Note: Descriptions are shown in the official language in which they were submitted.
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IRON ALUMINIDE COMPOSITE AND METHOD OF
MANUFACTURE THEREOF
The invention relates generally to iron aluminide composites and method of
manufacture thereof.
Iron base alloys containing aluminum can have ordered and disordered body
centered crystal structures. For instance, iron aluminide alloys having
intermetallic alloy
compositions contain iron and aluminum in various atomic proportions such as
FejAl,
FeAI, FeAl2, FeAl3, and FezAls. Fe3A1 intermetallic iron aluminides having a
body
centered cubic ordered crystal structure are disclosed in U.S. Patent Nos.
5,320,802;
5,158,744; 5,024,109; and 4,961,903. Such ordered crystal structures generally
contain 25
to 40 atomic % A1 and alloying additions such as Zr, B, Mo, C, Cr, V, Nb, Si
and Y.
An iron aluminide alloy having a disordered body centered crystal structure is
disclosed in U.S. Patent No. 5,238,645 wherein the alloy includes, in weight %
,
8-9.5 Al, s 7 Cr, s 4 Mo, s 0.05 C, _< 0.5 Zr and s 0.1 Y, preferably 4.5-5.5
Cr,
1.8-2.2 Mo, 0.02-0.032 C and 0.15-0.25 Zr. Except for three binary alloys
having 8.46,
12.04 and 15.90 wt % Al, respectively, all of the specific alloy compositions
disclosed in
the '645 patent include a minimum of 5 wt % Cr. Further, the '645 patent
states that the
alloying elements improve strength, room-temperature ductility, high
temperature oxidation
resistance, aqueous corrosion resistance and resistance to pitting. The '645
patent does not
relate to electrical resistance heating elements and does not address
properties such as
thermal fatigue resistance, electrical resistivity or high temperature sag
resistance.
Commonly owned U.S. Patent Nos. 5,595,706 and 5,620,651 disclose iron base
alloys containing aluminum which are useful for electrical resistance heating
elements.
Examples of heating element configurations can be found in commonly owned U.S.
Patent
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Nos. 5,530,225 and 5,591,368. Other examples of electrical resistance heating
elements
can be found in commonly owned U.S. Patent Nos. 5,060,671; 5,093,894;
5,146,934;
5,188,130; 5,224,498; 5,249,586; 5,322,075; 5,369,723; and 5,498,855.
A 1990 publication in Advances in Powder Metallurgy, Vol. 2, by J.R. Knibloe
et
al., entitled "Microstructure And Mechanical Properties of P/M Fe~AI Alloys",
pp. 219-
231, discloses a powder metallurgical process for preparing Fe3A1 containing 2
and 5% Cr
by using an inert gas atomizer. This publication explains that Fe~AI alloys
have a D03
structure at low temperatures and transform to a B2 structure above about
550°C. To make
sheet, the powders were canned in mild steel, evacuated and hot extruded at
1000 ° C to an
area reduction ratio of 9:1. After removing from the steel can, the alloy
extrusion was hot
forged at 1000 °C to 0.340 inch thick, rolled at 800 ° C to
sheet approximately 0.10 inch
thick and finish rolled at 650°C to 0.030 inch. According to this
publication, the atomized
powders were generally spherical and provided dense extrusions and room
temperature
ductility approaching 20% was achieved by maximizing the amount of B2
structure.
A 1991 publication in Mat. Res. Soc. Symp. Proc., Vol. 213, by V.K. Sikka
entitled "Powder Processing of Fe3A1-Based Iron-Aluminide Alloys," pp. 901-
906,
discloses a process of preparing 2 and 5 % Cr containing Fe3A1-based iron-
aluminide
powders fabricated into sheet. This publication states that the powders were
prepared by
nitrogen-gas atomization and argon-gas atomization. The nitrogen-gas atomized
powders
had low levels of oxygen (130 ppm) and nitrogen (30 ppm). To make sheet, the
powders
were canned in mild steel and hot extruded at 1000°C to an area
reduction ratio of 9:1.
The extruded nitrogen-gas atomized powder had a grain size of 30 ,um. The
steel can was
removed and the bars were forged 50 % at 1000 ° C , rolled 50 % at 850
° C and finish rol led
50 % at 650 ° C to 0. 76 mm sheet.
A paper by V.K. Sikka et al., entitled "Powder Production, Processing, and
Properties of Fe3A1", pp. 1-11, presented at the 1990 Powder Metallurgy
Conference
Exhibition in Pittsburgh, PA, discloses a process of preparing FejAl powder by
melting
constituent metals under a protective atmosphere, passing the metal through a
metering
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nozzle and disintegrating the melt by impingement of the melt stream with
nitrogen
atomizing gas. The powder had low oxygen (130 ppm) and nitrogen (30 ppm) and
was
spherical. An extruded bar was produced by filling a 76 mm mild steel can with
the
powder, evacuating the can, heating 1 1/2 hr at 1000°C and extruding
the can through a 25
mm die for a 9:1 reduction. The grain size of the extruded bar was 20 ~cm. A
sheet 0.76
mm thick was produced by removing the can, forging 50% at 1000°C,
rolling 50% at
850 ° C and finish rolling 50 % at 650 °C.
Oxide dispersion strengthened iron-base alloy powders are disclosed in U.S.
Patent
Nos. 4,391,634 and 5,032,190. The '634 patent discloses Ti-free alloys
containing 10-40%
Cr, 1-10% A1 and s 10% oxide dispersoid. The ' 190 patent discloses a method
of forming
sheet from alloy MA 956 having 75 % Fe, 20 % Cr, 4.5 % Al, 0.5 % Ti and 0.5 %
Y203.
A publication by A. LeFort et al., entitled "Mechanical Behavior of FeAI,~
Intermetallic Alloys" presented at the Proceedings of International Symposium
on
Intermetallic Compounds - Structure and Mechanical Properties (JIMIS-6), pp.
579-583,
held in Sendai, Japan on June 17-20, 1991, discloses various properties of
FeAI alloys (25
wt % Al) with additions of boron, zirconium, chromium and cerium. The alloys
were
prepared by vacuum casting and extruding at 1100°C or formed by
compression at 1000°C
and 1100°C. This article explains that the excellent resistance of FeAI
compounds in
oxidizing and sulfidizing conditions is due to the high A1 content and the
stability of the B2
ordered structure.
A publication by D. Pocci et al., entitled "Production and Properties of CSM
FeAI
Intermetallic Alloys" presented at the Minerals, Metals and Materials Society
Conference
(1994 TMS Conference) on "Processing, Properties and Applications of Iron
Aluminides",
pp. 19-30, held in San Francisco, California on February 27 - March 3, 1994,
discloses
various properties of Fe~AI intermetallic compounds processed by different
techniques such
as casting and extrusion, gas atomization of powder and extrusion and
mechanical alloying
of powder and extrusion and that mechanical alloying has been employed to
reinforce the
material with a fine oxide dispersion. The article states that FeAI alloys
were prepared
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having a B2 ordered crystal structure, an A1 content ranging from 23 to 25 wt
% (about 40
at % ) and alloying additions of Zr, Cr, Ce, C, B and Y203. The article states
that the
materials are candidates as structural materials in corrosive environments at
high
temperatures and will fmd use in thermal engines, compressor stages of jet
engines, coal
gasification plants and the petrochemical industry.
A publication by J. H. Schneibel entitled "Selected Properties of Iron
Aluminides",
pp. 329-341, presented at the 1994 TMS Conference discloses properties of iron
aluminides. This article reports properties such as melting temperatures,
electrical
resistivity, thermal conductivity, thermal expansion and mechanical properties
of various
FeAI compositions.
A publication by J. Baker entitled "Flow and Fracture of FeAI", pp. 101-115,
presented at the 1994 TMS Conference discloses an overview of the flow and
fracture of
the B2 compound FeAl. This article states that prior heat t1'eatmPntc etrnnolv
~ffo~r rt.e
mechanical properties of FeAI and that higher cooling rates after elevated
temperature
annealing provide higher room temperature yield strength and hardness but
lower ductility
due to excess vacancies. With respect to such vacancies, the articles
indicates that the
presence of solute atoms tends to mitigate the retained vacancy effect and
long term
annealing can be used to remove excess vacancies.
A publication by D.J. Alexander entitled "Impact Behavior of FeAI Alloy FA-
350",
pp. 193-202, presented at the 1994 TMS Conference discloses impact and tensile
properties
of iron aluminide alloy FA-350. The FA-350 alloy includes, in atomic % , 35 .
8 % Al,
0.2 % Mo, 0.05 % Zr and 0.13 % C.
A publication by C.H. Kong entitled "The Effect of Ternary Additions on the
Vacancy Hardening and Defect Structure of FeAI", pp. 231-239, presented at the
1994
TMS Conference discloses the effect of ternary alloying additions on FeAI
alloys. This
article states that the B2 structured compound FeAI exhibits low room
temperature ductility
and unacceptably low high temperature strength above 500°C. The article
states that room
temperature brittleness is caused by retention of a high concentration of
vacancies following
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high temperature heat treatments. The article discusses the effects of various
ternary
alloying additions such as Cu, Ni, Co, Mn, Cr, V and Ti as well as high
temperature
annealing and subsequent low temperature vacancy-relieving heat treatment.
5 The invention provides an iron aluminide composite comprising iron
aluminide, an
oxide filler and an additive which improves metallurgical bonding of the oxide
filler to the
iron aluminide. The oxide filler can comprise alumina, zirconia, yttria, rare
earth oxide
and/or beryIlia. The additive can comprise a refractory carbide such as TiC,
HfC and/or
ZrC. A preferred ratio of oxide: additive is 1 to 3. The composite can be used
for various
devices such as tool bits, structural components or electrical resistance
heating elements in
devices such as heaters. According to a preferred embodiment, the composite
comprises a
liquid phase sintered composite.
The iron aluminide preferably comprises a binary alloy of iron and aluminum or
an
alloy. For instance, the iron aluminide alloy can comprise, in weight % , 14-
32 % Al, <_
2.0%Ti,s2.0%Si,s30%Ni,s0.5%Y,slS%Nb,sl%Ta,s3%W,slO%Cr,
s 2.0% Mo; s 1 % Zr, s 1 % C and s 0.1 % B. The oxide filler preferably
comprises
alumina which can be present in any desired amount such as s 40% . The
additive
preferably comprises s 40% TiC.
According to various preferred aspects of the invention, the composite can be
Cr-
free, Mn-free, Si-free, and/or Ni-free. The composite can include non-oxide
filler ceramic
particles such as SiC, Si3N4, A1N, etc. Preferred iron aluminide alloys
include 20.0-3I.0%
Al, 0.05-0.15 % Zr, s 3 % W, s 0.1 % B and 0.01-0.2 % C; 14.0-20.0 % Al, 0.3-
1.5 % Mo,
0.05-1.0 % Zr, s 3 % W and s 0.2 % C, s 0.1 % B and s 2. 0 % Ti; and 20.0-31.
0 % AI,
0.3-0.5 % Mo, 0.05-0.3 % Zr, _< 0.2 % C, s 2 % W, s 0.1 % B and s 0.5 % Y.
The electrical resistance heating element can be used for products such as
heaters,
toasters, igniters, heating elements, etc. wherein the composite has a room
temperature
resistivity of 80-400, S2 ~ cm, preferably 90-200 ,u ~2 ~ cm. The composite
preferably heats
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to 900°C in less than 1 second when a voltage up to 10 volts and up to
6 amps is passed
through the alloy. When heated in air to 1000°C for three hours, the
composite preferably
exhibits a weight gain of less than 4 % , more preferably less than 2 % . The
composite
preferably exhibits thermal fatigue resistance of over 10,000 cycles without
breaking when
pulse heated from room temperature to 1000°C for 0.5 to 5 seconds.
With respect to mechanical properties, the composite has a room temperature
flexure strength of at least 300 MPa in the liquid phase sintered condition
and at least 1000
MPa in the hot forged condition.
The invention also provides a powder metallurgical process of making an iron
aluminide composite by forming a mixture of iron aluminide powder, oxide
powder and an
additive which promotes adhesion of the oxide powder to the iron aluminide,
forming the
powder mixture into a body and sintering the body. According to various
aspects of the
method, the body can be formed by hot or cold pressing and the sintering can
comprise
solid state, partial liquid or liquid phase sintering. For example, the
forming can be carried
out by placing the powder in a metal can, sealing the metal can with the
powder therein,
and hot pressing or hot extruding the metal can. Alternatively, the body can
be made by
liquid phase infiltration of an iron aluminide matrix into a mass of oxide
filler particles. In
order to densify and/or shape the sintered body, the sintered body can be hot
forged or
subjected to other working steps such as cold working, extrusion, rolling,
etc. If desired,
the powder mixture can be cold pressed prior to sintering and/or annealed
subsequent to
sintering.
Brief Descripition of the Dry
Figure 1 shows an X-ray diffraction pattern for an FeAI/A1z03 composite in
accordance with the invention;
Figure 2 shows an X-ray diffraction pattern for an FeAI/ZrOz composite in
accordance with the invention;
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Figure 3 shows a scanning electron microscope image of an FeAI/Zr01 composite
in
accordance with the invention;
Figure 4 shows exudation of FeAI during liquid phase sintering of an
FeAI/A>z03
composite which did not include a TiC additive in accordance with the
invention;
Figure 5 shows the effect of TiC on improving liquid infiltration of AIz03 of
iron
aluminide;
Figure 6 shows a scanning electron microscope image of a polished section of
an
FeAI/TiC/AI203 composite in accordance with the invention;
Figure 7 shows a hot forged coupon of Fe-lSTiC-l5AIz03 (vol. % ) in accordance
with the invention wherein the interior of the coupon is sound and some edge
cracking is
evident around the exterior of the coupon;
Figure 8 is an optical micrograph of a liquid phase sintered composite of FeAI-
16.STiC-16.SAI203 (vol. % ) in accordance with the invention;
Figure 9 is an optical micrograph of a hot forged composite of FeAI-lSTiC-
l5A1z03
(vol. % ) in accordance with the invention;
Figure 10 is a graph of stress versus crosshead displacement produced during a
flexure stress test of a composite of FeA1-lSTiC-15A1z03 (vol. % ) in
accordance with the
invention; and
Figure 11 is a graph of load versus crosshead displacement produced during a
fracture toughness test of a composite of FeAI-lSTiC-l5AIz03 (vol. % ) in
accordance with
the invention.
The present invention is directed to iron aluminide composites including iron
aluminide, an oxide filler and an additive which improves metallurgical
bonding of the
oxide filler to the iron aluminide. According to one aspect of the invention,
the iron
aluminide can include an iron concentration ranging from 4 to 32 % by weight
(nominal)
and the oxide filler can comprise one or more oxides such as alumina,
zirconia, yttria, rare
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earth oxide and/or beryllia. The additive preferably comprises at least one
refractory
carbide, refractory nitride or refractory boride such as TiC, HfC, ZrC, TiN,
HfN, ZrN,
TiB2, HfB2 and/or ZrBz.
The concentration of the alloying constituents used in forming the iron
aluminide is
expressed herein in nominal weight percent. However, the nominal weight of the
aluminum essentially corresponds to at least about 97% of the actual weight of
the
aluminum in the iron aluminide. For example, in a preferred composition, a
nominal 18.46
wt % may provide an actual 18.27 wt % of aluminum, which is about 99 % of the
nominal
concentration.
The iron aluminide can be processed or alloyed with one or more selected
alloying
elements for improving properties such as strength, room-temperature
ductility, oxidation
resistance, aqueous corrosion resistance, pitting resistance, thermal fatigue
resistance,
electrical resistivity, high temperature sag or creep resistance and
resistance to weight gain.
The iron aluminide composite can be used to make heating elements for various
devices
such as described in commonly owned U.S. Patent No. 5,530,225 or 5,591,368.
However,
the composite can be used for other purposes such as in thermal spray
applications wherein
the composite could be used as coatings having oxidation and corrosion
resistance. Also,
the composite can be used as oxidation and corrosion resistant electrodes,
furnace
components, chemical reactors, sulfidization resistant materials, corrosion
resistant
materials for use in the chemical industry, pipe for conveying coal slurry or
coal tar,
substrate materials for catalytic converters, exhaust pipes for automotive
engines, porous
filters, etc.
According to one aspect of the invention, in the case where the composite is
used
for heating elements of electrical smoking articles, the geometry of the
composite can be
varied to optimize heater resistance according to the formula: R = p (L/W x T)
wherein R
= resistance of the heater, p = resistivity of the heater material, L = length
of heater, W
= width of heater and T = thickness of heater. The resistivity of the heater
material can
be varied by adjusting the iron aluminide alloy composition and/or the amount
and/or type
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of filler material in the composite. The composite can optionally include
filler.such as
ceramic particles to enhance creep resistance and/or thermal conductivity. The
composite
may also incorporate particles of electrically insulating material for
purposes of making the
composite creep resistant at high temperature and also enhancing thenmal
conductivity
and/or reducing the thermal coefficient of expansion of the composite. The
electrically
insulating/conductive particles/fibers can be added to a powder mixture of Fe,
A1 or iron
aluminide or such particles/fibers can be formed by reaction synthesis of
elemental powders
which react exothermically during manufacture of the composite.
The composite can be made in various ways. For instance, the iron aluminide of
the
composite can be made from a prealloyed powder or by mechanically alloying the
alloy
constituents. The mechanically alloyed powder can be processed by conventional
powder
metallurgical techniques such as by canning and extruding, slip casting,
centrifugal casting,
hot pressing and hot isostatic pressing. Another technique is to use pure
elemental powders
of Fe, A1 and optional alloying elements and mechanically alloying such
ingredients. In
addition to the above, the above mentioned electrically insulating and/or
electrically
conductive particles can be incorporated in the powder mixture to tailor
physical properties
and high temperature creep resistance of the composite.
The composite is preferably made by powder metallurgy techniques. For
instance,
the composite can be produced from a mixture of powder having different
fractions but a
preferred powder mixture comprises particles having a size smaller than minus
100 mesh.
According to one aspect of the invention, the iron aluminide powder can be
produced by
gas atomization in which case the powder may have a spherical morphology.
According to
another aspect of the invention, the iron aluminide powder can be made by
water
atomization in which case the powder may have an irregular morphology. The
iron
aluminide powder produced by water atomization can include an aluminum oxide
coating
on the powder particles and such aluminum oxide can be broken up and
incorporated in the
composite during thermomechanical processing of the powder to form shapes such
as sheet,
bar, etc. The alumina particles are effective in increasing resistivity of the
iron aluminum
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alloy and while the alumina is effective in increasing strength and creep
resistance, the
ductility of the alloy is reduced.
When molybdenum is used as one of the alloying constituents of the iron
aluminide
it can be added in an effective range from more than incidental impurities up
to about 5.0%
with the effective amount being sufficient to promote solid solution hardening
of the iron
aluminide alloy and resistance to creep of the alloy when exposed to high
temperatures.
The concentration of the molybdenum can range from 0.25 to 4.25 % and in one
preferred
embodiment is in the range of about 0.3 to 0.5 % . Molybdenum additions
greater than
about 2.0 % detract from the room-temperature ductility due to the relatively
large extent of
solid solution hardening caused by the presence of molybdenum in such
concentrations.
Titanium can be added to the iron aluminide in an amount effective to improve
creep
strength of the iron aluminide alloy and can be present in amounts up to 3 % .
When
present, the concentration of titanium is preferably in the range of s 2.0 % .
When carbon and the carbide former are used in the iron aluminide alloy, the
carbon
is present in an effective amount ranging from more than incidental impurities
up to about
0.75 % and the carbide former is present in an effective amount ranging from
more than
incidental impurities up to about 1.0% or more. The carbon concentration is
preferably in
the range of about 0.03 % to about 0.3 % . The effective amount of the carbon
and the
carbide former are each sufficient to together provide for the formation of
sufficient
carbides to control grain growth in the iron aluminide alloy during exposure
thereof to
increasing temperatures. The carbides may also provide some precipitation
strengthening
in the iron aluminide alloy. The concentration of the carbon and the carbide
former in the
iron aluminide alloy can be such that the carbide addition provides a
stoichiometric or near
stoichiometric ratio of carbon to carbide former so that essentially no excess
carbon will
remain in the finished alloy.
Zirconium can be incorporated in the iron aluminide alloy to improve high
temperature oxidation resistance. If carbon is present, an excess of a carbide
former such
as zirconium in the iron aluminide alloy is beneficial in as much as it will
help form a
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spallation-resistant oxide during high temperature thermal cycling in air.
Zirconium is
more effective than Hf since Zr can form oxide stringers perpendicular to the
exposed
surface of the iron aluminide alloy which pins the surface oxide whereas Hf
forms oxide
stringers which are parallel to the surface.
The carbide formers include such carbide-forming elements as tungsten,
titanium,
zirconium, niobium, tantalum and hafnium and combinations thereof. The carbide
former
is preferably in a concentration sufficient for forming carbides with the
carbon present
within the iron aluminide alloy. The concentrations for tungsten, niobium,
tantalum,
titanium, zirconium and hafnium when used as carbide formers can be present in
amounts
up to 3 wt % each.
In addition to the aforementioned alloy elements the use of an effective
amount of a
rare earth element such as about 0.05-0.25 % cerium or yttrium in the iron
aluminide alloy
composition is beneficial since it has been found that such elements improve
oxidation
resistance of the alloy.
The oxide filler can be in the form of particles such as powder, fibers, etc.
For
example, the composite can include up to 40 wt % of oxide particles such as
Y203, A1203,
rare earth oxide, beryllia or combinations thereof. The oxide particles can be
added to a
melt or powder mixture of Fe, Al and other alloying elements. Alternatively,
the oxide can
be created in situ by water atomizing a melt of an aluminum-containing iron-
based alloy
whereby a coating of alumina or yttria on iron-aluminum powder is obtained.
During
processing of the powder, the oxides break up and are arranged as stringers in
the final
product. Incorporation of the oxide particles in the iron aluminide alloy is
effective in
increasing the resistivity of the alloy. For instance, by incorporating about
0.5-0.6 wt
oxygen in the alloy, the resistivity can be raised from around 100 ~u Sa ~ cm
to about
160 ~, ~2 ~ cm.
The additive for promoting bonding between the iron aluminide and oxide filler
can
comprise any element or compound which improves wetting of the iron aluminide,
i.e.
lowers surface tension and/or contact angle. For instance, the additive can
comprise a
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carbide which is stable in molten iron aluminide. A preferred additive is a
refractory
carbide such as TiC, HfC and/or ZrC. During liquid phase sintering wherein the
iron
aluminide is partially or fully melted, the refractory carbide remains solid
and promotes
bonding of the oxide filler to the molten iron aluminide matrix.
In order to improve thermal conductivity and/or resistivity of the iron
aluminide
alloy, particles of electrically conductive and/or electrically insulating
metal compounds can
be incorporated in the alloy. Such metal compounds include oxides, nitrides,
silicides,
borides and carbides of elements selected from groups IVb, Vb and VTb of the
periodic
table. The carbides can include carbides of Zr, Ta, Ti, Si, B, etc., the
borides can include
borides of Zr, Ta, Ti, Mo, etc., the silicides can include silicides of Mg,
Ca, Ti, V, Cr,
Mn, Zr, Nb, Mo, Ta, W, etc., the nitrides can include nitrides of Al, Si, Ti,
Zr, etc., and
the oxides can include oxides of Y, Al, Si, Ti, Zr, etc.
Additional elements which can be added to the iron aluminide alloy include Si,
Ni
and B. For instance, small amounts of Si up to 2.0% can improve low and high
temperature strength but room temperature and high temperature ductility of
the alloy may
be adversely affected with additions of Si above 0.25 wt % . The addition of
up to 30 wt
Ni can improve strength of the iron aluminide alloy via second phase
strengthening but Ni
adds to the cost of the alloy and can reduce room and high temperature
ductility thus
leading to fabrication difficulties particularly at high temperatures. Small
amounts of B can
improve ductility of the alloy and B can be used in combination with Ti and/or
Zr to
provide titanium and/or zirconium boride precipitates for grain refinement.
The invention is now described with reference to the following examples which
provide exemplary details of how to make low-cost FeAI-based composites.
FeAI-based composites reinforced with insulating oxide filler can be prepared
by a
variety of techniques including conventional casting and powder metallurgical
processes.
However, because oxides are oxidation resistant and have poor electrical
conductivity, their
presence in iron aluminide composites can be used to increase the electrical
resistivity of
the composite which is an advantage in resistance heater applications. In the
following
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examples, fabrication of iron aluminide-oxide composites was carried out using
powder
metallurgical techniques.
In the following examples, iron aluminide composites were prepared using A1z03
and/or Zr02 as the oxide particulates. Zr02, in particular, exhibits a high
coefficient of
thermal expansion, and has therefore a relatively small thermal mismatch with
the iron
aluminide matrix. The composites were made by hot-pressing as well as low-cost
techniques such as liquid phase sintering.
In order to fabricate FeAI/oxide composites, the following three issues were
addressed: (a) the thermodynamic compatibility between the oxides and the iron
aluminide
matrix, (b) the degree by which oxide particles are wetted by liquid iron
aluminides, and (c)
the extent to which the wetting behavior can be modified by alloying additions
to the iron
aluminide. It has been found that AIz03 is thermodynamically compatible with
FeAI
whereas Zr02 is not. Further, while liquid iron aluminide does not wet AIz03
adequately,
additions of TiC to FeAI/A1203 powder mixtures improves wetting and
fabricability. Hot
forging of FeAI-15 vol. % TiC-IS vol. % A1z03 composites improved the room
temperature
flexure strength more than threefold. For instance, room temperature flexure
strengths
exceeding 1000 MPa can be obtained with the hot-forged composites. Such
improvement
in mechanical properties mat be due to reduction in residual porosity in the
composites. In
addition, a dramatic improvement of the liquid phase sintering behavior can be
obtained by
incorporating an additive (e.g., TiC) which promotes wetting of the oxide
filler.
Experiments were carried out with FeAI/A1z03 and FeAI/Zr02 specimens prepared
by mixing Fe-40 at. % Al, A1z03 or Y203-stabilized Zr02 powders and liquid
phase
sintering them in vacuum at 1450°C or 1500°C. In the following
discussion, "FeAI" is
intended to denote Fe-40 at. % AI. As a result of X-ray diffraction data, it
has been
determined that the FeAI/A1203 composite included alpha-A1203 and FeAI and the
FeAI/Zr02 composite included cubic stabilized Zr02 as well as FeAI. However,
there was
also evidence of substantial amounts of alpha alumina suggesting a
displacement reaction of
the type: 3 Zr02 + 24 FeAI --~ 2 Fe3A1 + 3 Fe6Al~Zr + 2 A1z03, where Fe6A16Zr
is a
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14
ternary intermetallic phase. Consistent with the proposed reaction,electron
dispersive
spectroscopy (EDS) in a scanning electron microscope (SEM) verified the
presence of
FeAI, FeAIZr intermetallic and A1z03.
A hot-pressed FeAI/Zr02 specimen including 10 % A1203 and 10 % Zr02 was tested
to determine flexure strength. Optical microscopy of the sample revealed that
a reaction
occurred in the material and chipped edges of flexure bars ground from the
material
indicated that the material was brittle in nature. The flexure bars fractured
in a brittle
manner indicating that the iron aluminide had reacted to form more brittle
phases. The
material exhibited a room temperature flexure strength of 215 t 29 MPa. As a
result of
the tests it was determined that Zr02 is not thermodynamically stable in
contact with FeAI.
In the following experiments, prealloyed iron aluminide powders were mixed
with
oxide powders. The powder mixtures were then poured into alumina crucibles
which were
covered with an alumina lid. In most cases, the crucibles had an inner
diameter and inner
height of 38 mm and 8 mm, respectively. Although the powder mixtures were not
cold
pressed prior to sintering, cold pressing prior to sintering is expected to
improve the
fabricability significantly. The filled crucibles were usually pumped
overnight to an
indicated vacuum better than 10-5 Torr. Subsequently, the specimen was ramped
to 1450
or 1500°C over a period of 2 h, held for 0.2 to 0.3 h at that
temperature, followed by
furnace cooling. At 1450 or 1500°C, the iron aluminide melted and
liquid phase sintering
occurred. Attempts were also made to infiltrate oxide powders with liquid iron
aluminide
alloys. In a number of cases, elemental Ti or C was added to the binary iron
aluminides to
enhance the wetting. The best coupons were obtained when a fraction of the
oxide powders
was replaced by TiC powders. The metal alloy and oxide powders employed in the
examples are summarized in Table 1. Table 2 summarizes data obtained from the
specimens. Tables 1 and 2 will be used to discuss the various processing
experiments
carried out.
Figures 1 and 2 show powder X-ray diffraction patterns for specimens A003
(FeAl/A1203) and A004 (FeAI/Zr02). Consistent with thermodynamic stability,
the
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WO 99/39016 PCT/US99/02211
1S
diffraction pattern for the FeAI/A1203 composite indicates mostly a-A1203 and
FeAI. Two
small peaks at 21 and 30° could not be identified. The diffraction
pattern for the
FeAI/Zr02 composite indicates cubic stabilized Zr02 as well as FeAI. However,
there is
evidence for substantial amounts of a-alumina suggesting a displacement
reaction of the
S type:
3 Zr02 + 24 FeAI < --- > 2 Fe,3Al + 3 Fe6A16Zr + 2 A1203,
where Fe6A16Zr is a ternary intermetallic phase. The X-ray result is
substantiated by Figure
3, in which electron dispersive spectroscopy (EDS) in a scanning electron
microscope
(SEM) verifies the presence of FeAI, FeAIZr intermetallic, and A1z03. Clearly,
Zr02 is not
thermodynamically stable in contact with liquid FeAI. Once this was found out,
processing
with Zr02 was discontinued.
In iron aluminide composites containing carbides and borides, wetting by
liquid iron
aluminides is so effective that porous preforms made from these ceramics are
readily
infiltrated. The applicability of this approach to oxides was investigated
(Specimens AOOS,
1S A006, A011, A012). Iron aluminide powder was placed on a bed of A1z03 or
Zr02,
followed by heating to 1450°C in vacuum in order to melt the iron
aluminide. As expected
from the literature on the wetting of oxides by liquid metals, infiltration
did not occur. A
possible solution to this might be addition of a reactive element such as Ti.
However,
infiltration did also not occur when Ti was added to the iron aluminide powder
(AOOS,
A006). Additions of TiC particulates are expected to improve infiltration
behavior during
liquid phase sintering of FeAI/TiC/A1z03 mixtures.
Experiments were carried out with alumina powder A002, which had a particle
size
less than 38 ~.m. Liquid phase sintering of iron aluminides with alumina
resulted usually in
porous coupons and large amounts of exuded FeAI, which was expelled because of
its poor
2S wetting. This is illustrated by Figure 4. When the volume fraction was on
the order of 30
wt (specimens A020 and A041) the coupons were very fragile. When the content
was
lowered to values on the order of 20 wt (A014), the coupons tended to be
stronger. An
iron aluminide powder A040 gave poor results (specimen A044) apparently
because the
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16
powder had a larger particle size than the < 45 ~m powder used for other
samples (A032).
This larger size may have contributed to the poor sinterability. Additions of
Ti or C
(A007, A016, A018) did not cause noticeable improvements. These results are
consistent
with the infiltration experiments. However, as shown below, additions of TiC
improved
the fabricability dramatically.
Partial replacement of A1203 by TiC improved the fabricability substantially.
In
coupons A021, A022, and A023 the TiC/A1z03 ratio was systematically increased.
Once
the TiC content was increased to sufficiently high levels (Z 18 wt%), the
specimens
appeared dense and exhibited no or only few surface cracks. Figure 5 shows a
successfully
processed coupon containing TiC and A1z03. The raised patches on this coupon
appear to
be exuded iron aluminide. However, as compared to Figure 4 the wetting is
dramatically
improved. The microstructure of an FeAI/TiC/A1z03 coupon is depicted in Figure
6.
Although there is still some porosity, many A1z03 particles, such as that in
the center of
Figure 6, are fully surrounded by FeAI.
Surprisingly, additions of Ti were detrimental to the fabricability (A025,
A026,
A027). However, small amounts of C (0.3 wt% , specimens A028 and A030) did not
degrade the fabricability. Thus, optimized additions of C have the potential
to improve the
fabricability.
In summary, A1203 was found to be a suitable reinforcement in iron aluminide
cermets. Zr02, on the other hand, was unstable in contact with liquid FeAI,
and brittle Fe-
Al-Zr intermetallics formed instead. As expected, A1z03 was poorly wetted by
liquid iron
aluminide. Surprisingly, additions of either Ti or C to the iron aluminide did
not improve
the wetting of the A1203. However, the combined addition of Ti and C, in the
form of TiC
particulates, improved the wettability dramatically and resulted in much
denser coupons.
Various changes and modifications can be made to the process according to the
invention. For instance, cold pressing of the powder mixtures can be used to
reduce the
porosity of the final product. Optimization can be achieved by conducting
quantitative
density and porosity measurements to determine the concentrations of alloying
additions
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17
such as carbon. Further, it is expected that niobium additions will have a
beneficial
influence on the wetting and bonding of A1z03 by iron aluminides. Instead
af~prealloyed
FeAI powders, elemental Fe and A1 powders can also be used as well. In fact,
the
exothermic reaction between the elemental Fe and A1 may be beneficial. Also,
elemental
powders are softer than prealloyed FeAI powder (which is strongly hardened by
frozen-in
thermal vacancies) and will therefore result in higher green density. High
green densities
will lead to higher final densities associated with improved strength and
oxidation
resistance.
Table 1: Raw materials used in this research
I~~signatlc~tr Cmriposlttah ~~~
A001 Zr02 Powder Zr02-YZO3 -325 mesh
(93-7) (s45~.m)
A002 A 1203 PowderA 1203 -400 mesh
( s 38~.m)
A019 Graphite C /em range
Powder
A024 TiC Powder TiC 1.9/cm
A032 FeAI Powder Fe-40 at. % Al -325 mesh
(s45~,m)
A033 Ti Powder Ti, 99.5 % pure -200 mesh
(s75um)
A040 PeAI Powder
A045 TiC Powder TiC, 99 % metal 2.5-4 ~,m
basis
Table 2: Summary of nrn~.PCeina PvnPr;."o"t
,. ' . ..--~."
~ ~cin~aen..C r.F , .
F o sttx~n f~wdez;~ Purpose
~ ; used Pxndut
' s
g
~1 ,; (s: ;
uuoa, I ' se 'fable
fer . l )
A00 Fe40A1-22 wt% A1203 3 powdery-ray estimated
A032, A002
diffraction, porosity
metallography 20 %
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WO 99/39016 PCT/US99/02211
18
Cln7.~i1: ~ Iti~l
P~ . ~ ~'owders . Pulse Fttidi~gs
tt~d
N~ . ,..' ses ~'~I~I~
1) :: .
~:
A004 Fe40A1-30 wtl Zr02 powder x-ray ,
estimated
diffraction, porosity
metallography 20%
A005 Fe40A1-11 wt~ Ti/A1203A032, A033,Infiltration no
A002 attempt infiltration
found
A006 Fe40A1-11 wt~ Ti/A1z03A032, A033,Infiltration no
A001 attempt infiltration
found
A007 (Fe40A1-11 wt9o Ti)/20A032, A033,Liquid phase Porous
wt% A1z03 A002 sintering withpellet,
Ti
addition exuded
FeAI, Pellet
electrically
conductive
A008 (Fe40A1-11 wt% Ti)/28A032, A033.,Liquid phase Porous
wt~ Zr02 A001 sintering withpellet,
Ti
addition exuded
FeAI, Pellet
electrically
conductive
A009 Fe40A 1 / 15 wt % A032, A045,Liquid phase Dense
TiC/ 12
wt~ A1203 A002 sintering withappearance,
TiC
addition exuded
patches
on
top (see
macrograph)
A010 Fe40A 1 / 14 wt ~ A032, A045,Liquid phase Dense
TiC/ 14
wt~ Zr02 A001 sintering withappearance,
TiC
additions large surface
cracks
A011 Fe40A1/Zr02 A032, A001Infiltration No
attempt infiltration
found
A012 Fe40A 1 /A1203 A032, A002Infiltration No
attempt infiltration
found
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WO 99/39016 PCT/US99/02211
19
P de 'e F"tnudii~
en ~~mpoat~Qi~ ~ s. ..
: ~ ~s :. g
, :, Q
~
:. . . . ..... . ..
~um~r .. ~s~ ' ',
.. e'~'ltb
.
A013 Fe40A1/30 wtb Zr02 A032, A001Liquid phase Porous,
sintering fragile
pellet,
exuded
FeAI.
A014 Fe40A1/22 wt~ A1203 A032, A002Liquid phase Porous
sintering pellet,
exuded
FeAI
A015 (Fe40A1-11 wt~ Ti)/30A032, A033,Liquid phase Porous
wt~ Zr02 A001 sintering pellet,
exuded
FeA 1
A016 (Fe40A1-11 wt~ Ti)/22A032, A033,Liquid phase Porous
wt~ A1z03 A002 sintering pellet,
exuded
FeAl
A017 (Fe40A1-2.9 wtgb A032, A019,Liquid phase Porous
C)/30
wt ~ ZrOz A001 sintering pellet,
black
and silver
areas,
no
exuded
FeA 1
A018 (Fe40A1-2.9 wt% C)/22A032, A019,Liquid phase Porous
wt% A1203 A002 sintering pellet,
exuded
FeA 1
A020 FeAl/33 wt9~ A1z03 A032, A002Liquid phase Porous
sintering pellet,
fragile,
exuded
FeA 1
A021 Fe40A1/9 wt~O TiC/22A032, A024,Liquid phase Dense
wt~ A1203 A002 sintering appearance,
but many
surface
cracks.
Some exuded
FeA 1
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WO 99/39016 PCT/US99/02211
. . . . .. ,
~pecxmen:Cpzr~pos~tic~n '! P .. ' F
owd~rs ~'urp~se ui~ygs
tts~d: .:
.
.
~tu~nbex, ~sc~ ~aia~~ . .
> 1) .::
A022 Fe40AI/18 wt% TiC/I4A032, A024,Liquid phase Dense
wt% A1203 A002 sintering appearance,
a few
surface
cracks,
no
exuded FeA
1
A022B Fe40A1/18 wt% TiC/14A032, A024,Liquid phase Dense
wt% A1203 A002 sintering appearance,
a few
surface
cracks,
no
exuded FeA
1
A023 Fe40A1/27 wt% TiC/7 A032, A024,Liquid phase Dense
wt % A 1203 A002 sintering appearance,
a few
surface
cracks,
no
exuded FeA
1
A025 (Fe40A1-5 wt% Ti)/18A032, A033,Liquid phase Dense
wt% TiC/14 wt% A1203A024, A002sintering appearance,
several
surface
cracks,
exuded FeA
1
A026 (Fe40A1-1.4 wt% C)/18A032, A019,Liquid phase Dense
wt% TiC/14 wt% A1203A024, A002sintering appearance,
~Y
surface
cracks,
exuded FeA
1
A027 (Fe40A1-1.1 wt% Ti)/18A032, A033,Liquid phase Dense
wt% TiC/i4 wt% A1203A024, A002sintering appearance,
~Y
surface
cracks,
exuded FeAI
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21
:.
c' en ~vm ' st~vn ~owders:~s~d..::: s~ Ftndart
~p~ un pc~ ~ s
8
Nix (~~~ Tab~~e.
1) .
A028 (Fe40A1-0.3 wt~C)/18A032, A019,Liquid phase Dense
wt~ TiC/14 wt~ A120,A024, A002sintering appearance,
a few
surface
cracks,
exuded FeA
1
A029 Fe40A1/18 wt9b TiC/14A032, A024,Liquid phase Dense
wt6 A1203 A002 sintering appearance
A030 Fe40A1-0.3 wt%C/18 A032, A019,Liquid phase Dense
wt~
TiC/14 wt~ A1203 A024, A002sintering appearance
A031 Fe40A1/18 wt~6 TiC/14A032, A045,Liquid phase Dense
wt~ A1203 A002 sintering appearance,
a few
surface
cracks,
exuded FeA
1
A041 Fe40A1/30 wtb A12O3 A040, A024,Liquid phase Fragile
A002 sintering coupon,
exuded FeA
1
A042A Fe40A 1 / 18 wt6 A040, A024,Liquid phase Porous
TiC/ 14
wt% A1203 A002 sintering coupon with
many
surface
cracks,
exuded FeAI
A043A Fe40A1/24 wt% A1203 A032, A002Liquid phase Porous
sintering coupon,
a
few surface
cracks,
exuded FeA
1
A043B Fe40AI/24 wt% A1203 A032, A004Liquid phase Porous
sintering coupon,
very
fragile,
exuded FeAI
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WO 99/39016 PCT/US99/02211
22
x x ::
Sp~Cune~Co~riposa~on Powd~xs Pu .se Fundmgs
u~d
Number ~'~"~~le . : ' .....
1)
A044 FeAI/24 wt% A1z03 A040, A002 Liquid phase Porous
sintering coupon,
very
fragile,
exuded
FeAI
As a result of the above experiments it was determined that FeAI did not wet
A1z03
sufficiently well to fabricate FeAI/A1203 composites by liquid phase
sintering. In order to
improve the sintering behavior, some of the A1z03 powder was replaced by TiC
powder.
For example, specimen A009 was fabricated from Fe-40 at. % Al powder (-325
mesh or
<45,um), TiC powder (2.5-4um), and A1203 powder (s38 um), the sample having a
nominal composition of FeAI-16.5 vol. % TiC- 16.5 vol. % A1z03. Compositions
and
preparation techniques for specimen A009 and additional specimens are set
forth in Table
3. The same size powders used for specimen A009 were also used for A046.
Specimen
A062C was made from powders having the following sizes: 1-S,um Fe, 10 ,urn Al,
2.5-4
,um TiC and s38 ~cm A1z03. The liquid phase sintering was carried out as
follows: 0.3 h in
vacuum for specimen A009, 0.2 h in vacuum for specimen A046, 0.2 h in vacuum
for
specimen A047, 0.2 h in vacuum for specimen A050, and 0.2 h in vacuum for
specimen
A062C.
Table 3.
Specimen Number Composition Powders Used Processing
For
Iron Aluminide Technique
A009 FeAI-16. Svol Prealloyed FeAI,Liquid phase
% TiC-
l6.Svo1 % A1203TiC and A1203 Sintering at
1450
A046 FeAI-l6.Svo1 Prealloyed FeAI,Liquid phase
% TiC-
16. Svol % A1203TiC and A1203 Sintering at
1500 C
A047 FeAI-16. Svol Fe and A1 Liquid phase
% TiC-
l6.Svo1%A1203 Sintering at
1500C
A050 FeAI-9wt%Nb- Prealloyed FeAI Liquid phase
l6.Svo1%TiC- and Nb Sintering at
1500C
l6.Svo1 % A1203
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WO 99/39016 PCT/US99/02211
23
A055 FeAI-10%A1203 Prealloyed FeAIHot Pressed
Zr02
A062C FeAI-lSvol % Fe, Al, TiC Liquid phase
TiC- and
15vo1 % A1203 A1203 Sintering at
1500 C
and hot forging
at
1000C from 20
to
8 mm
Table 4.
Specimen Composition at. Flexure Strength
% MPa
A046E-1 FeAI-15vo1 % 304
TiC-
lSvol % A1203
AOSOA-1 FeAI-9wt % Nb- 189
l6.Svol % TiC-
16 . S vol %
A1203
AOSOA-2 FeAI-9wt % Nb- 185
l6.Svo1 % TiC-
16. Svol % A1203
AOSS# 1 FeAI-l Ovol % 212
A1203 -
l Ovol % Zr02
A055#1 FeAI-lOvol%A12032I7
-
l Ovol % Zr02
A055#1 FeAI-lOvol % 249
A1203 -
l Ovol % Zr02
A055#2 FeAI-l Ovol % 169
A1203 -
l Ovol % Zr02
A055#2 FeAI-lOvol % 226
A1203 -
lOvol % Zr02
A062C#i FeAI-15vo1%TiC- 996
15vo1 % A1203
A062C#1 FeAI-15vo1 % 1081
TiC-
1 Svol % A1203
A062C#1 FeAI-15vo1 % 1160
TiC-
15vo1 % A1203
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24
A062C#2 FeAI-15vo1%TiC- 1099
l5vol %A1203
A062C#2 FeAI-15vo1% TiC-1202
15V01 % A12O3
A062C#2 FeAI-l5vol%TiC- 1173
15vo1%A1203
A062C#3 FeAI-15vo1%TiC- 1056
15vo1 % AI203
A062C#3 FeAI-15vo1 % 981
TiC-
15vo1 % A1203
Specimens with the nominal composition FeAI-16.5vo1 % TiC-16.5vo1 % A1z03 were
also fabricated by cold-pressing and subsequent sintering for 12 minutes at
1500°C in
vacuum. Similar results were achieved using prealloyed FeAI (specimen A046) or
elemental Fe and A1 powders (specimen A047). However, the composite fabricated
from
elemental powders may have a slightly lower porosity level. In specimen A050,
elemental
Nb was added to the composite with the expectation that Nb would bond well to
the A~03
and improve fracture toughness.
In the experiments, it was found that fully dense material was not produced
during
liquid phase sintering even when TiC was added to the composite material. As
such,
secondary processing was utilized to remove the pores. Specimen A062C was made
by
mixing 60 g of Fe, AI, TiC and Aiz03 and liquid phase sintering the mixture in
an A1z03
crucible to provide a FeAI-lSTiC-15A1z03 (vol. %) composite. The sintered
cylinder was
hot forged at 1000°C from a height of 20 mm to approximately 8mm. The
hot forged
coupon is shown in Figure 7 wherein edge cracking can be seen around the
periphery of
the coupon and the interior of the coupon is sound.
Figure 8 is an optical micrograph of specimen A046 fabricated with prealloyed
Fe40Al powder. The bright TiC particles, dark A1z03 particles, with black
pores
surrounded by gray iron aluminide matrix are clearly visible. Processing with
elemental
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Fe and A1 powders, instead of prealloyed FeAI powder, gave similar results
except that the
porosity levels may have been lower. Figure 9 shows the microstructure of a
hot forged
coupon (A062C) wherein there is an absence of porosity.
Specimens for room temperature flexure tests were prepared by grinding samples
5 having a cross section of approximately 3x4 mm. The flexure tests were
carried out with a
span of 20 mm and a cross head speed of 10 ,um/s. The fracture stress q. was
calculated
from the linear-elastic equation: of = 1.5 L~P/(wtz), where L is the span, P
is the load at
fracture, w is the specimen width and t is the specimen thickness.
The strength of liquid phase sintered FeAI-16.STiC-16.SA1203 (vol. % )
exceeded
10 300 MPa (specimen A046E-1). Fracture occurred not catastrophically, but in
a gradual
manner by controlled crack propagation. The reason for the gradual fracture is
thought to
be the porosity of the material which did not permit sufficient storage of
elastic energy to
result in catastrophic fracture. The Nb-alloyed material A050 exhibited
fracture in a
gradual manner and had a much lower strength of 187 MPa which is presumably
due to its
15 higher porosity. Although the Nb may have strengthened the interfacial
AIz03/FeAI
bonding, this could not be verified because of the negative effect of the high
porosity
levels.
Hot forging resulted in a pronounced strength increase. Figure 10 shows three
stress displacement curves for bend bars machined from coupon A062C (FeAl-
ISTiC-
20 15A1203, vol. % ). The curves demonstrate not only a high strength, but
also a small
amount of ductility. The beneficial effect of the hot forging is attributed to
the removal of
porosity. Some specimens were annealed for 1 day at 500°C in order to
remove thermal
vacancies which were presumably frozen in during the hot forging. The removal
of excess
vacancies in iron aluminides results in a reduction of the high yield strength
and an
25 increase in ductility. Although the anneal was expected to reduce the flaw
sensitivity and
increase fracture strength, it was found that the anneal did not affect the
fracture strength
significantly.
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26
The room temperature fracture toughness of the hot forged FeAI-lSTiC-l5A1z03
composite was determined from the controlled fracture of chevron-notched
specimens.
Figure 11 shows a measured load-displacement curve. The fracture toughness was
evaluated from the equation: KQ = [(W/A)E']"2 where W is the absorbed energy
(which
corresponds to the area under the load-displacement curve), A is the area
traversed by the
crack, and E' is the plane strain Young's modulus, namely E/(1-uz). A value of
0.25 was
assumed for v. The Young's modulus E was estimated from the following
equation: E =
[(cEPEm+Em2)(1 +c)2-Em2+EPEm]/[(cEP+Em)(1 +c)2] where c = (1/Vp)"3-1. VP is
the
volume fraction of the ceramic particles, Ep and Em are the moduli of the
ceramic phases
(estimated to be 410 GPa) and the matrix (180 GPa). Using the above equations,
the
Young's modulus for FeAI-lSTiC-l5AIz03 (vol. % ) is estimated to be 228 Gpa.
The fracture toughness of two specimens evaluated in this manner are listed in
Table 5. Considering the relatively low fracture toughness of monolithic iron
aluminides
(30-50 MPa m"~), the composites exhibited satisfactory fracture toughnesses.
Table 5. Fracture Toughness of Hot Forged FeAI-lSTiC-l5A1z03 (vol. % )
Sample W H W A G, E KQ
A062C# 6.59 1.66 2.67 2.216 2973.7 228.0 26.9
A062C 7 1.7 2.8 2.38 2941.2 228.0 26.7
From the foregoing discussion it can be appreciated that A1~03 is not wetted
well
enough by liquid FeAI to allow the processing of composites by liquid phase
sintering. In
contrast to A1203, Zr02 is thermodynamically not stable in contact with iron
aluminides.
Since brittle intermetallic phases form during the reaction between ZrO~ and
FeAI, Zr02 is
less desirable as a filler in FeAI/ceramic composites. On the other hand, TiC
promotes
wetting of A1z03 by FeAI. Moreover, instead of prealloyed FeAI, elemental Fe
and Al
powders may be used for liquid phase sintering of FeAI/TiC/A1z03 composites.
Additions
of refractory metals such as Nb may improve the properties of the composites
provide
porosity can be reduced to acceptable levels. Room temperature flexure
strengths of
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27
approximately 300 MPa can be achieved for liquid phase sintered iron aluminide
composites containing TiC and A1203. Hot forging of liquid phase sintered FeAI-
TiC-
A1203 composites can increase the room temperature flexure strengths to
approximately
1000 MPa as well as provide fracture toughness on the order of 27 MPa m'~.
The foregoing has described the principles, preferred embodiments and modes of
operation of the present invention. However, the invention should not be
construed as
being limited to the particular embodiments discussed. Thus, the above-
described
embodiments should be regarded as illustrative rather than restrictive, and it
should be
appreciated that variations may be made in those embodiments by workers
skilled in the an
without departing from the scope of the present invention as defined by the
following
claims.
