Chapter 1

Introduction to Nanomaterials

 

 

1.1.        Introduction

Nanophase materials with an average grain size in the range of 1 to 50 nm have attracted research interest for more than a decade since their physical properties are quite different from that of their bulk micron-sized counterparts because of the large volume fraction of atoms that occupies the grain boundary area [1-3].  This new class of materials is used in important applications like high frequency transformers, ferrofluids, pigments in paints and ceramics, biomedical applications like drug delivery system, hyperthermia, NMR, high density magnetic recording, varistors and dye-sensitized solar cells [4-10]

            The surface area of the nanostructured materials is large as the grain sizes are small. The increase in the interfacial energy due to defects, dislocations and lattice imperfections leads to changes in various physical properties and hence one can tailor make the materials with specific properties.  Almost 50 % of the atoms reside in the grain boundary area when the grain size is reduced to less than 10 nm whereas it is only 1-3 % when the grain size is 100 nm [1, 11]. Since a large fraction of atoms is present at the grain boundaries, the nanocrystalline materials exhibit enhanced diffusivity.  The enhanced diffusion results in the formation of stable and metastable phases at low temperatures, lowering of the sintering temperature and increase in the solid solubility limits.  For example, the melting point of  2.5 nm particles of CdS is 673 K compared to the bulk value of nearly 1873 K [12] and the sintering temperature of 12 nm particles of TiO2 is 400-600 K which is lower than that of the micron-sized particles [13]. The phase transformation temperature is also reduced in the nanoparticles as in the case of ZnS where it transforms from zinc-blende structure to wurtzite  structure at 673 K itself compared to the bulk materials phase transformation temperature of 1296 K [14]. The materials with stable structure in the bulk form are found to exhibit metastable structures in the nanocrystalline state like hcp Co which transforms to e-Co [15]. The mechanical, band structure, electrical, optical and magnetic properties of the nanocrystalline materials are quite distinct from the microcrystalline materials. The increase in the hardening of the nanocrystalline materials is described by the Hall-Petch relationship [16,17]. The nanocrystalline Cu and Pd of grain sizes 3-50 nm are found to show an enhanced microhardness of an order of 2-5 higher than that of the coarse-grained materials [18]. High tensile ductility has also been reported for nanocrystalline Cu [19]. However, when the grain size is decreased below 10 nm, the average microhardness decreases due to flaws present in the grain boundaries. Thus, an inverse Hall-Petch relationship is found to occur as reported by Chokshi et al. [20]. The band gap of CdS nanoparticles of size about 2 nm is 4.5 eV compared to the bulk value of 2.4 eV [21]. The electron scattering at the grain boundaries has resulted in an increase in the resistivities of nanocrystalline materials. The electrical resistivities of Fe, Ni and Cu of average grain size 100-200 nm increases by nearly 15-55 % compared to the bulk value [22]. The emission spectrum of CdSe nanoparticles, on excitation with ultra-violet radiation (UV), can be tuned from the red end of the visible spectrum to almost the entire UV range by changing the particle size [23].

The remarkable properties of the nanocrystalline magnetic materials arise from the various magnetic phenomena occurring at the nanosize level and hence the author is interested in understanding the effect of particle/grain size on the magnetic properties of both soft and hard magnetic materials.  In view of this, he has chosen Mn-Zn ferrite, yittrium iron garnet and gadolinium iron garnet which are soft magnetic materials and SmCo7 and Nd2Fe14B/a-Fe which are hard magnetic materials for his study.  The present chapter gives an introduction to the techniques for the synthesis, the structure and magnetic properties of the nanocrystalline soft magnetic materials like ferrites and garnets and the permanent magnetic materials like Sm2Co17 and Nd2Fe14B/a-Fe that are studied and reported by the author in this thesis.

1.1.1. Effect of particle size on the magnetic properties of nanocrystalline materials

1.1.1.1. Soft magnetic materials

The magnetic materials are classified into two groups such as soft and hard magnetic materials. Nanocrystalline soft magnetic materials have low magnetocrystalline anisotropy resulting in reduced coercivity and high permeability such as FINEMET alloys of the melt-spun (Fe-Si-B-Nb-Cu) [24], NANOPERM (Fe-M-B-Cu) [25] alloys and high Curie temperature HITPERM (Fe-Co-Cu-Zr-B) alloys [26] used for high frequency applications. The coercivity, Hc is related to the average anisotropy <K> by                         

where pc is a dimensionless constant close to unity and Js is the saturation polarization [27].  The coercivity increases with grain size reduction down to the single domain size obeying the 1/D law where D is the grain size and when reduced further, it follows the sixth power of the grain size as represented in Fig. 1.1.(a). However, the stress anisotropy is found to increase the coercivity in the ball-milled FINEMET alloys [28]. The initial permeability of the ball- milled FINEMET alloy core has been improved upon by the substitution of Al due to the reduction in the magnetocrystalline anisotropy energy and the magnetostriction coefficient [29]. The permeabilities of the FINEMET and the Co- based amorphous alloys are also found to be superior to that of the otherbulk alloys as shown in Fig. 1.1.(b) [30]. The random anisotropy [31] realized in these materials results in an enhanced permeability and hence these alloys show a better frequency response.

The saturation magnetization for the nanocrystalline ferrites, in general, is found to be lower compared to their bulk value, which is attributed to surface spin effects [32].  In some cases an enhancement in the saturation magnetization is observed due to the change in cation distributions [33, 34] which depends on the crystal field stabilization energy of the cations.  Apart from grain size, the cation distribution, which depends on the synthesis condition, is found to play a major role in the observed changes in their magnetic properties [35, 36].  An enhancement in the Néel temperature and spin canting [37-39] are observed in ferrites. The coercivity for the nanocrystalline materials is small when they are in their multidomain state and will be maximum at the critical single domain size and decreases further with grain size reduction as it approaches superparamagnetism, as represented in Fig. 1.2 due to increase in the thermal energy compared to the anisotropy energy. The superparamagnetic behaviour, below a critical size, limits them for magnetic recording applications. Recently  it has been reported that the superparamagnetic limit could be overcome by increasing the exchange anisotropy [40]. However, the superparamagnetic property of ferrite nanoparticles can be utilized in treating cancers from the heat produced by superparamagnetic particles due to Néel and Brownian relaxations [9,41]. The core-shell morphology of the nanoparticles resulting in an enhanced anisotropy [42] and magnetization [43] is also studied by researchers. The surface anisotropy of the nanomagnetic materials increases with particle size reduction for some soft magnetic materials [44,45].

1.1.1.2. Hard magnetic materials

The high coercivity of the hard magnetic materials arises from the large magnetocrystalline anisotropy of the materials with a non-cubic structure [46]. FePt nanoparticles of f.c.t. structure with high coercivity, synthesized using chemical methods  have the potential for use as magnetic recording medium [47]. Apart from the structure, the mechanism for the enhanced magnetic properties at the nanoscale could be different like domain wall pinning or nucleation as in Sm-Co permanent magnetic materials or exchange coupling as in nanocomposite permanent magnetic materials. Nanocrystalline Sm-Co permanent magnetic materials with cellular microstructures are being developed for high temperature applications [48]. Nanocomposite Nd2Fe14B/a-Fe exhibits enhanced energy product due to the improved exchange coupling between the two phases when the grain size of the magnetic materials is reduced to a few nanometer [49].

 

1.2.      Synthesis of nanostructured magnetic materials

The nanophase materials can be synthesized with varying grain size using the techniques like mechanical milling [50,51], melt spinning [52], chemical vapor deposition [53], inert gas condensation [54], pulsed laser ablation [55], sputtering [56] etc. Nanoparticles can also be synthesized through coprecipitation [57], oxidation [58], reverse micelles [59] and polyol process [60] in addition to other chemical methods [61]. As some of the techniques like chemical, mechanical alloying and melt spinning are used by the author to synthesize the nanomagnetic materials, a brief description of these techniques is given in this chapter.

1.2.1.  Chemical methods

A variety of chemical methods like coprecipitation, oxidation, sol-gel, reverse micelle, polyol process etc. are available for the synthesis of nanomagnetic materials. The coprecipitation [57, 62] technique, which results in particle sizes smaller than 25 nm, is widely used in the synthesis of nanoparticles. The coprecipitation technique has been used in the synthesis of CdFe2O4 [63], ZnFe2O4 [57, 64], Ni-ZnFe2O4 [65] and MgFe2O4 [66] in addition to other spinel ferrites. The size-dependent magnetic properties of MnFe2O4 with particle size from 5 to 25 nm have been studied for the samples synthesized using the coprecipitation method [67]. The superparamagnetic behaviour of CuFe2O4 has been studied for the coprecipitated samples prepared in a polymer matrix [68]. The superparamagnetic behaviour in MgFe2O4 in the particle size ranging from 6 nm to 18 nm has been studied by Chen and Zhang [69]. The coprecipitation technique was also used in synthesizing Co-Cu powders by Chow et al. [70].

The reduction of metal salts to their corresponding metals in an alcohol medium like ethylene glycol or propylene glycol is employed in the synthesis of metals, compounds and ferrite nanoparticles [71-74]. The thermal decomposition of the carbonyl precursors was used to synthesize SmCo5 through chemical methods [75].  However, the oxidation of the rare earth compounds prevents them from being synthesized through chemical methods and there are no reports in the literature on thesynthesis of the Nd2Fe14B compound through chemical methods. Magnetic spinel ferrite nanoparticles in the size range of 10-20 nm are synthesized through reverse micelles methods [76,77].

 Sol-gel technique in which an amorphous gel is obtained from solution through hydrolysis and polycondensation reactions has been used in the preparation of nanocrystalline spinel ferrites [78,79] and  ferrite/SiO2 nanocomposites [80,81]. The superparamagnetic behaviour of Fe-SiO2 composite has been studied for the samples synthesized using the sol-gel technique [82].  Citrate precursor method is employed to synthesize rare earth iron garnets, spinel ferrites and hexaferrites [83-86]. Ultrafine barium ferrite was synthesized in the size range from 5 nm to 100 nm by thermal decomposition of a citrate precursor [87].   Hydrothermal synthesis of ferrite particles in the temperature range 293 – 423 K were employed to synthesize ultrafine ferrite particles [88, 89].

Oxidation method in which the ferrous ions are oxidised to ferric ions using air oxidation or an oxidant like KNO3 has been used in the synthesis of spinel ferrites [90,91].  The mechanism for the size-controlled synthesis of particles was described by LaMer and Dinegar [92] as shown in Fig. 1.3. The LaMer’s diagram represents the nucleation and growth of particles in three stages where the concentration of the soluble species generated insitu during a chemical reaction increases during the prenucleation period and reaches a saturation concentration and finally reaches the supersaturation concentration. The spontaneous nucleation arises at the supersaturation stage followed by the growth of particles at stage III represented in Fig. 1.3. A modified oxidation method has been utilized to synthesize ferrites with large particle sizes where a suitable concentration of ferric ion is used along with the precursor salts thereby separating the nucleation and growth stages by Chinnasamy et al. [93].

 

1.2.2.     Mechanical milling/alloying

Mechanical milling is effectively used in the reduction of bulk materials to the nanocrystalline or amorphous form and also in the synthesis of nanomaterials. Various types of mills are used in the synthesis of  nanomaterials like planetary mills, vibratory mills and attritor mills. The repeated welding and fracturing of powder particles in between the balls during milling changes the physical characteristics of the material. Hence the extent of the physical change depends on the mechanical properties of the materials especially their ductile and brittle behaviour. The energy imparted in a planetary mill is larger compared to that in vibratory mills [94]. The ball movement in a planetary ball mill is schematically shown in Fig. 1.4 [95]. During the collision of balls with the powder, the events like direct and indirect seizure and dynamic, forging and shear fractures occur [96], which alter the particle morphology. Mechanical milling has been extensively carried out for grain size reduction in the study of magnetic properties of nanocrystalline spinel ferrites [97-99] and rare earth permanent magnetic materials [100-102]. Ball milling has yielded interesting magnetic properties like changes in the cation distribution and Néel temperature in spinel ferrites [103-105]. Through ball milling not only the changes in the magnetic properties are observed, but also the structural properties are drastically altered like decomposition as in hexaferrites [106], reduction of ferrites to metals and ions [107] and phase changes [108].

            Extensive work has been done in the study of rare earth permanent magnetic materials using mechanical milling and alloying. Mechanical milling leads to the formation of amorphous phase [109] and nanocomposites can be synthesized using both mechanical milling and mechanical alloying [110].

1.2.3.     Melt spinning technique

The rapid solidification process through melt spinning technique is widely used to synthesize amorphous alloys. They are widely used in the synthesis of metglas, FINEMET and rare earth alloys [111-113]. In melt spinning technique the molten alloy is allowed to pass through a narrow orifice on to a rotating copper wheel which quenches the molten alloy at a cooling rate greater than 104 K/s. The resulting samples will be in the form of ribbons, typically a few microns thick. The schematic diagram representing the melt spinning technique is shown in Fig. 1.5. The final microstructure and phase of the ribbon sample is influenced by the melting temperature, ejecting pressure, wheel speed, nozzle diameter, distance between the wheel to orifice, conductivity of the wheel and the nature of the sample [114]. Since a lot of parameters are involved in the synthesis of nanomaterials through melt spinning technique, the material properties naturally vary with the processing condition. Melt spinning technique is efficiently used in the synthesis of rare earth permanent magnetic materials containing boron due to their glass forming nature, especially Nd2Fe14B compounds [115]. Depending on the composition of the material chosen, amorphous or crystallized ribbon is obtained. Nd2Fe14B permanent magnetic material with varying grain sizes were synthesized using melt spinning technique by altering the wheel speed. The influence of the cooling rate on the magnetic properties of Nd-Fe-B materials has been studied and the critical cooling rate for the amorphous phase formation is determined as 5.4 × 105 to 7.1 × 105 K/s [116].  A series of Nd9Fe91-xBx (x = 0-9) ribbons prepared using the melt spinning technique are found to show a metastable TbCu7 phase for x Ł 7 [117]. The cooling rate for the formation of Fe3B/Nd2Fe14B composite has been estimated to be 2x105 K/s above which the ribbons become amorphous [118]. The grain sizes in the range 10 to 500 nm can be obtained through melt spinning technique. The magnetic properties are found to be dependent on the processing conditions of the melt-spun ribbons [119].

 

1.3.      Characterization techniques for the nanophase materials

Characterization of nanocrystalline materials is required in order to understand the correlation between the properties and the particle size. Various characterization techniques are employed to study the structural and physical properties of nanomaterials. Some of the characteristics of the nanostructured materials like line broadening in the X-ray peak are used to estimate the average grain size using the X-ray diffraction [120]. The surface morphology and shape of the nanoparticles are examined using electron microscopy [121]. Other techniques used are atomic force microscopy [122], scanning tunneling microscopy [123] Raman spectroscopy, Rutherford backscattering [124,125], extended X-ray absorption fine structure (EXAFS), X-ray photoelectron spectroscopy (XPS), X-ray magnetic circular dichroism (XMCD) [126-128], positron annihilation [129] and muon spin rotation [130]. The relaxation effects occurring in the nanoparticles on size reduction are examined using Mössbauer spectroscopy [131] and the changes in the magnetic properties of nanocrystalline materials are studied using vibrating sample magnetometer (VSM) and superconducting quantum interference device (SQUID). They are also used in the determination of the blocking temperature of the small particles [132,133]. The sophisticated technique like atom probe field ion microscopy (APFIM) is also employed to understand the correlation of the microstructure to the properties [134]. Some of the techniques used by the author for characterizing the nanomaterials chosen for the present work are described in Chapter 2 of this thesis.

 

1.4.      Structure and magnetic properties of nanocrystalline soft and hard magnetic  
      materials

1.4.           1.         Spinel ferrites


Spinel ferrites have the f.c.c. structure with the molecular formula of the type AB2O4 where A and B are cations. The oxygen ions form the f.c.c. structure and the cations occupy the interstitial positions. There are two interstitial sites, one being the tetrahedral or A- site surrounded by four oxygen ions and the other, octahedral or B- site surrounded by six oxygen ions. Figure 1.6 shows two of the eight octants of the cubic unit cell structure of the spinel ferrite showing the tetrahedral, octahedral and oxygen sites. There are 64 tetrahedral and 32 octahedral sites in the unit cell of which only 8 tetrahedral and 16 octahedral sites are occupied by the metal ions. According to the site occupancy of the metal ions, the spinel ferrites are classified as (a) normal spinel; where the tetrahedral sites are occupied by divalent metal ions, (b) inverse spinel; where all the divalent ions are present in the octahedral site and (c) mixed spinel; where divalent ions are present both in tetrahedral and octahedral sites.

The exchange interaction in spinel ferrites where the antiparallel alignment of magnetic moments of A- site with B- site is mediated by oxygen ions is called superexchange interaction. The strength of the antiparallel coupling between the metal ions depends on the A-O-B bond angle with the strength being the maximum for an angle of 180o [135]. The superexchange interaction is quite significant in ferrites due to the presence of metal ions in the A- and B- sites. Interesting magnetic properties such as magnetization, Curie temperature etc. of various ferrites have been observed to depend on the superexchange interaction strength, which is determined by the site occupancy of metal ions in the A- and B- sites.

The size-dependent magnetic properties of ferrites have been investigated by many researchers for particle sizes less than 25 nm where the saturation magnetization decreases with particle size reduction due to surface spin effects as explained by Kodama and Berkowitz [32, 136]. The model proposed by Kodama et al. consists of ferrimagnetically aligned core spins and a spin-glass like surface layer as shown in Fig. 1.7. For the MnFe2O4 prepared, using the coprecipitation technique, a 50 % reduction in the saturation magnetization for the 7.5 nm particles has been observed compared to the bulk value [67]. Various nanocrystalline mixed spinel ferrites synthesized using the reverse micelle technique yielded very low saturation magnetization of the order of 5 to 25 A m2/kg and the low values are attributed to the core-shell structure with a spin-glass like surface layer [137,138].

Extensive work by Šepelák et al. have shown that the ball-milled samples exhibit spin canting and cation redistributions [34,37,98,99,139,140]. Chinnasamy et al. [104, 141] have shown that ZnFe2O4 and CdFe2O4, which are antiferromagnetic in the bulk state, exhibit ferrimagnetic ordering when the grain size is reduced to nanometer level using high-energy ball milling and this is attributed to the cation redistribution. The ferrite nanoparticles are found to exhibit an enhanced Curie temperature compared to the bulk value [105, 142]. Rath et al. [143-145] have observed an enhanced Curie temperature in Mn-Zn ferrite nanoparticles prepared using chemical methods, which is attributed to a metastable cation distribution. The observed variation in the magnetic properties in nanocrystalline ferrites is related to the cation distribution among the A- and B- sites which depends on the synthesis conditions [35, 146, 147]. A good review, covering the recent literature, discusses the various aspects of the changes in the magnetic properties of nanostructured spinel ferrites [36].

            Smaller ferrite particles are found to exhibit superparamagnetic behaviour with their coercivity approaching zero. The critical size for superparamagnetism has also been calculated for various ferrites with the sizes being 14 nm, 25 nm and 50 nm for the CoFe2O4, Fe3O4 and MnFe2O4 respectively [148]. The blocking temperature for the superparamagnetism of the nanoparticles depends on their magnetocrystalline anisotropy. A comparative study of CoFe2O4 and MgFe2O4 with 20 nm particle sizes has suggested that the blocking temperature of CoFe2O4 is higher than that of MgFe2O4 by 150 K which is due to the higher magnetocrystalline anisotropy of the former [149]. The blocking temperature not only depends on the magnetocrystalline anisotropy, but also on the particle size. The blocking temperature increases with particle size due to increase in the anisotropy energy, KV compared to the thermal energy, kBT.  However, surface anisotropy tends to increase with particle size reduction [150]. The anisotropy constant of CoFe2O4 particles of size about 3.3 nm is 3.15x106 J/m3 which is an order higher than that of the bulk materials [151].

Mn-Zn ferrites prepared through the hydrothermal route have resulted in smaller particle sizes [145]. The control over particle size to achieve better magnetic properties was attempted for Mn-Zn ferrites and a growth assisted coprecipitation had yielded magnetization as large as 50 A m2/kg compared to the coprecipitated samples [152] with an average crystallite size of 12 nm.

The size-dependent magnetic properties of various ferrites have been studied for particles with sizes less than 25 nm synthesized using aqueous methods. The synthesis of larger particles involve heat treatment which leads to particle agglomeration. Hence, in this thesis, a modified oxidation technique has been used to synthesize Mn0.67Zn0.33Fe2O4 nanoparticles with a wide range of particle sizes in the as-prepared state itself and their size-dependent magnetic properties have been studied.

 

1.4.2.  Magnetic garnets

Rare earth iron garnets (RIG) with the general formula R3Fe5O12 are an important class of ferrimagnetic materials because of their high frequency applications and have been extensively studied [153-157]. The rare earth garnets have three different crystallographic sites with 16 Fe3+ ions in the octahedral [a] sites, 24 Fe3+ ions in the tetrahedral (d) sites and 24 R3+ in the dodecahedral {c} sites. The cation arrangement in four octants of the unit cell of garnets is represented in Fig. 1.8.(a) and Fig. 1.8.(b) shows the position of the [a], {c} and (d) site cations at the centres of octahedra, dodecahedra and tetrahedra respectively for yttrium iron garnet. The cube represents one of the eight octants of the unit cell and the distance in Ĺ represents the interatomic distances to the nearest oxygen ion. The temperature dependent magnetic behaviour of garnets is governed by the relative magnitudes of the sublattice moments. The magnetic behaviour of some garnets with temperature is shown in Fig. 1.9.(a). The temperature dependent magnetization of the rare earth iron garnets shows a compensation temperature at which their magnetization becomes zero, which is followed by an increase in magnetization and becoming paramagnetic at the Curie temperature for all the rare earth garnets from Gd to Yb. Some of the rare earth ions like Y3+ do not have a permanent magnetic moment and hence they do not exhibit a compensation temperature. Figure 1.9.(b) shows the temperature dependent magnetic moment of [a], (d) and {c} sublattices in Gd3Fe5O12. The overall temperature dependent magnetic behaviour is represented by the thick line. The magnetic contribution arises from the antiparallel alignment of the magnetic moments in the {c} site to the resultant of the antiparallely coupled magnetic moments in the other two sites having Fe3+ ions. The net magnetization direction at 0 K is determined by the {c} sublattice magnetization due to itslarge magnetization at low temperatures. At high temperatures the weak coupling between the magnetic moments decreases the magnetization of the {c} sublattice rapidly. At the compensation temperature (Tcomp), the {c} sublattice magnetization is equal and opposite to the resultant of the [a] and (d) sublattice magnetizations. The compensation temperature of gadolinium iron garnet (GdIG) is exhibited at 290 K [158].

Unlike spinel ferrites, the size-dependent magnetic properties of the garnets is not studied in detail. The rare earth iron garnets requires sintering at high temperature (above 1473 K) for the phase formation [159] which results in grain growth. Even high temperature sintering had resulted in the formation of small traces of orthoferrites [160]. Y3Fe5O12 is obtained after sintering the amorphous precipitate. The phase transition temperature from amorphous to orthorhombic is 1123 K whereas the transition to cubic YIG phase occurs at 1373 K [161]. The cubic phase formation temperature is lowered to 1273 K when the precursors are ball-milled prior to sintering [162]. YIG nanoparticles have been synthesized with particle size as low as 20 nm with a focus on the structure and morphology [163]. The size-dependent magnetic properties of Y3Fe5O12 have been reported by Vaqueiro et al. [164] in the size range 50-700 nm and by Sanchez et al. [165] for the particle sizes from 45 nm to 450 nm. A decrease in the saturation magnetization with size reduction is observed due to surface spin effects. For an average size of 60 nm, the magnetization decreases by one third of the bulk magnetization due to surface spin effects [166].

Most of the techniques to synthesize the garnets involve heating the precursors at high temperatures. Although YIG has been synthesized with small grain sizes, the synthesis and magnetic properties of nanocrystalline Gd3Fe5O12 have not been, so far, reported in the literature. Also, the change in the structure and magnetic properties have not been reported so far for the garnets on ball milling. Since it is well known that ball milling could reduce the grain size, the author is interested in studying the changes in the structural and magnetic properties of magnetic garnets on mechanical milling.

 

1.4.3.  Sm-Co permanent magnetic materials

Sm-Co permanent magnetic materials are used for high temperature applications as their Curie temperature is higher compared to Nd2Fe14B. The Sm-Co compounds include SmCo5 having the CaCu5 structure and Sm2Co17 exhibiting structural forms like Th2Zn17

and Th2Ni17 [167] commonly represented as 2:17 structure and these structures are shown in Fig. 1.10.  In addition to these structural forms, a disordered form of CaCu5 structure having P6/mmm space group where one of the Sm atoms is replaced randomly with a transition metal dumbbell pair randomly in the structural form of TbCu7 is also commonly reported [168,169]. The substitution of Fe in the Sm-Co system increases the magnetization whereas the Zr or Cu addition helps in the stabilization of the TbCu7 phase and also increases the magnetic anisotropy [170,171]. The Sm-Co permanent magnetic materials can be synthesized through mechanical alloying or melt spinning techniques [172,173]. However, the synthesis through both the techniques had resulted in the formation of two or more phases [174-177], which could affect the magnetic properties. Chen et al. studied the magnetic properties of mechanically milled and annealed PrxCo100-x (x = 15.4-20.5) and found that the presence of two phases could affect the coercivity whereas a high coercivity can be obtained in single phase magnets with a higher anisotropy field [178,179]. An enhancement in the coercivity and squareness ratio for lower milling durations followed by a gradual deterioration, on further milling, has been observed in the case of SmCo5 permanent magnetic materials [109]. Ball milling SmCo5 with antiferromagnetic powders such as NiO in an appropriate ratio has resulted in an enhanced coercivity and squareness ratio due to the coupling between both the phases [180]. The increase in the coercivity is attributed to various reasons like grain size reduction and pinning of the domain walls by the phases such as SmCo5, FeZrCu [181] which segregate at the grain boundary. A high coercivity of 688 kA/m to 864 kA/m at a temperature of 723 K has been achieved throughproper heat treatments and development of cellular microstructure having Sm2Co17 as the main phase and SmCo5 as the grain boundary phase [182-184]. Figure 1.11 shows the cellular microstructure of the Sm-Co alloy in the parallel and perpendicular to the alignment direction. The mechanism for the increased coercivity is reported as due to pinning [182],  alignment effects [183] or nucleation [185]  but the exact mechanism is not yet understood. However, recent calculations have predicted that the temperature dependence of the coercivity is determined by repulsive pinning of domain walls at the cell walls at low temperatures, by attractive pinning of domain walls at the cell walls at intermediate temperatures and by nucleation at high temperatures [186] as represented in Fig. 1.12. Studies on the structural and magnetic properties of mechanically milled Pr-Co-Zr powders have shown that the TbCu7-type structure has high anisotropy fields up to 8000 kA/m [187]. The coercivity is found to decrease on grain size reduction for the TbCu7 phase of Sm2Co17, where the Sm2Co17 particles are homogeneously separated by Cu particles, which is obtained by co-milling the annealed Sm-Co powder along with Cu as reported by       Zhang et al. [188]. In this case the decrease has been attributed to the particles approaching superparamagnetism. Prolonged milling may also lead to amorphisation [189] resulting in a decrease in the coercivity which can be again increased by annealing. Although all the reports focus on increasing the coercivity with proper microstructure involving complex heat treatment conditions, the magnetic properties of these materials with the TbCu7 structure have not been understood. Hence the author is interested in studying the changes in the magnetic properties on mechanically milling the Sm(Co,Fe,Cu,Zr)7 permanent magnetic material with the TbCu7 structure.

 

1.4.4.     Nanocomposite permanent magnetic materials (Nd2Fe14B/a-Fe)

One of the important permanent magnetic materials offering high coercivity and energy product is Nd2Fe14B which was discovered by Sagawa et al. [190]. Nd2Fe14B magnetic material is the highest energy product material available till today and many reviews are available in the literature [191-193].  The higher anisotropy field value of this magnetic material has given rise to larger coercivity and hence its energy product is higher. The energy product is defined as the area enclosed in the second quadrant of the B-H loop and hence high magnetization and coercivity values give rise to high energy products. The energy product of the single phase magnet initially reported was 320 kJ/m3 and with a considerable research it has now reached a maximum of over 440 kJ/m3 [194]. Since the energy product is inversely proportional to the volume of the magnet, a higher energy product leads to miniaturized magnets. It is mainly used in applications like voice coil motors in hard disk drives, communication devices, MRI, traveling wave tubes and  in a variety of house hold products. However, the single phase Nd2Fe14B magnet has some drawbacks like very low Curie temperature of 585 K and also they are highly corrosive in nature.

When a hard magnetic material like Nd2Fe14B and a soft magnetic material like a-Fe with high magnetization are exchange coupled  higher energy product is achieved due to the combined effect of higher coercivity of the hard phase and higher magnetization of the soft phase. Such nanocomposite magnetic materials are called exchange spring magnets which are described in detail by Kneller and Hawig [49].

The schematic description of the exchange spring mechanism is shown in Fig. 1.13. The hard and soft magnetic phases are represented as k and m phases respectively. The hard magnetic phase (k) has a large uniaxial magnetocrystalline anisotropy constant K whereas the soft magnetic phase (m) has a low magnetocrystalline anisotropy and large saturation magnetization. The expression for the equilibrium wall thickness is given by

where A is the exchange stiffness constant. Due to the large magnetocrystalline anisotropy of the hard phase compared to the soft phase, the equilibrium wall thickness of the hard phase (dok) will be smaller than that of the soft phase (dom). Let us consider the k and m phases with their thicknesses corresponding to bm dom and bk dok as shown in              Fig. 1.13.(a) - (c). Let bcm be the critical thickness of the m phase below which the exchange coupling takes place. The Fig. 1.13.(a) - (c) represent the condition where bm > bcm and Fig. 1.13.(d) represents the condition where bm < bcm.  Figure 1.13.(a) represents the saturation remanence in an easy direction where all the moments of k and m phases face up. If a demagnetizing field is applied in a direction opposite to the easy direction, two 180 degree walls will form first in the m phase due to its soft magnetic nature as shown in (b). Further increase in the demagnetizing field will compress the walls towards the boundary with the k phase as shown in Fig. 1.13.(c). The compression of the walls towards the k phase will continue till a critical field at which the walls penetrate the k phase at reverse fields greater than the magnitude represented in the figure. At this critical field Hno, the irreversible magnetization of both the m and k phase reverses. For bm < bom as shown in Fig. 1.13.(d), the thickness of the 180o wall is confined to bm. The nucleation field does not change but the coercivity increases as the applied field is less than the coercive field.

Figure 1.14.(a) clearly shows that the exchange spring magnet with the critical m phase thickness shows a broader hysteresis loop whereas the presence of two independent phases would result in a constricted loop as shown in (b). The width of the hysteresis loop can be increased if the nucleation field is increased through exchange coupling. The enhanced remanence and Curie temperature, corrosion resistance and lower cost make exchange spring magnets potential candidates for permanent magnet applications. Theoretical calculations have shown that a multilayer composed of alternating 2.4 nm thick hard magnetic Sm2Fe17N3 layer and a soft magnetic Fe65Co35 layer of thickness 9 nm canhave an energy product as high as 1 MJ/m3 (120 MG Oe) with a rare-earth content of only 5 wt.% [195]. The range of exchange interactions is given by the Bloch wall width, dB given by

where K is the magnetocrystalline anisotropy constant of the hard magnetic phase. In nanocomposites with hard and soft magnetic phases, theoretical calculations have shown that the maximum coercivity can be achieved only when the size of the soft magnetic inclusion is less than twice the domain wall width of the hard magnetic phase [196] and hence for Nd2Fe14B/a-Fe, the size of the soft magnetic phase should be less than 10 nm. Although the theoretical energy product of the Nd2Fe14B/a-Fe exchange spring magnets is of the order of 880 kJ/m3, the experimentally achieved value is too low of the order of 192 kJ/m3 [197]. A considerable amount of work has been carried out to improve upon the energy product. Numerous modeling studies have been carried out, which suggest that the phases should be homogeneous and the volume fraction of the phases should be optimum [198-200]. Recently Zeng et al. [201] have succeeded in synthesizing FePt/Fe3Pt exchange coupled magnets, through chemical methods, with energy products greater than the single phase magnets. The synthesis of Nd2Fe14B through chemical methods has not been reported and they are commonly synthesized through mechanical alloying and melt spinning techniques. In the mechanical alloying technique, powders of the constituent elements or alloy ingots are milled together followed by a post annealing step. It is difficult to handle the mechanically alloyed powders since even a short exposure to air may deteriorate the magnetic properties as the large surface area of the powders leads to oxidization. Another technique used in the synthesis of nanocomposite magnetic materials is the melt spinning method. This method is less cumbersome and the desired grain size can be achieved by annealing the synthesized ribbons. The materials can be synthesized in the crystalline form either directly without the requirement for any post annealing step or in the amorphous form and subsequently crystallizing it by heat treatment. The nucleation and reversal of the magnetic domains control the coercive field during a magnetization reversal process. The nucleation field for an ideal homogeneously magnetized material was shown by Brown [202] as

where K1 is the anisotropy constant, Js is the saturation polarization, Np and are the demagnetization factors parallel and perpendicular to the axis of ellipsoidal particles respectively. Since the nucleation field is affected by the microstructure of the material, Kronmüller introduced microstructural parameters and thus the nucleation field becomes [203,204]

where the microstructural parameter aK describes the inhomogeneity in the magnetocrystalline anisotropy in the grain boundary region and ay  is the parameter for the misaligned grains. The nucleation field for which the reversible behaviour occurs increases with Nd content as reported by Fang et al. [205]. Enhanced remanence ratio has been observed in single phase Nd2Fe14B ribbons [206] and also in nanocomposites with grain size reduction [207-210]. The nanocomposite phase formation depends on the chosen quantity of Nd. The two phase composite can be obtained by choosing a Nd concentrationbelow 12 at.% as shown in Fig. 1.15 [211]. However, when the Nd concentration is below 4.5 at.% Fe3B soft phase is formed which is not desirable because of its lower saturation polarization compared to that of Fe. Hence by limiting the Nd concentration in the range from 4.5 to 12 at.%, the relative volume fraction of the soft and hard magnetic phases can be suitably adjusted to get the desirable energy product. It has been proved experimentally that a reduction in the magnetic anisotropy of the hard phase can enhance the exchange coupling between the hard and soft magnetic phases [212]. All these studies aim to enhance the exchange coupling between the phases in order to improve upon the energy product. As the exchange coupling is shown to enhance the remanence ratio, studies are focused to enhance the remanence ratio. However, the remanence ratio can be enhanced due to dipolar coupling also. Hence the author is interested in studying the nature of coupling in the nanocomposite Nd2Fe14B/a-Fe and also the factors that influence the exchange and dipolar couplings.

 

1.5.   Conclusion

In this chapter a brief introduction to the nanomaterials, especially the nanomagnetic materials, their physical properties and application is presented. The various synthesis techniques used by the author are described. A brief review on the structure and magnetic properties of ferrites, garnets and Sm-Co permanent magnetic materials is provided. Finally the exchange spring mechanism to improve the energy product of the nanocomposite Nd2Fe14B/a-Fe is discussed.

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