EP2480695A1 - Procédé de production d'un alliage fer-chrome - Google Patents

Procédé de production d'un alliage fer-chrome

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Publication number
EP2480695A1
EP2480695A1 EP10760237A EP10760237A EP2480695A1 EP 2480695 A1 EP2480695 A1 EP 2480695A1 EP 10760237 A EP10760237 A EP 10760237A EP 10760237 A EP10760237 A EP 10760237A EP 2480695 A1 EP2480695 A1 EP 2480695A1
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EP
European Patent Office
Prior art keywords
alloy
containing particles
max
air
inert gas
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Granted
Application number
EP10760237A
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German (de)
English (en)
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EP2480695B1 (fr
Inventor
Heike Hattendorf
Osman Ibas
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VDM Metals International GmbH
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ThyssenKrupp VDM GmbH
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Classifications

    • 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
    • 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
    • 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
    • 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
    • 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
    • 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
    • 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

Definitions

  • the invention relates to a molten metallurgically produced ferritic iron-chromium alloy.
  • DE 100 25 108 A1 discloses a high-temperature material comprising a chromium oxide-forming iron alloy 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 one Elements M from the group Mn, Ni and Co, which forms with chromium oxide at high temperatures a spinel phase of the type MCr 2 O 4 , up to 2 wt .-% of another element from the group Ti, Hf, Sr, Ca and Zr, which increases the electrical conductivity of Cr-based oxides.
  • the chromium content should be present in a concentration range between 12 and 28%. Areas of application 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 having 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% SE and not more than 1% Zr, balance iron and manufacturing impurities.
  • DE 10 2006 007 598 A1 discloses a creep-resistant ferritic steel comprising precipitations of an intermetallic phase of the Fe 2 (M, Si) or Fe 7 (M, Si) 6 type with at least one metallic alloying element M, which passes through the elements Niobium, molybdenum, tungsten or tantalum may be formed can.
  • the steel should preferably be used for a bipolar plate in a fuel cell stack.
  • EP 1 536 031 A1 has disclosed a metallic material for fuel cells, comprising 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%, balance iron and unavoidable impurities, the composition of the following equation should satisfy: 0, 1 -S Mo / Nb ⁇ 30.
  • EP 1 882 756 A1 describes a ferritic chrome steel, in particular 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%, SE max. 0.3%, AI max. 0, 1%, N max. 0.07%, remainder Fe and impurities caused by melting, the content of Zr + Ti being at least 0.2%.
  • the alloys listed in DE 10 2006 007 598 A1, EP 1 536 031 A1 and EP 1 882 756 A1 are optimized for use as interconnector plate for the high-temperature fuel cell: they have improved by using a ferritic alloy with 10 to 40% Chromium an expansion coefficient adapted to the ceramic components anode and electrolyte as far as possible.
  • the requirements of the reformer and the heat exchanger for the high-temperature fuel cell are as good as possible creep resistance, a very good corrosion resistance and low chromium evaporation.
  • the oxide does not have to be conductive for these components.
  • the requirements for components for example, for the exhaust system of an internal combustion engine or for steam boilers, superheated, turbines, and other parts of a power plant are the best possible creep resistance, a very good corrosion resistance.
  • the chromium evaporation does not produce any poisoning phenomena here, as in the fuel cell, and the protective oxide does not have to be conductive for such components.
  • the excellent corrosion resistance is achieved by forming a chromium oxide topcoat. Due to the fact that a spinel with Mn, Ni, Co or Cu is additionally formed on the chromium oxide topcoat, less volatile chromium oxides or chromium oxyhydroxides form, which poison the cathode.
  • a sample is subjected to a constant static tension at a constant temperature. For reasons of comparability, this tensile force is given as the initial tensile stress in relation to the initial sample cross-section.
  • the time to break t ß - the break time - the sample is measured. The test can then be carried out without strain measurement on the sample during the test. The elongation at break is then measured after the end of the experiment.
  • the sample is installed in the creeping machine at room temperature and heated to the desired temperature with no tensile force. After reaching the test temperature, the sample is held for one hour without temperature load for temperature compensation. Thereafter, the sample is loaded with the tensile force and the test time begins.
  • the break time can be taken as a measure of the creep resistance. The greater the break time at a given temperature and initial tensile stress, the more creep resistant the material is. Breaking time and creep resistance decrease with increasing temperature and increasing initial tensile stress (see eg "Bürgel" page 100)
  • the deformability is determined in a tensile test according to DIN 50145 at room temperature.
  • the yield strength R p o , 2, the tensile strength RM and the elongation to breakage are determined.
  • the elongation A is determined on the fractional sample from the extension of the original measurement distance l_o:
  • the elongation at break is provided with indices:
  • the amount of elongation A in the tensile 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 visualized on a metallographic grinding by etching with V2A pickling or electrolytic etching with oxalic acid do.
  • V2A pickling When etching with V2A pickling, the grains or grain boundaries are additionally visibly etched. When viewed in a light microscope, only particles of a size of approximately 0.5 ⁇ m are visible. Smaller ones can not be recognized, but they do exist. Therefore, metallography is only used as a supporting explanation Effectiveness of a measure is conveniently evaluated by break time or creep strength.
  • the task is to improve the creep resistance of the precipitation-hardened iron-chromium alloy
  • the invention has for its object to provide a method for producing a precipitation hardened iron-chromium alloy component, by means of which, while maintaining an acceptable deformability at room temperature, the high heat resistance or Creep resistance of a precipitation-hardened ferritic alloy over the prior art further increase.
  • thermomechanically treated consisting of an iron-chromium alloy component / semi-finished is to be provided, which can be used to achieve a high heat resistance or creep resistance, while maintaining an acceptable ductility at room temperature.
  • This object is achieved on the one hand by a method for producing a component, of a Laves phase and / or Fe-containing particles and / or Cr-containing particles and / or Si-containing particles and / or carbides ausscheidenden, iron-chromium alloy in that a semi-finished product made of 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 cooling in inert gas or air, inert (blown) inert gas or air or in water, in a second step, a mechanical deformation of the semifinished product in the range of 0.05 to 99% is carried out and in a subsequent step Laves phases Fe 2 (M, Si) or Fe 7 (M, Si) 6 and / or Fe-containing particles and / or Cr-containing particles and / or Si-containing particles and / or carbides are selectively and finely distributed thereby excreted, that the manufactured from the formed semi-finished component du a heating with from
  • this object is achieved by a method for producing a component, iron-chromium alloy leaving a Laves phase and / or Fe-containing particles and / or Cr-containing particles and / or Si-containing particles and / or carbides in that a semi-finished product made from the alloy is subjected to a thermomechanical treatment, wherein in a first step the alloy is solution annealed at temperatures> above the solution annealing temperature, followed by cooling in inert gas or air, agitated (blown) inert gas or air or in water, in a second step mechanical deformation of the semifinished product in the range 0.05 to 99% and in a subsequent step Laves phases Fe 2 (M, Si) or Fe 7 (M, Si) 6 and / or Fe-containing particles and / or Cr-containing particles and / or Si-containing particles and / or carbides targeted and be finely distributed by the fact that the formed semi-finished product for a time between t m j n and t max is subjected to
  • thermo-mechanical treatment has a deformed microstructure such that Laves phase (s) are incorporated in finely divided form into the microstructural dislocations of the microstructure, wherein in a creep test with e.g. 35 MPa at 750 ° C and at an elongation of at least 18% a fracture time is set in the microstructure, which exceeds the fracture time of a coarse-grained, fully recrystallized microstructure by a factor of at least 1, 5.
  • the temperatures for the creep test preferably being in the range between 500 and 1000 ° C.
  • a semifinished product is made from a
  • thermomechanically Alloy of the following composition (in% by weight) treated thermomechanically:
  • the method according to the invention makes it possible to produce semi-finished products in the form of sheets, strips, rods, forgings, pipes or wire and to manufacture components in a wide variety of shapes required for the respective application.
  • the solution annealing temperature preferably> 1050 ° C for more than 6 minutes, respectively> 1060 ° C for more than 1 minute, followed by cooling in inert gas or air, moving (blown ) Inert gas or air or in water at the initial state before deformation in the semi-finished product only few or no Laves phases Fe 2 (M, Si) or Fe 7 (M, Si) 6 and / or Fe-containing particles and / or Cr containing particles or Si-containing particles and / or carbides are present in the alloy.
  • the deformation of the semifinished product can be done by hot working. Alternatively, however, the deformation can also be brought about by cold working.
  • the semi-finished product with a starting temperature> 1070 ° C hot formed the last 0.05 to 95% of mechanical deformation between 1000 ° and 500 ° C are applied, advantageously the last 0.5 to 90% between 1000X 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 deformation of the semifinished product be 20 to 99% and thereafter the formed semi-finished product for a period between t m i n and t max a
  • alloy already mentioned is to be used as an interconnector for a solid oxide fuel cell, a content of 0.001-0.5% aluminum is advantageous.
  • a content of 2 to 6% aluminum is advantageous, since then can form a closed aluminum oxide, compared to a chromium oxide layer again a much lower growth rate and in addition still has lower Chromoxidabdampfung than a chromium-manganese spinel.
  • both variants can be considered. It should be remembered in particular that with increasing aluminum content, the processability and weldability of the alloy deteriorates, resulting in higher costs. Therefore, if an oxide layer consisting of a chromium oxide and a chromium-manganese spinel, a sufficient oxidation resistance by use of 0.001 - 0.5% aluminum can be guaranteed. If a higher oxidation resistance is required, as is ensured, 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 system of an internal combustion engine or for steam boilers, superheated, turbines and other parts of a power plant.
  • a preferred aluminum range is in particular the range of 2.5% to 5.0%, which is still characterized by good processability.
  • the contents of the additional elements which can be introduced in the alloy can be set 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) of one or more of the elements Ce, La, Pr, Ne, Sc, Y, Zr or Hf may contain, in contents, 0.02-0.3%.
  • the alloy (in% by weight) of one or more of the elements Ce, Pr, Ne, Sc, Y, Zr or Hf may contain in amounts 0.02-0.2%.
  • the Nb content is 0.3 to 1.0% and the Si content is 0.15 to 0.5%.
  • the element tungsten may be wholly or partly replaced by at least one of the elements Mo or Ta.
  • the alloy can also max. 0.2% V and / or max. 0.005% S included.
  • the oxygen content should not be greater than 0.01%.
  • the alloy can also max. Containing 0.003% boron.
  • the alloy should each have a maximum of 0.01% of the following elements: Zn, Sn, Pb, Se, Te, Bi, Sb.
  • Components / semi-finished products which consist on the one hand of the mentioned alloy composition and on the other hand produced by the method according to the invention can preferably be used as an interconnector in a fuel cell or as a material in a component, such as a reformer or a heat exchanger in an auxiliary unit of the fuel cell.
  • the component / semifinished product produced by the method according to the invention, or the alloy itself is also possible to use as a component in the exhaust system of an internal combustion engine or for steam boilers, superheaters, turbines and other parts of a power plant or in the chemical process industry.
  • the first step of a Laves phase and / or Fe-containing particles and / or Cr-containing particles and / or Si-containing particles and / or carbides ausschußenden iron-chromium alloy in the thermomechanical treatment must be an annealing above the solution annealing temperature in order that the Laves phases and / or Fe-containing particles and / or Cr-containing particles and / or Si-containing particles and / or carbides are brought into solution and are available for precipitation for the subsequent thermomechanical treatment.
  • the solution annealing temperature is alloy-dependent, but is preferably above 1050 ° C for more than 6 minutes, or above 1060 ° C for more than 1 minute, followed by quenching in inert gas or air, moving (blown) inert gas, or Air or in water.
  • the exact temperature control above this solution annealing temperature is not decisive for the properties.
  • the annealing can be done in air or under inert gas. It should be below the melting temperature, preferably ⁇ 1350 ° C. For reasons of cost, the annealing times should preferably be ⁇ 24 hours, but may also be longer depending on the performance.
  • the solution annealing is followed by cooling in inert gas or air, moving (blown) inert gas or air or in water, in which only a small Laves phase forms new.
  • an increased dislocation density must be introduced into the material.
  • Increased dislocation densities have reshaped structures or recovered microstructures, where the dislocations are located in small-angle grain boundaries.
  • the second step must therefore be a forming so that the dislocations are introduced into the material, which then in the subsequent annealing for a uniform distribution of Laves phases and / or Fe-haitigen particles and / or Cr-containing particles and / or Si -containing particles and / or carbides provide.
  • This deformation can be a cold forming, but also a hot forming, whereby the hot working must ensure that the structure is not fully recrystallized during rolling. This is done by limiting the deformation range for the last forming and the temperature at which it takes place. In the case of deformations above 1000 ° C., the material already tends to recrystallise or recover during the deformation, so that the transformation must preferably take place below 1000 ° C. At temperatures below 500 ° C are in the range occurring in ferrites 475 ° C embrittlement. There it has a lower elongation and an increased resistance to deformation, which makes a transformation less advantageous and reduces the efficiency.
  • One way 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 formed material is to produce the required components from the semifinished product and then To bring the manufactured component by heating with from 0.1 ° C / min to 1000 ° C / min to the application temperature between 550 ° C and 1000 ° C.
  • the Laves phases and / or Fe-containing particles and / or Cr-containing particles and / or Si-containing particles and / or carbides are excreted finely distributed in the microstructure.
  • the fine distribution is produced by nucleation in the lower temperature range, followed by some growth of the nuclei at the higher temperatures. Therefore, the heating rate must not be faster than 1000 ° C / min, otherwise the time for this process is too low. Heating rates less than 0.1 ° C / minute are uneconomical.
  • a second possibility is a separate heat treatment of the material.
  • the formed semifinished product / component is subjected to a heat treatment in the temperature range between 550 ° and 1060 ° C. under protective gas or air for a period between t m j n and t ma x, followed by cooling in inert gas or air, agitated (blown ) Inert gas or air or in water or for heat treatments up to 800 ° C in the oven, wherein
  • the desired component can be manufactured before or after this heat treatment.
  • Times shorter than t m i n are insufficient to form the Laves phases and / or Fe-containing particles and / or Cr-containing particles and / or Si-containing particles and / or carbides.
  • t max there is the danger of an excessive coarsening of the precipitates, as a result of which the particles can no longer noticeably contribute to creep resistance.
  • tmax in the upper temperature range of 550 ° C and 1060 ° C, there is the possibility that a recovered microstructure may arise, which may still be effective. However, as the recovery proceeds, the dislocation density is further reduced, so that the distribution of the precipitates becomes increasingly uneven and the positive effect on creep resistance finally disappears.
  • times above t max are additionally uneconomical.
  • the annealing can take place 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 cooling takes place for economic reasons in inert gas or air, moving (blown) inert gas or air or in water, especially for temperatures above 800 ° C, a furnace cooling should be avoided, at temperatures ⁇ 800 ° C, this is also possible.
  • the chromium content determines the oxidation resistance and the thermal expansion coefficient of the material.
  • the oxidation resistance of the material is based on the formation of a closed chromium oxide layer. Below 12%, especially at higher operating temperatures, iron-containing oxides are increasingly formed which impair the oxidation resistance.
  • the chromium content is therefore set to> 12%. Above 30% of chromium, the workability of the material and its usability by increased formation of embrittling phases, in particular the sigma phase, impaired.
  • the chromium content is therefore limited to ⁇ 30%. As the chromium content increases, the expansion coefficient decreases.
  • the coefficient of expansion can thus be set in a range which matches the ceramics in the fuel cell.
  • These are chromium contents around 22 to 23%.
  • this restriction does not exist.
  • manganese causes the formation of a chromium-manganese spinel on the chromium oxide layer that forms on the material for lower aluminum contents of less than 2%.
  • This chromium manganese spinel reduces chromium evaporation and improves contact resistance. It is necessary for at least a manganese content of 0.001%. More than 2.5% manganese affects the oxidation resistance by forming a very thick chromium-manganese spinel layer.
  • Niobium, molybdenum, tungsten or tantalum may be involved in the formation of precipitates in iron-containing alloys, such as carbides and / or M in the Laves phases Fe 2 (M, Si) or Fe 7 (M, Si) 6 , Molybdenum, tungsten or tantalum are also good mixed crystal hardeners and thus contribute to the improvement of creep resistance.
  • the lower limit is determined in each case by the fact that a certain amount must be present in order to be effective; the upper limit is determined by the processability.
  • the preferred range of Nb is 0.1-2%
  • W can also be completely or partially replaced by Mo and Ta: 0.1 - 5%
  • Silicon may be involved in the formation of precipitates in iron-containing alloys, for example in the Laves phase Fe 2 (M, Si) or Fe 7 (M, Si) 6. It ensures the increased precipitation and stability of these Laves phases and thus contributes to creep resistance. When the Laves phase is formed, it becomes complete tied into these. Thus, the formation of a silicon oxide layer below the chromium oxide layer no longer occurs. At the same time, the incorporation of M into the oxide layer is reduced, thereby preventing the negative influence of M on the oxidation resistance. At least 0.05% Si must be present for the desired effect to occur. If the content of Si is too large, the negative effect of Si may occur again. The Si content is therefore limited to 1%.
  • Aluminum deteriorates the oxidation resistance at levels below 1% because it leads to internal oxidation.
  • an aluminum content greater than 1% leads to the formation of an aluminum oxide layer below the chromium 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 generator is desired or its oxidation resistance is sufficient.
  • An example of this is e.g. for use as interconnector plate.
  • a certain aluminum content of at least 0.001% is necessary for deoxidizing the melt.
  • the alloy can 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%, so as not to affect the processability. However, it should be> 0.002% for an effect to occur.
  • the nitrogen content should not exceed 0.1% in order to avoid the formation of processability deteriorating nitrides. It should be greater than 0.002% to ensure the workability of the material.
  • the levels of sulfur should be as low as possible, since this surfactant affects the oxidation resistance. It will therefore max. 0.01% S set.
  • Oxygen-affinity elements such as Ce, La, Pr, Ne, Sc, Y, Zr, Hf improve oxidation resistance by reducing oxide growth and improving adhesion of the oxide layer. It is useful to have a minimum content of 0.02% of one or more of Ce, La, Pr, Ne, Sc, Y, Zr, Hf to obtain the oxidation resistance-enhancing effect of Y.
  • the upper limit is set for cost reasons at 0.3 wt .-%.
  • Titanium like any oxygen-affine element, is incorporated into the oxide layer during oxidation. In addition, it still causes internal oxidation. However, the resulting oxides are so small and finely divided that they cause a hardening of the surface and thus prevent bulging of the oxide layer and the inclusion of metallic areas during the oxidation (see DE 10 2006 007 598 A1). These bulges are unfavorable because the resulting cracks cause an increase in the oxidation rate. Thus, Ti contributes to the improvement of the oxidation resistance. For Ti content to be effective, at least 0.01% Ti must be present, but not more than 0.5%, since this no longer improves the effect, but increases the cost.
  • the content of phosphorus should be less than 0.030% since this surfactant affects the oxidation resistance. A too low P content increases the costs. The P content is therefore> 0.002%.
  • the contents of magnesium and calcium are set in the spread range of 0.0001 to 0.05 wt .-%, respectively 0.0001 to 0.03 wt .-%.
  • Table 1 shows the analyzes of the batches used for the following examples. These charges were melted in an electric arc furnace in an amount of about 30 t, then poured into a pan and subjected to a decarburization and deoxidation treatment and a vacuum treatment in a VOD plant and cast into blocks. These were then hot rolled and cold rolled with intermediate annealing, depending on the final thickness. After hot rolling, the oxide layer was removed by pickling
  • a material with an analysis as given in Table 1 precipitates mainly Laves phases Fe 2 (M, Si) or Fe 7 (M, Si) 6 and in significantly reduced amounts of carbides.
  • material from the charge 161061 shown in Table 1 was hot rolled after a solution annealing above 1070 ° C for a period of more than 7 minutes followed by cooling in still air to 12 mm thick sheet, the mechanical working with a starting temperature> 1070X was started and the last 78% of mechanical deformation between 500 ° C and 1000 ° C were applied by rolling.
  • Figure 1 shows the typical appearance of such a deformed structure.
  • the etched sections by means of electrolytic etching with oxalic acid, it can be clearly seen that only a few Laves phases have been excreted microscopically.
  • Figure 2 After annealing the material thus formed at 1075 ° C for 20 minutes with cooling in still air, we obtain a microstructure with only a few precipitates of Laves phase and a grain size of about 137 pm (Figure 2), which is a typical coarse grained Structure is.
  • the hot rolled material which is equivalent to a pre-forming, prepared a sample for a creep test as a simulation for a component and then heated at about 60 ° C / minute to an application temperature of 750 ° C and then a creep test with a Surprisingly, the sample tears only after 255 hours at an elongation A of 29%, which means an increase in the time to breakage by 20 times the initial stress of 35 MPa at a temperature of 750 ° C.
  • the production of the component is very easy, since the hot formed state, as described above, in the tensile test at room temperature has an elongation of 19%, which is a good value and makes the material easy to process.
  • the structure Up to 850 ° C (the maximum of the break time), the structure has the typical appearance of a deformed structure. From about 900 ° C, a clear recovery can be seen, but this means that the dislocation density compared to a globular recrystallized structure is still increased. In a recovered structure, the dislocations have partially rearranged into small-angle grain boundaries. It has a similar effect as a pre-forming. In the etched with electrolytic etching with oxalic acid grindings can be clearly seen that from about 750 ° C the Laves phase is excreted microscopically visible, being excreted up to 850 ° C (the maximum of the fracture time) more dense and uniform.
  • 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 considered to be satisfactory for a ferritic alloy and makes the material processable.
  • the elongation is lowest in the range of 700 ° C to 800 ° C and improves in each case to the lower or higher annealing temperatures, because at the lower temperatures, although Laves phase has been eliminated, which is microscopically not yet visible and thus has a lower volume fraction, but is very finely distributed.
  • a larger volume fraction is eliminated, but a bit coarser and recognizable on the small-angle grain boundaries and grain boundaries.
  • the break time is also only 10.4 hours.
  • the annealing time of 20 minutes at temperatures between 600 ° C and 950 ° C are in the range of the invention between t m j n and t max .
  • the break time according to the invention was significantly increased by more than a factor of 7 compared to the coarse-grained, globular recrystallized state of Example 1, which is produced after annealing at 1075 ° C. for 20 minutes followed by quenching in still air.
  • material from the batch 161061 indicated in Table 1 was hot rolled to a 12 mm thick sheet after a solution annealing of more than 1070 ° C for a period of more than 7 minutes followed by quiescent air cooling, forming with a starting temperature> 1070 ° C was started and the last 60% of mechanical deformation between 1000 ° C and 500 ° C are applied by rolling. If the sheet metal shaped in this way is then calcined in the continuous furnace at 920 ° C. for 28 minutes in air and cooled in still air, a tensile test sample made from this material has a creep test with an initial stress of 35 MPa at a temperature of 750 ° C.
  • the sheet thus treated has a very good elongation of 18%, a yield strength of 475 MPa, and a tensile strength of 655 MPa (see Table 4), which makes the material readily reshapeable.
  • material from lot 161061 and lot 161995 is after a solution anneal of more than 1070 ° C for a period of more than 7 minutes, followed by cooling in blown inert gas and hot rolling, and removing the oxide layer to a 1.5 mm thick sheet cold rolled, with a cold working of 53% was applied. This was followed by annealing at 1050 ° C. for 3.4 minutes under protective gas in a continuous furnace with subsequent cooling in the cold inert gas stream.
  • Both Lot 161061 ( Figure 5) and Lot 161995 show a recovered texture with elongated grains ( Figure 7) and excretion of Laves phase, albeit much less than Figure 4.
  • Tables 5a and 5b show the results of creep tests and tensile tests at room temperature. After annealing at 1050 ° C for 3.4 minutes, Batch 161061 in a creep test at 750 ° C with an initial load of 35 MPa has a break time of 25.9 hours at an elongation A of 50%, after additional annealing at 1050 ° C for 20 minutes, the very coarse grain produces a one-third break time of 7.9 hours at an elongation A of 83%.
  • material from batch 161061 was hot rolled after a solution annealing above 1070 ° C for a period of more than 7 minutes followed by quiescent air cooling to 12 mm thick sheet, forming with a starting temperature> 1070 ° C was started and the last 70% of mechanical deformation was rolled between 1000 ° C and 500 ° C.
  • the sample travels already after 104 hours, at an elongation A of 72.6%, ie a considerably shorter time than after the solution annealing of 1075 ⁇ for 22 minutes , If one carries out the creep test with 60 MPa, then tears the sample after 6.3 hours, at an elongation A of 63%, ie also after a much shorter time than after the solution annealing of 1075 ° C for 22 minutes.
  • Example 2 with the hot rolled material of Example 1 anneals between 750 ° C and 1000 ° C for 20 min and for some temperatures also 120 min, 240 min, 480 min, 960 min, 1440 min and 5760 min, (for t min and t ma x according to Equation 1 and 2, see Table 7) in air, followed by cooling in still air carried out.
  • samples were prepared from the sheets, and then the creep test was carried out at a tension of 40 MPa at 750 ° C as described above.
  • the higher voltage compared to Example 2 was chosen to shorten the experimental time.
  • the aim was to determine the heat treatment times suitable for the annealing. The results are summarized in Table 7.
  • an at least 10-fold increased break time of more than 100 hours is achieved at an elongation A of greater than 27%.
  • the highest break time for annealing with 20 minutes is achieved at 850 ° C with 296 hours.
  • the highest break time for 120 minute anneals is achieved at 800 ° C with 227 hours.
  • the highest break time for the 240 minute anneals is achieved at 750 ° C for 182 hours, with no value for 700 ° C.
  • the highest break time for the 480 minute anneals is achieved at 800 ° C with 169 hours.
  • For 960 minutes only a break time of 750 ° C would be determined, with a value of 139 hours at an elongation of 24.2%.
  • the sheet tempered for 20 minutes at each temperature has an elongation of at least 13% between 600 ° C and 900 ° C as in Example 2, which is still considered to be satisfactory for a ferritic alloy and makes the material processable.
  • a sample for a creep test is produced as a simulation for a component from the material converted to 1050 ° C / 2.8 min. Heated to an application temperature of 750 ° C and then subjected to a creep test with an initial stress of 35 MPa at a temperature of 750 ° C, so the elongation at break decreases with conversions between 5 and 40% on values around the 10 hours Elongation at break greater than 45%.
  • a sample for a creep test is produced from the material formed after 1050 ° C. for 10 minutes as a simulation for a component and then heated to an application temperature of 750 ° C. at about 60 ° C./minute and then a creep test with an initial stress of 35 MPa is carried out at a temperature of 750 ° C, the elongation at break during forming increases between 2.9 and 40% to values between 49 and 137 hours, which means an increase in the breaking time compared to that after 1050 ° C / 2.8 min Material by more than a factor of 4, with a maximum occurs at 10% and the elongation at break between 18.9 and 60%.
  • Table 1 Composition of the investigated alloy (all data in wt.
  • Example 1 Room temperature tensile tests for hot rolling and heat treatments in Example 1 for a 12 mm thick sheet.
  • R Reference according to the prior art
  • E According to the invention
  • Figure 8 Microstructure of batch 161995 in Example 5 after annealing at

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Abstract

L'invention concerne un procédé de production d'un élément à partir d'un alliage fer-chrome précipitant des phases de Laves et/ou des particules contenant du Fe et/ou des particules contenant du Cr et/ou des particules contenant du Si et/ou du carbure, un semi-produit réalisé dans l'alliage étant soumis à un traitement thermomécanique. Au cours d'une première étape, l'alliage est recuit en solution à des températures ≥ à la température de recuit en solution, puis refroidi dans du gaz de protection ou de l'air statiques, du gaz de protection ou de l'air en mouvement (soufflé) ou dans de l'eau, au cours d'une deuxième étape, une déformation mécanique du semi-produit de l'ordre de 0,05 à 99% est effectuée et, au cours d'une étape suivante, les phases de Laves Fe2(M, Si) ou Fe7(M, Si)6 et/ou les particules contenant du Fe et/ou les particules contenant du Cr et/ou les particules contenant du Si et/ou le carbure répartis de manière ciblée et fine sont précipités du fait que l'élément fabriqué à partir du semi-produit déformé est amené par un chauffage de 0,1 °C/min à 1000 °C/min à une température d'application comprise entre 550 °C et 1000 °C.
EP10760237.7A 2009-09-01 2010-08-18 Procédé de fabrication une alliage de fe-cr Active EP2480695B1 (fr)

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DE102009039552A DE102009039552B4 (de) 2009-09-01 2009-09-01 Verfahren zur Herstellung einer Eisen-Chrom-Legierung
PCT/DE2010/000975 WO2011026460A1 (fr) 2009-09-01 2010-08-18 Procédé de production d'un alliage fer-chrome

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JP6881119B2 (ja) * 2017-07-14 2021-06-02 大同特殊鋼株式会社 フェライト系ステンレス鋼及び耐熱部材
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CA2773708C (fr) 2015-03-17
CA2773708A1 (fr) 2011-03-10
DE102009039552A1 (de) 2011-03-24
KR101476753B1 (ko) 2014-12-26
DE102009039552B4 (de) 2011-05-26
JP2013503265A (ja) 2013-01-31
EP2480695B1 (fr) 2018-05-23
AU2010291651B2 (en) 2016-04-28
WO2011026460A1 (fr) 2011-03-10
US20120145285A1 (en) 2012-06-14
KR20120061851A (ko) 2012-06-13
CN102471817A (zh) 2012-05-23

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