DE102009039552B4 - Process for producing an iron-chromium alloy - Google Patents

Abstract

A process for the production of a component, of an iron-chromium alloy which precipitates Laves phases and / or Fe-containing particles and / or Cr-containing particles and / or Si-containing particles and / or carbides, by using an alloy produced from the alloy In a first step, the alloy is solution-annealed at temperatures ≥ the solution annealing temperature, preferably ≥ 1050 ° C for more than 7 minutes followed by cooling in inert gas or air, moving (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 performed 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 dispersed thereby excreted, that the manufactured from the formed semi-finished component by a heating is brought from 0.1 ° C / min to 1000 ° C / min to an application temperature between 550 ° C and 1000 ° C.

Description

The invention relates to a molten metallurgically produced ferritic iron-chromium alloy.

Of the DE 100 25 108 A1 is a high temperature material comprising comprising a chromium oxide-forming iron alloy with up to 2 wt .-% of at least one oxygen-affinity element from the group Y, Ce, Zr, Hf and Al, up to 2 wt .-% of an element 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 is the electrical conductivity of oxides increased on Cr basis. 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.

The 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.

Common to both alloys is low heat resistance and insufficient creep resistance for temperatures above 700 ° C. Especially in the range above 700 ° C to about 900 ° C, however, these alloys have excellent resistance to oxidation and corrosion.

By the DE 10 2006 007 598 A1 a creep-resistant ferritic steel has become known comprising precipitates of an intermetallic phase of the type Fe ₂ (M, Si) or Fe ₇ (M, Si) ₆ with at least one metallic alloying element M formed by the elements niobium, molybdenum, tungsten or tantalum can be. The steel should preferably be used for a bipolar plate in a fuel cell stack.

By the EP 1 536 031 A1 a metallic material for fuel cells has become known including 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 of the group Sc, Y, La, Ce, Pr, Nd, Pm, Sn, Zr and Hf ≤ 1%, the remainder being iron and unavoidable impurities, the composition satisfying the following equation: 0.1 ≤ Mo / Nb ≤ 30.

The 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%, Al max. 0.1%, N max. 0.07%, remainder Fe and impurities caused by melting, the content of Zr + Ti being at least 0.2%.

All these alloys have compared to DE 100 25 108 A1 and EP 1 298 228 A2 an improved heat resistance and increased creep resistance at temperatures above 700 ° C, by the formation of precipitates, which hinders the dislocation movements and thus the plastic deformation of the material. These excretions exist z. B. in the case of DE 10 2006 007 598 A1 from a Laves phase, an intermetallic compound having the composition Fe ₂ (M, Si) or Fe ₇ (M, Si) ₆ , where M may be niobium, molybdenum, tungsten or tantalum. In this case, a volume fraction of 1 to 8%, preferably 2.5 to 5% should be achieved. But there may also be other precipitates such as Missing particles and / or Cr-containing particles and / or Si-containing particles, such as. B. in the EP 1 536 031 A1 be described, or carbides with Nb, W, Mo. All these particles have in common that they complicate the deformation of the material.

It is known from the prior art described above that minor additions of Y, Zr, Ti, Hf, Ce, La and like reactive elements can greatly enhance the oxidation resistance of Fe-Cr alloys.

The in the DE 10 2006 007 598 A1 , of the EP 1 536 031 A1 and the EP 1 882 756 A1 These alloys are optimized for use as an interconnector plate for the high-temperature fuel cell: By using a ferritic alloy with 10 to 40% chromium, they have an expansion coefficient adapted to the ceramic components anode and electrolyte as far as possible.

Further requirements for the interconnector steel of a high-temperature fuel cell are, in addition to the above-mentioned creep resistance, a very good corrosion resistance, a good conductivity of the oxide layer and a low chromium evaporation.

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.

In the DE 10 2006 007 598 A1 For example, 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. By stably bonding Si in the Laves phase Fe ₂ (M, Si) or Fe ₇ (M, Si) ₆ , no non-conductive underlayer of silicon oxide also forms under the chromium oxide topcoat. The corrosion resistance is further improved by keeping the Al content low and thus avoiding the increase of corrosion by the internal oxidation of the aluminum. A low addition of Ti additionally solidifies the surface and prevents bulging of the oxide layer and the inclusion of metallic areas 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.

There are increased market demands for products that require increased hot strength and creep strength at an elongation of at least 20% at application temperature to avoid brittle failure with at least equal oxidation or corrosion resistance and a higher use temperature of the alloy, among others Maintaining an acceptable ductility, measured as a tensile plastic strain with an elongation> 13% at room temperature.

Furthermore, the following examination methods are used.

In a creep test, 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. In the creep test, the time to break t _B - the break time - of the sample is measured in the simplest case. 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 _p0.2 , the tensile strength R _M and the elongation to breakage are determined. Elongation A is determined on the fractional sample from the extension of the original test section L ₀ : A = (L _U - L ₀ ) / L ₀ 100% = ΔL / L ₀ 100%

With L _u = measuring length after the break.

Depending on the measuring length, the elongation at break is provided with indices:
A ₅ , measuring length L ₀ = 5 · d ₀ or L ₀ = 5.65 · √S ₀
A ₁₀ , measuring length L ₀ = 10 · d ₀ or L ₀ = 11.3 · √S ₀
or z. B. A _{L = 100} , for the freely selected measuring length L = 100 mm.
(d ₀ initial diameter, S ₀ initial cross-section of the flat sample)

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. When etching with V2A pickling, the grains or grain boundaries are additionally visibly etched. When viewed in a light microscope, only particles with a size of approx. 0.5 μm are visible. Smaller ones can not be recognized, but they do exist. Therefore, metallography is only used as an aid in explaining that the effectiveness of a measure is conveniently evaluated by break time or creep resistance.

In the handbook of high-temperature material technology, Ralf Bürgel, 3rd revised edition, Viehweg Verlag, December 2006, hereinafter referred to as "Bürgel", the "Possible measures to increase the creep resistance of metallic materials" are shown on pages 196 to 199 and in Table 3.7.

• – „Hoher Schmelzpunkt, kubischflächenzentrierter. Werkstoff”,
• – „hoher E-Modul”,
• – „Werkstoff mit niedriger Stapelfehlerenergie”,

lassen sich nicht zur Verbesserung der angesprochenen Parameter verwenden, da sie einen Wechsel des Werkstofftyps erfordern, was hier nicht möglich und auch nicht die Aufgabe ist.The measures

• - "High melting point, cubic surface centered. Material",
• - "high modulus of elasticity",
• - "material with low stacking fault energy",

can not be used to improve the mentioned parameters, since they require a change of the material type, which is not possible here or the task.

• – „Mischkristallhärtung,
• – „Teilchenhärtung”,
• – „hoher Teilchenvolumenanteil”,
• – „Teilchen mit geringem Diffusionskoeffizienten des betreffenden Legierungselementes”

sind in der DE 10 2006 007 598 A1 und/oder der EP 1 536 031 A1 und/oder der EP 1 882 756 A1 schon angewandt worden.The measures

• - "solid solution hardening,
• - "particle hardening",
• - "high particle volume fraction",
• - "particles with low diffusion coefficient of the alloying element in question"

are in the DE 10 2006 007 598 A1 and / or the EP 1 536 031 A1 and / or the EP 1 882 756 A1 already applied.

• – „Teichen mit geringer Löslichkeit in der Matrix”,
• – „kohärente Teilchen mit geringer Grenzflächenenthalpie zur Matrix”,

gelten nicht für die betrachten Ausscheidungen.The measures

• - "Ponds with low solubility in the matrix",
• - "coherent particles with low interfacial enthalpy to the matrix",

do not apply to the considered excretions.

• – „Karbide oder Boride als Korngrenzenausscheidungen; Oxide und Sulfide vermeiden”,
• – „Positiv wirksame Korngrenzenelemente in genau kontrollierter Dosierung zugeben z. B. B, C, Zr, Ce”,
• – „höherer Reinheitsgrad der Legierung”,
• – „Getterelemente (z. B. für S) zugeben”,
• – „hohe Korrosionsbeständigkeit”

schon in der DE 10 2006 007 598 A1 und/oder der EP 1 536 031 A1 und/oder EP 1 882 756 A1 beschrieben worden.Likewise, the measures

• - "carbides or borides as grain boundary precipitates; Avoid oxides and sulphides ",
• - "Add positively effective grain boundary elements in precisely controlled dosage z. B, C, Zr, Ce ",
• - "higher purity of the alloy",
• - "add getter elements (eg for S)",
• - "high corrosion resistance"

already in the DE 10 2006 007 598 A1 and / or the EP 1 536 031 A1 and or EP 1 882 756 A1 been described.

• – „geringe Dendritenarmabstände bei Gussgefügen”,
• – „gerichtete Kornstruktur in Hauptbelastungsrichtung”,
• – „Einkristall”,
• – „geringe Dichte bei eigengewichtbelasteten u. rotierenden Bauteilen”

lassen sich auf diesen Legierungstyp, bzw. auf den Herstellungsweg, bzw. die Verwendung nicht anwenden.The measures

• - "low dendrite arm distances in casted structures",
• - "directional grain structure in the direction of main load",
• - "single crystal",
• - "low density at self-weighted u. rotating components "

can not be applied to this type of alloy or to the production route or use.

• 1) „Grobkorngefüge”,
• 2) „Verzahnung der Korngrenzen durch Ausscheidungen”,
• 3) „Optimierte Wärmebehandlung (Optimalen Teilchendurchmesser einstellen, Seigerungen bei Gussgefüge beseitigen, evtl. Korngrenzenrauhigkeit gezielt einstellen)”,
• 4) „Kaltverformung vermeiden”

zu betrachten.The task is to improve the creep resistance of the precipitation-hardened iron-chromium alloy

• 1) "coarse grain texture",
• 2) "intermeshing of grain boundaries by precipitates",
• 3) "Optimized heat treatment (adjust the optimum particle diameter, eliminate segregations in the cast structure, adjust the grain boundary roughness if necessary)",
• 4) "avoid cold deformation"

consider.

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 ductility at room temperature, the high heat resistance or creep resistance of a precipitation-hardened ferritic alloy over the prior art further increase.

In addition, a thermomechanically treated made 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.

Finally, the component / semifinished product thus produced for specific technical applications in the temperature range above 550 ° C should be used.

This object is achieved on the one hand by a method for producing a, Laves phases and / or Fe-containing particles and / or Cr-containing particles and / or Si-containing particles and / or carbides ausscheidenden, iron-chromium alloy by a 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, preferably ≥ 1050 ° C for more than 7 minutes followed by cooling in inert gas or air, moving (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 Laues 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 deliberately and finely distributed thereby excreted, that from the formed half manufactured manufactured component by heating with from 0.1 ° C / min to 1000 ° C / min to an application temperature between 550 ° C and 1000 ° C.

On the other hand, 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 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, preferably ≥ 1050 ° C for more than 7 minutes, followed by cooling in inert gas or air, moving (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 performed 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 deliberately and finely distributed thereby excreted that um shaped semifinished / component is subjected to a heat treatment in the temperature range between 550 and 1060 ° C under inert gas or air for a time between t _min and t _max , followed by cooling in inert gas or air, moving (blown) inert gas or air or in water or for heat treatments up to 800 ° C in the oven, wherein t _{min is of} the formula log 10 = -0.0038 T + 3 and t _max log 10 = -0.0069 T + 8.1, where the desired component is before or after this Heat treatment is made.

The data t _min and t _max are given in minutes, the heat treatment temperature in ° C. Log 10 is the base 10 logarithm.

Advantageous developments of the method according to the invention can be found in the associated subclaims.

In addition, the object is also achieved 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 rest and the usual melting impurities, which has a deformed microstructure following a thermomechanical treatment, such that Laves phase (s) is embedded in finely divided form in the microstructural dislocations of the microstructure, wherein in a creep with 35 MPa at 750 ° C and at an elongation of at least 20% a fracture time is set in the microstructure that exceeds the fracture time of a coarse-grained, fully recrystallized microstructure by a factor of at least 1.5.

Now consider the measures 1 to 4 described above.

Surprisingly, it has been shown that in contrast to the measure 4 "cold deformation" a pre-shaping followed by a matched annealing can cause extensions of the fracture times of the sample in the creep, which is more than 1.5 times, preferably more than 3 times over the break times for a coarse grain structure (measure 1) go beyond.

Furthermore, it is proposed that for the third step - the precipitation of the Laves phase (s) - the formed semi-finished product or possibly the component produced 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 inert gas or air, followed by cooling in inert gas or air, moving (blown) inert gas or air or in water or for heat treatments up to 800 ° C in the furnace are cooled, then, if necessary, the desired component is manufactured, wherein t _{min of} the formula log10 = -0.0038 T + 3 and t _max log10 = -0.0069 T + 8 , 1 corresponds.

In addition, there is the possibility that for the third step - the precipitation of the Laves phase - the semifinished product or the component made therefrom for a time between t _min and t _max a heat treatment in the temperature range between 550 and 1060 ° C under inert gas or air , followed by cooling in quiescent inert gas or air, agitated (blown) inert gas or air or in water or for heat treatment to 800 ° C in the oven are cooled, then, if necessary, the desired component is manufactured, wherein t _{min of} the formula log10 = -0.0038 T + 3 and t _max log10 = -0.0069 T + 8.1 and then the finished component by heating at 0.1 ° C / min to 1000 ° C / min to an application temperature between 550 ° C and 1000 ° C is brought.

According to a further aspect of the invention, a semi-finished product made of 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% S Max. 0.01% Fe rest as well as the usual smelting-related impurities.

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.

It is of particular advantage that after the solution annealing at temperatures ≥ the solution annealing temperature preferably ≥ 1050 ° C for more than 7 minutes followed by cooling in inert gas or air, moving (blown) inert gas or air or in water at the initial state prior to deformation in the Semi-finished only a 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 the alloy are present.

The deformation of the semifinished product can be done by hot working. Alternatively, however, the deformation can also be brought about by cold working.

In the first case, the semi-finished product with a starting temperature> 1070 ° C is thermoformed, the last 0.05 to 95% of mechanical deformation between 1000 ° and 500 ° C are applied, advantageously the last 0.5 to 90% between 1000 ° C and 500 ° 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%.

According to a further aspect of the invention, it is proposed that the mechanical deformation of the semi-finished product is 20 to 99% and then subjected to the transformed semi-finished product for a period between t _min and t _max a heat treatment in the temperature range between 950 ° and 1060 ° C under inert gas or air is followed by cooling in inert gas or air, moving (blown) inert gas or air or in water and then the desired component is manufactured, with log10 (t _min) = -0.0038 T + 3 log10 (t _max ) = -0.0069 T + 8.1 wherein
t _min and t _max in minutes
the heat treatment temperature T in degrees Celsius
log 10 is the base 10 logarithm.

If the 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.

For other applications such. B. 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 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.

For the applications, which neither need a conductive oxide nor have special requirements for the chrome evaporation, 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, when an oxide layer consisting of a chromium oxide and a chromium-magan spinel, sufficient oxidation resistance can be ensured by using 0.001-0.5% of aluminum. Is a higher oxidation resistance required, as z. B. is ensured 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 following elements can additionally be introduced individually or in combination into the alloy already mentioned: 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 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%.

In addition, the alloy (in% by weight) of one or more of Ce, La, Pr, Ne, Sc, Y, Zr or Hf may contain 0.02-0.3% by weight.

If desired, the alloy (in% by weight) of one or more of Ce, La, Pr, Ne, Sc, Y, Zr or Hf may contain 0.02-0.2% by weight.

To achieve the desired effects, the Nb content is 0.3 to 1.0% and the Si content is 0.15 to 0.5%.

If necessary, the element tungsten may be wholly or partly replaced by at least one of the elements Mo or Ta.

If necessary, 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%.

If necessary, the alloy can also max. Containing 0.003% boron.

In addition, 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.

Alternatively, it is also possible to use the component / semifinished product produced by the method according to the invention, or the alloy itself, 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.

As a result of the thermomechanical treatment, in the case of alloys produced by fusion metallurgy, Laves phases can be excreted selectively and finely distributed at the dislocations of the microstructure by means of the method according to the invention.

The details and advantages of the invention will be more apparent from the following examples.

In the following, the method steps according to the invention are subjected to a closer examination.

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 ausschenden iron-chromium alloy in the thermomechanical treatment must be an annealing above the solution annealing temperature, in order to bring the Laves phases and / or Fe-containing particles and / or Cr-containing particles and / or Si-containing particles and / or carbides into solution and to precipitate for the subsequent thermomechanical treatment To be available. The solution annealing temperature is alloy-dependent, but is preferably above 1050 ° C for more than 7 minutes, followed by cooling in quiescent inert gas or air, agitated (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. 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 few Laves phase forms new.

In addition, it must be ensured that, in particular with thick-walled components, all parts of the component achieve the required minimum annealing time at the specified temperature. This must be taken into account when determining the starting point of the glow time.

In a second step, 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 transformation, 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-containing particles and / or Cr-containing particles and / or Si -containing particles and / or carbides provide.

This deformation can be a cold working, but also a hot working, whereby during the hot working it has to be ensured that the structure does not completely recrystallise during the 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 transformation, 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.

Excretions below a certain size are less effective (see eg "Bürgel", page 141). Therefore, the dislocation density generated by the deformation must not be too high, because then very many but too weak precipitates arise and the excess dislocations can move freely and thus the pre-deformation is harmful. Ie. Preferably, the highest deformation is 90% for the part of hot working ≤ 1000 ° C and 90% for cold working.

With a transformation in the range of 20 to 99%, an annealing between 950 ° C and 1050 ° C can cause a recovery of the structure. As a result, the dislocation density is reduced, so that again the positive effect on the distribution of the Laves phases and / or Fe-containing particles and / or Cr-containing particles and / or Si-containing particles and / or carbides adjusts.

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. During the heating, 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. For this, 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 _min and t _max , followed by cooling in inert gas or air, moving (blown) inert gas or Air or in water or for heat treatments up to 800 ° C in the oven, wherein log10 (t _min) = -0.0038 T + 3 log10 (t _max ) = -0.0069 T + 8.1

Specification of t _min and t _max in minutes and heat treatment temperature T in ° C, log 10 is the logarithm base 10, wherein the desired component can be made before or after this heat treatment.

In the case of annealing, it must be ensured that, in particular for thick-walled semi-finished products / components, all parts of the component achieve the required minimum annealing time at the specified temperature. This must be taken into account when determining the starting point of the glow time. Likewise, care must be taken that no area of the semifinished product / component exceeds the required maximum annealing time

Times shorter than t _{min are} sufficient to form the Laves phases and / or Fe-containing particles and / or Cr-containing particles and / or Si-containing particles and / or carbides. At times longer than _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. At the same time there is the danger of an excessive coarsening of the precipitates. At the lower temperatures in the range of 550 ° C and 1060 ° C, times above t _{max are} uneconomical.

The annealing can take place under 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% chromium, the workability of the material and its serviceability is impaired by increased formation of embrittling phases, in particular the sigma phase. The chromium content is therefore limited to ≤ 30%. As the chromium content increases, the expansion coefficient decreases.

In particular for use in a fuel cell, 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%. For other applications z. However, for reformers or in power plants, this restriction does not exist.

The addition of 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 adversely 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. This is the preferred range of Nb 0.1-2% W: 0.1-5%

W can also be replaced in whole or in part 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 completely tied up in 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. However, 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 z. B. for use as an interconnector plate. However, a certain aluminum content of at least 0.001% is necessary for deoxidizing the melt. If no conductive oxide is required and the demand for a still significantly higher oxidation resistance than that given by a chromium oxide layer is demanded, the alloy can form a closed aluminum oxide layer with an aluminum content of at least 2% ( DE 101 205 61 ). Aluminum contents above 6.0% lead to processing problems and thus to increased costs

Carbon leads to carbide precipitations and thus contributes to creep resistance. 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 hardening of the surface, thus preventing bulging of the oxide layer and the inclusion of metallic areas during 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 .-%.

It has been found that cobalt contents from 3% affect the oxidation resistance. The lower limit is set at 0.01 wt .-% for cost reasons. For nickel and copper, the same applies as for cobalt, boron is limited to max. 0.005% limited because this element reduces the oxidation resistance.

Based on embodiments of the subject invention will now be explained in more detail.

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 depending on the final thickness with Intermediate annealing cold rolled. After hot rolling, the oxide layer was removed by pickling. A material with an analysis as shown in Table 1 precipitates mainly Laves phases Fe ₂ (M, Si) or Fe ₇ (M, Si) ₆ and in significantly reduced amounts of carbides.

example 1

In this example, 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 quiescent air cooling to 12 mm thick sheet, the mechanical working with a starting temperature> 1070 ° C was begun 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. In the case of 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.

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 microns (Figure 2), which is a typical coarse grained Structure is.

If a creep test is carried out on this material as described above with an initial stress of 35 MPa at a temperature of 750 ° C., the sample breaks after 12.8 hours at an elongation A of 69.8%. (Table 2). At room temperature, the annealed material has an elongation of 35%, which is a very good value for a ferrit.

If, however, from 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 specimen ruptures only after 255 hours at an elongation A of 29%, which means an increase in the time to break 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.

This example clearly shows that the microstructure is superior to the pre-forming and the coarse-grained microstructure in terms of fracture time or creep resistance, which contradicts the state of the art as described in "Bürgel" pages 196 to 199 Table 3.7.

Example 2

In this example, with the hot rolled material of Example 1, anneals between 600 ° C and 1000 ° C for 20 min each, (for t _min and t _max according to Equation 1 and 2 see Table 3) in air, followed by cooling dormant air performed. After the heat treatment, samples were prepared from the sheet, and then the creep test was carried out at a tension of 35 MPa at 750 ° C as described above. The results are summarized in Table 3.

At 1000 ° C only a break time of 10.4 h is achieved at an elongation A of 79.5%. At 600 ° C to 950 ° C, an at least 10-fold increased fracture time of over 100 hours is achieved at an elongation A of greater than 22.7%. The highest break time is achieved after annealing at 650 ° C for 20 minutes with 286 hours. Figure 3 shows the structure after the different annealing. The microstructures in Figure 3 are not globularly recrystallized. Up to 850 ° C, 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 case of the etched sections by means of electrolytic etching with oxalic acid, it can be clearly seen that the Laves phase is excreted microscopically visible from about 750 ° C. From about 900 ° C, it is noticeable in addition to the precipitates in the grain and noticeable on the small-angle grain boundaries or grain boundaries. and thus ensures a toothing of the small-angle grain boundaries or grain boundaries, which corresponds to measure 2 for increasing the creep resistance (see above). At 1000 ° C., noticeably very large grains have been formed as a result of progressive recovery, so that the dislocation density is greatly reduced and thus no increase in the breaking time occurs any more.

At room temperature, the treated sheet at all temperatures between 600 ° C and 950 ° C 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. At the higher temperatures, a larger volume fraction is eliminated, but a bit coarser and recognizable on the small-angle grain boundaries and grain boundaries. The annealing temperatures between 600 ° C and 700 ° C have the advantage that simultaneously reaches the time to break in the creep test at about 650 ° C a maximum.

The annealing at 1000 ° C is with an annealing time of 20 minutes above t _max = 16 minutes. It is therefore not in the area of invention and serves as a reference. 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 _min and t _max . Thereafter, according to the invention, the fracture time was markedly increased by more than a factor of 9 compared to the coarse-grained, globularly recrystallized state of Example 1, which is produced after annealing at 1075 ° C. for 20 minutes followed by quenching in still air.

Example 3

In this example, 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. a break time of 391 hours at an elongation A of 38% (Table 4). The structure is not globular recrystallized but recovered. It has precipitates in the grain and on the small-angle grain boundaries or grain boundaries (Fig. 4). The fracture time is 30 times the time, which was achieved in Example 1 after annealing of 1075 ° C for 20 min with a globular recrystallized coarse grain with a grain size of 137 microns. The annealing at 920 ° C is with an annealing time of 28 minutes in the range of the invention between t _min = 0.32 minutes and t _max = 56 minutes.

At room temperature, 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 well formable.

Example 4

In this example, 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 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. A part of the material was then again annealed at 1050 ° C for 20 min in air with subsequent cooling in still air. Thereafter, both batches are globularly recrystallized, batch 161061 with a particle size of 134 μm (FIG. 6) and batch 161995 with a particle size of 139 μm. Eliminated Laves phase is only slightly to find.

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%.

Similarly, charge 161995 has a break time of 33.5 hours 89% in a creep test at 750 ° C with an initial load of 35 MPa, after the additional annealing at 1075 ° C for 20 minutes, the produces a very coarse grain, a fracture time of only 7.9 hours with a strain A of 92%. The elongation in the tensile test at room temperature is 1050 ° C and 3.4 minutes. Glow time for batch 161061 at 28% and batch 161995 at 26% for a ferritic very good, which makes the material very good formability. It becomes even higher with the coarse-grained texture for batch 161061 at 31% and batch 161995 at 29%.

This shows the influence of the annealing time at temperatures around 1050 ° C. For short-term annealing of a few minutes, dislocations (deformation) and sufficient Laves phase are still present in the material, which in this example results in a 3 to 4 times longer time to break in the creep test. For longer anneals, the Laves phase dissolves sufficiently, as shown in batch 161061, and the microstructure recrystallizes globular with correspondingly short times to break during the creep test.

The annealing at 1050 ° C for 20 min is with an annealing time of 20 minutes above t _max = 7.2 minutes. It is therefore not in the invention range and serves as a reference as well as the annealing at 1075 ° C for 20 min. The annealing at 1050 ° C. for 3.4 minutes is with an annealing time of 3.4 minutes in the range according to the invention between t _min = 0.1 minutes and t _max . = 7.2 minutes and shows according to the invention a significantly increased break time in the creep test.

Example 5

In this example, 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.

If a solution annealing is then carried out at 1075 ° C. for 10 minutes at 1075 ° C. with cooling in still air, a very coarse-grained microstructure is obtained with only a few precipitates of Laves phase and a particle size of about 134 to 162 μm ( Picture 9). If a creep test is carried out on this material with an initial stress of 40 MPa at a temperature of 700 ° C., the sample breaks after 228 hours at an elongation A of 51%. (Table 6) If the creep test is carried out at 60 MPa, the sample breaks after 8.1 hours, with an elongation A of 43%. At room temperature, the annealed material has an elongation of 35%, which is a very good value for a ferrit.

If, in addition, annealing at 700 ° C. for 4 hours at 700 ° C. with subsequent cooling in inert air is carried out on the material which has been solution-annealed at 1075 ° C. for 22 minutes, the Laves phase is distributed in the microstructure. (Picture 10). If the creep test is then carried out at 700 ° C. with an initial stress of 40 MPa, 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 ° C. 22 minutes. If the creep test is carried out with 60 MPa, the sample breaks after 6.3 hours, with an elongation A of 63%, ie also after a much shorter time than after the solution annealing of 1075 ° C. for 22 minutes.

This is the proof that the precipitation of the Laves phase (s) must take place in a structure with increased dislocation density, ie a reshaped or recovered structure, in order to obtain an extension of the fracture time. An excretion into a solution-annealed structure causes exactly the opposite namely a shortening of the break time. The reason for this is the more even distribution of very fine precipitates upon excretion into a structure with increased dislocation density, d. H. a deformed or recovered microstructure in comparison to an excretion into a dislocation-rich coarse-grained microstructure.

The labels / descriptions of the tables / images are shown as follows:
Table 1 Composition of the tested alloy (all figures in% by weight) Table 2 Results of the creep tests at 750 ° C and the tensile 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 prior art, E: According to the invention)
Table 3 Results of the creep tests at 750 ° C and the tensile tests at room temperature for the hot rolling of Example 1 and the heat treatment of Example 2 for a 12 mm thick sheet. (R: Reference according to the prior art, E: According to the invention)
Table 4 Results of creep tests and tensile tests at room temperature for Example 3 for a 12 mm thick sheet. (R: Reference according to the prior art, E: According to the invention)
Table 5 Results of creep tests and tensile tests at room temperature for Example 4 for 1.5 mm thick tape. (R: Reference according to the prior art, E: According to the invention)
TABLE 6 Results of creep tests and tensile tests at room temperature for Example 5 on 12 mm thick sheet. (R: Reference according to the prior art, E: According to the invention)

Figure 1 Microstructure of the hot-formed material in Example 1

Figure 2 Microstructure of the hot-formed material in Example 1 after annealing at 1075 ° C for 20 minutes and cooling in still air. Grain size 137 μm.

Figure 3 Structure of the material in Example 2 after annealing between 600 ° C and 1000 ° C for 20 minutes and cooling in still air.

Figure 4 Structure of the material in Example 3 after annealing at 920 ° C in a continuous furnace in air with subsequent cooling in still air for 20 minutes and cooling in still air. (Etching V2A stain)

Figure 5 Microstructure of batch 161061 in Example 4 after annealing at 1050 ° C / 3.4 minutes under inert gas in a continuous furnace with cooling in a cold inert gas stream.

Figure 6 Microstructure of batch 161061 in Example 4 after annealing at 1050 ° C / 3.4 minutes under inert gas in a continuous furnace with cooling in a cold inert gas stream and annealing of 1050 ° C / 20 min under air with subsequent cooling in still air, grain size 134 μm (etching with V2A stain)

Figure 7 Microstructure of charge 161995 in Example 5 after annealing at 1050 ° C / 3.4 minutes under inert gas in a continuous furnace with cooling in a cold inert gas stream.

Figure 8 Microstructure of charge 161995 in Example 5 after annealing at 1050 ° C / 3.4 minutes under inert gas in a continuous furnace with cooling in a cold inert gas stream and annealing of 1075 ° C / 20 min under air with subsequent cooling in still air, grain size 139 μm

Figure 9 Structure of the hot-formed material in Example 5 after annealing at 1075 ° C for 22 minutes and cooling in still air. Grain size 134 μm. up to 162 μm

Figure 10 Microstructure of the hot worked material in Example 5 after annealing at 1075 ° C for 22 minutes, followed by quiescent air cooling followed by annealing at 700 ° C for 4 hours followed by quiescent air cooling. Grain size 136 μm. charge charge 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 Not a word 0.02 0.02 Ti 0.07 0.06 Nb 0.51 0.49 Cu 0.02 0.02 Fe rest rest 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 O 0,004

Claims (44)

A process for the production of a component, of an iron-chromium alloy which precipitates Laves phases and / or Fe-containing particles and / or Cr-containing particles and / or Si-containing particles and / or carbides, by using an alloy produced from the alloy In a first step, the alloy is solution-annealed at temperatures ≥ the solution annealing temperature, preferably ≥ 1050 ° C for more than 7 minutes followed by cooling in inert gas or air, moving (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 performed 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 dispersed thereby excreted, that the manufactured from the formed semi-finished component by a heating is brought from 0.1 ° C / min to 1000 ° C / min to an application temperature between 550 ° C and 1000 ° C. A process for the production of a component, of an iron-chromium alloy which precipitates Laves phases and / or Fe-containing particles and / or Cr-containing particles and / or Si-containing particles and / or carbides, by using an alloy produced from the alloy In a first step, the alloy is solution-annealed at temperatures ≥ the solution annealing temperature, preferably ≥ 1050 ° C for more than 7 minutes followed by cooling in inert gas or air, moving (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 performed 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 formed semi-finished product for a time between t _min and T _{max is} subjected to a heat treatment in the temperature range between 550 and 1060 ° C under inert gas or air, followed by cooling in inert gas or air, moving (blown) inert gas or air or in water or for heat treatments to 800 ° C in the oven to be cooled in which t _min corresponds to the formula log 10 = -0.0038 T + 3 and t _max log 10 = -0.0069 T + 8.1, wherein the desired component is manufactured before or after this heat treatment. Process according to claim 1 or 2, wherein semifinished product of the following composition (in% by weight) is treated thermomechanically: Cr 12-30% Mn 0.001-2.5% Nb 0.1-2% W 0.1-5% Si 0.05-1% C 0.002 to 0.01% N 0.002 to 0.01% S Max. 0.01% Fe rest and the usual melting-related impurities, wherein at room temperature, a mechanical deformability, measured as a plastic strain in the tensile test> 15% is given. Method according to one of claims 1 to 3, characterized in that after the solution annealing at temperatures ≥ the solution annealing temperature, preferably ≥ 1050 ° C for more than 7 minutes followed by cooling in inert gas or air, moving (blown) inert gas or air or in water in the initial state before deformation in the semifinished 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 and / or or Si-containing particles and / or carbides are present in the alloy. Method according to one of claims 1 to 4, characterized in that the deformation of the semifinished product is carried out by hot forming. Method according to one of claims 1 to 5, characterized in that the hot working of the semifinished product begins with a starting temperature> 1070 ° C, wherein the last 0.05 to 90% mechanical deformation between 1000 ° C and 500 ° C are applied. Method according to one of claims 1 to 6, characterized in that hot working of the semifinished product begins with a starting temperature> 1070 ° C, wherein the last 0.05 to 95% mechanical deformation between 1000 ° C and 500 ° C are applied. Method according to one of claims 1 to 7, characterized in that the hot working of the semifinished product starts with a starting temperature> 1070 ° C, wherein the last 0.05 to 90% mechanical deformation between 1000 ° C and 500 ° C are applied. Method according to one of claims 1 to 8, characterized in that adjoins the hot working of the semifinished product cold forming. Method according to one of claims 1 to 4, characterized in that the deformation of the semifinished product takes place by a cold forming. A method according to claim 10, characterized in that the degree of cold working of the semifinished product is 0.05 to 99% Method according to one of claims 10 or 11, characterized in that the cold working of the semifinished product is 0.05 to 95%. Method according to one of claims 10 to 12, characterized in that the cold working of the semifinished product is 0.05 to 90%. Method according to claims 1 to 13, characterized in that the mechanical deformation of the semi-finished product is 20 to 90% and then the formed semi-finished product for a period between t _min and t _{max of} a heat treatment in the temperature range between 950 ° and 1060 ° C under inert gas or Air is subjected, followed by a cooling in inert gas or air, moving (blown) inert gas or air or in water and then the desired component is manufactured with log10 (t _min) = -0.0038 T + 3 log10 (t _max ) = -0.0069 T + 8.1 Specifying t _min and t _max in minutes and heat treatment temperature T in ° C, log 10 is the base 10 logarithm. A method according to any one of claims 1 to 14, wherein the alloy additionally contains (by weight) 0.02 to 0.3% La. A method according to any one of claims 1 to 15, wherein the alloy additionally contains (in wt%) 0.01 to 0.5% Ti. A method according to any one of claims 1 to 16, wherein the alloy additionally contains 0.02 to 0.3% of one or more of Ce, La, Pr, Ne, Sc, Y, Zr or Hf. A method according to any one of claims 1 to 17, wherein the alloy additionally contains (in wt%) 0.001 to 0.5% Al. A method according to any one of claims 1 to 17, wherein the alloy additionally contains (in wt%) 2.0 to 6.0% Al. The method of claim 19, wherein the alloy additionally contains (in wt%) 2.5 to 5.0% Al. A method according to any one of claims 1 to 20, wherein the alloy additionally contains one or more of the elements 0.0001 to 0.07% Mg, 0.0001 to 0.07% Ca, 0.002-0.03% P. The method according to any one of claims 1 to 21, wherein the alloy further contains 0.01 to 3.0% of one or more of Ni, Co or Cu. The method of any of claims 1 to 22, wherein the alloy further contains up to 0.005% B. A method according to any one of claims 1 to 23, wherein the thermo-mechanically treated and Laves phases in finely divided form precipitating iron-chromium alloy comprises the following composition (in 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% Fe rest as well as the usual smelting-related impurities. A method according to any one of claims 1 to 24, wherein the alloy additionally contains (in wt%) 0.02-0.2% of the element La. A method according to any one of claims 1 to 25, wherein the alloy additionally contains (in wt%) 0.02-0.2% Ti. A method according to any one of claims 1 to 26, wherein the alloy additionally contains (in wt%) 0.02-0.2% of one or more of Ce, Pr, Ne, Sc, Y, Zr or Hf. A method according to any one of claims 1 to 27, wherein the alloy additionally comprises one or more of the elements (in wt%) 0.0001-0.05% Mg, 0.0001-0.03% Ca, 0.002-0.03 Contains% P The method of any one of claims 1 to 28, wherein the alloy further contains (in wt%) up to 0.003% B. A method according to any one of claims 1 to 29, wherein (in wt%) the Nb content is 0.3 to 1.0% and the Si content is 0.15-0.5%. Method according to one of claims 1 to 30, wherein the W is replaced wholly or partly by at least one of the elements Mo and / or Ta. Method according to one of claims 1 to 31, wherein the alloy (in wt .-%) max. 0.2% V and / or max. 0.005% S contains. Method according to one of claims 1 to 32, wherein the alloy (in wt .-%) max. 0.01% O contains. A method according to any one of claims 1 to 33, wherein the alloy each contains (in wt%) max 0.01 of each of Zn, Sn, Pb, Se, Te, Bi and Sb. Method according to one of claims 1 to 34, characterized in that semifinished product is formed by sheet metal, strip, rod, forging, pipe or wire. Method according to one of claims 1 to 35, characterized in that the heat treatment is carried out only after completion of the component. Method according to one of claims 1 to 36, characterized in that the thermomechanical treatment of the semi-finished product, a particularly high creep strength of the semifinished product and / or the component is generated with simultaneous elongation> 13% in a tensile test at room temperature. Metallic component or semi-finished 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 rest and the usual melting impurities, which has a deformed microstructure following a thermomechanical treatment, such that Laves phase (s) is embedded in finely divided form in the microstructural dislocations of the microstructure, wherein in a creep with 35 MPa at 750 ° C and at an elongation of at least 20% a break time h in the microstructure is set which exceeds the fracture time of a coarse-grained, fully recrystallized microstructure by a factor of at least 1.5. Metallic component or semi-finished 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 rest and the usual melting impurities, which has a deformed microstructure following a thermomechanical treatment, such that Laves phase (s) is embedded in finely divided form in the microstructural dislocations of the microstructure, wherein in a creep with 35 MPa at 750 ° C and at an elongation of at least 20% a break time h in the microstructure is set which exceeds the fracture time of a coarse-grained, fully recrystallized microstructure by a factor of at least 3. Use of a component produced according to one of claims 1 to 39 as an interconnector in a fuel cell. Use of a component produced according to one of claims 1 to 39 as material in a component, such as a reformer or a heat exchanger or in a secondary unit of a fuel cell. Use of a component produced according to one of claims 1 to 39 in the exhaust gas line of an internal combustion engine. Use of a component manufactured according to one of claims 1 to 39 for steam boilers, superheaters, turbines and other parts of a power plant or in the chemical process industry.

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