Thin films for high-density magnetic recording

Chapter 10

THIN FILMS FOR HIGH-DENSITY MAGNETIC RECORDING Genhua Pan Centre for Research in Information Storage Technology, Department of Communication and Electronic Engineering, University of Plymouth, Plymouth, Devon PL4 8AA, United Kingdom

Contents 1.

2.

3.

4.

5.

Instruments for Magnetic Measurement . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1.1. BH and MH Loop Measurement . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1.2. Magnetoresistance Measurement: Four-Point Probe Method . . . . . . . . . . . . . . . . . . . 1.3. Schematic Frequency Permeameter . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Basic Principles of Magnetic Recording . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.1. The Write/Read Process . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.2. Write Field of Recording Heads . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.3. Written Magnetization Transition in a Recording Medium . . . . . . . . . . . . . . . . . . . . Thin Film Recording Media . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.1. Physical Limits of High-Density Recording Media . . . . . . . . . . . . . . . . . . . . . . . . 3.2. Considerations of Medium Design . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.3. Preparations of Recording Media . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.4. Characterization of Recording Media . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Thin Films for Replay Heads . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.1. Anisotropic Magnetoresistance Films . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.2. Giant Magnetoresistance Films . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.3. Properties of Exchange-Biased Spin-Valve Films . . . . . . . . . . . . . . . . . . . . . . . . . 4.4. Spin-Valve Head Engineering . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Films for Write Heads . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5.1. Basics of Soft Magnetic Films for Writers . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5.2. Basics of Thin Film Writers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5.3. Magnetic Domain Configurations in Film Heads . . . . . . . . . . . . . . . . . . . . . . . . . . 5.4. Soft Magnetic Films for Writers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Acknowledgment . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

1. INSTRUMENTS FOR MAGNETIC

495 495 497 498 499 499 500 500 502 502 507 507 511 514 514 517 526 531 538 538 540 543 547 550 550

m e a s u r e the v o l t a g e i n d u c e d b y a c h a n g i n g flux, and (2) f o r c e

MEASUREMENT

t e c h n i q u e s that m e a s u r e the f o r c e e x e r t e d on a m a g n e t i z e d s a m p l e in a m a g n e t i c field gradient. B H l o o p e r and V S M are

1.1. BH and MH Loop Measurement

c a t e g o r y (1) i n s t r u m e n t s , a n d A G F M b e l o n g s to c a t e g o r y (2).

T h e m o s t c o m m o n l y u s e d i n s t r u m e n t s for M H & B H l o o p m e a s u r e m e n t are the B H looper, the v i b r a t i n g s a m p l e m a g n e t o m e t e r ( V S M ) , a n d the a l t e r n a t i n g g r a d i e n t force m a g n e t o m e -

1.1.1. B H Looper

ter ( A G F M ) . T h e s e i n s t r u m e n t s c a n be d i v i d e d into t w o cat-

A B H l o o p e r u s u a l l y e m p l o y s a pair o f H e l m h o l t z coils driven

e g o r i e s b y the t e c h n i q u e s used: (1) i n d u c t i v e t e c h n i q u e s that

b y an ac field o f f r e q u e n c y f r o m 1 to 100 Hz. T h e test sam-

Handbook of Thin Film Materials, edited by H.S. Nalwa Volume 5: Nanomaterials and Magnetic Thin Films Copyright 9 2002 by Academic Press All rights of reproduction in any form reserved.

ISBN 0-12-512913-0/$35.00 495

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PAN

pie is placed in a pair of pickup coils, which are connected differentially so that the voltages induced by the flux change due to the applied ac field are canceled out and the voltages induced by the flux change of the test sample are enhanced. The induced voltage in the pick-up coil is given by Faraday's d . _ N A dB law, v = N-57 dt' where N is the number of turns of the pick-up coil, and A is the cross-sectional area. The integration of the induced voltage v over time t gives the value of flux density B, which is fed to the vertical axis of a B H loop plotter (usually an oscilloscope). The H field value can be measured from the voltage across a resistor (usually 1 f~) in series with the Helmholtz coil, or by a Hall probe next to the pick-up coil inside the Helmholtz coil. A B H looper is normally used for a quick measurement of BH loops of soft magnetic films. It is a useful tool for determining the H c of soft magnetic films. But it is not suitable for measuring B H loops of films with large coercivities because of the limited field value a Helmholtz coil can produce (usually smaller than 100 Oe). BH loopers are also very limited in sensitivity and can only be used with relatively large samples. Because of the difficulties in calibrating the magnetic moment, usually they are not usually used for magnetic moment measurement.

1.1.2.

VSM

VSM is an instrument based on the principle of electromagnetic induction. Its basic structure is schematically shown in Fig. 1, which consists of an electromagnet, a vibration unit with a sample holder, a pair of pick-up coils, and a field sensor. A maximum field of 30 kOe is obtainable from an electromagnet field system. If a field higher than this is required, a superconducting magnet can be used, which is capable of producing a field as high as 160 kOe. The test sample (usually less than 10 mm square in size) in a uniform magnetic field can be treated as a magnetic dipole. The magnetic fluxes from the sample, which vibrates in the direction perpendicular to the applied field, induces an emf in the pick-up coil positioned near by. This emf is proportional to the magnetic moment of

Fig. 1. Schematicdiagram of an VSM.

the sample and is given by 3 e = ~tzoMo~gl

47r

(1.1)

where M is the magnetic moment of the test sample, ~o is the angular velocity of the vibration, and gl is the instrument factor. The value of gl is determined by measuring a sample with a known magnetic moment (calibration sample). After calibration with a standard sample, the output from the pickup coil represents the value of the magnetic moment of the sample, and the output from the field sensor gives the applied field value. The output of the instrument is therefore a MH loop or, more accurately, a curve showing the field dependence of the magnetic moment of the test sample, from which many magnetic parameters can be obtained. VSM is one of the most commonly used magnetometers for the characterization of magnetic materials in both the research laboratory and the production factory. The ultimate sensitivity of a modem VSM is as high as 10 -6 emu. And the field resolution can be as small as 1 mOe. It is widely used for measurements of soft magnetic films (H e and Ms), thin film recording media (H e, M s, M r t , S*, A M curve), and gaint magnetoresistance (GMR) spin-valve films (MH loops, Hex, Hf). When equipped with a cryostat or a heating module, VSM can be used to measure magnetic properties of films at various temperatures ranging from 4 K (with a liquid He cryostat) to 1000 K (with a heating module). This extended capability of VSM is particularly useful for the measurement of GMR, spinvalve, and tunnelling magnetoresistive (TMR) films and antiferromagnet/ferromagnet exchange systems. Details of these measurements are discussed in the corresponding sections. 1.1.3. A G F M

The alternating gradient force magnetometer (AGFM) was invented by Flanders in 1988 [1]. Its major advantage is that it has a sensitivity of 10 -8 emu, which is 100 time more sensitive than a VSM. Its basic structure is schematically shown in Fig. 2, consisting of a field system that is the same as that of a VSM, two pairs of field gradient coils, and a piezo bimorph unit with an extension cantilever, on one end of which the

Fig. 2. Schematicdiagram of an AGFM.

HIGH-DENSITY MAGNETIC RECORDING

497

magnetic sample is mounted. The test sample is magnetized by the field from the electromagnet and is at the same time subject to a small alternating field gradient, on the order of a few Oe to a few tens of Oe/mm. The alternating gradient field exerts an alternating force fx on the sample, which is proportional to the magnitude of the field gradient d h x / d x and the magnetic moment of the sample m x, and is given by dhx fx -- mx d x

(1.2)

The alternating force exerted on the sample produces a displacement of the cantilever, which is converted to the voltage output of the piezo bimorph. By operating at or near the mechanical resonance frequency of the cantilever, the sensitivity of the system is greatly enhanced. For a detailed quantitative correlation between the voltage output and the sample magnetic moment, refer to the original paper by Flanders [ 1].

1.1.4. Comparisons between VSM and AGFM Although the AGFM is 100 times more sensitive than the VSM, the real advantage of using an AGFM in terms of sensitivity is not that different. This is because a VSM can accept much larger sample sizes than an AGFM. The lower sensitivity of a VSM can be compensated for by its larger sample size [2]. The disadvantages of an AGFM are associated with its alternating gradient field, which makes it unsuitable for certain measurements, such as low coercivity, magnetization decay, and remanence measurements [3]. Table I gives a brief comparison of VSM and AGFM. For a detailed technical and application review of the two types of magnetometers, refer to the article by Speliotis [2].

Fig. 3.

Schematic diagram of a linear four-point probe.

probe method. A set of four linear microprobes is placed on a flat surface of a material with their separation distances marked in the figure as $1, $2, and $3, respectively. A constant current source is connected to the two outer electrodes, and the floating potential is measured across the inner pair. This configuration has been treated in detail by Valdes [4], who showed that when the probes are placed on a material of semi-infinite volume, the resistivity is given by V p -- -- -I When

S1-

27r

S2 -

S 3 --- S,

(1.4)

If the material to be measured is a thin film with a thickness of 8 on an insulating substrate and 8 << S, it can be shown that [4] p --

V 67r

(1.5)

I ln2

or sheet resistance R s is

The four-point probe is a widely used method for nondestructive measurement of sheet film resistance or magnetoresistance. Figure 3 is a schematic setup of a four-point Table I.

Eq. (1.3) reduces to V p = --27rS I

1.2. Magnetoresistance Measurement: Four-Point Probe Method 1.2.1. Linear Four-Point Probe

(1.3)

1 I S 1 + 1 I S 3 - 1 / ( S 1 + 1 / S 2 ) - 1 / ( S 2 -.[- S3)

p V I/" V Rs = -8 = I In 2 = 4.532 -1

(1.6)

Equation (1.6) implies that the value of the measured sheet resistance is independent of the probe spacing.

Comparisons of VSM and AGFM VSM

Typical parameters Noise floor

Typical sample size Range of measurement Optimal field resolution Typical measurements Low-coercivity measurement Magnetization decay measurement Remanence measurement Very thin films with high M s Very thin films with very low M s

AGFM

10 - 6 e m u 6 t o 10ram Up to a few thousand emu 0.001 Oe

Affected by AFG

Fine Fine Fine Fine Not suitable

Not suitable Not suitable Not suitable Fine Fine

10 -8 emu 1 to 3 mm Up to 1 emu

1.2.2. Square Four-Point Probe A square four-point probe as shown in Fig. 4 may also used for sheet resistance measurements. In the square probe configuration, the current is fed in through any adjacent probes, and the voltage is measured from the other two. For an equally spaced square probe configuration, the sheet resistance is given by [5] Rs =

V27r I ln2

V -----9.06-I

(1.7)

or the resistivity p of the film is p - 8R s =

V8 2,rr

V8

= 9.06~ I ln2 I

(1.8)

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Fig. 4.

Schematic diagram of the square four-point probe.

The advantage of using a square probe configuration for magnetoresistance measurement is that the current direction during measurement can be changed by a simple switch in the circuit without the need to touch the sample or change the direction of the applied field. A particular use of this is in the measurement of anisotropic properties of GMR [7].

1.3. Schematic Frequency Permeameter Knowledge of the permeability of soft magnetic films over a wide frequency range is important for recording head applications. However, conventional permeability measurement methods, such as the toroidal method or the impedance bridge method, are not suitable for measurement of thin film head materials because thin films for heads are usually anisotropic, whereas these methods can only be used for isotropic materials. There has been considerable interest in the past two decades in the development of permeameters for the highfrequency permeability measurement of anisotropic thin film materials. One of most successful instruments of this type was developed by Grimes et al. [6], who used a shorted transmission line as a driving source and a "figure-eight" coil to sense the flux change in the sample. The output signal from one of the figure-eight coil loops is measured as the $21 forward transmission parameter with a HP network analyzer and is converted to the complex permeability of the sample. The instrument can operate in the frequency range between 0.1 and 200 MHz.

Fig. 5.

Figure 5 is a schematic diagram of the swept frequency permeameter. The high-frequency driving signal supplied by the network analyzer is fed into a shorted transmission line jig with a characteristic impedance Z 0 of 50 fI. In the operation frequency range, the driving signal is considered a traveling wave with its E field perpendicular to the drive sheets and H field transverse to the drive sheets. The sample is inserted into one of the loops of the figure-eight coil with its easy axis perpendicular to the field direction for hard axis permeability measurement. The network analyzer is capable of storing in its registers the complex phasors measured by the transmission line jig at different frequencies, which is critically important for producing a permeability spectrum. It can also be interfaced, via its HPIB (IEEE 488.2) interface, for control, data logging, and data processing to a PC with HP VEE visual programming software. According to Grimes [6], the forward transmission parameter $21 is given by $21 = Vout/Vin =

-iogtZo(tL r -

1)aft + F

(1.9)

where Vout and Vin are the output and input voltages of the transmission line jig, i is the complex operator, ~o is the angular frequency of the driving signal, Af is the cross-sectional area of the magnetic film, 13(= 1 / 2 W Z o, W is the width of the drive sheets) is a geometry constant related to the shorted transmission line driving jig, and F represents a signal arising from a circuit response and possible loop imbalance when the permeability of the sample is unity (a magnetic saturated sample or a plain substrate sample). In the permeability measurement, the $21 parameters of a sample are measured with and without a dc saturation field and are stored in the network analyzer as $2110 and S21/sat, respectively. The value of S21/sat is equal to F by Eq. (1.1) and is numerically subtracted by ASsample = S2l[0 -- S 2 1 / s a t - -i(.o]~Lo(],s r - 1)af/~

(1.10)

A reference sample with known permeance is required to determine the permeability of a test sample. A similar operation for ASref is carried out for the reference sample, and the permeability of the sample is calculated from the formulae R = mSsample/mSref

A swept frequency permeameter consisting of a shorted transmission line jig, a figure-eight coil, and a HP Network analyzer.

(1.1 1)

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499

Fig. 6. A typical hard axis permeability spectrum for thin film samples.

and e -- [ ( / ' / ' r - 1)Af]sample/[ld~r-

:~ [[-~rt]sample//[lZrt]ref

1)Af]ref (1.12)

for samples with/.z r ~>~>1 and with the same sample size as the reference. Therefore [-I,r -- e[IJ, rt]ref /tsample

(1.13)

where t is the thickness of the magnetic film. The measured permeability is a complex number and is given by /2 = ~' - i/x"

(1.14)

where/z' and/x" are the real and imaginary parts of the complex permeability, respectively. Figure 6 is a typical permeability spectrum measured by the swept frequency permeameter.

2. BASIC PRINCIPLES OF MAGNETIC R E C O R D I N G 2.1. The Write/Read Process

The basic principle of an inductive writing process in a computer disk drive recording system is depicted in Fig. 7. The recording system is composed of a thin film recording head and a thin film recording medium. The thin film head consists of three major parts, a yoke-type magnetic core, a conducting Cu coil, and an air gap. The magnetic core is composed of a bottom yoke (P1) and a top yoke (P2), which is usually made of NiFe Permalloy or other soft magnetic films, such as CoNbTa, FeAIN, or FeTaN films. The recording medium is a thin layer (10-30 nm thick) of CoCr-based film coated with other nonmagnetic layers on a glass or A1 disk substrate. When in operation (write or read), the disk spins at a speed of 3600-10,000 rpm, and the head flies on the disk surface at a typical flying height of 10-100 nm, which is also known as the head-medium spacing. The information to be stored in a medium is first encoded and converted by signal processing into write current (write pulses) of the recording head. The write current, I, in the Cu coil produces a magnetic flux circulating in the core. The

Fig. 7. Schematic illustration of the inductive write process. Reproduced with permission from [8], copyright 1999, ICG Publishing: Datatech.

direction of the magnetic flux depends on the polarity of the writing current, which is usually toggled from one polarity to the other to write digital information into the recording medium. The gap permits part of the magnetic flux in the core to fringe out, intercept, and magnetize the recording medium in close proximity to the air-bearing surface of the head, forming recorded magnetization patterns in the medium (as also shown in the inset of the figure). The recording density is determined by the length of the written bit and the track width. The information thus stored in the medium can be kept for a long period (usually over 10 years) if the medium is not exposed to magnetic fields strong enough to magnetize the medium. The recorded information can be reproduced by a replay head (an inductive replay head, a MR head, or a GMR spinvalve head). When an inductive replay head, for example, passes over a recorded magnetization pattern in the disk, the surface flux from the magnetized pattern is intercepted by the head core, and a voltage is induced in the coil, which, by Faraday's law, is proportional to the rate of change of this flux. The induced voltage is then reconstructed by signal processing into the original signal. The magnetization pattern in the medium can be recorded longitudinally or perpendicular to the film plane, depending on the recording mode used or, in other words, depending on the type of head and medium used. The combination of a ring head and a medium with in-plane anisotropy tends to produce a recorded magnetization that is predominantly longitudinal. This is known as longitudinal recording and is the predominant recording mode used so far. An alternative to this is perpendicular recording [11], by which the magnetization is perpendicularly recorded in the medium by with the combination of a single pole head and a double-layered medium [14]. A double-layered medium consists of a perpendicular anisotropy medium and a soft magnetic back layer.

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2.2. Write Field of Recording Heads

2.2.1. Write Field of Longitudinal Heads Various magnetic heads have been used for magnetic recording. The magnetic heads for longitudinal recording are designed to produce a greater longitudinal field component Hx and, for perpendicular recording, a greater perpendicular field component Hr. The precise calculation of the magnetic field produced by a real head requires numerical computation. The most commonly accepted analytical solution for the field in front of the write gap g of a ring head with infinite size is the Karlquist approximation [12]. The longitudinal component of the writing field H x at a point P(x, y) beneath the gap is given by

Hx= Hg[aractng(2/x-+ )aractng(2/+rtx]')

Y

(2.1)

where x and y are the horizontal and vertical coordinates of point R and Hg is the deep gap field of the recording head. It can be seen from the above equation that the write field of a head is proportional to the deep gap field Hg. The inset of Fig. 7 gives a schematic view of the head-medium interface and the location of Hg and H x. According to Mallinson [9], the deep gap field Hg in Oersteds for a head with write efficiency of r/w, a gap length of g (cm), a coil of N turns, and a write current of I (amperes) is given by

Hg=

NI 0.47rr/w ~ g

These criteria provide a simple guideline for the design of recording heads and the selection of head materials and medium parameters. For example, if the coercivity H e of a recording medium for an areal density of 20 Gbit/in 2 is 3000 Oe, the required deep gap field of the writer will be 9000 Oe, and the saturation flux density of the head material for the thin film head will be 15 kG.

2.2.2. Write Field of a Perpendicular Single Pole Head The combination of a single pole head with double-layered perpendicular media [14] effectively places the media in the write gap of the head. The deep gap field is thus given by Eq. (2.3). And the maximum writable coercivity of perpendicular recording media is equal to the deep gap field of the single pole head, which is given by

H c = Hg ~ 0.6B s

(2.9)

It can be seen from Eq. (2.9) that the maximum writable medium coercivity in perpendicular recording with a single pole head and double-layered media combination is almost three times the maximum writable coercivity in longitudinal media. High medium coercivity is essential to overcome the superparamagnetic problem at increasingly higher recording densities, which is discussed in Section 3.1.3. Perpendicular recording has a major advantage in this over its longitudinal counterpart.

(2.2)

The maximum deep gap field a head can produce is limited by the saturation magnetization M s (or saturation flux density Bs) of the head material. Pole tip saturation occurs when Hg >_ 0.6B s

(2.3)

Hg > 0.5B s

(2.4)

Hg >_ 0.8B s

(2.5)

for thin films heads [9],

2.3. Written Magnetization Transition in a Recording Medium

2.3.1. Longitudinal Transition In digital recording, a bit of information is stored in the magnetic medium as the write current drives the writer to magnetize the medium right beneath the head gap along one of the two possible directions. An ideal written magnetization pattern representing di-bit information is shown in Fig. 8a, and the variation of magnetization along the x direction in the di-bit

for ferrite heads, and

for metal-in-gap (MIG) heads [ 13]. In high-density recording, the required deep gap field for effective writing in a recording medium with a coercivity of H e is given by [9] Hg > 3H c

(2.6)

To achieve effective writing, the saturation flux density of the head material for a thin film head must satisfy B s > 5H c

(2.7)

or the maximum medium coercivity, H c <_ 0.2B s

(2.8)

Fig. 8. A di-bit written magnetization pattern (a), and its ideal step function magnetization variation across the transition (b).

HIGH-DENSITY MAGNETIC RECORDING

501

pattern is given in Fig. 8b. In such an ideal case, the magnetization written in the medium is at one of the remanent levels. The transition between the two remanent levels is a step function with zero transition length. A real transition is not a step function. The variation of recorded magnetization in the x direction of the transition for longitudinal recording is best described by the arctangent model in which the magnetization distribution in the transition is given by

M (x) = -~M~arctan

a

(2.10)

where a is the so-called transition parameter, by which the transition length l is defined as l = Ira. Fig. 9 illustrates the arctangent transition model. When a recording bit pattern is written by the magnetic field from a recording head, the magnetized bit pattern will generate an internal field opposing the magnetization of the written bit because of the magnetic charges arising from the divergence of the written magnetization. This field is termed the transition self-demagnetizing field IId and is given by

i--~dV

Hal=--

(2.11)

where r is a position unit vector, V is the volume of the written magnetization pattern, and p is the magnetic pole density, which is defined as the divergence of magnetization M, p = -V.M

(2.12)

2.3.2. Perpendicular Transition The perpendicular recording was first proposed by Iwasaki and Nakamura [14] from their study on the transition selfdemagnetization phenomenon in longitudinal recording media. One of the fundamental differences between perpendicular recording and its longitudinal counterpart is the distribution of the demagnetizing field in the written transition. The transition self-demagnetizing field for longitudinal and perpendicular transition can be calculated with Eqs. (2.11) and (2.12) if the transition magnetization distribution is known. Figure 10 shows the transition self-demagnetizing field distribution for a

M

j# .m

+Mr "I---"" t/

s t r a i g h t line

approximation ,

~ - - - - ~I I

/

~"---

t_/

II

- Mr

arcta,, e,,t

transition

Fig. 9. The arctangent transition. Adapted from [9], copyright 1996, Academic Press.

Fig. 10. Schematic illustration of demagnetizing field distribution along a longitudinal steplike transition (a) and a perpendicular steplike transition (b).

longitudinal step-like transition (a) and a perpendicular steplike transition (b). It can be seen that the demagnetizing field in a perpendicular steplike transition is zero in the center of the transition and 47rMr away from the center, whereas in the case of longitudinal transition, the demagnetizing field is at maximum (47rMr) in the center of the transition and almost zero away from the transition. This implies that it is possible to realize steplike transition (zero transition length) in perpendicular recording from the point of view of transition self-demagnetization. The ultimate recording density for perpendicular recording is therefore not limited by the transition self-demagnetizing effect. The demagnetizing field away from the perpendicular transition (Hd = 47rMs), which is inherent from its thin film geometry, shears the perpendicular MH loop. For low-coercivity media, the shearing of the MH loop results in a considerably reduced remanence squareness ratio (S << 1) and wide switching field distribution (SFD). This will cause the time decay of recorded magnetization away from transition and high media noise [16]. Figure 11 shows the sheared MH loops of highcoercivity media (solid line) and low-coercivity media (dashed line), in comparison with the unsheared loop (dotted line loop) for films without the effect of H d. A strong anisotropy field, [Hk• 47rMs] > 5-10 kOe [15], is required to overcome this demagnetizing field and to give a stable remanent magnetization state of the medium (unity remanence squareness + a relatively large negative nucleation field Ha, as shown in the figure) [148]. A soft magnetic back layer reduces the surface magnetic pole density and therefore reduces the effect of the demagnetizing field [129]. Middleton and Wright [128] have used the WilliamsComstock model [18] to compare the transition length for perpendicular and longitudinal recording and concluded that a very narrow transition length could be realized in perpendicular recording.

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PAN The transition parameter a derived from the WilliamsComstock model [18], is given by

I

M

Hn ~,,

i.....

VHc

/i

where d is the head/medium spacing. The significance of the Williams-Comstock model is that it revealed the correlation of the transition parameter with the head and medium parameters. It confirms the self-demagnetization limit and reveals the effect of head/medium spacing on the recording density. A higher density (smaller a) is achievable by using a medium with smaller Mr8 and higher H c, and a smaller head/medium spacing.

Fig. 11. MH loops of perpendicular media. Dashed line, lower coercivity MH loop sheared because of Ha and S << 1; solid line, high-coercivity MH loop sheared because of Ha and with a negative nucleation field and unity squareness; dotted line, MH loop without shearing.

3. THIN FILM RECORDING MEDIA 3.1. Physical Limits of High-Density Recording Media 3.1.1. Transition Self-Demagnetization Limit One of the fundamental physical limits of longitudinal recording is the transition self-demagnetization limit. The selfdemagnetizing field for an arctangent transition is given by [9] Ha =

M~Sx qr(X2 -+- a 2 )

(3.1)

where Mr6 is the product of remanent magnetization M r and thickness 6 of a medium. The maximum value of a self-demagnetizing field Hd, ma~ occurs at x = a,

mr~

Ha'max = 27ra

(3.2)

Equation (1.6) implies that the maximum self-demagnetizing field in a longitudinal transition increases with the increase in linear recording density (decrease in transition parameter a). It is possible that at higher density the self-demagnetization field associated with a recorded transition might be high enough to demagnetize the recorded magnetization pattern and make the stored information disappear. If the recording transition is self-demagnetization limited, the minimum transition length lmin, a recording medium can sustain is obtained from Eq. (1.6) by equating n d , max " - n c , lmin : "tra :

)

27rMr8

He

(3.3)

A higher density (smaller lmin) can be achieved by using a smaller Mr6 or a higher H e medium. A modem recording medium, for an areal density of 20 Gbit/in 2, for example, has a typical Mr8 value of 0.45 memu/cm 2 and a H c of 2900 Oe [17].

3.1.2. Signal-to-Noise Ratio Limit The overall reliability of a digital recording system is determined by its bit error rate. Noise causes bit shift [19] and is one of the major error sources in a digital recording system. The minimum required signal-to-noise ratio (SNR) of a current hard disk drive is about 20 dB. Noise in a recording system arises from three predominant sources: the playback electronic circuits, the playback heads, and the recording medium [40]. In a modern recording system, the percentage of noise from each of the above three sources is typically 10%, 20%, and 70%, respectively. Therefore, the ultimate SNR limit in magnetic storage is perceived as being related to the inherent signal-to-noise capability of the media [19]. Medium noise can be separated into three somewhat distinctive sources [40]: amplitude modulation noise, particulate noise, and transition noise. Detailed physics and mathematical analysis of the three types of noise sources can be found in [40]. Thin film recording media possess virtually no amplitude modulation or particulate noise. The predominant noise source in thin film recording media is the transition noise. Transition noise refers to the phase and amplitude fluctuations of playback voltage, which are related to the transition center of recorded bits. It increases with the increase in recording density for longitudinal media. The medium transition noise is normally described in terms of the total noise power (NP) and the noise power spectrum (NPS). The noise power spectrum can be compared with the signal power spectrum leading to the narrow-band signal-tonoise ratio (SNR)n and the wide-band signal-to-noise ratio (SNR)w [10]. Noise in thin film media has been intensively studied in the past decade, particularly for longitudinal media [130-134]. It has been concluded that intergranular exchange coupling, which significantly enhances the size of magnetic structures and results in large-sized domains away from saturation, is the main origin of thin film medium noise. According to Bertram [135], the NP reduced by the Mr6 product of a medium is given by Mr6

. f C*, ~--~K.g

,~

(3.5)

HIGH-DENSITY MAGNETIC RECORDING where D is the grain diameter,

3.1.3. Gyromagnetic Switching and Superparamagnetic Limit 3.1.3.1. Gyromagnetic Switching of Magnetization Under a Reversal Field

g,g d is the intergranular interaction energy function, g ( 36, 3) is the intergranular magnetostatic interaction energy function, d is the intergranular nonmagnetic separation, and C* is the intergranular exchange coupling constant, which is defined by [135]

C* =

503

A

*

KuD '

(3.6)

where A* is the effective intergranular exchange coupling energy constant and K u is the crystalline anisotropy constant. Equation (3.5) implies that, for a medium with a given grain diameter, the reduction of the intergranular exchange coupling constant C* will lead to the reduction of medium noise power. For a medium with a given thickness, zero intergranular exchange coupling, and large crystalline anisotropy (large HK), the noise power decreases as the grain diameter decreases. However, if the intergranular exchange coupling is not zero, or the crystalline anisotropy of the medium is not strong, the reduction of grain diameter will lead to very complicated effects on the noise power because the reduction of grain diameter also leads to an increase in intergranular exchange coupling (Eq. (3.6)) and intergranular magnetostatic d coupling g ( 3a , 3)" There are distinct differences in noise characteristics between longitudinal and perpendicular media. For longitudinal media, noise occurs mainly at the transition centers and increases with increasing recording density. In contrast to this, medium noise in perpendicular recording occurs away from the transition centers and decreases with increasing recording density [ 132]. Low-noise thin film media can usually be considered as media with very low or zero intergranular exchange coupling. In this case, the physical interpretation of media noise developed by Mallinson for particulate media also applies to thin film media. According to Mallinson, the (SNR)w at high bit density is given by [10]

When a reversal magnetic field H (say a field from a writer) is applied to a recording medium at t = 0, the magnetization vectors in the medium respond to the applied field in the form of gyromagnetic precession around the field axis at an angular velocity of about 1.76 x 107 rad/Oe-s before reaching their final equilibrium state. This is schematically shown in Fig. 12, where M(0) and M(oo) are the two equilibrium magnetization states before and after the reversal, respectively. M(t) is the magnetization state at observation time t. This magnetization precession process is governed by the LandauLifshitz-Gilbert equation of motion in the absence of thermal perturbation (i.e., T = 0 K) [27], dM dt

.----

~

T(M

x

c~ Heft) + ~ss M

x

dM dt

1,3~)'".""

where M is the magnetization vector, and Heft is the sum of all fields acting on the magnetization vector, including the external field, the anisotropy field, the demagnetizing field, and the exchange interation field, a is the Gilbert damping constant, which is material dependent, and y is the gyromagnetic ratio (y = (lel/2mec)g = 1.76 x 107 rad/Oe-s = 2.8 MHz/Oe, where e is the electron charge, m e is the electron mass, c is the speed of light, and g is the Landau factor, g = 2). The time required to complete the gyromagnetic switching process (the switching time) depends on the material properties (mainly anisotropy field and grain volume) and the strength of the effective field. In general, numerical computation with Eq. (3.8) is used to study the magnetization reversal dynamics. The typical switching time in thin film recording media is less than 1 ns for head fields as high as or greater than the anisotropy field of the media. This implies that thin film recording media are capable of supporting a data rate up to 1 Gbits/s. For a data rate faster than this, the dynamic gyromagnetic switching process will start to play a role [ 174].

I Easy I ~ ~ L axis! I

(SNR)w = n --rmnWA2" 27r

(3.7) t

where n is the number of particles (grains) per unit volume, W is the track width, and /~min is the minimum recording wave length. Equation (3.7) simply states that the (SNR)w for a particulate medium is equal to the number of particles in a volume of size W-(Arnin/Tr). (Amin/2). This volume is the effective medium volume sensed by a replay head at any instant. The SNR of a medium is proportional to the number of grains per unit volume. The SNR decreases with increasing recording density (decreasing effective medium volume).

t)

i

,,, -,

t"" '~',

"9

-,

I "~,~ s Trajectoryof precessionof ~ ', ,,"" I ~ t h e magnetizationvector P

~

II

..r"

t

I

Fig. 12. Schematicillustration of gyromagneticprecession in a reversal field H applied parallel to the easy axis of the magnetic material.

504

PAN

Thermal fluctuation terms can also be introduced into Eq. (3.8) to study thermally assisted switching [25, 26]. In special cases where the switching time is at a moderate scale (no less than a few nanoseconds) and the material consists of uniaxial single-domain particles, semiempirical Neel-Arrhenius formalism is applicable, which is further dicussed in the next section.

3.1.3.2. Thermal Fluctuation After-Effect and Superparamagnetic Phenomenon With the assistance of thermal energy, the magnetization of a particle can overcome an energy barrier and switch from one stable state to another. This phenomenon is known as the thermal fluctuation after-effect, which is one of the magnetic after-effects [20] and is inherent to ferromagnets. Thermally assisted switching will take a certain time compared with the quick switching of particle magnetization in a large external field. The basic concept of thermally activated magnetization switching was first studied by Street and Woolley [21]. For uniaxial single-domain Stoner-Wohlfarth particles, or grains with coherent rotation only, the thermally assisted switching rate f (the probability per unit time of successfully crossing the barrier energy Eb) can be described by Neel-Arrhenius formalism [22].

f = fo eeb/kbr

(3.9)

where the barrier energy E b is the energy needed to keep the magnetization of the particle in its original state or the energy required to switch the original magnetization of the particle, and f0 is the attempt frequency, whose dependence on field and temperature is negligible compared with that of the exponential factor, and it is typically taken to be on of the order of 109 s -~. kb is the Boltzmann constant and T is the absolute temperature. At temperature T, the electron spins in the magnetization of a single-domain particle are subject to thermal agitation, the energy of which is represented by kbT. The time required for the magnetization of a particle to switch, known as the relaxation time or time constant z, is given by 1 7" " -

--

f

1 -"

--e

Eb/kbT

(3.10)

f

At temperature T, the time constant 7" depends on the value of the barrier energy E b. The limitation of Eq. (3.10) is that the minimum time constant obtainable by the equation is 1 ns when E b = 0. Therefore Eqs. (3.9) and (3.10) are only applicable to moderate time scales and to high-energy-barrier cases [30]. The energy barrier E b depends on the anisotropy constant and volume of the particle, the applied field, and the relative orientation of the magnetization with respect to the magnetic easy axis. For a uniaxial anisotropy Stoner-Wohlfarth particle with magnetization aligned parallel to its magnetic easy axis, the energy barrier E b in a zero extemal field is given by

Eb= KuV

(3.11)

where K u is the uniaxial anisotropy constant of the particle, and V is the particle volume, which is also the smallest switching volume in the thermally activated switching process (also termed activation volume). A smaller anisotropy constant K u or a smaller particle volume V will result in a smaller energy barrier E b which makes the particle more vulnerable to thermally assisted switching. If the energy barrier KuV at temperature T is on the same order of magnitude as kbT, the magnetization of the particle is expected to be activated to overcome the energy barrier [21] under a zero external field. The ratio of the barrier energy E b to the thermal energy kbT, Eb/kbT, is known as the thermal stability factor. For noninteraction and single-domain particles, E b = K uV, the thermal stability factor becomes K uV/kbT. According to Charap et al. [25], the behavior of particles can be classified essentially into the following three categories, according to their thermal stability factors: 1. Particles with a thermal stability factor much smaller than unity. These particles are superparamagnetic (i.e., they are internally magnetically ordered, but they lose hysteresis). 2. Particles with a thermal stability factor close to unity. The magnetic behavior of these particles is time dependent. Their response to small applied fields is linear. 3. Particles with a thermal stability factor much greater than unity. These particles respond to applied fields according to their switching characteristics, quickly acquiring an appropriate stable magnetization. The time dependence of magnetization M(t) and coercivity He(t ) for a particle of category 2 are given by [22] [23].

M(t) = Moe-ft +Moo(1-e -ft)

(3.12)

where M 0 is the initial magnetization value and Moo is the ultimate value, and

V kbT n (fot)] 1/2} Hc(tl = Hk l l _ L~uV1

(313) .

where Hk = 2Ku/Ms is the effective anisotropy field of the grain (Ms is the saturation magnetization) and t is the characteristic time derived for a changing field measurement (an MH loop, for example). The time dependence of magnetization is termed magnetization decay, which is the origin of signal amplitude decay in high-density magnetic recording media. Time-dependent coercivity is known as dynamic coercivity. As can be seen from Eq. (5.9), dynamic coercivity is a function of the thermal stability factor and the characteristic time. "Characteristic time" refers to the field duration or pulse width if a constant field is used. In the case of a smoothly swept field, the characteristic time is derived from the rate of change of the field [24], which is discussed further in Section 3.4.3.1. Figure 13 is the dependence of media coercivity on the characteristic time and the thermal stability factor of a medium,

HIGH-DENSITY MAGNETIC RECORDING 18000

505

Substituting Eq. (3.14) into Eq. (3.10), we have

16000 14000

~

~

1 V/kbT)[I_H/Hk]2 T = - -1e eb/,kb T . _ _e(ru

I

fo

, 000

~~-

~

~__KuV/kT=200

O 10000 8000

- ~ 1 0 0

6000 4000 2000

1.E-08

1.E-06

1.E-04

1.E-02

1.E+00

1.E+02 1.E+04

1.E+06

Time (Sec) Fig. 13. Media coercivity as a function of characteristic time and thermal stability factor K u V / k T , plotted with Eq. (3.13) and assuming H k = 16 kOe.

plotted with Eq. (3.13). As shown in the figure, for particles with smaller thermal stability factors, the dynamic coercivity is more sensitive to the characteristic measurement time. When the measurement time scale is reduced, the dynamic coercivity of the media increases considerably. For a recording medium with a thermal stability factor of 40, the measured coercivity could vary from 2100 Oe on a time scale of 10,000 s to 12,000 Oe at 10 ns. This has two major implications for magnetic recording: the writability of recording medium at high frequency as the field duration decreases to the nanosecond regime and the long-term stability of the written data. The time dependence of medium coercivity is the consequence of the natural response (gyromagnetic switching) of the magnetization of a material to an applied magnetic field. When a magnetic field is applied to a sample, the magnetization in the sample responds to the field in the form of precession around the field axis at a relatively high frequency of 7 H ( = 2.8H MHz) in the case of free precession, where 3/ is the gyromagnetic ratio and H is the applied field. It will take some time to reach a static equilibrium state. The coercivity of a magnetic material is a measure of the field required to switch the magnetization from one state to another (strictly speaking, the field required to reduce the magnetization of the sample to zero). The role of the applied field is to provide the magnetostatic energy (or Zeeman energy = - M s H cos 0) required to overcome the energy barrier and to make the switching happen. In that sense, the applied Zeeman energy is equivalent to the thermal energy (= kbT ) in terms of making the magnetization switching happen. As a result of the applied field, the energy barrier term in Eq. (3.10) is reduced by the amount of the equivalent Zeeman energy [168]. For particles of uniaxial anisotropy, with the preferred axis aligned with the field direction (0 = 0), the new energy barrier at field H is given by

2K u

[

= KuV 1 -

(3.15)

fo

Equation (3.15) is valid for H / H k < 1. The figure shows the switching time as a function of H / H k plotted by Eq. (3.15) in the field range H / H k = 0.55-0.95. We can see that the time required for a magnetization to switch depends on the strength of the applied field. For a very short measurement time scale, the field required to switch the magnetization could be very large, resulting in large coercivity. On the other hand, if the time scale of measurement is very long, the field required to switch the magnetization is reduced, resulting in smaller coercivity. In low-frequency applications, the magnetization switching time is negligible in comparison with the duration of the applied field. However, as the recording density increases, the duration of the field from the writing head is only a few nanoseconds, and the dynamic process of magnetization reversal in such cases is no longer negligible. When the field duration is shortened, a significantly larger field is required to realize the switching within that duration (Fig. 14). The dynamic coercivity of a recording medium is the true coercivity in the medium the writing head has to overcome during the writing process. It is therefore also called writing coercivity. As the linear recording density increases (the field pulse width decreases), the writing coercivity of a medium will increase significantly because of the time-dependent nature of the coercivity. In addition, a higher density medium has smaller grain volumes dictated by the low noise and small transition parameter requirements. This implies that the writing coercivity of the medium becomes even more sensitive to the write field duration. 3.1.3.3. Superparamagnetic Phenomenon in Thin Film Recording Media

Current thin film recording medium consists of hcp Co grains of various grain sizes with an in-plane magnetic anisotropy. 100000

.mo

10000

KuV/kbT = 50 1000

~

~

100

lO

0.4

0.5

0.6

0.7

0.8

0.9

1

H/I-K

(3.14)

Fig. 14. Characteristic magnetization switching time as a function of H / H k for particles with thermal stability factor K uV / k bT = 50.

506

PAN

Because the grains in modern thin film recording media are isolated, to a large extent, by the segregation of nonmagnetic Cr in the grain boundaries (dictated by low medium noise), a thin film medium can be assumed to be composed of uniaxial single-domain Stoner-Wohlfarth grains capable of coherent switching. Therefore, the general principle of superparamagnetism and the formulae developed in Section 3.1.3.2 apply to thin film recording media. However, some corrections are required to use these formulae for quantitative evaluation of the superparamagnetic effect in recording medium because recording medium differs from a Stoner-Wohlfarth particle in the following:

Fig. 15. Schematicillustrationof grainsizedistributionsin recordingmedia. 1. Easy axis distribution. The easy axes of hcp Co grains in the medium are randomly oriented in the film plane, and, therefore, it is unlikely that in the recorded data pattern the magnetization of each grain has a parallel alignment with its magnetic easy axis. It has been shown by Bertram and Richter [29] that for thin film media consisting of noninteracting grains with their easy axes oriented at random in the plane, the time-dependent coercivity is given by

Hc(t)_O.566Hk[l_O.977[kb T ( fot k--~ln\ln2)]

2/3

]

(3"16)

2. Activation volume. Grains in thin film media may not be completely free from intergranular exchange coupling. The grain volume V in the formulas discussed in this section is the smallest volume of material that reverses coherently in the thermally activated switching process. Such a volume is also known as the activation volume. For thin film media, if the grains are completely decoupled, the activation volume is equal to the grain volume. However, intergranular exchange coupling leads to activation volumes greater than a single grain volume. The activation volume in exchangecoupled medium may represent a significant number of grains. Bertram and Richter have pointed out [29] that exchange interaction between the grains reduces the pre-factor (0.566) in Eq. (5.12). A reduced pre-factor implies that the coercivity of a medium with intergranular interaction is less time dependent, or a medium is thermally more stable. For a typical medium, the pre-factor was estimated to be 0.474 [29]. However, large activation volume is the main source of medium noise [40].

On the other hand, if the distribution curve is wide and shifted toward the larger grain region, the medium is thermally very stable but with low SNR, because the number of grains in each recorded bit is considerably reduced. There is always a trade-off in media thermal stability and SNR. Obtaining a narrow grain size distribution is one of the key techniques used in the modern thin film medium preparation process to improve the thermal stability of a medium. A more detailed mathematical treatment of the effect of grain volume distribution on time-dependent magnetic behavior can be found in [37].

4. Transition self-demagnetization field. In recorded data patterns of longitudinal media, as discussed in Section 2.3, a self-demagnetization field exists in each transition. The energy barrier term E b needs to be modified to include the effect of the demagnetization field from the recorded magnetization pattern, with the from [24] E b = KuV 1 - ~

where H a is the spatially varying demagnetization field from the recorded magnetization pattern. A formula for Ha of arctangent transition was given by Eq. (3.1). The track width averaged di-bit demagnetizing field at the di-bit center is given by [39]

n d = 8Mr(t)' [~1 + (~--)2 W

3. Grain size distribution. The grain volume in thin film media is not the same for every grain, therefore a grain volume distribution must be considered in determining V. For typical media, grain size distribution function is a log-normal function [28]. The grain volume V in the formulas shown above should be the mean grain volume of the entire medium if they are used to assess the thermal stability of a medium. As shown in Fig. 15, a preferred grain size distribution should have a very narrow peak (solid curve) with its maximum probability density above and away from the thermally unstable grains. If the distribution peak is wide and shifted toward the superparamagnetic grain size region, the medium is thermally unstable.

(3.17)

-1

]

(3.18)

where W is the recorded trackwidth and B is the bit length. It can be seen that Ha increases with increasing recording density (decreasing B and W). At high densities, the high demagnetizing field of recorded transitions in longitudinal recording reduces the stability factor from its "bulk" value, leading to a rapid decay of playback signal [36]. One of the fundamental merits of perpendicular recording over its longitudinal counterpart is its lower transition selfdemagnetizing field, H a, particularly at higher densities. This makes the perpendicular recording thermally more stable than the longitudinal recording. However, because of the inherent

HIGH-DENSITY MAGNETIC RECORDING perpendicular demagnetizing field (H d = 47rMs) in the perpendicular media as discussed in Section 2.3.2, perpendicularly recorded bits tend to be unstable at lower recording densities. Special requirements for perpendicular media, such as a large anisotropy field, large coercivity, unity squareness, and a negative nucleation field, are required to overcome this inherent demagnetizing field.

3.2. Considerations of Medium Design In summary, from the viewpoint of medium thermal stability, we require a thin film medium with large grain volume or activation volume V, a large anisotropy constant K u, and small transition demagnetization field H d. To maintain a high SNR (minimum 20 dB), fine grains with large K u and minimum intergranular interaction are required. It is also known from Section 3.1 that to overcome the transition selfdemagnetization limit, a medium with small Mr6 and high H c is required. In the design of high-density recording media, the limit of grain size is set by the thermal stability requirement, typically a medium thermal stability factor of 40, i.e., E b = 40k bT at its working temperature T. The lifetime of the stored data estimated by Eq. (3.10) is about 10 years [25]. By the definition of the thermal stability factor, we have two options for achieving the thermal stability factor of 40: using larger grain volumes or using media with a larger anisotropy constant K u. Because a larger grain volume requires either large grain diameter or thicker films, which conflict with the SNR and transition self-demagnetization limits, media with large Ku become a preferred option. To meet the transition self-demagnetization requirement, there has been a trend to use media with increasingly smaller Mr6 and increasingly higher H e. The typical values of Mr~ are smaller than 1 memu/cm 2 for MR media [9] and smaller than 0.5 memu/cm 2 for spin-valve GMR media [17]. The lower limit of the Mr~ product is set by the increasingly smaller sensitivity of the read heads, i.e., by the increasingly smaller GMR ratio for the case of spin-valve replay heads due to the scaling effect, which is discussed in Section 4.4.3. The limit of higher media coercivity is set by two factors: the available medium coercivity of the existing recording media and the available writing field of the writer due to the limitations of saturation magnetization of the head material. The fundamental origin of high coercivity in thin film recording media is the combination of high magnetocrystalline anisotropy of the grains, smaller grain sizes, and lower intergranular interaction. The highest coercivities are obtained from films composed of the so-called Stoner-Wohlfarth singledomain grains. For media with exchange isolated grains, the maximum coercivity is given by H e = Hk/2, where Hk(= 2Ku/Ms) is the anisotropy field of the media. For Cobased media with an H k of 8-10 kOe, the limiting coercivity is about 4-5 kOe. Coercivities higher than 5 kOe are obtainable from FePt thin media because of their extremely large magnetocrystalline anisotropy of the long-range ordered FePt face-centered tetragonal (fct) L 1(0) structure [117-126]. However, the application of such films as longitudinal media is

507

inadequate because of the limited writing field a longitudinal thin film head can produce. As discussed in Section 2.2, the maximum writable media coercivity by a longitudinal thin film head is given by H e < 0.2 • 47rM s, where M s is the saturation magnetization of the thin film head material. For a medium with a coercivity of 4000 Oe, the required 47rM s of the head material will be 20 kGauss, which is almost the highest 47rM s of the currently available head material. This imposes fundamental restrictions on the maximum writable medium coercivity in longitudinal recording and will eventually limit the maximum recording density of longitudinal recording due to the superparamagnetic effect. On the other hand, the combination of a single-pole head with double-layered media in perpendicular recording is much more advantageous than its longitudinal counterpart in terms of maximum writable medium coercivity. Because the use of high-coercivity media is essential to overcoming the superparamagnetic problem in high-density recording, perpendicular recording has recently been considered as the technology to take the recording density beyond the superparamagnetic limit of longitudinal recording.

3.3. Preparations of Recording Media

3.3.1. CoCr-Based Longitudinal Recording Media Media Layer Structure. The magnetic layer is the essential layer for providing the storage mechanism for a recording medium. However, a modem thin film medium usually consists of multiple layers with layer structures as shown in Fig. 21 for a typical longitudinal medium. The purpose of using the multiple-layer structure is to obtain a controlled grain size, grain isolation (Cr segregation), low intergranular exchange coupling, high coercivity, and preferred crystallographic orientation of the magnetic layer. Chemically toughened glass or ceramic disks are usually used as substrates. A1-Mg alloy substrates coated with a very thick electrolessplated NiP layer have also been used in the past as substrate materials. The surface of these substrates can be assumed to be perfectly smooth, although they are polished and textured before the deposition of films. These various layers are usually deposited on the substrates by vacuum sputtering. The role of the seed layer is to provide an initial texture and grain growth templates, while still providing a smooth surface to the underlayer [42]. It may also provide an epitaxial growth surface condition for obtaining the desired crystallographic orientation of the underlayer. This crystallographic orientation is transferred to the magnetic layer via epitaxial growth in the growth process of the magnetic layer. The grain size and grain size distribution of the magnetic layer are mainly determined by the grain sizes of the underlayer. The underlayer also has a function to promote the Cr segregation and grain isolation of the magnetic layer, which is essential for achieving high coercivity and low noise. The overcoat and lubricant layers are added to provide their protective functions.

508

Seed Layer. Numerous seed layer materials have been used, which include Ni-A1 [32, 33], MgO [35], ZnO [35], and Co-Ti [34]. The typical thickness of these seed layers is between 50 and 100 nm. The primary functions of these seed layers are to provide the necessary conditions for the desired crystallographic texture of the underlayers and to promote fine grains.

Underlayer. The underlayer material is usually Cr or Crbased alloys, such as CrV, CrW, CrMn, CrTi, CrNb, and CrMo. Cr, which has a bcc structure and grows either with a [001] or [110] texture normal to the film plane, is a preferred underlayer for longitudinal media because it can provide the crystallographic texture for the epitaxial growth of the CoCr-based magnetic layer with in-plane c axis orientation (Fig. 16). The use of various Cr alloys allows a more precise lattice match with the hcp Co grains in the magnetic layer and hence better crystallographic orientation of the magnetic layer. They usually follow the epitaxial growth patterns of [35] A1-Mg/NiP substrates/seed layer/Cr(001)/Co(11 2 0) or glass (ceramic) substrates/seed layer/Cr(110)/Co alloy(1010). The Cr underlayer also promotes Cr segregation into the grain boundaries of the hcp Co grains of the magnetic layer. The additives in the underlayers, such as V, Mn, Nb, or Ti, have a second function of promoting fine grains. Magnetic Layer. Co-based alloys, such as Co86Cr10Ta4, Co75Cr13Pt12, or Co77Cr13Pt6Ta4, are used as the magnetic layer. The Co-based alloy has a hcp crystal structure, which has a magnetic easy axis in the c axis direction. The longitudinal anisotropy of the medium is obtained by growing the Co-based hcp grains with their c axis lying in the film plane. This is achieved by epitaxial growth of hcp Co grains on Cr-based alloy underlayers into [117-0] or [10i0] textures. In both cases, the c axis of the Co grains is parallel to the film plane. The primary role of the Cr in Co-based media is to separate the Co grains through Cr segregation into the hcp Co grain boundaries. Such segregation can be further enhanced by the addition of a small amount of Ta, typically 4-6%. The addition of Pt will elongate the c axis of the hcp Co and therefore give riseto larger magnetocrystalline anisotropy [42] and hence larger coercivity.

Effect of Deposition Parameters. The goal of sputter deposition of thin film media is to obtain a medium with fine

PAN grains, good grain size distribution, excellent crystallographic orientation, zero intergranular exchange coupling, and high coercivity. In addition to the use of a seed layer, an underlayer, and the right composition of magnetic layers, sputtering parameters, such as substrate temperature, substrate bias, and Ar pressure, also play important roles. Deposition at elevated substrate temperature is one of the key process parameters for obtaining high coercivity. The optimal substrate temperature range for deposition Co-based alloys is between 200~ and 300~ Figure 17 shows a typical substrate temperature dependence curve of media coercivity for CoCrPtTa films. H e increases sharply with the increase in substrate temperature, reaching a maximum at about 240~ The increased grain size due to high temperature and enhanced Cr segregation due to higher atomic mobility are mainly responsible for the increased coercivity [43]. Both dc and RF substrate bias in the range from - 1 0 0 V to 400 V have been used for the deposition of thin film media to increase the coercivity. The optimal bias voltages vary with the medium layer structure, film composition, and sputtering systems. The negative substrate bias, which provides a moderate ion bombardment to the growing film surface, has the following two major effects on the film growth process: increased atomic mobility and improved film quality. The atomic mobility of the adatoms is increased because of the extra energy provided by the ion bombardment. This effect is similar to the effect of high substrate temperature. The film quality is improved because the moderate ion bombardment may effectively remove the gas impurities in the film. The effect of sputtering gas pressure on film microstructure can be explained by the classical structure zone model [ 115]. In the deposition of thin film recording media, higher Ar pressure is effective in promoting columnar growth, Cr segregation, and grain isolation, but at the expense of a rough medium surface.

Optimization of Medium Coercivity. The fundamental origin of high coercivity in thin film recording media is the combination of high magnetocrystalline anisotropy of the

3500 3000

\

f

2500 ~" 2000 0 1500 1000

500 0

50

100

150

200

250

300

350

Substrate temperature (~

Fig. 16. Schematicdiagram of the layer structure of a typical thin film longitudinal medium.

Fig. 17. Dependence of He on substrate temperature for CoCrPtTa films.

HIGH-DENSITY MAGNETIC RECORDING grains, smaller grain sizes, and lower intergranular interaction. The highest coercivities are obtained from films composed of the so-called Stoner-Wohlfarth single-domain grains. In thin film media, maximum coercivity occurs at the transition from a multidomain grain to a single-domain grain structure [42]. Therefore, the coercivity of a thin film medium usually increases with decreasing grain size because of the formation of more single-domain grains in the film. But the smallest allowable grain size (grain volume) is constrained by the superparamagnetic limit. When the grain size (volume) is smaller than a critical value, the media coercivity will decrease with the reduction in grain size, or the apparent medium coercivity will increase with the reduction in measurement time scale, as discussed in Section 3.1.3.2. This region is termed the superparamagnetic transition region. Figure 18 shows the Mr8 dependence of coercivity of Co84Cr10Ta6 films deposited on different underlayers with the Circulus sputtering system from BPS [117]. The general trend of the Mr6 dependence of medium coercivity is that, at very small Mr6 (or very small film thickness), the coercivity of the films increases with the increase in film thickness before reaching the maximum coercivity. This is due to the fact that the grain size in the very thin film is very small and increases with the film thickness. At the maximum coercivity point, the film is dominated by single-domain grains with the maximum allowable grain volume for sustaining the single domain state. When the film thickness (or Mr6 ) is greater than the critical thickness (at an Mr8 value of about 0.6), multidomain particles are formed, and the coercivity of the medium decreases with the increase in Mrr. As one of the requirements of highdensity recording media is to use increasingly smaller Mr8 media (Mr6 < 1 memu/cm 2 for MR media [9] and M r 6 < 0.5 memu/cm 2 for GMR media [17]), it is important to prepare a recording medium with its superparamagnetic transition region below the Mr6 of the actual medium and with a slow transition into the superparamagnetic region. As can be seen from Fig. 18, the use of different Cr alloy underlayers also plays a part in determining the medium coercivities

Fig. 18. Dependence of the medium coercivity of Co84Cr10Ta6 films on Mr6 and underlayer materials. Reproduced with permission from [117], copyright 1999, ICG publishing Ltd: Datatech.

509

Fig. 19. Effect of film composition on the Mr6 dependence of medium coercivity. Reproduced with permission from [117], copyright 1999, ICG Publishing Ltd: Datatech.

below the superparamagnetic transition. Among the three Cr alloys (Cr96Ti4, Cr80V20 , and Cr), the Cr96Ti 4 underlayer gives the highest maximum coercivity at superparamagnetic transition and a much slower transition into the superparamagnetic region. The Mr6 dependence of medium coercivity also varies with the composition or constituent materials of the magnetic layer. This is shown in Figs. 19 and 20, respectively. Figure 19 shows the results of CoCrTa films with various compositions deposited on Cr underlayers. These curves show that the maximum coercivity points occurs at much lower Mr8 values for magnetic layers with higher Cr concentration. The highest coercivity of just below 2700 Oe is obtained for the Cos1Crl4Ta 5 alloy at an Mr6 of 0.5. The addition of the Pt to the CoCrTa ternary system, as shown in Fig. 20, can significantly enhance the value of coercivity and improve the superparamagnetic transition. A medium coercivity of over 3000 Oe is obtainable for the quaternary alloy. The superparamagnetic transition occurs at a much smaller Mr8 value (0.35 memu/cm2), and the transition into the superparamagnetic region is much slower than the ternary systems. As

Fig. 20. Dependence of coercivity on Mr6 for CoCrPtTa alloy. A & B represent the two sides of the disk. Reproduced with permission from [117], copyright 1999, ICG publishing Ltd: Datatech.

510 discussed in Section 3.3, this is due to the fact that the addition of Pt to the CoCrTa ternary alloy will elongate the c axis of the hcp Co and therefore give rise to larger magnetocrystalline anisotropy [42] and larger coercivity. The increase in magnetocrystalline anisotropy K u is mainly responsible for the improved superparamagnetic behavior because K uV/kT is also increased.

3.3.2. CoCr-Based Perpendicular Recording Media Perpendicular recording media have been intensively investigated since the discovery of perpendicular recording technology by Iwasaki and Nakamura [14]. The first generation perpendicular media are CoCr-based alloys, such as CoCrTa and CoCrPtTa, which are also used as the magnetic layers of longitudinal recording media. The major difference in the preparation of perpendicular and longitudinal media is the underlayers used to control the crystallographic orientation of the hcp Co grains. As discussed in the previous section, the magnetic easy axis of the hcp Co is along the c axis direction. In longitudinal media, Cr or Cr alloys are used as underlayers, so that the hcp Co grains grown on them have their c axis lying in the film plane. In perpendicular media, underlayers such as Ti [137], TiCr [136], or Pt [138, 139] are used to provide the epitaxial growth conditions for the hcp Co grains growing with their c axis perpendicular to the film plane. The Ti or TiCr alloy has a hcp structure that can grow on various substrates with their c axis perpendicular to the film plane and therefore provide the basis of epitaxial growth for the perpendicularly oriented hcp Co grains. The Pt underlayer has a fcc structure that tends to grow on various substrates with a [ 111 ] texture, and the hcp Co grains grown on the Pt fcc (111) plane exhibit excellent [0001] texture [139]. The quality of the c axis orientation of perpendicular media can be examined by X-ray diffraction patterns and the rocking curve method. Figure 21 is a schematic diagram showing the layer structures of double-layered perpendicular media. The functions of seed layer and underlayer are very similar to those of longitudinal media, providing grain size and crystallographic orientation control of the magnetic layer. The soft magnetic backlayer, usually Permalloy or Co-based amorphous alloys [140], is deposited on the underlayer to form the essential

Fig. 21. Schematicdiagram of the layer structure of double-layered perpendicular media.

PAN part of the double-layered media used in combination with the single-pole head. An interlayer between the soft backlayer and the magnetic layer may be required in some cases to provide the ultimate grain size and crystallographic orientation control of the magnetic layer. The usual additives in longitudinal media, such as Ta or Pt, are used for the CoCr-based perpendicular magnetic layer to achieve the same purposes, such as promoting Cr segregation or enhancing the magnetocrystalline anisotropy. One of the major disadvantages of CoCr-based perpendicular media is their relatively weak magnetocrystalline anisotropy with typical values of Hk• between 5 and 10 kOe (Hk• = 2Ku/Ms). The Ilk• values of the CoCr-based media are almost equal to the 47rM s values of the media, which does not meet the conditions of [Hkl-47rMs] > 5-10 kOe [15] and therefore are not sufficient to overcome the perpendicular demagnetizing field, H d = 4zrM s, inherent to the thin film geometry. Because of this, the shearing of the MH loop due to the perpendicular demagnetizing field makes the squareness ratio S of a perpendicular medium substantially smaller than unity. Figure 22 shows a typical MH loop of a CoCrTa/Ti perpendicular medium, the squareness ratio of which is only 0.38. The small S causes time decay of recorded magnetization away from transition and high media noise [16]. The CoCr-based perpendicular media also exhibit a relatively small coercivity, limited by its small magnetocrystalline anisotropy, and a magnetically "dead" initial layer. All of these make the CoCr-based media unsuitable for high-density recording.

3.3.3. High-Coercivity Perpendicular Media So far only two major categories of high-coercivity perpendicular media have been studied: the Co-based multilayers [142-148] and the FePt thin films [118-126]. Both types exhibit a Hk• much higher than the 47rMs values of the

Fig. 22. Typicalperpendicular MH loops of CoCrTa/Ti films (solid curves) and Co/Pd multilayers (dotted curves).

HIGH-DENSITY MAGNETIC RECORDING media and can easily be made to meet the condition of [Hk• > 5-10 kOe; i.e., they have a sufficiently high magnetocrystalline anisotropy to overcome the inherent perpendicular demagnetizing field. The high Hk• value of these films made it possible to obtain high medium coercivity. Because high medium coercivity is essential to combating the medium thermal stability problem, and the combination of a single-pole head with perpendicular media has the capability to write media with much higher coercivities than in longitudinal recording, the development of high-coercivity perpendicular media has been receiving increasing attention in recent years.

Co/Pt and Co/Pd Multilayer Perpendicular Media. Co/Pt and Co/Pd multilayers were originally developed as magnetooptic recording media because of their high perpendicular anisotropy [ 141 ]. Considerable work has been done in recent years on Co/Pd or Co/Pt multilayers as perpendicular recording media. It has been found that Co/Pd multilayer perpendicular media exhibit much higher perpendicular anisotropy than CoCr-based films [150]. Their perpendicular anisotropy field (Hk• typically ranges from 15 to 30 kOe, and their perpendicular coercivity (He• from 2 to 10 kOe. The MH loops of these multilayers exhibit a unity squareness ratio because the perpendicular anisotropy energy of the multilayers is sufficiently large to overcome the perpendicular demagnetizing field. These multilayers also show little "dead" layer. Overall, they are much better perpendicular media than CoCr-based films. The only drawback found in the early work was the high medium noise associated with these multilayers [145], even when the squareness ratio of these media is unity. The high medium noise is associated with the locally reversed magnetic domains away from transition [149], which has the same origin as the medium noise in CoCr-based perpendicular media. Wu et al. [148] have found that low medium noise can be achieved in Co/Pd multilayered perpendicular media if the medium is prepared to exhibit a MH loop with a unity squareness ratio, a significantly large negative nucleation field (typically 1.5 kOe), and a small slope at coercivity. High Ar pressure and an optimized Co and Pd thickness ratio are the two key parameters for the preparation of the low-noise multilayer media. A desirable MH loop of a Co/Pd multilayer perpendicular medium is sketched in Fig. 22 (dotted curves). Wu et al. [148] also demonstrated by magnetic force microscope images that a Co/Pd multilayer with a unity squareness ratio but a zero nucleation field H n exhibits a large number of reversed domains between the recorded transitions, whereas a Co/Pd multilayer with unity squareness and a negative nucleation field H n exhibits no reversed domains between recorded transitions. The Co/Pt multilayers exhibit large perpendicular anisotropy, however, the media tend to be very noisy because of the large domain size originating from the strong intergranular exchange coupling in the multilayers [146]. Reduction of media noise in the Co/Pt multilayer media was achieved by using CoB/Pt or CoCrTaJPt multilayers [146].

511

FePt Perpendicular Media. The long range ordered equiatomic FePt film (among the FePt, CoPt, and (CoFe)Pt family) with face-centered tetragonal (fct) L 1(0) structure [117-126] are of particular interests because of their high magnetocrystalline anisotropy energy (7 • 107 erg/cm 3 [121]), high coercivity, and good corrosion resistance. The perpendicular anisotropy field of these films is between 30 and 40 kOe, and coercivities up to 5-15 kOe are obtainable. The long-range ordering of the L 1(0) structure in these films is the essential condition for high coercivity. The magnetic easy axis of the FePt fct L1 (0) grain is long the c axis or the [001] direction. The FePt films can be made into longitudinal as well as perpendicular media, depending on the underlayers or substrates used during film deposition. Cr(100)/MgO, Pt, Au, and Pt seeded MgO(001) are the commonly used underlayers for obtaining perpendicular anisotropy [118-126], and FePt films deposited directly on a ZrO2 substrate exhibit in-plane anisotropy [ 127]. The as-deposited FePt films usually exhibit the fcc phase. High-temperature annealing is required to obtain the phase transformation from fcc to fct L 1(0). The required annealing temperature, depending on the underlayers, sputtering Ar pressure, and annealing gas atmosphere, ranges typically from 400~ to 600~ Suzuki et al. found that the ordering temperature of FePt can be reduced to 450~ if a sputtering Ar gas pressure of 100 Pa is used [121]. Park et al. obtained the long-range ordered FePt fct LI(0) phase at a temperature of 350~ with the in-air annealing technique [124]. FePt films exhibit very fine magnetic domains (<80 nm) [120], and the domain size appears to be inversely proportional medium coercivity. Because of the huge perpendicular anisotropy, the MH loops of FePt films have a steep slope, a unity squareness ratio, and a large negative nucleation field, which are essential properties for low-noise and high thermal stability perpendicular media. Excellent recording performances at high densities for both single-layered and double-layered FePt media have been demonstrated [120]. 3.4. Characterization of Recording Media

3.4.1. MH Loop The magnetization versus applied magnetic field (MH loop) of a recording medium can be measured by VSM or AGFM, from which key magnetic parameters for recording media, such as coercivity H e, remanence coercivity He,r, remanent magnetization M r, Mr 8 , saturation magnetization Ms, and squareness ratio S, are directly obtained at various points of the loop, as shown in Fig. 23. For high-density digital recording media, it is always desirable to obtain a sharp switching between the two discrete magnetization states. The switching behavior of a medium is usually described in terms of SFD or coercive squareness S*. The SFD measures the narrowness of the range over which the medium magnetization reverses or switches as a function of the applied field H. It is defined as the half-amplitude width AHs0 normalized by He of the derivative dM/dH curve of an MH loop centered about H c, as

512

PAN

M

Similarly, we get a dc demagnetization remanence curve Md(H ), as shown in Fig 24b, by reversing a field t o - H i from the positive dc saturation state, with the remanent magnetization Md(Hi) recorded as a function of reversing field until the negative saturation has been reached. According to Wohlfarth [45], the two remanence curves of noninteraction particles obey the Wohlfarth relation

s = M.._.~ ~

M,

Mr \ lb.

H

Md(H ) -- I-2Mr(H )

Henkel [46] investigated the nature of particle interactions by plotting Md(H ) versus Mr(H ) and suggested that the particle interaction can be characterized by a parameter AM, which is the deviation from the Wohlfarth relation and is given by

k dM/dH SFD = AHs~ Hc

A M = M,~(H) - (1-2Mr(H))

H Fig. 23. MH loop of recording media showing the definition of media parameters and derivative dM/dH of a MH loop showing the definition of SFD.

also shown in Fig 23. As can be seen from the dM/dH curve in the figure, the maximum switching rate (dM/dH) of a medium in an applied field H happens around H = He. The coercive squareness S* is defined by the equation dM dH

I.= c =

Mr

(3.19)

A value of S* close to unity is desirable for sharp switching. 3.4.2. A M Curve Measurements AM curve is usually used to characterize the intergranular interaction of recording media. It is derived from the measurement of the isothermal remanent magnetization curve Mr(H ) and the dc demagnetization remanent magnetization curve Md(H).As shown in Fig 24a, when a positive field H 1 is applied to an initially ac demagnetized thin film medium and then removed, we obtain a remanence value of Mr(H1). For a larger applied field H i, we get Mr(Hi). If the process is progressively repeated until saturation is reached, we can plot a curve of Mr(H ) as a function of the applied field H. such a curve is called the isothermal remanence curve.

l

M

I

MXHi)~ M,(HI) I

H~H, (a)

Fig. 24. curves.

M

[?Y v

H

Hi HI

[

(3.20)

iv

H

(b)

Schematic illustration for the measurement of Mr(H ) and Md(H )

(3.21)

Positive values of AM are interpreted as a result of stabilizing interactions. Conversely, negative values of AM are due to destabilizing interactions [46]. Specifically, a positive value of AM indicates the existence of intergranular exchange interaction, and a negative value of AM indicates weak or no intergranular exchange interaction. 3.4.3. Thermal Stabifity Measurements o f Recording Media 3.4.3.1. Dynamic Coercivity Measurement The dynamic coercivity of recording media can be measured by varying the sweep rate of the applied field (sweep rate method) or by varying the field duration of a constant applied field (constant field method). In the early work for the experimental observation of time-dependent magnetic effects, the dynamic coercivity was mainly measured by the sweep fieldmethod because of the difficulties of applying a constant field of a very short duration. The drawback of the sweep field measurement is its relatively large characteristic time (a minimum of 10 -5 s for a 60-Hz sweep field BH loop measurement [24]) due to the available slow sweep rate, which is not comparable to the field pulse width of a writing head (typically 10 -8 s to 10 -9 s). Stinnett et al. [153] developed a pulse field system with microstrips capable of measuring dynamic coercivities on a time scale of 10 -8 s. It was only recently that the constant field method has shown its advantages over the sweep field method due to the application of the spin-stand test in dynamic coercivity measurements [154-158]. VSM and AGFM Methods. The Measurement of the time dependence of media coercivity by VSM or AGFM is the most straightforward method. However, the limitation of this method is that it cannot go down very far in terms of the characteristic time due to the limit on the field sweep rate. Sharrock et al. [ 167] and Oseroff et al. [165] have developed a method to determine the characteristic time from the sweep field measurement. It has been found that the characteristic time (~') depends not only on the field sweep rate (R), but also on the dynamic switching process of a specific material [167]. A series of measurements is required to determine the

HIGH-DENSITY MAGNETIC RECORDING relationship between the field sweep rate and the characteristic time. The experimental procedure is as follows: 1. The sample is first magnetized to saturation by a field

ns. 2. A reverse field - H 0 is applied at a sweep rate R 0, and the time t o is recorded when the magnetization of the sample reaches zero at field - H 0. 3. Steps 1 and 2 are repeated with a reverse field - ( H o + 6H) at the same sweep rate R 0, and the time t is recorded when M = 0. 4. Step 3 is repeated for several field values. 5. lg(t/to) is plotted as a function of 6H. The relationship between lg(t/to) and 6H is given by [165], T In-- = - a 6 H

(3.22)

70

6. Parameter a is determined from the slope of the lg(t/to)-6H line. And the equivalent characteristic time z at field sweep rate R for that sample is given by [165] 1

Ra[ 1 -

e-a(Hs-Hf)]

1

Ra

where H s and He are the saturation and final fields, respectively, and a is the parameter associated with the dynamic switching process of a particular material and is determined by steps 5 and 6. In a modem VSM or AGFM with a stepped field system, the field between measurement points can be changed at a rate as high as 2000 Oe/s without causing any significant measurement errors [1]. The characteristic time for that sweep rate with a medium of a = 0.5 will be approximately 1 ms.

Spin-Stand Method.

This is an indirect measurement of medium coercivity as a function of measurement time scalefield pulse width or pulse duration, also known as rotating disk magnetometry [156]. Several experimental techniques have been developed with this method [154-158]. Robin et al. [154] and Moser et al. [156] used the method of erasing di-bits with pulsed fields from a recording head, and the medium coercivity corresponding to a field pulse width is determined from the write current needed to erase 50% of the di-bit signal. Corrections for the dynamic behavior of the head and for the demagnetizing field associated with the di-bit transition are required in the interpretation of the measured data. The method developed by Richter et al. [155] is based on the idea of using the reverse dc eraser noise as an indicator for coercivity [ 151, 152]. A band of dc saturated track is recorded in one direction, then a dc writing field in the reverse direction is applied along the track with a writing head and the broadband medium noise is measured. It has been confirmed in their experiment that the peak of the noise occurs exactly at the coercivity. A demagnetizing field correction is also required. The pulse duration for every portion of the written track is approximated as t = g/v,

513

where g is the gap length of the writing head and v is the linear velocity. Because only dc currents are employed, the rise time effect is disregarded. It was argued that the effective pulse duration time is equal to the number of revolutions (also known as reptation) multiplied by the pulse duration (t) of a single revolution. In this way, the pulse duration can easily be varied from nanoseconds to 100 s (10-12 decade seconds). It was observed that the noise peaks shift toward smaller writing currents with the increase in the number of revolutions, indicating that the medium coercivity decreases with the measurement time scale (field pulse duration).

3.4.3.2. Magnetization Decay Measurement The magnetization decay can be measured by either VSM or AGFM. However, as discussed in Section 1.1.4, VSM is a more appropriate instrument for magnetization time decay measurement because the alternating gradient field (on the order of a few Oe/mm) from the gradient field coil of an AGFM may affect the stability of the magnetization of the sample, resulting in false readings of the measurement results. The magnetization decay for a longitudinal medium is measured by first applying a saturation field to the medium, followed by a reverse field just a few Oe smaller than the medium coercivity. While that field is held constant, the change in magnetization of the sample is measured. The so measured time decay of the magnetization will emulate the magnetization stability of the recorded bits. The reversing field applied during the time decay measurement is an equivalent to the transition self-demagnetizing field in the recording transitions. The measured time dependence of magnetization M(t) is usually represented as a logarithmic function of time t[25],

M(t) -- C + S lnt

(3.23)

where C and S are constants and S ( = d M / d In t) is known as the magnetic viscosity coefficient. The thermal stability of a medium can also be represented by an equivalent field known as the fluctuation field H e, which is derived from the measurement of S and the irreversible susceptibility Xirr, and is then given by [30] He =

S

(3.24)

Xirr For a modem medium, H e is typically a few Oe. The irreversible susceptibility Xirr is usually determined from the differential of the appropriate remanence curve [159]. H e can also be determined from the measurement of time dependence of remanence coercivity Hcr(t ) [160, 161]. The magnetization decay for perpendicular recording media is measured by applying a saturation field to the sample and measuring the change in magnetization of the sample after the field is removed (at zero field). A reversing field is not required because a perpendicular demagnetizing field already exists in the medium and a perpendicular transition is free from demagnetizing field at high densities.

514

3.4.3.3. Signal Decay Measurement The time decay of recorded magnetization causes a decease in the playback signal output of recorded information over time, which can be measured by read/write tests [38, 161-163], known as the signal decay measurement. In the signal decay measurement, a MR-read/inductive-write composite head is usually used. To eliminate the effect of changes in the sensitivity of the MR head on the head signal output, the MR head signal should be calibrated with a freshly recorded reference track of the same medium and same recording density just before each measurement [38, 163]. It is also mandatory to scan the MR head over the written track to accommodate the inevitable position changes of the head while data are taken[28].

4. THIN FILMS FOR REPLAY HEADS The history of replay heads is closely associated with the advances of magnetic recording density. Figure 25 shows the evolution history of magnetic recording heads. Early generations of heads were ring shaped heads made of laminated mu-mentals (NiFeMoCu) or ferrites (Ni, Zn-Fe203). Thin film inductive heads were first introduced to the hard disk drive market in 1979 by IBM (IBM 3370) for an areal recording density of 7.7 Mbits/in2[175]. As the recording density increased, thinner media with higher coercivity were introduced, which required heads with high writing fields for effective overwrite and heads with higher output for acceptable signal amplitude and signal-to-noise ratio. The inductivewrite/inductive-read technology, which dominated the disk drive industry up to the early 1990s, could meet both the write field and the signal-to-noise ratio requirements for the areal densities up to 500 Mbits/in 2. Although the anisotropic magnetoresistance (AMR) phenomenon was discovered as early as in 1857 and the first AMR head was invented in 1970

PAN by Hunt [49], the first-generation disk drive product using integrated inductive-write/MR-read heads did not appear until 1991. AMR head technology allowed the extension of areal recording density to 4-5 Gbits/in 2. The discovery of the giant magnetoresistance (GMR) phenomenon [50, 51], particularly as applied to spin-valve head technology [7], has provided the technology for current disk drive replay heads, which is capable of taking the areal density up to or beyond 100 Gbits/in 2. The focus of this section is a review of GMR multilayers and spin-valve head technology. A brief review of AMR films and AMR heads is also presented.

4.1. Anisotropic Magnetoresistance Films Early investigations of electron transport properties of nonmagnetic conducting wires revealed that there was a small difference in resistance when a conducting wire was placed parallel or perpendicular to the magnetic field direction. If p// and p• represent the resistivities when the current is parallel and perpendicular to the magnetic field, respectively, we have P_L > P// and Ap = p • Pli.AP in general proportional to H 2, except at high magnetic fields and low temperatures [44]. This phenomenon is known as the ordinary magnetoresistance (OMR), which can be explained by the effect of Lorentz force experienced by the conduction electrons in a magnetic field. On the other hand, in ferromagnetic materials (Fe, Co, Ni) and their alloys, the electrical resistance depends on the current direction relative to the direction of the spontaneous magnetization in the sample. In contrast to OMR, in most cases we have p•
(4.1)

1 2 where Pay-- 3Pll -+- 3P•

Experimentally, Pay is usually replaced with P0, which is the zero field resistivity. The difference in Pav and P0 is insignificant with regard to the magnetoresistivity ratio [44]. The AMR phenomenon can be explained by spin-orbit coupling, which tends to induce an anisotropic scattering of the conduction electrons in the exchange split spin 1" and spin $ 3d-subbands [47]. Comprehensive reviews of AMR in ferromagnetic 3d alloys and its basic physics can be found in [44] and [47].

4.1.1. Basic Parameters of AMR Films 4.1.1.1. MR Ratio and MR Amplitude Because, in practice, the magnetoresistivity ratio given by Eq. (4.1) can be directly obtained from A R / R o without the need to know the dimensions of the specimen ( A p / p o - AR/Ro), Eq. (4.1) can be rewritten as Fig. 25. Evolution of magnetic recording heads. Reproduced with permission from [8], copyright 1999, ICG Publishing: Datatech.

AR R0

=

R(H)-R o R0

(4.2)

HIGH-DENSITY MAGNETIC RECORDING 4.1.2. AMR Measurement

where R 0 is the isotropic magnetoresistance of the sample, which is equivalent to the magnetoresistance measured when the current is perpendicular to the internal magnetization of is the resistance measured at field the film, i.e., R 0 = R• H. The zero level is defined as R ( H ) = R o. The maximum MR ratio is known as the MR amplitude and is given by (A~oR)

-

RII - R~

max

Because the MR change AR depends on the relative orientation of the measuring current and the internal magnetization in the film, the simplest measurement of AR of a uniaxial anisotropy film is to measure the difference in resistance along the magnetic easy axis and hard axis, i.e., AR = Rh.a.- Re.a, where Re.a. and Rh.a. are the resistance measured when the current is along the easy axis and hard axis, respectively. The dependence of the magnetoresistance on the magnetic field can be measured by starting either with the current parallel or perpendicular to the easy axis of the film, but the external field must be perpendicular to the easy axis, which will give a 90 ~ rotation of magnetization with the application of a saturation field. Figure 26 illustrates a typical four-point probe configuration, in which the measuring current is parallel to the easy axis (the M direction represents the easy axis of the film and the H direction is the applied field direction). At zero external field, the measuring current (the linear fourpoint probe) is aligned parallel to the internal magnetization of the film. At this point, the measured R value represents the value of RII, i.e., R = RII. As the perpendicular external field increases, the magnetization vectors in the film rotate progressively toward the field direction until saturation occurs. At saturation, the magnetization is perpendicular to the measuring current and this gives R = R1. A typical AR/R o ~ H curve measured by the configuration shown in Fig 26a is given in Fig. 26b. The AMR amplitude of Permalloy film is thickness dependent, ranging from 4% (1000-A,-thick film) to 1.5% (150-Athick film). The decrease in AMR amplitude with the decrease in film thickness is attributed to the increasing importance of surface scattering of the conduction electrons, which causes the increase in R 0, the isotropic resistance and the denominator of Eq. (4.3).

(4.3)

Ro

i.e., the difference in resistance, A R - RII- R• for currents flowing parallel (RII) and perpendicular (R• to the internal magnetization of the material, normalized by the perpendicular resistance (R0).

4.1.1.2. Field Sensitivity The differential form of the field sensitivity of an AMR film is defined by

S~/=

(4.4)

6H

The averaged field sensitivity of an AMR film is given by the ratio of the M R amplitude and the saturation field, i.e.,

AR/Ro) S. =

max

//s

515

(4.5)

At a saturation field, the AMR film is magnetized to magnetic saturation. The averaged field sensitivity is obviously proportional to the AMR amplitude and inversely proportional to H s. For a uniaxial anisotropy Permalloy film with infinite size (zero demagnetizing field), the saturation field H s is equivalent to the anisotropy field H k, which is typically 5-10 Oe. If the Permalloy film has a MR amplitude of 2.5%, the averaged field sensitivity is 0.5-0.25%/Oe.

4.1.3. AMR Replay Heads The sensing element of an AMR head (MRE) is a 150-250 A- thick sputtered Permalloy (81Ni/19Fe, by weight) film,

2-:5 !

0

-1.5

-I

-0.5

0

0.5

1

(b)

Fig. 26. Schematicdiagram showing a typical set-up for AMR measurement (a), and its corresponding MR curve (b).

1.5

516

PAN Substituting Eq. (4.7) into Eq. (4.6) gives a quadratic field dependence of magnetoresistance of a single-domain MRE [48], R=R 0+AR

Fig. 27.

A single-domain MRE placed in a magnetic field.

patterned into strips by a photolithographic fabrication process. Because the NiFe alloy film exhibits almost zero magnetocrystalline anisotropy and magnetostriction with this composition and is also very easy to make into a uniaxial anisotropy film by applying an orientation magnetic field during film deposition, it has become the unique choice for AMR and even GMR head material [9]. Figure 27 is a schematic diagram of a MRE with width W, height h, and thickness t, placed in a magnetic field H. The induced magnetic easy axis of the Permalloy film is usually aligned parallel to the ABS during head fabrication. The field (or flux) from the recording medium is perpendicular to the easy axis of the film. If we assume that the MRE is a single domain element, its anisotropic magnetoresistance can be expressed as a function of the angle 0 between the magnetization of the MRE and the sensing current and is given by R -- R o + AR cos 2 0

(4.6)

where R 0 is the isotropic resistance of the MRE (independent of 0), the value of which is equal to R• AR is the maximum resistance change, AR = RII- R• The 0 dependence of the magnetoresistance of a single-domain MRE can be sketched as in Fig. 28, which is similar in shape (but with different axis notations)to the A R - H curve in Fig. 26b. If the MRE has an infinite size, then the effect of the demagnetizing field is negligible. Simple energy arguments show that, for H < H k [48], H

sin 0 =

(4.7)

i eq

-100

-50

0 (degree)

50

100

Fig. 28. Magnetoresistance change in a MRE as a function of the angle between magnetization and current.

1-

Hkk

(4.8)

The magnitude of H k for Permalloy is typically 5 0 e . A real MRE is usually patterned into a rectangular shape with a finite width W and height h, as shown in Fig. 27. Because of the finite width and height of a real MRE, there are demagnetizing fields in both the vertical and horizontal directions. The values of two demagnetizing fields depend on the MRE dimensions and the magnetization angle 0, and they vary spatially. The average values of the two demagnetizing fields are given by [48] t

nd,y -- 47rMs. ~ 9sin 0

(4.9)

for a vertical demagnetizing field and t

Ha, x -- 47rMs. ~ .cos 0

(4.10)

for a horizontal demagnetizing field, where M s is the saturation magnetization of the MRE film. It can be seen from the above equations that the average demagnetizing field is proportional to the film thickness and inversely proportional to the size of the MRE. For a saturation flux density 4zrM s - 10 kG, a MRE height of 2/xm, and a thickness of 20 nm, the estimated average vertical demagnetizing field Hd,y by Eq. (4.9) is about 100 Oe, which is 20 times the anisotropy field (Hk) of the Permalloy film. According to Markham and Jeffers [48], because of the effect of a demagnetizing field the quadratic field dependence of Eq. (4.8) is retained over two-thirds of the resistance change, but H k is replaced by 4(H k + nd,y)]3, resulting in

[ (

R -- R 0 + AR 1 -

H (4/3)(Hk + Ha'y )

(4.11)

where nd,y is the average vertical demagnetizing field. The resultant MR response curve for a sensor with finite size, as shown in Fig. 29, is a much broader one in comparison with that of a single-domain film with infinite size. The demagnetizing fields significantly increase the saturation field and reduce the field sensitivity of the MRE. For this reason, the thickness of the MR film for high-density magnetoresistive heads must be as thin as possible, usually less than 15 nm. The demagnetizing field is nonuniform across the MRE and is much larger at the sensor edge than in the center. This causes the center of the sensor to saturate first and gives rise to the inflection points and skirts on the R ( H ) curve [48], which is also shown in Fig. 29. The inflection points occur at the average angle 0 -------55 ~ or H = H k + HO,y. Because of the quadratic field dependence of the MR response of a MRE, in magnetic recording head applications it is necessary to give an appropriate vertical bias field so that

HIGH-DENSITY MAGNETIC RECORDING .'l .~~_.~ / ill ~ ~| ~ll

MRE of infinite size Inflection point ~ ~ / ~ J

GMR multilayer exhibits maximum resistance in antiparallel magnetization configuration and minimum resistance in parallel configuration. The GMR ratio, which is field dependent, is defined as [51 ]

MRE of ) n i t e size ~ Bias ~ / point -Y

!] !

]~k

Skirt due

O

fields |

,

,

|

H/Ilk

.

.

.

.

.

.

517

|

Hb ~i00 Oe

Fig. 29. MR-H curves showing the effect of a demagnetizing field due to the finite size of a MRE.

the MRE will operate in the linear region. The ideal vertical bias condition is to set the vertical bias field at the inflection point. This will give an almost linear response of the MRE to vertical magnetic flux. A typical value for a vertical bias field is around 100 Oe, or H b - - H k d-Hd,y. Various vertical bias schemes were employed to obtain the bias fields for MR heads, which include permanent magnet bias, shunt current bias, soft adjacent layer (SAL) bias, exchange bias, self-bias, double-element bias, split-element bias, barber-pole bias, and servo-bias schemes. It is also necessary to have small a horizontal bias field (typically -~50e) along the easy axis of the MR film to eliminate the Barkhauson noise. For a comprehensive review of the various bias schemes and other topics related to MR heads, readers are encouraged to refer to [48] and [9].

The giant magnetoresistance (GMR) phenomenon was first discovered by Baibich et al. [50] and Binasch et al. [51] in 1988 in Fe/Cr multilayers grown by molecular beam epitaxy (MBE). The significance of the discovery is that the value of the observed magnetoresistance change is as high as 200% [50] in comparison with the 2-3% in AMR films. In 1990, Parkin et al. [52] demonstrated that GMR could be obtained in sputter-deposited Fe/Cr, Co/Cr, and Co/Ru multilayers and that the GMR amplitude could be even larger in sputtered multilayers than in MBE-grown samples. Since then, the magnetotransport properties of magnetic multilayers were under intensive investigation by many research groups, and many multilayer systems, such as Co/Cu [54], Fe/Cu [55], CoFe/Cu, CoNiFe/Cu [58], NiFe/Ag [56], etc. have been found to exhibit GMR.

-

Rsat

Rsat

-

(4.12)

Rsat

where R(H) is the resistance at external field H. Rsat is the resistance for a saturation field, i.e., when the magnetization vectors of the adjacent layers have a parallel configuration, here Rsat = Rtt = R~. The zero level is defined by R(H) =Rsat. The maximum GMR ratio is known as the GMR amplitude and is given by mR) -~sat max--

R1"$ -

esat

(4.13)

Rsat

whereRt+ is the resistance measured with antiparallel magnetization configuration, which is usually equal to the resistance at zero field, R 0. A typical GMR ratio (AR/Rsat) versus H curve is sketched in Fig. 30. The relative magnetization configurations of the adjacent layers at zero field and at saturation fields are also shown in the figure. The GMR ratio, AR/Rsat, is usually measured by the four-point probe method as described in Section 5.2.2. The definition of the GMR ratio is given by Eq. (4.12). Another commonly used definition of the GMR ratio is the one normalized by the resistance at zero field, R 0, and is given by

AR Ro

4.2. Giant Magnetoresistance Films

R(H)

AR

=

R(/-/) - Ro Ro

(4.14)

The two different definitions of the GMR ratio given by Eqs, (4.12) and (4.14) are related to each other by [66]

aR) _

(aR/R0)

(4.15)

However, the definition given by Eq. (4.12) is a more commonly used formula. This is due to the fact that the zero field

4.2.1. Basic Parameters of GMR Films 4.2.1.1.

GMR Ratio and GMR Amplitude

Giant magnetoresistance refers to the difference in resistance, AR = R t + - Rtt, or resistivity A p - P t + - Ptt, for a multilayer with its magnetization vectors in the neighboring layers aligned antiparallel (Rt+) and parallel (Rtt). A

Fig. 30. A typical AR/R -~ H curve of GMR multilayers, also shown the magnetization orientations in the adjacent layers.

518 magnetic configuration is often poorly defined. If the antiferromagnetic coupling is weak or if the hysteresis of the magnetic layers is large, R 0 can depend on the magnetic history of the sample [66].

4.2.1.2. Field Sensitivity Similar to the field sensitivity for AMR film, the averaged field sensitivity of a GMR film is defined by the ratio of the GMR amplitude and the saturation field, i.e., (AR/gsat)/n s. At saturation field, the GMR multilayer is magnetized to magnetic saturation and therefore has the lowest resistance. The averaged field sensitivity is obviously proportional to the GMR amplitude and inversely proportional to H s. For a magnetic sensor application, a GMR film with high field sensitivity is required, i.e., a GMR film with high GMR amplitude and a small saturation field. The differential form of field sensitivity is given by Eq. (4.4).

4.2.2. Multilayer GMR Measurement Figure 31 shows the two probe configurations for GMR measurement. In general, GMR multilayers (superlattice) are magnetically isotropic. Therefore, the four-point probes can be placed along any direction on the film. The difference in the two configurations is as follows. In Fig. 3 l a the applied field is parallel to the current, whereas in Fig. 3 l b the applied field is perpendicular to the current. It was initially thought that GMR was isotropic [50], i.e., the GMR amplitude is independent of the angle between the applied field and the sensing current. Therefore it can be measured by either of the two configurations shown in Fig. 31. However, Dieny et al. [7] have demonstrated that GMR is also anisotropic. A difference in GMR amplitude on the order of 7% was obtained by using a square probe configuration for a (Co/Cu/NiFe/CU)l 0 multilayer. The GMR amplitude measured with current perpendicular to the applied field, as in Fig. 31b, is 7% higher than that obtained with current parallel to the applied field, as in Fig. 31a. Some GMR multilayers (typically sandwiched films) are magnetically anisotropic. In such cases, the applied field for

PAN GMR measurement must be along the easy axis of the film to exclude the AMR contribution. If the applied field is perpendicular to the magnetic easy axis of the film (i.e., using the same probe configuration for AMR measurement as shown in Fig. 26), the GMR-H curve will be a hard axis response curve and the GMR amplitude will also include the contribution of AMR. A typical example of this can be found in [51].

4.2.3. Types of GMR Multilayer Structures 4.2.3.1. Anti-Ferromagnetically Coupled Multilayers (Superlattice) The GMR phenomenon was first discovered in antiferromagnetically (AF) coupled multilayers. A GMR multilayer consists of a number of very thin ferromagnetic layers separated by thin layers of nonmagnetic material, normally expressed in the form of (F/NM)n, in which F designates a ferromagnetic layer (transition metal Fe, Co, Ni, or their alloys), NM is a nonferromagnetic transition metal or a noble metal (V, Cr, Mn, Nb, Mo, Ru, Re, Os, Ir, Cu, Ag, or Au), and n is the number of bilayers. The archetypal system in this category is (Fe/Cr)n multilayers. At zero magnetic field, the magnetization vectors in the adjacent ferromagnetic layers of a multilayer system are ordered antiparallel (antiferromagnetic coupling), and the film exhibits the highest resistance. With the application of a sufficiently large magnetic field, the magnetization vectors in the multilayer become aligned in the same direction as the applied field (parallel magnetization configuration) and the film exhibits the lowest resistance. The GMR amplitude is determined by the maximum and minimum resistance, as defined by Eq. (4.13). The thickness of the magnetic layer is between 10 and 100 A, with an optimum thickness (maximum GMR amplitude) between 10 and 30 A. And the thickness of the nonmagnetic spacer can vary from 5 to 100 A. A higher GMR amplitude is usually obtained at smaller spacer thicknesses. But the GMR amplitude does not increase monotonically with the reduction of spacer thickness. The oscillatory dependence of GMR on spacer thicknesses is usually observed.

Fig. 31. Schematicdiagrams of the four-point probe configurationsfor GMR measurement (the H field is parallel to the magnetic easy axis). (a) Current parallel to magnetic field. (b) Current perpendicular to magnetic field.

HIGH-DENSITY MAGNETIC RECORDING

519

Table II. GMR Characteristics of Some Antiferromagnetically Coupled Multilayers Multilayer system (Fe 120/~/Cr 10t~k/Fe 120/~)

(Fe 30A/Cr 9/~,)60 Cr 100]k/(Fe 14]k/Cr8/~)50 Cr 100]k/(Fe 14/~/Cr8A)50 Fe 50A/(Co 8~dCu 9]k)60/Fe (Co/Cu 25~)10 (Co 18~/Ru 8/~)2o (Co 15~dCr4/~)3o (CogoFelo//Cu 12/~)1o (Ni76Fe10Co14/eu)l 0

Co70Fe30 4/~/Ag 15]k) Fe 50/~/(NiFe 10/~dCu 10/~)20 (NiFe 20A/Ag 10/~) (Ni81Fe~912/~dAg l lA)~ NisoFe2o/Ag

GMR amp. (%)

Measuring temp. (k)

1.5 50 150 30 115 18.5 2.5 2.5 23 11 100 19 50 17 6

Saturation field(kOe)

1.5 4.2 4.2 RT 4.2 RT 4.5 4.5 RT RT 4.2 RT 4.2 RT RT

20 20 13 0.5 10 6 2 0.05 3 3 1 0.3 0.01

Preparation method

MBE [51] MBE [50] Sputter [53] Sputter [53] Sputter [54] Sputter [58] Sputter [52] Sputter [52] Sputter [58] Sputter [58] Sputter, 77K [66] IBD [59] Sputter, 77K [62] Sputter/ann. [61 ] Sputter/ann. [57]

.....

Adapted with permission from [66], copyright 1994, Elsevier Science.

Typical GMR multilayer systems of this category, together with their GMR amplitude, saturation field, and preparation methods, are listed in Table II. It is worth noting that the GMR amplitudes and saturation fields vary considerably from one system to another. The reasons for these variations can be attributed to the influences of the combinations of magnetic/spacer materials, the interlayer coupling strength, the characteristics of the so-formed interfaces, and the effect of thicknesses of both materials on the transport properties of a particular multilayer system. Detailed discussions on the factors influencing the GMR amplitudes of GMR multilayers are given in Section 4.2.5. The saturation fields in early GMR superlattices such as Fe/Cr or Co/Cr are very large, typically ranging from a few kOe to a few tens of kOe. Although the multilayer has a much higher GMR amplitude than does an AMR film, the field sensitivity of a GMR multilayer is much smaller than that of an AMR film because of the high H s. The high H s is due to two factors: the strong interlayer antiferromagnetic exchange coupling in the Fe/Cr and Co/Cr multilayer systems and the high anisotropy field of the Fe or Co films. A high external field is therefore required to magnetize the sample to saturation because it has to overcome both the large anisotropy field of the magnetic material and the antiferromagnetic exchange coupling field. It was later found that the interlayer exchange coupling exhibits an oscillatory behavior [54]. Thicker spacer layers can be used to reduce the strength of interlayer antiferromagnetic coupling and hence to obtain low saturation fields, but at the expense of lower GMR amplitudes. The use of low-anisotropy materials such as Permalloy, CoFe, or CoNiFe as the magnetic material can also reduce the saturation field. The combination of these techniques has produced GMR multilayers with quite small saturation fields, for example, (Ni76Fe10COla/CU)l0 [58] and Ni80Fe20/Ag multilayers [57] with H s on the order of 10-50 Oe.

4.2.3.2. Uncoupled GMR Multilayers In this type of multilayer, the GMR effect is obtained in the absence of interlayer antiferromagnetic coupling. The antiparallel orientation of the magnetic moment in the successive layers is achieved by using alternate layers with different coercivities (e.g., Permalloy (low Hc) and Co (higher Hc) in a multilayer of (Permalloy/Cu/Co/Cu)n). When the external magnetic field sweeps between the positive and negative saturation of the multilayer, the magnetic configuration of the successive layers changes from parallel at large field to antiparallel at fields between the two coercivities of the two layers. The uncoupled GMR multilayers have the advantages of lower saturation field and higher GMR amplitude at room temperature than antiferromagnetically coupled systems [67].

4.2.3.3. Exchange Biased Spin-Valve Sandwiches This is the most successful type of GMR structure so far for recording head applications. An exchange-biased spin-valve sandwich is composed of essentially four films: a free layer (sensing layer), a conducting spacer, a pinned layer, and an AF exchange pinning layer. Seed and capping layers are also used in the deposition of the sandwich structure. A typical spinvalve with a layer structure of substrate/seed/F1/NM/F2/AF/ cap is schematically shown in Fig. 32. The pinned layer F2, free layer F 1, and Cu conducting layer are very thin, allowing conduction electrons to frequently move back and forth between the sensing and pinned layers via the conducting spacer. The magnetic orientation of the pinned layer is fixed and held in place by the adjacent antiferromagnetic exchange pinning layer, and the magnetic orientation of the sensing layer changes in response to the magnetic field, H, from the disk. A change in the magnetic orientation of the sensing layer will cause a change in the resistance of the combined sensing and pinned layers. When the magnetization of the free layer is parallel to the magnetization of the pinned layer, the spin-valve

520

PAN

Fig. 32.

Schematic diagram of a typical spin-valve layer structure.

film exhibits minimum resistance, and when the magnetization of the free layer is antiparallel to the pinned layer, it exhibits maximum resistance. The GMR amplitude is defined as the maximum resistance change normalized by the saturation resistance (minimum resistance of the film). A more detailed discussion on exchange-biased spinvalves is given in Section 4.3.

4.2.4. Basic Physics of GMR Phenomenon The interpretation for the origin of GMR, as first proposed by Baibich et al., was based on spin-dependent scattering of conduction electrons [51]. Usually it is assumed that the conductivity of a ferromagnet is due to the 4s band electrons and 3d band holes (empty band states). But the 4s electrons are the primary electric current carriers because the 4s bands are broad and the 4s electrons have low effective mass, in contrast to the narrow 3d bands and high effective mass of the 3d holes. When a conduction electron passes through a ferromagnetic metal, scattering will occur because of impurities, structural defects, phonons, and magnons, etc. The probability of occurrence of a scattering event will depend both on the effectiveness of the scattering centre and on the availability

of a terminal state into which an electron can scatter [65]. The 3d bands in ferromagnetic materials play a very important role in providing the terminal states (holes) into which the 4s electrons can be scattered. The relaxation time of the scattered electrons in the empty states contributes to the different conductivities (resistivities) of magnetic metals. Fig. 33 is a schematic illustration of the electronic band diagrams for the 3d ferromagnetic transition metals. The density of states function, N(E), for a 3d transition metal can be divided into two components: N(E) = N "~ (E) + N $ (E), where N 1" (E) and N $ (E) represent the density of states of the spin1" and spin,[ electrons, respectively. The spin1" direction usually refers to the direction of the magnetic moment (or spontaneous magnetization), and the spins direction is opposite to the magnetic moment. For a nonmagnetic 3d transition metal, N t (E) and N $ (E) are equal within any band. This is shown in Fig. 33a, where the 3d1" and 3d$ subbands have equal numbers of occupied states as well as empty states (holes). So do the 4s1" and 4s$ subbands. In ferromagnetic metals, as shown in Fig. 33b, because of the exchange interaction, the two 3d subbands of a ferromagnetic transition metal are split, resulting in more occupied states in 3d1" than in 3d$ bands and more empty states in 3d$ than in 3d1". The number of occupied and empty states in each 3d subbands varies with different metals. Table III is a list of the electron states for the three ferromagnetic elements (Fe, Co, and Ni) calculated by the electronic band theory. The value of the saturation magnetic moment of a ferromagnetic material depends on the difference in the densities of the occupied states in the 3d1" and 3d$ subbands and is given by m = (N t - N ~ ) / x B, where/x 8 is the Bohr magneton. The band theory predicts that (N t - N ~ ) is usually a nonintegral number, which agrees with experimental data. The 4s conduction electrons are distinguished into two families according to their spins (two-current model), the spin1" ( S - +~1 ) and spins ( S - _ 1 ) electrons. It is commonly assumed that the two families have equal numbers of electrons and that at temperatures comparatively lower than the ferromagnetic ordering temperature the two families of electrons are quite independent, i.e., there is no spin mixing between

Fig. 33. Schematicillustration of the electronic band diagrams, (a) before exchange interaction and (b) after exchange interaction. As a result of exchange interaction, the 3d band of transition-metal ferromagnetic materials (Co, Ni, Fe, and their alloys) is split into two subbands of opposite spins and with different densities of empty states at the Fermi level.

HIGH-DENSITY MAGNETIC RECORDING

521

Table III. Electron States of 3d Ferromagnetic Transition Metals [73] Electron

Empty states

Occupied states

Element

configuration

3dt

3d$

4st

4s$

3dr

3d$

Fe Co Ni

3d64s 2 3d74s 2 3d84s 2

4.8 5.0 5.0

2.6 3.3 4.4

0.3 0.35 0.3

0.3 0.35 0.3

0.2 0 0

2.4 1.7 0.6

the two spin channels. When the conduction electrons pass through a 3d ferromagnetic transition metal, they are likely to be scattered into empty energy bands of the same spins at the Fermi level. As can be seen from Fig. 33, because of the exchange interaction in 3d ferromagnets, the densities of empty states of the two 3d subbands are different at the Fermi level. The 3d$ subband has more empty states than the 3dr band. This results in different scattering rates for spin t and spins conduction electrons when they pass through the ferromagnetic metal; i.e., the spin'l" electrons will be less likely to be scattered (they can only be scattered into the empty 4s]" band), whereas the spins conduction electrons are more likely to be scattered (they can be scattered into both 3d$ and 4s$ empty bands) at the Fermi level. The difference in scattering rates results in different mean free paths (MFPs, or A) of the spin]" and spins conduction electrons in ferromagnetic materials. For example, in Permalloy (Ni80Fe20), the mean free path of the spin-up electrons, At = 50-100,~, is about five times longer than that of spin-down electrons, A+ - 10--20 ,~ [66]. The resistivity of each channel is inversely proportional to the MFP of electrons, therefore we have P t < < P+" The electron transport behaviors in a GMR multilayer can be explained by the phenomenological model shown in Fig. 34 [65] when the electron mean free path is much larger

Nt- N$

2.2 1.7 0.6

than the thickness of the layers. As shown in Fig. 34a, at zero magnetic field, a GMR superlattice has an antiparallel magnetization configuration due to the interlayer antiferromagnetic coupling, which means that the two channels of electrons are altemately strongly and weakly scattered as they cross the successive ferromagnetic layers or interfaces. The mean free paths for both species of electrons are very short. The resistance of the multilayer in such an antiparallel magnetization configuration is a maximum. If an extemal field H > H s is applied, as shown in Fig. 34b, the superlattice will be in a magnetic saturation state, i.e., the internal magnetization vectors are aligned parallel. The spin t electrons (spins parallel to the magnetization vectors) are then weakly scattered in all layers as they pass through the multilayer and therefore have very long mean free paths (fast electrons) and low resistivity, whereas the spins electrons (electron spins antiparallel to the magnetization vectors) are strongly scattered in every layer and therefore have very short mean free paths (slow electrons) and high resistivity. The electron transport properties in multilayers can be further simplified with a parallel resistive network of the two conduction electron channels [65, 67], as also shown in Fig. 34. In the case of an antiparallel configuration (Fig. 34a), the two channels have both low and high resistivity in altemate layers.

Fig. 34. A phenomenologicalmodel of the spin-dependent scattering for conduction electrons in GMR multilayers. (a) Antiparallel magnetization configuration and its correspondingresistance model. (b) Parallel magnetization configuration and its corresponding resistance model.

522

PAN

If we designate Pt and p+ as the low resistivity (when the spin of electrons is parallel to the magnetic moment, i.e., spin1") and high resistivity (when the spin of electrons is antiparallel to the magnetic moment, i.e., spin,k), respectively, the total resistivity of the multilayer in an antiparallel magnetic configuration can be estimated as [66]

where a - p~,/pt represents the contrast of the spin-dependent scattering of the conduction electrons when their spins are parallel or antiparallel to the magnetic moment.

However, the interfacial SDS is predominant in Co/Cu, Co/Ag, and Fe/Cr GMR multilayers. It is believed that in Co/Cu multilayers, the volume SDS is masked by the predominant interfacial SDS. When the ferromagnetic layers in the multilayer are alloy instead of pure metal, in particular with Permalloy, the volume SDS seems to be rather large and to dominate the SDS. For bulk scattering dominated systems, much thicker magnetic layers are required to alter the GMR amplitude. In contrast, in interfacial SDS-dominated systems, GMR amplitude can be significantly increased by the insertion of a very thin interfacial layer. Typical examples of the use of interfacial SDS to improve GMR amplitude is the insertion of two very thin Co layers (510/~), known as Co dusting, into NiFe/Cu multilayers [68] or into a spin-valve sandwich with a layer structure of NiFe(60 /k)/Cu(25 ,~)/NiFe(40 A)/FeMn(100 ,~). The GMR amplitude of the spin valve is almost doubled as a result of Co dusting [64]. For interfacial SDS-dominated systems, the quality of the interfaces (no interdiffusion) is vitally important for high GMR amplitude. The interfacial SDS characteristic (maximum increase in GMR amplitude) is usually obtained within a minimum thickness of two to three monolayers of dusting materials. The insolubility of the interfacial dusting material with the host magnetic and nonmagnetic layers is a prime requirement.

4.2.5. Parameters Influencing the GMR Amplitude

4.2.5.2. Interfacial Roughness and Film Quality

Excellent reviews of the influence of various parameters on the GMR characteristics can be found in [66-68]. The following discussions are adapted mainly from [66].

Interfacial roughness plays various roles in influencing the GMR amplitude. In Co/Cu or Fe/Cr multilayers for which the interfacial SDS is mainly responsible for their GMR amplitudes, some interfacial roughness may increase the density of scattering centers at interfaces and may therefore lead to a larger GMR amplitude. In contrast, in multilayers for which the interfacial SDS is not as important as the volume SDS (NiFe/Cu multilayers, for example), an increase in the interfacial roughness leads to a decreased GMR amplitude [66]. In the limiting case, the interfaces of GMR multilayers cannot be too rough or too smooth. If the interfaces are too rough, the mean free path of the conduction electrons becomes too small compared with the period of the multilayer. Consequently, the coherent interplay of SDS between the successive magnetic layers cannot occur efficiently. In the opposite limit of very smooth interfaces, the specular reflection of the conduction electrons on the interfaces may be increased, leading to a channeling of electrons within each layer. Such a channeling effect will also prevent the SDS from occurring [66]. The quality of the films also influences the GMR amplitude. In particular, the existence of pinholes through the nonmagnetic spacer may significantly reduce the GMR by introducing localized ferromagnetic coupling between magnetic layers [66]. However, the quality of the crystal growth does not appear to be a necessary requirement for obtaining GMR because the largest GMR amplitudes (in Co/Cu or Fe/Cr [54]) are obtained in multilayers grown by sputtering and not by MBE.

Pt~ =

P+ + Pt

(4.16)

2

In the case of a parallel magnetic configuration, the spin1" channel will exhibit very small resistivity and the spins down channel will exhibit very large resistivity, as schematically shown in the equivalent circuit diagram in Fig. 34b. The resulting total resistivity of the multilayer is given by Ptt=

2p~Pt

(4.17)

P, + Pt

and the GMR amplitude is given by Ap__pt~-Ptt__ Pt~ Pt~

(

P~-Pt P~ + Pt

)2( )2 =

a-1 a+ 1

(4.18)

4.2.5.1. Volume versus lnterfacial Spin-Dependent Scattering It is widely accepted that GMR originates from spindependent scattering (SDS). However, SDS in GMR multilayers may occur at the interfaces of F/NM layers, in bulk ferromagnetic films (volume), or in both. A question under contention has been the whereabouts of the conduction electron scattering, whether within the bulk of the magnetic layers or at the magnetic/nonmagnetic interfaces. The nature and location of the scattering centers that give rise to the SDS are important for GMR optimization. It is possible to experimentally distinguish these two types of scattering centres by introducing very thin layers of transition metal impurities at the F/NM interfaces or in the bulk of the ferromagnetic layers at various distances from the interfaces and looking at the influence of these impurities on the amplitude of the GMR [66]. Alternatively, they can be separated by measuring the transport properties of the multilayers with current perpendicular to the film plane (CPP) [76]. The variation of GMR amplitude with the thickness of magnetic layers also contains the potential information of the proportion of bulk and interfacial SDS [67]. The major findings on this topic are summarized as follows. Sputtered Co layers seem to give rise to significant volume SDS, whereas Fe layers have very weak volume SDS [71].

HIGH-DENSITY MAGNETIC RECORDING

4.2.5.3. Thickness of Magnetic Layer The dependence of GMR amplitude on the thickness of magnetic layers was extensively studied by many research groups [66]. The general observation is that the GMR amplitude always shows a maximum when the thickness of the magnetic layers (tF) is varied from the thickness of a monolayer to a few tens of monolayers. Typical experimental results for this are shown in Fig. 35. The optimal magnetic layer thickness IF, at which maximum GMR amplitude occurs, depends mainly on the type of GMR structures, the location of the SDS centers, the total number of periods (n) in the case of multilayers, and the thickness of the nonmagnetic spacer and even of the buffer and capping layers. For thicker multilayers (n > 10), as shown in Fig. 35a, the optimal magnetic layer thickness is typically 10-30/~, in contrast to 40-100 ,~ for spin-valve sandwiches, as shown in Fig.35b. For interfacial SDS-dominated systems, maximum GMR amplitudes were usually obtained at thinner magnetic layers. Dieny [66] has found from the study of spinvalve sandwiches in the form of FtF/CU 25 ,~JNiFe 50 ! ~ e M n 100 ,~, with F - NiFe, Fe, or Co, that the maximum GMR occurs at a Fe layer thickness of 50 ,~ for spin valves with Fe free layers in contrast to 80/~ for spin valves with NiFe free layers.

25 20 9

~-

9

9

]5

5 0

0

20

40

60

80

100

te(~.) 9

~

,

.

.

.

.

.

i

.

.

.

.

i

.

.

.

.

,i

,,

, - , . , . ,

4

g3 1 ~

NisoFeN -

0

* - ~ ' - - - - - ~ o Fe 0

100

200

300

400

500

t F (A) Fig. 35. Dependence of GMR amplitude on the thickness of ferromagnetic layers at room temperature in structures of (a) a multilayer of glass/(Ni80Fe20 50 ]VCu 10/~dNi80Fe20tF/CU 10 A)30, and (b) a spin-value sandwich of the form FtF/CU 25 ]~NiFe 50 JJFeMn 100 /~, with F = Ni8oF%o, Co, or Fe. The solid lines in both figures are fits of experimental data according to the phenomenological expression of Eq. (4.19). Reproduced with permission from [66], copyright 1994, Elsevier Science.

523

A phenomenological expression describing the dependence of GMR amplitude of multilayers or spin-valve sandwiches on the thickness of the ferromagnetic layers is given by [66] --~- (tF) --

0

]~FT~ ~

(4.19)

where t F is the thickness of the ferromagnetic layer, l F is the critical thickness at which the GRM amplitude is at maximum, to is an effective thickness that represents the shunting of the current in the rest of the structure (i.e., in all layers except the ferromagnetic layer, whose thickness t F is varied), and (AR/R)o is a normalization coefficient that depends on the combination of F and NM materials and on the thickness of the spacer layer. The fits of experimental data with the phenomenological expression are shown in Fig. 35 by the solid lines. As pointed out by Dieny [66], a limitation of Eq. (4.19) is that it does not properly account for the various incident angles 0 of the conduction electrons with respect to the plane of the layers. The effective thickness of the ferromagnetic layers seen by the conduction electrons varies as tel COS0. The decrease in GMR amplitude above the optimal thickness, as summarized by Dieny [66], is mainly due to the increased shunting current in the inner part of the magnetic layers. On the other hand, the decrease in GMR amplitude below the optimal thickness has two possible origins, depending on whether the multilayer structure has a large number of periods or is just a sandwich. In a multilayer with a large number of periods, where the scattering on the outer surfaces (substrate, buffer, or capping layers) can be ignored, the decrease in GMR at small tF is due to insufficient scattering of the normally strongly scattered species of electrons. As shown by Eq. (4.18), large GMR amplitude originates from the large contrast in SDS of the two species of electrons in the antiparallel and parallel magnetic configurations. Strong scattering of electrons in magnetic layers whose magnetization is antiparallel to the electron spins (spinS) leads to a short MFP (mean free path, A $) and high effective resistivity (Pt-/), whereas weak scattering of electrons in magnetic layers whose magnetic moment is parallel to the electron spins (spin1") leads to a long MFP (A 1") and low effective resistivity (PL)" The effectiveness of SDS of electrons in the antiparallel magnetic configuration is an important factor in determining the GMR amplitude. In the case of bulk SDS, the critical thickness below which insufficient SDS will occur is determined by the shorter of the two MFPs (A $ and A 1"), which is on the order of 10-20 ,~. In the case of spin-valve sandwich structures, the decrease of GMR amplitude at small t F is not due to the insufficient SDS of the spins electrons since the optimal thickness for spin-valve sandwich is always significantly larger than A $. However, if the magnetic layer thickness is smaller than A 1" (the longer of the two MFPs), the spin1" electrons penetrate the magnetic layer and are scattered by the outer surfaces (substrate, AF layer, buffer layer or capping layer), which results in small A 1", increased low effective resistivity (PL), or reduced SDS contrast and hence reduced GMR amplitude.

524

PAN

4.2.5.4. Thickness of Nonmagnetic Layers

Effect of Spacer Thickness on Multilayer GMR-Oscillatory Interlayer Coupling. The general observation is that the mul-

The spacer thickness dependence of GMR amplitude should be discussed separately with respect to multilayer GMR and exchange-biased spin-valve sandwich-type GMR because of the different mechanisms used to obtain GMR in the two systems. Although the GMR effect in the two types of systems originates from the spin-dependent scattering and it is commonly observed that the GMR amplitudes in the two systems vary with the spacer thickness, their mechanisms for obtaining the GMR are different, however. In multilayer films, the GMR phenomenon is closely related to the interlayer antiferromagnetic coupling, and the interlayer antiferromagnetic coupling is the essential requirement for obtaining GMR, whereas in spin-valve sandwiches, the interlayer coupling between the free and pinned layers is not related to GMR, and in sensor applications the interlayer coupling is not desirable.

L'

1.0

9

,

,

,

tilayer GMR amplitudes and the saturation fields decrease with an increase in spacer thickness. However, they do not decrease monotonically with increasing spacer thicknesses. Instead, they oscillate in magnitude as a function of the spacer thickness with a period ranging from 12 /~ in the Co/Ru system to 10-21 /k in the Fe/Cr and Co/Cr systems [52]. The GMR amplitude oscillation reflects the oscillations of the interlayer antiferromagnetic coupling through the spacer layer as its thickness varies. Figure 36 [68] shows the typical MH loops of a series of multilayers of the structure 100 ~ Ru/[30 /~ Ni81Felg/RU(tRu)]20 as the Ru spacer thickness varies from 4 ,~ to 44/k. The MH loops of these samples clearly show an oscillatory variation of saturation field with Ru thickness. For the Ru spacer layer thicknesses of 4, 12, 24, and 37 A shown in Fig. 36, the magnetization of the multilayer is saturated in

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. . . .

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.

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Fig. 36. Typical MH loops at room temperature for several multilayers of the structure 100 /~ Ru/[30 /~ Ni81Fe19/Ru (tRu)]20 as a function of Ru spacer thickness. Reproduced with permission from [68], copyright 1994, Springer-Verlag.

HIGH-DENSITY MAGNETIC RECORDING a very low field of about 10 Oe. For intermediate Ru thickness the saturation fields are larger, although they decay with increasing Ru thickness. A detailed dependence of saturation field on Ru thickness is shown in Fig. 37 for multilayers with the same structures as in Fig. 36. Five oscillations in the saturation field are shown in Fig. 37 with an oscillation period of about 11 A. In the limit of very thin Ru thickness, the coupling is antiferromagnetic. Even for a Ru layer that is just 3 thick, strong AF coupling is observed. This is in contrast to the Fe/Cr multilayer system, where the interlayer coupling becomes very weak as the Cr thickness is decreased below 8

A [52]. The GMR amplitude in multilayer systems is closely related to the interlayer coupling and hence to the saturation field. The highest GMR amplitude is always obtained at the first peak, which also gives the highest saturation field. However, the optimum thickness of a nonmagnetic spacer for obtaining the largest field sensitivity ( A R / R ) / H s is normally not at the first GMR peak. There is always a compromise between the highest GMR amplitude and the highest field sensitivity. But trcM must be chosen in a range for which the coupling is antiferromagnetic for any observable GMR. Interlayer antiferromagnetic coupling and oscillations in the coupling have been found in numerous multilayer systems, including ferromagnet/antiferromagnet (Cr, Mn) multilayers, ferromagnet/transition metal (Ru, Mo) multilayers, and ferromagnet/noble metal (Cu, Ag, Au) multilayers. A systematic study of the interlayer coupling in sputtered Co-based multilayers with 3d, 4d, and 5d transition metal spacers was carried out by Parkin [63]. The results are shown in Fig. 38. It was found from Parkin's work (see Fig. 38) that the period of oscillatory coupling is similar for most metals (typically 5-6 ML, where ML refers to monolayer), with the exception of Cr, for which the period is significantly longer (12.5 ML). Among all spacer metals, the Co/Ru system has the strongest antiferromagnetic coupling strength at its first peak (5 erg/cm 2, at 3/k). 10 I -

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This unique system has recently found application in synthetic spin-valve structures, which are discussed in Section 4.3.3. Note the following: 9 There are four cells beneath each element in the periodic table, marked A 1, J1, A1, AA1 and P. A 1 (/~) is the spacer layer thickness corresponding to the position of the first peak in antiferromagnetic exchange coupling strength as the spacer layer thickness is increased; J1 (erg/cm2) is the magnitude of the antiferromagnetic exchange coupling strength at this first peak; AA 1 (A) is the approximate range of spacer layer thickness of the first antiferromagnetic region; 9 and P (A) is the oscillation period. 9 The most stable crystal structures of the various elements are included for reference. Note that no dependence of the coupling strength on crystal structure or any correlation with electron density is found. 9 For the elements Nb, Ta, and W, only one AF coupled spacer layer thickness region was observed, so it was not possible to directly determine P. For Ag and Au no oscillatory coupling was observed in Co-based multilayers. Pd and Pt show strong ferromagnetic coupling, with no evidence of oscillatory coupling. The oscillatory behaviour in interlayer exchange coupling has been experimentally confirmed by magnetic domain imaging of a Fe/Cr/Fe superlattice sample with a wedge-shaped Cr spacer by scanning electron microscopy with polarization analysis (SEMPA) [60]. For a detailed review of the SEMPA technique, readers are encouraged to read reference [69] by Pierce et al. The interlayer antiferromagnetic coupling in GMR multilayers can be explained by the RKKY-like coupling, which is a type of indirect coupling between the successive magnetic layers through the spacer and is characterized by oscillations of coupling strength and phase. For a complete review of the RKKY coupling theory of GMR multilayers, please refer to reference [70] by Hathaway.

300 K

ql

0 v.x -10

525

2'o Ru s p a c e r

'

' 4O

layer thickness

6O (~)

Fig. 37. Dependence of saturation field on Ru spacer layer thickness for several series of Ni8~Fe~9/Ru multilayers with structure 100 ~ Ru/[30 Nis~Fe~9/RU(tRu)]20, where the topmost Ru layer thickness is adjusted to be about 25 A for all samples. Reproduced with permission from [68], copyright 1994, Springer-Verlag.

Effect of Spacer Thickness on Exchange-Biased Spin-Valve GMR. In contrast to multilayer GMR, the GMR amplitudes in exchange-biased spin valves decrease monotonically with increasing conductor spacer thickness. Typical examples of this are shown in Fig. 39 for Co/NiFe spin-valve sandwiches with Cu and Au conductor spacers [66]. As pointed out by Dieny [66], qualitatively this decrease is due to two factors: (i) the increasing scattering of the conduction electrons as they traverse the spacer layer (which reduces the flow of electrons crossing the spacer from one ferromagnetic layer to the next and therefore reduces the efficiency of the spin-valve mechanism) and (ii) the increased shunting current in the thicker spacer layers. The variation of the GMR amplitude with spacer thickness can be represented by the following phenomenological

526

PAN

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Oscillatory exchange coupling period is P (A), Coupling strength at first antiferromagnetic peak is J; (erg/cm'). Position of first antiferromagnetic peak is A, (,/k). Width of first antiferromagnetic peak is aA; (A). +No coupling is observed with Co Fig. 38. Compilation of data on various polycrystalline Co/TM multilayers with magnetic layers composed of Co and spacer layers of the transition and noble metals. Reproduced with permission from [68], copyright 1994, Springer-Verlag.

expression, similar to Eq. (4.19), which is derived for the ferromagnetic layer thickness dependence of GMR amplitude.

AR ( A R ) ( e-tNM/lNM) ----R-(tNM)= --R 1 l + tNM/to

(4.20)

where tNM is the thickness of the nonmagnetic spacer, INM is related to the MFP of the conduction electrons in the spacer layer, to is an effective thickness that represents the shunting of the current in the rest of the structure (i.e., in all layers except the nonmagnetic spacer, whose thickness, tNM is varied), and (AR/R)I is a normalization coefficient that depends on the combination of F and NM materials and on the thickness of the spacer layer. The fits of experimental data with the phenomenological expression are shown in Fig. 35 by the solid lines.

It is worth noting that in exchange-biased spin-valve films there also exists the oscillatory interlayer exchange coupling of the two magnetic layers through the conductor spacer as the thickness of the spacer varies. However, the existence of this oscillatory interlayer coupling does not affect the spinvalve GMR amplitude, as the GMR spin valve is obtained through a different mechanism. This has been experimentally demonstrated by Speriosu [72] for a series of spin-valve films of (Co 30 AJCu tcu/Co 30 ~ e M n ) ; the result is shown in Fig. 40. As shown in Fig. 40a, for tcu ranging from 18 to 40/k, the coupling between the Co layers oscillates between parallel (ferromagnetic) and antiparallel (antiferromagnetic) coupling with a period of 9 A, which is consistent with that obtained in GMR multilayers [68]. However, the GMR amplitude of these spin valves, as shown in Fig. 40b, decreases with increasing spacer thickness, which is independent of the oscillatory coupling of the two Co layers.

4.3. Properties of Exchange-Biased Spin-Valve Films

es

4.3.1. Basic Principle of Operation of Spin Valves o

<3

o

0

0

20

40

60

80 100 120 140 160 tNM (A)

Fig. 39. Dependence of GMR amplitude at room temperature on the thickness of the nonmagnetic spacer for a spin-value sandwich of the form Si/Co 70 /~dNMtNM/NiFe 50 A./FeMn 80 A with NM = Cu, Au. The solid lines are fits of experimental data to Eq. (4.20). Reproduced with permission from [66], copyright 1994, Elsevier Science.

In Section 4.2.3.3, we briefly introduced the basic structures of a spin-valve sandwich. The principle of operation of a spinvalve is again based on the spin-dependent scattering (SDS) of conduction electrons. Similar to the antiferromagnetically coupled GMR multilayers, spin-valve sandwiched films exploit the quantum nature of the conduction electrons, which have two spins, 1'and $. The sensing current, I, is composed of these two families of electrons. As shown in Fig. 41, when the magnetization of the free layer is parallel to the magnetization in the pinning layer; i.e., the electron spins in the magnetization of both layers are all spin1", the spin'l" electrons

HIGH-DENSITY MAGNETIC RECORDING

527

Fig. 42. Four-pointprobe measurement of spin-valve GMR. entation of the free layer. When the magnetic moments in the free layer and the pinned layer are antiparallel, the two channels of conduction electrons will be alternately strongly and weakly scattered as they pass across the free layer and pinned layer. This is equivalent to a two-resistor network in which both resistors have high resistance. The total resistance of the spin valve in this case is high.

4.3.2. Spin-Valve GMR Measurement Fig. 40. (a) Coupling field and (b) magnetoresistance vs. Cu thickness for spin valves of underlayer/Co 30 iVCutcu/Co 30 A,/FeMn. The two sets of points correspond to different underlayer thickness, giving different shunting. Reproduced with permission from [72], copyright 1991, American Physical Society.

Four-point probe method. Spin-valve GMR is characterized by low-field and high-field R(H) loops together with a series of characteristic fields. The low-field and high-field R(H)

in the sensing current will pass through the multilayer structure without any scattering, whereas the spins electrons in the sensing current are subject to heavy scattering by the spin]" electrons in the magnetization vectors of the pinned and free layers. This is equivalent to a two-resistor parallel network in which one of the resistors has a very small resistance. The total resistance of the spin valve is therefore very low. A change in the external field will cause a change in the magnetic off-

loops can be measured by the four-point probe method at low or high applied fields, respectively, for sheet films, or measured directly from patterned stripes for sensors. The low field is usually created with a pair of Helmholtz coils, and the high field is created With a pair of electromagnets. Figure 42 is a schematic diagram of the four-point probe setup for R(H) loop measurement, in which the four-point probe (or current) is in line with the magnetic moments. The four-point probes can be arranged either parallel or perpendicular to the magnetic moments. However, as pointed out by Dieny [7], the

Fig. 41.

528

PAN

spin-valve GMR is intrinsically anisotropic, i.e., the two configurations (parallel or perpendicular) will give different GMR values. It has been observed by Dieny et al. that the GMR amplitude is larger when the magnetic moments are perpendicular to the current in comparison with the situation where the magnetic moments are parallel to the current. They attribute this to the anisotropy in the a = Pt/P+ ratio caused by spinorbit coupling [7]. In their measurement, the contribution of AMR to the GMR ratio was excluded.

(%)

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./

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9

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Low-Field R(H) Loop and Coupling Field He. Typical lowand high-field R(H) loops of spin-valve films together with their field parameters are shown in Figs. 43 and Fig. 44, respectively. At a low applied field (typically H < 100 Oe), only the free layer of the spin-valve film responds to the applied field. The spin-valve film exhibits the lowest resistance when the magnetic moment in the free layer is parallel to the magnetic moment of the pinned layer, and it exhibits the highest resistance when the two magnetic moments are antiparallel. The low-field R(H) loop is offset from a zero field. This is due to the presence of an interlayer exchange coupling field (He) between the free and the pinned layers. As discussed in 4.2.5.4, the coupling field is spacer thickness dependent and is oscillatory in nature, i.e., it may be ferromagnetic (positive) or antiferromagnetic (negative), depending on the spacer thickness, but it has no effect on the GMR amplitude of the spin valves. The ferromagnetic coupling field He and the coercivity of the free layer Hc(free ) can be obtained from the low field R(H) loop, as shown in Fig. 43. However, one must note that the offset field of the R(H) loop measured by the four-point probe method represents roughly the ferromagnetic coupling field between the pinned and free layers (neglecting the field induced by measuring current in the free layer). If the R(H)

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1200

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loop is measured from a patterned sensor, the offset field is no longer equal to the ferromagnetic coupling field. In fact, it is equal to the sum of the ferromagnetic coupling field, the magnetostatic stray field from the pinned layer, and the field induced by the measuring current.

High-Field R(H) Loop and Pinning Field npi n. When a field higher than the exchange pinning field (npin) between the AF and the pinned layer is applied, the pinned layer will also respond to the applied field. The reversal of magnetic moment of the pinned layer (red arrows in Fig. 44) results in a second R(H) loop, which is offset from the zero field position by the unidirectional exchange pinning field, npi n. Before the reversal of the pinned layer, as shown in Fig. 44, the spin valve exhibits the highest resistance because the free and pinned layers have an antiparallel magnetic configuration. When the magnetic moment of the pinned layer is completely reversed, the spin valve returns to the lowest resistance state. The highfield R(H) loop gives a measurement of the exchange pinning field, npin, and the coercivity of the pinned layer, He(pinned).

VSM Measurement. The characteristic fields of spin valves can also be obtained from the MH loops measured by VSM. Figure 45 is a schematic illustration of the MH loops for spin-valve films, their field parameters, and the magnetization orientation of the pinned and free layers. Assuming that the saturation magnetizations of the free and pinned layers are equal, and in zero applied field the magnetic moments of the free and pinned layers have a parallel configuration, the spinvalve films are then at negative saturation state, as shown in Fig. 45. With the application of a positive field, the magnetization of the free layer will reverse first. When the free layer is completely reversed, the magnetization of the free layer is antiparallel to that of the pinned layer, and the spin-valve film

HIGH-DENSITY MAGNETIC RECORDING

M(arb. u.) i

npin

I

using different materials for F 1 and F 2 is that the material of F 1 can be chosen to obtain an excellent soft magnetic behavior, and F 2 can be chosen to obtain the highest exchange pinning field between the F 2 and the AF layers and maximum spindependent scattering. Ni81Fe19 (Permalloy) is one of the most frequently used free-layer materials because of its magnetic softness at very small thicknesses. Co, FeCo films are also used as free layers, but with a slightly larger saturation field. The largest spin-valve GMR values are obtained with Co and Permalloy layers separated by a noble metal (Cu, Ag, or Au) [66]. Cu is the most commonly used conductor spacer. Typical AF materials used in spin valves include FeMn, NiMn, IrMn, PtMn, RuRhMn, PtPdMn, NiO, a-Fe20 2, etc. Properties of AF exchange-coupled systems are discussed further in Section 4.4.4.

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l

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400

I

,

l

800

529

,

1200

H (Oe) Fig. 45. Schematic illustration of MH loops of spin-valve films, their characteristic fields, and the relative orientation of the magnetic moment of the pinned layer (red arrows) and the free layer (black arrows) at various points in the loops.

will exhibit zero magnetization. As the applied field increases further, the magnetization of the pinned layer will start to reverse, giving rise to the top curve of the second MH loop. In the maximum applied field, the pinned layer is magnetized to saturation, and the spin-valve film is saturated by the applied field and has a parallel magnetic configuration in the direction opposite that of the original.

4.3.3. Types of Spin Valves Since the discovery of the first spin-valve structure [71 ], many different types of spin valves have been introduced by various researchers. Here is a list of some of the spin-valve layer structures reported so far.

1. Basic Spin Valves. The basic layer structure of a spin valve consists of four function layers grown on a substrate coated with a Ta seed layer. One usually finishes the deposition of spin-valve multilayers with a Ta protection layer. The structure is either in the sequence of substrate/Ta/F1/NM/F2/AF/Ta or substrate/Ta/AF/F2/NM/F1/Ta, where F 1 refers to a free layer, NM to a noble metal conductor spacer, F 2 to a pinned layer, and AF to an antiferromagnetic layer. The first sequence is known as a top spin valve and the second sequence is called a bottom spin valve. Because of the deposition sequence dependence of exchange coupling between the F2/AF layers, not all of the spin valves can be directly made into top and bottom spin-valve structures by the sputtering process. For example, if FeMn AF material is used, only top spin valves can be made because no exchange coupling between F2/AF can be obtained if the FeMn AF layer is deposited before the ferromagnetic layer. The magnetic materials for F 1 and F 2 can either be the same or be chosen differently. The advantage of

2. Interface Engineered Spin Valves. This type of spin valve has a layer structure of substrate/Ta/F1/Co/NM/Co/F2/AF/Ta, i.e., with two extra very thin Co layers (about 5-20/~) inserted into the interfaces of free layer/Cu and Cu/pinned layer [78]. The very thin Co layers are known as "dust layers." FeCo may also be used as an interface-dust layer. As discussed in Section 4.2.5.1, dust layers are used to enhance the interfacial spin-dependent scattering and hence to enhance the GMR amplitude of the spin valves. 3. Symmetric Spin Valves. A symmetric spin valve [88, 89] has a layer structure of substrate/Ta/AF/F2/NM/F1/ NM/F2/AF/Ta; i.e., it is a combination of a top and a bottom spin valve, but with a shared free layer. The two pinned layers of a symmetric spin valve are biased in the same direction (parallel magnetization configuration). As shown in Fig. 46a, when the magnetic moments in the free and pinned layers are parallel, the symmetric spin valve exhibits a low resistance state. When the magnetic moments of the free and pinned layers are antiparallel (Fig. 4.6b), the symmetric spin valve is in a high resistance state. Because the symmetric spin valve allows double spin-dependent scattering, as shown in Fig. 4.6b, its GMR amplitude can be increased to about 130% of that of a simple spin valve, but at the expense of an almost doubled sensor thickness. Another minor drawback of this structure is that, as the free layer is between two pinned layers, it feels a double coupling strength [79].

4. Synthetically Pinned Spin Valves. The synthetically pinned spin valve [91, 93, 110] uses an alternative pinning mechanism to replace the AF/F exchange pinning system in conventional spin valves. Its layer structure is schematically shown in Fig. 47. It consists of an AF pinning layer, a synthetic ferrimagnet (or antiferromagnet) pinned layer (SF), a conductor spacer, and a free layer (F1). The synthetically pinned layer has a structure of Fa/Ru/F b, where F a and F b are ferromagnetic materials (Co or FeCo). The Ru layer usually has a thickness of 3-5 /~. The name "synthetic ferrimagnet" comes from the fact that the two ferromagnetic layers have different thicknesses and therefore different saturation magnetic moments. As discussed in Section 4.2.5.4, the Co/Ru/Co

530

PAN

Fig. 46. Schematicdiagram of a symmetric spin valve. (a) Low-resistance state. (b) High-resistance state. or FeCo/Ru~eCo multilayers are unique in that they exhibit very strong antiferromagnetic coupling (ncoupling is up tO a few thousand Oe) when the Ru thickness is in the range of 3-5 ,~. Synthetically pinned spin valves exhibit the following three major advantages over conventional spin valves: (i) the antiferromagnetic pinning system becomes very rigid because of the strong RKKY indirect coupling of the two ferromagnetic layers and the reduction of the self-demagnetizing field within the pinned layer; (ii) the thermal stability of the AF exchange pinning system is much inproved [94]; (iii) there is a significant reduction of the demagnetizing field in the free layer arising from the pinned layer stray field, which allows for the reduction of the free layer thickness. All of these features are important for very small (submicron-sized) spin-valve heads, which are further discussed in Section 4.4.3.2.

5. Synthetic Free-Layer Spin Valves. In this type of spin valves [90], the free layer consists of a synthetic ferrimagnet (or antiferromagnet) in the form of Fa/Ru/F b, similar to the ones used in synthetic pinned layers, where F a and F b are ferromagnetic materials (Permalloy, Co, CoFe). The advantage of such a free-layer structure is the reduced self-demagnetizing field due to flux closure within the free layer and hence

Fig. 47. Schematicdiagram of a symmetric spin valve. The pinned layer SF has a layer structure of F/Ru/F. The two F layers are antiferromagnetically coupled through the very thin Ru layer.

reduced magnetostatic interaction between the free layer and the pinned layer. The reduced self-demagnetizing field of the free layer is important to maintaining the field sensitivity of submicron-sized spin-valve sensors, which are discussed in Section 4.4.3.2. It was claimed [90, 170] that the used of a synthetic free layer can reduce the effective magnetic thickness of the free layer (tef f - - ( M a t a + Mbtb)/MNiFe) while maintaining the large physical thickness.

6. Specular Spin Valves. A specular spin valve may have a layer structure similar to that of any of the above, but with an insulation antiferromagnetic layer as a specular reflection layer. In spin valves with metallic AF and capping layers, some of the conduction electronsare lost because of the shunting effect via the AF and capping layers. As a result, the number of electrons participating in the spin-dependent scattering (contributing to the GMR effect) is reduced. An insulation AF layer functions as an electron mirror, reflecting (scattering) the conduction electrons reaching the Fz/AF interface back to the multilayer and preventing them from escaping the spinvalve multilayers. Specular reflection can be further enhanced by the use of an insulation capping layer [82] or a very thin (two monolayers) Au capping layer [83]. However, the specular reflection of electrons at the interfaces must not affect the electron spins. Otherwise the GMR ratio will not be enhanced. Egelhoff et al. have obtained a GMR amplitude of 15% at room temperature for NiO bottom spin valves with a 4-/k Au capping layer [83] and 24.8% for specular symmetric spin valves [81]. Sugita et al. [84, 85] have reported that a GMR amplitude of 28% was obtainable for specular symmetric spin valves with metal/insulator (ce-Fe203/Co and Co/ce-Fe203) interfaces. 7. Spin-Filter Spin Valves. Spin-filter spin valves [171-173] allow a reduction of the free layer thickness without a significant reduction of GRM amplitude. Therefore, ultrathin free layers can be used in such a spin-valve structure. As discussed in Section 4.2.5.3, if the free layer thickness is smaller than

HIGH-DENSITY MAGNETIC RECORDING A t (the longer of the two MFPs), the spin 1" electrons will penetrate the free layer and be scattered by the outer surfaces (substrate, AF layer, buffer layer, or capping layer), which results in smaller effective A 1", increased low effective resistivity (PL), or reduced SDS contrast and hence reduced GMR amplitude. The spin-filter spin valve uses an extra high conductance layer on the other side of the free layer (the free layer is sandwiched by the Cu conducting layer and the highconductance layer). Because of the low scattering rate of the extra high conductance layer, the A 1' of the spint electrons is no longer restricted by the free layer thickness. As a result of this, the GMR amplitude of the spin-filter spin valve does not decrease significantly as the free layer thickness is reduced to below the value of A t.

8. Dual Spin Valves. As in to the dual-element MR heads [9], various types of dual spin-valve structure can be formed, depending on the combinations of magnetization pinning directions, sensing current directions, types of spin valves, and the voltage-sensing mode. Figure 48 shows two typical types of dual spin-valve structures, namely the antiparallel and the parallel dual spin valves [92]. According to [92], both types of dual spin-valve heads are thermal noise immune because of differential voltage sensing. The thermal noise immunity of heads is accomplished by connecting a differential amplifier to the output terminal of each spin valve. The antiparallel dual

531

spin-valve head exhibits greater linear dynamic range than a single spin-valve head and an almost doubled sensitivity of a single spin-valve head. It also has an ideal transfer curve passing through the origin that is independent of the bias state of its single spin valves. However, the drawback of dual spinvalve heads is that they are three terminal devices and are therefore more expensive to fabricate. In spin-valve design, it is possible to combine various types of spin-valve structures to form a new spin valve to meet the design requirements. For example, a GMR amplitude of 15.4% was obtained for synthetically pinned symmetric spin valves with the composition Ta/Seed/IrMn/CoFe/Ru/CoFe/Cu/CoFe/ NiFe/CoFe/Cu/CoFe/Ru/CoFe/IrMnffa, which was used in the demonstration of a 20 Gb/in 2 recording system [80].

4.4. Spin-Valve Head Engineering 4.4.1. Considerations of Spin- Valve Head Design The major considerations for the design of high recording density read heads are as follows [86, 87].

1. Pulse Amplitude. The desired pulse amplitude of the reproduced signal from the increasingly smaller recorded bits must be maintained as the recording density increases. Pulse amplitude is normally measured as the track averaged amplitude (TAA), which is simply given by Ohm's law [87], AV=IsAR

(4.21)

where I s is the sensing current and AR is the sensor magnetoresistance change. To achieve high pulse amplitude from the increasingly smaller flux from the media, spin valves with a higher GMR ratio and higher field sensitivity are required. Unfortunately, the spin-valve field sensitivity decreases with the reduction of sensor size (or increasing t/h or t/w ratio) because of the enhanced magnetostatic interactions; this puts further demands on the increase in GMR ratio of spin-valve films.

2. Pulse Amplitude Symmetry. It is important to have equal amplitude of the positive and negative output pulses. To achieve this, the read sensor must be properly biased to have a linear transfer curve with a bias point at the origin of the coordinator system, as shown in Fig. 50.

3. Stability. Stability issues for spin-valve heads include magnetic domain stability and thermal stability. Magnetic domain stability is achieved by hard magnet bias, and thermal stability is optimized by the choice of antiferromagnetic exchange materials, which requires AF materials of high pinning field, high blocking temperature, and high temperature stability.

Fig. 48. Schematicdiagram of the antiparallel (a) and parallel (b) dual spin valve structure. Reproducedwith permissionfrom [92], copyright 1999,IEEE.

4. Lifetime Reliability. The lifetime reliability of a spin-valve head is determined by the operating temperature of the heads, which in turn depends on the joule heating effect due to I2R. Low sensor resistance and sensing current are therefore preferred options.

532

PAN

5. ESD (Electrostatic Damage) Immunity. The ESD immunity of a spin-valve head decreases with decreasing sensor size. For example, the ESD threshold voltage for an areal density of 1 Gb/in 2 is about 15 V, but it is reduced to about 1.5 V for 10 Gb/in 2 heads [87]. For submicron-sized sensors, ESD becomes a very serious problem. ESD can cause pinned layer reversal, sensor interdiffusion, amplitude loss, and physical damage of the sensor [86].

=9

r

AR W (cos(0, - 02) ) AR ~ - - R s h t l , R h 2

(4.22)

where A R / R -- ( R t + - Rsat)/Rsa t is the intrinsic magnetoresistance of the spin valve as measured on infinite samples, and Rshl,t is the sheet resistance measured in the parallel magnetic state. W is the length of the sensor active region (track width), and the notation ( . . . ) denotes averaging over the sensor height, h. An ideal transfer curve of a spin valve with infinite size is shown in Fig. 50. Because the uniform external field acts on the magnetic hard axis of the sensor, the response of the sensor to the field is linear. Here we only consider the ideal spinvalve, i.e, the magnetic moment of the free layer is at a fight angle with the pinned layer when the external field is zero, and the ferromagnetic coupling field and the demagnetizing field are both zero. In such a case, the linear transfer curve will pass through the origin of the coordinate system and is equidistant from saturation in both positive and negative fields. The sensor exhibits minimum resistance (parallel magnetic configuration)

Fig. 49. Schematicdiagram showing the magnetization orientations in the free and pinned layers of a spin-valve sensor.

uoe

0.5

|

0 0

"~

-3

-2

....... Low R

The magnetization orientations of free and pinned layers of a spin-valve read sensor in an external field H are schematically shown in Fig. 49. In spin-valve replay heads, the free layer is usually a uniaxial anisotropy film with easy axis aligned longitudinally (parallel to the ABS). The magnetic moment of the pinned layer is pinned along the vertical axis, which is ideally at a fight angle with the easy axis of the free layer. This initial easy axis alignment is achieved during film deposition or postdeposition annealing. For such a magnetic configuration, the sensor magnetoresistance is proportional to the cosine of the relative angle between the magnetization vectors in the free and pinned layers and is given by [205]

,ooOOOO

A

V

4.4.2. Basics of Spin-Valve Replay Heads: Sensor Bias Point

1

A

rllgn K

-

3 ~

11 | .,..J

Bas

5.

"

Uniform field, H/Hk Fig. 50. Ideal spin-valve transfer curve for an infinite sensor in a uniform hard axis field (solid line) [75], and transfer curve due to the effect of nonuniform stray field from the pinned layer (dotted line).

when saturated by a negative field and maximum resistance (antiparallel magnetization configuration) when saturated by a positive field. Such a sensor is intrinsically linear, and the linear dynamic range of such an ideal sensor is ( c o s ( 0 1 02)) = 4-1, or A(cos(01-- 02)) = 2 [75]. In a real spin-valve film, the free layer is acted on by the ferromagnetic coupling field from the pinned layer He, the stray field from the transversely saturated pinned layer 1-Is, and the field generated by the sensing current H I [205]. Fig. 51 is a schematic illustration of various fields acting on the free layer for a spin-valve replay head. Both H e and I-Is originate from the pinned layer. These two fields are always opposite in direction and therefore are partially canceled out by each other. As discussed in Section 4.2.5.4, a e originates from the RKKY interlayer coupling, is nonmagnetic spacer thickness dependent, and may exhibit oscillatory behavior. I-If must be optimized to an acceptable value (typically H e < 20 Oe) during the spin-valve film deposition process, so that the free layer is

Fig. 51. Schematicillustration of the magnetic fields acting on the free layer of a spin-valve head.

HIGH-DENSITY MAGNETIC RECORDING

Uoffset = Hr +Hs

,u~

i

12

+HI

t8

[/--'~0 = n f "l'ns \

-30

-20

] A

f

....... 4j_ . . . . . . . . . . . . . . . . . . . . .

-~0r

533

0

10

T...............

20

ij'~ . . . .- 4 .- - . . . . . . . .

30

4

I . . . . . . . . . . . ~..................

~

:

i

1

i

:

I

-12 '

'

'

I

H/Hk Fig. 52.

R(H) curve of a patterned spin-valve sensor. The offset field Hoffset = Hf + H s + H t and Hoffset1/~0 ---

reasonably decoupled from the pinned layer. H s is generated by the stray flux from the transversely saturated pinned layer. Its value depends on sensor height, thickness, and saturation magnetization of the pinned layer and thickness of the free layer. The value of H I depends on the magnitude of the sensing current and is given by Ampere's circuital law, H I = f I dl. The direction of the field H I , which depends on the direction of the current, can be in the direction of I-If or I-Is . Because of the effect of these fields on the magnetization orientation of the free layer, the R(H) transfer curve of a real spin-valve film is usually offset from the origin. If the sensor size is reduced to the submicron scale, the effect of these fields on the offset of the transfer curve of the free layer will be pronounced. Figure 52 shows a typical R(H) curve of a spin-valve sensor. The offset field Hoffset of the R(H) curve represents the vector summation of the three fields. The value of the offset depends not only on the magnitude of the sensing current during measurement, but also on the direction of the current. If the measuring sensing current is negligibly small, the offset field is equal to the vector sum of I-If and I-Is . In spin-valve head design, sensing currents with different magnitudes and in different directions are employed to balance out coupling fields from the pinned layers to obtain transfer curves with zero offset, i.e., to make Hf+I-I/+H s = 0. However, this is only feasible if the sum of H e and I-Is is not so large, because the maximum sensing current density that can be used in a spin-valve head is about 6 x 107 A/cm 2 [98]. Too great a current leads to overheating of the sensor, which causes reduction of the GMR ratio and destabilization of the pinned layer magnetic configuration. Too large a current will also produce a strong H I in the pinned layer opposing the pinned magnetization, as shown in Fig. 51, leading to a reduced pinning effect [110]. The stray field I-Is from the pinned layer is highly nonuniform [205]; it is usually not possible to achieve 01 = 0

I-If

-11- I-I s 9

everywhere over the entire free layer. 01 will be significantly nonzero at some points of the sensor [205]. This will cause a reduction of the linear dynamic range of the sensor [75]. However, within a considerable dynamic range, the spin-valve sensor is still linear. This is shown in Fig 50 by the dotted curve.

4.4.3. Effect of Finite Sensor Size 4.4.3.1. Effect of Self-Demagnetizing Fields Similar to the finite-sized AMR sensors discussed in Section 4.1.3, self-demagnetizing fields also exist in spinvalve sensors in both the vertical and horizontal directions of pinned and free layers. Equations (4.8) and (4.9) can be used for estimation of the average values of the vertical and horizontal self-demagnetizing fields of spin valves. But the self-demagnetizing fields are highly nonuniform across a spinvalve sensor and are much larger at the edge than at the center [9]. The self-demagnetizing field in the pinned layer tends to rotate the pinned magnetization from its original direction (perpendicular to the ABS) to the longitudinal direction [95]. This effect is pronounced when the sensor t/w or t/h ratio increases (sensor size decreases), which will lead to amplitude asymmetry of the replay waveform. Therefore, for smaller sensor sizes, AF materials with higher exchange pinning fields are required. The use of a synthetic pinned layer is another effective solution to this. As discussed in Section 4.3.3, the self-demagnetizing field of the pinned layer can be significantly reduced because of the flux closure within the two antiferromagnetically coupled layers. The major effect of the self-demagnetizing field in the free layer is on its transfer curve. As can be seen from Eqs. (4.8) and (4.9), the average self-demagnetizing fields are proportional to the sensor thickness and inversely proportional to the sensor width (for a horizontal demagnetizing field) or sensor

534

PAN

height (for a vertical demagnetizing field). Because the thickness of the free layer of a spin-valve head is much smaller than an AMR film (about 1/5 to 1/10 of the thickness of an AMR film), for large-sized spin-valve sensors, the selfdemagnetizing field is not significant. However, as the recording density increases, the spin-valve sensor size becomes increasingly smaller, and the effect of the self-demagnetizing field will become serious if the ratios t/h and t/w are not kept constant. For example, if a spin-valve sensor has a free layer thickness of 4 nm and a 47rM s of the free layer of 10 kG, the estimated average vertical demagnetizing field of the sensor from Eq. (4.8) is about 20 Oe and 100 Oe, for sensor heights of 2 /xm and 0.4 /xm, respectively. The self-demagnetizing fields of the free layer will have two major effects on the transfer curves of a spin-valve sensor, the linear dynamic range and the field sensitivity. This is schematically illustrated in Fig. 53. Because of the self-demagnetizing field, more external field is required to rotate the magnetic moment of the sensor to the same angle, and a large saturation field is required to magnetize the sensor to saturation, which means that the field sensitivity is considerably reduced as the sensor size decreases. The self-demagnetizing field has a nonuniform spatial distribution and is much larger at the sensor edge than at the center. This causes the center of the sensor to saturate first. At lower fields, the sensor linearity is still preserved. However, because the edge magnetic moments are more difficult to saturate, the transfer curve is not linear when approaching saturation. This causes a reduction in the linear dynamic range of the sensor, i.e., A(cos(01 - 0 2 ) ) <~<~2, or the maximum linear AR/R is less than the value obtained from a sheet spin-valve film. To compensate for the loss of field sensitivity and maximum linear AR/R, the required GMR amplitude of spin-valve films increases with the reduction in sensor size (increase in the t/h ratio).

Ideal transfer curve ^

~"

I

I

[

I

i;

I

I

I

I

i/-

I0

40/

o 40

i I

!l

Ii |

" / " ~ ~ ~

/t

t0

~

Transfer curve due to f'mite sensor size

30 i

4

I-I/Hk Fig. 53. Schematic illustration of spin-valve transfer curves showing the effect of a self-demagnetizing field due to the finite size of the sensor. The maximum linear dynamic range of A R / R is reduced as marked by the two dots on the curve (1/2 slope points [75]). Perfect bias is assumed.

4.4.3.2. The Use of Thinner Free Layers Equations (4.8) and (4.9) imply that one of the effective ways of reducing the average self-demagnetizing field is to keep t/W or t/h down or at a constant value. This means that as the sensor dimension (W, h) decreases, the thickness of the free layer must also scale down. However, the reduction of free layer thickness will also cause lower GMR amplitude and higher H s arising from the pinned layer. As discussed in Section 4.2.5.3, the GMR amplitude of spin-valve sandwiches decreases with the reduction in free layer thickness (as shown in Fig. 39). This is due to the fact that when the free layer thickness is smaller than h 1' (the longer of the two MFPs, typically ~ 50,~ [96]), the spin 1" electrons will penetrate the free layer and be scattered by the outer surfaces (substrate, AF layer, buffer layer, or capping layer), which results in increased low effective resistivity (PL) or reduced SDS contrast (a = At/A ~ = P~/Pt) and hence reduced GMR amplitude (see Eq. (4.18)). Various new spinvalve structures, such as spin-filter spin valves, spin valves with ferrimagnet free layers, and specular spin valves, have been employed to overcome this problem. The second problem of using a thinner free layer is the increased H s arising from the pinned layer. For a given pinned layer, the magnetostatic stray field in the free layer from the pinned layer increases with the decrease in free layer thickness (the same flux passing through a smaller cross-sectional area). The increased H s makes it difficult to obtain an optimum bias point for submicron-sized sensors in conventional spin valves. Spin valves with synthetically pinned layers have been successfully used to solve this problem [93, 110]. With the use of a synthetic pinned layer, the H s from the pinned layer can be significantly reduced. Leal and Kryder [110] have experimentally measured the offset field of synthetic spin-valve sensors with track widths of 10 and 1 k~m, as a function of sensor height h, and compared the results with conventional spinvalve sensors. As shown in Fig. 54, for sensors with relatively large sizes (h > 4/xm), the effective coupling fields of the two types of spin valves are quite similar and are on an order of magnitude of smaller than 20 Oe. As the sensor height decreases, the effective coupling field of spin-valve sensors with a conventional AF exchange bias system increases considerably, reaching 125 Oe at a sensor height of 1 /zm. It is not possible to obtain a zero offset bias point for such spinvalve sensors if only the field from the sensing current is used to cancel out the offset field, because the maximum sensing current is limited by the joule heating effect and by the ESD immunity requirement. On the other hand, the effective coupling field of spin-valve sensors with synthetic pinned layers remains almost constantly below 20 Oe, even when the sensor height is reduced to 0.5/xm. The advantages of synthetic spin valves over conventional spin valves in the application of submicron-sized sensors can be clearly seen from this work. A zero offset bias point for such a sensor can easily be obtained by varying the sensing current values. This is shown in Fig. 55 by the two R(H) loops of a spin-valve sensor with a 10/xm

HIGH-DENSITY MAGNETIC RECORDING T .... A

1,,~

i :-

-r-, ~ . .~- ~..,,_

.

,

535

M (arb. u.)

6

tip

O

"--

J

I

-50

I

HUA 1 9

tI

r J

E r

I

-100 o

tO

2,10 -150

'--

' 2

0

"'

'-4

'

' . . . . . 6 8

' 10

12

h (p,m) Fig. 54. Effective coupling field measured as a function of stripe, height for a synthetic spin-valve head with W = 10 mm (open dots) and W = 1 mm (solid square). The dashed line shows the expected effective coupling field as a function of stripe height for simple spin valves. Negative values are associated with antiferromagnetic coupling. Reproduced with permission from [110], copyright 1999, IEEE.

track width and a 0.5-/zm sensor height, measured at two different sensing current values. The R(H) loop with a current density of 3.3 x 107 A/cm 2 shows a zero offset field and therefore is a properly biased transfer curve. The slight reduction in the GMR ratio at higher current density is due to the heating effect of the current, which leads to an increase in sensor temperature and hence an increase in film resistivity and a decrease in the GMR ratio [110].

-I-"

tl 4~)0

6(/

i /I

_22

, 30

If H (Oe)

Fig. 56. Schematic illustration of the MH loop of an antiferromagnet/ferromagnet exchange-biased system, together with the definition of Hug.

offset from the origin of the MH loop of the exchangebiased system, as shown in Fig. 56 for an antiferromagnet/ferromagnet bilayer system. This field pins the magnetic moment of the ferromagnet to saturation in one direction. It is therefore also known as a pinning field, npin. Other names used for the same field include bias field, exchange field, or exchange pinning field, nua is a convenient term to use because it describes the interface exchange coupling strength in terms of an effective magnetic field strength. However, nua is ferromagnetic thickness dependent., which decreases with increasing ferromagnetic thickness.

4.4.4. Optimization of AF Exchange-Biasing Materials 4.4.4.1. Basic Parameters olAF Exchange-Biased @stems An antiferromagnet/ferromagnet exchange bias material is characterized by the following parameters: a unidirectional exchange anisotropy field (Hua), interfacial coupling energy (Jk), a blocking temperature (TB), and a critical thickness (6c).

Unidirectional Exchange Anisotropy Field (nua). The unidirectional exchange anisotropy field nua (Oe) is a measure of the exchange coupling strength of the antiferromagnet/ferromagnet interface. It is measured by the unidirectional

r

'

'

""

9

'

' I

,

, r ,, T_

,

6

2 0 0

describing the interface exchange coupling strength is to use the interfacial coupling energy Jk (erg/cm2) 9 Because Jk is a term of surface energy density, it is simply equal to the product of Hua, M s, and t F and is given by [109] Jk = nuaMstF

(4.23)

where M s and t F a r e the saturation magnetization and thickness of the ferromagnetic layer, respectively. The value of Jk does not depend on the ferromagnetic thickness.

Blocking Temperature ( Ta) and Thermal Stability. The

,

-" - 5x10 A/cm

-600

Interfacial Coupling Energy (Jk)" An alternative way of

600

H (Oe) Fig. 55. R(H) loops for an unshielded spin-valve head with a sensor width of 10 mm and a height of 0.5 mm, measured at two different current densities. Reproduced with permission from [110], copyright 1999, IEEE.

exchange anisotropy pinning field of an antiferromagnet/ferromagnet system usually decreases with increasing temperature and can be plotted as a temperature dependence curve as shown in Fig. 57. The blocking temperature is the temperature at which the exchange pinning field becomes zero, which mainly depends on the crystal phase structure of the material. The blocking temperature (TB) is one of the measures of the thermal stability of an exchange-coupled system. However, a high blocking temperature does not necessarily mean a good thermal stability. The blocking temperature distribution is more closely related to the thermal stability of the exchange-coupled system. There are various ways to characterise the blocking temperature distribution of an AF/F system [ 113, 114]. Here we introduce the method used by Nagai et al

536

PAN

800

200 180

700 600

160

w 9149149 9149149149149149

S

140 500

~" 120

9-~ 400

,<

Blocking temperature

89 O9

300

r

200

..

100

I

80

I I

60

I

I

'/

40

100

'

20

8c

1

0 0

50

100

150

200

250

300

350

400

0

50

100

150

200

250

300

350

AF layer thickness (~)

Temperature ("C) Fig. 57. Temperature dependence of the exchange pinning field and definition of the blocking temperature Tb.

Fig. 59. Schematic illustration of the definition of critical thickness 6 c of AF films.

[113]. Samples are heated to a pre-set temperature in a magnetic field (typically 3 ,~ 10 k Oe) applied against the original pinning direction, soaked at that temperature for 10~ 15 minutes, and then cooled down to room temperature for magnetic measurement. The experiment is repeated at different soak temperatures up to the blocking temperature. The measured exchange pinning field is normalised by the original value of Hua and plotted versus the soak temperature, as shown in Fig. 58. The differential of the curve represents the blocking temperature distribution. Because there is a large field applied to the sample against its pinning direction during annealing, part or all of the pinning field will be reversed as the temperature increases. As shown in Fig. 58, if the curve stays at the value of the normalized pinning field of 1.0 at temperature T, the original pinning field is stable at that temperature. As the temperature increases, the curve will start to fall off, changing sign, and, finally, at the blocking temperature, reaching - 1 . 0 of the normalized field value, which means a complete reversal of the pinning field. The thermal stability of the pinning field can therefore be characterized by the peak temperature Tp of the blocking temperature distribution curve together with

its half peak width Wso. For good thermal stability, an AF/F system must have a small Ws0 and a high Tp.

2 1.5

-

Wso .

.

.

\

.

0.5 O

~.

/

0

N -0.5

i,

[[*, 9 [ ":/

010

)

...........

1 10

9-

-r .....

2t~

...... 3(10

,~

~ Z

-1 -1.5

TB distribution curve normalised differential of the solid curve.

~

Critical AF Layer Thickness (6c). The exchange pinning field is AF layer thickness dependent [115]. The thickness dependence usually has the general tendency shown in Fig. 59. For a particular AF/F exchange coupling system, there is always a minimum AF layer thickness, below which the exchange pinning field will start to fall off. This thickness is termed the critical thickness (6c) of the AF material. 4.4.4.2. Requirements for AF Materials for Spin-Valve Heads Antiferromagnetic films play a vital role in spin valves. The basic requirements of AF materials for spin-valve head applications are high Hua (or high Jk), high TB, high thermal stability, small 6c, high resistivity, high corrosion resistance, and ease of fabrication. The required nua for the desired output and linearity of spin-valve heads is sensor size dependent. For a sensor size of 1/zm, the required minimum nua is about 300 Oe [98] at sensor operation temperature. Smaller exchange pinning fields cause read-back waveform distortion [98], and make the output smaller [99]. High TB and high thermal stability are important for ensuring the temperature stability of the spin-valve heads at operating temperatures (typically 150~ A good thermal stability is also essential for improved ESD immunity. A small critical thickness 6c of the AF layer is required to fit the spin-valve element into the increasingly smaller read gap at high densities. For example, the projected read gap length for an areal density of 40 Gb/in 2 is 50 nm if the track width of the head is 0.2/zm [97], and the required 6 c of the AF material must be smaller than 100 ,~.

4.4.4.3. Properties of AF Materials Soak Temperature (~ Fig. 58. Soak temperature dependence of the normalized exchange pinning field (solid curve) and its Tb distribution (dotted curve).

Various AF materials have been studied [97-108]. Figure 60 presents the temperature dependence curves of the exchange pinning fields of some typical AF materials [111]. Typical

HIGH-DENSITY MAGNETIC RECORDING

1000 900 i

800 700 " " ~ -- ,JrMn -~ " ~

600 u

-,,yln

500

............

""~

.... .,

'~'~

~

..... "';-~ OtlkA 2

NiNln

400 "%

300

~

~

~

'~

"'.

\ C r P t ~ in "-.

200 FeMn

100

"..._

~

~ q

0 0

50

100

%

150

\

""

,~ \ ',

200

~,

250

".

\

"-,

300

350

400

Temperature ("C) Fig. 60. Temperature dependence of exchange pinning field for some typical AF materials. Sketched base mainly on the data from [111], copyright 2000, American Institute of Physics.

properties of AF materials commonly used in spin-valve heads are also listed in Table IV. For a modern disk drive, the operating temperature of the spin-valve heads is typically around 150~ and the minimum required exchange pinning field is about 200-300 Oe [87]. It can be seen that NiMn is the best AF material discovered so far in terms of JK, TB, and thermal stability. However, the antiferromagnetic 0-NiMn phase is obtainable only after postdeposition annealing [104, 105] at high temperatures and o~er a long period, which is not a preferred process for spin valves, because of its detrimental effect on multilayer interfaces and AR/R ratio [79]. PtMn is the second best material but is easier to fabricate than NiMn systems. Both AF materials are suitable for use in current spinvalve heads with areal densities up to 10 Gb/in 2. But their 6c must be considerably reduced to be used in future generation spin-valve heads. Oxide AF films such as NiO and CoO have the advantage of high corrosion resistance, zero shunting loss, and enhanced AR/R due to specular reflection [100]. However, the JK and TB of the existing oxide AF materials are too

537

low. The properties of AF materials depend on material, structure, underlayer, grain size, AF layer thickness, fabrication process, etc. Elemental AF materials have been well studied. Considerable work has been done for binaries. However, only very limited cases for ternaries and beyond have been studied [105]. The major technical challenge for future AF material will be to obtain high JK and Ta and other desired properties at very small 6 c. As discussed in Section 4.3.3, exchange-bias systems with synthetically pinned layers exhibit much better properties than AF/F bilayer systems.

4.4.5. Hard Bias As in to AMR heads, horizontal biasing is necessary for spinvalve heads, so that a single domain state is obtained in the free layer of the sensor and the single-domain state is stable against all reasonable perturbations [9]. The most commonly used horizontal biasing technique for spin-valve heads is permanent magnet hard biasing, as shown in Fig. 61. Patterned high-coercivity magnetic strips (CoCrPt, for example) are abutted against the ends of the spin-valve sensor, which provides a magnetic field that magnetizes the sensor along the horizontal direction to a single-domain state. The magnitude of the field produced by the hard magnet bias film depends on the Mrt of the film. Figure 62 shows the effect of relative layer thickness of hard bias films on spin-valve transfer curves, where the relative layer thickness RL is defined as Mrt RL

(4.24)

--

Ms,Ftv

Schematic illustration of the permanent magnet hard bias technique.

Fig. 61.

Table IV. Typical Exchange Systems Used in Spin Valves and Their Properties Exchange systems Ta/NiFe/FeMn [64, 106] NiO/NiFe/Co [100, 106] Ta/NiFe/IrMn [97, 107] Co/CrMnPt [97, 101] NiFe/RuRhMn [103] Ta/NiFe/PtMn [97] Ta/NiFe/NiMn [ 104, 106]

JK (erg/cm 2)

0.13 0.09 0.15 0.19 0.17 0.2 0.24

Hua (Oe) 420 (/NiFe = 4 rim) 360 (tNiFe = 4 nm) 400 (tNiFe = 2 nm) 320 (/Co = 3 nm) 350 (/NiFe = 2.5 nm) 600 (tNiFe = 3 nm) 650 (tNiFe = 3 nm)

TB

(~

6c (nm)

Corrosion resistance

Thermal stability

Annealing

Poor

Poor

No

Very good

Poor

No

Good

Acceptable

Yes

30

Very good

Acceptable

Yes

250

10

Good

Good

No

340

30

Good

Very good

Yes

380

30

Good

Very good

Yes

150

7

200

40

280

8

320

538

PAN . . . . . . . . . . . . . . . . . . .

i

I

~6- t 0.4 [

- 11.O.,,fT-

~

-0.8

.

.

.

.

~

1

,i 90

.

.

.

.

.

.

.

.

.

.

.

.

.

.

.

.

.

.

RL=0.5

,-It- RL=1.0

010

1,0

2i0

i

.

,

-e- RL=2.0

Medla_MrT

Fig. 62. Effect of hard bias layer thickness on the transfer curves of spin-value sensors. Reproduced with permission from [86], copyright 1999, IDEMA.

where Ms,F and tF are the saturation magnetization and thickness of the free layer, respectively. RL = 1 when Mrt of the hard biasing film is equal to Ms,FtF of the free layer. As can been seen from Fig. 62, when RL = 0.5, the single-domain state is not formed in the free layer, and the transfer curve is affected by the moving domain walls in the free layer. When RL = 1, a smooth transfer curve is obtained, which indicates that the free layer is in a single-domain state. However, the bias field also causes a reduction of amplitude of the transfer curve. A further increase in the thickness of the hard bias layer (RL = 2) causes further reductions in the amplitude. This is due to the fact that the excessive bias field will lock the free layer in the bias field direction and the free layer will lose its sensitivity to the field from the recording media.

5. F I L M S F O R W R I T E H E A D S

a magnetic material. The hard axis usually exhibits a straightline loop, reflecting the fact that the magnetization reversal process along the hard axis is predominantly coherent rotation rather than domain wall motion. The effective anisotropy field (Hk) of a soft magnetic film is defined from the hard axis loop, as also shown in the figure. In thin film heads, the magnetic film is so oriented that the magnetic flux path is always along the hard axis. The uniaxial anisotropy of a soft magnetic film can be obtained by applying a magnetic field during the film deposition or annealing process. The origin of the induced anisotropy is the short-range directional ordering, in which atomic pairs in a film tend to align with the local magnetization. A magnetic field that is strong enough to saturate the magnetization of the film (typically 50-100 Oe) is required to induce uniaxial anisotropy.

5.1. Basics of Soft Magnetic Films for Writers

5.1.1. Hysteresis Loop, Coercivity Hc, and Anisotropy Field H k Soft magnetic films used for film heads are usually required to have an in-plane uniaxial anisotropy. The typical easy and hard axis BH loops of a uniaxial anisotropy film are shown in Fig. 63. The BH loop along the easy axis is a nearly rectangularshaped open loop, which reflects the fact that the magnetization reversal process along the easy axis is predominantly domain wall motion. The area S (= f H dB) enclosed by the loop represents the hysteresis loss per cycle in unit volume material (joules/m 3 cycle). For most magnetic applications it is always desirable to make the hysteresis loss as small as possible. The coercivity H c is a measure of the openness of the loop and therefore is the representation of the softness of

is j

1 .D r s

axi o H ./i

~.

Ho /

m .: 0

-10

A

,,~

10

t #

i 1 E

H (Oe)

Fig. 63. TypicalBH loops of soft magnetic films.

HIGH-DENSITY MAGNETIC RECORDING

5.1.2. Permeability and Its Frequency Dependence

Y

/

The permeability/z of a soft magnetic material is defined as B

].s = ]./,0]./,r-- ~

539

(5.1)

Bm

v

where /z 0 is the permeability of free space (/z 0 - 1 in CGS units), ].Lr is the relative permeability, B is the flux density (Gauss), and H is the magnetic field (Oe). It can be seen that /.~-/'/'r in CGS units. Because of the nonlinear nature of the hysteresis loop, the permeability of a magnetic film varies with the magnitude of the applied field. Terms like initial permeability (/'/'i = dB/dHll4=O,B=o), maximum permeability (/~max- (B/H)max), and differential permeability (/x d = dB/dH) are frequently used to describe the field dependence of permeability. In a high-frequency magnetic field, the permeability of a magnetic thin film is a complex number. If the high frequency field, given by H-

H mcos tot -- n m eit~

(5.2)

is applied to a magnetic thin film, the flux density B of the magnetic film will lag behind the field H by a phase of because of the effect of hysteresis, eddy current, magnetic viscosity, etc. and can be represented by B

--

B m

c o s ( t o t - 6)

=

Bm ei(~

(5.3)

According to the definition of permeability of Eq. (5.1), we have

fz = - -B- - ~ e Bm H Hm

-i6

=

[.1,t

- itz"

(5.4)

where/z' and ~" are the real and imaginary parts of the permeability, respectively, and are given by ~,

-

Bm

~

Hm

cos ~

(5.5)

/z" -- Bm sin 6

(5.6)

Hm

v

X

Fig. 64. Schematicillustration of the phase relationship between B m and Hm at high frequency.

and

lf0

P~oss = -~

H . dB - 7rftz"H2

(5.9)

The ratio of/x' and/z" is known as the quality factor of the soft magnetic material. Figure 65 is a schematic illustration of a typical frequency spectrum of the hard-axis complex permeability of a soft magnetic film [74]. In considering the effect of frequency on permeability, it has been assumed that the true permeability /z is independent of the frequency of the field and that only the apparent permeabilities (/z' and/z") vary with frequency [216]. The permeability spectrum in the frequency domain can be divided into five zones. In zones I and II, the change in/z' and /z" with frequency is insignificant. In zone III, the real part of the permeability starts to drop and the imaginary part starts to rise. The primary reasons for this could be the effects of hysteresis and eddy current damping. The resonance-type curve in zone III may be caused by the magnetic after-effect (or magnetic lag or magnetic viscosity). As the frequency further increases (zone IV), magnetic resonances (ferromagnetic resonance, domain wall resonance, and dimensional resonance) may occur, which result in a resonance-type spectrum. The corresponding frequencies for each zone may vary with different materials. The challenge of magnetic materials science and engineering is to extend zones I and II to frequencies that are as high as possible.

and

I~1- J~'~ + ~"~

(5.7)

Figure 64 is a schematic representation of the phase relationship between B and H. As can be seen from the figure, the real part of the complex permeability/z' is the ratio of the in-phase component of the amplitude of the flux density B m with the applied field and the amplitude of the applied field H m. The imaginary part of the complex permeability/z" is the ratio of the quadrature components of the amplitude of the flux density B m and the amplitude of the applied field H m. In fact,/x' is associated with the energy storage rate Wstorag e o f a magnetic material, and/z" is associated with the power loss rate/~ of the material. These are given by Wst~

lfrl -- -T J0 -2I - I . B d t -

--

1 , 2 2/xHm

(5.8)

Fig. 65. Typical curves of the frequency dependence of complex permeability.

540

PAN

5.1.3. Magnetostriction Magnetostriction is defined as the fractional change in length of a sample induced by the change in magnetization state of the sample. In the one-dimensional case, the magnetostriction coefficient A is given by Al

~=-y

(5.~o)

where Al is the change in length of a sample and l is the total length of the sample. )t can be positive or negative and is usually a very small number (typically on the order of magnitude of 10-6). The value of )t depends on the applied magnetic field (magnetization states). The saturation magnetoresistance As applies if the sample is magnetized from a demagnetized state to saturation. For single-crystal materials, the magnetostriction is direction dependent, and their magnetostriction constants usually refer to the corresponding crystal axis. For example, )tl00 and )till refer to the magnetostriction constants of a cubic single crystal along the (100) and (111) axes, respectively. If a material is polycrystalline, the magnetostriction of the sample along any direction is the averaged value of magnetostriction of the single crystals along that direction and is given by

X=

=SX~ cos~0-5

(5.11)

where 0 is the angle between the field direction and the measurement direction, and A0 is the magnetostriction constant when 0 is zero. For a polycrystalline material, )to is isotropic and is given by

A0 = 2A10o + 3Alll 5

(5.12)

5.2. Basics of Thin Film Writers

A recording head can be analyzed with a magnetic circuit model in a way similar to that in which an electrical circuit is analyzed. The reluctance of a magnetic circuit R m is the analogue of resistance R in an electrical circuit. It is therefore also known as magnetic resistance. Reluctance is defined as l /zA

5.2.2. Write Head Efficiency A magnetic recording head can be represented by a magnetic circuit model, with reluctances or permeances. In the simplest case, the head reluctance can be reduced to three reluctances: yoke reluctance Rmy, gap reluctance Rmg, and leakage reluctance Rml. And the magnetic head can be represented with the equivalent circuit, as shown in Fig. 66. The write head efficiency r/w is defined as r/w =

(5.15)

(I)y

Using Kirchhoff's law around the bottom loop gives ((I)y -- (I)g) gml -- (I)ggmg -- 0

(5.16)

From Eqs. (5.15) and (5.16), r/w =

Rml Rn~ + Rmg

=

1 1+

Rmg/Rml

(5.17)

It can be seen from Eq. (5.17) that a smaller gap reluctance results in a high write efficiency. It is not obvious from Eq. (5.17) that the write head efficiency actually depends on the yoke reluctance of the permeance of the head. In fact, the high permeance or the low reluctance of the yoke is critically important for high write efficiency, because if the permeance of the yoke is low, there will be little magnetic flux reaching the gap. Most of the flux will be lost in air before reaching the gap. If ~g is small, by Eq. (5.15) r/w will be small.

The expression of inductance of a coil with N turns can be derived from Faraday's law (e = N ~d4, ) and Lenz's law (e = L d~) and is given by d4,

L -- N ~ di

(5.13) 9..

where I is the length of the magnetic circuit or magnetic material,/z is the permeability, and A is the cross-sectional area of the magnetic material or magnetic circuit. The unit of reluctance, as can be seen from Eq. (5.13), is simply cm -1 in the CGS unit system. Like electric resistance, two or more reluctances in a series magnetic circuit add together. In analogy to conductance in an electrical circuit, the inverse of reluctance is known as permeance P and is given by

p=tZA l

~g

5.2.3. Head I n d u c t a n c e

5.2.1. R e l u c t a n c e a n d P e r m e a n c e

R m --

The unit for permeance in the CGS system is cm. Like conductance, two or more permeances in a parallel magnetic circuit add together.

,v

,I

(5.18)

Rmc

]

m.m.f. = Ni Rml

]

a~ '

.,,

R.~ (5.14)

Fig. 66. A simple magnetic circuit model for a recording head.

HIGH-DENSITY MAGNETIC RECORDING For a circuit with constant reluctance (or constant permeability), Equation (5.18)can be written as

L = N ~. l

(5.19)

Equation (5.19) implies that the inductance of a coil is a measure of its flux linkages per ampere. From the simple magnetic circuit model (as shown in Fig. 66) the current circulating in the coil of a head is given by

Ni = 4) y~ R m

(5.20)

where ~ R m is the total reluctance of a magnetic circuit. The inductance of a recording head is then given by N2 L=

(5.21) Rm

Head inductance can be divided into two parts [8]: the coil inductance in the absence of soft magnetic material in the yoke, and the magnetic yoke inductance contributed by the flux linkages due to the presence of the soft magnetic yoke materials. The coil reluctance varies with the coil diameter, and the magnetic yoke reluctance is proportional to the volume of the yoke material and inversely proportional to the separation of the two poles [8]. At high frequencies (high data rate recording), a high head inductance results in a high head impedance, which is not desirable because it requires a high power voltage supply.

541

whole head, showing the pancake-shaped coil structure (C), the electrical contact leads of the coil, and the top magnetic pole. The pole tip structure on the ABS surface of the head is shown in Fig. 67c. The top pole or trailing pole (P2) is slightly narrower than the bottom pole (P1). The writing track width is defined by the width of P2, which was 38/xm for the IBM 3370 heads. The gap length of the head was 0.6/xm. The IBM 3370 film heads were fabricated using the socalled wet process; i.e., metal films (Permalloy and copper) were deposited by electroplating. This wet process has since been used by most thin film head manufacturers because of its economical viability. A ceramic wafer of alumina and titanium carbide [218] was used as the head substrate, which was chosen because of its high yields of chip-free rails and good durability in start-stop operation [175]. The first layer deposited on the substrates was a thin film of alumina, which was used to provide a smooth surface for further layer processing and to insulate the heads from the slightly conductive ceramic substrate. The bottom Permalloy pole (P~) was then photolithographically patterned and electroplated, followed by the sputter deposition and patterning of the gap layer (D) (alumina) and a hard-cured photoresist insulation/planarization layer (E). The copper coil layer (C), together with the first coil lead, was then patterned and formed by electroplating, followed by another hard-cured photoresist insulation/planarization layer (E). Next, the top pole (P2), together with the second coil lead, was patterned and electroplated on top of the hard-cured photoresist insulation layer. The same

5.2.4. Structure and Microfabrication Process of Film Heads As has already been mentioned in Section 2, a thin film head consists of three principal functional parts, the pancake-shaped copper coil, the yoke (poles), and the front gap. They are all made of thin films. The function of the coil, which is sandwiched by the upper and lower poles, is to convert the electrical signals (write current) into magnetic fields. The yoke, which consists of two poles (P1 and P2) partially separated by the gap layer, the coil, and insulation layers, is used to form a low-reluctance (high permeance) magnetic circuit and to deliver the magnetic flux generated by the coil to the front gap. Only the stray field from the air-bearing surface (ABS) of the front gap is used for writing. The writing track width is defined by the strip width of the trailing pole (usually P2), and the writing bit length is determined by the front gap length, the linear speed of the media, and the write current pulse duration. A schematic drawing of an IBM 3370 thin film head is shown in Fig. 67 [175]. Figure 67a is an expanded crosssectional view of the lower half of the head (from the ABS to the center of the copper coil), showing the two poles (A), insulation layers (E), the eight-turn copper coil (C), and the gap layer (D). Both the copper coil layer and the Permalloy yoke are electroplated. Insulation layers between the coil and the poles are usually made of hard-cured photoresist. The most commonly used material for the gap layer (D) is sputterdeposited A1203 film [175]. Figure 67b is a planar view of the

.~176

., . , " .... : ' 9.1/ 9

PI

(~)

~--

P2 i

0

i

10 j~.m

tO

.:" ..', "~ .'.'~,

~)

, '.... 9 (']~) (D)

....,*"

,," :o

01):.'~..-~"~'-:. -7", I

(c)

". ,.

"-,

Air-

Fig. 67. IBM 3370 film head. (a) Schematic cross section showing the magnetic layers (A), pole tips (B), conductor turns (C), gap layers (D), and insulation layers (E). (b) Planar view of the film head. (c) Pole tip structure at the air-bearing surface. Reproduced with permission from [175], copyright 1996, IBM.

542

PAN Table V. IBM Inductive Film Head Products

3370 3375 3380 3380E 9335 3380K 3380K 9332 3390* Aptos* 3390-3* Tanba Turbo* Tanba-3* Ritz- 1 Ritz-2

3370 3375 3380 3380E 9335 3380K 9332 3390 Aptos 3390-3 Tanba Turbo Tanba-3 Ritz-1 Ritz-2

GA

TPI

KBPI

AD (Mb/ln.)

PI (/xm)

P2 (/xm)

G (/xm)

1979 1980 1981 1985 1986 1987 1987 1988 1989 1990 1991 1992 1993 1994 1994

635 800 801 1,386 1,600 1,600 2,089 2,017 2,242 2,242 2,984 2,436 3,041 4,000 4,000

12.1 12.1 15.2 16.2 16.2 16.2 15.2 23.6 27.9 27.9 30.0 59.8 61.0 80.0 80.0

7.3 9.7 12.2 22.5 25.9 25.9 31.7 47.5 62.6 62.6 89.4 146 186 320 320

1.6 2.O 1.7 1.6 1.6 1.6 1.6 3.0 0.9 0.9 0.9 3.0 3.0 3.5 3.5

1.9 1.9 2.0 1.9 1.9 1.9 1.9 3.0 1.0 1.2 1.2 3.1 3.1 3.5 3.5

0.60 0.7O 0.60 0.60 0.60 0.60 0.60 0.55 0.55 0.55 0.55 0.35 0.32 0.25 0.25

Plw

P2w

R

L

(/xm)

(/zm)

Turns

Layers

(11)

(nil)

38.0 27.5 29.5 15.0 13.5 8.5 10.5 8.0 8.0 5.6 8.4 6.4 6.3 6.5

34.0 24.5 26.5 12.0 11.0 6.5 8.5 8.0 8.0 5.6 8.4 6.4 4.8 5.0

8 8 8 18 18 31 31 31 31 37 44 44 45 36

1 1 1 1 1 2 2 2 2 2 2 2 3 3

7 7 7 15 15 24 24 24 24 31 39 39 29 20

80 80 80 350 350 650 800 800 800 950 1,400 1,200 800 500

Efficiency (LN3) 1.25 1.25 1.25 1.55 1.55 0.67 0.83 0.83 0.83 0.69 0.72 0.61 0.39 0.39

Slider length (nun)

Slider width (mm)

Slider height (mm)

4.0 4.0 4.0 4.0 4.0 4.0 4.0 4.0 4.0 4.0 2.5 2.5 2.0 2.0

3.2 3.2 3.2 3.2 3.2 3.2 3.2 3.2 3.2 3.2 1.6 1.6 1.6 1.6

0.850 0.850 0.850 0.850 0.850 0.850 0.850 0.850 0.850 0.850 0.425 0.425 0.425 0.425

*Ion milled. Reproduced with permission from [175], copyright 1996, IBM.

Permalloy layer as the top pole layer was used as the second coil lead just for process simplicity. The top pole has a physical contact with the bottom pole through a pre-patterned hole in the center of the head to ensure low magnetic reluctance of the magnetic circuit. The second coil lead was also made to have a good electrical contact with the central contact pad of the coil layer. The photoresist insulation layer (E) was limited to the region of the copper coil as shown in Fig. 67a, and consequently was not exposed, as the head was lapped to final throat height [175]. Thick (> 25/zm) and large copper connection studs were electroplated through photoresist masks to provide electrical connections for the head. The last layer of the film head is a very thick ( ~ 2 5 / x m ) alumina overcoat, which is necessary for chemical and mechanical protection of the heads during further wafer processing and use. The alumina overcoat over the studs was mechanically lapped away after deposition, exposing the copper for electrical connections [175]. The performance of thin film heads depends on various head parameters and head materials. A higher areal density and a high date rate of magnetic recording are achieved with narrower tracks and a smaller gap length. Techniques employed to achieve

higher recording densities and higher data rates with inductive head technology included the used of multilayer coil structures, thick poles, separate designs for reader and writer, FIB pole trimming, and novel yoke structures [175, 186]. Typical head parameters of IBM inductive head products from 1979 to 1994 are given in Table V. Detailed reviews of the technical evolution of film head technology can be found in [175, 186].

5.2.5. Factors Determining the Yoke Permeance of Film Heads It can be seen from Eq. (5.14) that the permeance of a film head is proportional to the permeability of the poles and their cross-sectional area and inversely proportional to the length of the magnetic circuit. Therefore there are three main factors determining the head permeance at high frequencies: the domain configuration, eddy current, and yoke length.

Domain Configuration and Control. One of the key factors in realizing high yoke permeance, particularly at high frequencies, is the control of magnetic domains of the yoke, particularly in the pole tip of the head. In film heads, the induced easy

HIGH-DENSITY MAGNETIC RECORDING

543

axis of the magnetic films of the yoke is always perpendicular to the flux path, so that the magnetic domains in the poles are aligned perpendicular to the flux propagation direction (easy axis domains), and the magnetization reversal process in the yoke is predominately coherent rotation rather than domain wall motion. This ensures that the permeability of the yoke is high at high frequencies and is essential for obtaining a linear transfer curve for the head. Details about domain configuration and control in thin film heads including the use of lamination to obtain an easy axis single-domain state and the effect of magnetostriction on head domain structures are discussed in Section 5.3.

Eddy Current. At high frequencies, the time-varying flux in magnetic materials will induce eddy currents in the yoke, which in turn induce magnetic fields opposing the original flux. The eddy currents act as a screen to prevent the magnetic flux from penetrating the magnetic material and consequently reduce the effective cross-sectional area of the yoke for flux carrying. The depth of a magnetic material at which the highfrequency magnetic field can penetrate is known as the skin depth (~), which depends on the frequency of the field (to), and the permeability (/g) and conductivity (o') of the magnetic material, and is given by

to/go"

(5.22)

As a result of the skin depth effect, the permeance of the yoke will decrease from the outer surface to inner depths reaching its minimum at the skin depth. To overcome the eddy current effect, soft magnetic films with low conductance are desirable. For very high-frequency heads, laminated yoke materials with insulation spacers are usually used. As discussed in Section 5.3, another major characteristic of lamination is that an easy axis single-domain state of the yoke can be obtained with proper lamination, which can further improve the highfrequency response of film heads.

Yoke Length. Film heads with two or more coil layers are typical examples in which the length of magnetic circuit is reduced to achieve high permeance. These heads are more efficient than the ones with a single coil layer, but at the expense of process complexity [ 175]. 5.2.6. Required Head Field Properties for a High-Density Inductive Writer The primary function of a writer is to produce a writing field strong enough to reverse the magnetization in the medium. Figure 68 is a typical spatial distribution of the longitudinal field component of a film head simulated by computer. The magnitude of the field, which depends on the deep gap field Hg of the head, must be sufficiently high to achieve effective writing on recording media. As discussed in Section 2.2, the deep gap field Hg of a film head must be greater than three times the medium coercivity H c for effective writing without

Fig. 68. A typical spatial distribution of the longitudinal write field of a film head, simulated by computer, its field mapping on a recording medium, and the resulting bit length, track width, and erase band.

pole-tip saturation. The continued increase in recording density can have two major impacts on head materials. One is the demand for higher medium coercivity to overcome the transition self-demagnetizing field (see Section 3.1) and to maintain the thermal stability of the recorded information (see Section 3.1.3). The other is that the increase in the data rate will substantially reduce the write head field duration, which makes the dynamic writing coercivity of the media even higher. Both require an increasingly high writing field, which is primarily limited by the saturation magnetization M s or the saturation flux density B s of the head material. High Bs head materials are therefore indispensable for meeting the requirements of high medium coercivity and high data rate. Sharp spatial distribution of the write field is also important for achieving high track and linear densities. The sharpness of the field distribution is improved by maximizing the derivatives o f the head field H x with respect to x (downtrack direction) and z (cross-track direction) at the point where H x = H e. A maximum 8Hx/rXIHx=nc will ensure a sharp written transition, and a maximum 8Hx/SZlnx=Uc will ensure a narrow erase band, which are of critical importance to high linear and track densities [8].

5.3. Magnetic Domain Configurations in Film Heads One of the key factors in realizing high yoke permeance, particularly at high frequencies, is the control of magnetic domains of the yoke, particularly in the pole tip of the head. In film heads, the induced easy axis of the magnetic films of the yoke is always perpendicular to the flux path, so that the magnetic domains in the yoke are aligned perpendicular to the flux propagation direction. This is essential for obtaining a linear transfer curve for the head and high permeability and hence high permeance at high frequencies. However, as a consequence of energy minimization of the magnetic system, as discussed below, typical domain patterns in a film head yoke may look like those shown in Fig. 69, where the majority of

544

PAN where 0 is the angle between M and H. E H is at minimum when 0 is 0 ~ and maximum when 0 is 180 ~ The demagnetization energy E d measures the interaction between the magnetic film and the demagnetizing field H d (= NdM, where Nd is the demagnetization factor), which originates from the magnetic charges arising from the divergence of the magnetization given by Eq. (2.12). The general formula for E d is given by E d = - f I-I~. dM = - N d

Fig. 69. Typical domain configurations in a film head. Reproduced with permission from [ 186], copyright 1994, IEEE.

the domains are aligned parallel to the easy axis of the film, but domains along the edge of the strips are aligned parallel to the edges. The possible domain configurations in film heads will be analyzed separately for the following two cases: single-layered poles and laminated poles.

f

M. dM

(5.25)

E~ is sample shape, size, thickness, and magnetic domain configuration dependent, and it is not always easy to calculate, particularly for multidomain samples. However, in most cases, it is one of the dominant factors determining the domain configuration of a magnetic film, which will be further discussed in Fig. 70. In film heads, the magnetic films are usually very thick, and therefore all of the domain walls are treated as Bloch walls, the wall energy of which is given by Ew = SwYBloch

(5.26)

where Sw is the total area of walls per unit of magnetic film area, and Tmoch is the total wall energy per unit wall area (wall energy density), given by [187]

5.3.1. Domain Configurations in Single-Layered Films In a ferromagnetic thin film, the atomic magnetic moment tends to have a perfect parallel alignment to minimize the exchange energy Eex, which is given by Eex -- - 2 A S 2 cos t~

(5.23)

where A is the exchange integral and A > 0 for a ferromagnet, S is the atomic spin angular momentum, and r is the angle between the neighboring atomic moments. The condition for minimum exchange energy is obviously r = 0, i.e., perfect parallel alignment of all atomic moments in the ferromagnet (single-domain state). However, as is well known, in a ferromagnetic material magnetic moments are broken into localized small-volume regions, known as magnetic domains. The domains are separated by domain walls. Only the atomic moments in each domain have a perfect parallel alignment. The formation of magnetic domains in a magnet is due to the energy minimization of the magnetic system. In addition to the exchange energy Eex, other energy terms also have to be considered, such as the magnetostatic e n e r g y Emag, the domain wall energy E w, and the anisotropy energy Ea, which include the uniaxial anisotropy energy E k and the stressinduced anisotropy energy or magnetoelastic energy Eela. The magnetostatic energy includes the external field energy E H (or Zeeman energy) and demagnetization energy Eo (or shape anisotropy energy). For a given ferromagnet with a magnetization of M s in an applied field H, the Zeeman energy is given by E H = - f H. dM = -MsH

d

cos 0

(5.24)

-2-3/~ 6

60)27r62MZ6+t

T00+

(5.27)

+

where To is the wall energy density of the bulk material (% = 0.1 erg/cm 2 for Permalloy), 6 is the total wall thickness, 6 o is the bulk wall width (6 o - 2 • 10 -4 cm for Permalloy), t is the film thickness, and M s is the saturation magnetization of the film. The uniaxial anisotropy energy E k is given by E k -- K u sin 2 0

(5.28)

where K u is the uniaxial anisotropy constant (K u = 1.5 • 103 ergs/cm 3 for Permalloy [187]) and 0 is the angle between the magnetization vector and the easy axis. The minimum E k is obtained when 0 is zero or 180 ~ i.e., when the domains are oriented along the easy axis. The stress-induced magnetoelastic energy Eel a is given by the following equation for polycrystalline films with isotropic magnetostriction constant As, and with positive (tensile) stresses o"x and O'y in the x - y film plane and induced easy axis along the y axis, Eela -- --23 hs(Orx cos 2 0 + O'y sin 2 0)

(5.29)

The angle 0 is the angle between the magnetization vector and the easy axis. The combined anisotropy energy term is E k -- [ g u - 3~s(O" x - O ' y ) ] c o s 2 0

(5.30)

HIGH-DENSITY MAGNETIC RECORDING

545

Magnetic charges Easy axis

i-,,--+~- 4

I

......

(a) Single domain Fig. 70.

180" wall

0" walls

(b) Open structure

(c) Closure domain

Typical magnetic domain configurations of a uniaxial anisotropy film.

Summation of Eqs. (5.24) to (5.30) gives the total system energy, and the minimum total energy condition gives the most stable domain configurations. However, because the values of each energy term are interrelated, such a minimum energy condition usually can only be found with the aid of computer simulationma subject of micromagnetics. Here we give only a rough analysis of the possible roles each energy term may play in the formation of a domain configuration. Figure 70 shows the possible domain configurations in a uniaxial anisotropy soft magnetic film. Figure 70a is a single-domain state, in which the atomic moment of the film is aligned perfectly parallel to meet the minimum exchange energy requirement, and the domain magnetization is aligned along the easy axis to meet the minimum anisotropy energy requirement (assuming zero external field and zero magnetostriction). However, when a ferromagnet is in a singledomain state, the demagnetization energy could be very high because of the high magnetic charges in such a domain configuration, which makes the total energy of the system high. In Fig. 70b, the single-domain state of Fig. 70a is broken up into two subdomains, with their magnetizations aligned antiparallel but still along the easy axis. The walls between the two domains are 180 ~ B loch walls. In such a case, the anisotropy energy term is unchanged. The demagnetization energy is considerably reduced because fewer charges are formed, because of partial flux closure in the two antiparallel domains. The exchange energy is increased slightly, and an extra energy term, the domain wall energy, is introduced into the system. Figure 70c shows another typical type of domain configuration, the closure domain structure. Domains with 90 ~ domain walls are formed along the strip edge. In this case, the demagnetization energy term almost vanished because of the complete flux closure within the magnetic domains. There is a slight increase in exchange energy and domain wall energy because more domains are formed. The anisotropy energy term is no longer zero because the edge domains are aligned at a 90 ~ angle with respect to the easy axis. If the uniaxial anisotropy of the film is not sufficiently strong, the edge domains will expand further to form the so-called diamond structure. In the worst case, the edge domains will occupy the whole strip to form the hard-axis aligned domains, which are further discussed in Section 5.3.2. The effect of the signs of magnetostriction on domain configurations can be understood from Eq. (5.30). In the case of

o"x > O'y, a positive As will reduce the effect of K u, where as a negative As will enhance the effect of K u. A typical example of this was demonstrated by Narishige et al. [217], as shown in Fig. 71, for domain patterns in the pole and yoke of thin film heads with positive and negative magnetostrictive Permalloy films. Because the y dimension (easy axis) is considerably smaller than the x dimension (hard axis) in a thin film head, the tensile stresses follow trx > O'y. Therefore, the Permalloy yoke with negative As exhibits more stable easy axis aligned domains than the one with positive As.

5.3.2. Domain Configurations in Multilayered Films Multilayered films or laminated films are frequently used in film heads to improve high-frequency performance. Lamination of soft magnetic layers with an insulation spacer brings two major advantages [176]. One is that lamination can suppress eddy currents. If the thickness of each magnetic sublayer is made smaller than the eddy current skin depth, the highfrequency magnetic flux is carried by the whole thickness of the yoke, and the flux-carrying efficiency of the yoke at high frequency is greatly improved. The other is that proper lamination can eliminate domain walls and achieve a single-domain yoke, in which the contribution to the flux reversal is mainly from magnetization rotation rather than a domain wall motion mechanism. Hence, the high-frequency permeability of the

L

9

/], = + 8x 10 -7

~. = - 5x 10 -7

Fig. 71. Domain configurations of film heads with positive and negative Permalloy films, Reproduced with permission from [217], copyright 1984, IEEE.

546

PAN

head is maintained. The possible domain structures and micromagnetic analysis of laminated film heads were well studied by Slonczewski et al. [ 176]. Here a brief summary of the work is given. According to Slonczewski et al. [176], a two-layer laminated Permalloy strip may have one of the three domain configurations shown in Fig. 72. Figure 72a is the periodic stripe domain pattem with hard-axis closure domains along the edge, which is similar to the domain configuration in a singlelayered film, as discussed in the previous section. The formation of the edge domains is due to the strong demagnetizing field (or shape anisotropy field) along the strip edge, which overcomes the uniaxial anisotropy field and forces the magnetization aligned parallel to the edge (parallel to the shape anisotropy field). Such a shape anisotropy field is strip width and film thickness dependent and highly nonuniform in nature. When the strip width decreases, the shape anisotropy field increases and the edge closure domains expand to form the diamond pattern, as shown on the fight of Fig. 72a, in which vertical 180~ walls no longer exist. This is a typical problem in single-layered narrow track film heads, where diamond patterns are formed in the narrow pole tip, and the hard-axis edge domains occupy a large proportion of the pole tip but play no part in the flux reversal process. The energy per unit strip length of each magnetic sublayer for a closure domain pattern is [176] Ecp = D. v/2ru TBlochW

')/Bloch is the Bloch wall energy density given by Eq. (5.27) or, alternatively, by a simple numerical expression [176], 21.43A ')/Bloc h =

D

+ 0.581KuD +""

(5.32)

where A is the exchange constant of the magnetic material (A = 10 -6 erg/cm for Permalloy). The energy per unit length of each magnetic sublayer for a diamond pattern is [ 176] / EDp

: D - /\

\ KuW + 7mo~h/ / 2

(5.33)

Figure 72b shows the easy axis state domain pattern of a two-layer laminate, where each layer exhibits a single domain state with domain magnetizations aligned antiparallel with each other along the easy axis direction and with two edge walls (known as edge curling walls or ECW) to form a closed magnetic circuit. In such a domain configuration, the energy per unit strip length of one magnetic sublayer is [176] EEA -- 7rKuDA

(5.34)

where A is termed the wall shape energy and is given by [176]

(5.31) A = Ms

Ku

(5.35)

where D is the thickness of the magnetic sublayer, K u is the uniaxial anisotropy constant, W is the width of the strip, and where b is the thickness of the nonmagnetic spacer. Figure 72c shows the hard-axis state domain pattern of a two-layer laminate, in which each layer also exhibits a singledomain state, but with domain magnetizations aligned antiparallel with each other along the hard axis. This situation occurs when the strip width is very small and the energy density required to maintain the easy-axis state becomes too high in comparison with the energy density required for a hard-axis state domain, which is given by [176]

T

Wy

(a)

L

p-x--~

/ / / / (b)

l'

•t

/ / / / '

EHA -- KuDW

(5.36)

"Y

_.

)

Edge walls

_

Fig. 72. Domain pattern classific.ation of laminated strips. (a) Closuredomain pattern. (b) Easy-axis state. (c) Hard-axis state. Reproduced with permission from [176], copyright 1988, IEEE.

"Phase diagrams" showing the conditions under which the three domain configuration are stable can be constructed by using the above energy density equations for each domain pattern and using the basic material and geometry parameters such as H k, D, b, and W. Some typical results are shown in Fig. 73 for two-layer laminated narrow Permalloy strips. These phase diagrams give a simple and clear indication of the effect of the strip width W, thicknesses (D and b) of magnetic and nonmagnetic layers, and the anisotropy field (Hk) on the possible domain configuration in film heads, which can be used as a rough guide in the design of thin film heads. This simple model can be extended to multilayer systems [176].

HIGH-DENSITY MAGNETIC RECORDING 3O

'

!

!

!

i

Hx = 3 0 e W = I 0 /zm

20

closuredomain ;tale

\

E C v

.13

easyaxis stale

I

I

0

I

0.2

0.4

os

oe

io

o (m-) (a) 30

!

i

I

i

2O

E C: v

.Q

10

'

~

9

!

!

!

!

Hx = 3 0 e W = 6/zm

hardaxis slate

closu r e domaln

state

~

1

o

I

l. . . .

0.2

I

I

0.6

0.8

I

0.4

9

,

O (/zm)

(b) 3O HK= lOOe W = 6/zm. \ \

2O

closuredomaln slate

E

.

C v

e a s y - ~ axis state

.Q I0

0

l

0

l

0.2

1

.

I

1

0.4

frequencies (for high efficiency), zero or near-zero magnetostriction (for reduced domain noise/head instability), high resistivity (for reduced eddy current damping), and good thermal stability and corrosion resistance (to survive the head fabrication process and to allow for long life). Table VI lists the properties of some typical soft magnetic materials used in film heads. Soft magnetic thin films for recording head applications can be roughly classified into the following four categories: Ni-Fe alloys, FeSi-based alloys, cobalt-transition metal (CoTM) amorphous alloys, and FeN-based nanocrystalline films. However, there are still a large number of soft magnetic materials that are not included in these four categories.

5.4.1. Ni-Fe Alloys

easy-

0

547

I

0.6

1

I

0.8

.

, l

Ni8]Fei9 (wt %), known as Permalloy, was the most popular head material for the early generation of film heads because of its excellent soft magnetic properties. The uniaxial anisotropy of Permalloy films can be obtained by applying an external magnetic field during film deposition or by postdeposition annealing [ 187]. One of the major advantages of using Permallay in film heads is that it can be deposited by electroplating (wet process), which is a fast and cheap deposition process particularly suitable for mass production. Detailed descriptions of the electrodeposition process of Permalloy for film heads can be found in [187, 189]. In comparison with the vacuum deposition process (dry process), electrodeposition is an additive process for feature definition, which allows for much easier lithographic definition and control of small features [ 184]. This can considerably reduce the cost of fabrication of film heads. The major disadvantages of Permalloy are its low Bs and low resistivity, which is not suitable for heads used in high-density disk drives with a high data rate and high H e media. The electroplating process also has its limitations. For example, it is difficult to make laminated films with insulation spacers, which is essential for domain control of narrow track heads and for eddy current suppression in high data rate recording. An alternative material in the Ni-Fe alloy family is the NiasFe55alloy [184, 185], which exhibits higher Bs (17 kG) and higher resistivity than Permalloy, and which can also be deposited by electroplating. Its major disadvantage is its relatively high saturation magnetostriction ('~ 10 -5).

I,o

D (,u,m)

(c) Fig. 73. Theoretical phase boundaries for narrow stips, assuming three combinations of W and H k. Reproduced with permission from [176], copyright 1988, IEEE.

5.4. Soft Magnetic Films for Writers

In addition to the high Bs requirement, film head materials usually have a well-defined uniaxial anisotropy (for domain structure control) with small coercivity (for low hysteresis loss), reasonably high permeability, particularly at high

5.4.2. FeSi.Based Alloys Sendust (Fe-Si-A1 alloy) is one of the popular materials in the FeSi-X family, which is mainly used in metal-in-gap (MIG) heads because of its superior thermal stability. To obtain desirable soft magnetic properties, the composition of Fe-Si-A1 alloy must be within the so-called Sendust composition, which is typically Fe73.6A19.95i16.5 . Such a film exhibits higher Bs (11 kG), higher resistivity (84/xIl-cm), and higher permeability than Permalloy films. The addition of N Sendust can improve its soft magnetic properties. For example, films with a composition of Fe76.6All.3Si16.33N5. 8 prepared in m r ] N 2 ambient have been

548

PAN Table VI.

FeSi-X alloys

Ni-Fe alloys Permalloy Properties

(Ni81Fel9) [ 186]

Properties of Some Typical Soft Magnetic Films in Recording Heads

Ni45 Fe55 [ 184, 185]

Co-TM amorphous alloys [ 177]

FeN-X nanocrystalline

Sendust (Fe-Si-A1)

Co96 Zr4

C091 N b 6 Z r 3

C095.2 Ta3.6Zrl. 2

FeTaN

[1881

[197]

[194]

[193]

[178, 182]

FeA1N

[179--181,183]

Bs(KG )

10

17

11

~ 13

~ 15

~ 20

~ 20

H k (Oe) H e (Oe) Is As

2.5 0.1 ~, 0.5 ~' 1000 - 1 x 10 -6

9.5 0.4 ~ 1700 2 x 10 .5

N/A --- 0.2 ~ 2000 ~ 10 .6

24 0.5 1500 2 x 10 -6

3 ~ 16 0.01 1 5 0 0 ~ 5000 ~ 10 .8

3 ~ 16 0.01 1 5 0 0 ~ 6000 ~ 10 .8

~ 10 1.5 ~ 2500 ~ 10 .6

~ 5 0.6 ~ 2500 ~ 10 -6

24

48

85

>100

125

>100

30 ~ 120

20 ~ 130

~ 425

N/A

~ 700

450

~ 450

~ 450

300

Excellent

Excellent

Excellent

270~

~ 270~

~ 270~

200~

Good

Poor

N/A

Fair

Good

Good

Good

N/A

N/A

p (/212 - cm) Recrystalization temperature (~ Thermal stability Corrosion resistance

16

shown to have a B s of 13.6 kG, an H e of 0.3 Oe, a ].Lr of 2500, and a As of 5 x 10 -7 [199]. FeRuGaSi alloy (SOFMAX) is another type of soft magnetic material that belongs to the FeSi-X family [190]. It exhibits higher B s (up to 13.6 kG) with improved soft magnetic properties, thermal stability, and wear and corrosion resistance compared with Sendust alloys.

5.4.3. Co-TM Amorphous Soft Magnetic Films Co-TM (TM = Ti, Zr, Nb, Mo, Hf, W, or Ta) binary or ternary amorphous films are found to exhibit higher saturation magnetization and better soft magnetic properties than Permalloy and Sendust films. An excellent review of these films can be found in [177]. These films can be deposited by sputtering. One of the advantages of these films is that zero or nearsaturation magnetostriction can be obtained by using different glass-forming elements or by adjusting the film composition. As shown in Fig. 74, the saturation magnetostriction of Corich binary amorphous alloys can be either positive or negative, depending on the type of glass-forming element used. As also varies with film composition. This provides us with a large number of possibilities for new ternary alloys with zero or near-zero magnetostriction. Examples of these films include Co87NbsZr8 [192], Co85Nb7.5Ti8.5 [191], Co92TanZr 4 [193], and Co91Nb6Zr 3 [194], It is also possible to add a ferromagnetic element such as Fe Co-TM binary systems to form zero or near-zero magnetostriction amorphous films. A typical example of this is Co89.8Nb8Fe2.2 [195,196], which exhibits zero magnetostriction and a high B s of 14.3 kG. The Co-Nb binary system exhibits a negative As. The addition of Fe atoms changes As toward positive passing through zero at an Fe concentration of 2.2 at % [1771. The dependence of the saturation magnetic flux density B s of Co-TM amorphous alloys on the concentration (y) of the glass-forming element is shown in Fig. 75 [177]. Almost all

500

of the values of B s follow a common straight line increasing with the decrease in y, except for the Co-Ta alloys. The arrows in the figure indicate the critical concentration (Yc) of the glass-forming element, below which the alloy becomes crystalline. The highest B s for each Co-TM amorphous alloy is determined by the value of Yc. Among them, the Co96Zr 4 amorphous alloy exhibits the highest B s (16 kG), but with nonzero magnetostriction. The values of B s for nonmagnetostrictive amorphous alloys are always smaller than this because additional glass-forming elements other than Zr are required to achieve the zero magnetostriction. The nonmagnetostrictive condition in these amorphous alloys is achieved at the sacrifice of B s. For example, an optimized composition of Co-Ta-Zr amorphous alloy to make B s as high as possible within the condition of zero magnetostriction is found in Co95.2Ta3.6Zrl. 2 [193]. The B s of such an alloy is about 15 kG [193].

......

t

.......

I

4

"

'

I

.........

M=Zr

t .

.

.

.

.

I

I

"

" '"

,,

.-.

Ti

'o

'

t

Nb

v r ,,<

"

- z, -

Ta

"/'XI, / 1 B

-6-

~

o

]

/ "8

--

0

t

5

...... I

10

l

15

I ........... l

20

25

t

.

.

.

.

.

.

.

l

30

V [m%] Fig. 74. Saturation magnetostriction (As) of Co-TM amorphous alloys. TM is Zr, Hf, Ti, Nb, Ta, W, or B. Reproduced with permission from [177], copyright 1984, American Institute of Physics.

HIGH-DENSITY MAGNETIC RECORDING I 6

v'

l .....

v

.......

I

I

about 1500. The anisotropy field is reduced to about 3 Oe, and permeability is increased to 5000 by rotational field annealing (RFA) [194]. The major drawback of amorphous soft magnetic films is the thermal stability of the uniaxial anisotropy. The easy axis of most Co-TM amorphous alloys becomes unstable above 200~ which makes it difficult for them to survive the head fabrication process without degradation of easy axis orientation. However, the easy axis thermal stability of these amorphous films can be improved by an initial magnetic annealing after deposition. Ishi et al. [198] have found that CoZrTa amorphous films exhibit stable easy-axis orientation up to 270~ if the film was initially annealed in a magnetic field at 350~

.....

= Zr

15 6 H!

N ~ r - ' - ra

549

9 Nb

~0

\

\ \

,,

,

.

't, \ %

0 ~ 0

1 5

:

1

1

10

15

,

1

20

\

1

i

25

30

(Co)

Fig. 75. Saturation magnetic flux density (Bs) of Co-TM amorphous alloys. TM is Zr, Hf, Nb, Ta, or Ti. Reproduced with permission from [177], copyright 1984, American Institute of Physics.

The amorphous state of the Co-TM alloys is metastable. The recrystallization temperature of these amorphous alloys is between 400 and 500~ varying with the type and concentration of the constituent glass-forming element. The soft magnetic properties of these films disappear completely after recrystallization. Hayashi et al. [190] have formulated an expression based on the available data for the estimation of recrystallization temperatures (Tx) for Co-TM (TM = Zr, Nb, or Ta) amorphous alloys Tx(~

= (3.43 +0.35)Xco + (16.20+2.13)XTa + (13.44 + 2.07)X~

(5.37)

+(19.79 i 2.59)Xzr where Xco, Xra, XNb, and Xzr are the at % of the Co, Ta, Nb, and Zr elements in the alloy, respectively. Equation (5.37) is valid for 75 < Xco < 95, 5 < Xra and XNb and 2 < Xzr. Obviously, in terms of the contribution to the higher recrystallization temperatures, the three TM elements are in the order Zr, Ta, and Nb. The highest recrystallization temperatures are obtainable in the Co-Zr amorphous alloys with higher Zr concentrations. The uniaxial anisotropy of Co-TM amorphous films can be obtained by using a magnetic substrate holder with a field strength of 50-100 Oe during film deposition, or by subsequent magnetic annealing after deposition. The temperature for magnetic annealing is in the range of 200-350~ Typical values of the anisotropy field H k of as-deposited Co-TM amorphous films is between 13 Oe and 20 Oe. The hardaxis initial relative permeability of Co-TM amorphous films is

5.4.4. FeN-Based Nanocrystalline Films The FeN-X (X = A1, Ta, or Ti) nanocrystalline films are the only category of soft magnetic films developed so far with the highest saturation magnetization (47rMs up to 20 kG, only slightly lower than that of pure Fe), good soft magnetic properties (well-defined uniaxial anisotropy, near-zero magnetostriction, and small He), and good thermal stability for recording head applications. Thin films heads with FeA1N [183] poles have been fabricated and used in high-density recording. The soft magnetism of FeN-X films originates from their nanocrystalline structure and can be explained by Hoffmann's ripple theory [203,204]. These films are usually prepared by sputter deposition techniques. The key factors determining the soft magnetism of these films are grain size, their morphology, and the nature of grain boundary materials [201]. The addition of a small amount (2-3 at %) of a third element (A1, Ta, Ti, or Si) to the Fe-N binary system shows a marked effect in preventing the grain growth and in improving the soft magnetism and thermal stability of the films [179]. It was found that when the grain size of these films is reduced to below 15 nm, the magneto-crystalline anisotropy vanishes and the coercivity is considerably reduced to below 1 0 e . However, it is necessary that the grain boundary be ferromagnetic and very thin [ 188], so that strong intergranular exchange coupling is maintained and the nanosized grains are not magnetically isolated from each other. Grain isolation will result in high coercivity because of the single-domain particle behavior of the isolated grains. Columnar structure is not desirable because it plays the key role in promoting the perpendicular anisotropy component (K• Under noncolumnar growth conditions, the incorporation of nitrogen into the film first decreases K• to a critical value, then the N acts as a "grain refiner," and excellent soft magnetic properties can be obtained [201]. Nitrogen incorporation in FeTa films is much higher than in Fe films [201]. However, it was found that the saturation magnetization of the films did not change with nitrogen addition up to 2% lattice dilation [200]. A reversible degradation of soft magnetic properties above a critical temperature (T0,c) was reported in FeA1N films, and To,c decreases with the increase of N contents in the films [202]. The weakened intergranular exchange coupling due to the change of ferromagnetic phase

550

PAN

to the paramagnetic phase of the grain boundary materials at temperatures above To,c was believed to be responsible for the degradation of soft magnetic properties [202]. The effect of magnetic annealing on the behavior of soft magnetic properties of FeTaN films was studied by Viala et al. [206]. It was found that the soft magnetic properties of FeTaN films annealed in a field applied parallel to the original easy axis were stable up to annealing temperatures of 300~ Higher annealing temperatures lead to grain growth and TaN precipitation, along with substantial reduction of saturation magnetization and poorer soft magnetism. Changes in stress and magnetostriction were also observed. When the FeTaN films were annealed in a transverse field at 150~ as a function of annealing time, change and switching of H k were observed. H k decreases with increasing of annealing time, switching at 60 min and finally stabilizing at 120 min, aligning with the annealing field. Minor and Barbard [205] have studied the thermal stability of FeTaN films and found that the thermal stability of FeTaN films was improved because of the Ta. The Ta content for optimum thermal stability was about 10 wt%.

5.4.5. Giant Moment Soft Magnetic Films The giant magnetic moment phenomenon in Fe-N films was first reported by Kim and Takahashi in 1972 [208]. They discovered that the saturation magnetization (Ms) of polycrystalline Fe-N films deposited by evaporation in low-pressure nitrogen was as high as 2050 emu/cc (47rMs = 25.8 kG). Fig. 76 shows the dependence of saturation magnetization of the as-deposited Fe-N films on the nitrogen pressure during deposition. The giant moment of the Fe-N films was attributed to the a"-martensite phase (a"-Fel6N2) in the Fe-N films, which has a bct structure, and the atomic magnetic moment associated with Fe atoms of c~" -Fel6N 2 was deduced to be

3.0/z B, compared with 2.4/-s per atom for pure Fe. The c~"martensite phase was discovered in 1951 by Jack [207]. However, its magnetic properties were not investigated until 1972. There has been increased interest in this material over the past decade because of its potential applications in film heads. Giant magnetic moment c~"-Fel6N2 (001) single-crystal phase (47rMs -- 29 kG) was obtained by Komuro et al. by molecular beam epitaxy (MBE) on InGaAs (001) substrates [209]. The atomic moment measured by VSM and Rutherford backscattering was 3.5/-/~a [215]. The resistivity of the pure c~" -Fel6N 2 phase is 30/zf~-cm [215]. Giant magnetic moment (up to 25 kG) multiphased Fe-N films have also been obtained by ion beam deposition on S i ( l l l ) substrates [210], and dc and rf magnetron sputtering on Si(001) and glass substrates [211, 212]. Most recently, a giant moment ct"-(Fe,Co)]6N 2 phase was also reported by Wang and Jiang, who used facing target sputtering on Si(001) substrates [213]. Although results published to date have not been so consistent and are incompatible with theoretical predictions [214], and the reproducibility of the film was also in question, they appeared to indicate that the formation of the giant magnetic moment a" phase depends upon the substrate material, nitrogen partial pressure, film thickness, deposition process, and in situ and post deposition heat treatment. Questions regarding the mechanism of giant magnetic moment have yet to be answered. With regard to applications, reproducibility, soft magneti3m, and thermal stability of the a" phase have to be further investigated.

Acknowledgment The author is grateful to John Mallinson for the review of the GMR spin valve chapter.

REFERENCES Ms (emu/cm3) 2500

-

"

1

9

2t00

"

_ ~ - - ~ 3 5 % Co-t:e - ----

1900

x ,,,=,,,

.

'

L

_

f X ~ )

/o

~

"

1

"

"' " "

5oo

--

_

t"--

~ . . .

v_-'" t.~ "''-x

.N. 'N

k.,x

.

1700

1500

15OO

I100 ~r .

2

!

'~'

Io"

PRESSURE

..,

.

I

,

i

jo.3

5

icr=

2

~,

DURING

DEPOSITION

(Tort)

Fig. 76. Saturation magnetization of Fe-N films vs. pressure during deposition. O, deposited in nitrogen atmosphere; x, deposited in air (usually vacuum). Reproduced with permission from [208], copyright 1972, American Institute of Physics.

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