AU2010291651B2 - Method for producing an iron-chromium alloy - Google Patents
Method for producing an iron-chromium alloy Download PDFInfo
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- AU2010291651B2 AU2010291651B2 AU2010291651A AU2010291651A AU2010291651B2 AU 2010291651 B2 AU2010291651 B2 AU 2010291651B2 AU 2010291651 A AU2010291651 A AU 2010291651A AU 2010291651 A AU2010291651 A AU 2010291651A AU 2010291651 B2 AU2010291651 B2 AU 2010291651B2
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- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
- C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
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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
- C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
- C22C38/22—Ferrous alloys, e.g. steel alloys containing chromium with molybdenum or tungsten
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- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D6/00—Heat treatment of ferrous alloys
- C21D6/002—Heat treatment of ferrous alloys containing Cr
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- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D6/00—Heat treatment of ferrous alloys
- C21D6/02—Hardening by precipitation
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- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
- C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
- C21D8/0221—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
- C21D8/0226—Hot rolling
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- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
- C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
- C21D8/0221—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
- C21D8/0236—Cold rolling
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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
- C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
- C22C38/26—Ferrous alloys, e.g. steel alloys containing chromium with niobium or tantalum
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Abstract
The invention relates to a method for producing a component, made of an iron-chromium alloy that precipitates Laves phases and/or particles containing Fe and/or particles containing Cr and/or particles containing Si and/or carbides, by subjecting a semi-finished product made of the alloy to a thermomechanical treatment, wherein in a first step, the alloy is solution heat treated at temperatures ≥ the solution heat treatment temperature and it subsequently quenched in stationary protective gas or air, moving (blown) protective gas or air, or water. In a second step, a mechanical forming of the semi-finished product in a range from 0.05 to 99% is performed, and in a subsequent step, Laves phases Fe
Description
WO 2011/026460 PCT/DE2010/000975 METHOD FOR PRODUCTION OF AN IRON-CHROMIUM ALLOY The invention relates to a ferritic iron-chromium alloy produced by melting metallurgy. DE 100 25 108 Al discloses a high-temperature material comprising an iron alloy forming chromium oxide with up to 2% by weight of at least one oxygen-affine element from the group Y, Ce, Zr, Hf and Al, up to 2% by weight of an element M from the group Mn, Ni and Co, which forms a spinel phase of MCr 2 0 4 type at high temperatures, up to 2% by weight of a further element from the group Ti, Hf, Sr, Ca and Zr, which increases the electrical conductivity of Cr-based oxides. The chromium content should lie within a concentration range between 12 and 28%. Areas of use for this high-temperature material are bipolar plates in a high temperature fuel cell. EP 1 298 228 Al relates to a steel for a high-temperature fuel cell that has the following composition: not more than 0.2% C, not more than 1% Si, not more than 1% Mn, not more than 2% Ni, 15 - 30% Cr, not more than 1% Al, not more than 0.5% Y, not more than 0.2% REM and not more than 1% Zr, the remainder being iron and production-related impurities. Low hot strength and inadequate creep strength at temperatures of 700'C and above are common to these two alloys. Precisely in the range above 700*C up to approximately 900*C, however, these alloys have excellent oxidation and corrosion resistance. A creep-resistant ferritic steel, comprising precipitates of an intermetallic phase of Fe 2 (M, Si) or Fe 7 (M, Si) 6 type with at least one metallic alloying element M, which may be formed by WO 2011/026460 PCT/DE2010/000975 2 the elements niobium, molybdenum, tungsten or tantalum, has become known from DE 10 2006 007 598 Al. The steel is preferably intended to be used for a bipolar plate in a fuel-cell stack. EP 1 536 031 Al discloses a metallic material for fuel cells, containing C 5 0.2%, 0.02 to 1% Si, 2% Mn, 10 to 40% Cr, 0.03 to 5% Mo, 0.1 to 3% Nb, at least one of the elements from the group Sc, Y, La, Ce, Pr, Nd, Pm, Sn, Zr and Hf 5 1, the remainder being iron and unavoidable impurities, wherein the composition is supposed to satisfy the following equation: 0.1 Mo / Nb 30. EP 1 882 756 Al describes a ferritic chromium steel, especially usable in fuel cells. The chromium steel has the following composition: C max. 0.1%, Si 0.1 - 1%, Mn max. 0.6%, Cr 15 25%, Ni max. 2%, Mo 0.5 - 2%, Nb 0.2 - 1.5%, Ti max. 0.5%, Zr max. 0.5%, REM max. 0.3%, Al max. 0.1%, N max. 0.07%, the remainder being Fe and melting-related impurities, wherein the content of Zr + Ti is at least 0.2%. In comparison with DE 100 25 108 Al and EP 1 298 228 A2, all of these alloys have better hot strength and elevated creep strength at temperatures of 700 0 C and above, specifically due to formation of precipitates, which hinder the dislocation movements and thus plastic deformation of the material. In the case of DE 10 2006 007 598 Al, for example, these precipitates consist of a Laves phase, an intermetallic compound with the composition Fe 2 (M, Si) or Fe 7 (M, Si) 6 , wherein M may be niobium, molybdenum, tungsten or tantalum. Therein a proportion by volume of 1 to 8%, preferably 2.5 to 5%, should be reached. However, WO 2011/026460 PCT/DE2010/000975 3 there may also be other precipitates such as Fe-containing particles and/or Cr-containing particles and/or Si-containing particles, such as described, for example, in EP 1 536 031 Al, or there may be carbides containing Nb, W, Mo. It is common to all of these particles that they make deformation of the material difficult. From the state of the art described above, it is known that small additions of Y, Zr, Ti, Hf, Ce, La and similar reactive elements can influence the oxidation resistance of Fe-Cr alloys very positively. The alloys cited in DE 10 2006 007 598 Al, EP 1 536 031 Al and EP 1 882 756 Al are optimized for the application as interconnector plates for the high-temperature fuel cells: By use of a ferritic alloy containing 10 to 40% chromium, they have an expansion coefficient adapted as well as possible to the ceramic components anode and electrolyte. Further requirements on the interconnector steel of a high temperature fuel cell are, besides the creep strength already mentioned above, very good corrosion resistance, good conductivity of the oxide layer and little chromium volatilization. The requirements on the reformers and the heat exchangers for the high-temperature fuel cell are the best possible creep strength, very good corrosion resistance and little chromium volatilization. The oxide for these components does not have to be conductive.
WO 2011/026460 PCT/DE2010/000975 4 The requirements for components for the exhaust-gas line of a combustion engine, for example, or for steam boilers, superheaters, turbines and other parts of a power plant, are best possible creep strength and very good corrosion resistance. In these cases chromium volatilization does not cause any poisoning phenomena as in the fuel cell, and the protecting oxide does not have to be conductive for such components. In DE 10 2006 007 598 Al, for example, the excellent corrosion resistance is achieved by formation of a chromium oxide top layer. By the fact that a spinel containing Mn, Ni, Co or Cu is additionally formed on the chromium oxide top layer, fewer volatile chromium oxides or chromium oxyhydroxides that poison the cathode are formed. By the fact that Si is stably bound in the Fe 2 (M, Si) or Fe 7 (M, Si) 6 Laves phase, a nonconductive subsurface layer of silicon oxide is also not formed under the chromium oxide top layer. The corrosion resistance is further improved by the fact that the Al content is kept low and so the increase of the corrosion due to the internal oxidation of the aluminum is avoided. A small Ti addition additionally favors strengthening of the surface and thus prevents swelling of the oxide layer and the inclusion of metallic zones in the oxide layer, which increases the oxidation. In addition, the addition of oxygen-affine elements such as La, Ce, Y, Zr or the like further increases the corrosion resistance. From the market, increased requirements are being imposed on products, necessitating elevated hot strength and creep strength together with an elongation of at least 18% at application temperature for avoidance of brittle failure together with at least equally good oxidation or corrosion resistance and a WO 2011/026460 PCT/DE2010/000975 5 higher service temperature of the alloy, specifically while retaining acceptable deformability, measured as plastic deformation in the tension test with an elongation of > 13% at room temperature. Furthermore, the following investigation methods are used. In a creep test, a specimen is subjected to a constant static tensile force at a constant temperature. For the purpose of comparability, this tensile force is expressed as an initial tensile stress relative to the initial cross section of the specimen. In the creep test, the time tB until break - the time to break - of the specimen is measured in the simplest case. The test can then be performed without measurement of the elongation of the specimen in the course of the test. The elongation at break is then measured after the end of the test. The specimen is mounted at room temperature in the creep-testing machine and heated to the desired temperature without loading by a tensile force. After reaching the test temperature, the specimen is maintained for one hour without loading for temperature equilibration. Thereafter the specimen is loaded with the tensile force and the test time begins. The time to break can be taken as a measure of the creep strength. The longer the time to break is at a specified temperature and initial tensile stress, the greater the creep strength of the material is. The time to break and the creep strength decrease with increasing temperature and increasing initial tensile stress (see, for example, "Burgel", page 100).
WO 2011/026460 PCT/DE2010/000975 6 The deformability is determined in a tension test according to DIN 50145 at room temperature. In the process, the offset yield strength Rpo.
2 , the tensile strength Rm and the elongation at break are determined. The elongation A is determined on the broken specimen from the elongation of the original gauge length LO: A = (Lu-Lo)/Lo 100% = AL/Lo 100% Where Lu = gauge length after break. Depending on gauge length, the elongation at break is denoted by subscripts:
A
5 , gauge length Lo = 5-do or Lo = 5.65 - So Aio, gauge length Lo = 10-do or Lo = 11.3 qSo or, for example, AL=loo, for the freely chosen gauge length L = 100 mm. (do initial diameter, So initial cross section of the flat specimen) The magnitude of the elongation A in the tension test at room temperature can be taken as a measure of the deformability. The Laves phase(s) or the Fe-containing particles and/or Cr containing particles and/or Si-containing particles and/or carbides can be made visible on a metallographic ground section by etching with V2A pickling fluid or electrolytic etching with oxalic acid. During etching with V2A pickling fluid, the grains or grain boundaries also are additionally etched visibly. Only particles with a size of approximately 0.5 pm and larger are visible by viewing in an optical microscope. Smaller particles may not be recognized, but are definitely present. Therefore WO 2011/026460 PCT/DE2010/000975 7 metallography is used only in support of the explanation, while the efficacy of a measure is assessed more practically by the time to rupture or creep strength. In the Manual of High-Temperature Materials Technology, Ralf Burgel, 3rd Revised Edition, Viehweg Verlag, December 2006, hereinafter referred to as "BUrgel", the "Possible measures for increasing the creep strength of metallic materials" are presented on pages 196 to 199 and in Table 3.7. The measures - "High melting point, face-centered cubic material", - "High modulus of elasticity", - "Material with low stacking fault energy", cannot be used for improvement of the cited parameters, since they necessitate a change of material type, which is not possible here and also is not the task. The measures - "Solid solution hardening" - "Particle hardening" - "High particle volume fraction" - "Particles with small diffusion coefficients of the alloying element in question" have already been employed in DE 10 2006 007 598 Al and/or in EP 1 536 031 Al and/or in EP 1 882 756 Al. The measures - "Particles with low solubility in the matrix" - "Coherent particles with low interfacial enthalpy relative to the matrix", WO 2011/026460 PCT/DE2010/000975 8 are not applicable for the precipitates under consideration. Likewise, the measures - "Carbides or borides as grain-boundary precipitates; avoid oxides and sulfides", - "Add positively active grain-boundary elements in precisely controlled dosage, for example, B, C, Zr, Ce", - "Higher purity of the alloy" - "Add getter elements (for example, for S)", - "High corrosion resistance" have already been described in DE 10 2006 007 598 Al and/or in EP 1 536 031 Al and/or in EP 1 882 756 Al. The measures - "Small dendrite arm spacings in cast microstructures", - "Small grain structure in the main loading direction", - "Single crystal" - "Low density of components loaded by their own weight and rotating" cannot be applied to this alloy type, or to the production route, or the use. For the task of improving the creep strength of the precipitation-hardened iron-chromium alloy, the measures 1) "Coarse grain microstructure", 2) "Jaggedness of the grain boundaries due to precipitates", 3) "Optimized heat treatment (adjust optimum particle diameter, eliminate segregations in cast microstructure, purposefully adjust possible grain boundary roughness)", 4) "Avoid cold working", are to be considered.
WO 2011/026460 PCT/DE2010/000975 9 The task of the invention is to provide a method for production of a component made from a precipitation-hardened iron-chromium alloy, by means of which the high hot strength or creep strength of a precipitation-hardened ferritic alloy can be further increased compared with the state of the art while retaining acceptable deformability at room temperature. It is also intended to provide a thermomechanically treated component/semifinished product consisting of an iron-chromium alloy, which can be used for achievement of high hot strength or creep strength while retaining acceptable deformability at room temperature. Finally, it is intended that the component/semifinished product produced in this way can be used for specific technical applications in the temperature range above 550 0 C. This task is accomplished on the one hand by a method for production of a component from an iron-chromium alloy precipitating Laves phases and/or Fe-containing particles and/or Cr-containing particles and/or Si-containing particles and/or carbides, in that a semifinished product produced from the alloy is subjected to a thermomechanical treatment, wherein in a first step the alloy is solution annealed at temperatures > the solution-annealing temperature, followed by quenching in stationary protective gas or air, moving (blown) protective gas or air or in water, in a second step mechanical working of the semifinished product in the range from 0.05 to 99% is performed and in a subsequent step Fe 2 (M, Si) or Fe 7 (M, Si) 6 Laves phases and/or Fe-containing particles and/or Cr-containing particles WO 2011/026460 PCT/DE2010/000975 10 and/or Si-containing particles and/or carbides are precipitated purposefully and in finely dispersed form by the fact that the component made from the worked semifinished product is brought to an application temperature between 550*C and 1000'C by heating at 0.1 0 C/min to 1000*C/min. This task is accomplished on the other hand by a method for production of a component from an iron-chromium alloy precipitating Laves phases and/or Fe-containing particles and/or Cr-containing particles and/or Si-containing particles and/or carbides, in that a semifinished product produced from the alloy is subjected to a thermomechanical treatment, wherein in a first step the alloy is solution annealed at temperatures 2 the solution-annealing temperature, followed by quenching in stationary protective gas or air, moving (blown) protective gas or air or in water, in a second step mechanical working of the semifinished product in the range from 0.05 to 99% is performed and in a subsequent step Fe 2 (M, Si) or Fe7 (M, Si) 6 Laves phases and/or Fe-containing particles and/or Cr-containing particles and/or Si-containing particles and/or carbides are precipitated purposefully and in finely dispersed form by the fact that the worked semifinished product is subjected for a time between t'in and t,, to a heat treatment in the temperature range between 550 0 C and 1060 0 C under protective gas or air, followed by quenching in stationary protective gas or air, moving (blown) protective gas or air or in water or for heat treatments up to 800 0 C is quenched in the oven, wherein tmn and t~m are calculated according to the following formulas: tmin = Ta * 1 0 (6740/Ta-9.216) and tmax = Ta * 1 0 (17960/Ta-15.72) where Ta = T +273.15, and wherein the desired component is made before or after this 11 heat treatment. The times tmin and tmax are expressed in minutes and the heat treatment temperature T in 0 C. According to another aspect there is provided a method for the production of a component, from an iron-chromium alloy precipitating Laves phases and/or Fe-containing particles and/or Cr-containing particles and/or Si-containing particles and/or carbides, consisting of the following composition (in % by wt.) Cr 12 - 30% Mn 0.001 - 2.5% Nb 0.1 - 2% W 0.1 - 5% Si 0.05 - 1% C 0.002 - 0.1% N 0.002 - 0.1% S max. 0.01% wherein the alloy can optionally also contain the following elements: La 0.02 - 0.3% Ti 0.01 - 0.5%, one or more of the elements Ce, Pr, Ne, Sc, Y, Zr, Hf 0.02 - 0.3% Al 0.001 - 0.5% or 2.0 - 6.0%, one or more of the elements Mg 0.0001 - 0.07% Ca 0.0001 - 0.07% P 0.002 - 0.03%, one or more of the elements Ni, Co, Cu 0.01 - 3.0% 11A B up to 0.005%, V max. 0.2% and/or S max. 0.005% 0 max. 0.01% Zn max. 0.01% Sn max. 0.01% Pb max. 0.01% Se max. 0.01% Te max. 0.01% Bi max. 0.01% Sb max. 0.01% wherein if necessary the W-content can be replaced totally or partly by at least one of the elements Mo and/or Ta, Fe remainder as well as the usual melting-related impurities, in which a semi-finished product produced from the alloy is subjected to a thermomechanical treatment, wherein in a first step the alloy is solution annealed at temperatures the solution-annealing temperature, at a temperature 1050'C for longer than 6 minutes or at a temperature 1060'C for longer than 1 minute, followed by quenching in stationary protective gas or air, moving (blown) protective gas or air, or in water, in a second step mechanical working of the semi-finished product in the range from 0.05 to 99% is performed and in a subsequent step Fe 2 (M, Si) or Fe (M, Si) 6 Laves phases and/or Fe-containing particles and/or Cr containing particles and/or Si-containing particles and/or carbides are precipitated purposefully and in finely dispersed form, by the fact that the component made from the worked semi finished product is brought to an application temperature 11B between 550 0 C and 1000 0 C by heating at 0.1 0 C/min to 1000 0 C/min. Advantageous further developments of the inventive method are to be inferred from the associated dependent claims pertaining to the method. For the first step, the following temperature ranges and times are practical for solution annealing: > 1050 0 C for longer than 6 minutes > 1060 0 C for longer than 1 minute The resulting changes of the material characteristics are explained in more detail in the course of the further description. Furthermore, the task is also accomplished by a metallic component or semifinished product consisting of the following chemical composition (in % by weight) Cr 12 - 30% Mn 0.001 - 2.5% Nb 0.1 - 2% W 0.1 - 5% Si 0.05 - 1% C 0.002 - 0.1% N 0.002 - 0.1% S max. 0.01% Fe remainder as well as the usual melting-related impurities, WO 2011/026460 PCT/DE2010/000975 12 which at the end of a thermomechanical treatment has a deformed microstructure, to the effect that Laves phase(s) is or are embedded in finely dispersed form in the microstructural dislocations of the microstructure, wherein, in a creep test with, for example, 35 MPa at 750 0 C and an elongation of at least 18%, a time to break that exceeds the time to break of a coarse grained, completely recrystallized microstructure by a factor of at least 1.5 is established in the microstructure. A comparable result is achieved for creep tests with different stresses and temperatures, wherein the temperatures for the creep test preferably lie in the range between 500 and 1000 C. Measures 1 to 4 described above will now be considered. Surprisingly it has been found in this connection that, in contrast to measure 4 "cold deformation", preworking followed by an adapted annealing treatment can bring about prolongations of the times to break of the specimen in the creep test that go more than 1.5 times, preferably more than 3 times beyond the times to break for a coarse-grained microstructure (measure 1). Furthermore, it is proposed that, for the third step - the precipitation of the Laves phase(s) - the worked semifinished product or if applicable the component made therefrom, by a combination of heating at 0.1*C/min to 1000'C/min to a heat treatment temperature between 550 0 C and 1060 0 C with subsequent heat treatment for a time between tm n and tmax at this temperature under protective gas or air, followed by quenching in stationary protective gas or air, moving (blown) protective gas or air or in water or for heat treatments up to 800 0 C is WO 2011/026460 PCT/DE2010/000975 13 quenched in the oven, after which the desired component is made if applicable, wherein tmn and tm,, are calculated according to the following formulas: tin = Ta * 1 0 (6740/Ta-9.216) and tmax = Ta 1 0 (17960/Ta-15.72) where Ta = T +273.15. In addition, the possibility exists that, for the third step the precipitation of the Laves phase(s) - the worked semifinished product or the component made therefrom is subjected for a time between tmin and tm,, to a heat treatment in the temperature range between 550 0 C and 1060 0 C under protective gas or air, followed by quenching in stationary protective gas or air, moving (blown) protective gas or air or in water or for heat treatments up to 8000C is quenched in the oven, after which the desired component is made if applicable, wherein tmin and tax are calculated according to the following formulas: tin = Ta * 1 0 (6740/Ta-9.216) and tax = Ta * 1 0 (17960/Ta-15.72) where Ta = T +273.15 and then the component that has been made is brought by heating at 0.1"C/min to 1000*C/min to an application temperature between 5500C and 10000C. According to a further concept of the invention, a semifinished product from an alloy of the following composition (in % by weight) is treated thermomechanically: Cr 12 to 30% Mn 0.001 to 2.5% Nb 0.1 to 2% W 0.1 to 5% Si 0.05 to 1% C 0.002 to 0.03% N 0.002 to 0.03% WO 2011/026460 PCT/DE2010/000975 14 S max. 0.01% Fe remainder as well as the usual melting-related impurities. With the inventive method it is possible to produce semifinished products in the form of sheets, strips, bars, forgings, pipes or wire and to make components in the most diverse forms needed for the respective application. It is of special advantage that only little or even no Fe 2 (M, Si) or Fe 7 (M, Si) 6 Laves phases and/or Fe-containing particles and/or Cr-containing particles and/or Si-containing particles and/or carbides are still present in the semifinished product after solution annealing at temperatures 2 the solution-annealing temperature, preferably 1050 0 C for longer than 6 minutes or > 1060 0 C for longer than 1 minute, followed by quenching in stationary protective gas or air, moving (blown) protective gas or air or in water, in the initial state before deformation. The working of the semifinished product can take place by hot working. Alternatively, however, the forming can also be brought about by cold working. In the first case, the semifinished product is hot-worked with a starting temperature > 1070 0 C, wherein the last 0.05 to 95% of mechanical deformation is applied between 1000* and 500 0 C, advantageously the last 0.5 to 90% between 1000 0 C and 500 0 C. In the second case, the degree of cold working of the semifinished product is 0.05 to 99%, advantageously 0.05 to 95% or 0.05 to 90%.
WO 2011/026460 PCT/DE2010/000975 15 As a further concept according to the invention, it is proposed that the mechanical working of the semifinished product be 20 to 99% and then the worked semifinished product be subjected for a time between tm., and tma to a heat treatment in the temperature range between 950 0 C and 1060 0 C under protective gas or air, followed by quenching in stationary protective gas or air, moving (blown) protective gas or air or in water, after which the desired component be made, wherein tmn and t~m are calculated according to the following formulas: tmin = Ta e 1 0 (6740/Ta-9.216) and tem = Ta e 10o(17960/Ta-15.72) where Ta = T +273.15 with ten and tma in minutes and the heat-treatment temperature T in degrees Celsius. If the already indicated alloy is used as an interconnector for a solid oxide fuel cell, then a content of 0.001 - 0.5% aluminum is advantageous. For other areas of use, such as, for example, in the reformer or heat exchanger for the fuel cell, for which no conductive oxide layer is necessary, a content of 2 to 6% aluminum is advantageous, since then a closed aluminum oxide layer can form, which once again has a much slower growth rate compared with a chromium oxide layer and additionally has much less chromium oxide volatilization than a chromium-manganese spinel. For the areas of use that neither need a conducive oxide nor have special requirements on chromium volatilization, both variants may be considered. In this connection, it is to be kept WO 2011/026460 PCT/DE2010/000975 16 in mind in particular that the processability and weldability of the alloy deteriorate with increasing aluminum content and so higher costs are incurred. Therefore, when an oxide layer consisting of a chromium oxide and a chromium-manganese spinel, adequate oxidation resistance can be assured by use of 0.001 0.5% aluminum. If greater oxidation resistance is necessary, as is assured, for example, by the formation of an aluminum oxide layer, a content of 2.0 - 6.0% aluminum is advantageous. These two alloy variants can be used, for example, as components for the exhaust-gas line of a combustion engine or for steam boilers, superheaters, turbines and other parts of a power plant. A preferred aluminum range is in particular the range from 2.5% to 5.0%, which is still characterized by good processability. In the already indicated alloy, the following elements may be additionally used individually or in combination: La 0.02 to 0.3% Ti 0.01 to 0.5% Mg 0.0001 to 0.07% Ca 0.0001 to 0.07% P 0.002 to 0.03% Ni/Co/Cu 0.01 to 3% B up to 0.005%. The contents of the elements that can be additionally introduced in the alloy may be adjusted as follows: Mg 0.0001 to 0.05%, Ca 0.0001 to 0.03%, P 0.002 to 0.03%. Furthermore, the alloy (in % by weight) may contain one or more WO 2011/026460 PCT/DE2010/000975 17 of the elements Ce, La, Pr, Ne, Sc, Y, Zr or Hf in contents of 0.02 - 0.3%. If necessary, the alloy (in % by weight) may contain one or more of the elements Ce, Pr, Ne, Sc, Y, Zr or Hf in contents of 0.02 - 0.2%. To achieve the desired effects, the Nb content is 0.3 to 1.0% and the Si content 0.15 to 0.5%. If necessary, the element tungsten may be replaced entirely or partly by at least one of the elements Mo or Ta. If necessary, the alloy may also even contain max. 0.2% V and/or max. 0.005% S. In this case the oxygen content should not be greater than 0.01%. If necessary, the alloy may also even contain max. 0.003% boron. Furthermore, the alloy should have a maximum of 0.01% of the following elements respectively: Zn, Sn, Pb, Se, Te, Bi, Sb. Components/semifinished products that on the one hand consist of the cited alloy composition and on the other hand have been produced by the inventive method may preferably be used as interconnector in a fuel cell or as material in a component, such as a reformer or a heat exchanger in an ancillary aggregate of the fuel cell. Alternatively, the possibility also exists of using the component/semifinished product produced according to the WO 2011/026460 PCT/DE2010/000975 18 inventive method or the alloy itself as a structural element in the exhaust-gas line of a combustion engine or for steam boilers, superheaters, turbines and other parts of a power plant or in the chemical process industry. By means of the inventive method, Laves phases, by virtue of the thermomechanical treatment, can be precipitated purposefully and in fine dispersion at the dislocations of the microstructure in alloys produced by melting metallurgy. The details and the advantages of the invention will be explained in more detail in the following examples. In the following, the inventive method steps will be subjected to closer examination. The first step for the thermomechanical treatment of an iron chromium alloy precipitating Laves phases and/or Fe-containing particles and/or Cr-containing particles and/or Si-containing particles and/or carbides must be annealing above the solution annealing temperature, so that the Laves phases and/or Fe containing particles and/or Cr-containing particles and/or Si containing particles and/or carbides are dissolved and are available for precipitation in the subsequent thermomechanical treatment. The solution annealing temperature is alloy-specific, but preferably lies above 1050 0 C for a period of longer than 6 minutes or above 1060 0 C for longer than 1 minute, followed by quenching in stationary protective gas or air, moving (blown) protective gas or air or in water. The exact temperature control above this solution annealing temperature is not determining for the characteristics. The annealing may be carried out in air or WO 2011/026460 PCT/DE2010/000975 19 under protective gas. It should lie under the melting temperature, preferably < 1350 0 C. Also, for cost reasons, the annealing times should preferably be < 24 hours, but may also be longer depending on performance. The solution annealing follows quenching in stationary protective gas or air, moving (blown) protective gas or air or in water, during which only little Laves phase is newly formed. In addition, care is to be taken that, especially for thicker walled components, all parts of the component reach the required minimum annealing time at the specified temperature. This is to be considered in the determination of the starting point of the annealing time. In a second step, an elevated dislocation density must be introduced into the material. Elevated dislocation densities have worked microstructure or recovered microstructure, wherein the dislocations there are arranged at small-angle grain boundaries. The second step must therefore be working, so that the dislocations are introduced into the material, which then, in the subsequent annealing treatment, ensure a homogeneous dispersion of the Laves phases and/or Fe-containing particles and/or Cr-containing particles and/or Si-containing particles and/or carbides. This deformation may be cold working, but also hot working, wherein it must be ensured during hot working that the microstructure is not already recrystallized during rolling. This is achieved by restricting the deformation range for the WO 2011/026460 PCT/DE2010/000975 20 last working and the temperature at which this is carried out. For deformations above 1000'C, the material already tends to recrystallization or recovery during working, so that the working must preferably be carried out below 1000 0 C. At temperatures below 500*C, exist in the range of the embrittlement that occurs in ferrites at 475 0 C. There this has a smaller elongation and an elevated working resistance, which makes working less advantageous and reduces the economic benefit. Precipitates smaller than a certain size are less effective (see, for example, "BUrgel", page 141). Therefore the dislocation density generated by the deformation should not be too high, since then very many precipitates are indeed formed but are too fine, and the excess dislocations can move freely and in this way the preworking becomes harmful. This means that preferably the greatest deformation is 90% for the part of the hot working 1000 0 C and 90% for the cold working. During working in the range of 20 to 99%,annealing between 950 0 C and 1050'C may cause recovery of the microstructure. Thereby the dislocation density is reduced, so that the positive effect on the dispersion of the Laves phases and/or Fe-containing particles and/or Cr-containing particles and/or Si-containing particles and/or carbides is established again. The one possibility of introducing the Laves phases and/or Fe containing particles and/or Cr-containing particles and/or Si containing particles and/or carbides into the worked material is to make the needed components from the semifinished product and then to bring the component that has been made to the WO 2011/026460 PCT/DE2010/000975 21 application temperature between 550"C and 1000 0 C by heating at 0.1 0 C/min to 1000*C/min. During heating, the Laves phases and/or Fe-containing particles and/or Cr-containing particles and/or Si-containing particles and/or carbides are precipitated as a fine dispersion in the microstructure. The fine dispersion is generated by nucleation in the lower temperature range, followed by some growth of the nuclei at the higher temperatures. Therefore the heating rate should not be slower than 1000'C/min, because otherwise the time for this process is too short. Heating rates slower than 0.1*C/min are uneconomical. A second possibility is a separate heat treatment of the material. For this purpose, the worked semifinished product/component is subjected for a time between tin and t., to a heat treatment in the temperature range between 550 0 C and 1060 0 C under protective gas or air, followed by quenching in stationary protective gas or air, moving (blown) protective gas or air or in water or for heat treatments up to 800*C is quenched in the oven, wherein tmin = Ta e 1 0 (6740/Ta-9.216) and tax = Ta * 1 0 (17960/Ta-15.72) where Ta = T +273.15, with indication of tmin and tax in minutes and heat treatment temperature T in *C. In this connection, the desired component can be made before or after this heat treatment. In the annealing steps, care is to be taken that, especially for thicker-walled semifinished products/components, all parts of the component reach the required minimum annealing time at the specified temperature. This is to be considered in determination of the starting point of the annealing time. Likewise, care is WO 2011/026460 PCT/DE2010/000975 22 to be taken that no region of the semifinished product/component exceeds the required maximum annealing time. Times shorter than tmn are not sufficient for formation of the Laves phases and/or Fe-containing particles and/or Cr-containing particles and/or Si-containing particles and/or carbides. For times longer than tax, the danger exists of too great coarseness of the precipitates, whereby the particles can no longer contribute markedly to the creep strength. For times longer than tma, the possibility exists in the upper temperature range of 550 0 C and 1060 0 C that a recovered microstructure will be formed, which certainly can still be effective. However, with increasing recovery, the dislocation density is further reduced, so that the dispersion of the precipitates becomes increasingly inhomogeneous and the positive effect on the creep strength ultimately vanishes. In addition, at the lower temperatures in the range of 550'C and 1060*C, times longer than tm,, are uneconomical. The annealing step may be carried out under protective gas (argon, hydrogen and similar atmospheres with reduced oxygen partial pressure). For economic reasons, the quenching step is carried out in stationary protective gas or air, moving (blown) protective gas or air or in water, while oven quenching should be avoided in particular for temperatures above 800 0 C but is also possible at temperatures < 800 0 C. The oxidation resistance and the thermal expansion coefficient of the material are determined via the chromium content. The oxidation resistance of the material is based on the formation of a closed chromium oxide layer. Below 12%, iron-containing WO 2011/026460 PCT/DE2010/000975 23 oxides, which impair the oxidation resistance, are formed increasingly, especially at higher operating temperatures. The chromium content is therefore adjusted to 12%. Above 30% chromium, the processability of the material and its usability are impaired by increased formation of embrittling phases, especially the sigma phase. The chromium content is therefore limited to 30%. The expansion coefficient decreases with increasing chromium content. Especially for use in a fuel cell, therefore, the expansion coefficient can be adjusted in a range that matches the ceramics in the fuel cell. These are chromium contents around 22 to 23%. For other applications, however, for example for reformers or in power plants, this restriction does not exist. The addition of manganese brings about formation of a chromium manganese spinel on the chromium oxide layer, which is formed on the material at low aluminum contents below 2%. This chromium manganese spinel reduces the chromium volatilization and improves the contact resistance. A manganese content of at least 0.001% is necessary for this. More than 2.5% manganese impairs the oxidation resistance by formation of a very thick chromium manganese spinel layer. Niobium, molybdenum, tungsten or tantalum can participate in the formation of precipitates in iron-containing alloys, such as, for example, carbides and/or the M in the Fe 2 (M, Si) or Fe7(M, Si) 6 Laves phases. Molybdenum, tungsten or tantalum are also good solid-solution hardeners and thus contribute to the improvement of the creep strength. In this connection the lower limit is determined in each case by the fact that a certain content must WO 2011/026460 PCT/DE2010/000975 24 be present in order to be effective, while the upper limit is determined by the processability. Thus the preferred ranges are, for Nb 0.1 - 2% W: 0.1 - 5% W may also be replaced entirely or partly by Mo and/or Ta: 0.1 5%. Silicon can participate in the formation of precipitates in iron-containing alloys, for example in the Fe 2 (M, Si) or Fe,(M, Si) 6 Laves phases. It favors the increased precipitation and stability of these Laves phases and in this way contributes to the creep strength. During formation of the Laves phases, it is completely bound in these. Thus the formation of a silicon dioxide layer no longer takes place under the chromium oxide layer. At the same time, the incorporation of M in the oxide layer is reduced, whereby the negative influence of M on the oxidation resistance is prevented. At least 0.05% Si must be present for the desired effect to occur. If the content of Si is too high, the negative effect of the Si may reappear. The Si content is therefore limited to 1%. Aluminum in contents below 1% impairs the oxidation resistance, since it leads to internal oxidation. However, an aluminum content higher than 1% leads to formation, under the chromium oxide layer, of an aluminum oxide layer, which is not electrically conductive and thus reduces the contact resistance. Therefore the aluminum content is limited to 0.5% when a chromium oxide former is desired or its oxidation resistance is sufficient. An example of this is, for example, for the application as interconnector plate. However, a certain aluminum WO 2011/026460 PCT/DE2010/000975 25 content of at least 0.001% is necessary for deoxidation of the melt. If no conductive oxide is necessary and at the same time the requirement of much higher oxidation resistance than is given by a chromium oxide layer is still required, the alloy may form a closed aluminum oxide layer by a content of aluminum of at least 2% (DE 101 20 561). Aluminum contents above 6.0% lead to processing problems and thus to increased costs. Carbon leads to carbide precipitates and thus contributes to the creep strength. The carbon content should be < 0.1%, in order not to impair the processability. However, it should be > 0.002%, so that an effect can occur. The nitrogen content should be 0.1% maximum, in order to avoid formation of nitrides, which impair the processability. It should be higher than 0.002%, in order to assure the processability of the material. The contents of sulfur should be made as low as possible, since this interface-active element impairs the oxidation resistance. A maximum of 0.01% S is therefore stipulated. Oxygen-affine elements such as Ce, La, Pr, Ne, Sc, Y, Zr, Hf improve the oxidation resistance by reducing the oxide growth and improving the adherence of the oxide layer. A minimum content of 0.02% of one or more of the elements Ce, La, Pr, Ne, Sc, Y, Zr, Hf is practical in order to obtain the oxidation resistance-increasing effect of the Y. For cost reasons, the upper limit is set at 0.3% by weight. As with every oxygen-affine element, titanium is bound in the WO 2011/026460 PCT/DE2010/000975 26 oxide layer during oxidation. In addition, it also causes internal oxidation. However, the resulting oxides are so small and finely dispersed that they cause hardening of the surface and thus prevent swelling of the oxide layer and inclusion of metallic zones during the oxidation (see DE 10 2006 007 598 Al). These swellings are unfavorable, since the resulting cracks cause an increase of the oxidation rate. Thus Ti contributes to the improvement of the oxidation resistance. For effectiveness of the Ti content, at least 0.01% Ti must be present, but not more than 0.5%, since this does not improve the effect further but increases the costs. The content of phosphorus should be lower than 0.030%, since this interface-active element impairs the oxidation resistance. Too low P content increases the costs. The P content is therefore 0.002%. The contents of magnesium and calcium are adjusted in the spread ranges of 0.0001 to 0.05% by weight and 0.0001 to 0.03% by weight respectively. It has been found that cobalt contents of 3% and higher impair the oxidation resistance. For cost reasons, the lower limit is set at 0.01% by weight. For nickel and copper, the same applies as for cobalt. Boron is limited to max. 0.005%, since this element reduces the oxidation resistance. The subject matter of the invention will now be explained in more details on the basis of exemplary embodiments.
WO 2011/026460 PCT/DE2010/000975 27 The analyses of the batches used for the following examples are presented in Table 1. These batches were melted in the arc furnace in an amount of approximately 30 metric tons, thereafter cast in a ladle and subjected to a decarburization and deoxidation treatment as well as to a vacuum treatment in a VOD system and cast to ingots. These were then hot-rolled and, depending on final thickness, cold-rolled with intermediate annealing steps. After the hot-rolling, the oxide layer was removed by pickling. A material with an analysis as indicated in Table 1 precipitates mainly Fe 2 (M, Si) or Fe7 (M, Si) 6 Laves phases and, in much smaller contents, carbides. Example 1 In this example, material from the batch 161061 listed in Table 1 was hot-rolled to 12 mm thick sheet after solution annealing above 1070 0 C for a period of longer than 7 minutes followed by quenching in stationary air, wherein the mechanical working was begun with a start temperature > 1070 0 C and the last 78% of mechanical deformation was applied by rolling between 500 0 C and 1000*C. Figure 1 shows the typical appearance of a microstructure deformed in this way. In the microsections etched by means of electrolytic etching with oxalic acid, it can be clearly seen that only little Laves phase has been precipitated in microscopically visible form.
WO 2011/026460 PCT/DE2010/000975 28 When the material formed in this way is then annealed at 1075 0 C for 20 minutes with quenching in stationary air, a microstructure is obtained with only few precipitates of Laves phase and a grain size of approximately 137 pm (Figure 2), which is a typical coarse-grained microstructure. When a creep test as described above is performed on this material with an initial stress of 35 MPa at a temperature of 750'C, the specimen breaks after 12.8 hours at an elongation A of 69.8%. (Table 2). At room temperature, the material annealed in this way has an elongation of 35%, which is a very good value for a ferrite. If, in contrast, from the hot-rolled material, which is synonymous with preworking, a specimen for a creep test is made as simulation for a component and this is then heated at approximately 60*C/minute to an application temperature of 750 0 C and then a creep test is performed with an initial stress of 35 MPa at a temperature of 750'C, the specimen surprisingly breaks only after 255 hours at an elongation A of 29%, which means a prolongation of the time to break by a factor of 20. Making the component is very easily possible, since the hot-worked condition, as was described above, has an elongation of 19% in the tension test at room temperature, which is a good value and makes the material readily processable. This example clearly shows that the microstructure with the preworking and the coarse-grained microstructure is superior with respect to time to break or creep strength, which contradicts the state of the art as described in "Birgel" pages 196 to 199 Table 3.7.
WO 2011/026460 PCT/DE2010/000975 29 Example 2 In this example, annealing steps were performed on the hot rolled material from Example 1 for 20 minutes each between 600 0 C and 1000 0 C or for some temperatures also for 240 or 1440 minutes (see Table 3 for tma and tax according to Equation 1 and 2) in air, followed by quenching in stationary air. After the heat treatment, specimens were made from the sheet and then the creep test was performed with a stress of 35 MPa at 7500C as described above. The results are compiled in Table 3. After 20 minutes at 1000*C, a time to break of only 10.4 hours is reached at an elongation A of 79.5%. After 20 minutes at 600 0 C to 9500C, a time to break of longer than 100 hours, at least 7 times longer, is achieved at an elongation A of greater than 22.7%. The longest time to break for annealing steps of 20 minutes is achieved at 8500C, with 564 hours. The longest time to break for annealing steps of 240 minutes is achieved at 8000C, with 396 hours. After 1440 minutes at 7000C, a time to break of 645 hours is even reached. Figure 3 shows the microstructure after the various annealing steps for 20 minutes. The microstructures in Figure 3 are not globularly recrystallized. Up to 8500C (the maximum of the time to break), the microstructure has the typical appearance of a deformed microstructure. Starting from approximately 9000C, recovery can be clearly recognized, but this means that the dislocation density is still increased compared with a globularly recrystallized microstructure. In a recovered microstructure, the dislocations have become partly reordered at small-angle grain boundaries. This has an effect similar to that of WO 2011/026460 PCT/DE2010/000975 30 preworking. In the microsections etched by means of electrolytic etching with oxalic acid, it can be clearly recognized that the Laves phase is precipitated in microscopically visible form starting from approximately 750*C, wherein it is precipitated increasingly more densely and more homogeneously up to 850 0 C (the maximum of the time to break). From approximately 9000C on, small-angle grain boundaries or grain boundaries can also be recognized markedly besides the precipitates in the grain, thus assuring jaggedness of the small-angle grain boundaries or grain boundaries, which corresponds to measure 2 for increasing the creep strength (see above). At 1000 0 C, very large grains have recognizably formed due to advancing recovery, such that the dislocation density is greatly reduced and so no further increase of the time to break occurs. The maximum of the time to break occurs in the deformed microstructure with dense homogeneously precipitated Laves phases. At room temperature, the sheet annealed for 20 minutes at all temperatures between 600*C and 9500C has an elongation of at least 13%, which is still to be regarded as satisfactory for a ferritic alloy and makes the material processable. The elongation is smallest in the range of 700'C to 8000C and is improved at the lower or higher annealing temperatures respectively, because at the lower temperatures Laves phase is certainly already precipitated but is not yet microscopically visible and therefore has a smaller proportion by volume, but in return is very finely dispersed. At the higher temperatures, a larger proportion by volume is precipitated, but in return is somewhat coarser and recognizable at the small-angle grain boundaries and grain boundaries.
WO 2011/026460 PCT/DE2010/000975 31 The annealing at 10000C for an annealing time of 20 minutes exceeds tmax = 19.6 minutes. Thus it is not in the range of the invention and is used as reference. Also, the time to break is only 10.4 hours. The annealing time of 20 minutes at the temperatures between 6000C and 950 0 C lies in the inventive range between tmn and tma. Accordingly, the time to break was clearly prolonged according to the invention by more than a factor of 7 compared with the coarse-grained, globularly recrystallized condition from Example 1, which is obtained after annealing at 1075*C/20 minutes followed by quenching in stationary air. Example 3 In this example, material from the batch 161061 listed in Table 1 was hot-rolled to 12 mm thick sheet after solution annealing above 10700C for a period of longer than 7 minutes followed by quenching in stationary air, wherein the working was begun with a start temperature > 10700C and the last 60% of mechanical working was applied by rolling between 10000C and 500 0 C. If the sheet worked in this way is then annealed industrially in the continuous furnace at 9200C for 28 minutes in air and quenched in stationary air, a tension specimen made from this material has a time to break of 391 hours at an elongation A of 38% (Table 4) in the creep test with an initial stress of 35 MPa at a temperature of 750*C. The microstructure is not globularly recrystallized but instead is recovered. It has precipitates in the grain and at the small-angle grain boundaries or grain boundaries (Figure 4). The time to break is 30 times the time achieved in Example 1 after annealing at 10750C for 20 minutes with a globularly recrystallized coarse-grained microstructure WO 2011/026460 PCT/DE2010/000975 32 with a grain size of 137 jim. The annealing at 9200C for an annealing time of 28 minutes lies in the inventive range between tmn = 0.32 minutes and tmax = 162.6 minutes. At room temperature, the sheet treated in this way has a very good elongation of 18%, an offset yield strength of 475 MPa and a tensile strength of 655 MPa (see Table 4), which makes the material readily workable. Example 4 In this example, material from batch 161061 and batch 161995 was cold-rolled to 1.5 mm thick sheet after solution annealing at above 1070 0 C for a period longer than 7 minutes followed by quenching in blown air and hot rolling as well as removal of the oxide layer, wherein cold working of 53% was applied. Subsequently annealing at 10500C was carried out for 3.4 minutes under protective gas in the continuous furnace with subsequent quenching in the cold stream of protective gas. Thereafter both batch 161061 (Figure 5) and batch 161995 exhibit a recovered microstructure with elongated grains (Figure 7) and precipitation of Laves phase, although much less than recognizable in Figure 4. Thereafter part of the material was annealed once again at 10500C for 20 minutes under air with subsequent quenching in stationary air. After this, both batches were globularly recrystallized, batch 161061 with a grain size of 134 gm (Figure 6) and batch 161995 with a grain size of 139 gm. Only slight precipitated Laves phase can still be found. Tables 5a and 5b show the results of the creep tests and of the WO 2011/026460 PCT/DE2010/000975 33 tension tests at room temperature. After the annealing at 1050*C for 3.4 minutes, batch 161061 has a time to break of 25.9 hours at an elongation A of 50% in a creep test at 750 0 C with an initial load of 35 MPa, and after additional annealing at 1050 0 C for 20 minutes, which produces very coarse grain, a time to break of only one third, 7.9 hours, at an elongation A of 83%. Similarly, batch 161995 has a time to break of 33.5 hours 89% in a creep test at 750 0 C with an initial load of 35 MPa, and after additional annealing at 10750C for 20 minutes, which produces very coarse grain, a time to break of only one third, 7.9 hours, at an elongation A of 92%. The elongation of 28% in the tension test at room temperature for batch 161061 for 10500C and 3.4 minutes of annealing time and of 26% for batch 161995 is very good for a ferrite, which makes the material very readily workable. For the coarse-grained structure, it is even higher, with 31% for batch 161061 and 29% for batch 161995. This shows the influence of the annealing time at temperatures around 10500C. In short-time annealing steps of a few minutes, dislocations (deformation) and adequate Laves phase are present in the material, which in this example has the consequence of a 3 to 4 times longer time to break in the creep test. For longer annealing steps, the Laves phase is sufficiently dissolved, as batch 161061 shows, and the microstructure recrystallizes globularly with correspondingly short times to break in the creep test. The annealing at 1050'C for 20 minutes lies with an annealing time of 20 minutes above tma = 6.0 minutes. Thus it does not fall within the range of the invention and is used as reference, WO 2011/026460 PCT/DE2010/000975 34 just as the annealing at 1075 0 C for 20 minutes. The annealing at 1050'C for 3.4 minutes lies with an annealing time of 3.4 minutes in the inventive range between tmn = 0.32 minutes and tmx = 6.0 minutes and according to the invention exhibits a clearly improved time to break in the creep test. Example 5 In this example, material from the batch 161061 was hot-rolled to 12 mm thick sheet after solution annealing above 1070*C for a period of longer than 7 minutes followed by quenching in stationary air, wherein the working was begun with a start temperature > 1070*C and the last 70% of mechanical deformation was applied by rolling between 1000 0 C and 500 0 C. When the material worked in this way is then subjected to solution annealing at 1075 0 C for 22 minutes with quenching in stationary air, a very coarse-grained microstructure is obtained with only few precipitates of Laves-phase and a grain size of approximately 134 to 162 pm (Figure 9). When a creep test is performed on this material with an initial stress of 40 MPa at a temperature of 700'C, the specimen breaks after 228 hours at an elongation A of 51%. (Table 6) When the creep test is performed at 60 MPa, the specimen breaks after 8.1 hours, at an elongation A of 43%. At room temperature, the material annealed in this way has an elongation of 35%, which is a very good value for a ferrite. If the material solution annealed at 1075 0 C for 22 minutes is additionally subjected to annealing for 4 hours at 700 0 C with subsequent quenching in stationary air, Laves phase dispersed in WO 2011/026460 PCT/DE2010/000975 35 the microstructure is precipitated. (Figure 10). When the creep test is then performed at 700*C with an initial stress of 40 MPa, the specimen already breaks after 104 hours, at an elongation A of 72.6%, therefore a much shorter time than after the solution annealing at 1075 0 C for 22 minutes. When the creep test is performed at 60 MPa, the specimen breaks after 6.3 hours, at an elongation A of 63%, therefore also after substantially shorter time than after the solution annealing at 1075*C for 22 minutes. This is the proof that the precipitation of the Laves phase(s) must take place in a microstructure with elevated dislocation density, therefore in a worked or recovered microstructure, in order to achieve prolongation of the time to break. Precipitation in a solution annealed microstructure has exactly the opposite effect, namely shortening of the time to break. The cause of this is the more homogeneous dispersion of very fine precipitates in the case of precipitation in a microstructure with elevated dislocation density, or in other words a deformed or recovered microstructure, in comparison with precipitation in a dislocation-poor coarse-grained microstructure. Example 6 In this example, as in Example 2, annealing steps were performed on the hot-rolled material from Example 1 for 20 minutes each between 750 0 C and 1000 0 C or for some temperatures also for 120 minutes, 240 minutes, 480 minutes, 960 minutes, 1440 minutes or 5760 minutes (see Table 7 for tmin and tm,, according to Equation 1 and 2) in air, followed by quenching in stationary air. After the heat treatment, specimens were made from the sheets and then WO 2011/026460 PCT/DE2010/000975 36 the creep test was performed with a stress of 40 MPa at 750 0 C as described above. The higher stress in comparison with Example 2 was chosen for shortening of the test time. The objective was to find heat-treatment times suitable for the annealing steps. The results are compiled in Table 7. After 20 minutes at 10000C, a time to break of only 8.8 hours is reached at an elongation A of 78.7%. In Example 2, after 20 minutes at 10000C and a creep test at 7500C and 35 MPa, a time to break comparable with that after solution annealing at 1075 0 C for 20 minutes with quenching in stationary air was reached, and so this value can be taken as reference for the time to break of the solution annealed condition. Inventive variants should also exceed this break time once again by a factor of at least 1.5. After 20 minutes at 7500C to 9000C, a time to break of longer than 100 hours, at least 10 times longer, is achieved at an elongation A of greater than 27%. The longest time to break for annealing steps of 20 minutes is achieved at 85000 with 296 hours. The longest time to break for annealing steps of 120 minutes is achieved at 8000C with 227 hours. The longest time to break for annealing steps of 240 minutes is achieved at 75000 with 182 hours, but in this connection no value exists for 7000C. The longest time to break for annealing steps of 480 minutes is achieved at 8000C with 169 hours. For 960 minutes, only one time to break was determined, for 7500C, with a value of 139 hours at an elongation of 24.2%. After 1440 minutes and 5760 minutes at 7500C and 8000C, only times to break clearly shorter than the maximum times to break achieved at these temperatures are still achieved. At 8000C, for example, the time to break after a treatment time of 480 minutes drops from 169 WO 2011/026460 PCT/DE2010/000975 37 hours to a value of 46 hours after a treatment time of 1440 minutes, to a value of 17.5 hours after a treatment time of 5760 minutes, although this is still in the inventive range. A further prolongation of the treatment time should shorten the time to break further, so that tmx of 7059 minutes is logically somewhat longer than the time of 5750 minutes. All elongations for the heat treatment temperatures from 750 to 900'C and times from 20 minutes to 5760 minutes lie between 24.2% and 43% and therefore are larger than 18%, as required, in order to avoid brittle failure. For the microstructure after 20 minutes of annealing, what was said in Example 2 is applicable, since the annealing steps are the same. Even at the higher stress of 40 MPa in the creep test, the maximum of the time to break occurs in the deformed microstructure, with dense homogeneously precipitated Laves phase. At room temperature, the sheet annealed for 20 minutes at all temperatures between 600 0 C and 900 0 C in Example 2 has an elongation of at least 13%, which is still to be regarded as satisfactory for a ferritic alloy and makes the material processable. Example 7 In this example the 1.5 mm thick material of batch 161995, which was annealed after cold working of 53% at 1050'C for 3.4 minutes under protective gas in the continuous furnace with subsequent quenching in the stream of cold protective gas, was used once again. In the same manner, 2.5 mm thick material from batch 161995 was produced by annealing it, after cold working of 40%, at 1050 0 C for 2.8 minutes under protective gas in the continuous WO 2011/026460 PCT/DE2010/000975 38 furnace with subsequent quenching in the stream of cold protective gas. Even the 2.5 mm thick material then exhibits a recovered microstructure with elongated grains (Figure 11), just as the material from Example 4 in Figure 7, and precipitation of Laves phase, albeit clearly less than recognizable in Figure 4. Part of the material was then annealed once more at 10500C for 10 minutes under air with subsequent quenching in stationary air. After this the material had globularly recrystallized structure, with a grain size of 108 pm. Only little precipitated Laves phase is still to be found. The material with 1050 0 C/2.8 minutes and the material with 10750C/10 minutes as the last heat treatment was then rolled with degrees of working between 2.8 and 40%. Thereafter creep tests were performed at 750*C and 35 MPa and tension tests were performed at room temperature. The results are summarized in Table 8. After the annealing at 10500C for 3.4 minutes, batch 161995, in a creep test at 750"C with an initial load of 35 MPa, had a time to break of 33.5 hours at an elongation A of 89% and, after the additional annealing at 1050 0 C for 10 minutes, which produces very coarse grain, it had a time to break amounting to only one third, 10.8 hours, at an elongation A of 50.4%. If, from the material worked after 1050 0 C/2.8 minutes, a specimen for a creep test is made as simulation for a component and this is then heated at approximately 60*C/minute to an application temperature of 7500C and then a creep test is performed with an initial stress of 35 MPa at a temperature of 7500C, the elongation at break for degrees of working between 5 and 40% drops to values around the 10 hours with elongations at break greater than 45%.
WO 2011/026460 PCT/DE2010/000975 39 If, in contrast, from the material formed after 1050*C/10 minutes, a specimen for a creep test is made as simulation for a component and this is then heated at approximately 60*C/minute to an application temperature of 750'C and then a creep test is performed with an initial stress of 35 MPa at a temperature of 750 0 C, the elongation at break for degrees of working between 2.9 and 40% increases to values between 49 and 137 hours, which means an increase by more than a factor of 4 in the time to break compared with the material worked after 1050 0 C/2.8 minutes, wherein a maximum occurs at 10% and the elongations at break lie between 18.9 and 60%. From 10% degree of working on, however, the elongation at break in the tension test at room temperature becomes smaller than 8%, so that the material is increasingly more poorly processable. In other words, preferred degrees of forming for cold shaping lie between 0.05 and 10%. Thus this example shows that the increase of the elongation at break after working does not occur if the annealing before working was carried out at too low temperatures or for too short times. (Here at 1050 0 C for 2.8 minutes) The increase of the elongation at break occurred after annealing was carried out at > 1050 0 C for > 6 minutes (here at 1075 0 C for 7 minutes). The titles/descriptions of the tables/figures are reproduced as follows: Table 1 Composition of the investigated alloy (all values in % by weight) WO 2011/026460 PCT/DE2010/000975 40 Table 2 Results of the creep tests at 7500C with 35 MPa and of the tension tests at room temperature for the hot rolling and the heat treatments in Example 1 for a 12 mm thick sheet. (R: reference according to the state of the art, I: according to the invention) Table 3 Results of the creep tests at 750'C with 35 MPa and of the tension tests at room temperature for the hot rolling from Example 1 and the heat treatment from Example 2 for a 12 mm thick sheet. (R: reference according to the state of the art, I: according to the invention) Table 4 Results of the creep tests at 750*C with 35 MPa and of the tension tests at room temperature for Example 3 for a 12 mm thick sheet. (R: reference according to the state of the art, I: according to the invention) Table 5 Results of the creep tests at 750'C with 35 MPa and of the tension tests at room temperature for Example 4 for a 1.5 mm thick strip. (R: reference according to the state of the art, I: according to the invention) Table 6 Results of the creep tests at 7000C and of the tension tests at room temperature for Example 5 on 12 mm thick sheet (R: reference according to the state of the art, I: according to the invention) Table 7 Results of the creep tests at 7500C with 40 MPa and of the tension tests at room temperature for the hot rolling and the heat treatments in Example 6 for a 12 WO 2011/026460 PCT/DE2010/000975 41 mm thick sheet. (R: reference according to the state of the art, I: according to the invention) Table 8 Results of the creep tests at 750 0 C with 35 MPa and of the tension tests at room temperature for Example 7 for 1.5 mm to 2.5 mm thick strip from batch 161995. (R: reference according to the state of the art, I: according to the invention) Figure 1 Microstructure of the hot-worked material in Example 1 Figure 2 Microstructure of the hot-worked material in Example 1 after annealing at 1075*C for 20 minutes and quenching in stationary air, grain size 137 gm. Figure 3 Microstructure of the material in Example 2 after annealing between 600 0 C and 1000 0 C for 20 minutes in each case and quenching in stationary air. Figure 4 Microstructure of the material in Example 3 after annealing at 920*C in the continuous furnace in air with subsequent quenching in stationary air for 20 minutes in each case and quenching in stationary air. (etching with V2A pickling fluid) Figure 5 Microstructure of batch 161061 in Example 4 after annealing at 1050"C/3.4 minutes under protective gas in the continuous furnace with quenching in the stream of cold protective gas. Figure 6 Microstructure of batch 161061 in Example 4 after WO 2011/026460 PCT/DE2010/000975 42 annealing at 1050*C/3.4 minutes under protective gas in the continuous furnace with quenching in the stream of cold protective gas and annealing at 1050*C/20 minutes under air with subsequent quenching in stationary air, grain size 134 pim (etching with V2A pickling fluid) Figure 7 Microstructure of batch 161995 in Example 5 after annealing at 1050*C/3.4 minutes under protective gas in the continuous furnace with quenching in the stream of cold protective gas. Figure 8 Microstructure of batch 161995 in Example 5 after annealing at 1050'C/3.4 minutes under protective gas in the continuous furnace with quenching in the stream of cold protective gas and annealing at 1075*C/20 minutes under air with subsequent quenching in stationary air, grain size 139 pim Figure 9 Microstructure of the hot-worked material in Example 5 after annealing at 1075'C for 22 minutes and quenching in stationary air, grain size 134 pm to 162 pm Figure 10 Microstructure of the hot-worked material in Example 5 after annealing at 10750C for 22 minutes followed by quenching in stationary air and subsequent annealing at 7000C for 4 hours followed by quenching in stationary air. Grain size 136 pm. Figure 11 Microstructure of batch 161995 in Example 7 after WO 2011/026460 PCT/DE2010/000975 43 annealing at 1050 0 C for 2.8 minutes under protective gas in the continuous furnace with quenching in the stream of cold protective gas. Figure 12 Microstructure of batch 161995 in Example 7 after annealing at 1050 0 C for 2.8 minutes under protective gas in the continuous furnace with quenching in the stream of cold protective gas, followed by annealing at 1075*C for 10 minutes under air with subsequent quenching in stationary air. Grain size 108 pm.
WO 2011/026460 PCT/DE2010/000975 44 Table 1 Batch Batch Element 161061 161995 C 0.007 0.009 S < 0.002 < 0.002 N 0.015 0.018 Cr 22.9 22.6 Ni 0.30 0.22 Mn 0.43 0.43 Si 0.21 0.24 Mo 0.02 0.02 Ti 0.07 0.06 Nb 0.51 0.49 Cu 0.02 0.02 Fe Remainder Remainder P 0.014 0.017 Al 0.02% 0.019 Mg 0.0006 < 0.01 Pb < 0.001 < 0.001 Sn < 0.01 < 0.01 Ca 0.0002 0.01 V 0.05 0.02 Zr < 0.01 < 0.01 W 1.94 1.97 Co 0.04 0.02 La 0.08 0.05 Ce < 0.01 0 0.004 Lc) Cuf 0) 0C E C 0 - a.y .0 - 0) 0 I-R LO N00 p0) U) -0 U) 0) C' ) CL M 00 0oL) 0 CD Lo IA 04 - C0 C') -) S0 0.0 -E U) (D M N 0 A m 00 EE
CC
0) Cu 0 I - - - m - - I CV 0 c. T- 0C CU N I~ N CC N 0) N CO c 0 L)L 11 V0) 0r 0 V 00 CDi Cin u ) r- mn m~ v w 0D0 CU NO Lo NN (D 0R N CC E 8 0 =. cO CD 0 LO M V co 0~ 0 m COj cm0 vNOi ',n NI- N N C' I E - N NNNNNNN 0 0 0 000 0 a00C N I-I-I- 0 0 0 w N CD 0- 0-~ 0 (x~I a I N w EucE 0 I T 0 CCV E0E wCD 0 m E Cu x w E 0 0 .) v C) Cv) ,*- 0 C 0 C *0 oE F~r7- nixC r C D Cu 0 L E) a E ( 2_ CLC 0 cmCf CW) w 0 C) CC) 0 CL r CY) (j~ Q D Lo 04 0) C' EC uo E Eu = a)V a) CDC d) a) c I 0 ;T a) wE
-(D
) 00 EE 0ME 0 (D 2 0 a a. m 0n C -= 0 0C -- 0 o C DI 0 C)C U) 0 C') mUC) c a) E E , - E or )m 2 02 0o m* 00 C_ Dc a,.EVL0 0 0 C a) a)C____ _____ - . , Ep I of I 0p 0D 0 co CF c0 0 0 (A D LO~ Ic C ~) I LO I - C 0 0 CY) CL Cv) CV D0 (3 w 0 0 :3 - C 0o C CC 0 0 D .m v It) En 0D 0- 0 LO c 0 0) c (lr) Cf) I-CU 0 w a) 2 oB Lo WUN 0 0 U) CC a) C 0 Co a, a, m m 0 0 m CD U o 0 o w 0 0 - 0 -O UT)-c a) '- Co a,.a (D N C-4 Co ata CV a.00 0 0D cm m .
C D Co 0 o Coc coC a,~~ 0~ o C C3o T- *E 0 0 co cc E N CN v CN CN 4 0 U) U') 0 C) 0 C C13 r- I- U r'- V) 0 0 CD t-. Y O 0 li 1~ o U, C: 0Q mDCV-rcoE E 2 0 E _) 0 LD :5 C4 0 CO) ) LO 0 .9 CC 0 m r-C4 C 0 E cc 0V C) 03 C) D N o 0 *~0 a - - - - (NJ CD 0) LO L m C -(CD 0D ) 1- If) N 0) 0) ~ ~ ~ C C-~ ~ - N N N~ E 204 04 C1 2o C)~ 0 0C~) rL .l- 00 c '-) r- . r-I CO C)
EE
C 0 CU 0 CLC E_ E (D 0Cm 0 -25 4) C g w - - - - - - E c i 0 0 0 *C CU C1 o C% L l c. CVU Nl qT LO t Ul) 0 w C~ )N ~Lo Co 0 a) -1 M0 0 14 c 0. GoJ in Nn CO *' CV - m co CD co C) C *-U) C~ ") 0 0 0 ) d).- CD -C' E CM ( 720 - o a 8 a) 10 - E LO N O~0
E)I
C) C0 2 C .14()0 u) C u) u) C) w ) w ) CUo 0: O 0.CL c - CO 00 t'- CO C0 1~- I'- CO r- OD E w C 0 M c CD co C 0) 1-L N ' OD CD IT 0 2 < 0.U E E a) cu 0 a-O. co co Lo) Lo r4 o C Oh - CD V~ M' 0 CDt)ODco c - CD
(D
0 ci CD 0U 0)CD 'tO CD CM u* a)0) CL m t CD co - co 0 0 -o ro - C - CDC ) ~ 0.0 L() LO- W 0 c (D S00 *0 ,o- (.4 .j CD S 0 0* 0 CDi CD CD CD a) E i'N .JJ E m C) a - 0M 0 U L 0 E 0 (D a)C F PCuo CL 0 CY U* InC U) ~ 0 )Il - C A t - r- w c m (n w 0. 0 =- N) - ' U CL I- a, C O m c n c mo U) L D r- r 0 c CO N W) a, a U) ) U)C D C 0 0 U) N ) CD - CD ) N aD N C. 0 Co E o 0. OW (D 00 0 0
E--
55A Throughout this specification and the claims which follow, unless the context requires otherwise, the word "comprise", and variations such as "comprises" or "comprising", will be understood to imply the inclusion of a stated integer or step or group of integers or steps but not the exclusion of any other integer or step or group of integers or steps. The reference in this specification to any prior publication (or information derived from it), or to any matter which is known, is not, and should not be taken as, an acknowledgement or admission or any form of suggestion that that prior publication (or information derived from it) or known matter forms part of the common general knowledge in the field of endeavour to which this specification relates.
Claims (23)
1. Method for the production of a component, from an iron chromium alloy precipitating Laves phases and/or Fe containing particles and/or Cr-containing particles and/or Si-containing particles and/or carbides, consisting of the following composition (in % by wt.) Cr 12 - 30% Mn 0.001 - 2.5% Nb 0.1 - 2% W 0.1 - 5% Si 0.05 - 1% C 0.002 - 0.1% N 0.002 - 0.1% S max. 0.01% wherein the alloy can optionally also contain the following elements: La 0.02 - 0.3% Ti 0.01 - 0.5%, one or more of the elements Ce, Pr, Ne, Sc, Y, Zr, Hf 0.02 - 0.3% Al 0.001 - 0.5% or 2.0 - 6.0%, one or more of the elements Mg 0.0001 - 0.07% Ca 0.0001 - 0.07% P 0.002 - 0.03%, one or more of the elements Ni, Co, Cu 0.01 - 3.0% B up to 0.005%, V max. 0.2% and/or S max. 0.005% 57 0 max. 0.01% Zn max. 0.01% Sn max. 0.01% Pb max. 0.01% Se max. 0.01% Te max. 0.01% Bi max. 0.01% Sb max. 0.01% wherein if necessary the W-content can be replaced totally or partly by at least one of the elements Mo and/or Ta, Fe remainder as well as the usual melting-related impurities, in which a semi-finished product produced from the alloy is subjected to a thermomechanical treatment, wherein in a first step the alloy is solution annealed at temperatures the solution-annealing temperature, at a temperature 1050 C for longer than 6 minutes or at a temperature 1060'C for longer than 1 minute, followed by quenching in stationary protective gas or air, moving (blown) protective gas or air, or in water, in a second step mechanical working of the semi-finished product in the range from 0.05 to 99% is performed and in a subsequent step Fe 2 (M, Si) or Fe 7 (M, Si) 6 Laves phases and/or Fe-containing particles and/or Cr containing particles and/or Si-containing particles and/or carbides are precipitated purposefully and in finely dispersed form, by the fact that the component made from the worked semi-finished product is brought to an application temperature between 550 0 C and 1000 0 C by heating at 0.1 0 C/min to 1000 0 C/min. 58
2. Method according to claim 1, wherein the semi-finished product has a mechanical deformability at room temperature of > 13%, measured as plastic elongation in the tension test.
3. Method according to claim 1 or 2, wherein only little or even no Fe 2 (M, Si) or Fe 7 (M, Si) 6 Laves phases and/or Fe containing particles and/or Cr-containing particles and/or Si-containing particles and/or carbides are still present in the alloy after solution annealing at temperatures the solution-annealing temperature, or 1050 0 C for longer than 6 minutes, or 1060 0 C for longer than 1 minute, followed by quenching in stationary protective gas or air, moving (blown) protective gas or air or in water, in the initial state before deformation in the semi-finished product.
4. Method according to any one of claims 1 to 3, wherein the working of the semi-finished product takes place by hot working.
5. Method according to any one of claims 1 to 4, wherein the hot working of the semi-finished product begins with a starting temperature >1070 0 C, wherein the last 0.05 to 90% of mechanical deformation is applied between 1000'C and 500 C.
6. Method according to any one of claims 1 to 4, wherein hot working of the semi-finished product begins with a starting temperature >1070 0 C, wherein the last 0.05 to 95% of mechanical deformation is applied between 1000 0 C and 5000C. 59
7. Method according to any one of claims 1 to 6, wherein the hot working of the semi-finished product begins with a starting temperature >1070 0 C, wherein the last 0.5 to 90% of mechanical deformation is applied between 1000'C and 500 OC.
8. Method according to any one of claims 1 to 7, wherein the hot working of the semi-finished product is followed by cold working.
9. Method according to any one of claims 1 to 7, wherein the working of the semi-finished product is carried out by cold working.
10. Method according to claim 9, wherein the degree of cold working of the semi-finished product is 0.05 to 99%.
11. Method according to either of claims 9 or 10, wherein the cold working of the semi-finished product is 0.05 to 95%.
12. Method according to any one of claims 9 to 11, wherein the cold working of the semi-finished product is 0.05 to 90%.
13. Method according to any one of claims 1 to 12, wherein the alloy contains (in % by wt.) 2.5 to 5.0% Al.
14. Method according to any one of claims 1 to 13, wherein the alloy contains (in % by wt.) 0.02 - 0.2% of the element La. 60
15. Method according to any one of claims 1 to 14, wherein the alloy contains (in % by wt.) 0.02 - 0.2% Ti.
16. Method according to any one of claims 1 to 15, wherein the alloy contains (in % by wt.) 0.02 - 0.2% of one or more of the elements Ce, Pr, Ne, Sc, Y, Zr or Hf.
17. Method according to any one of claims 1 to 16, wherein the alloy additionally contains (in % by wt.) one or more of the elements 0.0001 - 0.05% mg, 0.0001 - 0.03% Ca, 0.002 0.03% P.
18. Method according to any one of claims 1 to 17, wherein the alloy further contains (in % by wt.) up to 0.003% B.
19. Method according to any one of claims 1 to 18, wherein (in % by wt.) the Nb-content is 0.3 to 1.0% and the Si-content is 0.15 - 0.5%.
20. Method according to any one of claims 1 to 19, wherein the iron-chromium alloy which is thermomechanically treated and which precipitates Laves phases in finely dispersed form consists of the following composition (in % by wt.) Cr 12 - 30% Mn 0.001 - 2.5% Nb 0.1 - 2% W 0.1 - 5% Si 0.05 - 1% C 0.002 - 0.03% N 0.002 - 0.03% S max. 0.005% 61 Fe remainder as well as the usual melting-related impurities.
21. Method according to any one of claims 1 to 20, wherein semi-finished product is formed by sheet, strip, bar, forging, pipe or wire.
22. Method according to any one of claims 1 to 21, wherein the heat treatment is carried out only after finishing of the component.
23. Method according to any one of claims 1 to 22, wherein, by the thermomechanical treatment of the semi-finished product, a particularly high creep strength is produced in the semi-finished product and/or in the component with simultaneous elongation > 13 % in the tension test at room temperature.
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DE102012004488A1 (en) * | 2011-06-21 | 2012-12-27 | Thyssenkrupp Vdm Gmbh | Heat-resistant iron-chromium-aluminum alloy with low chromium evaporation rate and increased heat resistance |
JP5344103B1 (en) | 2011-11-22 | 2013-11-20 | 新日鐵住金株式会社 | Ferritic heat resistant steel and method for producing the same |
EP2840158A1 (en) * | 2013-08-21 | 2015-02-25 | MTU Aero Engines GmbH | Ferritic FeAlCr alloy with ternary Laves phases and with oxides and/or carbides for components of a gas turbine |
CN103667997A (en) * | 2013-11-08 | 2014-03-26 | 张超 | Refractory steel material for pump valves and preparation method thereof |
CN103667996A (en) * | 2013-11-08 | 2014-03-26 | 张超 | Wear-resistant low-carbon steel material for pumps and preparation method thereof |
CN103695792B (en) * | 2013-11-14 | 2016-01-13 | 安徽荣达阀门有限公司 | A kind of High carbon alloy steel wear-resistance valve material and preparation method thereof |
EP2876177B1 (en) * | 2013-11-22 | 2017-07-12 | MTU Aero Engines GmbH | Material from laves phase and ferritic fe-al phase |
FR3027032B1 (en) * | 2014-10-08 | 2021-06-18 | Air Liquide | ALLOY MICROSTRUCTURE FOR REFORMING TUBE |
DE102016122603B4 (en) * | 2016-11-23 | 2020-04-09 | Lob Gmbh | Process for producing a heat exchanger element and heat exchanger element |
JP6881119B2 (en) * | 2017-07-14 | 2021-06-02 | 大同特殊鋼株式会社 | Ferritic stainless steel and heat resistant members |
CN109988975A (en) * | 2017-12-29 | 2019-07-09 | 中国核动力研究设计院 | The regulation method of disperse nanometer precipitated phase is obtained in a kind of FeCrAl alloy |
US10883160B2 (en) | 2018-02-23 | 2021-01-05 | Ut-Battelle, Llc | Corrosion and creep resistant high Cr FeCrAl alloys |
CN110273114B (en) * | 2019-08-06 | 2021-04-09 | 华北理工大学 | Wear-resistant iron-silicon-chromium alloy and preparation method thereof |
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CN111440993B (en) * | 2019-12-12 | 2021-06-18 | 广东省钢铁研究所 | Iron-chromium-aluminum alloy bar and preparation method thereof |
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