Background of the Invention 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 Fe3Al, FeAl, FeAl2, FeAl3, and Fe2Al5. Fe3AI 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, < 7 Cr, < Mo, Mo, : 05 C, C, < 0.5 Zr and < 0. Y, preferably preferably 5-5.5-5. 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

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 Fe3Al Alloys", pp. 219- 231, discloses a powder metallurgical process for preparing Fe3Al containing 2 and 5 % Cr by using an inert gas atomizer. This publication explains that Fe, AI alloys have a DO3 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 Fe3Al-Based Iron-Aluminide Alloys,"pp. 901-906, discloses a process of preparing 2 and 5 % Cr containing Fe3Al-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 rolled 50% at 650°C to 0.76 mm sheet.

A paper by V. K. Sikka et al., entitled"Powder Production, Processing, and Properties of Fe3Al", pp. 1-11, presented at the 1990 Powder Metallurgy Conference Exhibition in Pittsburgh, PA, discloses a process of preparing Fe ; Al powder by melting constituent metals under a protective atmosphere, passing the metal through a metering

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, um. 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% Al and 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 FeAl40 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 FeAl 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 FeAl compounds in oxidizing and sulfidizing conditions is due to the high Al content and the stability of the B2 ordered structure.

A publication by D. Pocci et al., entitled"Production and Properties of CSM FeAl 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 Fe4oAl 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 FeAl alloys were prepared

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 find 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 FeAl compositions.

A publication by J. Baker entitled"Flow and Fracture of FeAl", 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 treatments strongly affect the mechanical properties of FeAl 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 FeAl 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 FeAl", pp. 231-239, presented at the 1994 TMS Conference discloses the effect of ternary alloying additions on FeAl alloys. This article states that the B2 structured compound FeAl 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

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.

Summary of the Invention 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 beryllia. 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, # 2.0% Si, < 30 % Ni, < 0. 5 % Y, < 15 % Nb, < 1% Ta, < 3% W, < 10% Cr, < 2. 0% Mo, < 1 % Zr, # 1 % C and # 0.1% B. The oxide filler preferably comprises alumina which can be present in any desired amount such as # 40%. The additive preferably comprises < 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, AIN, etc. Preferred iron aluminide alloys include 20.0-31.0% Al, 0.05-0.15% Zr, # 3% W, # 0.1% 0. and 2% 14.0-20. 0% 0-20.0.3-1.0.3-1.

0.05-1.0% Zr, # 3% W and # 0.2% C, 0. 1% C, and # 2.0% Ti 1% and 20.0-31. 0% Al, 0.3-0.5% Mo, 0.05-0.3% Zr, # 0.2% C, # 2% W, # 0.1% B and # 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 80-400y Q cm, preferably 90-200, u Q cm. The composite preferably heats

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 Description of the Drawings Figure 1 shows an X-ray diffraction pattern for an FeAl/AtO3 composite in accordance with the invention; Figure 2 shows an X-ray diffraction pattern for an FeAI/ZrQ composite in accordance with the invention;

Figure 3 shows a scanning electron microscope image of an FeAI/ZrQ composite in accordance with the invention; Figure 4 shows exudation of FeAl during liquid phase sintering of an FeAl/AI203 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 Al, 0g of iron aluminide; Figure 6 shows a scanning electron microscope image of a polished section of an FeAl/TiC/Al203 composite in accordance with the invention; Figure 7 shows a hot forged coupon of Fe-lSTiC-15A103 (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 FeAl- 16.5TiC-16. SA1203 (vol. %) in accordance with the invention; Figure 9 is an optical micrograph of a hot forged composite of FeAl-15TiC-15Al203 (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 FeAl-15TiC-15A03 (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 FeAl-15TiC-15Al203 (vol. %) in accordance with the invention.

Detailed Description of the Preferred Embodiments 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

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 ZrB2.

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

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 thermal 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, Al 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, Al 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

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 < 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

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 A 0-cm to about 160 Q'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

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 VIb 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 FeAl-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

examples, fabrication of iron aluminide-oxide composites was carried out using powder metallurgical techniques.

In the following examples, iron aluminide composites were prepared using Auto3 and/or ZrO2 as the oxide particulates. ZrO2, 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 FeAl/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 Alto3 is thermodynamically compatible with FeAl whereas ZrO2 is not. Further, while liquid iron aluminide does not wet AtO3 adequately, additions of TiC to FeAl/Al203 powder mixtures improves wetting and fabricability. Hot forging of FeAl-15 vol. % TiC-15 vol. % A1203 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 FeAl/AtO3 and FeAI/ZrO2 specimens prepared by mixing Fe-40 at. % Al, Al203 or Y203-stabilized ZrO2 powders and liquid phase sintering them in vacuum at 1450°C or 1500°C. In the following discussion,"FeAl"is intended to denote Fe40 at. % Al. As a result of X-ray diffraction data, it has been determined that the FeAl/AtO3 composite included alpha-Al203 and FeAl and the FeAl/ZrO2 composite included cubic stabilized ZrO2 as well as FeAl. However, there was also evidence of substantial amounts of alpha alumina suggesting a displacement reaction of the type: 3 ZrO2 + 24 FeAl-2 Fe3Al + 3 Fe6Al6Zr + 2 Al203, where Fe6Al6Zr is a

ternary intermetallic phase. Consistent with the proposed reaction, electron dispersive spectroscopy (EDS) in a scanning electron microscope (SEM) verified the presence of FeAl, FeAlZr intermetallic and Al203.

A hot-pressed FeAl/ZrO2 specimen including 10% A1203and 10% ZrO2was 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 29 MPa. As a result of the tests it was determined that ZrO2 is not thermodynamically stable in contact with FeAl.

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/Al203) and A004 (FeAl/ZrO2). Consistent with thermodynamic stability, the

diffraction pattern for the FeAl/Al203 composite indicates mostly a-Al203 and FeAl. Two small peaks at 21 and 30° could not be identified. The diffraction pattern for the FeAl/ZrO2 composite indicates cubic stabilized ZrO2 as well as FeAl. However, there is evidence for substantial amounts of a-alumina suggesting a displacement reaction of the type: 3 ZrO2 + 24 FeAl <---> 2 Fe ; Al + 3 Fe6Al6Zr + 2 Au203, where Fe6Al6Zr 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 FeAl, FeAlZr intermetallic, and AtO3. Clearly, ZrO2 is not thermodynamically stable in contact with liquid FeAl. Once this was found out, processing with ZrO2 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 A005, A006, A011, A012). Iron aluminide powder was placed on a bed of At03 or ZrO2, 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 (A005, A006). Additions of TiC particulates are expected to improve infiltration behavior during liquid phase sintering of FeAl/TiC/AtO3 mixtures.

Experiments were carried out with alumina powder A002, which had a particle size less than 38, um. Liquid phase sintering of iron aluminides with alumina resulted usually in porous coupons and large amounts of exuded FeAl, which was expelled because of its poor 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

powder had a larger particle size than the < 45, um 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/AtO3 ratio was systematically increased. Once the TiC content was increased to sufficiently high levels (2 18 wt%), the specimens appeared dense and exhibited no or only few surface cracks. Figure 5 shows a successfully processed coupon containing TiC and A1203. 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 FeAl/TiC/AtO3 coupon is depicted in Figure 6.

Although there is still some porosity, many AtO3 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. ZrO2, on the other hand, was unstable in contact with liquid FeAl, and brittle Fe- Al-Zr intermetallics formed instead. As expected, AtO3 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 A12O3. 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

such as carbon. Further, it is expected that niobium additions will have a beneficial influence on the wetting and bonding of A1203 by iron aluminides. Instead of prealloyed FeAl powders, elemental Fe and Al powders can also be used as well. In fact, the exothermic reaction between the elemental Fe and Al may be beneficial. Also, elemental powders are softer than prealloyed FeAl 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 Designatian Compositio Size AOO1ZrO2 Powder ZrO2-Y2O3-325 mesh (93-7) (<45ym) A002 A1203 Powder A1203-400 mesh (38m) A019 Graphite Powder C s4m range A024 TiC Powder TiC 1.9ym A032 FeAl Powder Fe-40 at. % Al-325 mesh (_45, m) A033 Ti Powder Ti, 99.5% pure-200 mesh (<75ym) A040 FeAI Powder A045 TiC Powder TiC, 99% metal basis 2.5-4 zm Table 2: Summary of processing experiment !es ! s !Q : rnp Ttiinber : : (see Table 1) : A003 Fe40A1-22 wt (% o Al203 A032, A002 powder x-ray estimated diffraction, porosity metallography 20%

Specun. en : CoiipositiQZawdxs used'Puiose FriIins s d P d* A004 Fe40A1-30 wt% ZrO2 powder x-ray estimated diffraction, porosity metallography 20 % A005 Fe40A1-11 wt% Ti/Al203 A032, A033, Infiltration no A002 attempt infiltration found A006 Fe40Al-ll wt% Ti/Al203 A032, A033, Infiltration no A001 attempt infiltration found A007 (Fe40A1-11 wt% Ti)/20 A032, A033, Liquid phase Porous wt% Al203 A002 sintering with Ti pellet, addition exuded FeAl, Pellet electrically conductive A008 (Fe40Al-ll wt% Ti)/28 A032, A033., Liquid phase Porous wt% ZrO2 A001 sintering with Ti pellet, addition exuded FeAl, Pellet electrically conductive A009 Fe40A1/15 wt% TiC/12 A032, A045, Liquid phase Dense wt% Al203 A002 sintering with TiC appearance, addition exuded patches on top (see macrograph) A010 Fe40Al/14 wt% TiC/14 A032, A045, Liquid phase Dense wt% ZrO2 A001 sintering with TiC appearance, additions large surface cracks A011 Fe40Al/zrO2 A032, A001 Infiltration No attempt infiltration found A012 Fe40Al/Al203 A032, A002 Infiltration No attempt infiltration found

! sSpseuo on.... P A013 Fe40A1/30 wt% ZrC2 A032, A001 Liquid phase Porous, sintering fragile pellet, exuded FeAI. A014 Fe40A1/22 wt% Au203 A032, A002 Liquid phase Porous sintering pellet, exuded FeAI A015 (Fe40A1-11 wt% Ti)/30 A032, A033, Liquid phase Porous wt% ZrO2 A001 sintering pellet, exuded FeA1 A016 (Fe40A1-11 wt% Ti)/22 A032, A033, Liquid phase Porous wt% A1203 A002 sintering pellet, exuded FeA1 A017 (Fe40A1-2.9 wt% C)/30 A032, A019, Liquid phase Porous wt% ZrO2 A001 sintering pellet, black and silver areas, no exuded FeA1 A018 (Fe40A1-2.9 wt% C)/22 A032, A019, Liquid phase Porous wt% A1203 A002 sintering pellet, exuded FeA1 A020 FeA1/33 wt% A1203 A032, A002 Liquid phase Porous sintering pellet, fragile, exuded FeA1 A021 Fe40A1/9 wt% TiC/22 A032, A024, Liquid phase Dense wt % A 1203 A002 sintering appearance, but many surface cracks. Some exuded FeA1

COMDO$ifi ow ers., use r Number' (see Table 13.. A022 Fe40A1/18 wt% TiC/14 A032, A024, Liquid phase Dense wt% AlOj A002 sintering appearance, a few surface cracks, no exuded FeA1 A022B Fe40Al/18 wt% TiC/14 A032, A024, Liquid phase Dense wt% Au203 A002 sintering appearance, a few surface cracks, no exuded FeA1 A023 Fe40A1/27 wt% TiC/7 A032, A024, Liquid phase Dense wt% A1203 A002 sintering appearance, a few surface cracks, no exuded FeA1 A025 (Fe40A1-5 wt% Ti)/18 A032, A033, Liquid phase Dense wt% TiC/14 wt% A1203 A024, A002 sintering appearance, several surface cracks, exuded FeA1 A026 (Fe40A1-1.4 wt% C)/18 A032, A019, Liquid phase Dense wt% TiC/14 wt% A1203 A024, A002 sintering appearance, many surface cracks, exuded FeAl A027 (Fe40A1-1.1 wt% Ti)/18 A032, A033, Liquid phase Dense wt% TiC/14 wt% A1203 A024, A002 sintering appearance, many surface cracks, exuded FeA1

de: : : rs d in Number (see Table 1) A028 (Fe40Al-0.3 wt% C)/18 A032, A019, Liquid phase Dense wt% TiC/14 wt% A1203 A024, A002 sintering appearance, a few surface cracks, exuded FeA1 A029 Fe40A1/18 wt% TiC/14 A032, A024, Liquid phase Dense wt% A1203 A002 sintering appearance A030 Fe40A1-0.3 wt% C/18 wt% A032, A019, Liquid phase Dense TiC/14 wt% A1203 A024, A002 sintering appearance A031 Fe40A1/18 wt% TiC/14 A032, A045, Liquid phase Dense wt% A1203 A002 sintering appearance, a few surface cracks, exuded FeA1 A041 Fe40A1/30 wt% A1203 A040, A024, Liquid phase Fragile A002 sintering coupon, exuded FeA1 A042A Fe40A1/18 wt% TiC/14 A040, A024, Liquid phase Porous wt% A1203 A002 sintering coupon with many surface cracks, exuded FeAl A043A Fe40A1/24 wt% A1203 A032, A002 Liquid phase Porous sintering coupon, a few surface cracks, exuded FeA1 A043B Fe40A1/24 wt% A1203 A032, A004 Liquid phase Porous sintering coupon, very fragile, exuded FeA 1 ... pecl. mdg :...... N : : umbe A044 FeAI/24 wt% A1203 A040, A002 Liquid phase Porous sintering coupon, very fragile, exuded FeAl

As a result of the above experiments it was determined that FeAl did not wet Auto3 sufficiently well to fabricate FeAl/Al203 composites by liquid phase sintering. In order to improve the sintering behavior, some of the A1203 powder was replaced by TiC powder.

For example, specimen A009 was fabricated from Fe-40 at. % Al powder (-325 mesh or # 45µm), TiC powder (2.5-4, um), and A1203 powder (<38, um), the sample having a nominal composition of FeAl-16.5 vol. % TiC-16.5 vol. % A1203. 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: 1-5, xm Fe, 10, um Al, 2.5-4 µm TiC and <38, um A1203. 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 For Processing Iron Aluminide Technique A009 FeAl-16.5vol% TiC-Prealloyed FeA1, Liquid phase 16. 5vol%Al2O3 TiC and A1203 Sintering at 1450° A046 FeAl-16.5vol % TiC-Prealloyed FeAl, Liquid phase 16.5vol%Al203TiC and A1203 Sintering at 1500°C A047 FeAl-16.5vol% TiC- Fe and Al Liquid phase 16. 5vol%Al2O3 Sintering at 1500°C A050 FeAlLiquidphasePrealloyed 16.5vol% TiC-and Nb Sintering at 1500°C 16. 5vol % A1203 A055 FeAl-10%Al2O3 Prealloyed FeAl Hot Pressed ZrO2 A062C FeAl-15vol% TiC-Fe, Al, TiC and Liquid phase Sinteringat1500°C15vol%Al2O3Al2O3 and hot forging at 1000 °C from 20 to 8 mm

Table 4. Specimen Composition at. % Flexure Strength MPa A046E-1 FeAI-l5vol % TiC-304 15vol%Al2O3 A050A-1 FeAl-9wt%Nb-189 16.5vol% TiC- 16.5vol%Al2O3 A050A-2 185 16.5vol% TiC- 16.5vol% Al203 -212A055#1FeAl-10vol%Al2O3 lOvol% ZrO2 A055#1 FeAl-10vol%Al2O3 217 lOvol% ZrO2 A055#1 FeAl-10vol% Al203-249 10vol%ZrO2 A055#2 169- 10vol%ZrO2 A055#2 FeAl-10vol% Al203-226 10vol%ZrO2 996A062C#1FeAl-15vol%TiC- 15vol%Al2O3 A062C#1 FeAl-15vol% TiC-1081 15vol%Al2O3 A062C#1 FeAl-15vol% TiC- 1160 15vol % Al203

A062C#2FeAl-15vol% TiC- 1099 15vol % Al203 A062C#2FeAl-15vol% TiC- 1202 15vol % Al203 A062C#2FeAl-15vol% TiC- 1173 15vol % Al203 A062C#3FeAl-15vol% TiC- 1056 15vol % Al203 A062C#3FeAl-15vol % TiC-981 15vol % Al203 Specimens with the nominal composition FeAl-16. 5vol% TiC-16. 5vol% Al203 were also fabricated by cold-pressing and subsequent sintering for 12 minutes at 1500°C in vacuum. Similar results were achieved using prealloyed FeAl (specimen A046) or elemental Fe and Al 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 At03 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, Al, TiC and Alto3 and liquid phase sintering the mixture in an Alto3 crucible to provide a FeAl-15TiC-15Al203 (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 Alto3 particles, with black pores surrounded by gray iron aluminide matrix are clearly visible. Processing with elemental

Fe and Al powders, instead of prealloyed FeAl 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 having a cross section of approximately 3x4 mm. The flexure tests were carried out with a <BR> <BR> <BR> <BR> span of 20 mm and a cross head speed of 10, um/s. The fracture stress q was calculated<BR> <BR> <BR> <BR> <BR> from the linear-elastic equation: (if 1.1.5 L-P/ (we), 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 FeAl-16.5TiC-16. SA1203 (vol. %) exceeded 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 higher porosity. Although the Nb may have strengthened the interfacial AOg/FeAl 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-15TiC- 15Al203, 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 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.

The room temperature fracture toughness of the hot forged FeAl-15TiC-15A10 :, 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'l 1/2 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-v2). A value of 0.25 was <BR> <BR> <BR> <BR> assumed for v. The Young's modulus E was estimated from the following equation: E = Em2)(1+c)2-Em2+EpEm]/[(cEp+Em)(1+c)2]wherec=(1/Vp)1/3-1.Vpis the[(cEpEm+ <BR> <BR> <BR> volume fraction of the ceramic particles, l 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 FeAl-15TiC-15AI203 (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 mu'2), the composites exhibited satisfactory fracture toughnesses.

Table 5. Fracture Toughness of Hot Forged FeAl-15TiC-15A03 (vol. %) HWAG1EKQSampleW 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 A1203 is not wetted well enough by liquid FeAl to allow the processing of composites by liquid phase sintering. In contrast to Alto3, ZrO2 is thermodynamically not stable in contact with iron aluminides.

Since brittle intermetallic phases form during the reaction between ZrO2 and FeAl, ZrO2 is less desirable as a filler in FeAl/ceramic composites. On the other hand, TiC promotes wetting of Al203 by FeAl. Moreover, instead of prealloyed FeAl, elemental Fe and Al powders may be used for liquid phase sintering of FeAl/TiC/AOg 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

approximately 300 MPa can be achieved for liquid phase sintered iron aluminide composites containing TiC and Al203. Hot forging of liquid phase sintered FeAl-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 i/2.

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 art without departing from the scope of the present invention as defined by the following claims.