US20120145285A1 - Method for producing an iron-chromium alloy - Google Patents

Method for producing an iron-chromium alloy Download PDF

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US20120145285A1
US20120145285A1 US13/389,677 US201013389677A US2012145285A1 US 20120145285 A1 US20120145285 A1 US 20120145285A1 US 201013389677 A US201013389677 A US 201013389677A US 2012145285 A1 US2012145285 A1 US 2012145285A1
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alloy
semifinished product
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containing particles
protective gas
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Heike Hattendorf
Osman Ibas
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VDM Metals GmbH
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    • CCHEMISTRY; METALLURGY
    • C21METALLURGY OF IRON
    • C21DMODIFYING 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
    • C21D8/00Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • CCHEMISTRY; METALLURGY
    • C22METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22CALLOYS
    • C22C38/00Ferrous alloys, e.g. steel alloys
    • C22C38/18Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/22Ferrous alloys, e.g. steel alloys containing chromium with molybdenum or tungsten
    • CCHEMISTRY; METALLURGY
    • C21METALLURGY OF IRON
    • C21DMODIFYING 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
    • C21D6/00Heat treatment of ferrous alloys
    • C21D6/002Heat treatment of ferrous alloys containing Cr
    • CCHEMISTRY; METALLURGY
    • C21METALLURGY OF IRON
    • C21DMODIFYING 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
    • C21D6/00Heat treatment of ferrous alloys
    • C21D6/02Hardening by precipitation
    • CCHEMISTRY; METALLURGY
    • C21METALLURGY OF IRON
    • C21DMODIFYING 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
    • C21D8/00Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0221Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
    • C21D8/0226Hot rolling
    • CCHEMISTRY; METALLURGY
    • C21METALLURGY OF IRON
    • C21DMODIFYING 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
    • C21D8/00Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0221Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
    • C21D8/0236Cold rolling
    • CCHEMISTRY; METALLURGY
    • C22METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22CALLOYS
    • C22C38/00Ferrous alloys, e.g. steel alloys
    • C22C38/18Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/26Ferrous alloys, e.g. steel alloys containing chromium with niobium or tantalum

Definitions

  • the invention relates to a ferritic iron-chromium alloy produced by melting metallurgy.
  • DE 100 25 108 A1 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 O 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 A1 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.
  • 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 the elements niobium, molybdenum, tungsten or tantalum, has become known from DE 10 2006 007 598 A1.
  • the steel is preferably intended to be used for a bipolar plate in a fuel-cell stack.
  • EP 1 536 031 A1 discloses a metallic material for fuel cells, containing C ⁇ 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 ⁇ 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 A1 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%.
  • these alloys have better hot strength and elevated creep strength at temperatures of 700° C. and above, specifically due to formation of precipitates, which hinder the dislocation movements and thus plastic deformation of the material.
  • 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.
  • the alloys cited in DE 10 2006 007 598 A1, EP 1 536 031 A1 and EP 1 882 756 A1 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.
  • 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.
  • the excellent corrosion resistance is achieved by formation of a chromium oxide top layer.
  • 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.
  • 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.
  • oxygen-affine elements such as La, Ce, Y, Zr or the like further increases the corrosion resistance.
  • a specimen is subjected to a constant static tensile force at a constant temperature.
  • this tensile force is expressed as an initial tensile stress relative to the initial cross section of the specimen.
  • the time t B 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, “Bürgel”, page 100).
  • the deformability is determined in a tension test according to DIN 50145 at room temperature. In the process, the offset yield strength R p0.2 , the tensile strength R M and the elongation at break are determined. The elongation A is determined on the broken specimen from the elongation of the original gauge length L 0 :
  • 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 ⁇ m and larger are visible by viewing in an optical microscope. Smaller particles may not be recognized, but are definitely present. Therefore 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.
  • 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.
  • 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.
  • 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 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
  • 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 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 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 aw , and t max to a heat treatment in the
  • the times t min and t max are expressed in minutes and the heat treatment temperature T in ° C.
  • the task is also accomplished by a metallic component or semifinished product consisting of the following chemical composition (in % by weight)
  • temperatures for the creep test preferably lie in the range between 500 and 1000° C.
  • 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).
  • 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° C. and 1060° C. with subsequent heat treatment for a time between t min and t max 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° C. is quenched in the oven, after which the desired component is made if applicable, wherein t min and t max are calculated according to the following formulas:
  • the working of the semifinished product can take place by hot working. Alternatively, however, the forming can also be brought about by cold working.
  • the semifinished product is hot-worked with a starting temperature ⁇ 1070° C., wherein the last 0.05 to 95% of mechanical deformation is applied between 1000° and 500° C., advantageously the last 0.5 to 90% between 1000° C. and 500° C.
  • the degree of cold working of the semifinished product is 0.05 to 99%, advantageously 0.05 to 95% or 0.05 to 90%.
  • the mechanical working of the semifinished product be 20 to 99% and then the worked semifinished product be subjected for a time between t min and t max to a heat treatment in the temperature range between 950° C. and 1060° 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 t min and t max are calculated according to the following formulas:
  • 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.
  • both variants may be considered.
  • 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.
  • 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%.
  • the alloy in % by weight may contain one or more of the elements Ce, La, Pr, Ne, Sc, Y, Zr or Hf in contents of 0.02-0.3%.
  • 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%.
  • the Nb content is 0.3 to 1.0% and the Si content 0.15 to 0.5%.
  • the element tungsten may be replaced entirely or partly by at least one of the elements Mo or Ta.
  • 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%.
  • the alloy may also even contain max. 0.003% boron.
  • 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.
  • 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 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° C. for a period of longer than 6 minutes or above 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.
  • the exact temperature control above this solution annealing temperature is not determining for the characteristics.
  • the annealing may be carried out in air or under protective gas. It should lie under the melting temperature, preferably ⁇ 1350° 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.
  • 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 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° C. At temperatures below 500° C., exist in the range of the embrittlement that occurs in ferrites at 475° 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, “Bürgel”, 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° C. and 90% for the cold working.
  • annealing between 950° 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 application temperature between 550° C. and 1000° C. by heating at 0.1° C./min to 1000° C./min.
  • 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.
  • the worked semifinished product/component is subjected for a time between t min and t max to a heat treatment in the temperature range between 550° C. and 1060° 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
  • the desired component can be made before or after this heat treatment.
  • Times shorter than t min 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.
  • t max the danger exists of too great coarseness of the precipitates, whereby the particles can no longer contribute markedly to the creep strength.
  • t max the possibility exists in the upper temperature range of 550° C. and 1060° C. that a recovered microstructure will be formed, which certainly can still be effective.
  • 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.
  • times longer than t max are uneconomical.
  • the annealing step may be carried out under protective gas (argon, hydrogen and similar atmospheres with reduced oxygen partial pressure).
  • protective gas argon, hydrogen and similar atmospheres with reduced oxygen partial pressure.
  • 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° C. but is also possible at temperatures ⁇ 800° 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 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.
  • 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%.
  • this restriction does not exist.
  • 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 Fe 7 (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.
  • the lower limit is determined in each case by the fact that a certain content must be present in order to be effective, while the upper limit is determined by the processability.
  • the preferred ranges are, for
  • 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 7 (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.
  • 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.
  • 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%.
  • An example of this is, for example, for the application as interconnector plate.
  • 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.
  • 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.
  • 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.
  • the upper limit is set at 0.3% by weight.
  • titanium is bound in the 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 A1). These swellings are unfavorable, since the resulting cracks cause an increase of the oxidation rate.
  • 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%.
  • 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.
  • a material with an analysis as indicated in Table 1 precipitates mainly Fe 2 (M, Si) or Fe 7 (M, Si) 6 Laves phases and, in much smaller contents, carbides.
  • material from the batch 161061 listed in Table 1 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 mechanical working was begun with a start temperature >1070° C. and the last 78% of mechanical deformation was applied by rolling between 500° C. and 1000° C.
  • FIG. 1 shows the typical appearance of a microstructure deformed in this way.
  • 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.
  • 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° 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.
  • annealing steps were performed on the hot-rolled material from Example 1 for 20 minutes each between 600° C. and 1000° C. or for some temperatures also for 240 or 1440 minutes (see Table 3 for t min and t max 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 750° C. as described above. The results are compiled in Table 3.
  • FIG. 3 shows the microstructure after the various annealing steps for 20 minutes.
  • the microstructures in FIG. 3 are shown in FIG. 3 .
  • the microstructure 3 are not globularly recrystallized. Up to 850° C. (the maximum of the time to break), the microstructure has the typical appearance of a deformed microstructure. Starting from approximately 900° C., 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 preworking.
  • 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° C. (the maximum of the time to break).
  • 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).
  • 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.
  • the sheet annealed for 20 minutes at all temperatures between 600° C. and 950° C. 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 800° C. 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.
  • the annealing time of 20 minutes at the temperatures between 600° C. and 950° C. lies in the inventive range between t min and t max . 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.
  • material from the batch 161061 listed in Table 1 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 60% of mechanical working was applied by rolling between 1000° C. and 500° C.
  • 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 ( FIG. 4 ).
  • the time to break is 30 times the time achieved in Example 1 after annealing at 1075° C. for 20 minutes with a globularly recrystallized coarse-grained microstructure with a grain size of 137 ⁇ m.
  • 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.
  • material from batch 161061 and batch 161995 was cold-rolled to 1.5 mm thick sheet after solution annealing at above 1070° 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 1050° C. 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 ( FIG. 5 ) and batch 161995 exhibit a recovered microstructure with elongated grains ( FIG. 7 ) and precipitation of Laves phase, although much less than recognizable in FIG. 4 .
  • Tables 5a and 5b show the results of the creep tests and of the 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° C. with an initial load of 35 MPa, and after additional annealing at 1050° 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%.
  • batch 161995 has a time to break of 33.5 hours 89% in a creep test at 750° C. with an initial load of 35 MPa, and after additional annealing at 1075° 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 92%.
  • the elongation of 28% in the tension test at room temperature for batch 161061 for 1050° C. 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.
  • 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° C. and 500° C.
  • Example 2 annealing steps were performed on the hot-rolled material from Example 1 for 20 minutes each between 750° C. and 1000° C. or for some temperatures also for 120 minutes, 240 minutes, 480 minutes, 960 minutes, 1440 minutes or 5760 minutes (see Table 7 for t min and t maX according to Equation 1 and 2) in air, followed by quenching in stationary air.
  • the creep test was performed with a stress of 40 MPa at 750° 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.
  • Example 2 After 20 minutes at 1000° C., a time to break of only 8.8 hours is reached at an elongation A of 78.7%.
  • Example 2 After 20 minutes at 1000° C. and a creep test at 750° C. and 35 MPa, a time to break comparable with that after solution annealing at 1075° 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.
  • 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 850° C. with 296 hours.
  • the longest time to break for annealing steps of 120 minutes is achieved at 800° C. with 227 hours.
  • the longest time to break for annealing steps of 240 minutes is achieved at 750° C. with 182 hours, but in this connection no value exists for 700° C.
  • the longest time to break for annealing steps of 480 minutes is achieved at 800° C. with 169 hours.
  • Example 2 At room temperature, the sheet annealed for 20 minutes at all temperatures between 600° C. and 900° 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.
  • 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.
  • 2.5 mm thick material from batch 161995 was produced by annealing it, after cold working of 40%, at 1050° C. for 2.8 minutes under protective gas in the continuous 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 ( FIG. 11 ), just as the material from Example 4 in FIG. 7 , and precipitation of Laves phase, albeit clearly less than recognizable in FIG. 4 .
  • batch 161995 After the annealing at 1050° C. 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° 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%.
  • 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° C., 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%.
  • 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° 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° C./2.8 minutes, wherein a maximum occurs at 10% and the elongations at break lie between 18.9 and 60%.
  • FIG. 1 Microstructure of the hot-worked material in Example 1
  • FIG. 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 ⁇ m.
  • FIG. 3 Microstructure of the material in Example 2 after annealing between 600° C. and 1000° C. for 20 minutes in each case and quenching in stationary air.
  • FIG. 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)
  • FIG. 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.
  • FIG. 6 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 and annealing at 1050° C./20 minutes under air with subsequent quenching in stationary air, grain size 134 ⁇ m (etching with V2A pickling fluid)
  • FIG. 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.
  • FIG. 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
  • FIG. 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 ⁇ m to 162 ⁇ m
  • FIG. 10 Microstructure of the hot-worked material in Example 5 after annealing at 1075° C. for 22 minutes followed by quenching in stationary air and subsequent annealing at 700° C. for 4 hours followed by quenching in stationary air. Grain size 136 ⁇ m.
  • FIG. 11 Microstructure of batch 161995 in Example 7 after annealing at 1050° C. for 2.8 minutes under protective gas in the continuous furnace with quenching in the stream of cold protective gas.
  • FIG. 12 Microstructure of batch 161995 in Example 7 after annealing at 1050° 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 ⁇ m.
  • n.m. n.m. I 650 240 11.26 31.8*10 5 293 36.2 n.m. n.m. n.m. I 700 240 5.00 33.6*10 4 233 32.8 n.m. n.m. n.m. I 750 240 2.41 44300 224 23.2 n.m. n.m. n.m. I 800 240 1.25 7059 396 43.2 n.m. n.m. n.m. I 850 240 0.69 1328 181 35.2 n.m. n.m. n.m. I 900 240 0.40 289 45.6 55.7 n.m. n.m. n.m.
  • I 750 120 2.41 44300 150 31.7 n.m. n.m. n.m. I 800 120 1.25 7059 227 26.2 n.m. n.m. n.m. I 850 120 0.69 1328 133 29.7 n.m. n.m. n.m. I 900 120 0.40 289 32.7 40 n.m. n.m. n.m. n.m. not measured
  • I 750 480 2.41 44300 152 43.4 n.m. n.m. n.m. I 800 480 1.25 7059 169 26.5 n.m. n.m. n.m. I 850 480 0.69 1328 35 33 n.m. n.m. n.m. I 750 960 2.41 44300 139 24.2 n.m. n.m. n.m. I 750 1440 2.41 44300 82 25.5 n.m. n.m. n.m. I 800 1440 1.25 7059 46 46.1 n.m. n.m. n.m. I 750 5760 2.41 44300 54 52.9 n.m. n.m. n.m. I 800 5760 1.25 7059 17.5 50.3 n.m. n.m. n.m. not measured

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