theory and advances in magnetic materials

Acta Materialia 51 (2003) 5907–5939www.actamat-journals.com

Recent advances and future directions in magneticmaterials�

D.C. Jiles∗

Materials and Engineering Physics Division, Ames Laboratory, Iowa State University, Ames, IA 50011-3020, USA

Accepted 31 August 2003

Abstract

This paper reviews recent developments in four important categories of magnetic materials that are currently oftopical interest: soft magnets, hard magnets, magnetomechanical and magnetoelectronic materials. For each categorythe various properties of prime interest are discussed, how these differ from one class to another, and how theseproperties can be controlled. Recent developments in materials are highlighted through the consideration of a fewselected new magnetic materials which are at the leading edge of current research, including: amorphous magneticfibers, nanocrystalline permanent magnet materials, ferromagnetic shape memory alloys and spintronic materials. 2003 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

Keywords: Magnetic properties; Hard magnets; Magnetoresistance; Magnetostriction; Soft magnets

1. Introduction

1.1. Magnetic properties of materials

The magnetic properties of materials are meas-ured from certain defined points and derivativesobtained from the variation of magnetization withmagnetic field as shown inFig. 1. Magneticmaterials are broadly classified into two maingroups with either hard or soft magnetic character-istics. Soft magnetic materials can be magnetizedby relatively low-strength magnetic fields, andwhen the applied field is removed, they return to

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Science and Engineering: Past, Present and Future, edited byS. Suresh.

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a state of relatively low residual magnetism. Softmagnetic materials typically exhibit coercivitiesvalues of approximately 400 A m�1 (5 Oe) to aslow as 0.16 A m�1 (0.002 Oe). Soft magneticbehavior is important in any application involvinga change in magnetic induction. Hard magneticmaterials retain a large amount of residual magnet-ism after exposure to a magnetic field. Thesematerials typically have coercivities, Hc, of 10kA/m (125 Oe) to 1 MA/m (12 kOe). The materialsat the high coercivity end of this range are knownas permanent magnets. These materials are usedprincipally to supply a magnetic field.

The magnetic properties of materials can be div-ided into two general categories: those that arestructure sensitive and those that are structureinsensitive. Structure insensitive refers to proper-ties not markedly affected by changes in materials

5908 D.C. Jiles / Acta Materialia 51 (2003) 5907–5939

Fig. 1. Magnetic properties of materials as defined on the BHplane or flux density B versus magnetic field H, (or the MHplane of magnetization M versus magnetic field H). Theseinclude coercivity Hc, remanence BR (MR), hysteresis loss WH,initial permeability min (initial susceptibility cin), maximum dif-ferential permeability mmax (maximum differential susceptibilitycin) and saturation flux density Bs (saturation magnetizationMs).

processing (heat treatment or mechanicaldeformation) or by small changes in composition,including small amounts of certain impurities.Structure-insensitive properties include the satu-ration magnetization and resistivity. These proper-ties are largely dependent on the composition ofthe particular alloy and are not changed substan-tially in the process of manufacturing a componentfrom the alloy.

Structure-sensitive properties are those that aredrastically affected by impurities. Small amountsof elements such as carbon, oxygen, nitrogen, andsulfur are commonly found in small quantities inmagnetic materials. These elements tend to locateat interstitial sites in the crystalline lattice andconsequently the lattice can be severely strained.As a result small concentrations of these elementscan have large effects on some of the magneticproperties of the materials. Permeability, coerciv-ity, hysteresis losses, remanence, and magneticstability are all considered to be structure sensitive.The structure sensitive properties are controlledthrough processing of the material including mech-anical and thermal treatments.

1.2. Units in magnetism

One can not go far in the subject of magnetismwithout encountering the problem of units, whichis rather unique to this discipline. The reasons forthis problem are hardly understood by those out-side magnetism, but are a cause for continuedcontroversy by the practitioners in the field. Thereare currently three main systems of units in wide-spread use in magnetism and several other systemsof units which are variants of these. The three unitsystems are the Gaussian or cgs system and twoMKS unit systems, the Sommerfeld conventionand the Kennelly convention. The basic units andequations of these three systems are summarizedin Table 1. Each of these unit systems has certainadvantages and disadvantages. The cgs and SI sys-tems of magnetic units have different philosophies.The cgs system is based on an approach usingmagnetostatics and the concept of the ‘magneticpole’ , while the SI system takes an electrodynamicapproach to magnetism based on electric currents.The SI system of units was adopted at the 11thGeneral Congress on Weights and Measures(1960). The Sommerfeld convention was sub-sequently the one accepted for magnetic measure-ments by the International Union for Pure andApplied Physics (IUPAP).

2. Soft magnetic materials

Soft ferromagnetic materials are those magneticmaterials with high permeability, low coercivityand low hysteresis loss [1, p. 600], which can beused to amplify the flux density generated by amagnetic field. The available range of magneticproperties of soft magnetic materials is continuallybeing expanded. This amounts to reduction in coer-civity, increase in permeability and consequently adecrease in hysteresis loss. Normally this can beaccomplished through reduction in anisotropy andreduction in domain wall pinning. The latter canbe strongly affected by the impurity content.

5909D.C. Jiles / Acta Materialia 51 (2003) 5907–5939

Table 1Principal unit systems currently used in magnetism and the corresponding equations for flux density B, energy E and torque τ

SI SI EMUQuantity (Sommerfeld) (Kennelly) (Gaussian)

Field H A/m A/m oerstedsInduction B tesla tesla gaussMagnetization M A/m - emu/ccIntensity of magnetization I - tesla -Flux � weber weber maxwellMoment m Am2 weber metre emuPole strength p Am weber emu/cmField equation B = m0(H + M) B = m0H + I B = H + 4pMEnergy of moment (in free space) E = �m0m·H E = �m·H E = �m·HTorque on moment (in free space) t = m0m × H t = m × H t = m × H

Note: The intensity of magnetization I used in the Kennelly system of units is merely an alternative measure of the magnetizationM, in which tesla is used instead of A/m. Under all circumstances therefore I = m0M.

2.1. Magnetic properties

2.1.1. PermeabilityPermeability is the most important parameter for

soft magnetic materials since it indicates howmuch magnetic induction B is generated by thematerial in a given magnetic field strength H.Initial permeabilities of all magnetic materialsrange from mr = 1,000,000 [1, p. 603] in materialssuch as amorphous alloys down to as low as mr =1.1 in some of the permanent magnets. It is knownthat initial permeability and coercivity have inbroad terms a reciprocal relationship, so thatmaterials with high coercivity necessarily have alow initial permeability and vice versa.

2.1.2. CoercivityCoercivity is the parameter which is used to dis-

tinguish hard and soft magnetic materials. Tra-ditionally a material with a coercivity of less than1000 A m�1 is considered magnetically ‘soft’ . Amaterial with a coercivity of greater than 10,000A m�1 is considered magnetically ‘hard’ . Lowcoercivities are achieved in nickel alloys such aspermalloy in which the coercivity can be as lowas 0.4 A m�1 [2].

2.1.3. Saturation magnetizationThe highest saturation magnetization available

in bulk magnetic materials is Ms = 1.95 × 106 A/m(Bs = 2.43 T) which is achieved in an iron-cobaltalloy containing 35% cobalt. The possible valuesof saturation magnetization then range downwardcontinuously to effectively zero. There has beenlittle progress in improving the range of saturationmagnetization of materials for about 100 years.

2.1.4. Hysteresis lossThe hysteresis loss is the area enclosed by the

hysteresis loop on the B, H plane. It represents theenergy expended per unit volume during one cycleof the hysteresis loop. The hysteresis loss increasesas the maximum magnetic field reached during thecycle increases. This loss is closely related to thecoercivity so that processing of materials to reducecoercivity also reduces the hysteresis loss. Gener-ally low hysteresis loss is a desirable characteristicof soft magnetic materials.

2.1.5. Energy dissipation and power lossesA related property is the power loss which arises

when a soft magnetic material is subjected to atime dependent magnetic field. The hysteresis lossis only one component of the power loss, being the


power loss obtained when the field is cycled veryslowly under quasi-static conditions. The totalpower loss depends on the frequency of excitation,the amplitude of magnetic induction, the hysteresisloss, the physical dimensions of the material beingmagnetized and the eddy current dissipation. Inaddition there is usually a discrepancy between themeasured power loss and the loss expected fromthe sum of hysteresis and eddy current losses andthis is usually referred to as the “anomalous loss” .The total electrical power loss can be expressed asthe sum of these various components. Total powerlosses can be reduced if the conductivity of thematerial is reduced.

2.1.6. Barkhausen effectThe Barkhausen effect comprises a series of dis-

continuous changes in magnetization, sometimesknown as ‘Barkhausen jumps’ , that are caused bythe sudden motion of domain walls or by the sud-den rotation of magnetization within a domain.These can be observed along the magnetizationcurve provided sufficient resolution is used in mea-suring changes in flux density or magnetization.

2.2. New soft magnetic materials

2.2.1. Nanocrystalline magnetic materialsIn recent years the range of available soft mag-

netic materials has been significantly increased bythe development of nanocrystalline magneticmaterials beginning with Yoshizawa et al. [3]. Thematerials are mostly based on iron alloys and havegrain diameters of typically 10–15 nm. The excep-tional properties of these materials, which havecoercivities below 1 A m�1 (0.0125 Oe) and highrelative permeabilities of typically 105 combinedwith relatively high saturation magnetization ofMs = 1.05 × 106 A m�1 (13 kG) and resistivitiesas high as 1.15 × 106 ohm m. The most widelyinvestigated alloy is Fe73.5Si13.5B9Nb3Cu1, which isproduced by rapid solidification and is thenannealed above its crystallization temperature toproduce the nanocrystalline structure. Develop-ments in these nanocrystalline materials have beendescribed by Herzer [4].

2.2.2. Artificially structured magnetic materialsRecent developments in soft magnetic materials

have included what may be termed artificiallystructured materials in which the structure of het-erogeneous materials is carefully controlled to pro-duce the desired magnetic properties. A hetero-geneous magnetic material consisting ofnanocrystalline particles embedded in an amorph-ous magnetic matrix phase, as shown in Fig. 2, inwhich the magnetic properties of the particulateand matrix phases are essentially different, allowsconsiderable scope for controlling the magneticproperties through structural changes broughtabout by annealing.

One example of this has been described by Her-nando et at. [5], in which iron-silicon-boron nanoc-rystals are embedded in an amorphous matrix ofthe same material. This resulted in a material witha desirable low magnetic anisotropy, and hencelow loss. The Curie temperature of the material,which is normally low for the amorphous alloys,can be controlled by the presence and size of thenanocrystalline particles. Curie temperatures in therange 200–350 °C can be obtained by annealing atdifferent temperatures and for different times [6].These are multiphase materials consisting of par-ticulate and matrix materials with different mag-netic and structural phases in which the eventualproperties of the entire material can be controlledby structural changes to produce a material with

Fig. 2. Schematic of an artificially nanostructured magneticmaterial consisting of nanocrystalline particles embedded in anamorphous matrix, as described by Hernando et al. [5].


the desired combination of anisotropy and Curietemperature.

2.2.3. Amorphous magnetic ribbonsSoft magnetic alloys can be prepared in amorph-

ous, noncrystalline form, known as metallicglasses. These have low anisotropy combined withhigh resistivity. As a result they have high per-meability and low power losses. In the preferredplanar-flow casting method of production, themetal is rapidly quenched from the melt ontocooled rotating drums to form long ribbonsapproximately 30–50 µm thick which can be up to1 m wide. Those produced for soft magneticmaterials are based on iron-, nickel-, and cobalt-alloys (and sometimes a combination of two ofthese magnetic elements). They also contain vary-ing amounts of metalloids (boron, silicon, carbon,and phosphorus) which help to form the amorph-ous “glassy” phase, and other alloying elementssuch as chromium, molybdenum and aluminum.

These materials are characterized by low hyster-esis loss and low coercive force. However, theCurie temperature is limited to approximately 475°C and magnetic saturation is limited to approxi-mately Ms = 1.3 × 106 A m�1 (Bs = 1.7 T or 17kG).

The low electrical conductivity of metallicglasses compared with crystalline soft magneticmaterials, such as grain-oriented silicon steels,makes amorphous metals extremely attractive forreducing eddy current losses. The low coerciveforces compared with competing crystallinematerials also results in lower hysteresis loss inamorphous metals.

2.2.4. Amorphous magnetic fibersOne of the most important recent developments

in soft magnetic materials has been amorphousmagnetic fibers. Like the amorphous magnetic rib-bons these are produced by rapid solidificationfrom the melt, but in this case the material isejected in a jet from a nozzle and quenched in astream of liquid, usually water, but in some casesoil can be used for particularly reactive metals. Thefinal product is a fiber or wire typically 50 µm indiameter which has no long range crystalline struc-ture.

The compositions of amorphous magnetic fibersthat have been most widely investigated are TMx

Siy Bz, where TM is one of the transition metalsFe, Co or Ni. Typically x = 0.7–0.8, while y andz are each in the range 0.1–0.2. In some cases othertransition metals are added such as Cr whichimproves the corrosion resistance. Fibers withpositive magnetostriction based on Fe can be pro-duced as well as fibers with negative magnetostric-tion based on Co. Those containing Co:Fe in theratio of 16:1 result in an alloy with nearly zeromagnetostriction.

Some of the early work on these amorphousmagnetic fibers was by Ohnaka et al. [7] and byMohri et al. [8]. The first application identified byMohri for these materials was in ‘ jitterless’ pulsegenerator elements, in which the sharp magnetiz-ation changes that occur in these materials over awide range of frequency of applied field, are ideal.These rapid magnetization changes are due to largeBarkhausen effect jumps inside the material whichcan be induced by field strengths as low as 10 Am�1 (0.12 Oe).

In most cases the amorphous fibers that havebeen produced and studied are iron-basedmaterials, such as the Fe81Si4B14C1 composition.Other compositions based on the FeSiB glass-for-ming alloy include FeCoSiB, FeCrSiB and FeNi-SiB as reported by Mohri et al. [9], who also sug-gested a domain model for the as-cast fibers whichcan explain the unusually large Barkhausen jumpsin these materials.

Accordingly the basic model consists of an axi-ally oriented ‘core’ domain running along thelength of the fiber, surrounded by outer or ‘shell’domains as shown in Fig. 3. The large Barkhausenjumps are caused by a sudden reorientation of the

Fig. 3. Magnetic domain structures in amorphous wires withpositive and negative magnetostrictions, after Mohri et al. [9].


core domain. The orientation of the ‘shell’ domainsdepends crucially on the magnetostriction of thematerial because the quenching process induceshigh levels of radial stress and hence stress-induced anisotropy in the radial direction. Formaterials with positive magnetostriction thisresults in radially directed magnetization in the‘shell’ domains, whereas for materials with nega-tive magnetostriction the ‘shell’ domains orient cir-cumferentially.

The magnetic properties of these amorphousmagnetic fibers are strongly affected by post-pro-duction heat treatment, as demonstrated by Atkin-son et al. [10], in which FeSiB alloy fibers werestudied. Annealing at 425 °C relieves the castingstresses in the fibers and results in an alteration ofthe ratio of core domain volume to shell domainvolume. This has dramatic effects on the coerciv-ity, remanence and field-induced magnetostriction.For example, the field-induced magnetostrictioncan be reduced by 50%, from 30 × 10�6 to 15 ×10�6, by simply annealing the fibers for 600 s at425 °C. The coercivity was also reduced from 7.0to 2.9 A m�1 and the remanence ratio Mr/Ms wasincreased from 0.46 to 0.77.

Mohri et al. [11] have discussed the possibleuses of the new amorphous magnetic wirematerials. Many uses are directed principallytowards sensors, which arises because of their highsensitivity of the magnetization to magnetic field,the rapid response and low powerconsumption/dissipation. The two main conceptsare the giant magnetoimpedance effect (seen inamorphous wire materials with saturation mag-netostriction ls ~ 0) and the stress impedance effect(seen in amorphous wire materials with ls � 0).The giant magnetoimpedance effect (GMI) inamorphous wires of FeCoSiB was first reported byPanina and Mohri [12]. Values of 60% wereobserved which was attributed to the skin effectand the strong field dependence of the circumfer-ential permeability. This GMI effect is perhaps oneof the most promising magnetotransport phenom-ena in new nanostructured magnetic materials.

Chiriac et al. [13] have studied saturation mag-netostriction in amorphous wires and hasdeveloped nearly zero magnetostriction Co richamorphous magnetic wires. Interest in the zero

magnetostriction wires is high because of theiralmost reversible magnetization curves. They alsostudied the variation of saturation magnetostrictionwith stress which gave a derivative of dl/ds = –0.9 × 10�10 per MPa. Annealing of the materialcauses the sign of the saturation magnetostrictionto change from –0.11 × 10�6 to +0.3 × 10�6.Change in the sign of magnetostriction causes dra-matic changes in the domain structures because ofthe radially oriented tensile residual stress in theas-cast wires. This in combination with the mag-netostriction determines the orientation of the outershell domains, which results in differences in per-meability and initial magnetization curves asshown in Fig. 4. Chiriac et al. [14] also investi-gated the magnetoelastic behavior in these Co-richamorphous glass covered wires. In this the effectsof iron content and stress annealing on magnetoel-astic properties of the wires were studied. The mainmaterial of interest was Co68Fe4Si13B15 which hasnearly zero magnetostriction. The magnetizationdata showed that the hysteresis loops of thesematerials are highly sensitive to stress, whichoffers excellent opportunities for the materials instress sensors.

Chiriac and Ovari [15] also studied the GMIeffect in families of soft magnetic wires as indi-cated in Fig. 5. It was found that magnetostrictionof the materials plays a crucial role because in theabsence of long range crystalline anisotropy thereare very few factors, other than magnetoelastic andmagnetostatic effects that determine magneticproperties. The circularly polarized shell domainstructure is important in defining the properties.These circular domains exist at the surface of thels � 0 amorphous wires. The result was wires withvalues of fractional change in impedance �Z/Z of60%. The effects were found to be different in dif-ferent “ families” of wires.

Zhukova et al. [16] have investigated ways tomaximize the giant magnetoimpedance in amorph-ous wires. A large sensitivity of up to �Z/Z = 300%resulted from the application of a magnetic field.This can be explained in terms of the classical skineffect, using the model for the material thatincludes an outer shell of circumferential mag-netization, and an inner core of axial magnetizationas described previously by Mohri. In some of the


Fig. 4. Stress dependence of the magnetization curves inamorphous metal wires as described by Chiriac et al. [13]. Thelabels 1-6 refer to different tensile stress levels:- for ‘as cast’material 1) a = 0, 2) a = 83, 3) a = 180 MPa, and for‘stress/current annealed’ material 4) a = 0, 5) a = 83, 6) a =180 MPa.

materials a magnetoimpedance of �Z/Z = 615%was reported.

Betancourt et al. [17] undertook theoretical stud-ies to investigate a domain model for understand-ing the observed magnetoimpedance of amorphousmagnetic wires. The material of most interest is125 µm diameter amorphous wires ofCo68Fe4Si13B15, with ls � 0 which has the circum-ferential shell domains. Betancourt performed cal-culations to determine the complex circular per-

Fig. 5. Frequency dependence of the giant magnetoimpedanceeffect in amorphous magnetic wires, Chiriac and Ovari [15].

meability of the wires from a model of inductance.Real components of inductance are associated withpermeability while imaginary components areassociated with dissipative processes. The complexcircular permeability calculated in this way was ingood agreement with experimental observations.Ciurenau et al. [18] has measured the circumfer-ential and longitudinal permeabilities of Co richamorphous wires at high frequencies up to theGHz range.

Li et al. [19] investigated GMI and magnetoelas-tic properties of stress annealed nanocrystallineFeCuNbSiB amorphous magnetic wires. Theeffects of stress annealing were studied using sim-ultaneous applied stress and current. Values of�Z/Z of typically 150% were observed, as shownin Fig. 6. The behavior of the wires was explainedin terms of a simple effective field model wherebythe anisotropy of the material is altered by bothannealing and stress which controlled the magneticfield at which the maximum in �Z/Z occurred.Zhukova et al. [20] has also shown the effects ofstress on the magnetization and magnetic proper-ties of FeSiB amorphous wires. The application ofstress dramatically affects the magnetic propertiesof these materials because they have low ani-sotropy, but sufficiently high magnetostriction


Fig. 6. Stress dependence of the giant magnetoimpedance(GMI) effect in nanocrystalline FeCuNbSiB magnetic wire afterLi et al. [19].

which ensures that the magnetoelastic couplingstrongly affects the easy direction of magnetiz-ation. They measured the effect of tensile stress onthe hysteresis curves of these materials. The resultscan be understood in terms of two factors—mag-netostatic and magnetoelastic energies. Magneticremanence increased monotonically with tensilestress. Large Barkhausen jumps between two stablestates were observed. They also studied the closuredomain structure in the core. The slope of the hys-teresis curves varied strongly with the length ofthe wires.

Losin et al. [21] have looked at the effects oftorsional stress on the magnetic properties ofamorphous magnetic wires. The material of princi-pal interest was again Co68Fe4Si13B15. The tor-sional effect was measured showing that resistanceand reactance of a circuit including these materialsare only slightly affected by stress, whereas thesecond harmonic voltage amplitude and per-meability are strongly affected by stress. ThereforeGMI is itself not useful for stress detection, but onthe other hand hysteresis or permeability measure-ments on the material are useful for stress detec-tion.

Zhukov [22] have studied the magnetic proper-ties of glass coated microwires. These wires haddiameters of typically 100 µm and exhibited highsensitivity of the magnetic properties to stress.Because the magnetic properties of the zero mag-netostriction materials were almost hysteresis free

they have nearly reversible, one-to-one propertiesthat may be useful for sensors. Some of thematerials also showed large Barkhausen jumps andthe giant magnetoimpedance (GMI) effect.Amorphous magnetic wire materials showing mag-netic bistability can be useful for switches.

2.2.5. Other nanostructured soft magneticmaterials

Kulik et al. [23] have also looked at magneti-cally soft nanostructured materials for high tem-perature applications; in this case amorphous rib-bons consisting of nanocrystalline a-Fe/Co in anamorphous matrix. This was achieved by partialrecrystallization of the material to form the nanoc-omposite. Addition of hafnium was found to inhibitcrystal growth so that even after annealing for 700hours at 500 °C the material was still nanocrystal-line. The result was materials with coercivities ofless than 50A/m.

In the search for new magnetic materials withdifferent properties Fert and Piraux [24] investi-gated magnetic nanowires that were electrodepos-ited into the pores of membranes with 30–500 nmdiameters. These were quite different in type fromthe amorphous soft magnetic wires investigated byMohri, Chiriac and Hernando. They were multi-layer wires, with coercivities in the range of hun-dreds of kA m�1 (~ a few kOe), which exhibitedgiant magnetoresistance of 15–25% depending onthe temperature of measurement.

Petzold [25] investigated the properties of mod-ern nanocrystalline soft magnetic ribbons for elec-tronic device applications and these are comparedwith the magnetic properties of other soft magneticmaterials in Fig. 7. The materials were producedfrom the rapidly solidified amorphous ribbons bya partial recrystallization anneal. The result wasmaterials with a unique combination of low coer-civity, high permeability and high thermal stability.

Amorphous magnetic materials provide a rela-tively simple model system for investigatingdomain wall motion. Himeno et al. [26] studied thedynamics of magnetic domain walls in amorphousmagnetic wires and ribbons with an artificial con-striction. The trilayer wires were prepared by elec-tron beam lithography. The wires were similar tosome of the layered structures that had already


Fig. 7. Comparison of the initial permeabilities and saturationmagnetizations of a variety of soft magnetic materials after Pet-zold [25].

been investigated for GMR in films and multilay-ers. The magnetization reversal process was stud-ied using giant magnetoimpedance measurements.A constriction in the wire acted as a pining site fordomain walls. Constriction widths of 0.25, 0.30and 0.35 µm were studied and it was found thatthe strength of the required depinning fielddecreased with neck width. However the fractionalchange in resistance �R/R was found to be only0.12%.

Moulin et al. [27] have studied nanocompositesof Fe and Fe/Ni dispersed in a Mn-Zn ferrite. Theconcept of a nanocomposite that can take advan-tage of the best properties of both materials. Thesimple idea here was to improve the saturationmagnetization of the material by forming a nanos-tructured composite. A combination of nanostruc-tured permalloy and ferrite gave a material withhigher magnetization than the ferrites.

3. Hard magnetic materials

The term hard magnetic material is used todescribe materials that have sufficiently high resist-ance to demagnetizing fields. Coercivity is there-fore used to distinguish between hard and soft

magnetic materials. The hard magnetic materialsare classified as those with coercivities above 10kA m�1 (125 Oe). The class of hard magneticmaterials includes permanent magnets. The mostrecent permanent magnet materials have coercivi-ties two orders of magnitude greater than this.

Permanent magnets are used for generating amagnetic field. The energy needed to maintain themagnetic field is stored when the permanent mag-net is initially ‘charged’ (i.e. magnetized initiallyin a high field strength and then to remanencewhen the applied field is removed). Permanentmagnet materials are based on the cooperation ofa large number of magnetic moments within amagnet body to produce a high magnetic inductionwhich is then retained because of a strong resist-ance to demagnetization, usually as a result of highanisotropy. Ferromagnetic elements such as ironand cobalt are widely used, and these are oftenalloyed with high anisotropy rare earth elements toproduce high coercivity. Other elements, such asmanganese or chromium, can be made ferromag-netic by alloying. Ferromagnetic metals combinewith other metals or with oxides to form ferrimag-netic materials. Ceramic (“oxide” ) magnets are ofthis type.

Permanent magnet materials are developed fortheir characteristic magnetic properties of: highremanent magnetic induction, high resistance todemagnetization, and high maximum energy pro-duct. Magnetic induction is limited by compo-sition; the highest saturation induction is found inbinary iron-cobalt alloys. Resistance to demag-netization is conditioned less by composition thanby shape or crystal anisotropies and the mech-anisms that subdivide materials into microscopicregions. Precipitations, strains and other materialimperfections, and fine particle technology are allused to obtain a characteristic resistance to demag-netization.

Several different types of permanent magnetmaterials are now available and improvements inproperties such as coercivity and maximum energyproduct continue to be made in rare earth transitionmetal permanent magnets and the development oftwo-phase ‘exchange spring’ magnets. The mag-netization state of a permanent magnets is normallynot deliberately changed. This need for stability of


output means that insensitivity to temperatureeffects, mechanical shock, and demagnetizingfields are important considerations in permanentmagnets.

Permanent magnet materials include a variety ofalloys, intermetallics, and ceramics. The mostwidely studied materials are: cobalt-rare earthalloys (SmCo5 or Sm2Co17), and neodymium-iron-boron (Nd2Fe14B), iron-platinum, cobalt-platinumand hard ferrites (SrO-Fe2O3 or BaO-6Fe2O3).


When ferromagnetic materials are used as per-manent magnets they must operate in conditionswhere they are subject to their own demagnetizingfield and may even be subjected to various demag-netizing effects of magnetic fields in their vicinity.It is essential that they are not easily demagnetized.They must therefore have high coercivity, reman-ence and energy product. Maximum energy pro-duct is an important parameter because permanentmagnets are used primarily to produce a magneticflux density. Maximum energy content and certainother characteristics of materials used for perma-nent magnets, are best described by its hysteresisloop in the second quadrant which is sometimesreferred to as the “demagnetizing curve” .

3.1.1. CoercivityThe ability of permanent magnets to resist

demagnetization is an important property andconsequently a high coercivity is essential. Overthe years there has been continual progress in thediscovery of new permanent magnet materials withhigher coercivities [28, p. 7]. For example, theintrinsic coercivity of neodymium-iron-boron istypically 1.1 MA m�1 (14 kOe) and for samarium-cobalt it is typically 0.69 MA m�1 (8.7 k Oe). Inpermanent magnets the magnetization is not simplyan approximately linear function of the flux den-sity, unlike soft magnetic materials in which B isapproximately equal to moM. This is because thevalues of magnetic field H used with permanentmagnets are generally much larger than in softmagnetic materials. The result of this is that thecoercivity can be defined as either the field atwhich the magnetization M is zero, the intrinsic

coercivity MHc, or the field at which the magneticflux density B in the material is zero BHc. Thesequantities have quite different values in hard mag-netic materials, and the greater the difference thebetter the material is as a permanent magnet. Itshould be noted that MHc is always greater thanBHc.

3.1.2. RemanenceIn order to act as a source of magnetic field a

permanent magnet is only of use if it has a rela-tively high magnetization when removed from anapplied magnetic field. Therefore a high remanenceis essential and this inevitably means a high satu-ration magnetization. The remanence MR (or BR)is the maximum residual magnetization (or mag-netic flux density) that can be obtained in a closed-loop configuration in which there is no demag-netizing field. Since, to be of any use, all perma-nent magnets must be operated in an ‘open circuit’configuration, the residual magnetization, or fluxdensity, at which the permanent magnet operatesin open circuit will always be below the remanencevalue. Therefore a combination of high remanenceand high coercivity is essential to permanent mag-nets. The remanence in some of the most recenthigh performance Nd2Fe14B is for example typi-cally MR = 1.2 × 106 A m�1 (BR = 1.5 Tesla) [29]and for Sm2Co17 it is MR = 0.72 A m�1 (BR =0.9 Tesla) [30].

3.1.3. Saturation magnetizationThe remanence is of course dependent on the

saturation magnetization, and for this reason thesaturation magnetization of a permanent magnetshould be large. While this condition is necessaryit is not sufficient, since the squareness ratioMR/Ms must also be as close to 1 as possible inorder to ensure a large remanence. The saturationmagnetization in neodymium-iron-boron is MS =1.27 × 106 A m�1 (BS = 1.6 Tesla) and in sam-arium-cobalt it is MS = 0.99 × 106 A m�1 (BS =1.24 Tesla) [31].

3.1.4. Maximum energy productOne parameter that is often of interest to specify

the performance characteristic of a permanentmagnet is the maximum energy product, which is


the maximum value of BH obtained in the secondquadrant. This is closely related to the total hyster-esis loss or area enclosed by the hysteresis loop.The interest of researchers in the maximum energyproduct of permanent magnets stems from the totalenergy associated with a permanent magnet. Themaximum energy product is simply a measure ofthe maximum amount of useful work that a perma-nent magnet is capable of doing outside the mag-net. There have been continual improvements inthe maximum energy product of various permanentmagnet materials, from tungsten steel to rare earthtransition metal alloys, over the last hundred years,as indicated in Fig. 8. The highest values of BHmax

have been obtained in neodymium-iron-boron per-manent magnets [32,33]. The maximum energyproduct is often given in alternative non-SI unitsof MegaGauss-Oersted (MGOe). The conversionfactors between MGOe, J m�3 and ergs/cc are asfollows: 1 MGOe = 7.96 kJ m�3 = 79.6 Kerg/cc.

3.1.5. Demagnetization curveThe maximum energy product only gives limited

information about the properties of a permanentmagnet. A more useful way of displaying the mag-netic properties of a permanent magnet is to plot

Fig. 8. Progress in the maximum energy product in permanentmagnet materials [2, p.367].

the portion of the hysteresis loop in the secondquadrant from the remanence point to the coerciv-ity. This demagnetization curve indicates the mag-netization under various demagnetizing fields.Such a curve contains information about themaximum energy product. The strength of thedemagnetizing field of a permanent magnet inopen-circuit configuration depends on the shape ofthe permanent magnet. Therefore the performanceof permanent magnet material as a field source isdependent as much on the geometry as on theintrinsic material properties.

3.2. New hard magnetic materials

3.2.1. Rare earth transition metal permanentmagnets

Rare Earth Transition Metal (RETM) permanentmagnets, based either on neodymium-iron-boron orsamarium-cobalt alloys, have been the main focusof attention in hard magnetic materials in recentyears. One of the disadvantages of Nd2Fe14B is itsrather low remanent magnetization in the isotropicform of the material. This is about MR = 0.64 ×106 A m�1 (0.8 T) compared with 1.05 × 106 Am�1 (1.3 T) in the anisotropic or oriented form ofthe material. The resultant energy product is 100kJ m�3 (12.5 MG Oe) compared with 320–400 kJm�3 (40–50 MG Oe) in some of the orientedmaterial. In a theoretical paper Kneller and Hawig[34] proposed an idea for improving both the satu-ration and remanent magnetizations in permanentmagnet materials by using a two-phase nanostruc-tured material consisting of a magnetically hardphase embedded in a high saturation magneticallysoft matrix. Such a material allows good exchangecoupling between the two phases so that the highmagnetization of the matrix phase helps to main-tain a high magnetization in the hard phase. Thisidea became known as the ‘exchange spring’ mag-net. An interesting prediction was that permanentmagnet nanocomposites could be formed in thisway consisting of 90% soft phase (such as α-Fe)and only 10% hard phase (such as Nd2Fe14B).

It was found experimentally that a structure offine and regularly dispersed particles with grainsize typically 10 nm does lead to enhanced MS andMR. Hadjipanayis et al. [35,36] were able to pro-


duce isotropic remanence enhanced magnets basedon these ideas. The material consisted of a fine-grained composite of Nd2Fel4B in a matrix of α-Fe which comprised 50–70 wt% of the material.The remanence ratios MR/MS in these materialsimproved from 0.5–0.78, although the coercivitywas much reduced, 320 kA m�1 (4 kOe) comparedwith 1.2 × 106 A m�1 (15 kOe) in the normal neo-dymium-iron-boron material. Davies [37] has alsoreported that remanence enhancements from 0.8–1.2 T are achieved in NdFeB/a-Fe nanocompositeswith grain sizes below 40 nm.

The exchange coupling between the hard andsoft phases becomes relatively more important asthe ratio of interfacial area to volume of the grainsincreases. While the intrinsic coercivity of thesematerials is lower than in normal NdFeB, the pres-ence of the a-Fe on an ultrafine scale apparentlydoes not result in serious deterioration of themaximum energy product, which remains at about160 kJ m�3 (20 MGOe), a level comparable withNdFeB. These “ remanence enhanced” magnetshave received much attention and the idea has beenextended to other materials. Ding, McCormick andStreet [38] have reported large remanenceenhancements in SmFeN using a nanocompositewith an a-Fe matrix.

Theoretical modeling by Schrefl [39] has shownexcellent agreement with experimental results,confirming that the proposed mechanism is the rea-son for the enhancement in the properties. Thesecalculations provided some of the most advancedmaterials property modeling for establishing adirect relationship between magnetic propertiesand the structure of the material.

3.2.2. Nanophase magnets based on rare earthtransition metal alloys

Hadjipanayis [40] has considered variousaspects of “nanophase hard magnets” . These arenanocomposite magnets consisting of a very finenanoscale mixture of exchange coupled hard andsoft phases. As is generally known coercivityincreases as particle size is reduced, reaching amaximum at the single domain size, thereafterdecreasing to zero at the superparamagnetic limit,as shown in Fig. 9. The properties also dependstrongly on whether the demagnetization process

Fig. 9. Dependence of magnetic coercivity on particle size inpermanent magnet materials as described by Hadjipanayis [40].

is principally via a nucleation mechanism or via adomain wall pinning mechanism, as indicated foran example of each in Fig. 10. In exchange coupledmagnets an attempt is made to get the best of bothtypes of materials that comprise the structure. Theygive higher BR than the hard phase, and a higher Hc

than the soft phase. This only applies if the phasesconsist of volumes that are small enough to beexchange coupled. Therefore particle size is a criti-cal factor. Certain trade offs can be made to tailor

Fig. 10. Initial magnetization and demagnetization curves, andfield amplitude dependence of coercivity (inset) in pinning type(eg. Sm2(CoFeCuZr)17) and nucleation type (eg.Nd2Fe14B)magnets, after Hadjipanayis [40].


the properties of the materials, since a larger frac-tion of magnetically hard material gives highercoercivity, whereas a larger fraction of soft giveshigher remanence. The future direction of hardmagnetic materials is towards structures with veryfine particle sizes and a regular distribution of par-ticles with uniform sizes.

Stamps [41] has investigated the mechanism ofexchange bias in hard magnets in which he focusedon magnetic thin film interfaces and spin elec-tronics. Exchange bias is a phenomenon closelyassociated with magnetization and hysteresis inferromagnetic/antiferromagnetic bilayers. Theinteractions between the antiferro- and ferro-mag-netic phases are of great current theoretical interest,since the magnetic properties of the ferromagneticfilm can be controlled by the presence of a nearbyantiferromagnetic layer without the need for amagnetostatic field. Biased hysteresis loops resultfrom the exchange coupling. It was found that therate dependent magnetization processes aredetermined by thermal fluctuations in the materialstrong enough to unpin domain walls.

Skomski et al. [42] investigated exchange in per-manent magnets propagated through a non-mag-netic insulating matrix. In these cases the embed-ding of single domain magnetic particles inparamagnetic matrix should allow the coercivity toapproach anisotropy field H = 2K /moMs. Theexchange coupling decreases exponentially with acharacteristic decay length, which does not exceedseveral interatomic spacings. The characteristicdecay length is affected by the nature of the matrixso that unless the matrix is close to ferromagneticthe decay length will be even shorter and is at mosta few atomic spacings. If the interactomic spacingincreases a strong decrease in exchange energyoccurs. The small characteristic decay length, ~0.1nm, emphasizes the relative weakness of inter-atomic (long range ordering) interactions comparedwith intra-atomic interactions.

Gutfleisch et al. [43] have considered nanocrys-talline RETM permanent magnets. In thesematerials it is the grain size and the intergranularphases that combine to give properties that are dif-ferent from bulk materials. Single domain particlesgive the highest Hc values and for most RETM per-manent magnet materials the single domain size is

�1 µm. In situations where there is an excess ofrare earth metal beyond the stoichiometric compo-sition (for example Nd2Fe14B) the rare earth richmaterial on the grain boundaries results in hardmagnetic grains separated by paramagnetic layersand this gives high coercivity while stoichiometircmaterial RE2Fe14B results in a single phase withexchange coupled grains. On the other hand rareearth deficient material gives RE2Fe14B grainscoupled through Fe3B or other iron rich phasessuch as α-Fe in which the results can be materialswith remanences higher than those expected on thebasis of the Stoner-Wohlfarth model. Increasing(BH)max in this way can give 85% of the theoreti-cal maximum for NdFeB magnets. Neverthelesscoercivity values rarely exceed 30% of the ani-sotropy field.

Chiriac et al. [44] have also investigated glassyhard magnets based on NdFeAl. These were pro-duced as amorphous ribbons with thickness ~150µm and cast rods of the same material, with diam-eters �2 mm. The amorphous NdFeAl and NdFeSiternary alloy systems can be produced in a widerange of compositions. Although structurally thesecan be formed as glasses, their magnetic propertiesare more characteristic of granular materials It wasfound that the coercivity could be increased to 0.6MA/m at 200 K, as shown in Fig. 11, in a structurewhich consisted of a packing of nanometer sizedclusters. The high coercivities of these alloys ofNdFeAl result from the formation of small meta-stable or non-equilibrium magnetic clusters.Results are similar to those obtained on rare earth-iron binary amorphous alloys but different fromthose obtained on materials such as NdFeB inwhich the presence of the amorphous phase leadsto soft magnetic characteristics.

Rodewald et al. [45] have investigated NdFeBsintered magnets with a minimization of RE richconstituents giving a very fine grained structure.These were subjected to pulsed field treatment toalign the particles, resulting in a BHmax of greaterthan 56 MGOe Property values were: Br = 1.5Tesla, (15 KG) MHc = 7.8 KA/m (9.8 KOe), andBHmax = 451 KJ/m3 (56.7 MGOe).

Wang et al. [46] have studied the effects of Prcontent on amorphous Pr(FeCo)B alloys. Exchangeenhanced remanence and energy product obtained


Fig. 11. Temperature dependence of the magnetization (upperfigure) and coercivity (lower figure) of Nd50Fe40Al10 amorphousalloy, after Chiriac et al. [44].

in nanocomposites of this alloy gave higher coer-civity than the corresponding NdFeB alloys due tohigher anisotropy. Annealing of the alloys gavehigh Hc (755 kA/m) and BH max (160 KJ/m3), butthis was obtained only for high Pr contents. Forsub-stoichiometric alloys annealing caused deterio-ration of the hard magnetic properties. However Tc

was lower for the Pr alloy than the Nd alloy.Zhou et al. [47] reported on the effects of struc-

ture chemical composition and mechanical alloyingon the magnetic hysteresis in SmCo nanocrystal-line powders. A room temperature coercivity of 41kOe was observed in Sm2Co7 compound. In sam-

arium cobalt permanent magnets larger coercivityand enhancement of remanence can be achievedthrough control of grain size. Effects of theaddition of Cu to the SmCo alloys were studied.Coercivity was found to increase with Cu content.In addition mechanical alloying and mechanicalmilling lead to improvements in hard magneticproperties.

Betancourt and Valenzuela [48] investigatedenhanced coercivity in REFeAl based alloys. Thespecific alloys studied were NdFeAl and PrFeAlalloys, with small amounts of B, Cu, Dy added.These alloys are of interest because of the simplefabrication and processing. Coercivity in thesealloys is attributed to short range order within theamorphous materials in which the structure actu-ally consisted of an amorphous matrix with a smallvolume of fraction of hcp Nd crystallites embeddedin the matrix. Coercivity was enhanced by theaddition of either B, Cu or Dy. This is attributedmostly to domain wall pinning on the crystallites,but also to a lesser extent to alteration in thenucleation sites. The maximum coercivity obtainedwas 388 KA/m.

Branagan et al. [49] studied high hysteresis ina homogeneous TbAl metallic glass. The materialproduced was completely amorphous. Intrinsiccoercivity approached 23 kOe and BHmax was 12.4MGOe. The domain size in this glass was muchlarger than the structural cluster size. The high Hc

was explained on the basis of exchange bias. Noidentifiable nuclei or second phases present. Aftercrystallization the hard magnetic propertiesdeteriorated, as shown in Fig. 12. A three phasestructure formed, containing both hard and softphases resulting in a wasp-waisted hysteresis loopand associated poor hard magnetic properties.

Wang et al. [50] have looked at the effect of Cosubstitution on (PrTb)2 (FeNbZr)14B alloys. Thisalloy, which is of the “2–14–1” type, gives an indi-cation of how complicated the situation is becom-ing in the search for new and improved permanentmagnet materials as more chemical additions arebeing tried to enhance the range of properties. TheCurie temperatures of the rare earth transitionmetal alloys sometimes precludes their use in cer-tain applications. Therefore a search for permanentmagnets with higher Curie temperatures has been


Fig. 12. Differences in the demagnetization curve of as spunand crystalline (heat treated) Tb60Al40 alloy after Branagan etal. [49].

undertaken. Nanocomposites were formed of bothRE2Fe14B/Fe3B and RE2Fe14B/α-Fe by melt spin-ning. Addition of Co greatly enhances the Curietemperature, and in this case typically 30% of theFe was replaced with Co, so that Tc was increasedfrom 775C in the iron alloy to 950 °C in the Cosubstituted alloy. The highest coercivities obtainedwere 6.8 KOe, with a maximum energy product ofBHmax = 11.8 MGOe.

Ono et al. [51] studied crystallization and mag-netic properties of NdFeCoVB and a-Fe/ NdFeBnanocomposite hard magnetic ribbons. As theamorphous volume fraction increased the crystalgrain size became smaller and hard magneticproperties improved. An optimum quenching ratewas found however which indicates that shortrange order is useful in order to inhibit graingrowth of a-Fe crystals. The completely amorph-ous state was found to be less conducive to hardmagnetic properties than one with a small volumefraction of nanocrystalline clusters containing Nd.

3.2.3. Nanophase magnets based on cobalt oriron and platinum

CoPt and FePt have emerged as candidates forhigh performance exchange spring magnets withenhanced magnetic properties. Both the CoPt and

FePt materials have high anisotropies (4.9 × 106

J/m3 and 6.6 × 106 J/m3 respectively). The FePtsystem has the advantage of complete solid solu-bility at higher temperatures. At lower tempera-tures the crystal structure transforms from fcc to aface centered tetragonal structure which hasimproved hard magnetic properties due to theincreased anisotropy.

Huang et al. [52] have studied hysteresis in CoPtnanoparticles embedded in a non-magnetic matrixconsisting of C and/or BN. The effects of the par-ticle size on properties were investigated. Thinfilms of these materials were prepared by sputter-ing. It was found that coercivity depended princi-pally on the particle size. Annealing time and tem-perature were used to produce optimum particlesizes. Whereas the “as-fabricated” fi lms had Hc �8 kA/m (100 Oe), annealing caused an increase inparticle size and a consequent increase in Hc fromless than 80 KA m�1 (1 kOe) to 1000 KA m�1

(12 kOe).Barmak et al. [53] investigated ferromagnetic

exchange spring nanocomposite films. These con-sisted of CoPt films 10–25 nm thick, produced bydc magnetron sputtering. Various annealing treat-ments were used subsequently. The magneticproperties of the material changed from nucleationdominated behavior at low volume fractions ofmagnetically hard phase, to pinning dominatedbehavior at high volume fractions of hard phase.The particle size in the films increased withannealing time, resulting in increased coercivity.The reason was that the coercivity mechanismbecomes increasingly pinning controlled as thevolume fraction of the high coercivity phaseincreases. Even though the films were in all casestwo phase structures the properties exhibited whatwas described in the paper as “single phase” mag-netization behavior at all stages, meaning that therewere no obvious macroscopic discontinuities in themagnetization curves that could be indicative ofsecond phases.

Crew and Stamps [54] have described ferromag-netic resonance in exchange coupled thin films viatheoretical calculations, using the micromagneticLandau-Lifschitz-Gilbert model. For the experi-mental part of the magnetic resonance studyCo/CoPt exchange coupled thin films were made.


It is known that the resonance frequency is determ-ined in part by the perpendicular anisotropy in thefilm, and that such measurements can be used tofind the perpendicular anisotropy. Howeverexchange coupling between the films caused sig-nificant changes in the resonant frequency. It wasalso shown that magnetostatics plays an importantrole in determining the behavior.

Zhou et al. [55] have investigated improvementsin the properties of FePt through exchange coupledFePt–Fe layered films prepared by magnetron sput-tering, In FePt thin films a maximum energy pro-duct of up to BHmax = 15 MGOe, and in CoPt thinfilms a maximum energy product of up to BHmax

= 50 MGOe were observed as shown in Fig. 13.The fabrication process produced exchangecoupled FePt–Fe multilayers and it was found thatin combination with a rapid thermal anneal animproved energy product of 19 MGOe could beobtained in these multilayers.

Thang et al. [56] also investigated magneticproperties of FePt alloys, in this case with Nband/or Al doping. They studied the combinedeffects of annealing and the presence of thirdelements on the magnetic properties of thesealloys. The materials consisted of exchangecoupled grains, with domain sizes of typically 1–20 nm. A high coercivity resulted principally frompinning of domain walls. In the Al doped materialsthe observed magnetic properties were: BHmax =120 KJ/m3, Hc = 300 KA/m, Br =1.025 Tesla

Fig. 13. Demagnetization curves of annealed FePt single layerfilm, and annealed FePt/Fe multilayer film after Zhou et al. [54].

(10.25 KG, Mr = 0.82 × 106 A m�1). The Nb dopedmaterial had similar properties except for a lowerremanence of Br = 0.975 Tesla (9.75 KG, Mr = 0.78× 106 A m�1)

Huang et al. [57] investigated fabrication andmagnetic properties of FePt/C films containingFePt nanoparticles in order to increase the coerciv-ity compared with the normal FePt thin films. Thefilms were produced with FePt particle sizes in therange 3–20 nm. For the films with the smallest par-ticle sizes the blocking temperature, which marksthe boundary between hysteretic and non-hystereticbehavior, was found to be 200 K as shown in Fig.14. For films with larger particles sizes the block-ing temperature is higher. Once the particle sizeswere such that the blocking temperature was aboveroom temperature it was found that the coercivitieswere greater for the films with larger particles. Acoercivity of 1.84 × 106 A m�1 (23 kOe) wasobtained for the 8 nm diameter particles and acoercivity of 2.72 × 106 A m�1 (34 kOe) for the15 nm diameter particles. It was also found thatby depositing FePt simultaneously with carbon thestructure of the films can be controlled so thatstrong perpendicular anisotropy can be producedin the films, suitable for magnetic recordingmedia applications.

Ding et al. [58] have looked at phase transform-ation in FePt nanoparticles. As is known thematerial transforms from fcc to face centered tetra-gonal with high anisotropy and for this reason the

Fig. 14. Dependence of magnetization on temperature for fieldcooled (F.C.) under 2 KOe (160 KA/m) applied field and zerofield cooled (Z.F.C) nanoparticles of Fe50Pt50 alloy, after Huanget al. [57].


material has potential for use in ultra-high densityrecording media applications. However annealingto cause the desired phase transformation is usuallyaccompanied by sintering of the material whichdegrades the magnetic properties. Differentannealing conditions were tried in order to get thebest magnetic properties by achieving the phasetransformation without the unwanted sintering. Itwas found that particle coalescence occurred atabout 550 °C and once started proceeded very rap-idly. Since this is also the temperature what whichthe phase transformation occurs it remains prob-lematic. One solution that was suggested was toalloy the FePt with some other elements to try toreduce the transformation temperature below thesintering temperature—or alternatively to increasethe particle separation in the film.

Simizu et al. [59] have also investigatedexchange coupling in FePt magnets, the propertiesof which are given in Table 2. Those alloys in therange Fe40Pt60 to Fe60Pt40 alloys, exhibit the phasetransformation from fcc to face centered tetragonal.Both phases can coexist, so this leads to the pros-pect of an exchange coupled magnet consisting ofone material in two different structural and mag-netic phases. The alloy Fe60Pt40 has a (BH)max of120 KJ/m3 (15 MGOe) and high squareness ratioof Mr/Ms = 0.68, a value which is consistent withthat predicted in the exchange coupled model. Thisenhanced remanence was shown not to be due totexturing of the material, since magnetizationcurves were identical in three orthogonal direc-tions. Also it was found that the “as quenched”samples had the hardest magnetic properties, andany annealing procedure that was attempted in

Table 2Energy product, remanence, saturation, remanence to saturation ratio, and coercivity of three samples (x, y and z) of ‘as cast’ materialand two heat-treated samples (16 hours at 600 °C and 16 hours at 650 °C) of Fe0.6Pt0.4

Fe0.6 Pt0.4 As cast 16 h/600 °C 16 h/650 °C

x y z z z

BHmax (MG Oe) 4.87 5.33 5.94 5.36 4.15Br (kG) 6.86 7.34 7.63 7.13 6.49Bs (kG) 10.37 10.92 11.20 10.60 10.36Br/Bs 0.66 0.67 0.68 0.67 0.63Hc (kOe) 2.37 2.34 2.36 2.40 2.11

order to promote formation of the hard magneticphase was found to reduce Hc and Br. Paradoxicallytherefore the best hard magnetic properties wereobtained by suppressing the formation of thehard phase.

4. Magnetomechanical materials

These are materials in which the strong mag-netomechanical coupling causes strain in thematerial when magnetized. The class of magneto-mechanical materials is rather broad, but includesmagnetoelastic materials in which the recovery ofstrain is reversible when the field is removed, andother materials for which the induced strain doesnot necessarily recover simply on removal of themagnetic field. These latter materials could becalled magneto-plastic materials, but are oftendescribed as magnetic shape memory alloys. Highperformance magnetomechanical materials includecobalt ferrite, iron-gallium alloys, rare earth metalssuch as dysprosium and terbium, rare earth-ironalloys such as Terfenol (terbium-dysprosium-iron),and shape memory alloys such as NiMnGa andGdSiGe.


4.1.1. MagnetostrictionMagnetostriction is the fractional change length

of a material when magnetized, either spon-taneously by virtue of a magnetic phase transition(spontaneous magnetostriction) or under theaction of a magnetic field (field induced


magnetostriction). The magnetostriction λ is thefractional change in length under the action of amagnetic field

l � ��

�o� (1)

where �o is the original length and �� is the changein length. For convenience the length �o is usuallytaken as the length in the demagnetized state andthe changes are therefore measured by conventionrelative to the demagnetized state. This effect isonly significant in magnetically ordered materialssuch as ferromagnets, ferrimagnets and antiferrom-agnets. Typical magnetostrictive strains of mag-netic materials are of the order of � � 10�5�10�6,[60] although more recent advanced materials,such as terbium dysprosium iron, have magneto-strictions as high as 10�3 [61,62].

4.1.2. Motor coefficient (strain derivative)The strain derivative dl/dH for magnetostrictive

materials, d, is a useful representation of materialsproperties, since it indicates how rapidly the strainchanges with the relevant applied field,

d � �dldH�s (2)

This is also sometimes known as the strain coef-ficient, strain derivative, or simply as the “d coef-ficient” . When changes expressed by differentialsare small and reversible this d coefficient is alsoequal to (dB/ds)H where B is the magnetic fluxdensity and s is the stress. It is a useful measureof the performance of a material, since it is closelyrelated to the energy conversion efficiency of thematerial. However the dependence strain on mag-netic field is also hysteretic. It is often quoted in therange of nanometers per amp in magnetostrictivematerials. Typical values of the d coefficient forvarious materials are: CoO.Fe2O3 ~1 × 10�9 mA�1, Gd5Si2Ge2, ~1 × 10�9m A�1, Fe ~0.3 × 10�9

m A�1, Ni ~�2 × 10�9 m A�1, Co ~�0.2 × 10�9

m A�1, polycrystalline (Tb0.3Dy0.7)Fe2 ~1.3 × 10�9

m A�1.

4.1.3. Generator coefficientWhen a magnetic material generates a change in

magnetic field by application of a stress the para-meter of interest is the rate of change of H fieldwith stress, dH/ds. This is called the generatorcoefficient g since it generates a field from appliedstress at constant magnetic flux density

g � ��dHds�B

� �dMds�B

(3)

This is measured in units of AN�1m. It is alsoequal to the derivative of the magnetization M withrespect to stress at constant magnetic flux densityB as given in the equation. When the changes rep-resented by the differentials are small and revers-ible this g coefficient is also equal to �(dl/dB)σ.

4.1.4. Coupling coefficientAn important consideration in magnetomechan-

ical materials is often the efficiency for convertingone form of energy into another (eg. mechanicalinto magnetic or vice versa). The input energy isusually supplied by an external influence such asan applied stress or a magnetic field. The responseof the material, is in the form of strain or magnetiz-ation.

The energy conversion efficiency can beexpressed via the square of the coupling coef-ficient, k which is defined as k2 = (energyoutput)/(energy input). In the case of a mag-netostrictive material k2 is related to the materialsproperties by the equation

k2 �EY

momrd2 (4)

where EY is the appropriate elastic modulus, mr isthe relative permeability, mo is the permeability offree space and d is the d coefficient given by Eq.(2) above. However it must be remembered thatin ferromagnetic materials mr is not constant. Thismeans that the energy conversion efficiencychanges with applied field strength.

4.1.5. PiezomagnetismUnder certain conditions the application of stress

to a magnetic material can cause a change in mag-netization. Only a few examples are known of


magnetic materials which can change from ademagnetized to a magnetized state under theaction of a stress. One of these is the antiferromag-net CoF2 in which a small magnetization of theorder of 103 A/m can be produced by large shearstresses [63]. There are no known cases of ferrom-agnets which become magnetized under the actionof a stress on the unmagnetized state. A related andmuch more common phenomenon is the tendencyof magnetic materials to change their magnetiz-ation under the action of an applied stress whenalready magnetized [64]. This effect is only of sig-nificant size in ferromagnets and ferrimagnets. Jiles[65], and later Li and Jiles [66] have studied theseeffects. In these phenomena the order of appli-cation of stress and magnetic field is critical. Theeffects of change in stress on the magnetizationhave been modeled. The hysteretic nature of thebehavior dictates that simple linear differentialequations based on the “small amplitude” Maxwellrelations are not valid for describing these effects.

4.2. New magnetomechanical materials

4.2.1. Gaint magnetostrictive rare earthtransition metal alloys

Much progress has been made in recent years inthe development of giant magnetostrictivematerials such as Terfenol-D (TbxDy1�xFe2 with x~ 0.3). Reliable model descriptions of thesematerials have been developed at the same time,and these can now provide valuable predictivesimulations of materials properties and perform-ance. Calkins, Smith and Flatau [67] havedeveloped a magnetostrictive model based onenergy considerations which gave good agreementwith performance of Terfenol-D. The modeldescribed the B, H hysteresis loop of the material,and then incorporated a non-hysteretic quadraticrelationship between magnetostriction and mag-netization. In this way a hysteretic relationshipbetween magnetostrictive strain and applied mag-netic field was obtained. The model was able todescribe strain under both major and minor loopexcursions of magnetization versus magnetic field.The model is also compact because it has only asmall number of model parameters and can be

operated under a wide variety of conditions includ-ing non-linear regimes.

Dapino, Smith and Flatau [68] have shown howto include non-linear and hysteretic effects directlyinto the description of strain. In this case, both theelastic compliance and the piezomagnetic coef-ficient d were allowed to be variable and pathdependent. This generalization of the magnetostric-tive model is suitable for describing the behaviorof magnetostrictive materials at high fields.

4.2.2. Magnetic shape memory alloysShape memory alloys can exhibit deformations

of 5–10% on application of a magnetic field.Ullakko et al. [69] were first to report unusuallylarge magnetically induced strains in ferromagneticshape memory alloys based on the Heusler alloycomposition. Ni2MnGa is a member of the familyof Heusler alloys which undergoes a thermomech-anical martensitic phase transformation on coolingbelow Tm = 202 K. Since the ferromagnetic tran-sition is at Tc = 376 K the structural transformationoccurs while the material is in its ferromagneticstate. Both transition temperatures Tm and Tc aresensitive to the chemical composition. Substitutionof Mn for Ni can cause Tm to increase and Tc todecrease. These two temperatures coincide forcompositions of x ~ 0.18–0.20 in Ni2+xMn 1�xGa,according to Vassiliev [70], whereupon it trans-forms from a high temperature paramagnetic phaseto a low temperature ferromagnetic phase simul-taneously with the structural transformation fromhigh temperature cubic to low temperature tetra-gonal. The c/a ratio is 0.94, which gives rise to a6% deformation.

The martensitic transition can also be inducedby a magnetic field. The application of a magneticfield leads to a redistribution of the structural stateof the material. In order to obtain a colossal defor-mation under the action of a magnetic field the fol-lowing criteria have been identified: i) the tran-sition has to be thermoelastic, ii) the martensiticphase must be ferromagnetic iii) the martensiticphase must have sufficient anisotropy.

Sozinov et al. [71] have investigated crystalstructures and magnetic anisotropy of Ni2MnGawith giant field induced strain. In this study non-stoichiometric alloys were investigated. It was


found that, depending on the exact details of thecrystal structure, a 6–10% strain can be induced byan applied field of less than 1 Tesla.

Pasquale et al. [72] have also studied magneticand mechanical properties of Ni2MnGa single crys-tals. Large field induced strains were produced inoriented single crystals by pre-treating thematerials by subjecting them to a sequence ofmechanical cycles and thermal treatments. Typicaltreatment needed to “stabilize” the samplesinvolved a load of 20 MPa and a thermal treatmentof 200 °C. Stress and field were applied sub-sequently. A repeatable sequence of “pseudoplas-tic” behavior was obtained which gave a 1.2% irre-versible strain under a field of 450 kA/m (5.6 kOe)with an equivalent stress value of 12.5 MPa. The“pseudoplastic” region is defined as that where thestrain changes with no additional stress due to theinduced phase transformation, as shown in Fig. 15.

Khovailo et al. [73] studied the entropy changeat the martensitic transformation in Ni2MnGa.They investigated �S between the high tempera-ture cubic phase and the low temperature tetragonalphase. The lowest value of �S was obtained for

Fig. 15. Compressive stress v. compressive strain curves at20C for Ni2MnGa after Pasquale et al. [72] showing the pseudoplastic regime once the stress exceeded12 MPa (unfilled circlesand unfilled squares). The filled triangles and filled diamondsshow portions of the stress-strain curve obtained dduring mag-netomechaincal characterization, in which the sample was sub-jected to a 20MPa load, aan anneal at 200C and then a resto-ration of the load to 12 MPa.

the stoichiometric alloy composition. Thereafter,as shown in Fig. 16, �S was found to increase withincreasing nickel content beyond the stoichio-metric composition. This is presumably due to anincrease in the magnetic contribution to theentropy change.

4.2.3. Magnetostructural phase transformationalloys

Samolyuk and Antropov [74] studied the mag-netostructural phase transformation inGd5(Si1�xGex)4. They investigated the stability ofthe magnetic structures theoretically. They showedthat the effective Heisenberg exchange interactiondecreased with deformation. This decrease inexchange coupling gave rise to the first order mag-netostructural phase transition.

Casanova et al. [75] investigated the change inentropy at the first order magnetoelastic phase tran-sition in Gd5(Si1�xGex)4. This was performed usingdifferential scanning calorimetry under the simul-taneous action of a magnetic field. It was shownthat the variation of M with both field H and tem-perature T in each phase outside the immediate“phase transition region” gave an additionalchange in entropy to add to that observed duringthe transition.

Han et al. [76] have looked at the effects of sam-ple purity on the thermal expansion of polycrystal-

Fig. 16. Entropy change �S at the martensitic phase trans-formation in Ni2+xMn1-xGa shape memory alloy as a functionof the nickel content, after Khovailo et al. [73].


line Gd5(Si1�xGex)4 at the magnetostructural phasetransformation temperature of 270 K. Measure-ments were taken on two samples made from dif-ferent purity gadolinium. Thermal expansionresults confirmed the existence of a first orderphase transition with a discontinuity of 8000 ppmin the curve of strain versus temperature at theCurie point. It was found that the Curie tempera-ture was lower for the lower purity sample, but thestrain was larger in the lower purity material.

Han et al. [77] have also looked at the magneto-elastic properties of single crystal Gd5(Si1�xGex)4

alloys. The order-disorder transition can beinduced by either field or temperature. Measure-ments were made for the first time of thermalexpansion on single crystal specimens. Resultsshowed a hysteresis in Tc of about 2 degreesdepending on whether the temperature was beingincreased or decreased, as shown in Fig. 17.Results also showed an unusual field dependenceof the transition temperature indicated in Fig. 18.Jiles et al. [78] have also looked at extraordinarymagnetoelastic coupling in single crystal speci-mens of these Gd5(Si1�xGex)4 materials. Theeffects of application of a magnetic field on theCurie point transition temperature was demon-strated and the transition temperature was found to

Fig. 17. Hysteresis in the Curie temperature of Gd5Si2Ge2 asseen from thermal expansion measurements, after Jiles et al.[78].

Fig. 18. Unusual dependence of the Curie temperature onapplied magnetic field in Gd5Si2Ge2 , after Jiles et al. [78].

be linearly dependent on the magnetic field with aderivative of about 5K/T.

Han et al. [79] have found another surprisingresult in this material concerning the field depen-dence of the Curie temperature. It was shown thatTc is different when the magnetic field is appliedalong different crystallographic axes and the differ-ence in the values of Tc along the different direc-tions depends on the field strength.

4.2.4. Iron gallium alloysClark et al. [80] investigated extraordinary mag-

netoelasticity and lattice softening in bulk, singlecrystal bcc Fe–Ga alloys. Both Al and Ga havehigh solid solubility in Fe. As the Ga content isincreased the material retains the local bcc struc-ture and the saturation magnetostriction increases.The single crystal magnetostriction coefficient λ100

exhibited two peaks as the Ga content wasincreased, showing 265 × 10�6 at 19% Ga and 235× 10�6 at 27% Ga as indicated in Fig. 19. The othersingle crystal magnetostriction coefficient λ111 wasfound to be much smaller than λ100. and thereforedid not play a significant role in the compositiondependent behavior of the bulk magnetostriction.The variation of the magnetostriction as a functionof Ga content was interpreted in terms of the vari-ation of the magnetoelastic coupling coefficient b1

and a strongly temperature dependent shear modu-


Fig. 19. Variation of the magnetostriction of Fe100-xGax alloyswith composition x as described by Clark et al. [80].

lus c11–c12. Surprisingly the magnetoelastic coef-ficient b1 showed only one peak as the Ga concen-tration increased.

Cheng et al. [81] has also investigated the struc-ture and properties of melt spun FeGa mag-netostrictive materials with different alloy compo-sitions. The objective of the investigation was totry to retain the disordered phase at higher Ga com-positions by melt spinning of the alloy. Alloys ofthe type Fe1�xGax in the range (x = 0.17, 0.21, 0.25,0.3) were rapidly quenched to form ribbons. It wasfound that the Curie temperature Tc of Fe79Ga21

was 681 °C with saturation magnetostriction of ls

= 98 ppm, whereas the Tc of Fe83Ga17 was 715 °Cand ls = 130 ppm. Magnetostriction was measuredalong the ribbon length. The addition of Ga to ironincreases l100 from 20 × 10�6 to 207 × 10�6 at acomposition of Ga ~ 17–19%. Above this concen-tration a structural change occurs in the materialand the value of magnetostriction decreases.

Kellogg et al. [82] studied the dependence ofsaturation magnetostriction on texture and grainmorphology in polycrystalline Fe83Ga17. The FeGaalloys exhibit magnetostrictive strains of �100ppm in polycrystalline form whereas for [100] ori-ented materials the bulk strain can be up to 300ppm. Control of crystallographic texture can beachieved through rolling deformation in order tomaximize the resulting magnetostrictive strain.

However the results of the measurements disagreedwith the predictions of finite element model calcu-lations (37 ppm was observed, whereas 157 ppmwas predicted).

5. Magnetoelectronic materials

Magnetic materials can be used in informationtechnology either as the medium for storing thedata or as the sensors needed for detecting themagnetized regions on the medium. To make thisof any practical use it must be possible to storelarge amounts of data in as small a space as poss-ible. Consequently there is a drive to improve datastorage densities, and currently these are doublingin less than a year. The ultimate goal is to establishfully integrated spintronics in which the devicesthat are selected for their magnetic functions arean integral part of the structure instead of an “addon” needed to achieve the desired functionality.Once this is achieved a new class ofmagnetoelectronic/spintronic materials and devices[83] will emerge which will offer a range ofenhanced possibilities, including the prospect of allmetallic transistors with very high charge carrierdensities as a replacement for semiconductor tran-sistors, and significant improvements in data stor-age densities.


5.1.1. Hysteresis: saturation, remanence andcoercivity

The hysteresis of magnetization versus magneticfield in ferromagnets and ferrimagnets can be usedfor data storage. As a result of hysteresis the rem-anent magnetization of the medium acts as a mem-ory of the last field maximum. Therefore data,either in digital or analog form can be stored asmagnetic ‘ imprints’ on magnetic media. Therequirements are that the magnetic material musthave high saturation magnetization to give as largea signal as possible. The coercivity must be suf-ficient to prevent unanticipated erasure, but smallenough to allow the local regions for the materialmaterial to be easily remagnetized when necessary.Coercivities in the range of 20–100 kA m�1 are


common for magnetic recording media, although inthe future in order to continue to improve storagedensities the coercivity of the media will need toincrease further. Therefore materials with coercivi-ties in the range 200–240 kA m�1 (2.5–3 kOe) willbe used [84]. For magnetic materials used in readand write heads the requirements are different.These materials have soft magnetic propertyrequirements to allow the magnetic flux density inthe read head or the write head to be as high aspossible for a given field strength. Therfore per-meability should be as high as possible, while coer-civity and remanence should be as low as possible.

5.1.2. MagnetoresistanceThe electrical resistance of some magnetic

materials can change quite dramatically when amagnetic field is applied. Normally this effect ismeasured in terms of the fractional change inresistance �R/R versus magnetic field. In the caseof materials such as permalloy �R/R can be about2% at magnetic saturation.

Giant magnetoresistance (GMR) occurs in somemultilayers when the relative orientations of themagnetic moments are changed. The change inresistance usually amounts to a decrease when astrong enough magnetic field is applied. If themagnetic moments in the successive layers are par-allel the material has relatively low resistance, butwhen the moments are antiparallel the material hasrelatively high resistance. This is due to differencesin the scattering of the electrons depending onwhich of the two polarizations they have, asexplained below.

The first layered magnetic material to exhibitthis ‘giant magnetoresistance’ was Fe/Cr/Fe [85],as shown in Fig. 20, and this was soon followed byother materials including Fe/Au/Fe [86], Co/Au/Co[87] and a range of other layered structures includ-ing Co/Cu/Co [88]. The magnitude of the effectdepends on several factors including the thicknessof the layers. The physical effect that causes thischange in resistance with field is the spin-depen-dent scattering of electrons. This means that whenmoving between the magnetic layers electrons withspins in one direction are scattered more than elec-trons with spins in the other direction [89]. Theband theory of magnetism provides an explanation

Fig. 20. Giant magnetoresistance �R/R in Fe/Cr/Fe multilay-ers of different thicknesses, showing up to 45% decrease in anfield of 2 Tesla, after Baibich et al. [85].

of this effect as described by Fert et al. [90]. Thechange in resistance with change in relative orien-tation of the directions of magnetization in the suc-cessive magnetic layers can be attributed to thespin-dependent conduction of the electrons in theseitinerant electron ferromagnets. The majority(“spin up” ) and the minority (“spin down” ) elec-trons have different densities of states andmobilities at any given energy due to the exchangesplitting of the energy bands [91] as indicated inFig. 21. Consequently the conductivities of thesetwo groups of electrons can be different. Thereforeelectrons with a particular energy and spin orien-tation which pass from a layer where the magnetiz-ation, and hence the majority spins, are pointing tothe right, to a layer where the magnetization, andhence the majority spins, are pointing to the left,experience a dramatic change in their conductivity.

In situations where the electron mean free pathis larger than the thickness of the layers, in theantiparallel magnetization configuration all elec-trons which move across the layer boundaries findthat they are strongly scattered in alternate layers,whereas in the parallel configuration one group ofelectrons will be weakly scattered in both magneticlayers, while the other group will be strongly scat-tered in both magnetic layers. The resistivity is


Fig. 21. Schematic of the electron band for a tunnel junction with the magnetizations in both electrodes parallel (upper figure) andantiparallel (lower figure), as described by Leib et al. [91].

therefore much lower in the parallel configurationthan in the antiparallel configuration.

This change of resistance when subjected to amagnetic field can be used to develop high resol-ution sensors, such as read heads, which can detectthe magnetic state of regions on the magneticmedium. Rapid progress in storage densitiesrecently has been due principally to the develop-ment of giant magneto-resistive multilayers, whichhave led to the production of a new generation ofhigh-sensitivity, high spatial resolution read heads[92]. These magnetic multilayers can exhibitchanges in resistance �R/R of 80%.

The magnetoresistive material is used to detectthe leakage flux from the locally magnetized

regions of the recording medium by means of atwo-point resistance measurement. To obtain anantisymmetric response that is approximately lin-ear with strength of the leakage field, the magneto-resistive material also needs to be biased withanother constant applied field. This is often sup-plied by using an adjacent magnetized film.

5.2. New magnetoelectronic materials

5.2.1. Magnetoresistive spin valvesGMR devices can be constructed in which one

of the ferromagnetic layers is strongly pinned in aparticular direction through exchange coupling toan additional nearby antiferromagnetic layer. In


this way an exchange biased magnetic electrodecan be used to create what is known as a “spinvalve” whereby the magnetization of one layer (thefree layer) can be more easily oriented by appli-cation of a magnetic field than the magnetizationof the other layer (the pinned layer). Thesematerials have high magnetoresistance and lowresistance that make them suitable for read headsensors.

The spin valves can employ either current in theplane “cip” or current perpendicular to the plane“cpp” configurations. Bass and Pratt [93] have dis-cussed structures that employ current perpendicularto plane “cpp” magnetoresistance because the“cpp” magnetoresistance is usually larger than the“cip” magnetoresistance, as shown in Fig. 22, andmany layered magnetoresistive structures can beoperated in either mode. New directions of investi-gation for these materials include ballistic trans-

Fig. 22. CPP (upper figure) and CIP (lower figure) magnetore-sistances for a Co/Ag/Co multilayer at 4.2 K after Bass andPratt [93], showing that the cpp magnetoresistance is larger.

port, in which the free path length of the electronsis greater than the dimensions of the device.

Generation and control of spin polarized cur-rents in devices such as spin valves has been dem-onstrated in a metallic mesoscopic spin valvedevice by Jedema et al. [94] in which ferromag-netic electrodes of permalloy (Ni–Fe) were used todrive a spin polarized current into copper. Thisallows control of spin polarized currents in solidstate devices at room temperature through the useof ferromagnetic contacts and marks a significantdevelopment towards “all metallic” electronicdevices, such as diodes and transistors, as an alter-native to conventional semiconductor devices. Inthese all-metallic spin valve structures the resist-ances are lower than in the tunnel junctions andthe fractional change in resistance is about 10%.The all-metallic devices have the possibility ofhigh storage densities, which are comparable tomagnetic tunnel junctions. However because of thelower resistance of the all metallic structures theyhave low read out voltages. In addition a slowresponse time means longer read/write times ofabout 100 ns. Together these two factors have pre-vented them reaching performance levels that areacceptable for magnetic random access memory(MRAM) applications. Therefore these devicesseem to have better prospects for sensors than fordata storage.

Schumacher et al. [95] has investigated the pre-cessional magnetization reversal in microscopicspin valve cells. The main interest is to enablerapid reversal of the magnetization in these struc-tures. Possible uses in ultra fast, low power mag-netic memory cells such as MRAM are envisagedwhere the data can be written rapidly. The rapidreversal of magnetization was achieved by theapplication of field pulses along the hard axis.Short, intense field pulses were applied with thedirection oriented perpendicular to the initial mag-netization direction of the domains. This resultedin stable and reversible switching of the magnetiz-ation direction with switching reliability greaterthan 99.8%.

Spin valves in which the pinned or free layerhave been oxidized have shown enhancement ofthe magnetoresistance over conventional unoxid-ized spin valves as described by Veloso et al. [96].


A theoretical study of the anomalous magnetore-sistance behavior of these oxidized spin valvesusing a modified Stoner-Wohlfarth approach, hasbeen performed by Ventura et al. [97] to explainthe behavior. In particular they investigated thefield dependence of the magnetoresistance at dif-ferent temperatures of CoFe spin valves consistingof a structure with MnIr/CoFe/nano-oxidizedCoFe/CoFe/Cu/CoFe/Ta. They have extended thetheoretical understanding of a spin valve by separ-ating the pinned layer into two sublayers which areferromagnetically coupled across an intermediatenano-oxide layer. The results showed that thebehavior of these materials, including the anomal-ous behavior of the magnetoresistance at low tem-peratures, can be explained in terms of the strengthof the ferromagnetic coupling between the upperand lower magnetic sublayers of the pinned layer,and the coupling between the antiferromagneticMnIr layer and the bottom magnetic sublayer ofthe pinned layer, both of which control the reversalin magnetization of the pinned layer. Changing theapplied field from the negative maximum at whichparallel alignment occurs first causes the reversalof the pinned layer. This occurs gradually in stagesbecause of the composite structure of the pinnedlayer, until high resistance antiparallel alignmenthas been established. As the field is increased tosmall positive field strengths the free layer sud-denly switches from parallel to antiparallel con-figuration, resulting in a discontinuous reductionin magnetoresistance.

In order to improve the performance of spinvalve devices Yuasa et al. [98] have studied theeffect of inserted Cu on the current perpendicularto plane GMR in Fe50Co50 spin valves formed bydeposition of multilayers of Cu and FeCo. Thesolid solubility of Cu in the Fe50Co50 magneticelectrode material is an important factor in increas-ing the spin dependent scattering. The result is thatthe change in resistance area product of the spinvalve is enhanced by the use of the Cu layer, typi-cally 5nm thick between the free and the pinnedlayers.

5.2.2. Magnetic tunnel junctionsThese devices have similarities to spin valves

and that most device designs are analogous to those

using spin valves. These devices also use the cur-rent perpendicular to plane (cpp) configuration.The tunneling current that passes across the junc-tion also differs depending on the relative orien-tations of the magnetization vectors in the twomagnetic layers, as shown in Fig. 23. In magnetictunnel junctions the magnetoresisitivities arehigher than in GMR and spin valve devices, whichis due to the use of an intermediate insulating layerinstead of a metallic layer. These devices are usefulfor both data storage and sensors, in which theirmain characteristics are: large signal to noise ratios(sensors) and non-volatility (data storage).

Moodera et al. [99] developed the first magnetictunnel junctions. Magnetic tunnel junctions in theirsimplest form consist of a three layer sandwich oftwo magnetic layers separated by a thin, 0.5–1.0nm, non-magnetic, non-conducting layer. Mooderaand Mathon [100] have reviewed spin polarizedtunneling in ferromagnetic junctions. Tunnelingmagnetoresistance values of up to 30% have beenachieved repeatably. Spin polarization (n↑�n↓) / (n↑

+ n↓), where n↑ is the number density of “spin up”electrons and n↓ is the number density of “spindown” electrons, within the conduction band of themagnetic electrodes is an important factor indetermining the tunneling magnetoresistance sincethis gives the direct measure of the excess popu-

Fig. 23. Variation of resistance with magnetic field for aCo/Al2O3/Ni80Fe20 tunnel junction at room temperature 295Kand at liquid nitrogen temperature 77 K, after Moodera andMathon [100].


lation of one spin type over the other. This has avalue of 50% in a material such as Fe50Co50.

Tunnel junctions have several attractive featuresincluding the possibility of fabricating MRAMwith extremely high data storage densities [101].Random access memories (RAM) that are cur-rently in use for data storage in computers arebased solely on semiconductor technology andtherefore data that is stored in these memories islost when the power supply is interrupted. On theother hand magnetic memories offer the prospectof non-volatile RAM whereby the magnetic statesof the memory bits are not altered by the voltageinterruptions and therefore data is preserved [102].Prototype MRAM devices have already beenshown to have improved on the speed and powerconsumption of conventional non-volatile sem-iconductor memory by several orders of magni-tude, and have reached data storage densities andspeeds that are comparable with those of conven-tional volatile dynamic RAM.

Reiss et al. [103] have reported on the propertiesand applications of magnetic tunnel junctions, inparticular the strong dependence of tunneling cur-rent on external field at constant applied voltage.Device sizes that have been scaled down to below0.01 mm2 are of interest in order to achieve highdata storage densities and high spatial resolutionsensors. Considerations of single domain behavior,switching properties and edge coupling effects areimportant in trying to understand the behavior ofthe magnetization of the material at these lengthscales.

For MRAM type applications the mostimportant consideration is a “clean” rapid reversalof magnetization. Size and aspect ratio of thedevices play an important role in determining theswitching field, as shown in Fig. 24. Warot et al.[104] have studied the magnetic properties of pat-terned tunnel junctions with a view to both MRAMand read head sensor applications. Magnetizationreversal was studied using Lorentz microscopy.Several energy terms need to be considered inorder to understand the properties. These includeNeel coupling, stray field magnetostatic couplingand shape anisotropy. It was found that the highsaturation magnetization in NiFeCo was princi-pally responsible for the difference in behavior

Fig. 24. Switching fields for MRAM tunnel junctions, com-prising of a multiplayer of Ta/NiFe/MnFe/NiFe/Al2O3/NiFeCoas a function of device aspect ratio for three different widths:a) 0.7 mm, b) 1 mm and c) 2 mm after Warot et al. [104]. P-AP represents the parallel to antiparallel switching field, AP-Prepresents the antiparallel to parallel switching field.

between tunnel junction fabricated with it andsimilar tunnel junction fabricated with NiFe freelayers.

Nozaki and Matsuyama [105] have consideredsize dependence of switching current and energybarrier in the magnetization reversal of MRAMcells. To do this they investigated rectangular cellshapes. The switching current I and the energy bar-rier �E show strong dependence on width w andthickness t of the cells as indicated in Fig. 25. Iwas found to be proportional to t/w, while �E wasproportional to wt2. A sufficient energy barrieragainst thermal activation is essential in the practi-cal applications of these materials to MRAM. This


Fig. 25. Calculated variation of switching current Iw andenergy barrier �E with width w and thickness t of a perrmalloyMRAM cell after Nozaki and Matsuyama [105].

amounts to �E � 80 kBT. Larger values of t areneeded because w is being reduced to increase datastorage density. However larger values of t alsoincrease the switching current, which is undesir-able. Optimization of aspect ratio of the cellresulted in length to width ratio of l/w ~ 1.5 anda width of w ~ 0.1 µm which gave a data densityof 5 Gbit/sq.in (0.8 Gbit/cm2), for thermal stabilitya thickness of t � 3.5 nm was needed.

Junction reliability is related to the electricalbreakdown mechanism. Kim et al. [106] studiedthe electrical breakdown in tunnel junctions andhow this depended on the area of the junction. A“constant voltage stress” was applied to tunneljunctions with cross sectional areas of 200 µm2 and0.5 µm2. The mechanism of the electrical break-down appeared to be quite different depending onthe area of the junction. For large area junctionsthe MR decreased gradually with time, as shownin Fig. 26, while for small area junctions the MRdecreased abruptly with time. The breakdown volt-age was found to be a function of the junction bar-rier thickness. Higher breakdown voltages were

Fig. 26. Magnetoresistance curve of a 200 mm2 tunnel junc-tion (upper figure) and variation of magnetoresistance (MR) andresistance area product (RA) with time at a constant appliedvoltage of 0.75 V showing the progress in electrical breakdown(lower figure) after Kim et al. [106].

found for thicker barriers, and for smaller crosssectional areas. The breakdown voltage alsodepended on the barrier layer oxidation time,which relates to the junction barrier quality. Timeneeded for electrical breakdown to occur underconstant voltage stress was found to be shorter forlarger area junctions, because the process is depen-dent on defects such as pinholes, process induceddefects and side wall damage, all of which increasewith area. For small area junctions no changeseemed to occur in MR prior to electrical break-down.

Low resistance tunnel junctions of a few � mm2

are needed for read head sensor applications. Wanget al. [107] have investigated different materials forcontinuous thin barriers in order to produce lowresistance spin dependent tunnel junctions as


shown in Fig. 27. However with thinner barrierlayers the probability of “pinholes” increases andtheir presence in the insulating layer degrades theperformance of tunnel junctions, including the tun-neling magnetoresistance. Various different oxidebarriers were studied, including HfOx, HfAlOx,ZrAlOx and AlOx. It was found that HfOx madethe thinnest continuous insulating barriers. Aresistance area product of 0.4 � mm2 was obtainedwith a 0.4 nm HfOx barrier with 5.5% tunnelingmagnetoresistance. A resistance area product of0.65 � mm2 was obtained with a 0.2 nm Hf + 0.3nm AlOx barrier with 9% tunneling magnetoresist-ance.

Neel (“orange peel” ) coupling between thinfilms is caused by magnetostatic interactions acrossthe interface arising from free poles generated onthe surface when it is corrugated. Surface rough-ness together with high magnetization in the sur-face layer is the root cause of Neel coupling. Thisis undesirable in magnetic tunnel junctions becausethere is a shift, known as the bias field, in theresponse of the free layer to an applied magneticfield. Particularly in the case of low field devicesand applications it is highly desirable to reduce oreliminate Neel coupling. Wang et al. [108] havelooked at spin dependent tunnel junctions with anew structure for the pinned layer giving reduced

Fig. 27. Tunnel magnetoresistance and resistance area productfor tunnel junctions constructed with different barrier layermaterials, after Wang et al. [107].

Neel coupling between the free and pinned layers.They used a five layer modified synthetic antiferro-magnetic composite as the pinned electrode, withthree layers of ferromagnet and two intermediatelayers of ruthenium. The new structure that hasbeen developed gives a strong antiferromagneticcoupling between the magnetic layers in the com-posite pinned electrode. This antiparallel magnetiz-ation configuration is maintained under normaloperating conditions of the tunnel junction. The netbulk magnetic moment of the composite pinnedelectrode is thereby kept close to zero to minimizethe coupling even in the presence of surface rough-ness. The new structure reduced the bias field toless than 80 A/m (1 Oe) as shown in Fig. 28.

Fig. 28. Magnetic hysteresis loop for magnetic tunnel junctionusing conventional structure (upper figure) and for the newmodified synthetic antiferromagnet composite pinned layer(lower figure). This shows the reduction in the bias or offsetfield due to reduction in Neel coupling, after Wang et al. [108].


5.2.3. Nanoscale magnetic switching materialsFuture high density recording media appli-

cations with data storage densities in the range of1012 bits/sq.in will require the development ofmaterials for nanoscale magnetic switching struc-tures. Nanodots in the single domain state withwell controlled anisotropy and low hysteresis havepotential for i) high resolution field sensors and ii)high density low power magnetic logic gates. Multidomain structures are undesirable from the view-point of data storage, but the formation of multido-mains is suppressed in very small structures, whichis therefore doubly advantageous. Cowburn [109]formed nanometer scale magnets which have dif-ferent magnetic properties from the bulk magnetsformed of the parent bulk material. These nanos-cale magnets exhibit suppression of multidomainconfigurations and configurational anisotropy.

Configurational anisotropy is a form of ani-sotropy that arises from the details of the magneticmoment configuration in a small structure. Thisanisotropy is the difference in energies betweenstates with different deviations of the magneticmoment configuration from complete uniformity.It is different from shape anisotropy, but has somesimilar characteristics. For example the energy ofperfectly uniformly magnetized square is inde-pendent of the direction of magnetization, butimmediately the non-uniformity of magnetizationis allowed energy differences arise between the dif-ferent directions of magnetization. It was foundthat the competition between exchange and demag-netizing energy is the determining factor in decid-ing the actual magnetic moment configuration inthese nanostructures, and slight differences in theseconfigurations can lead to very different depen-dence of magnetization on applied field Thereforeattention has to be given to geometry, even at thesesmall scales, in order to obtain the desired mag-netic moment configuration.

Fidler et al. [110] have modeled fast switchingbehavior in square nanoelements 10 × 10 × 1 nm,both with and without anisotropy. The modeling ofswitching dynamics of these structures was perfor-med using a combination of finite element mode-ling and micromagnetic Landau-Lifschitz-Gilbertmodeling. Switching time and critical field werefound to depend strongly on the Gilbert damping

parameter and the field sweep rate. All other fac-tors being equal the longest switching times werefound in the elements with uniaxial anisotropy,while the shortest switching times were in materialwithout anisotropy.

Schrefl et al. [111] have modeled coercivity andremanence in annealed and monodisperse selfassembled nanoparticles of FePt. These “as synthe-sized” particles have a disordered fcc structure.Heat treatment can cause a phase transformation at530 °C which results in an ordered face centeredtetragonal structure with high anisotropy of typi-cally 106 J/m3 and a coercivity that can reach 716kA/m. Model calculations were made using amodified Stoner-Wolfarth model (to allow forinteractions between the particles) and these resultswere compared with a combined finite elementmodel together with a Landau-Lifschitz-Gilbertmicromagnetic calculation. It was found that rem-anence enhancement occurred only when theexchange interaction was reduced, as indicated inFig. 29, in contrast to the usual result where reman-ence increased with exchange interaction.

Fig. 29. Calculated variation of reduced remanence (ratio ofremanence to saturation) and coercivity of FePt nanoparticleassemblies with exchange interaction for different volume frac-tions of the disordered fcc phase after Schrefl et al. [111].


6. Conclusions

In recent years there has been an explosion ofinterest in new magnetic materials. These materialsprovide a fertile subject area for scientific investi-gations which transcend the boundaries of con-densed matter physics, materials science and elec-trical and computer engineering. They also havecurrent and future technological applications indiverse areas such as energy conversion devices,field generation, sensors, actuators, data storageand memory devices.

In this review recent scientific and technologicaladvances have been described in four importantcategories of magnetic materials that are currentlyof high interest. These include: soft magnets, hardmagnets, magnetomechanical and magnetoelec-tronic materials. It is seen that the requirements ineach of these areas are quite different and yet thereis also much common ground, so that an under-standing of one of these groups of materials inevi-table leads to a better understanding of the othergroups. Although the range of magnetic materialswill continue to diversify in the future the samebasic concepts described here will remain essentialto the understanding and future development ofmagnetic materials.

Improvements and completely new develop-ments in the various types of magnetic materialshave been highlighted through the consideration ofa few selected new magnetic materials which areat the leading edge of current research and devel-opment, including: amorphous magnetic fibers,nanocrystalline permanent magnet materials, ferro-magnetic shape memory alloys and spintronicmaterials. For each type of material the variousmagnetic properties of prime interest have beendiscussed and how these properties can be con-trolled and improved. In most cases it can be seenthat in the future the greatest progress is likely tobeing made through emphasis on control of thematerials on the smallest scale that is possible.Currently this emphasis is on “nanostructured”materials. The interesting result is that the bound-ary between engineering applications and funda-mental science is beginning to be erased as the twofields of activity meet on the scale of nanostruc-tured materials. For the development of better mag-

netic materials in the future, the emphasis will shiftdown to an even smaller scale.

Acknowledgements

This was supported by the US Department ofEnergy, Office of Basic Energy Sciences, MaterialsSciences Division. Ames Laboratory is operated byIowa State University for the US Department ofEnergy under Contract No.W-7405-ENG-82.

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