tg/TC. This fact indicates that, first, the cluster spin-glass state (CSG) is insulating, and second, that the applied magnetic field of 14.5 kOe is not able to induce a metallic state, although the resistivity is remarkably reduced. The magnetoresistance is strongly enhanced as the magnetic field increases, reaching 2000% at 14.5 kOe, which is a typical value for CMR effects. However, the resistance decreases monotonously as the field increases and we have not identified any discontinuous magnetoresistance change as a function of the field. This observation is noticed better in Figure 3.28B, where the isothermal magnetoresistance at selected temperatures is shown. We point out that our resistance measurement limit is 5 GΩ, which disallows measuring ρ(0,T) below 35 K.

0.015

0.004 AC Susceptibility (emu g–1 Oe–1)

AC Susceptibility (emu g–1 Oe–1)

Functional Magnetic Materials: Fundamental and Technological Aspects

0.012 0.009 0.006

3 kOe 4 kOe 6 kOe

0.003 0.002 0.001 0.000

0.003

0

100 T (K)

147

Figure 3.29 Real part of AC susceptibility versus temperature under superimposed DC fields of 0, 0.02, 0.05, 0.5, 1, 2, 3, 4, 5 and 6 kOe (from top). The inset enlarges the susceptibility curves for applied field values of 3, 4 and 6 kOe [83].

200

0.000 0

50

100 T (K)

150

200

Now we give a fundamental understanding of the observed magnetoresistance behaviour for the x 5 0.243 sample. Figure 3.29 depicts the real part of the AC susceptibility under a superimposed DC magnetic field. Under zero DC field, a cusp-like peak is observed at 45.0 K (5 Tg, the cluster spin-glass freezing temperature), indicating a cluster spin-glass state at lower temperatures. Interestingly, with increasing DC field, the peak starts rounding up without any observable shift in the peak temperature (Tg). However, at further higher fields (H $ 4 kOe), the susceptibility curve loses its peak and becomes flat at the lower temperature region. This type of flat susceptibility behaviour is typical of a ferromagnetic system. Now we present neutron diffraction results to gain further understanding of the observed magnetoresistance behaviour for the x 5 0.243 sample. Figure 3.30A shows the field dependence of the peak profile at 5 K of some fundamental Bragg peaks of the x 5 0.243 compound. At applied fields of 0 and 2 kOe, no ferromagnetic contribution to the signal is found. However, magnetic Bragg intensity appears for H $ 6 kOe, giving direct evidence for the development of a static longrange ferromagnetic order at these higher fields. The width of the magnetic Bragg peaks was found to remain unaltered and limited by the instrumental resolution, indicating a true long-range magnetic order for all fields down to 6 kOe. This result demonstrates that at 5 K an external parameter, the magnetic field, can induce a cluster-glassferromagnetic transition. The temperature dependence of the Bragg peak profiles under the 40-kOe applied field, shown in Figure 3.30B, indicates that at that large magnetic field the paramagnetic to ferromagnetic phase transition occurs at B118 K (5TC, the Curie temperature). Even though a magnetic field above 4 kOe induces ferromagnetic interactions, the overall sample volume occupied by the ferromagnetic phase does not suffice to surpass the percolation threshold for electrical conduction; therefore the x 5 0.243 compound retains its insulating behaviour up to, at least, 14.5 kOe. This is not surprising, because, as illustrated in Figure 3.30A, under 15 kOe, the magnetic Bragg intensity ascribed to the ferromagnetic phase with long-range magnetic interactions is less than half that under magnetic fields large enough to melt the CSG state completely.

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Functional Materials

Neutron counts/10 min

(B)

Bragg intensity

5K 60

80 60 40 5K

20 0

0 10 20 30 40 50 H (kOe)

45 30 15 40 kOe 60

45

Bragg intensity

Neutron counts/10 min

(A) 75

80 40 kOe

60 40

0 kOe 20 0

50 100 150 250 T (K)

30

15 112

114 Scattering angle (°)

116

Figure 3.30 (A) Magnetic field evolution of the fundamental Bragg peaks (020), (112) and (200) with very close Q-values of around ˚ 21 at 5 K, under 0, 2, 6, 10, 20, 2.3 A 30, 40 and 50 kOe applied fields (starting from bottom) for (La1 2 xDyx)0.7Ca0.3MnO3 with x 5 0.243 compound. The solid line is the fitted Bragg curve by a Gaussian function for each field. The inset shows the integrated intensity of the Bragg peaks as a function of H. (B) The observed convoluted profile of the Bragg peaks (020), (112) and (200) under 40 kOe applied field at 5, 17, 29, 41, 52, 63, 74, 96, 123, 118, 141, 164, 188, 212 and 253 K (starting from top). The inset shows the integrated intensity of the Bragg peaks as a function of temperature under 0 and 40 kOe fields [83].

We also present here the results of the small-angle neutron scattering (SANS) study on the x 5 0.243 compound. Figure 3.31A depicts the Q-dependences of the background corrected SANS signals at various applied fields. The SANS signal is found to depend strongly on temperature, indicating its magnetic origin. The Q-dependences of the SANS signals clearly show a steep deviation for H $ 6 kOe from the SANS patterns obtained at lower fields (H 5 0 and 2 kOe). Also it is noticed that the SANS patterns for H 5 0 and 2 kOe overlap each other, showing a consistent behaviour with the field-dependent Bragg peak profile where no magnetic Bragg intensity is found for H 5 0 and 2 kOe. For H 5 0 and 2 kOe, we could analyse the observed SANS curves by considering the Lorentzian expression I1/[Q2 1 κ2], which corresponds to the form factor for the short-range magnetic ordered regions. The Lorentzian-function term implies that the magnetic correlations have the OrnsteinZernike form ,M(0). M(r).B(1/r) exp(2κr) with κ as an inverse of the correlation length ξ [91]. For H 5 0, the temperature dependence of the correlation length is shown in Figure 3.31B. A finite correlation length ˚ ) is found to exist at all higher temperatures in the paramagnetic state. (ξ9 A A similar behaviour was first found for the compound La2/3Ca1/3MnO3 where it was connected to the formation of polarons [92]. With decreasing temperature, ξ starts to increase significantly at TB75 K, and a saturation behaviour with ˚ is found at T # 12 K. For H $ 6 kOe, the Q-dependence of the SANS ξ16.5 A profile has been analysed by considering the expression, I1/[Q2 1 κ2] 1 I2/Q4,

Functional Magnetic Materials: Fundamental and Technological Aspects

(A) 80

60

40

40 –1

0 kOe 2 kOe 6 kOe 10 kOe 20 kOe 30 kOe 40 kOe

ISANS (Q = 0.08 Å )

SANS intensity/60 min

5K

0 kOe 30 20 10 40 kOe 0

20

0

100 200 T (K)

300

149

Figure 3.31 (A) SANS intensity as a function of Q for various H for (La12xDyx)0.7Ca0.3MnO3 with x50.243 compound. The inset depicts the SANS intensity at ˚21 versus T for 0 and Q50.08 A 40 kOe. (B) Temperature evolution of correlation length ξ under H50 [83].

0 0.1

0.2

0.3

0.4

Q (Å ) –1

(B) 16 0 kOe

ξ (Å)

14 12 10 8 0

100

200

300

T (K)

where the second term (Porod law) is the tail part of the form factor of the spherical ferromagnetic domains and dominates the ‘lower-Q’ region of our measurements (those ferromagnetic domains are also responsible for the observed magnetic Bragg scattering, shown in Figure 3.30). The first term (Lorentzian term), originated from the short-range magnetic clusters, dominates the ‘higher-Q’ region of our measurements. Under increasing applied field, the second term enhances at the expense of the first term. This indicates that the size of the smaller spin clusters grows under an external magnetic field leading to a formation of larger ferromagnetic domains. A decrease of SANS intensity at lower temperatures (below B118 K) under 40 kOe, shown in the inset of Figure 3.31A, and the corresponding rise of Bragg intensity, shown in Figure 3.30B, support this argument. For the series discussed here, with increasing Dy concentration the tolerance factor t for the perovskite structure decreases owing to a lower average ionic size of the La site hrRi, which causes bending of MnaOaMn bonds away from 180 and, as a consequence, decreases the electron transfer between Mn ions and hence the ferromagnetic double-exchange interaction. Second, owing to the increasing La-site ionic size-disorder, i.e. the width σ of the distribution of the ions on the La-site, the electron transfer is also suppressed. These effects, plus the quenched disorder,

150

Functional Materials

possibly take the system (x 5 0.243) to a cluster-glass state (quantum Griffiths phase). In the present study we have shown for the first time that the Griffiths phenomenon could be observed by substitution on the rare-earth site of CMR manganite La0.7Ca0.3MnO3. The results obtained in the (La12xDyx)0.7Ca0.3MnO3 manganites mimic the theoretically proposed phase diagram by Dobrosavljevi´c and Miranda [93] for itinerant systems in the presence of RKKY interactions, where, with increasing quenched disorder, the system evolves from a magnetically ordered phase to a cluster-glass phase and then to a quantum disorder phase. Castro Neto et al. [94] proposed a similar possibility to explain the non-Fermi-liquid behaviour in f-electron systems, where the existence of the Griffiths phase was explained by interplay among disorder and the competing RKKY and Kondo interactions. For the present manganite system with increasing Dy concentration, the two competing interactions (the ferromagnetic double-exchange interaction and the antiferromagnetic superexchange) and the presence of bond disorder including inherent disorder (caused by the chemical substitution) enhance the Griffiths features and take the system from a ferromagnetic to a cluster spin-glass state. This cluster-glass phase in the doped manganites has been characterized by inhomogeneous magnetic states at the nanometric length scale that bear local magnetic moment and strongly interact with each other inhomogeneously over the entire sample. Under higher H, the static magnetic order develops on these rare regions and these local regions start growing with H; eventually, a long-range ferromagnetic order develops at finite temperature for H . 3 kOe. Our study indicates that when the doping level is close to the critical one, the external magnetic field can be used to fine tune the magnetically disordered state. The present example, therefore, shows that understanding the role of an external magnetic field on such manganites is extremely important because of their CMR effect under field. When field is applied, we demonstrate that the reduced spin fluctuations lead to the formation of a long-range ferromagnetic phase.

3.6

Summary and Conclusion

Magnetic properties of various functional magnetic materials, namely CMR manganites, high magnetocaloric materials, hexacyanide-based molecular materials and magnetic nanoparticles having various functionalities, are described in this chapter. We have particularly brought out the tunability aspects of the magnetic properties of these functional magnets. These materials have the potential for use in information storage and processing, spintronics, drug delivery, cooling technology, etc. The influence of an external magnetic field in controlling their functionalities, which are of practical importance, has been discussed in detail. In particular, the structural and magnetic properties of high magnetocaloric materials, namely TbCo22xFex and La0.67Ca0.33MnO3 compounds, have been presented because of their usefulness in magnetic cooling at near room temperature and in low temperature regimes. The control of magnetization as well as its polarity in (CuxMn12x)1.5[Fe(CN)6]
zH2O

Functional Magnetic Materials: Fundamental and Technological Aspects

151

hexacyanide compounds as a function of magnetic field and temperature has been studied. The possible applications of the magnetic pole reversal phenomenon in magnetoelectronic and magnetocaloric devices such as magnetic memory and magnetic cooling/heating-based constant temperature baths have been revealed. The thickness- and stoichiometry-dependent magnetic properties of electrochemically prepared crystalline thin films of PBAs KjFek[Cr(CN)6]l
mH2O have been discussed. The role of structural and magnetic properties of nanoparticles in designing magnetic nanoparticle systems for their use in spintronics, high-density magnetic recording, biological applications, radionuclide separation, etc. has been brought out through our investigation of many magnetic nanoparticle systems. The ionic size effects to tune the important CMR effect in the Dy-substituted La12xCaxMnO3 CMR manganites have been addressed.

Acknowledgements This chapter is based on several collaborative research works. The author acknowledges the valuable contributions of his collaborators. The author also acknowledges M.D. Mukadam and Amit Kumar for their help in finalizing the article.

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