Buschow K.H.J. (Ed.) Concise Encyclopedia of Magnetic and Superconducting Materials

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where l

100

and l

111

are the magnetostriction con-

stants reﬂecting the elastic property of l, m, and m

Here, Einstein’s summation rule is adopted for the

repeated subscripts as in Sect. 2.1. In order to deter-

mine g

ðxÞ, the total energy must be minimized and

this is written as

½f

þ f

dv ¼ 0 ð11Þ

where the integration is taken over the specimen.

Under a high ﬁeld of near saturation, since directions

of the M

are nearly parallel to applied magnetic

ﬁeld, the partial differential equations derived from

Eqn. (11) can be solved by means of Fourier trans-

formation under a linear approximation (Seeger and

Kronmu

ller 1960, Kronmu

ller and Seeger 1961,

Umakoshi and Kronmu

ller 1981a). The spin distri-

bution, g

ðfxgÞ, is most conveniently visualized in

the dislocation coordinates (Seeger and Kronmu

ller

1960, Kronmu

ller and Seeger 1961, Kronmu

ller 1967,

Gessinger 1970, Umakoshi and Kronmu

ller 1981a,

Yasuda et al. 2000). This is exempliﬁed in Fig. 3 for

an edge dislocation (after Kronmu

ller 1967). Here,

the deviation of g

ðfxgÞ from the direction of the

magnetic ﬁeld applied along b is shown by arrows.

A method to obtain u

ðfxgÞ and s

ðfxgÞ around a

dislocation in an anisotropic material is given in the

literature (Stroh 1958, Nabarro 1967, Hirth and

Lothe 1968, Mura and Mori 1976, Mura 1987). For

simplicity though, m

is set to zero in most of the

calculations of s

ðfxgÞ used to analyze the crystallo-

graphic characteristics of magnetism. Since disloca-

tions may pile up against obstacles, the equilibrium of

linear arrays of dislocations must be known and this

problem was solved by (Eshelby et al. 1951).

(a) B.c.c. Fe

Plastic deformation in b.c.c. transition metals and in

some covalent crystals is controlled by the motion of

screw dislocations that overcome a large Peierls–

Nabarro stress (Nabarro 1967) by thermally acti-

vated nucleation of double kinks (Celli et al . 1963,

Dorn and Rajnak 1964). Such a motion is assisted by

dispersed misﬁt strain centers at low temperatures,

whereas the motion of edge portions is opposed by

the same strain centers at high temperatures (Sato

et al. 1970, Sato and Meshii 1973). This explains why

TEM observations show that the dislocation struc-

ture strongly depends on deformation temperature

(Foxall and Statham 1970).

Another characteristic feature to be noted in plastic

deformation is that as dislocations are generated at

some multiplication sources (Read 1953), they tend to

pile up against obstacles. In the b.c.c. Fe, dislocation

interaction with other obstacles takes over the double

kink nucleation mechanism at a higher temperature

so that dislocations are introduced as groups. In

general, the susceptibility under a high magnetic ﬁeld

is proportional to Bb

N where N is the total number

of dislocations but this is modiﬁed by dislocation

pileups as follows. In a pileup containing n disloca-

tions, b is effectively increased to nb and N is

decreased to N=n, thus changing the proportionality

factor into ðnbÞ

ðN=nÞ¼nb

N. Consequently, w

increased by a factor n (Kronmu

ller and Seeger 1961,

Kronmu

ller 1967, Gessinger 1970, Willke 1972,

Schroeder 1978, Umakoshi and Kronmu

ller 1981a,

ltz et al. 1986).

This property has been utilized to identify the

number of dislocations contained in pileups. Accord-

ing to the orientation dependence of the high-ﬁeld w

measured on a disk sample cut either parallel to the

slip plane or perpendicular to the Burgers vector,

Schroeder (1978) showed that n ¼ 16217 for the pri-

mary dislocations and n ¼ 12 for the secondary dis-

locations in Fe single crystals deformed to stage III at

room temperature. On the other hand, isolated edge

dislocations are identiﬁed to be responsible for the

orientation dependence of w

after deformation to

stage I.

Umakoshi and Kronmu

ller (1981a) have examined

the changes of w

in b.c.c. Fe due to primary edge and

conjugate screw dislocations at successive stages of

Figure 3

Distribution of the spontaneous magnetization (spins)

around an edge dislocation in Ni. The magnetic ﬁeld of

H ¼ 100 Oe is parallel to the Burgers vector. The arrows

indicate directions magniﬁed along the vertical axis by a

factor of 5 (Kronmu

ller and Seeger 1961) (reproduced

by permission of Kronmu

ller (1967) from Can. J. Phys.

45, 633).

530

Magnetic Properties of Dislocations

deformation at 77 K after pre-straining at 295 K.

Figure 4 shows an example of their study. The ori-

entation dependence of w

measured on a disk sample

sliced perpendicular to the primary Burgers vector

(after Umakoshi and Kronmu

ller 1981a) shows

asymmetry due to introduction of conjugate screw

dislocations. By comparison with the calculated con-

tribution shown in (a), it can be seen that the meas-

ured w

in (b) is contributed by all dislocation

characters. For a better understanding of the ﬁgure,

it is required to incorporate the results of other disks

sliced along different orientations (Umakoshi and

Kronmu

ller 1981a).

Long-range internal stresses due to piled-up dislo-

cations may alter the magnetic domain patterns.

Umakoshi and Kronmu

ller (1981b) have examined

the domain structures of Fe single crystals deformed

at 77 K and 195 K by using the magneto-optical Kerr

effect (Fowler and Fryer 1954). They showed that the

structure is ﬁner after deformation at 77 K than at

195 K in accordance with the operation of a double

kink nucleation mechanism at low temperatures.

Schroeder (1978) has examined domain patterns in

Fe single crystals deformed at room temperature and

explained the structure as a compromise between (i)

the stray-ﬁeld energy that forces spins to align par-

allel to each other and (ii) the Bloch wall energy that

governs spin rotation in domain walls under the in-

ﬂuence of stress ﬁeld from piled-up dislocations.

Magnetic measurements, thus, provide a wide range

of information about dislocations introduced into

ferromagnetic b.c.c. Fe by plastic deformation with-

out damaging a sample.

(b) F.c.c. Ni and Ni

In 1960s, it was questioned if the forest dislocations

contribute to the work hardening in f.c.c. metals. In

his review on the question, Kronmu

ller (1967) point-

ed out that magnetic measurements are useful as a

nondestructive testing method to delineate the dislo-

cation arrangements and to get some insight into the

work-hardening mechanism. In spite of the fact that

investigations are restricted to ferromagnetic materi-

als and that interpretation of experimental results

require some solid theoretical background, several

researchers have taken advantage of magnetic meas-

urements to identify the dislocation types, densities,

and their distribution in f.c.c. Ni in a wide range of

plastic deformation.

According to Eqn. (11), w

and H

areshowntobe

proportional to N

1=2

and N

1=2

, respectively (Seeger

et al. 1964, Willke 1972). Since the work-hardening Ds

is proportional to N

1=2

in a polycrystal (Taylor 1934) or

in stage II of a single crystal, w

and H

are propor-

tional to Ds

1

and Ds, respectively (Seeger et al. 1964),

as is observed in Ni single crystals (Hayashi 1967).

Hayashi (1967) also shows experimentally that w

1

proportional to Ds

1=2

in stage I in agreement with the

work-hardening theory (Mott 1952). Dislocation struc-

tures in deformed Ni single crystals have been studied

in more details by analyzing the high ﬁeld w

due to

edge, screw, and 601 dislocations (Gessinger 1970). The

information about dislocation distribution can also be

obtained from the domain structure by differentiating

the dislocation-induced anisotropy of f

from the

magnetocrystalline anisotropy of f

on the ground of

theoretical analysis of the high-ﬁeld w

(Willke 1972).

Figure 4

Comparison of the calculated (a) and measured (b) high-ﬁeld w shown by polar diagrams on a (111) plane

perpendicular to the primary Burgers vector. It can be noted that w conjugate Burgers vector [1

11]/2 becomes notable

in addition to the contribution of the primary screw dislocations at a large strain in (b) (reproduced by permission of

Umakoshi and Kronmu

ller (1981a) from Phys. Stat. Sol. (a) 66(2), 509–520).

531

Magnetic Properties of Dislocations

In ordered alloys, dissociated superdislocations

may also contribute to the magnetic anisotropy. In

a deformed Ni

Fe, the anti-phase boundary the

(APB) strip alters the orientation dependence of w

measured on the disks sliced along the low index

planes in the dislocation coordinates (Yasuda et al.

2000). The magnetic anisotropy induced by the APBs

differs from that due to the s

ðfxgÞ around disloca-

tions. More importantly, introduction of the APBs

alters the absolute magnitude of the M

as described

below.

3.3 Role of Atomic Rearrangements

The motion of partial dislocations affects magnetic

properties in various ways depending on crystal

structure, chemical composition, deformation tem-

perature, strain rate, etc. The Ne

el and Curie tem-

peratures are different in deformed specimens,

basically due to the change in spin–orbital interac-

tions (Kramers 1934, Van Vleck 1937, Anderson

1950). Microscopic observations, diffraction analy-

ses, and spectroscopy are useful in studying magnetic

properties that can be strongly affected by inter-

changes of atoms as exempliﬁed below for some pop-

ular intermetallics.

(a) B.c.c.-based superlattice

There are three main types of b.c.c.-based superlat-

tices as shown in Fig. 5. Plastic deformation increases

magnetization most evidently for B35 at.% Al in

Fe–Al intermetallics (Taylor and Jones 1958, Huffman

and Fisher 1967, Gengnagel et al. 1972, Takahashi

et al. 1996). In this composition, although both the

B2 and D0

phases coexist, dislocations generated in

the B2 phase control the plastic deformation as well

as magnetic properties (Yamashita et al.1998).The

local change in ordering introduced by deformation

increases the magnetization (Chikazumi et al. 1957,

Takahashi and Umakoshi 1990, Takahashi et al.

1996). Mo

ssbauer spectroscopy shows that the number

of ﬁrst-nearest neighbor (FNN) Fe–Fe pairs contrib-

utes to M

(Huffman and Fisher 1967).

In contrast, plastic deformation reduces the mag-

netization of a Heusler alloy of L2

-ordered Pd

MnSi

(Takahashi and Shinohara 1982), whose structure

sketched in Fig. 5 was ﬁrst found in a Cu–Mn–Al

system (Heusler 1934). NMR analyses of various

Heusler alloys show that antiferromagnetic and ferro-

magnetic properties come from the interaction

between atoms in the second- (SNN) and the third-

nearest neighbor positions (TNN), namely by B–C

and B–B interactions where B is Mn (Khoi et al.

1978). The motion of a partial dislocation /111S/2

on a {110} plane in the L2

structure introduces the

FNN Mn–Mn pairs at the expense of the TNN Mn–

Mn pairs reducing ferromagnetism. This is because

the Mn–Mn magnetic interaction is antiferromagnetic

when they are located close (Ne

el 1948) and ferro-

magnetic when they are located apart (Kasper and

Roberts 1956).

(b) F.c.c.-based superlattice

and D0

are the most popular structures as

shown in Fig. 6. The L1

structure originally found in

an Au–Cu system (Johansson and Linde 1925) has

gained much attention because of the good symmetry

and ductility. Although slip takes place on {111}

planes in general at a low temperature, limited mo-

tion of /110S/2 superpartial dislocations on cube

planes introduces D0

structure locally. Plastic

deformation reduces the susceptibility of Pt

Fe and

Figure 5

B2, D0

, and L2

in b.c.c.-based superlattices. The ﬁgure

contains eight unit cells for the B2 structure. A Heusler

alloy assumes the L2

structure.

Figure 6

and D0

in f.c.c.-based superlattices. The motion of

a [110]/2 superpartial dislocation on a (001) plane

induces the phase transformation L1

2 )

, which

does not involve interchanges of the FNN atoms (Flinn

1960).

532

Magnetic Properties of Dislocations

introduces spontaneous magnetization (Bacon et al.

1963, Takahashi and Umakoshi 1990). On the con-

trary, spontaneous magnetization in ferromagnetic

Mn, Ni

Al, and Ni

Ge disappears by plastic

deformation (Taoka et al. 1959, Bacon et al. 1963,

Takahashi and Ikeda 1983, Izumi 2003). This subject

will be discussed more in Sect. 3.4(b).

3.4 Contribution of APB Tubes

Vidoz and Brown (1962) have introduced the concept

of APB tubes to explain the work hardening in su-

perlattices. APB tubes were then observed and dis-

cussed by many researchers in both B2 (Crawford

1976, Chou and Hirsch 1981, 1983, Yamashita et al.

1998, Yamamoto et al. 2003) and L1

alloys (Chou

et al. 1982, Shi et al. 1996, Murakumo et al. 2000a,

2000b). The APB tubes disappear upon annealing at

a relatively low temperature, thus providing a simple

means to discriminate their contribution to the mag-

netization (Yang and Baker 1997, Yamashita et al.

1998, Yang et al. 1999) from that due to APB ribbons

bound by superpartials within a dissociated super-

dislocation (Takahashi et al. 1996, Takahashi and

Umakoshi 1990).

APB tubes can be formed in any given ordered

alloy whenever the leading and trailing partials that

belong to a perfect dislocation are forced to move

on different slip planes, and this may occur locally

at jogs or at kinks. It is important to notice that

APB tubes are free from dislocations except at their

ends. Another point to be recalled is that change in

ferromagnetism is essentially determined by spin–

orbital interactions (Kramers 1934, Van Vleck 1937,

Anderson 1950) related strictly to the atomic ele-

ments at APB tubes and APB ribbons in the ordered

intermetallics. Formation of the APB tubes is favored

by high deformation rates (Murakumo et al. 2000b,

Yoneyama 2002) as well as by cyclic deformation

(Yamamoto et al. 2003, Yasuda et al. 2000, 2003,

Izumi 2003). Since the latter case is regarded as heavy

deformation in the sense of accumulated strain, it

is speculated that APB tubes might have also been

introduced at high densities in crashed Fe–Al

(Gengnagel et al. 1972) and in ﬁled Pt

Fe (Bacon

et al. 1963).

(a) APB tubes in b.c.c.-ordered superlattices

The introduction of APB tubes by plastic deforma-

tion in an Fe–Al alloy increases magnetization sub-

stantially (Yang and Baker 1997, Yamashita et al.

1998, Yang et al. 1999). Furthermore, M

increases

markedly by cyclic deformation as shown in Fig. 7

(after Yamamoto et al. 2003). Here, the cyclic defor-

mation corresponds to 33.7% of accumulated strain

achieved by changing the mutually orthogonal com-

pression axes. A constraint was imposed on the

side surfaces to limit the operative shear systems

(Murakumo et al. 2000a) and to hold the specimen

dimension unchanged.

The observed evolution of H

can be attributed to

the introduction of APB tubes. The small value of H

is reasonable since APB tubes are rather narrow

without accompanying long-range stress. H

is ap-

proximately the same as that in Fe

Ge, but substan-

tially larger than that in Ni

Ge, as shown later in

Figs. 9 and 11. The APB tube density is several times

higher than the dislocation density as shown in Fig. 8

(after Yamamoto et al. 2003). M

decreases by 80%

after annealing at 573 K for 30 min, which, as shown

by TEM observations, coincides with the disappear-

ance of APB tubes (Yamamoto et al. 2003).

(b) APB tubes in f.c.c.-ordered superlattices

In contrast to what is observed in B2-ordered Fe–Al,

plastic deformation reduces M

markedly in Ni

cyclically deformed at 4.2 K as shown in Fig. 9 (after

Izumi 2003). Superpartial dislocations are forced to

glide on both {111} and {001} planes as a result of

compression being applied along [100] and [011] axes

alternatively. Generation of a high density of APB

tubes by cyclic deformation is illustrated in Fig. 10,

conﬁrming that introduction of APB tubes is the

main origin of the reduction in M

(Izumi 2003).

At this point, it is interesting to compare the above

result on Ni

Ge with that on Ni

Al (Takahashi and

Ikeda 1983) that shows a similar decrease in M

4.2 K. In the latter, M

decreases proportionally to

the strain in the early stages and to the square root of

Figure 7

Strong magnetization in B2-ordered Fe–Al induced by

cyclic deformation to an accumulated strain of 33.7%.

Note that the small value of H

is about the same as that

in deformed and nondeformed Fe

Ge shown later in

Fig. 11 (Yamamoto et al. 2003).

533

Magnetic Properties of Dislocations

the strain at higher strains. It is known that APB

tubes are also formed in Ni

Al (Chou et al. 1982, Shi

et al. 1996, Murakumo et al. 2000b) and it is likely

that they contribute to the ferromagnetic to para-

magnetic transition through changes in environmen-

tal atoms as discussed for APB ribbons by Takahashi

and Ikeda (1983).

It should be emphasized that plastic deformation

does not reduce M

in Fe

Ge at 300 K as shown in

Fig. 11 (after Taniguchi 2003). The effect of plastic

deformation on Fe

Ge, thus, differs from that ob-

served in Ni

Ge and Ni

Al, despite that they are all

ferromagnetic at the temperatures examined. The

origin may be ascribed to the difference in activated

Figure 9

The spontaneous magnetization in Ni

Ge measured at

4.2 K decreases notably by cyclic deformation given at

77 K to accumulated strain of 10.0% (Izumi 2003).

Figure 8

APB tubes introduced into the deformed Fe–Al of

Fig. 7. The beam direction is nearly along [01

1] and the

dark-ﬁeld image is taken under the 111 order spot

(Yamamoto et al. 2003).

Figure 10

APB tubes introduced into the deformed specimen of

Fig. 9. Irradiation-induced secondary defects are formed

preferentially along the APB tubes which facilitates the

visualization of the latter (Murakumo et al. 2000). The

inset shows the 111 ordered spot used in the dark-ﬁeld

image. The beam direction is taken nearly along [01

(Izumi 2003).

Figure 11

Magnetization curves obtained in nondeformed and

cyclically deformed Fe

Ge at room temperature. Plastic

deformation at room temperature does not change

either M

or H

in visible amount, but it increases the

high-ﬁeld w

notably (Taniguchi 2003).

534

Magnetic Properties of Dislocations

slip planes. The preference for {111} versus {001} slip

has been discussed and the anomalous strengthening

of L1

-ordered alloys is commonly attributed to the

K–W locks (Kear and Wilsdorf 1962). More precise

arguments are founded on the dislocation core struc-

ture (Yamaguchi et al. 1982, Paider et al. 1984,

Umakoshi et al. 1984, Mryasov et al. 2002), APB

energies (Flinn 1960, Saada and Veyssie

re 1992,

1994), elastic anisotropy (Stroh 1958, Yoo 1987), su-

per intrinsic stacking fault (SISF) (Fang et al. 1994),

and complex stacking fault (CSF) (Hemker and Mills

1993, Balk et al. 2001).

It is interesting that Fe

Ge is one of a few L1

in-

termetallics, which does not exhibit anomalous

strengthening and this has been interpreted as a result

of the operation of cube slips in a wide tempera-

ture range (Ngan et al. 1992). Recently, production of

/110S/2 {001} partial dislocations and their multipli-

cation have been conﬁrmed by semi-in-situ experi-

ments of a thin foil of stoichiometric Fe

(Taniguchi 2003). Since the structural change from

to D0

in Fig. 6 or equivalently motion of

/110S/2 superpartial dislocations on {001} planes

does not involve interchanges of the FNN pairs (Flinn

1960), it is natural that plastic deformation does not

affect M

in Fe

Ge if FNN pairs are important. Con-

currently, the high-ﬁeld w

increases due to introduc-

tion of superdislocations as shown in Fig. 11.

On the other hand, plastic deformation induces

spontaneous magnetization in Pt

Fe (Bacon et al.

1963, Takahashi and Umakoshi 1990) and in Ni

(Yasuda et al. 2000) mainly by introduction of the

FNN Fe–Fe pairs as evidenced in neutron diffraction

study (Bacon et al. 1963). The spontaneous magnet-

ization can be attributed to the introduction of APB

tubes and APB ribbons as in the case of Fe–Al

(Sect. 3.4.(a)). The dissimilar properties of M

Fe and Ni

Ge (Fe

Ge) can be ascribed to the dif-

ference in the sites A and B occupied by ferromag-

netic atoms: B for the former and A for the latter as

shown in Fig. 6. Ferromagnetic FNN pairs are absent

before deformation in Pt

Fe and introduced by su-

perpartial dislocation motion, whereas FNN pairs of

ferromagnetic atoms are readily present in Ni

Ge and

Ge. The FNN pairs are destroyed by {111} slip in

Ge but not by {001} slip in Fe

Ge. The APB tubes

are not found in Fe

Ge, whereas /110S/2 superpar-

tial dislocations are found in {001} planes after both

bulk and thin foil deformation (Taniguchi 2003), in

support of the predominant operation of {001} slip in

Ge (Ngan et al. 1992).

3.5 Magnetic Transition Involving Large Strain

Both for practical and academic purposes, ferromag-

netic shape-memory alloys (SMAs) have gained much

attention in the recent years. Interest started after the

discovery of a large strain in a tetragonal Ni–Mn–Ga

induced by application of a magnetic ﬁeld at a low

temperature (Ullako et al. 1996). The elongation

strain becomes as large as 6% in Ni

(Murray et al. 2000), 0.6% in Fe

(James and

Watting 1998), and 2.3% in Fe

Pt (Kakeshita et al.

2000). These values are strikingly larger than usual

strain values for magnetostriction, e.g., 0.001%

(Rhoads 1901, Nagaoka and Honda 1902) B0.17%

(Clark et al. 1988). The effects of a magnetic ﬁeld on

the cubic-to-tetragonal phase transformation have

also been discussed in the presence of hydrostatic

(Kurisu et al. 1987) and unidirectional (Muto et al.

1987) stresses, raising a question of a distinction

between mechanical and magnetic forces.

Steuwer et al. (2003) have compared the energies

dissipated during 6% elongation in the two experi-

ments of Ni

(Murray et al. 2000), one

under a stress and the other under a magnetic ﬁeld.

The estimated energy dissipation is markedly larger

for the stress–strain closure than for the magnetic

hysteresis curves. The discrepancy may result from

combination of the following two factors: (i) Evolu-

tion of such a large strain may be ascribed to motion

of a line defect that carries uniform displacement or

uniform rotation of the tetragonal lattice. The latter

line defect is called disclination for which the eigen-

distortion b

ðfxgÞ is written in different manner from

a dislocation (see Mura 1987). (ii) The variables con-

trolled in the two experiments are different in terms

of thermodynamics (Callen 1960); the parameter un-

der control in the magnetic measurements is a mag-

netic ﬁeld, H, which is intensive, whilst it is the length

of the sample that is extensive, that is controlled in

the mechanical tests.

Regarding the point (i), the so-called ‘‘giant mag-

netostriction’’ expressed by f

would be localized

at the line defect. The large strain induced either by

magnetic or mechanical force must be attributed to

reorientation of the tetragonal lattice variants either

from the magnetically hard to soft variants, or from

the stress-free to stress-induced variants. Application

of a sufﬁcient magnetic ﬁeld will move the line defect

to reduce the magnetocrystalline anisotropy energy

by the variant conversion, whereas application of a

stress will do work to change the material dimension.

These motions will accompany energy dissipation.

The magnetic driving force required for the motion

would be supplied by the asymmetric distribution of

spins around the line defect under a high magnetic

ﬁeld, similar to that calculated for a dislocation

(Seeger and Kronmu

ller 1960, Kronmu

ller and Seeger

1961, Seeger et al. 1964, Umakoshi and Kronmu

ller

1981a, Go

ltz et al.1986).Thepoint(ii)comesinto

the evaluation of the magnetic energy dissipated by

motion of the line defect. In practice, the variant con-

version takes place from one domain to another in

discrete steps under a high magnetic ﬁeld (Steuwer

et al. 2003). The sudden change in the magnetic ﬁeld

caused by rotation of the spin direction would be

535

Magnetic Properties of Dislocations

adjusted instantaneously by the power supply to

maintain the magnetic ﬁeld set in the experiment.

Steuwer et al. (2003) suggest that the energy dissi-

pation may be found only by measuring the energy

consumed in the power supply that generates

magnetic ﬁeld.

From practical purposes, the large strain induced

by a magnetic force is useful, for example, as an

actuator that responds quickly by application of a

magnetic ﬁeld. The property differs from the usual

SMA, which requires heat for the shape change. The

quick response certainly comes from the characteris-

tic properties of magnetism controlled by motion of

electrons. However, it must be reminded that the

shape change occurs by motion of a line defect, ac-

companying a friction. The energy dissipated in the

shape deformation measurable on a stress–strain

curve should provide useful information about the

switching when such a material is used as an actuator.

4. Conclusion

Magnetic properties are affected by plastic deforma-

tion as a result of the introduction of perfect and

imperfect dislocations and of the motion of these.

Magnetic moments interacting with stress and strain

around dislocations change w and H

offering a non-

destructive method to investigate dislocation struc-

tures in ferromagnetic materials. The observation of

APB tubes by electron microscopy shows that the

large effects of plastic deformation on the magnetic

properties come from the interchanges of atoms in

the wake of jogs and kinks, and this is not necessarily

linked to the dislocation density generated by defor-

mation. Giant magnetostriction in a ferromagnetic

SMA is a current problem whose interpretation is

certainly also founded on the motion of a line defect.

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Magnetic Recording Devices: Future

Technologies

The ability to electronically access stored programs

and data was the enabling technology in the devel-

opment of general purpose computers a little over 50

years ago. One of the ﬁrst practical electronic mem-

ory devices was a tank of mercury, roughly the size of

an oil barrel, in which data was ‘‘stored’’ dynamically

in the form of acoustic pulses passing through the

mercury from a transmitter on one end to a receiver

on the other. Less than 10 000 bits could be stored for

only a few milliseconds at a cost of several dollars per

bit and, of course, the data was lost when the power

was shut off. Now for less than US$1000, one can

obtain a desktop computer operating at 700 MHz

with 128 Mbytes of high-speed semiconductor ran-

dom access memory (RAM) supported by a 30 Gbyte

magnetic hard drive.

This incredible improvement in performance and

cost has been driven in both semiconductor and

magnetic technology primarily by dimensional scal-

ing of critical geometries, the sizes of which are now

rapidly approaching fundamental limits. This article

surveys new concepts which have been proposed to

replace traditional magnetic disk systems. Only those

which are nonvolatile are considered. These include

patterned media, probe-based storage, and optical

storage.

1. Storage Metrics and Recording Limits

Storage density is the most important metric driving

both capacity and cost. Conventional longitudinal

disk storage, for which new world records are being

announced on the internet every few weeks, has

been demonstrated at 450 Gbit in

2

. Recent projec-

tions (Bertram and Williams 2000) suggest limits

of 100 Gbits in

2

for longitudinal recording and

500 Gbits in

2

for perpendicular recording. In fact,

values as high as 1000 Gbits in

2

seem theoretically

possible for perpendicular recording (Wood 2000). At

100 Gbits in

2

, the bit cell size is only 80 nm

. Thus,

new technologies which should realistically promise

factors of 10 or more greater density to justify in-

vestment must immediately confront dimensions be-

yond the limits of established processing technology.

Capacity is the number of bytes (8 bits) which can

be accessed directly. Drives with capacities 470 Gby-

tes are being sold today and by extrapolation should

reach capacities of 4700 Gbytes. The range of ca-

pacities, however, will extend as low as 1–10 Gbytes

in physically smaller drives targeted for increasingly

important portable applications.

The ultimate metric is the cost/byte. In 1999, the

average cost per Gbyte for hard disk storage was

B$20 and by 2001 had dropped to B$4. Again by

extrapolation, the cost of disk storage will certainly

be reduced to o$1/Gbyte for high capacity drives.

Access time is the time required on average to

reach a random bit. Because they are mechanical,

even the highest performance hard drives have rela-

tively slow access time of 46 ms. Even when rotation

speeds are increased to 20 000 rpm or higher, the cost

of reducing the head positioning time will probably

limit access times to B4 ms.

The rate at which data can be either written or read

is the same in hard drives but in other technologies

may not be. For fundamental reasons relating to the

dynamics of magnetization switching and available

head materials, data rates 45 GHz will be difﬁcult to

achieve in drives using only a single head. Multiple

head systems could, of course, provide high data

rates at a cost and power penalty.

539

Magnetic Reco rding Devices: Future Technologies