EP0587960B1 - Production of iron aluminide materials - Google Patents

Production of iron aluminide materials Download PDF

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Publication number
EP0587960B1
EP0587960B1 EP92810713A EP92810713A EP0587960B1 EP 0587960 B1 EP0587960 B1 EP 0587960B1 EP 92810713 A EP92810713 A EP 92810713A EP 92810713 A EP92810713 A EP 92810713A EP 0587960 B1 EP0587960 B1 EP 0587960B1
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Prior art keywords
alloy
temperature
dispersoids
dispersoid
rolling
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German (de)
French (fr)
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EP0587960A1 (en
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Emad Dr. Batawi
John Antony Dr. Peters
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Sulzer Markets and Technology AG
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Sulzer Innotec AG
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Priority to AT92810713T priority patent/ATE166112T1/en
Priority to DE59209325T priority patent/DE59209325D1/en
Priority to US08/120,718 priority patent/US5346562A/en
Publication of EP0587960A1 publication Critical patent/EP0587960A1/en
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    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/06—Ferrous alloys, e.g. steel alloys containing aluminium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C32/00—Non-ferrous alloys containing at least 5% by weight but less than 50% by weight of oxides, carbides, borides, nitrides, silicides or other metal compounds, e.g. oxynitrides, sulfides, whether added as such or formed in situ
    • C22C32/0047—Non-ferrous alloys containing at least 5% by weight but less than 50% by weight of oxides, carbides, borides, nitrides, silicides or other metal compounds, e.g. oxynitrides, sulfides, whether added as such or formed in situ with carbides, nitrides, borides or silicides as the main non-metallic constituents
    • C22C32/0068—Non-ferrous alloys containing at least 5% by weight but less than 50% by weight of oxides, carbides, borides, nitrides, silicides or other metal compounds, e.g. oxynitrides, sulfides, whether added as such or formed in situ with carbides, nitrides, borides or silicides as the main non-metallic constituents only nitrides

Definitions

  • the invention relates to a method for producing Iron aluminide materials and according to the process manufactured materials.
  • Iron aluminides which mainly consist of Fe 3 Al, are characterized by an ordered crystal structure with DO 3 symmetry: half of the lattice sites that form a cubic lattice are occupied by Fe atoms; the other half of the lattice sites, which are centered on the cubes of the first lattice, have a checkerboard arrangement of Fe and Al atoms.
  • the iron aluminide base alloy is an ordered intermetallic alloy. It is called Fe 3 Al-based alloy in the following.
  • the Fe 3 Al-based alloy also has a B2 structure (or CsCl structure) or a disordered body-centered Alfa structure.
  • Fe 3 Al-based alloys In the case of known Fe 3 Al-based alloys to which up to 10 at.% Chromium and small amounts of molybdenum, niobium, zirconium, yttrium, vanadium, carbon and / or boron are mixed, no low-melting eutectics are formed.
  • Fe 3 Al-based alloys have a protective aluminum oxide layer covering the surface.
  • iron aluminides and many of the Fe 3 Al-based alloys have a very low ductility at room temperature. Only when the great brittleness of these materials can be overcome can they be used as materials.
  • Ductility can usually be improved if the graininess of the structure is refined using alloy additives.
  • alloy additives are known, for example, from US-A-3 026 197 and US-A-5 084 109.
  • a Fe 3 Al-based alloy is known from a publication (SADavid et al (1989), Welding Research Sup., P.372), in which an increase in ductility at room temperature has been achieved by adding titanium diboride (TiB 2 ) .
  • TiB 2 titanium diboride
  • heat cracking was observed in welding tests (with electron beam, arc). Investigations with secondary ion mass spectrometry showed that enriched boron and titanium occur on the crack surface.
  • the titanium diboride dissolves in the melt; it has no influence on the grain formation. Titanium and boron are not incorporated into the crystal structure of the grains, so these components can finally be found at the grain interfaces after the Fe 3 Al-based alloy has solidified. Due to the influence of heat during welding, the adhesion between neighboring grains is greatly reduced due to the titanium diboride (due to the local melting point lowering at the grain boundaries), which can result in heat cracking. As a result, it is advisable not to add titanium diboride or substances that lead to similar phenomena, despite the improvement in ductility.
  • the dispersoids must be very small (in the range of 100 nm), it is recommended to pass these particles through Precipitate from the melt.
  • To components of the dispersoid compound are mixed in the melt to which ones go into solution first. While a hold time between 100 and 1000 seconds then the dissolved components together, whereby it precipitates the dispersoid-shaped compound form.
  • dispersoids were obtained with a size distribution in which the dispersoid diameters are mostly between 50 and 200 nm.
  • the alloy FA-129 known from WO 90/10722 composition: 28% Al, 5% Cr, 0.5% Nb, 0.2% C, rest Fe) was used as the starting alloy.
  • the melt has a considerable increase in its viscosity due to the dispersoids. For this reason, the melt must be poured when the superheat is relatively high (around 200 K) - in contrast to the pouring of the dispersoid-free melt. As a result, the microstructural grains of small samples, despite the dispersoids, turn out to be roughly the same size as the original Fe 3 Al-based alloy; in the case of large castings, even much larger grains form.
  • Metallurgical studies have shown that dispersoids are embedded in the monocrystalline phase thanks to the good coherence of the crystal structures. When forming by hot rolling, the grains formed during solidification are reduced to fine grains by breaking up new grain boundaries at the points where the dispersoids are embedded in the phase. Annealing the hot-rolled alloy at temperatures between 800 and 1000 ° C results in a stable high-temperature material.
  • dispersion hardening also takes place. This is confirmed by hardness measurements.
  • the hardness (Vickers hardness HV, test load 1 kg) is 260 after casting, 280 after hot rolling (900 ° C, 90%) and further 280 after annealing (600 ° C, 24 h) ; the corresponding values for the dispersoid-free alloy are: 230, 275 and 255. Thanks to the dispersion hardening, the creeping capacity of the material is advantageously reduced.
  • FIG. 1 is shown in FIG schematic form and on a smaller scale in Fig.2 recognizable.
  • the square section 2 in Fig.1 is shown enlarged in Fig.3.
  • the dispersoids unfold an important one Effect: As with hot rolling from 1 to 2 kg heavy, dispersoid-containing castings, grains that are 25 microns wide (and 0.5 mm long), while the corresponding transformation at a particle-free alloy to 60 micron grains Width (length also 0.5 mm) leads. After this Hot rolling are the grains of the invention Material significantly finer than that of the dispersoid-free Alloy, despite the fact that after the Pour the conditions were just reversed.

Abstract

The process makes possible the production of iron aluminide materials comprising an Fe3Al-based alloy with 18-35% of Al, with 3-15% of Cr, with 0.2-0.5% of B and/or C, and with a total of 0-8% of the following alloy additives: Mo, Nb, Zr, Y and/or V, and also Fe as predominant remainder. According to the invention, additives are added to the melt of a known alloy from which dispersed crystallites, so-called dispersoids, form and which are, thanks to good wettability, bedded into the monocrystalline phase on solidification. A fine-grained microstructure can be generated from the solid alloy by hot rolling at a temperature between 650 and 1000 DEG C. <IMAGE>

Description

Die Erfindung betrifft ein Verfahren zum Herstellen von Eisenaluminid-Werkstoffen sowie nach dem Verfahren hergestellte Werkstoffe. The invention relates to a method for producing Iron aluminide materials and according to the process manufactured materials.

Aus der Patentanmeldung WO 90/10722 ist bekannt, dass gewisse Eisenaluminidbasis-Legierungen als Material für die Ausführung von technischen Konstruktionen geeignet sind, insbesondere für Konstruktionen, die bei erhöhter Temperatur (bis 650°C) und in aggressiver Umgebung (beispielsweise H2S + H2 + H2O) eine gute Korrosionsbeständigkeit sowie eine gute mechanische Festigkeit aufweisen müssen. Solche Legierungen bieten sich beispielsweise als kostengünstiger Ersatz für Nickelbasis-Legierungen oder hochlegierte Stähle an. Eisenaluminide, die hauptsächlich aus Fe3Al bestehen, zeichnen sich durch eine geordnete Kristallstruktur mit DO3-Symmetrie aus: die eine Hälfte der Gitterplätze, die ein kubisches Gitter bilden, sind von Fe-Atomen besetzt; die andere Hälfte der Gitterplätze, die bezüglich den Kuben des ersten Gitters raumzentriert liegen, weisen eine schachbrettartige Anordnung von Fe- und Al-Atomen auf. Die Eisenaluminidbasis-Legierung ist eine geordnete intermetallische Legierung. Sie wird im folgenden Fe3Al-basis-Legierung genannt. Der Anteil des Aluminiums dieser Legierung mit DO3-Struktur weist einen Wert im Bereich zwischen 18 und 35% (at.% = Atomprozent) auf. Neben der DO3-Struktur liegt in der Fe3Al-basis-Legierung teilweise auch eine B2-Struktur (oder CsCl-Struktur) oder eine ungeordnete raumzentrierte Alfa-Struktur vor. From patent application WO 90/10722 it is known that certain iron aluminide-based alloys are suitable as a material for the execution of technical constructions, in particular for constructions which are carried out at elevated temperature (up to 650 ° C.) and in an aggressive environment (for example H 2 S + H 2 + H 2 O) must have good corrosion resistance and good mechanical strength. Such alloys offer themselves, for example, as a cost-effective replacement for nickel-based alloys or high-alloy steels. Iron aluminides, which mainly consist of Fe 3 Al, are characterized by an ordered crystal structure with DO 3 symmetry: half of the lattice sites that form a cubic lattice are occupied by Fe atoms; the other half of the lattice sites, which are centered on the cubes of the first lattice, have a checkerboard arrangement of Fe and Al atoms. The iron aluminide base alloy is an ordered intermetallic alloy. It is called Fe 3 Al-based alloy in the following. The proportion of aluminum in this alloy with a DO 3 structure has a value in the range between 18 and 35% (at.% = Atomic percent). In addition to the DO 3 structure, the Fe 3 Al-based alloy also has a B2 structure (or CsCl structure) or a disordered body-centered Alfa structure.

Bei bekannten Fe3Al-basis-Legierungen, denen bis zu 10 at.% Chrom und in kleineren Mengen Molybdän, Niobium, Zirkonium, Yttrium, Vanadium, Kohlenstoff und/oder Bor zugemischt sind, bilden sich keine niedrigschmelzenden Eutektika aus. Fe3Al-basis-Legierungen weisen eine schützende, die Oberfläche überziehende Aluminiumoxidschicht auf. Allerdings haben Eisenaluminide und viele der Fe3Al-basis-Legierungen eine sehr geringe Duktilität bei Raumtemperatur. Erst wenn die grosse Sprödigkeit dieser Materialien überwunden werden kann, sind sie als Werkstoffe verwendbar. In the case of known Fe 3 Al-based alloys to which up to 10 at.% Chromium and small amounts of molybdenum, niobium, zirconium, yttrium, vanadium, carbon and / or boron are mixed, no low-melting eutectics are formed. Fe 3 Al-based alloys have a protective aluminum oxide layer covering the surface. However, iron aluminides and many of the Fe 3 Al-based alloys have a very low ductility at room temperature. Only when the great brittleness of these materials can be overcome can they be used as materials.

Die Duktilität kann in der Regel verbessert werden, wenn mittels Legierungszusätzen die Körnigkeit des Gefüges verfeinert wird. Solche Legierungszusätze sind z.B. aus US-A-3 026 197 und US-A-5 084 109 bekannt. Aus einer Druckschrift (S.A.David et al (1989), Welding Research Sup., p.372) ist eine Fe3Al-basis-Legierung bekannt, bei der mittels Zusatz von Titandiborid (TiB2) eine Erhöhung der Duktilität bei Raumtemperatur erzielt worden ist. Allerdings wurde bei Schweissversuchen (mit Elektronenstrahl, Lichtbogen) eine Wärmerissbildung beobachtet. Untersuchungen mit Sekundärionen-Massenspektrometrie ergab, dass an der Rissfläche Bor und Titan angereichert auftreten. Dieser Befund führte zu folgender Einsicht: Das Titandiborid geht in der Schmelze in Lösung; es hat auf die Kornbildung keinen Einfluss. Titan und Bor werden nicht in die Kristallstruktur der Körner eingebaut, daher sind diese Komponenten schliesslich nach dem Erstarren der Fe3Al-basis-Legierung an den Korngrenzflächen vorzufinden. Durch den Wärmeeinfluss beim Schweissen wird der Kraftschluss zwischen benachbarten Körnern wegen des Titandiborids stark reduziert (wegen lokaler Schmelzpunkterniedrigung an den Korngrenzen), wodurch sich eine Wärmerissbildung einstellen kann. Folglich ist es ratsam, trotz Verbesserung der Duktilität auf den Zusatz von Titandiborid oder Stoffen, die zu ähnlichen Erscheinungen führen, zu verzichten. Ductility can usually be improved if the graininess of the structure is refined using alloy additives. Such alloy additives are known, for example, from US-A-3 026 197 and US-A-5 084 109. A Fe 3 Al-based alloy is known from a publication (SADavid et al (1989), Welding Research Sup., P.372), in which an increase in ductility at room temperature has been achieved by adding titanium diboride (TiB 2 ) . However, heat cracking was observed in welding tests (with electron beam, arc). Investigations with secondary ion mass spectrometry showed that enriched boron and titanium occur on the crack surface. This finding led to the following insight: The titanium diboride dissolves in the melt; it has no influence on the grain formation. Titanium and boron are not incorporated into the crystal structure of the grains, so these components can finally be found at the grain interfaces after the Fe 3 Al-based alloy has solidified. Due to the influence of heat during welding, the adhesion between neighboring grains is greatly reduced due to the titanium diboride (due to the local melting point lowering at the grain boundaries), which can result in heat cracking. As a result, it is advisable not to add titanium diboride or substances that lead to similar phenomena, despite the improvement in ductility.

Es ist Aufgabe der Erfindung, durch Zugabe geeigneter Stoffe und Durchführung geeigneter Verfahrensschritte die Kornbildung in Eisenaluminidbasis-Legierungen solcherart zu beeinflussen, dass eine verbesserte Duktilität bei Raumtemperatur erzielbar ist, wobei der erfindungsgemässe Werkstoff nebst einer hohen Festigkeit bei erhöhter Temperatur, eine gute Schweissbarkeit aufweisen soll. Diese Aufgabe wird durch die im Kennzeichen des Anspruchs 1 genannten Massnahmen gelöst. Die Anspruche 2 bis 4 betreffen bevorzugte Ausführungsformen des beanspruchten Verfahrens. It is an object of the invention, by adding more suitable ones Substances and implementation of suitable process steps Grain formation in iron aluminide base alloys of this type to influence that improved ductility Room temperature can be achieved, the inventive Material in addition to high strength with increased Temperature that should have good weldability. This task is accomplished by the in the hallmark of the claim 1 mentioned measures resolved. Claims 2 to 4 relate to preferred embodiments of the claimed method.

Die ursprüngliche Idee der Erfindung hat darin bestanden, kleine Partikel - sogenannte Dispersoide - in der aufgeschmolzenen Fe3Al-basis-Legierung zu dispergieren, welche als Keimbildner wirken. Bei der Suche nach geeigneten Stoffen ist von folgenden Forderungen auszugehen:

  • 1. Die Dispersoide sollen stabile, kristalline Partikel sein, die sich bei der Giesstemperatur nicht in der Schmelze auflösen. Der Schmelzpunkt der für die Dispersoide verwendeten Verbindung muss wesentlich grösser als die Liquidustemperatur (rund 1450°C) der Fe3Al-basis-Legierung sein.
  • 2. Die Dispersoide sollen gut benetzbar sein, d.h. die Grenzflächenenergie zwischen den kristallinen Partikeln und der Schmelze soll klein sein. Damit die Dispersoide mögliche Keimbildner sind, müssen an deren Oberfläche Gitterebenen vorhanden sein, für welche die Gitterkonstante ungefähr gleich gross wie die Gitterkonstante von Fe3Al bei der Erstarrung (CsCl-Struktur) ist, nämlich etwa 0,4 nm.
  • 3. Die Dichte der Dispersoide soll sich wenig von der Dichte der Schmelze (rund 6 bis 6,5 g/cm3) unterscheiden, damit eine inhomogene Verteilung der Dispersoide aufgrund von Sedimentation im wesentlichen ausbleibt.
    • The original idea of the invention was to disperse small particles - so-called dispersoids - in the melted Fe 3 Al-based alloy, which act as nucleating agents. When searching for suitable substances, the following requirements can be assumed:
  • 1. The dispersoids should be stable, crystalline particles that do not dissolve in the melt at the casting temperature. The melting point of the compound used for the dispersoids must be significantly higher than the liquidus temperature (around 1450 ° C) of the Fe 3 Al-based alloy.
  • 2. The dispersoids should be readily wettable, ie the interfacial energy between the crystalline particles and the melt should be small. In order for the dispersoids to be possible nucleating agents, lattice planes must be present on their surface, for which the lattice constant is approximately the same size as the lattice constant of Fe 3 Al during solidification (CsCl structure), namely approximately 0.4 nm.
  • 3. The density of the dispersoids should differ little from the density of the melt (around 6 to 6.5 g / cm 3 ), so that an inhomogeneous distribution of the dispersoids essentially does not occur due to sedimentation.

    • Da die Dispersoide sehr klein sein müssen (im Bereich von 100 nm), empfiehlt es sich, diese Partikel durch Ausfällung aus der Schmelze entstehen zu lassen. Dazu mischt man der Schmelze Komponenten der Dispersoid-Verbindung zu, welche zunächst in Lösung gehen. Während einer Haltezeit zwischen 100 und 1000 Sekunden reagieren anschliessend die gelösten Komponenten miteinander, wobei sie unter Ausfällung die dispersoidförmige Verbindung bilden. Since the dispersoids must be very small (in the range of 100 nm), it is recommended to pass these particles through Precipitate from the melt. To components of the dispersoid compound are mixed in the melt to which ones go into solution first. While a hold time between 100 and 1000 seconds then the dissolved components together, whereby it precipitates the dispersoid-shaped compound form.  

    Ein Versuch, Dispersoide in der Schmelze der Fe3Al-basis-Legierung herzustellen, wurde erfolgreich mit einer Verbindung ausgeführt, nämlich mit Titan/Zirkonium-Nitrid, (Ti,Zr)N. Ti und Zr (2 - 10 g/kg) wurden zur Bildung von 2 - 10 Volumenprozent Dispersoiden als Metallgranulat in die überhitzte Schmelze eingetragen, während der atomare Stickstoff (N) mittels eines Trägers, nämlich in Form einer N-haltigen Fe-Cr-Legierung, in die Schmelze befördert wurde. Damit der Stickstoff nicht ausgaste, wurde die Dispersoiderzeugung bei einem Druck von 0.5 bar durchgeführt, welcher mittels einer Schutzgasatmosphäre aus Argon hergestellt wurde (Es wäre auch ein anderer Druck zwischen 0.2 und 1 bar möglich). Während einer Haltezeit von 300 s und bei 1650°C ergaben sich Dispersoide mit einer Grössenverteilung, bei der die Dispersoiddurchmesser grösstenteils zwischen 50 und 200 nm liegen. Als Ausgangslegierung wurde die aus der WO 90/10722 bekannte Legierung FA-129 (Zusammensetzung: 28% Al, 5% Cr, 0.5% Nb, 0.2% C, Rest Fe) verwendet. An attempt to produce dispersoids in the melt of the Fe 3 Al-based alloy was successfully carried out with a compound, namely with titanium / zirconium nitride, (Ti, Zr) N. Ti and Zr (2-10 g / kg) were introduced into the superheated melt as metal granules to form 2-10 volume percent dispersoids, while the atomic nitrogen (N) was carried out by means of a carrier, namely in the form of an N-containing Fe-Cr- Alloy into which the melt has been conveyed. So that the nitrogen did not outgas, the dispersoid was produced at a pressure of 0.5 bar, which was produced from argon using a protective gas atmosphere (a different pressure between 0.2 and 1 bar would also be possible). During a holding time of 300 s and at 1650 ° C, dispersoids were obtained with a size distribution in which the dispersoid diameters are mostly between 50 and 200 nm. The alloy FA-129 known from WO 90/10722 (composition: 28% Al, 5% Cr, 0.5% Nb, 0.2% C, rest Fe) was used as the starting alloy.

    Durch die Dispersoide erfährt die Schmelze eine beträchliche Vergrösserung ihrer Viskosität. Deshalb muss das Giessen der Schmelze bei relativ grosser Überhitzung (rund 200 K) - im Gegensatz zum Giessen der dispersoidfreien Schmelze - vorgenommen werden. Dies hat zur Folge, dass bei kleinen Proben die Gefügekörner trotz der Dispersoide ungefähr gleich gross wie bei der ursprünglichen Fe3Al-basis-Legierung ausfallen; bei grossen Gussstücken bilden sich sogar weit grössere Körner aus. Metallurgische Untersuchungen haben ergeben, das innerhalb der Körner Dispersoide dank guter Kohärenz der Kristallstrukturen in die monokristalline Phase eingebettet sind. Beim Umformen durch Warmwalzen verkleinern sich die beim Erstarren entstandenen Körner zu feineren Körnern, indem an den Stellen, an denen die Dispersoide in die Phase eingebettet sind, neue Korngrenzen aufbrechen. Durch Glühen der warmgewalzten Legierung bei Temperaturen zwischen 800 und 1000°C ergibt sich ein stabiler Hochtemperaturwerkstoff. The melt has a considerable increase in its viscosity due to the dispersoids. For this reason, the melt must be poured when the superheat is relatively high (around 200 K) - in contrast to the pouring of the dispersoid-free melt. As a result, the microstructural grains of small samples, despite the dispersoids, turn out to be roughly the same size as the original Fe 3 Al-based alloy; in the case of large castings, even much larger grains form. Metallurgical studies have shown that dispersoids are embedded in the monocrystalline phase thanks to the good coherence of the crystal structures. When forming by hot rolling, the grains formed during solidification are reduced to fine grains by breaking up new grain boundaries at the points where the dispersoids are embedded in the phase. Annealing the hot-rolled alloy at temperatures between 800 and 1000 ° C results in a stable high-temperature material.

    Durch das Einbringen der Dispersoide in die Fe3Al-basis-Legierung findet auch eine Dispersionshärtung statt. Dies wird durch Härtemessungen bestätigt. Beim erwähnten Beispiel mit den Nitrid-Dispersoiden beträgt die Härte (Vickershärte HV, Prüflast 1 kg) 260 nach dem Giessen, 280 nach dem Warmwalzen (900°C, 90%) und weiterhin 280 nach dem Glühen (600°C, 24 h); die entsprechenden Werte bei der dispersoidfreien Legierung sind: 230, 275 bzw. 255. Dank der Dispersionshärtung verringert sich vorteilhafterweise das Kriechvermögen des Werkstoffs. By incorporating the dispersoids into the Fe 3 Al-based alloy, dispersion hardening also takes place. This is confirmed by hardness measurements. In the example mentioned with the nitride dispersoids, the hardness (Vickers hardness HV, test load 1 kg) is 260 after casting, 280 after hot rolling (900 ° C, 90%) and further 280 after annealing (600 ° C, 24 h) ; the corresponding values for the dispersoid-free alloy are: 230, 275 and 255. Thanks to the dispersion hardening, the creeping capacity of the material is advantageously reduced.

    Der Werkstoff gemäß den Ansprüchen 5 and 6, hergestellt nach dem erfindungsgemässen Verfahren, der nach dem Erstarren der dispersoidhaltigen Schmelze vorliegt, wird anhand von Zeichnungen näher erläutert. Es zeigen:

    Fig. 1
    eine Probe einer erfindungsgemässen Legierung (500-fach vergrössert, nach einem Rasterelektronenmikroskopie-Bild gezeichnet),
    Fig. 2
    eine schematische Darstellung der gleichen Probe wie in Fig.1, bei kleinerer Vergrösserung (200-fach), und
    Fig. 3
    einen Ausschnitt aus der Probe von Fig.1 mit Dispersoiden (5000-fach vergrössert).
    The material according to claims 5 and 6, produced by the method according to the invention, which is present after the dispersoid-containing melt has solidified, is explained in more detail with reference to drawings. Show it:
    Fig. 1
    a sample of an alloy according to the invention (magnified 500 times, drawn according to a scanning electron microscope image),
    Fig. 2
    a schematic representation of the same sample as in Fig.1, at a smaller magnification (200 times), and
    Fig. 3
    a section of the sample of Fig.1 with dispersoids (magnified 5000 times).

    Der in Fig.1 dargestellte Bildausschnitt 1 ist in schematischer Form und bei kleinerem Massstab in Fig.2 erkennbar. Der quadratische Ausschnitt 2 in Fig.1 ist vergrössert in Fig.3 gezeigt. 1 is shown in FIG schematic form and on a smaller scale in Fig.2 recognizable. The square section 2 in Fig.1 is shown enlarged in Fig.3.

    Der in Fig.1 strichpunktiert gezeichnete Streckenzug 3, der dem ausgezogenen Streckenzug 3' in Fig.2 entspricht, trennt eine monokristalline Eisenaluminid-Phase 5 von einem eutektischen Gebiet 6. Im Gebiet 6 befinden sich skelettartige Kristalle 30, die reich an Eisen, Chrom und Niobium sind. Die Fig.2 bietet einen besseren Überblick über die Verteilung von eutektischen Gebieten 6 und Eisenaluminid-Phase 5. In der Phase 5 sind Titan/Zirkonium-Nitrid-Dispersoide 20 eingebettet, die in Fig.1 als strukturlose Punkte erscheinen. (Der Nachweis, dass die beobachteten Partikel tatsächlich aus der angegebenen Verbindung (Ti,Zr)N bestehen, ist mittels energiedispersiver Elektronenstrahlanalyse erfolgt.) Die vier Kristallite 20 des Ausschnitts 2 sind in der Vergrösserung der Fig.3 als kleine Kreise dargestellt. Der grösste Durchmesser eines Dispersoids 20 beträgt rund 0,3 Mikrometer. Über die Gestalt der Dispersoide lässt sich aufgrund der mit dem Rasterelektronenmikroskop gemachten Bilder keine Aussage machen. The line 3 drawn in dot-dash lines in FIG. 1, which corresponds to the extended line 3 'in Figure 2,   separates a monocrystalline iron aluminide phase 5 from a eutectic area 6. Located in area 6 skeletal crystals 30 rich in iron, chromium and Are niobium. Fig. 2 offers a better overview on the distribution of eutectic areas 6 and Iron aluminide phase 5. Are in phase 5 Titanium / zirconium nitride dispersoids 20 embedded in Fig.1 appear as structureless points. (The proof, that the observed particles are actually from the specified compound (Ti, Zr) N exist, is by means of Energy dispersive electron beam analysis is carried out.) four crystallites 20 of section 2 are in the Enlargement of Figure 3 shown as small circles. The largest diameter of a dispersoid 20 is round 0.3 microns. About the shape of the dispersoids yourself due to using the scanning electron microscope taken pictures do not make a statement.

    Bei thermomechanischen Umformungen der partikelhaltigen Legierung entfalten die Dispersoide eine wichtige Wirkung: Wie sich beim Warmwalzen von 1 bis 2 kg schweren, dispersoidhaltigen Gussstücken gezeigt hat, entstehen Körner, die 25 Mikrometer breit (und 0.5 mm lang) sind, während die entsprechende Umformung bei einer partikelfreien Legierung zu Körnern mit 60 Mikrometer Breite (Länge ebenfalls 0.5 mm) führt. Nach dem Warmwalzen sind die Körner des erfindungsgemässen Werkstoffs bedeutend feiner als jene der dispersoidfreien Legierung, und dies trotz der Tatsache, dass nach dem Giessen die Verhältnisse gerade umgekehrt gewesen sind. With thermomechanical forming of the particle-containing Alloy, the dispersoids unfold an important one Effect: As with hot rolling from 1 to 2 kg heavy, dispersoid-containing castings, grains that are 25 microns wide (and 0.5 mm long), while the corresponding transformation at a particle-free alloy to 60 micron grains Width (length also 0.5 mm) leads. After this Hot rolling are the grains of the invention Material significantly finer than that of the dispersoid-free Alloy, despite the fact that after the Pour the conditions were just reversed.

    Beim Warmwalzen des dispersoidhaltigen Gussstücks soll bis zu einer Querschnittsreduktion von mindestens 80% gewalzt werden. When hot-rolling the dispersoid-containing casting up to a cross-section reduction of at least 80% be rolled.

    Claims (6)

    1. A process for manufacturing iron aluminide materials from a Fe3A1-based alloy, having 18-35 at% A1, having 3-15 at% Cr, having 0.2-5 at% B and/or C, and having total adhesion 0-8 at% of the following alloy additions, Mo, Nb, Zr, Y and/or V, likewise having Fe as the predominant residue, which process encompasses the following steps:
      melting the alloy in vacuum,
      manufacturing a protective gas over the molten mass, having a gas pressure of between 0.2 and 1 bar,
      addition of Ti, Zr, and an N-containing Fe-Cr-alloy at a temperature of 200-400 K above the liquidus temperature to form 2-10 percent by volume (Ti,Zr)N-dispersoids,
      pumping out the protective gas after a pause lasting between 100 and 1000 s,
      casting the dispersoid-containing molten mass at a temperature which is 100 - 200 K above the liquidus temperature, and allowing it to solidify,
      hot rolling at a temperature between 650 and 1000 C, and where it is rolled until a reduction in cross-section of at least 80% is achieved.
    2. A process according to claim 1, characterised in that after hot rolling, the alloy is annealed at a temperature between 400 and 1000°C.
    3. A process according to claims 1 or 2, characterised in that rolling takes place at about 900°C.
    4. A process according to claims 1 to 3, characterised in that subsequent to the rolling, annealing takes place at about 600°C during about 24 hours.
    5. A material produced in one of the processes according to claims 1 to 4, characterised in that the smallest diameter of the granular structure is smaller than about 30 micrometers.
    6. A material according to claim 5, characterised in that the dispersoid-free Fe3A1-based alloy of 26-30 at% A1, 3-10 at% Cr, 0.3-0.8 at% Nb, 0.1-0.5 at% C, residue Fe, is used as the initial alloy.
    EP92810713A 1992-09-16 1992-09-16 Production of iron aluminide materials Expired - Lifetime EP0587960B1 (en)

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    EP92810713A EP0587960B1 (en) 1992-09-16 1992-09-16 Production of iron aluminide materials
    AT92810713T ATE166112T1 (en) 1992-09-16 1992-09-16 PRODUCTION OF IRON ALUMINIDE MATERIALS
    DE59209325T DE59209325D1 (en) 1992-09-16 1992-09-16 Manufacture of iron aluminide materials
    US08/120,718 US5346562A (en) 1992-09-16 1993-09-13 Method of production of iron aluminide materials

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    EP92810713A EP0587960B1 (en) 1992-09-16 1992-09-16 Production of iron aluminide materials

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    EP0587960B1 true EP0587960B1 (en) 1998-05-13

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    DE19603515C1 (en) * 1996-02-01 1996-12-12 Castolin Sa Spraying material used to form corrosive-resistant coating
    DE19735217B4 (en) * 1997-08-14 2004-09-09 SCHWäBISCHE HüTTENWERKE GMBH Composite material with a high proportion of intermetallic phases, preferably for friction bodies
    US6030472A (en) 1997-12-04 2000-02-29 Philip Morris Incorporated Method of manufacturing aluminide sheet by thermomechanical processing of aluminide powders
    US6114058A (en) * 1998-05-26 2000-09-05 Siemens Westinghouse Power Corporation Iron aluminide alloy container for solid oxide fuel cells
    US6816934B2 (en) * 2000-12-22 2004-11-09 Hewlett-Packard Development Company, L.P. Computer system with registered peripheral component interconnect device for processing extended commands and attributes according to a registered peripheral component interconnect protocol
    US6524405B1 (en) * 2000-02-11 2003-02-25 Hui Lin Iron base high temperature alloy
    DE102009020922A1 (en) 2009-05-12 2010-11-18 Christoph Henrik Sterzel Use of liquid sulfur containing hydrogen sulfide and polysulfane or chlorine, as heat transfer- and heat storage liquid for transporting and storing of thermal energy, preferably in solar thermal power plants
    EP2733829A1 (en) * 2012-11-15 2014-05-21 Siemens Aktiengesellschaft Component for an electrical dynamo machine
    AT513255B1 (en) * 2012-12-28 2014-03-15 Miba Gleitlager Gmbh Multilayer plain bearings
    CN108603257B (en) * 2016-01-20 2021-02-26 蒂森克虏伯钢铁欧洲股份公司 Flat steel product and method for the production thereof

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    US1990650A (en) * 1932-06-25 1935-02-12 Smith Corp A O Heat resistant alloy
    US2726952A (en) * 1954-05-05 1955-12-13 Ford Motor Co Method of preparation of iron aluminum alloys
    US2768915A (en) * 1954-11-12 1956-10-30 Edward A Gaughler Ferritic alloys and methods of making and fabricating same
    US3026197A (en) * 1959-02-20 1962-03-20 Westinghouse Electric Corp Grain-refined aluminum-iron alloys
    US4961903A (en) * 1989-03-07 1990-10-09 Martin Marietta Energy Systems, Inc. Iron aluminide alloys with improved properties for high temperature applications
    US5084109A (en) * 1990-07-02 1992-01-28 Martin Marietta Energy Systems, Inc. Ordered iron aluminide alloys having an improved room-temperature ductility and method thereof
    DE59007276D1 (en) * 1990-07-07 1994-10-27 Asea Brown Boveri Oxidation and corrosion-resistant alloy for components for a medium temperature range based on doped iron aluminide Fe3Al.

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    US5346562A (en) 1994-09-13
    DE59209325D1 (en) 1998-06-18
    ATE166112T1 (en) 1998-05-15

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