Note : Les descriptions sont présentées dans la langue officielle dans laquelle elles ont été soumises.
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WITH TWO METALLIC MA(~NFT1(' PHAfiFfi
The present invention relates to materials for use in magnetic recording
sensors and the like, and more particularly relates to a magnetoresistive
material with
two metallic magnetic phases.
Brief Description of the Prior Art
Materials which exhibit a change in resistance when exposed to a
magnetic field are of use in preparing magnetic recording sensors, such as
those used,
for example, in computer disk drives. At the present time, state-of the-art
computer
disk drives employ sensor materials which exhibit the anisotropic
magnetoresistance
effect (AMR). Materials which exhibit the AMR have a magnetoresistance which
depends on how the magnetic field is applied with respect to the direction of
current
flow.
In other types of materials, which exhibit the giant magnetoresistance
effect (GMR), a non-magnetic metallic conductor, such as copper, is necessary
to
create a disordered state of electron spins in a ferromagnet. However, the
nonmagnetic metallic conductor does not contribute to desired scattering of
the
conduction electrons, and in fact, may act as a low resistance shunt path
which
decreases the magnetoresistance. These prior GMR materials include a magnetic
metal and a non-magnetic metal.
Macroscopic ferrimagnets are a new class of phase separated magnetic
materials which have been recently discovered, and are described, for example,
in R.J.
Gambino et al., 75 J. A~1. Phys_. 1871 (1994). The macroscopic ferrimagnets
include
two magnetic phases with a negative magnetic exchange at the phase boundary. A
prototypical example is the Co-EuS system which has 100 A particles of EuS in
a
cobalt matrix. The EuS is exchange coupled antiferromagnetically to the cobalt
at
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least at the Co/EuS interface. In the Co-EuS system, the small size of the EuS
particles results in a large fraction of the EuS being in close proximity to
the interface
which is influenced by the strong Co/EuS exchange. It has been found that
these
materials display unusual magneto-optical properties, as described in R.J.
Gambino
and P. Fumagalli, 30 IEEE Trans. Magn. 4461 (1994), and magneto-transport
properties, as described in R.J. Gambino and J. Wang, 33 Scr. Melall. Mater.
1877
(1995) and R.J. Gambino et al., 31 IEEE Trans. MaQn. 3915 (1995).
Magnetization
and Kerr hysteresis loops have confirmed the macroscopic ferrimagnetic model
for
these systems. In measurements of the optical and magneto-optical properties
of Co-
EuS thin films, polar Kerr rotations of up to 2 ° have been observed in
Co-rich films at
photon energies of 4.5 eV, as described in P. Fumagalli et al, 31 IEEE Trans.
Magn.
3319 (1995). Transport measurements show that the magnetoresistance of Co-EuS
behaves like that of the widely studied granular giant magnetoresistance
effect (GMR)
materials, as described in S. Zhang, 61 0,nnl. P y, .~ 1855 (1992), which
include
particles of a ferromagnetic metal in a conductive, nonmagnetic matrix. In
contrast,
Co-EuS includes semiconducting, ferromagnetic particles in a conductive,
ferromagnetic matrix of cobalt. As a consequence, the temperature dependence
of the
magnetoresistance is very different in the Co-EuS system as compared to the
ordinary
granular GMR materials. With respect to the magnitude of the effect, the
magnetoresistivity change (8p) of the Co-EuS system is 8 x 10'5 S2-cm at room
temperature in a field of 1 T, which is larger than other magnetoresistive
materials.
Even though the magnetoresistivity change of this system is large, the
magnetoresistance defined as 8p/p is small, typically 2 ~ 3 %, because of the
high
resistivity of the material caused by a large volume fraction of the
semiconducting
EuS phase.
While materials exhibiting the AMR effect have enhanced the
performance of computer disk drives, and while the aforementioned Co-EuS
systems
are promising, it would be desirable to develop materials having a larger
change in
resistance as a function of applied magnetic field strength, that is, a larger
30 magnetoresistance effect. Such materials could permit the development of
more
sensitive magnetic recording sensors. It would be desirable to develop such
materials
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which would exhibit the GMR as opposed to the AMR. In materials exhibiting the
GMR, the resistivity decreases with the applied magnetic field independent of
the
direction of the applied field with respect to the direction of current flow.
In addition
to this desirable isotropy, the GMR is usually stronger than the AMR.
SUMMARY QF THE INVENTION
It is an object of the present invention to provide an improved material
suitable for manufacturing more sensitive magnetic recording sensors.
It is another object of the present invention to provide such a material
which exhibits the GMR.
10 It is a further object of the present invention to provide such a material
which includes two ferromagnetic phases which are exchange coupled
antiferromagnetically.
It is yet another object of the present invention to provide a method of
manufacturing such a material.
15 It is a further object of the present invention to provide a method of
sensing magnetic fields using such a material.
It is still another object of the present invention to provide a digital
magnetic recording system which utilizes such a material in a read head.
In accordance with one form of the material of the present invention, a
20 magnetoresistive material exhibits the GMR and has two phases. The first
phase
includes a matrix of an electrically conductive ferromagnetic transition metal
or an
alloy thereof. The second phase is a precipitate phase of an electrically
conductive
rare earth pnictide which exhibits ferromagnetic behavior when precipitated
out of the
matrix. The second phase is antiferromagnetically exchange coupled to the
first
25 phase. In a preferred form of the first embodiment, the matrix comprises
cobalt and
the precipitate phase comprises terbium nitride.
Thus, the present invention provides a new macroscopic ferrimagnet,
in the system Co-TbN, which exhibits the GMR. The Co-TbN system demonstrates
typical macroscopic ferrimagnet properties: a magnetic compensation point and
30 negative GMR. The Co-TbN system with 32 % TbN by volume composition shows
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0.72 % GMR under an applied field of 8 kOe at room temperature and 9 % GMR at
250° K under an applied field of 40 kOe. In the Co-TbN system, the
temperature
dependence of the GMR is quite different from that of ordinary GMR materials,
where the negative magnetoresistance decreases with increasing temperature.
The
GMR in the Co-TbN system increases with increasing temperature, which is due
to
the increase of ferromagnetic alignment of the Co and TbN with an applied
field
caused by the decrease of exchange coupling by temperature.
In an alternative form of material according to the present invention,
the second precipitate phase comprises an electrically conductive Heusler
alloy such
as CoZMnSn or CozTiSn.
The present invention also provides a method of manufacturing a
magnetoresistive material of the types described above. The method includes
the
steps of providing a target (for example, a sheet metal or thin film target)
of an
electrically conductive ferromagnetic transition metal or an alloy thereof;
locating a
1 S plurality of pellets of an electrically conductive rare earth element (or
constituents of a
Heusler alloy) on a surface of the target; sputtering the target and the
pellets with ions
in a suitable plasma, such as an argon plasma, to cause the film and the
pellets to form
an amorphous alloy of the electrically conductive ferromagnetic transition
metal or
alloy thereof, and the electrically conductive rare earth element (or
constituents of a
Heusler alloy); and subsequently annealing the amorphous alloy to cause
formation of
the precipitate phase within the matrix of the ferromagnetic transition metal
or alloy
thereof. Techniques other than sputtering can also be employed.
The present invention fiuther provides a method of detecting magnetic
field strength of a magnetic field associated with a magnetization pattern
recorded in a
medium. The method includes the steps of providing a sensing head which
includes a
portion of a magnetoresistive material of the type described above; exposing
the
sensing head to the magnetic field of the magnetization pattern in the
magnetic
recording medium; sensing the electrical resistivity of the portion of
magnetoresistive
material exposed to the magnetic field of the magnetization pattern in the
medium;
and converting the electrical resistivity of the portion into a signal which
is indicative
of the magnetic field strength of the magnetization pattern in the medium. It
will be
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appreciated that any of the materials of the present invention described above
can be
used for the portion of magnetoresistive material in the read head.
The present invention yet further provides a magnetic recording system
which is adapted for use with a magnetic recording medium having a
characteristic
coercive force and which has a plurality of stored data in it. The data is
stored in the
form of a magnetization pattern in the medium. The magnetic recording system
includes a write head (which is optional) and a read head. The read head
includes a
portion of the magnetoresistive material of the present invention which is
located in
proximity to the magnetic recording medium and also includes a resistivity
sensor
which detects the resistivity of the portion of material according to the
present
invention corresponding to magnetic field strength levels of the magnetization
pattern
in the medium. The system also includes a controller which controls the read
head
(and optional write head) and which converts the detected resistivity of the
portion of
material according to the present invention into a signal which is indicative
of the
15 stored data in the recording medium. Again, the portion of material can be
any of the
magnetoresistive materials according to the present invention.
These and other features and advantages of the present invention will
be pointed out in the following specification, taken in connection with the
accompanying drawings, and the scope of the invention will be set forth in the
appended claims.
BRIEF DESCRIPTION OF THE DRAWINGS
Figure 1 is a plot of percent change in resistivity (GMR) against
applied magnetic field at room temperature for a Coo.ba-(TbN)o.3z material
according to
the present invention;
25 Figure 2 is a graph of resistance versus temperature at two different
values of applied magnetic field strength for the material of Figure 1;
Figure 3 is a plot of percent change in resistivity (GMR) vs. the
volume percent of terbium nitride present in a sample;
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Figure 4 is a plot of percent change in resistivity (GMR) vs. applied
magnetic field at room temperature for a system employing a Co/Co~TiSn thin
film
annealed in a vacuum;
Figure 5 is a figure similar to Figure 4 except showing results for the
material when annealed in nitrogen;
Figure 6 is a schematic diagram of a facing targets magnetron
sputtering system suitable for producing materials in accordance with the
present
invention;
Figure 7A is a diagram showing a uniform magnetic field in the system
of Figure 6;
Figure 7B is a diagram showing a divergent magnetic field in the
system of Figure 6;
Figure 8 is a diagram of deposition rate against argon pressure at
different substrate distances;
Figure 9 is a plot of composition vs. argon pressure at different
substrate distances;
Figure 10 is a plot of sputter yield ratio of terbium to cobalt against
argon pressure;
Figure 11 shows an x-ray diffraction pattern of a TbCo film annealed
at 650°C in a nitrogen gas atmosphere;
Figure 12 shows a digital magnetic recording system which employs
materials of the present invention;
Figure 13 shows the magnetization of Coo.bB-(TbN)a,32 as a function of
temperature at various applied magnetic fields;
Figure 14 shows the magnetization of a Co-TbN macroscopic
ferrimagnet with applied magnetic field at various temperatures;
Figure 15A shows the change in resistivity of a COp.bg-(TbN)0.32 f lm as
a function of magnetic field at room temperature;
Figure 1 SB shows the changes in resistivity for the same film, also as a
function of applied magnetic field, at room temperature; and
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Figure 16 is a plot of the percent change in resistivity (GMR) of a
Coo.bg-(TbN)o.sz film as a function of temperature in a relatively high
magnetic field of
40 kOe.
5 One form of material in accordance with the present invention is a
magnetoresistive material exhibiting the giant magnetoresistance effect (GMR)
and
which has two phases. The first phase comprises a matrix of an electrically
conductive ferromagnetic transition metal or an alloy thereof. The second
phase is a
precipitate phase of an electrically conductive rare earth pnictide which
exhibits
10 ferromagnetic behavior when precipitated out of the matrix. The second
phase is
antiferromagnetically exchange coupled to the first phase.
The electrically conductive ferromagnetic transition metal or alloy
thereof can include at least one of iron, cobalt and nickel or an alloy of one
or more of
those elements, such as FeCo, FeNi, NiCo, CoNiFe, NiCu, FeCr, CoAI, and the
like.
15 The alloys can be formed from the ferromagnetic metal and non-magnetic
materials
(e.g., the NiCu, FeCr, and CoAI). Ferromagnetic manganese and chromium
compounds can also be employed. The rare earth pnictide can include a rare
earth
element (or a compound or alloy) thereof, selected to anti-parallel couple to
the
transition metal or alloy thereof, and (in turn) compounded with one of
nitrogen,
20 phosphorous, arsenic, antimony and bismuth. Terbium is presently believed
to be the
preferred rare earth element.
The precipitate phase may exhibit independent ferromagnetic behavior,
that is, it may behave ferromagnetically by itself. Alternatively, the
precipitate phase
may not show any meaningful independent ferromagnetic behavior, but may become
25 ferromagnetic, that is, may exhibit meaningful ferromagnetic behavior, when
it is
precipitated into the matrix. In this latter case, the exchange coupling
between the
phases will boost the ferromagnetism in the small particles of the precipitate
phase.
In a preferred material according to the present invention, the matrix
comprises cobalt, and the precipitate phase comprises terbium nitride. The
terbium
30 nitride should be present in a volume percent which is sufficient to
produce the GMR
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in the material without causing undesirable discontinuities in the cobalt
matrix. That
is to say, the cobalt should be substantially continuous throughout the
device: there
should not be cobalt regions surrounded by TbN. The cobalt matrix can have a
volume percent of from about, for example, 30% to about 75%. The terbium
nitride
5 precipitate phase can have a volume percent of from about 25% to about 70%.
As
presently understood, it is believed that at about 70% terbium nitride, the
cobalt
matrix would begin to have the aforementioned undesirable discontinuities.
Preferably, the cobalt matrix has a volume percent of from about 61%
to about 70%, and the terbium nitride precipitate phase has a volume percent
of from
10 about 30% to about 39%. As discussed below, tests were conducted from about
31
volume percent to about 39 volume percent of terbium nitride and indicated
that the
GMR continued to increase with increasing volume percentage of terbium
nitride.
However, the rate of increase appeared to lower between about 35% and about
39%
volume percent of terbium nitride.
1 S The Co-TbN system is a macroscopic fernmagnet and has TbN
precipitates in a Co matrix. The TbN has the same magnetic moment as pure Tb
and
the rock salt structure, as described in R.J. Gambino and J.J. Cuomo, 113 ,L
Electrochem. Soc. 401 (1966), the same as EuS. The TbN precipitates also
provide
the higher Curie temperature and thus stronger antiparallel exchange coupling
with
20 the Co matrix than EuS. These stronger exchange effects are caused by
conduction
electron mediated exchange of the RKKY type, as described in C. Kittlel,
Introduction
to Solid State Physics 628 (72" Ed., John Wiley and Sons, 1996), which is weak
in
semiconducting EuS. Another difference is the single ion anisotropy of the Tb
ion
which is a non-S-state ion. In contrast, EuS contains divalent europium which
is a S-
25 state ion and thus has zero single ion anisotropy. Furthermore, the TbN is
a conductor
rather than semiconductor so the resistivity of Co-TbN is much less than that
of Co-
EuS, which can improve the magnetoresistance, 8p/p. The Co-TbN differs with
the
granular GMR materials in that both phases are magnetic and also differs from
the
Co-EuS in that both phases are conductors.
30 In an alternative type of material according to the present invention,
the first phase is substantially similar to the first phase discussed above,
and the
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second precipitate phase can include an electrically conductive Heusler alloy
which
exhibits ferromagnetic behavior when precipitated out of the matrix. Again,
the
second phase is antiferromagnetically exchange coupled to the first phase. In
both
cases, a preferred material for the matrix is cobalt. The precipitate phase
can be, for
example, Co~MnSn or CoZTiSn. Compositions in the range of 16.7 to 50 mole
percent Co,MnSn in a Co matrix and 16.7 to 50 mole percent CozTiSn in a Co
matrix
have been tested. In both systems, the 50 mole percent compositions are
currently
believed to be preferable. The mole % is determined by dividing the number of
moles
of CoZMnSn or Co2TiSn by the total number of moles and multiplying by 100. For
example, Co3MnSn = Co/Co,MnSn = 50 mole % CoZMnSn and Co7 MnSn =
SCo/Co,MnSn = 16.7 mole % CoZMnSn. Other operable ranges of the components,
exhibiting desirable GMR properties, which can be easily ascertained by one
skilled
in the art, are within the scope of the invention. The aforementioned
discontinuities in
the cobalt matrix should be avoided in the Heusler alloy systems as well.
Reference should now be had to Figure 1 which plots the GMR against
applied magnetic field strength for a film made of a material according to the
present
invention which includes 68 volume percent cobalt and 32 volume percent
terbium
nitride. The horizontal axis is the applied magnetic field strength H, in
oersteds,
while the vertical axis shows the GMR as measured by the change in resistivity
divided by the resistivity in the saturated state, p(Hsa~), expressed as a
percent. More
specifically:
GMR = 8p/p = LP(0) - P(Hm)~/P(Hsa~~ (1)
where p(0) and p(Hsa~ are the resistivities in zero field and in a saturating
magnetic
field, respectively.
Figure 2 shows a plot of the resistance of a sample of the same material
as Figure I, measured in ohms, as a function of temperature measured in
degrees
Kelvin with an applied field of H = 40 kOe and with an applied field of H = 0.
In
determining the GMR, using the equation R=pL/A where R = resistance, p =
resistivity, L = length, and A = cross sectional area, ~R/R = ~p/p for a given
geometry.
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Figure 3 is a plot of the GMR, again expressed as a percent change in
resistivity, against the volume percent of terbium nitride in the cobalt-
terbium nitride
material. As can be seen, the GMR continues to increase with increasing volume
percentage of terbium nitride but the rate of change of the increase is less
at the higher
5 volume percentages of terbium nitride, that is, about 35 - 39 volume percent
of
terbium nitride.
Figure 4 shows a plot of GMR, again as a percent change in resistivity,
versus the applied fields in oersteds for a Co/Co2TiSn thin film annealed in
vacuum.
The plot is in the form of a series of points. Figure 5 is a similar plot but
wherein the
10 thin film has been annealed in nitrogen gas. It can be seen that the GMR
increases
significantly when the annealing is performed in the nitrogen gas. Similar
results are
expected for the materials employing Heusler alloys.
The materials described above are macroscopic ferrimagnets which
comprise two ferromagnetic phases which are antiferromagnetically exchange
15 coupled. Both of the ferromagnetic phases are metallic conductors. The
electrical
resistivity of the materials decreases when a magnetic field is applied, and
the change
in electrical resistance is independent of the direction of the magnetic field
with
respect to the direction of the current flow. As noted, this type of behavior
is
characteristic of the giant magnetoresistance effect (GMR) which has
previously only
been observed in metallic materials comprising a ferromagnetic metal such as
cobalt
and a nonferromagnetic metal such as copper. The giant magnetoresistance
depends
on scattering of conduction electrons by spins in a ferromagnet in a
disordered state.
The prior art materials employed a nonmagnetic metallic conductor such as
copper
which was necessary to create the disordered state but which did not
contribute to the
scattering. As noted, depending on the geometry, the copper conductors might
in fact
act as low resistance shunt paths which decrease the magnetoresistance. In the
materials of the present invention, spin scattering can occur in both
ferromagnetic
phases so that larger GMR effects are possible, in turn enabling more
sensitive
magnetic recording sensors.
The GMR effect for an exemplary composition is about 8% at room
temperature in a field of 4 teslas and it increases slightly with decreasing
temperature
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down to 20 degrees kelvin. Figure 2 shows data from about 40° K to
about 240° K
for H = 0 and H = 40 kOe. It will be appreciated that, in c.g.s. units, B = H
+ 4nMs,
where B = magnetic flux density in kilogauss, H = magnetic field strength in
kilooersteds and MS = magnetization in kilogaus. Outside a ferromagnet, B in
5 kilogauss is normally equal to H in kilooersteds and thus 40 kOe implies 40
kilogauss
= 4T.
Methods of manufacturing the materials of the present invention will
now be described. The magnitude of the GMR and the field sensitivity of the
materials can be controlled by post-deposition heat treatment as set forth
below in the
Example. Reference is now made to Figure 6, which depicts a representative
sputtering apparatus, to be used in accordance with the present invention,
designated
generally as 8. A method of manufacturing a magnetoresistive material
exhibiting the
giant magnetoresistance effect (GMR) and having two phases, in accordance with
the
present invention, includes the step of providing a suitable target (e.g.,
sheet metal, a
disk, sheet, or plate) of an electrically conductive ferromagnetic transition
metal or an
alloy thereof. The target is designated as item 10 in Figure 6. The method
includes
the additional step of locating a plurality of pellets 12 of an electrically
conductive
rare earth element on a surface 14 of the target 10. In a preferred form of
the present
invention, a second target 16 (which can also be, e.g., a disk, sheet, or
plate), also of
an electrically conductive ferromagnetic transition metal or alloy thereof, is
employed. The process illustrated in Figure 6 is known as facing target
magnetron
sputtering (FTMS) with a composite target. The second target can initially be
a pure
conductive ferromagnetic transition metal or alloy thereof without any pellets
and
unalloyed with the rare earth element. FTMS with pellets on both targets
(i.e., both
25 targets are composite targets) is known in the art, as described in Naoe et
al., 23 IEF~
Trans. Mag_n. 3429 ( 1987).
The method also includes sputtering the target 10 (and target 16, when
present) and the pellets 12 so as to cause the targets 10 and 16 and the
pellets 14 to
form an amorphous alloy on substrate 18. The sputtering can be performed using
a
30 suitable plasma of a noble gas such as argon (preferred), helium or krypton
in a
plasma region 20.
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Other physical vapor deposition (PVD) techniques can be employed
besides sputtering, for example, co-evaporation or ion beam deposition.
Sputtering is
believed to be preferable. Further, chemical vapor deposition (CVD) can also
be
performed. Using gases of organometallic precursors of Co and Tb and
decomposing
them to deposit the metals.
The method can also include the step of annealing the amorphous alloy
in an atmosphere containing nitrogen to form a nitride. Other gases such as
phosphine, PH3, or arsine, AsH3, are used to form phosphides or arsenides,
respectively. Gaseous vapors of the elements P, As, Sb or Bi can also be used.
It is
desirable to exclude oxygen and have a slightly reducing atmosphere. The
nitrogen
family element is used to bond with the rare earth element in the amorphous
alloy and
thus form a precipitate phase of a rare earth pnictide within a matrix of the
ferromagnetic transition metal or alloy thereof. One suitable atmosphere (for
nitride
formation) in which the annealing can be carned out is so-called forming gas
which is
a mixture of 90% nitrogen and 10% hydrogen by volume. It should be appreciated
that any suitable gas containing nitrogen or any other element selected from
column
VA of the Chemical Abstracts Service (CAS) version of the periodic table, or a
compound thereof, can be employed. The annealing time required varies with the
temperature and in general will be less at higher temperatures. For example,
for
20 nitrogen, 10-12 hours at 650° C is sufficient. At 450° C, a
detectable reaction occurs
in 12 hours but is not complete.
Still with reference to Figure 6, the illustrated apparatus 8 can also
include first and second magnets 22, 24. The density of the plasma in the
plasma
region 20 is enhanced by the presence of the magnetic field induced by magnets
22,
24. The dotted line 26 represents a common axis for the magnets 22, 24.
Reference should now be had to Figures 7A and 7B which depict
close-up views of the magnets 22, 24 with certain other elements of the
apparatus 8
omitted for clarity. In Figure 7A, a configuration is shown wherein opposite
magnetic
poles of magnets 22,24 face each other. For example, a north pole of magnet 22
is
30 shown facing a south pole of magnet 24. The magnetic field produced by
magnets 22,
24 in the configuration shown in Figure 7A is represented by lines of force
28. In the
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case shown in Figure 7A, the plasma density in the plasma region 20 is
enhanced by
the magnetic field produced by magnets 22, 24. The field is axial from top to
bottom,
generally along axis 26. The field helps to confine the electrons produced
during the
sputtering process to the plasma region 20. These electrons collide with the
atoms of
the gas which forms the plasma, for example argon, to form additional ions,
for
example argon ions. This process is known as electron collision ionization.
As is known in the sputtering art, the ions hit the targets 10, 16 and
cause the emission of atoms as well as secondary electrons. In order to
maximize
yield, it is undesirable that the secondary electrons hit a surface at a large
negative
potential. Accordingly, it is desirable to expose the secondary electrons to
both a
confining magnetic field as shown in Figure 7A and to a suitable electric
field. As is
well known, electrons exposed to electric and magnetic fields move in spiral
paths
between collisions. Referring back to Figure 6, a suitable electric field can
be applied
from a radio frequency (RF) supply 30 connected to target holders 32, 34
through a
suitable matching network 36. RF supply 30 can have any appropriate power
level
and frequency known to those skilled in the plasma sputtering art. For
example, a
power of approximately 2 kW can be employed. In the United States, a frequency
of
13.56 MI~iZ, as allocated by the government for scientific experimentation,
can be
employed.
Referring back now to Figure 7A, the magnetic field produced by
magnets 22, 24 also serves to confine the plasma to a more limited plasma
region 20
which results in more frequent collisions between electrons and neutral gas
atoms, a
more intense plasma, and generation of more ions for bombardment of the
targets.
This in turn leads to a higher deposition rate for the material on substrate
18.
Note that the apparatus 8 can include first and second ground shields
38, 40, as known in the sputtering art for purposes of preventing bombardment
of the
fixtures. Further, a suitable load lock chamber 42 can be employed. Rare earth
elements oxidize easily. Although the oxidation could be removed by
sputtering, the
removed oxide would then tend to deposit on the substrate 18, which would be
undesirable. Load lock 42 can be employed to minimize or eliminate oxygen
contamination. The substrate 18 is placed into the load lock, the load lock is
then
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evacuated and then the load lock is connected to the main chamber which
contains the
plasma region 20.
Under certain conditions, it may be desirable to bombard the substrate
18 where the desired film is formed. In conventional processes, a bias voltage
can be
applied to the substrate. This need not be done with the present invention. A
self bias,
due to the plasma potential, accelerates positive ions towards all surfaces at
ground
potential.
In the present invention, bombardment of substrate 18 can be
controlled by adjusting the distance Lg between the magnet center line axis 26
and the
substrate 18. Different properties are produced in the finished product
depending on
whether or not the substrate 18 is itself bombarded. The composition, magnetic
anisotropy, density and electrical resistivity can be controlled by ion
bombardment of
the substrate during deposition of the material. Refer to the discussion of
Figures 8
and 9 below.
As was stated above, target 10 is employed together with pellets 12,
while second target 16, when employed, is normally formed only of the
electrically
conductive ferromagnetic transition metal. However, in the facing target
magnetron
sputtering process depicted in Figure 6, since targets 10 and 16 face each
other, there
will be a target-to-target interchange. Accordingly, the rare earth element,
even
though initially only on the bottom target 10, will eventually achieve some
steady-
state level in both targets due to inter-target exchange.
Throughout this discussion, it should be noted that the electrically
conductive ferromagnetic transition metal or alloy thereof can be any of the
metals or
alloys discussed above, including, for example, cobalt. Further, the rare
earth element
can be, for example, terbium, or any other suitable rare earth element.
Referring again now to Figures 7A and 7B, the magnet configuration
shown in Figure 7A is conventional in facing target magnetron sputtering. It
is
believed that this configuration is preferred for production since it will
result in a
relatively high deposition rate. Figure 7B shows an alternative configuration
for the
magnets 22, 24. In this case, like poles (illustrated as both north poles, but
alternatively could be both south poles) face each other. This configuration
has the
CA 02341857 2001-02-27
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undesirable effect of allowing secondary electrons to escape from the plasma
region,
thus lowering the deposition rate. However, it can be used to induce
perpendicular
anisotropy. This is because the growing film on the substrate 18 would be
bombarded
due to the divergence in the secondary electron paths caused by the
alternative lines of
5 magnetic force 44 as shown in Figure 7B. Thus, it will be appreciated that
the
secondary electrons in Figure 7A describe a confined path, while those in
Figure 7B
are not confined and would tend to follow a path as designated by arrow 46
The sputtering process just described can be performed so as to yield
the amorphous alloy having a composition of from about 10 atomic % terbium to
10 about 50 atomic % terbium, with the balance comprising cobalt, for example.
The
parameters to produce desired percentages of the rare earth element and the
ferromagnetic transition metal will be discussed further with respect to
Figure 10
below.
The aforementioned method for manufacturing the materials
15 containing the rare earth element can also be adapted to manufacture the
materials
containing the Heusler alloys. For example, Co, Mn and Sn or Co, Ti and Sn can
be
co-sputtered and then annealed to precipitate Co2MnSn or CoZTiSn respectively.
The
annealing or heat treating step can be performed, for example, in vacuum or an
inert
atmosphere (e.g., flowing argon) and the like. The forming gas annealing
atmospheres mentioned for the rare earth containing materials can be employed.
Throughout this application, for both the rare-earth-containing embodiments
and the
Heusler alloy embodiments, it is to be understood that " precipitated out"
generally
refers to the precipitation, out of an alloy, of the second phase referred to,
as described
throughout the application.
Other methods of forming precipitates are also within the scope of the
application. For example, rapid liquid quenching could be employed to form an
amorphous ribbon, wherein a molten mixture is cooled at a very high rate, for
example, about 106 °C/sec for example, by impinging on a water cooled
copper
surface or wheel. Subsequent annealing is then required to from the
precipitate.
Further, the aforementioned vapor deposition techniques can include
decomposition
of carbonyl compounds such as Fe(CO)5 to metals. A suitable mixture of metal
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16
carbonyl can be decomposed to form an amorphous alloy which precipitates a
magnetic phase out of a magnetic metallic matrix by annealing.
In the present invention using facing target magnetron sputtering
(FTMS), the magnetic field can be divergently applied to the plasma in plasma
region
20, which induces a high plasma potential as secondary electrons escape to the
chamber wall and other grounded parts of the system (refer to Figure 7B). When
amorphous TbCo thin films are deposited in a region of high plasma exposure in
FTMS, the deposition rates exhibit unusual changes with Ar pressure. The
changes of
deposition rate with pressure in the films deposited at different substrate
distances
from the system center line (LS) are shown in Figure 8. Circles represent
LS=5.4 cm
and squares represent LS 6.4 cm. The deposition rate (Angstroms per minute) of
the
films deposited in the low plasma-exposed region linearly increases with Ar
pressure.
In the films deposited in a high plasma-exposed region, the changes of
deposition
rates are not significant in the Ar pressure range from 6 mTorr to 12 mTorr.
The
difference between the pressure trend line and the observed rate is called
Or~p.
Figure 9 shows the film composition changes with Ar pressure at
different substrate distances from the system center line (LS). The Tb content
of the
films deposited in a low plasma-exposed region increases with Ar pressure. On
the
other hand, the Tb content in the films deposited in a high plasma region
suddenly
decreases and increases in the Ar pressure range which shows unusual changes
of
deposition rate. On the basis of quantitative analysis, the unusual changes
are due to
the preferential resputtering of Tb atoms in amorphous TbxCo,.x thin films. A
discussion of resputtering can be found in R.J. Gambino & J.J. Cuomo, 15 ~
V~s~, Sci.
Technoi. 296 (1978). Resputtering refers to the emission of atoms from the
material
forming on the substrate 18 due to ion bombardment thereof. The sputter yield
ratio
of Tb atoms to Co atoms in the resputtered films were estimated using the
model
proposed by Harper and Gambino, as described in J.M.E. Harper & R.J. Gambino,
Combined Ion Beam Deposition and Etching for Thin Film Studies, 161. V"~c.
Sci.
Technol. 1901 ( 1979).
30 The Harper and Gambino model is known to those of skill in the
plasma sputtering art. It provides a model for determining the composition of
the
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17
accumulating film AS,B,_y of two components on a target which is undergoing
resputtering. The value of y is determined from the equation:
y = [a + (az + 4x~3)'~] / 2(3, {2)
where:
a is a parameter given by:
a=(z+xE~-x-E~z-1), (3)
~i is a parameter given by:
~ _ {z + Er - Erz -1 ), {4)
x is the target composition AxB,_X and z is the fraction resputtered given by:
z = (RA + RB)/(FA + FB) {5)
where:
Rn = EaJY
Rs = EsJ{I-Y)
where EA is the sputter yield of component A in the film; EB is the sputter
yield of
component B in the film; J is the flux of the etching beam which bombards the
substrate on which the film grows (ions/cm2/sec); FA is the atom flux of
component A
arnving at the substrate (atoms/cm2/sec); FH is the atom flux of component B
arriving
at the substrate; and ET is the ratio of sputter yields defined by:
E, = Ep/Eg.
(8)
The fraction resputtered is strongly dependent on Ar pressure in FTMS
with a divergent magnetic field, as seen in Figure 7B. In order to apply the
above-
discussed Harper and Gambino model to resputtering effects in FTMS with a
composite target, certain modifications must be made. More specifically: the
parameter x in equations (2) - (4) must be redefined as film composition
without
25 resputtering instead of target composition; and the parameter E~ used in
equations (3)
and {4) above and defined in equation (8) must be redefined as the ratio of
sputter
yields in the resputtered film instead of the ratio of sputter yields. Thus,
E~ in the
model for the present invention is defined by:
Er = ~rA / E~ B, {9)
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18
where s~~ and e~B are the sputter yields in the resputtered film of components
A and B
respectively.
Therefore, combining the above equations, the sputter yield ratio, E~,
can be expressed as a function of compositions of the films with and without
resputtering (x and y) and fraction resputtered (z):
e~ _ $~n'I~rCo = ~(Y + x)( 1 + x - z) + y2(z - 1 )]/~(Y + x)(x - z) + Y2(z - 1
)]~ (10)
where E~'b and ~rc~ are sputter yields of Tb and Co, respectively. The
fraction
resputtered is obtained from the graph showing the unusual changes of
deposition rate
with pressure at Figure 8. The fraction resputtered {z) can be defined as the
ratio of
the deviation of the deposition rate from an ideal incremental line (Or~p) to
the ideal
deposition rate at a certain pressure (r~dealO at which it is assumed that the
ideal
deposition rate line is parallel to the other deposition rate line without
resputtering.
The film composition with and without resputtering can be determined from the
graphs on Tb contents with Ar pressure at different deposition distances. The
values
of parameters used to obtain resputter yield ratio using Eq. (9) are shown in
Table I
below.
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19
Pressure (mTorr) 5.6 7.3 8.3 9.6 11.6 12.6
Film composition 0.256 0.261 0.263 0.27 0.29 0.293
(x)
Film composition 0.246 0.22 0.226 0.257 0.281 0.288
(y)
Fraction resputtered0.09 0.125 0.13 0.13 0.125 0.1
(z)
Figure 10 exhibits sputter yield ratio of Tb to Co with pressures at
which resputtering takes place. The resputter yield ratio of Tb to Co are much
higher
than 20 in the range of pressure at which resputtering takes place. Therefore,
it can be
seen that the resputtering effect in amorphous TbXCo,.x thin films is mainly
caused by
Tb atoms. The resputter yield ratio has the largest value at 9.6 mTorr Ar
pressure
with high plasma potential and high plasma density. Note that, throughout the
foregoing, if a ferromagnetic transition metal other than Co or a rare earth
element
other than Tb were employed, a similar procedure could be used for all
calculations,
using different superscripts on the variables s~~' and e~c°, for
example.
Figure 11 shows the x-ray diffraction pattern of TbCo film annealed at
650°C in a nitrogen gas atmosphere. Sharp diffraction peaks are
observed showing
that the material is crystalline, not amorphous. The peaks correspond to
diffraction
from the (200), (220) and (311 ) crystal planes of TbN. Several peaks of face
centered
cubic (fcc) and hexagonal close packed (hcp) cobalt are also observed showing
that
crystalline cobalt is present. The diffraction peaks of (hcp) cobalt are sharp
indicating
larger crystallites.
Reference should now be had to Figure 12 in connection with a
magnetic recording system of the present invention and a method of detecting
magnetic field strength according to the present invention. The magnetic field
strength can, in some embodiments, correspond to digital information stored in
a
magnetic recording medium. Analog systems are also possible. As shown in
Figure
12, a magnetic recording system, designated generally as 100, is adapted for
use with
a magnetic recording medium 102 having a coercive force and adapted to store
data
(for example, digital data in the form of a plurality of bits (not amenable to
illustration)). The bits or other digital or analog data are stored as a
magnetization
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pattern in the medium 102; for example, the magnetization pattern may be
stored so
as to give rise to a z component of the H field at two different field
strengths to
represent binary logic levels. Medium 102 is shown as a tape-like element for
convenience, but it should be understood that medium 102 could also be a disk
5 element of a disk drive, or the like.
System 100 can include a write head 104 for writing the bits to the
medium 102 by producing a magnetic field of N ~ S or S ~ N polarity which
exceeds
the coercive force of the medium 102. Write head 104 can be of a conventional
type
made of a magnetic material and having a plurality of windings 106 and an air
gap
10 108. In some embodiments of the present invention, the write head 104 can
be
eliminated when it is desired to use the present invention in a read-only
situation, for
example, in a playback-only type of video cassette recorder.
System 100 further includes a read head designated generally as 110.
Read head 110 includes a portion 112 of a material exhibiting the GMR of any
of the
15 types discussed in this application. The portion 112 of magnetoresistive
material is
located in proximity to the medium 102. Portion 112 is coupled to a
resistivity (or
resistance) sensor 114 which detects changes in resistivity (or bulk
resistance) of the
portion 112 corresponding to magnetic field strength levels of the
magnetization
pattern in the medium 102 associated with the data such as the bits. As
discussed
20 above, OR/R = Op/p so resistivity and resistance are effectively
interchangeable with
respect to sensor 114. System 100 can further include a controller 116, of
which the
sensor 114 can be part, which controls the write head 104 and the read head
106 and
which converts the detected resistivity/resistance of the portion of GMR
material 112
to a signal (for example, a digital signal) which is indicative of the stored
data on the
25 medium 102 (for example, stored bits).
System 100 can further include a drive mechanism, depicted
schematically by arrows and boxes 118, which causes the medium 102 to move
past
the read head 110 and the write head 104. It is to be understood that drive
mechanism
118 can be any suitable mechanism known in the art and appropriate to the
character
30 of the medium 102; for example, a tape drive type mechanism, a disk drive
type
mechanism, and the like.
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It will be appreciated that resistivity/resistance sensor 114 can employ
any of a number of well-known techniques to sense the resistivity/resistance
of
portion 112 of the GMR material. For example, suitable bridge circuits can be
employed in order to determine the resistance of the portion 112, and knowing
the
shape and dimensions of the portion 112, the resistivity can then be
determined. As
discussed above, the percent change in both quantities is identical for a
given
geometry. The construction of resistivity/resistance sensor 114 and controller
116 can
be similar to those currently employed in prior-art systems utilizing the AMR
effect.
Typically, in digital applications, a comparator circuit is employed to
examine the
voltage drop across the portion 112 of GMR material for a suitable clipping
level.
Appropriate clocking circuitry, as is known in the art, can also be included
in
controller 116.
It should be noted that digital magnetic recording systems employing
the magnetoresistance effect do not need to rely on motion to produce a time-
changing magnetic flux as in older systems. Prior-art inductive read heads
produce a
voltage in the pickup coil which is proportional to the time rate of change of
magnetic
flux (d~/dt). Since the z-component of the H field present on magnetic
recording
medium 102 is sensed directly, system 100 can work with no motion or with slow
motion. This would be desirable, for example, in a video cassette recorder
wherein it
would no longer be necessary to spin the head to produce a large d~/dt in
order to
achieve "freeze-frame." As noted above, read heads according to the present
invention can be employed in many different types of devices, including
computer
disk drives, video cassette recorders, digital and analog audio tape decks ,
"minidisk"
playback systems, magnetic card readers such as credit card readers, and the
like.
Any type of digital or analog magnetic storage readout can be accomplished.
Other
applications unrelated to magnetic storage include servo systems such as
automatic
braking systems for autos, non-destructive testing (detection of eddy currents
around
defects), and the like.
Still with reference to Figure 12, a method of detecting magnetic field
strength according to the present invention will now be discussed. The
magnetic field
strength can be that of a magnetic field associated with any medium, for
example, the
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22
magnetic recording medium 102. One step of the method includes providing a
sensing head, such as read head 110 which comprises a portion 112 of a
magnetoresistive material exhibiting the GMR. The material can be any of the
materials discussed above. The method can also include exposing the sensing
head
5 such as read head 110 to the magnetic field of the medium such as magnetic
recording
medium 102 and sensing electrical resistivity of the portion 112 of the
material
exposed to the magnetic field. The resistivity varies due to the magnetic
field.
It is to be understood that throughout this discussion and the foregoing
discussion of the magnetic recording system 100, no explicit value for the
resistivity
10 of portion 112 need necessarily be calculated; for example, as discussed
above, a
voltage drop which depends on the resistivity can be used as a clipping level
or
threshold without explicit calculation of the resistivity. The method can also
include
converting the sensed electrical resistivity (or bulk resistance) of the
portion 112 to a
signal which is indicative of the magnetic strength of the magnetic field (or
15 magnetization pattern) associated with the medium such as magnetic
recording
medium 102. For example, an analog signal corresponding to the voltage drop
across
the portion 112 can be processed using the aforementioned comparator circuitry
to
produce a digital signal which corresponds to the sequence of bits (i.e., 1's
and 0's) on
magnetic recording medium 102. Instead of a magnetic recording medium 102, the
20 medium could be a material in an auto braking system or a material
undergoing non-
destructive testing(NDT). The magnetization pattern could be a series of bits
but can
be, for example, an analog pattern, a pattern for an auto brake system, or a
pattern
associated with the eddy currents around defects in NDT.
Still with reference Figure 12, it is to be appreciated that portion 112
25 can be coupled to resistivity sensor 114 using suitable leads 120, 122. The
ends of
leads 120, 122 can be located on portion 112 such that there is a defined
geometry
between the ends of the leads, permitting easy determination of the resistance
(and
corresponding voltage drop) due to changes in the resistivity. Windings 106 of
write
head 104 can also be coupled to controller 116 as shown.
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Amorphous TbXCo,.x thin films with the compositions of 25 - 32
atomic % of Tb
were prepared using facing target magnetron sputtering with a composite
target. The
parameter "x" is the atomic fraction of Tb, ranging from 0.25 to 0.32. The
film
compositions were controlled by changing the Ar sputtering gas pressure (5 ~
15
mTorr). In order to induce the phase separation of Co and TbN, nitrogen was
introduced into amorphous TbCo thin films by annealing at 650 °C for 12
hours with
a continuous flow of 10 % HZ-balance Nz gas mixture. (Annealing at a lower
temperature, such as 400°C, would be possible with a longer dwell
time). The
pressure of the gas mixture in the annealing furnace was about 1 atmosphere
and the
gas was flowed at a rate of approximately 500 standard cc/minute. The pressure
and
flow rate of the gas are not critical control parameters. Phase analysis was
made with
x-ray diffraction (See Figure 11 ) and with secondary electron images on a
field
emission SEM. Magnetization loops were made at room temperature in fields up
to
13 kOe using a vibrating sample magnetometer (VSM). The magnetization loops
and
magnetization vs. temperature with an applied field up to 30 kOe were measured
at
temperatures from 20 K to 300 K using a SQUID magnetometer. Magnetoresistance
at room temperature was measured with applied fields up to 8.5 kOe with a DC
electromagnet using Van der Pauw geometry. Electrical contacts were made with
fine
wires attached using silver paste at the corners of a square sample. The
magnetoresistance versus temperature was measured from 20 K to 250 K at 40 kOe
in
a superconducting coil cryostat. The Hall effect is large before annealing
when the
alloy is in the amorphous state.
Secondary electron SEM imaging shows that TbN becomes the
primary phase and precipitates out of a Co matrix. The TbN forms large (> 20
~c m)
surface patches and small(< 0.5 ~c m) precipitates in the Co matrix. The small
precipitates probably segregate to the grain boundaries of the Co matrix. The
X-ray
diffraction pattern from the film annealed at 650 °C exhibits strong
peaks of TbN and
the main peaks of the fcc and hcp Co structures which indicates that the Co
matrix
includes two structures, as seen in Figure 11. Both the Co and TbN phases are
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24
crystalline after the heat treatment, though the crystal structure of the Co
is somewhat
imperfect, possibly due to beginning with an amorphous material.
Various features of the magnetic and magnetotransport properties
obtained for the Co-TbN with 32% TbN composition are shown in Figures 1, 13,
14,
S 1 SA and 1 SB. The MR data
are typical of the behavior for all compositions from 25 mole % to 32 mole %
TbN.
The increase in magnetization with temperature and the broad minimum in
magnetization at an applied field of 1 kOe are indications of ferrimagnetic
behavior
(Fig. 13). The small jump in the magnetization curves at fields of 10, 20 and
30 kOe
at 50 K indicate that at these high fields the magnetic moments of the TbN
precipitates are ferromagnetically aligned with the Co matrix. The Curie
temperature
of TbN can be estimated as about 75 K by extrapolating from the break in the
magnetization curve at 30 kOe field. Figure 14 shows the magnetization curves
of
CO°,6g-(TbN)°,3z film with an applied field at various
temperatures. The magnetization
is not fully saturated even at the highest applied field of 30 kOe. Figures
15A and
15B show the resistivity (p) and resistivity change (Sp) of Co°.68-
(TbN)°.3z as a
function of magnetic field at room temperature. The curve shows a cusp-type
negative magnetoresistance at room temperature: the decrease of resistivity
with
increasing applied field. Considering that the electrical properties of the
rare-earth
nitrides DyN, HoN, and ErN are all metallic, N. Sclar, J. Ap~~ 1534 (1964),
and that the TbN has the same electronic structure as those nitrides, it can
be inferred
that the TbN precipitate is also an ordinary resistivity metal. The
magnetoresistivity
(8p) and magnetoresistance (8p/p) of Coo.b$-(TbN)°.3z are about 1. 12 x
10'7 S~cm and
0.72 % at room temperature up to 8 k0e field, respectively, where the sign of
the
magnetoresistance is negative (Figure 1).
The GMR effect of Co-TbN macroscopic fernmagnets can be
described in terms of the scattering of spin polarized conduction electrons by
the
antiparallel exchange coupled spins at the phase boundary between the TbN
precipitates and the Co matrix. R.J. Gambino and J. Wang, s_u~ra. In the Co-
TbN
30 system the matrix and precipitate are both metallic. The matrix is
ferromagnetic and
the precipitates are magnetically ordered through the exchange with the Co
matrix. In
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the Co-EuS system, the Co matrix is metallic but the EuS in the particles is a
semiconductor. Therefore, the carriers are mainly confined to the Co matrix,
which is
the main difference between Co-EuS and Co-TbN systems. In the Co-EuS system
the
conduction electrons are scattered mainly at the Co-EuS interface whereas in
the Co-
5 TbN system scattering can occur both at the Co/TbN interface and in the TbN
precipitates. That scattering depends on the magnetic alignment of EuS with
respect
to the cobalt. In the Co-TbN system, when the Garners pass through the phase
boundary between the two metallic phases, Co and TbN, they are scattered by
the
antiparallel exchange coupled spin and the resistivity is high in zero or low
fields. In
10 high fields with the Co and TbN ferromagnetically aligned, this spin
scattering
contribution is expected to disappear.
The resistance change of Coo,b$-(TbN)o.3z with temperature at H=0 and
40 kOe displays a behavior typically observed in metals, as shown in Figure 2.
The
GMR of Coo.68-(TbN)o.32 in the high field of 40 kOe is around 9 % at 250 K (
Figure
15 16), which is due to the increase of the ferromagnetic alignment between
the Co and
TbN by the high field. The GMR of Co-TbN shows an increase with temperature
(Figure 16). These data were obtained from the temperature dependence of
resistance
shown in Figure 2. In contrast, ordinary GMR materials have a negative
magnetoresistance, where the magnetoresistance decreases with increasing
20 temperature, as described in R.J. Gambino et al., 75 ;~Phys_. 6909 (1994).
Based on the magnetization curve with temperature at an applied field of 1 kOe
in
Figure 13 , the magnetization decreases with decreasing temperature, which can
be
explained by the increase of magnetization of the antiparallel exchange
coupled TbN
phase. The antiparallel exchange coupling between the Co and TbN with
different
25 magnetic moments may also become stronger with decreasing temperature. As a
result, the ferrimagnetic behavior between the two different magnetic moments
increases with decreasing temperature and thus the magnetization decreases.
The magnetization behavior with applied field can be divided by two
different regions around 150 K. Even though the magnetization curves shown in
30 Figure 14 are extrapolated to higher fields, the magnetization may not be
saturated in
a 40 kOe field at temperatures less than 150 K. On the other hand, at high
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26
temperatures the magnetization approaches saturation in a 40 kOe field. When
the
temperature increases above 150 K, the antiparallel exchange coupling of the
Co and
TbN decreases with increasing temperature and the Co and TbN are more easily
aligned ferromagnetically in a field. Therefore, the spin scattering
contribution is
expected to decrease with high applied fields with increasing temperature.
Thus the
magnetoresistivity, bp, and the GMR, 8p/p, both increase with increasing
temperature.
Figure 3 shows the dependence of GMR on the volume % of TbN
calculated from the composition assuming the normal densities for TbN and Co.
The
room temperature GMR (8p/p) of a Co,_x (TbN)x system at 8 kOe field increases
with
the TbN volume fraction. The 32 atomic % TbN composition which corresponds to
39 volume % of TbN has the largest room temperature GMR.
Thus, the new macroscopic ferrimagnet, Co-TbN, including TbN
precipitates in a cobalt matrix, has been formed by the transformation of
amorphous
TbCo to crystalline Co and TbN phases induced by annealing in an NZ gas
atmosphere. The fully transformed films annealed at 650 °C demonstrate
typical
macroscopic ferrimagnetic properties: evidence of negative exchange, magnetic
compensation and negative giant magnetoresistance. The antiparallel exchange
coupling at the phase boundary between the TbN precipitates and the Co matrix
can
explain all of these observations. It was found that the TbN magnetization
and/or the
ferrimagnetic exchange coupling at the phase boundary was increased with
decreasing
temperature. The temperature dependence of resistivity of the macroscopic
ferrimagnet COO_6g-(TbN)0.32 SNOWS the typical temperature dependence of a
metal.
The Co°.68-(TbN)°.sZ system has the largest values of 8p and
8p/p in the composition
range of 25 to 32 % TbN. The GMR effect was observed to increase with
increasing
temperature in the range 30 to 230 K, which was believed to be ascribed to the
increase of ferromagnetic alignment of the Co and TbN with a field caused by
the
weakening of exchange coupling by temperature.
Although the present invention has been described with reference to
specific exemplary embodiments, it should be understood that various changes,
substitutions and alterations can be made to the disclosed embodiments without
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27
departing from the spirit and scope of the invention as defined by the
appended
claims.
