U.S. patent application number 14/128163 was filed with the patent office on 2016-10-20 for nanostructured mn-al permanent magnets and method of producing same.
This patent application is currently assigned to THE TRUSTEES OF DARTMOUTH COLLEGE. The applicant listed for this patent is Ian Baker, Qi Zeng. Invention is credited to Ian Baker, Qi Zeng.
Application Number | 20160307677 14/128163 |
Document ID | / |
Family ID | 43229012 |
Filed Date | 2016-10-20 |
United States Patent
Application |
20160307677 |
Kind Code |
A1 |
Baker; Ian ; et al. |
October 20, 2016 |
Nanostructured Mn-Al Permanent Magnets And Method of Producing
Same
Abstract
A bulky consolidated nanostructured manganese aluminum alloy
includes at least about 80% of a magnetic .tau. phase and having a
macroscopic composition of Mn.sub.XAl.sub.YDo.sub.Z, where Do is a
dopant, X ranges from 52-58 atomic percent, Y ranges from 42-48
atomic percent, and Z ranges from 0 to 3 atomic percent. A method
for producing a bulky nanocrystalline solid of formula MnxAlyDoz is
provided. The method includes melting a mixture of metals to form a
substantially homogenous solution, casting the solution to form
ingots, measuring compositions of the ingots; crushing the ingots
to form crushed powders, and milling the crushed powders to form
nanocrystalline powders. The method further includes verifying the
presence of .tau. phase and determining the amount of the .tau.
phase, and simultaneously consolidating the nanocrystalline powders
into a bulky nanocrystalline solid and undergoing phase
transformation from .epsilon. phase to at least 80% .tau.
phase.
Inventors: |
Baker; Ian; (Etna, NH)
; Zeng; Qi; (Sinking Springs, PA) |
|
Applicant: |
Name |
City |
State |
Country |
Type |
Baker; Ian
Zeng; Qi |
Etna
Sinking Springs |
NH
PA |
US
US |
|
|
Assignee: |
THE TRUSTEES OF DARTMOUTH
COLLEGE
Hanover
NH
|
Family ID: |
43229012 |
Appl. No.: |
14/128163 |
Filed: |
March 20, 2012 |
PCT Filed: |
March 20, 2012 |
PCT NO: |
PCT/US2012/029812 |
371 Date: |
April 15, 2016 |
Related U.S. Patent Documents
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Application
Number |
Filing Date |
Patent Number |
|
|
13165595 |
Jun 21, 2011 |
8669822 |
|
|
14128163 |
|
|
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|
Current U.S.
Class: |
1/1 |
Current CPC
Class: |
H03H 9/02448 20130101;
H01F 1/08 20130101; B22D 7/005 20130101; B22F 3/24 20130101; C22C
1/02 20130101; C22C 22/00 20130101; B22F 2998/10 20130101; B22F
1/0003 20130101; B22F 9/04 20130101; B22F 2003/202 20130101; B22F
2003/248 20130101; H01F 1/047 20130101; B22F 2009/041 20130101;
H03H 9/2447 20130101; Y10T 29/42 20150115; B22F 3/20 20130101; B22F
2304/05 20130101; H03H 3/0072 20130101 |
International
Class: |
H01F 1/08 20060101
H01F001/08; B22F 1/00 20060101 B22F001/00; B22F 3/24 20060101
B22F003/24; B22D 7/00 20060101 B22D007/00; B22F 3/20 20060101
B22F003/20; C22C 22/00 20060101 C22C022/00; C22C 1/02 20060101
C22C001/02; H01F 1/047 20060101 H01F001/047; B22F 9/04 20060101
B22F009/04 |
Goverment Interests
GOVERNMENT INTERESTS
[0001] This invention was made with government support under
contract number 60NANB2D0120 awarded by the National Institute of
Standards and Technology (NIST). The government has certain rights
in the invention.
Foreign Application Data
Date |
Code |
Application Number |
Jul 2, 2010 |
EP |
10251200.1 |
Claims
1. A bulky consolidated nanostructured manganese aluminum alloy
comprising at least about 80% of a magnetic .tau. phase and having
a macroscopic composition of Mn.sub.XAl.sub.YDo.sub.Z, wherein Mn
is manganese, Al is aluminum, Do is a dopant, X ranges from 52-58
atomic percent, Y ranges from 42-48 atomic percent, and Z ranges
from 0 to 3 atomic percent.
2. The bulky consolidated nanostructured manganese aluminum alloy
of claim 1, wherein the manganese aluminum alloy further comprising
carbon, and having a macroscopic composition of 51 atomic percent
manganese, 46 atomic percent aluminum and 3 atomic percent
carbon.
3. The bulky consolidated nanostructured manganese aluminum alloy
of claim 2, wherein the permanent magnetic properties comprise
coercive forces of about 5.2 kOe.
4. The bulky consolidated nanostructured manganese aluminum alloy
of claim 1, wherein the manganese aluminum alloy has a macroscopic
composition of 54 atomic percent manganese, 46 atomic percent
aluminum.
5. The bulky consolidated nanostructured manganese aluminum alloy
of claim 4, wherein the permanent magnetic properties comprise
coercive forces of about 4.8 kOe.
6. A method for producing a bulky nanocrystalline solid comprising
melting a mixture of metals comprising between 52-58 atomic percent
manganese and between 42-48 atomic percent aluminum to form a
substantially homogenous solution; casting the solution to form
ingots; measuring compositions of the ingots; crushing the ingots
to form crushed powders; milling the crushed powders to form
nanocrystalline powders; verifying the presence of .tau. phase and
determining the amount of the .tau. phase; and simultaneously
consolidating the nanocrystalline powders into a bulky
nanocrystalline solid and undergoing phase transformation from
.epsilon. phase to at least 80% .tau. phase.
7. The method of claim 6, further comprising characterizing
microstructure of the bulky nanocrystalline solid and measuring
magnetic properties of the bulky nanocrystalline solid.
8. The method of claim 6, further comprising annealing the
nanocrystalline powders to determine conditions for consolidating
the nanocrystalline powders.
9. The method of claim 8, further comprising annealing at
temperatures between 200.degree. C. and 600.degree. C. to maximize
the amount of the magnetic metastable .tau. phase transformed from
a milled nanocrystalline unstable high-temperature .epsilon. phase,
thereby minimizing presence of non-magnetic equilibrium .beta. and
.gamma..sub.2 phases.
10. The method of claim 8, wherein the annealing time is shortened
for higher annealing temperature to avoid decomposition of the
.tau.-phase into .gamma..sub.2 and .beta. phases.
11. The method of claim 6, the step of consolidating the milled
powders comprising backpressure assisted equal channel angular
extrusion (ECAE).
12. The method of claim 11, further comprising increasing
backpressure to consolidate the nanocrystalline powders.
13. The method of claim 11, further comprising controlling a
temperature of the nanocrystalline powders within 200.degree. C. to
600.degree. C. during the ECAE to maximize presence of the .tau.
phase.
14. The method of claim 11, further comprising decreasing rate of
extrusion with increasing temperature to shorten annealing time at
higher temperature to avoid decomposition of the .tau.-phase into
.gamma..sub.2 and .beta. phases.
15. The method of claim 11, wherein the bulky nanocrystalline solid
is in a form of rod shapes.
16. The method of claim 15, wherein the bulky nanocrystalline solid
has a cross-section in one of square, rectangular, and circular
shape.
17. The method of claim 6, further comprising repeating the step of
consolidating the nanocrystalline powders until the bulky
nanocrystalline solid having minimum defects.
18. The method of claim 6, wherein the bulky nanocrystalline solid
is machinable.
19. The method of claim 6, wherein the mixture of metals further
comprises a dopant comprising at least one of carbon and boron.
20. The method of claim 6, wherein the mixture of metals comprises
54 atomic percent manganese, and 46 atomic percent aluminum.
21. The method of claim 6, wherein the bulky nanocrystalline solid
has a macroscopic composition of 54 atomic percent manganese and 46
atomic percent aluminum and coercive forces of about 5.2 kOe.
22. The method of claim 6, wherein the bulky nanocrystalline solid
has a macroscopic composition of 54 atomic percent manganese and 46
atomic percent aluminum and coercive forces of about 4.8 kOe.
23. A method for producing a bulky nanocrystalline solid
comprising: melting a mixture of metals comprising 51 atomic
percent manganese, 46 atomic percent aluminum and 3 atomic percent
carbon to form a substantially homogenous solution; casting the
solution to form ingots; measuring compositions of the ingots;
crushing the ingots to form crushed powders; milling the crushed
powders to form nanocrystalline powders; verifying the presence of
.tau. phase and determining the amount of the .tau. phase; and
simultaneously consolidating the nanocrystalline powders into a
bulky nanocrystalline solid and undergoing phase transformation
from .epsilon. phase to at least 80% .tau. phase.
24. The method of claim 21, wherein the bulky nanocrystalline solid
has a macroscopic composition of 51 atomic percent manganese, 46
atomic percent aluminum and 3 atomic percent carbon and coercive
forces of about 5.2 kOe.
25. The method of claim 21, wherein the bulky nanocrystalline solid
has a macroscopic composition of 54 atomic percent manganese and 46
atomic percent aluminum and coercive forces of about 4.8 kOe.
26. The method of claim 21, the step of consolidating the
nanocrystalline powders comprising backpressure assisted equal
channel angular extrusion (ECAE) and controlling a temperature of
the nanocrystalline powders within 200.degree. C. to 600.degree. C.
during the ECAE to maximum presence of the .tau. phase.
27. A permanent magnet comprising bulky consolidated nano
structured manganese aluminum alloy comprising at least about 80%
of a magnetic phase and having a macroscopic composition of
MnXAlYDoZ, wherein Mn is manganese, Al is aluminum, Do is a dopant,
X ranges from 5251-58 atomic %, percent, Y ranges from 42-48 atomic
%, percent, and Z ranges from 0 to 3 atomic percent %.
28. The magnet of claim 27 further comprising 3 atomic percent
carbon.
Description
RELATED APPLICATIONS
[0002] This application claims priority to U.S. patent application
Ser. No. 13/164,495 filed on Jun. 20, 2011, which is incorporated
by reference herein.
BACKGROUND
[0003] Magnets may be broadly categorized as temporary or
permanent. Temporary (soft) magnets become magnetized or
demagnetized as a direct result of the presence or absence of an
externally applied magnetic field. Temporary magnets are used, for
example, to generate electricity and convert electrical energy into
mechanical energy in motors and actuators. Permanent (hard) magnets
remain magnetized when they are removed from an external field.
Permanent magnets are used in a wide variety of devices including
motors, magnetically levitated trains, MRI instruments, and data
storage media for computerized devices.
[0004] High-performance permanent magnets, such as Sm--Co
(H.sub.C=10-20 kOe) and Nd--Fe--B (H.sub.C=9-17.5 kOe), are
generally intermetallic alloys made from rare earth elements and
transition metals, such as cobalt. Demand for high-performance
permanent magnets for motors is increasing rapidly for applications
such as wind turbine generators and motors in electric and hybrid
cars. Rare-earth magnets are generally used for such challenging
applications with, for example, each Toyota Prius using 1 kg of Nd
and a typical wind turbine generator using 250 Kg of Nd. These rare
earth magnets have highest energy product (BH).sub.max of any
material, where B is magnetic flux density and H is magnetic field
strength.
[0005] By way of example, Nd.sub.2Fe.sub.14B has the highest
(BH).sub.max at 45 MGOe. However, this material is not without
problems. Sintered Nd.sub.2Fe.sub.14B is vulnerable to grain
boundary corrosion, requiring nickel or copper/nickel plating or
lacquer coating. Although polymer-bonded Nd.sub.2Fe.sub.14B magnets
do not suffer from this grain boundary corrosion problem, they have
a significantly lower energy product due to the polymer matrix.
[0006] In another example, Sm--Co magnets have a (BH).sub.max of
28-30 MGOe, which is lower than that of Nd.sub.2Fe.sub.14B.
Although Sm--Co magnets do not suffer from corrosion and can be
used at higher temperatures (up to .about.350.degree. C.), they are
quite brittle, prone to chipping and can fracture from thermal
shock. Further issues with these materials are that over 95% of
rare earths are produced in one country and there are no US-owned
manufacturers of rare earth magnets. The high cost of rare earth
elements and cobalt makes the widespread use of high-performance
magnets commercially impractical.
[0007] Less expensive magnets are more commonly used, but these
magnets generally have lower coercivity H.sub.C i.e., their
internal magnetization is more susceptible to alteration by nearby
fields. For example, ferrites, which are predominantly iron oxides,
are the cheapest and most popular magnets, but they have both low
H.sub.C (1.6-3.4 kOe) and low values of M.sub.S. Similarly, Alnico
alloys, which contain large amounts of nickel, cobalt and iron and
small amounts of aluminum, copper and titanium, have H.sub.C in the
range of 0.6-2 kOe, which makes exposure to significant
demagnetizing fields undesirable.
[0008] The ferromagnetic .tau. phase in a MnAl magnet was first
reported by H. Kono, On the Ferromagnetic Phase in
Manganese-Aluminum Systems", Journal of the Physics Society of
Japan, 13 (1958) 1444 and Koch et al. "New Material for Permanent
Magnets on a Base of Mn and Al", Journal of Applied Physics, 31
(1960) 75S. This material is not used commercially in bulk form,
but continues to attract attention since it has an attractive
combination of magnetic properties for technological applications.
MnAl magnet has a theoretical energy product, (BH).sub.max, of 12
MGOe together with a relatively low density of 5200 kg/m.sup.3, and
thus a high density-compensated (BH).sub.max. In contrast, Sm--Co
magnets have a relatively high (BH).sub.max, but also a relatively
high density of .about.8300 kg/m.sup.3. Therefore, the MnAl magnet
has a comparable ratio of (BH).sub.max,/density as compared to the
Sm--Co magnet. Table 1 lists a comparison of the estimated maximum
energy product, density and density-compensated maximum energy
product for several classes of permanent magnets.
TABLE-US-00001 TABLE 1 ((BH).sub.max Density (BH).sub.max/Density
Magnet (MGOe) (kg/m.sup.3) (KGOe m.sup.3/kg) Ni.sub.2Fe.sub.14B 45
7600 5.92 Sm--Co 30 8300 3.61 Mn--Al 12 5200 2.31 AlNiCo 6 7000
0.86 Ferrites 4.5 5000 0.90
[0009] While the magnetic properties of MnAl magnets are superior
to conventional hard ferrites, Alnicos, and Fe--Cr--Co alloys, MnAl
is not as good as the rare-earth magnets. .tau.-MnAl does not
suffer from the issues associated with the rare-earth magnets
outlined above. Advantages of MnAl magnets include low costs and
excellent availability of the Mn and Al materials, good
machinability, high specific strength, high modulus of elasticity,
and excellent corrosion resistance.
[0010] More recently, Mn--Al--(C) alloys have been produced by
mechanical alloying processes. D. C. Crew, P. G. McCormick and R.
Street, Scripta Metall. Mater., 32(3), p. 315, (1995) and T. Saito,
J. Appl. Phys., 93(10), p. 8686, (2003) have shown that adding
small amounts of carbon (e.g., about 2 atomic % or less) to certain
Mn--Al alloys stabilizes the metastable .tau. phase and improves
magnetic properties and ductility. Crew et al. (1995) produced
Mn.sub.70Al.sub.30 weight % and Mn.sub.70.7Al.sub.28.2C.sub.1.1
weight % alloys by consolidating ball milled powders, annealing at
1050.degree. C. and then quenching, after which the materials were
no longer nanocrystalline. The resulting alloys had grain sizes of
about 300-500 nm and exhibited coercivities, H.sub.C, of 1.4 kOe
and 3.4 kOe, respectively. Saito (2003), produced mechanically
alloyed Mn.sub.70Al.sub.30 weight % and
Mn.sub.70Al.sub.29.5C.sub.0.5 weight % alloys that had grain sizes
of about 40-60 nm and coercivities of 250 Oe and 3.3 kOe,
respectively. In this study, the low coercivities reflected the
limited formation of the magnetic .tau. phase, which was determined
to be 10% in Mn.sub.70Al.sub.30 and 40% in
Mn.sub.70Al.sub.29.5C.sub.0.5. K. Kim, K. Sumiyama and K. Suzuki,
J. Alloys Comp., 217, p. 48, (1995), produced MnAl alloys that were
ball milled, but never annealed. The alloys displayed no hard
magnetic properties with H.sub.C of 130 Oe.
[0011] Mechanical milling (MM) has been used to synthesize a number
of rare earth permanent magnet alloys, including
Nd.sub.2Fe.sub.14B, Nd(Fe,Mo).sub.12N.sub.x and SmCo.sub.5. This
processing technique may be used to produce a nanocrystalline
microstructure having beneficial effects on the magnetic
properties. So far, only a few studies have dealt with the magnetic
behavior of MM Mn--Al. In one study by Satio, "Magnetic Properties
of Mn--Al System Alloys Produced by Mechanical Alloying", Journal
of Applied Physics, 93 (2003) 8686, a relatively high H.sub.C of
3.3 kOe was obtained, but the maximum M.sub.S was only 20 emu/g.
The same author Satio studied MM MnAl--C and obtained B.sub.r=60
emu/g and H.sub.C=1.95 kOe, see "Magnetic Properties of Mn--Al--C
Alloy Powders Produced by Mechanical Grinding", Journal of Applied
Physics, 97 (2005) 1. Commercial products with magnetic properties
of B.sub.r=5.75 kG, H.sub.C=3.0 kOe and (BH).sub.max=7 MGOe are
made by hot extrusion of annealed gas-atomized Mn--Al--C powders by
Sanyo Special Steel Co. Ltd., which has been working on MnAl--C
alloys for many years. These Mn--Al alloys are made from relatively
inexpensive materials, but the low coercivities remain a
problem.
SUMMARY
[0012] The subject matter of the present disclosure advances the
art and overcomes the problems outlined above by providing
nanostructured Mn--Al alloys and a method for their manufacture.
Constituents of these alloys may be mechanically milled and
heat-treated to form permanent room temperature magnets with high
coercivities and relatively high saturation magnetization
values.
[0013] In an embodiment, a bulky consolidated nanostructured
manganese aluminum alloy includes at least about 80% of a magnetic
.tau. phase and having a macroscopic composition of
Mn.sub.XAl.sub.YDo.sub.Z, where Do is a dopant, X ranges from 52-58
atomic %, Y ranges from 42-48 atomic %, and Z ranges from 0 to 3
atomic %. In a particular embodiment, the manganese aluminum alloy
includes carbon having a macroscopic composition of 51 atomic %
manganese, 46 atomic % aluminum and 3 atomic % carbon and has
coercive forces of about 5.2 kOe. In another particular embodiment,
the manganese aluminum alloy has a macroscopic composition of 54
atomic % manganese, 46 atomic % aluminum and coercive forces of
about 4.8 kOe. In an embodiment, a method for producing a bulky
nanocrystalline solid is provided. The method includes melting a
mixture of metals comprising between 52-58 atomic % manganese and
between 42-48 atomic % aluminum to form a substantially homogenous
solution. The method also includes casting the solution to form
ingots, measuring compositions of the ingots, crushing the ingots
to form crushed powders, and milling the crushed powders to form
nanocrystalline powders. The method also includes verifying the
presence of .tau. phase and determining the amount of the .tau.
phase. The method further includes simultaneously consolidating the
nanocrystalline powders into a bulky nanocrystalline solid and
undergoing phase transformation from .epsilon. phase to at least
80% .tau. phase, .beta. and .gamma..sub.2 phases.
[0014] In a particular embodiment, the method further includes
characterizing microstructure of the bulky nanocrystalline solid
and measuring magnetic properties of the bulky nanocrystalline
solid. The method further includes annealing the nanocrystalline
powders to determine conditions for consolidating the
nanocrystalline powders. The method further includes annealing at
temperatures between 200.degree. C. and 600.degree. C. to maximize
the amount of the magnetic metastable .tau. phase transformed from
a milled nanocrystalline unstable high-temperature .epsilon. phase,
thereby minimizing the presence of non-magnetic equilibrium .beta.
and .gamma..sub.2 phases. The annealing time is shorter for higher
annealing temperature to avoid decomposition of the .tau.-phase
into the .gamma..sub.2 and .beta. phases. The step of consolidating
the milled powders includes backpressure assisted Equal channel
angular extrusion (ECAE). The method further includes increasing
backpressure to consolidate the nanocrystalline powders. The method
further includes controlling a temperature of the nanocrystalline
powders within 200.degree. C. to 600.degree. C. during the ECAE.
The method also includes decreasing rate of extrusion with
increasing temperature to shorten annealing time at higher
temperature to avoid decomposition of the .tau.-phase into the
.gamma..sub.2 and .beta. phases.
[0015] In a particular embodiment, the bulky nanocrystalline solid
is in a form of rod shapes. The bulky nanocrystalline solid has a
cross-section in one of square, rectangular, and circular shape.
The method further includes repeating the step of consolidating the
nanocrystalline powders until the bulky nanocrystalline solid
having minimum defects. The bulky nanocrystalline solid is
machinable. The mixture of metals includes a dopant comprising at
least one of carbon and boron. The mixture of metals has 54 atomic
% manganese, and 46 atomic % aluminum. The mixture of metals
includes 51 atomic % manganese, 46 atomic % aluminum and 3 atomic %
carbon.
[0016] The benefits of the nanostructured Mn--Al and Mn--Al--C
permanent magnets include high coercivities (.about.4.8 kOe and 5.2
kOe, respectively) and high saturation magnetization values. The
benefits of the magnets also include relatively low cost and
readily available raw material supplies compared to rare earth
magnets.
BRIEF DESCRIPTION OF THE DRAWINGS
[0017] FIG. 1 is a flowchart illustrating a method of producing
magnetic alloys according to one embodiment.
[0018] FIG. 2 shows an X-ray diffraction pattern of
Mn.sub.54Al.sub.46 prior to annealing.
[0019] FIG. 3 shows an X-ray diffraction pattern of
Mn.sub.54Al.sub.46 annealed at 400.degree. C. for thirty
minutes.
[0020] FIG. 4 shows an X-ray diffraction pattern of
Mn.sub.54Al.sub.46 annealed at 500.degree. C. for thirty
minutes.
[0021] FIG. 5 shows an X-ray diffraction pattern of
Mn.sub.54Al.sub.46 annealed at 600.degree. C. for thirty
minutes.
[0022] FIG. 6 shows room temperature dependence of saturation
magnetization and coercive field on annealing temperatures for bulk
un-milled samples.
[0023] FIG. 7 shows room temperature dependence of saturation
magnetization and coercive field on annealing temperatures for
mechanically milled samples.
[0024] FIG. 8 shows room temperature hysteresis loops in a 15 kOe
field for mechanically milled (MM) and bulk Mn.sub.54Al.sub.46
powders annealed at 400.degree. C. for ten minutes.
[0025] FIG. 9 shows room temperature hysteresis loops in a 50 kOe
field for mechanically milled (MM) and bulk Mn.sub.54Al.sub.46
powders annealed at 400.degree. C. for ten minutes.
[0026] FIG. 10 shows room temperature isothermal remanence
magnetization (IRM) and dc demagnetization (DCD) curves for
mechanically milled Mn.sub.54Al.sub.46 annealed at 400.degree. C.
for ten minutes.
[0027] FIG. 11 shows room temperature isothermal remanence
magnetization (IRM) difference curves for bulk and mechanically
milled Mn.sub.54Al.sub.46 annealed at 400.degree. C. for ten
minutes.
[0028] FIG. 12 shows the room temperature dependence of the
coercive field on the magnetic field strength for mechanically
milled and bulk Mn.sub.54Al.sub.46 powders annealed at 400.degree.
C. for ten minutes.
[0029] FIG. 13 shows dependence of saturation magnetization on
annealing temperatures for mechanically milled and bulk samples of
various composition.
[0030] FIG. 14 shows dependence of coercivity on annealing
temperatures for mechanically milled and bulk samples of various
composition.
[0031] FIGS. 15(a) and (b) illustrate billets produced by ECAE of
commercial Al powder (a) four passes and (b) two passes.
[0032] FIGS. 16 (a) and (b) illustrate polished surface of (a) four
pass consolidated commercial Al powder and (b) two pass
mechanically-milled, nanocrystalline Al powder.
[0033] FIG. 17 illustrates X-ray diffraction patterns from
as-milled and ECAE-consolidated Al powder.
[0034] FIG. 18 illustrates the steps of production and
characterization of bulk fully dense nanocrystalline .tau.-MnAl
phase.
[0035] FIG. 19 illustrates a crystal structure of .tau. phase of
MnAl.
[0036] FIG. 20 illustrates a section phase diagram of MnAl.
DETAILED DESCRIPTION
[0037] Methods for producing mechanically milled, nanostructured
Mn--Al and Mn--Al--C alloys will now be shown and described. High
room temperature coercivities and saturation magnetization values
have been achieved for Mn--Al alloys that are produced by the
presently described methods, and it has been shown that the
addition of small amounts of carbon dopant (e.g., about 3 atomic %
or less) to Mn--Al alloys stabilizes the metastable .tau. phase and
improves magnetic properties.
[0038] Mechanically milled Mn--Al alloys possessing a
nanostructured ferromagnetic .tau. phase, with H.sub.C=4.8 kOe and
M.sub.S=87 emu/g at room temperature, were obtained by annealing
Mn.sub.54Al.sub.46 powders at 400.degree. C. for 10 minutes. The
coercivity value of this alloy is the highest ever reported for
Mn--Al materials. The amount of magnetic phase present in the
annealed product is estimated from the saturation magnetization to
be about 80% on the basis that M.sub.S of the pure .tau. phase is
about 110 emu/g. In another embodiment, a Mn--Al--C alloy,
Mn.sub.51Al.sub.46C.sub.3, prepared by the same method displayed a
coercivity that is the highest ever reported for Mn--Al--C
materials, H.sub.C=5.2 kOe.
[0039] The macroscopic formulas presented herein, e.g.,
Mn.sub.54Al.sub.46, pertain to the overall composition, but the
materials have nanostructure or microstructure of localized phase
variation (e.g., .gamma., .beta., and/or .tau. phases). As used
herein, a "nanostructured" material is a bulk solid characterized
by localized variation in composition and/or structure such that
the localized variation contributes to the overall properties of
the bulk material.
[0040] The large coercive forces observed are believed to result
from small grains of the magnetic .tau. phase (.about.30 nm) being
magnetically isolated from one another. This lack of magnetic
exchange coupling may result from non-magnetic phases (e.g.,
.beta., .gamma.) inhibiting changes in the alloy's internal
magnetization when an external magnetic field is applied (i.e., the
non-magnetic phase(s) act as magnetic domain wall pinning
sites).
[0041] The alloys disclosed herein are resistant to corrosion and
may, for example, be used in applications currently utilizing known
permanent magnets. In one embodiment, small particles or powders of
the alloys may be produced in a resin or plastic bonded form
according to known methods. The small grain size of the alloys may
provide improved ductility relative to materials with larger
grains.
[0042] FIG. 1 is a flowchart illustrating a method 100 of producing
magnetic alloys according to one embodiment. In a first step 102, a
mixture of metals, which may be in the form of ingots, powders,
ribbons, pellets or the like, is melted to provide a liquid
solution. In a second step 104, the liquid solution is quenched to
form a solid solution. Steps 102 and 104 may be repeated to ensure
that adequate mixing results in the formation of a substantially
homogeneous solid solution. A "substantially homogenous" solution
has a uniform structure or composition throughout, such that in a
randomized sampling of the solution at least 95% of the samples
would have consistent compositions. In step 106, the substantially
homogenous solid solution is reheated to a diffusion temperature
that is just below the melting temperature of the solid. The solid
is held at the diffusion temperature for a period of time that is
sufficient for the solid diffusion process to reach completion. For
example, the solid may be held at the diffusion temperature for
twenty hours. In step 108, the solid is quenched, e.g., with water,
to halt the diffusion process, and isolate the solid without
structural rearrangement that would otherwise occur in a slow
cooling process. In steps 110 and 112, the quenched solid is
crushed and milled to repeatedly fracture and cold weld the
particles in order to form a nanostructured material. The milling
is sufficient to cause a rupture of the crystals of the alloy as
well as to allow sufficient interdiffusion between the elementary
components. In step 114, the milled solid is annealed to ensure
complete formation of the nanostructured magnetic alloy.
Example 1
Production of Mn.sub.54Al.sub.46
[0043] Mn.sub.54Al.sub.46 alloy ingots were prepared by arc-melting
stoichiometrically balanced quantities of Mn and Al in a
water-cooled copper mold (T.sub.m.apprxeq.1250-1350.degree. C.).
The melted metallic solution was then heated until melted.
Quenching was performed by allowing the alloy to rapidly cool in
the copper mold to a temperature of .about.30.degree. C. in
approximately 10 minutes. Ingots were flipped and melted a minimum
of three times under argon to ensure mixing. Ingots were
subsequently heated to and held at 1150.degree. C. for 20 hours
followed by water quenching to retain the .epsilon. phase. The
ingots were then crushed and milled for eight hours in a hardened
steel vial using a SPEX 8000 mill containing hardened steels balls
with a ball-to-charge weight ratio of 10:1. The vials were sealed
under argon to limit oxidation. Both the as-milled powders and the
quenched bulk samples were annealed at temperatures from
350-600.degree. C. for 10-30 minutes to produce the ferromagnetic
.tau. phase.
[0044] The magnetic properties were measured at a room temperature
of about 20.degree. C. using a LakeShore 7300 vibrating sample
magnetometer (VSM) under an external magnetic induction field of 15
kOe. Some samples were also measured with an Oxford superconducting
quantum interference device (SQUID) magnetometer under a field of
50 kOe. Accuracy of the magnetic measurements is within .+-.2%.
Therefore, magnetic data may be reported as "about" a particular
value to account for ubiquitous sources of error (e.g., magnetic
fields within or near the magnetometer and errors associated with
weighing samples). Microstructural characterization was performed
using a Siemens D5000 diffractometer with a Cu X-ray tube and a
KeVex solid state detector set to record only Cu K.alpha.
X-rays.
[0045] FIGS. 2-5 show X-ray diffraction patterns of
Mn.sub.54Al.sub.46 annealed at various temperatures. X-ray
diffraction patterns for as-milled alloys showed peaks
corresponding to the .epsilon. phase of the MnAl alloy, where the
.epsilon. phase has a crystal structure of hexagonal close packed
(h.c.p). As shown in FIG. 2 the diffraction peaks were broad and of
low intensity, indicative of a nanocrystalline grain structure. The
grain size of the .epsilon. phase calculated from the (111) X-ray
peak using the Scherrer formula was 8 nm. Annealing the as-milled
sample of Mn.sub.54Al.sub.46 at 400.degree. C. for 30 minutes
caused the .epsilon. phase to transform to the .tau. phase which
has a crystal structure of face centered tetragonal (f.c.t.). FIG.
3 shows peaks indicative of the .tau. phase marked by asterisks.
The calculated .tau. phase grain size was .about.27 nm, which is
much smaller than that produced by conventional casting, grinding
or extruding. Without being bound by theory, the smaller grain size
appears to result from the .tau. phase forming from the
nanocrystalline .epsilon. phase. Increasing the annealing
temperature to 500.degree. C. for 30 minutes caused decomposition
of the .tau. phase, as shown in FIG. 4 by a decrease in intensity
of the .tau. phase peaks. Annealing at 600.degree. C. for 30
minutes resulted in a minimal presence of the .tau. phase in the
final product, as shown in FIG. 5.
[0046] These results show that the improved magnetic performance
may be related to small grain sizes, where the nanostructured
.epsilon. phase material is transformed to the ferromagnetic .tau.
phase at anneal conditions characterized by the 400.degree. C.
anneal which produced the results of FIG. 3. The effective
temperature range for this anneal is between 300.degree. C. and
600.degree. C., and more preferably from 350.degree. C. to
500.degree. C., and most preferably from 350.degree. C. to
450.degree. C. The smaller grain sizes are facilitated by the
milling that occurs just prior to the anneal.
[0047] FIGS. 6 and 7 show the sensitivity or dependence of
saturation magnetization, M.sub.S, and coercivity, H.sub.C, upon
annealing temperatures for both bulk (FIG. 6) and mechanically
milled (FIG. 7) Mn.sub.54Al.sub.46. For bulk samples, the M.sub.S
tends to increase with increasing annealing temperature from
300.degree. C. to 500.degree. C. The M.sub.S for mechanically
milled Mn.sub.54Al.sub.46 increases from 350.degree. C. to
400.degree. C., then decreases with increasing annealing
temperature from 400.degree. C. to 600.degree. C. This is
consistent with the X-ray diffraction data (FIGS. 3-5) that showed
the volume fraction of the magnetic .tau. phase decreasing with
annealing temperatures above 400.degree. C. The H.sub.C changes
relatively little from 350.degree. C. to 500.degree. C. for
mechanically milled samples. The optimal magnetic properties for
mechanically milled samples, H.sub.C=4.8 kOe, and M.sub.S=87 emu/g,
were obtained for Mn.sub.54Al.sub.46 powders annealed at
400.degree. C. for 10 minutes. The coercivity value of the
mechanically milled alloy is the highest reported to date for
Mn--Al magnetically isotropic powders. In general, the M.sub.S
obtained for annealed, mechanically milled samples was lower than
that obtained in bulk samples, while the H.sub.C was higher, due to
the small phase grain size.
[0048] FIGS. 8 and 9 show room temperature magnetic hysteresis
loops for mechanically milled (solid squares) and bulk (open
squares) Mn.sub.54Al.sub.46 powders annealed at 400.degree. C. for
10 minutes. FIG. 8 shows hysteresis loops in a 15 kOe field.
Coercivity is measured as the distance along the x-axis from the
origin to the intersection of the curve with the x-axis. It can be
seen that the mechanically milled sample has a much larger
coercivity (.about.5 kOe) than the bulk sample (.about.1 kOe).
Remanent magnetization, M.sub.r, is the intrinsic field of the
sample when the applied field is zero. M.sub.r of the mechanically
milled sample is approximately 35 emu/g, while that of the bulk
sample is approximately 25 emu/g. FIG. 9 shows hysteresis loops in
a 50 kOe applied field. Magnetic saturization, M.sub.S, has not
been reached, as evident from the increasing magnetization at high
fields. For the mechanically milled sample, the remanence ratio,
M.sub.r/M.sub.S, is about 0.5 when the applied field is 50 kOe,
which is characteristic of materials that are not
exchange-coupled.
[0049] FIGS. 10 and 11 show isothermal remanence magnetization
(IRM), dc demagnetization (DCD) and difference curves for
mechanically milled Mn.sub.54Al.sub.46 annealed at 400.degree. C.
for 10 minutes. FIG. 10 shows the IRM and DCD curve for the
mechanically milled sample, and FIG. 11 shows the .delta.M curves
for both mechanically milled and bulk samples annealed at
400.degree. C. for 10 minutes. Remanence curves and .delta.M plots
were used to determine the interaction between the .tau.-phase
grains. The DCD curve shows the progress of the irreversible
changes in magnetization. The IRM curve contains contributions from
both reversible and irreversible magnetization processes. The
change of magnetization .delta.M is defined as:
.delta.M=M.sub.d(H)-[M.sub.r(H.sub.sat)-2M.sub.r(H)]
Equation(1)
where M.sub.d (H) is the demagnetic remanent magnetization, M.sub.r
is remanent magnetization, and H.sub.sat is magnetic field strength
that saturates the magnet.
[0050] A plot of .delta.M versus H therefore gives a curve
characteristic of the interactions present. The overall negative
and small .delta.M for the mechanically milled sample indicates
that most of the .tau. phase nanograins are isolated with only
small dipolar interactions between them. No exchange coupling
exists in this nanostructured material, which explains why the
remanence ratio is close to 0.5.
[0051] FIG. 12 shows the dependence of the coercive field on the
magnetic field strength for mechanically milled and bulk
Mn.sub.54Al.sub.46 powders annealed at 400.degree. C. for 10
minutes. The bulk sample curve rises steadily to near saturation.
In contrast, the mechanically milled sample curve rises gradually
at low fields until the field strength approaches H.sub.C (5 kOe),
then it rises quickly to near saturation. This behavior indicates
that the mechanism for the magnetization process of the
mechanically milled material is controlled by domain wall pinning,
and that the applied field gradually removes the domain walls from
their pinning sites. The non-magnetic phase(s) that are present
could act as the pinning sites.
Example 2
Alloy Content Sensitivity
[0052] The manufacturing process of Example 1 was repeated by
varying the content of the Mn and Al metals, and doping with
carbon. FIGS. 13 and 14 show the dependence of saturation
magnetization and coercivity on annealing temperatures for
mechanically milled and bulk samples of various composition after
the samples had been annealed for thirty minutes. The legends of
FIGS. 13 and 14 show Mn content, and optionally C content, where
the remainder of the sample is Al. All samples are mechanically
milled, except for those labeled "bulk". It can be seen that 1-3
atomic % carbon decreased M.sub.S but increased H.sub.C in some
cases. In particular, Mn.sub.51Al.sub.46C.sub.3 had the highest
H.sub.C observed to date for a Mn--Al--C alloy, 5.2 kOe. Dopants
other than carbon may include boron. Generally, it can be noted
that because the .tau. phase is the only ferromagnetic phase in the
Mn--Al or Mn--Al--C systems, the saturation magnetization is
proportional to the percentage of the .tau. phase in the alloys.
When the Mn content is 50 atomic percent or less, little .epsilon.
phase can be developed, and therefore only a small amount of .tau.
phase can be produced. Also, when the Mn content is high, excess Mn
is used to stabilize the metastable .tau. phase. In this case, some
Mn atoms occupy lattice sites where they are coupled
antiferromagnetically to other nearby Mn atoms, thereby reducing
the magnetization. Thus, the Mn content is preferably between 52
and 58 atomic percent and the alloys may be described according to
Formula (1):
Mn.sub.XAl.sub.YDo.sub.Z, [0053] where Do is a dopant that may
include carbon, boron, X ranges from 52-58 atomic %, Y ranges from
42-48 atomic %, and Z ranges from 0 to 3 atomic %. In preferred
embodiments, Do is carbon, X ranges from 53-56 atomic %, Y ranges
from 44-47 atomic %, and Z ranges up to 3 atomic %. In a most
preferred embodiment, X is 54, Y is 46, and Do is not necessarily
present.
Example 3
Powder Consolidation by ECAE
[0054] The disclosure presents methods for producing
nanocrystalline .tau.-phase MnAl with a coercivity, H.sub.C, of 4.8
kOe, and a saturation magnetization, M.sub.S, of 87 emu/g by
mechanically-milling powders of the unstable, high-temperature
s-phase until they were nanocrystalline and then annealing the
milled powders. These values are the highest reported for Mn--Al
based powders (bulk magnets have M.sub.S of 110 emu/g). Good
magnetic properties resulted from producing the nanocrystalline
.epsilon.-phase first and then transforming the s-phase into the
i-phase rather than producing the .tau.-phase and then milling
i-phase to produce nanocrystalline material as conventionally
undertaken. The origin of the high H.sub.C may be the nanostructure
and/or the presence of small amounts of the equilibrium
.gamma..sub.2 and .beta. phases, which could pin the magnetic
domain walls.
[0055] While this improvement in magnetic properties is exciting,
the powders should be consolidated while retaining their good
magnetic properties in order for the material to be of practical
value. The heat treatment to which the milled .epsilon.-phase
powders have to subject in order to produce the superior magnetic
behavior provides a processing window to allow the warm
consolidation of the powders. This disclosure presents methods for
consolidating nanocrystalline MnAl powders into a fully dense solid
while producing the superior magnetic properties that can be
obtained in the powders.
[0056] The present disclosure further includes methods of producing
anisotropic magnets by warm extrusion, determination of the origin
of the high H.sub.C so that the heat treatment and composition of
the powders can be optimized to develop even better magnetic
properties. The milled nanocrystalline .epsilon.-phase powders,
with this heating, simultaneously consolidate to form a bulky solid
and undergo the necessary phase transformations to form the
required microstructure.
[0057] Equal channel angular extrusion (ECAE) is a powder
consolidation technique to consolidate microcrystalline powders by
Kim et al, "Equal Channel Pressing of Metallic Powders", Materials
Science Forum, 437-438 (2003) 89 and Xiang et al, "Microstructure
and mechanical properties of PM 2024Al-3Fe-5Ni alloy consolidated
by a new process, equal channel angular pressing", Journal of
Materials Science Letters, 16 (1997) 1725. Further, the use of
backpressure has been reported to be effective in aiding the
compaction of Mg alloy powders and shavings to full density by
Lapovok, "The positive role of back-pressure in equal channel
angular extrusion", Materials Science Forum, 503-504 (2006) 37.
[0058] ECAE was used to consolidate commercial Al powders, eutectic
Al--Si powders, unalloyed Mg powders and nanocrystalline Al
powders. The powders were cold-compacted into annealed copper cans,
onto which lids were pressed. Then, the cans were extruded at about
200.degree. C. through a well-lubricated hardened-steel ECAE jig. A
key to successful consolidation was to have a brass plug in front
of the copper can to provide backpressure during extrusion.
Extrusion without the plug led to poor powder consolidation,
particularly for the milled Al.
[0059] FIGS. 15(a) and 15(b) show billets that are 12.5 mm diameter
cylinders produced from Al powder extrusion through the ECAE jig
after (a) four ECAE passes and (b) two ECAE passes. FIGS. 16(a) and
16(b) show polished surfaces after (a) four ECAE passes and (b) two
ECAE passes. Note that FIG. 16 (a) shows a smoother polished
surface than FIG. 16(b). By increasing the number of passes, fewer
pores were observed in FIG. 16(a) than in FIG. 16 (b). After four
passes the Al was almost at full theoretical density, and the Al
was close to its full theoretical density even after two
passes.
[0060] An important feature of the consolidation of the
mechanically-milled Al powder is that the nanocrystalline grain
size of 55 nm produced by mechanical milling only marginally
increased during the ECAE extrusion to a grain size of 63 nm (see
FIG. 17) demonstrating the usefulness of this approach for
producing bulk nanocrystalline materials.
[0061] ECAE was also used to consolidate Mg-alloys without using a
can, i.e. the powder is loaded directly into the jig, reported by
Baker et al, "Containerless Consolidation of Mg Powders using
ECAE", Materials and Manufacturing Processes, 25(12) (2010)
1381-1384, which is incorporated by reference herein. Commercial Mg
can be consolidated to full density after two passes through the
ECAE jig at 200.degree. C.
[0062] ECAE can also be used to consolidate MnAl powders to form a
bulky MnAl solid in a way similar to consolidation of Mg powders or
Al powders. Extrusion may be performed at relatively lower
temperatures than is the case with conventional hot extrusion, in
order to avoid allowing the .epsilon. phase to fully transform to
the equilibrium .beta. and .gamma.2 phases during consolidation.
The temperature is a compromise between being as low as possible to
prevent significant grain growth and high enough to allow powder
consolidation. Variation in ECAE for consolidating MnAl powders may
include temperatures at which extrusion occur, the powders are
consolidated with and without a can, with a plug to provide
backpressure etc.
[0063] It should be noted that the nanocrystalline powders are
simultaneously consolidated and undergo the necessary phase
transformations from .epsilon. phase to magnetic phase and
non-magnetic .beta. and .gamma.2 phases to form the required
microstructure. The magnetic ti phase is at least 80%. The
coercivity H.sub.C of the bulky nanocrystalline solid remains
substantially unchanged from the nanocrystalline powders.
Specifically, MnAl bulky nanocrystalline solid can have a H.sub.C
of 5.2 kOe for 51 atomic % manganese, 46 atomic % aluminum and 3
atomic % carbon, and a H.sub.C of 4.8 kOe for 54 atomic %
manganese, 46 atomic % aluminum.
Example 4
Production of Bulk Fully Dense Nanocrystalline MnAl Solid
[0064] The relevant MnAl phase transformations simultaneously
consolidate nanocrystalline .epsilon.-MnAl phase powders and
transform them to the magnetic .tau.-MnAl phase. Understanding the
origin of the high H.sub.C in nanocrystalline .tau.-MnAl phase
assists processing and composition of the powders to optimize
magnetic properties. A working hypothesis is that it is the
nanoscale distribution of the non-magnetic .beta. and .gamma..sub.2
phases that produces the high H.sub.C, via domain-wall pinning, in
nanostructured .tau.-MnAl magnets.
[0065] FIG. 18 illustrates steps of production and characterization
of bulk fully dense nanocrystalline .tau.-MnAl phase. The steps
include the production of nanocrystalline MnAl powders, their
subsequent heat treatment; their consolidation using
backpressure-ECAE. The bulky solids are then characterized for
their magnetic properties using a vibrating sample magnetometer
(VSM), and for their microstructure using scanning electron
microscopy (SEM) coupled with electron backscattered patterns,
transmission electron microscopy (TEM) coupled with energy
dispersive X-ray spectroscopy (EDS) and convergent beam electron
diffraction (CBED), and atom probe tomography (APT).
[0066] At step 1802, casted ingots are produced from as-received
mixtures in the form of powders, flakes, pellets, ribbons or the
like (.about.100 g) of several compositions of Mn--Al near the
Mn.sub.54Al.sub.46 composition with or without carbon. The
as-received mixtures are arc melted and cast into a chilled copper
mold to form an ingot. Melting and casting each ingot three times
and flipping the ingot over in between may improve the homogeneity.
Since Mn has a low vapor pressure and can easily be lost during
melting, the compositions of the alloys is measured at step 1804,
including the oxygen and carbon content, after melting using a wet
chemistry approach. Typically, excess Mn is added to the starting
materials to compensate for any losses of Mn. The ingot contains
.epsilon.-MnAl.
[0067] At step 1806, the ingot is crushed into powders to prepare
for mechanical mill. The crushed powders have relatively large
grain sizes in the range of microns. Mechanical mill reduces the
grain sizes of the crushed powders to nanometer ranges, while the
s-phase of the MnAl alloy remains unchanged during crushing and
mechanical milling.
[0068] At step 1808, nanocrystalline powders are produced by
mechanically milling the crushed powders using a water-cooled Union
Process Szegvari attritor. Typically a rotation speed of
approximately 700 rpm is used with hardened steel balls for 30-40
hrs under an argon atmosphere. 50-100 g of powders are milled, with
a ball-to-powder ratio of 10:1. A SPEX mill may be used to more
rapidly (.about.8 h) produce nanocrystalline material than the
attritor, but only around 5 g of material are produced and more
contamination of the powder tends to occur with the SPEX. The
milled powders are now nanocrystalline .epsilon.-MnAl powders.
[0069] An understanding of the kinetics of the transformation may
inform about the conditions to use for the consolidation of the
nanocrystalline powders. It is much easier to run many different
anneals on the powders and to determine the phases present and the
magnetic properties than to have many consolidation runs and
determine microstructure and magnetic properties.
[0070] Nanocrystalline .epsilon.-MnAl powders may be optionally
annealed at different temperatures and times to determine the
optimum processing conditions for the best magnetic properties at
step 1810. The step 1810 is bounded in dashed lines to indicate
that this is an optional step. The starting point for the heat
treatments is a 30 min anneal at 400.degree. C. in a particular
embodiment. However, powders may be annealed for very short time at
higher temperature and for longer time at lower temperature.
Powders may also go through two-step anneals of different times at
different temperatures. Annealing temperature may be from
200.degree. C. to 600.degree. C. These anneals intend to maximize
the amount of the magnetic metastable .tau. phase transformed from
the as-milled, nanocrystalline unstable high-temperature .epsilon.
phase, thereby minimizing the amounts of the non-magnetic
equilibrium .beta. and .gamma..sub.2 phases present. If it is the
nanoscale distribution of the non-magnetic .beta. and .gamma..sub.2
phases that produces the high H.sub.C in MnAl, via domain-wall
pinning, a small amount of finely distributed .beta. and
.gamma..sub.2 has also to be present.
[0071] At step 1812, the nanocrystalline .epsilon.-MnAl powders are
consolidated into a bulky nanocrystalline solid, which includes at
least 80 % of .tau. phase and presence of .beta. and .gamma..sub.2
phases. The solid may be in a shape of rod. The rod may have a
cross-section in any shape including square, rectangular, circular,
etc. It will be appreciated by those skilled in the art that the
shape and dimension may vary for the nanocrystalline solid.
[0072] For consolidation of nanocrystalline powders, the ECAE
system includes the ECAE jig, cartridge heaters with thermocouple
feedback to control the temperature, and forward and backpressure
pistons. A variety of temperatures are used, based on the powder
annealing results outlined above, in order to consolidate the
powder. Higher backpressure is generally better for powder
consolidation. The powders may be consolidated in one pass or two
or more passes if a previous pass is unsuccessful. The consolidated
bulky nanocrystalline solid should have minimum defects or
substantially free of defects.
[0073] It is noted that the ECAE jig produced billets that are
about 15 mm diameter and 40 mm long. In principle, ECAE processing
can simply be scaled up using a larger jig and a larger extrusion
piston system, similar to a traditional direct extrusion set up. In
reality, the scale up probably requires lower pressures that simply
scale with the size of the billet. The reason for this is that in
small specimens the surface frictional forces in the billet have a
larger effect in small diameter specimens than large diameter
specimens due to the larger surface-to-volume ratio.
[0074] At step 1816, the annealed powders may be optionally
characterized for their microstructure and magnetic properties. The
step 1816 is bounded in dash lines to indicate that this is an
optional step. The size and morphology of the unmilled, milled and
annealed powders are determined using secondary electron imaging in
a field emission gun (FEG) XL30.TM. scanning electron microscope by
FEI company. The phases present in each of these annealed powders
are determined using a computer-controlled Riagku DMax rotating
anode X-ray diffractometer with a Cu target. The average grain size
(and lattice strain) are calculated from the corrected full width
at half maximum of each diffracted peak, .beta..sub.sample, for the
powders using a Hall-Williamson method:
.beta. sample cos .theta. = k .lamda. .delta. + 2 sin .theta.
Equation ( 2 ) ##EQU00001##
Where k is Scherrer constant, .delta. is the grain size, .lamda. is
the wavelength, .epsilon. is the internal strain introduce by
milling, and .theta. is the Bragg angle. .beta..sub.sample is
obtained from
.beta..sub.sample.sup.2=.beta..sub.measured.sup.2-.beta..sub.instrument.s-
up.2, where .beta..sub.instrument and .beta..sub.measured are the
Full Width Half Maximums (FWHM) of a well-annealed and a milled
specimen, respectively.
[0075] At step 1814, the bulky MnAl solid is characterized for its
microstructure and magnetic properties. For microstructural
characterization of extruded bulky MnAl solid, a FEI Tecnai F2 FEG
200 keV transmission electron microscope (TEM) is used. Energy
dispersive X-ray microanalysis (EDS) and convergent beam electron
diffraction (CBED) are performed using this instrument. Several
microstructural features are analyzed to determine the crystal
structures of the phases present using CBED, including whether any
.epsilon.-MnAl is left over after the processing. The chemistry of
the phases are measured using EDS. The grain size of the .tau.-MnAl
phase is determined using bright field imaging. The sizes of the
.beta. and .gamma..sub.2 particles are determined using dark field
imaging. The distribution of the .beta. and .gamma..sub.2
particles, i.e. whether they are homogeneously distributed, lie on
the .tau.-MnAl grain boundaries or lie in lines along the extrusion
direction, are determined using dark field imaging. The orientation
relationships between the .tau.-MnAl matrix and the .beta. and
.gamma..sub.2 particles are determined. The coherency of the .beta.
and .gamma..sub.2 particles with the .tau.-MnAl matrix are
determined. The dislocation density in the .tau.-MnAl matrix are
determined using a standard point line-intersection counting method
coupled with measurements of the foil thickness using the standard
two-beam CBED technique. The Burgers vectors of the dislocations
are determined using tilting experiments and the standard gb=0
(where g is the diffraction vector and b is the dislocation Burgers
vector) invisibility criterion.
[0076] When the phases present are small, such as within a few
nanometers, EDS may not be very useful for determining their
chemistry. In this case, a Cameca Local electrode atom probe (LEAP)
located at Oak Ridge National Laboratory (ORNL) is used. In the
LEAP, a high electric field is applied to a sharp needle-shaped
specimen, held in a high vacuum, to strip individual atoms from the
specimen atom layer by atom layer over thousands of layers and
identify them--including their location--by time-of-flight mass
spectrometry. The LEAP is somewhat complementary to the TEM since
it does not provide diffraction data.
[0077] Fine precipitates can be identified by using atom probe
tomography (API). The compositions of fine particles can be
determined using the APT. The extruded MnAl alloy may have a
texture. Thus, the grain orientations and grain misorientations are
determined using electron backscatter patterns (EBSPs) using the
FEI FEG XL30 SEM.
[0078] For magnetic measurements, the quasi-static magnetic
behavior of the powders and consolidated MnAl may be measured using
a Lakeshore Instruments 7300 VSM that can apply fields up to 15
kOe. This field strength cannot saturate the alloys, thus the
M.sub.S will be determined by extrapolation of H.sup.2.fwdarw.0,
from a plot of the saturation law M=M.sub.S (1-a/H.sup.2). The warm
extrusion via ECAE may produce anisotropic magnets which have
improved magnetic properties compared to isotropic magnets.
[0079] To interpret the measurements, a key question is whether the
.tau.-phase has to be nanocrystalline to get the superior magnetic
properties as seen in annealed powders, or whether it is sufficient
to have the equilibrium .beta. and .gamma..sub.2 phases distributed
on the nanoscale if these .beta. and .gamma..sub.2 phases are the
cause of domain wall pinning and, hence, the high H.sub.C.
Extruding the nanocrystalline powders at different temperatures may
lead to a wide variety of microstructures, i.e. different .tau.
grain structures and different distributions of the non-magnetic
equilibrium .beta. and .gamma..sub.2 phases. From analyzing these
microstructures and correlating them with the measured magnetic
properties, processing of the material can be improved and better
compositions can be determined.
[0080] FIG. 19 illustrates a crystal structure of .tau. phase of
MnAl. The crystal structure is tetragonal. The c/a ratio is close
to 1. The arrows indicate the magnetic dipoles. FIG. 20 illustrates
a section phase diagram of MnAl. Note that MnAl has a
high-temperature .epsilon.-phase, low-temperature equilibrium
.gamma..sub.2 and .beta. phases, and a metastable .tau. phase.
[0081] It is known that Mn metal is ordinarily antiferromagnetic.
By increasing the atomic distance between Mn atoms to 2.96 .ANG. or
more, the element becomes ferromagnetic. The ferromagnetism of the
.tau.-phase in MnAl occurs because the magnetic moments of Mn atoms
in 0, 0, 0 sites are parallel to one another (see FIG. 19).
[0082] Although various mechanisms have been proposed for
.tau.-phase formation, the generally accepted one is that the
high-temperature non-magnetic .epsilon.-phase (h.c.p.) transforms
into a non-magnetic .epsilon.'-phase (orthorhombic) by an ordering
reaction, and then transforms into a ferromagnetic .tau.-phase by a
martensitic phase transition, i.e.
.epsilon..fwdarw..epsilon.'.fwdarw..tau.. However, the
transformation from .epsilon. to .tau. may also involve diffusion,
and a nucleation and growth process or a massive transformation, as
suggested by recent electron microscopy observations and kinetic
analysis. The high density of lattice defects within the
.tau.-phase that develops during the phase transformation is
attributed to growth faults produced during atomic attachment at
the migrating interface. Practically, the tetragonal .tau. phase,
which is metastable, is usually produced either by a rapid
quenching of the high temperature .epsilon. phase followed by
isothermal annealing between 400.degree. C. and 700.degree. C., or
by cooling the .epsilon. phase at a rate of .about.10.degree.
C./min. Prolonged annealing and elevated temperatures result in
decomposition of the .tau. phase into the equilibrium cubic
.gamma..sub.2 and .beta. phases (see FIG. 20).
[0083] In order to stabilize the .tau.-phase, carbon has been
introduced into Mn--Al. Carbon reduces the Curie temperature
T.sub.C, and the magnetic anisotropy field, but increases M.sub.S
with a larger resultant Mn moment. It should be pointed out that
Mn--Al--C magnets have a very low Curie temperature T.sub.C of
.about.290-300.degree. C., compared with .about.700-800.degree. C.
for Alnico and .about.450.degree. C. for ferrite. The workability
is also improved due to the small C atoms that relieve internal
lattice stresses. In Mn.sub.53.6Al.sub.44.6 alloys, the best
magnetic properties were obtained for a carbon content just above
the solubility limit of carbon atoms (1.7 atomic %) because of the
formation of non-magnetic Mn.sub.3AlC precipitates. The magnetic
hysteresis behavior of Mn--Al--C is extremely sensitive to the
microstructure and defects introduced during the formation of the
.tau. phase within the high temperature .epsilon. phase. So far,
useful permanent magnets have been obtained only by doping the
alloy with carbon and extruding them. Anisotropic Mn--Al--C magnets
have been produced by subjecting the ternary alloys to warm
extrusion. The properties of the extruded material are a result of
the high anisotropy, grain size reduction and carbide
precipitations.
[0084] It is desirable that the resulting bulk material is able to
be machined for engineering applications. In order to determine the
utility of the material for this purpose and determine its
structural integrity, tensile tests may be performed at
1.times.10.sup.-4 s.sup.-1 in air on specimens from the
ECAE-processed rods after different extrusions to determine their
yield strengths and elongation at room temperature. The fracture
surfaces may be examined in the SEM. Their behavior will be
compared to that of the non-nanocrystalline MnAl. Additionally,
correlating the magnetic properties and phase transformations with
the microstructure via modeling help understand the magnetic
behavior and further refine the processing and alloy compositions.
Modeling that relates the observed phases to the temperature and
time at temperature as well as the extrusion pressures may be used
to guide the processing to obtain the optimum conditions. Software
Thermo-Calc and DICTRA may be used to model the phase
transformations.
[0085] The above description of the specific embodiments may be
modified and/or adapted for various applications or uses that do
not depart from the general scope hereof. Therefore, such
adaptations and modifications should and are intended to be
comprehended within the meaning and range of equivalents of the
disclosed embodiments. It is to be understood that the phraseology
or terminology employed herein is for the purpose of description
and not limitation.
[0086] This specification contains numerous citations to references
such as patents, patent applications, and publications. Each is
hereby incorporated by reference.
* * * * *