METHOD FOR PRODUCING A COMPONENT MADE OF A NICKEL-CHROMIUM-ALUMINUM ALLOY AND PROVIDED WITH WELD SEAMS

Abstract

In a method for producing, and/or installing into a system, a component with one or more weld seams containing a nickel-chromium-aluminum alloy, with (in wt. %) >18 to 33% chromium, 1.8-4.0% aluminum, 0.01-7.0% iron, 0.001-0.50% silicon, 0.001-2.0% manganese, 0.00-0.60% titanium, respectively 0.0-0.05% magnesium and/or calcium, 0.005-0.12% carbon, 0.0005-0.050% nitrogen, 0.0001-0.020% oxygen, 0.001-0.030% phosphorus, max. 0.010% sulfur, max. 2.0% molybdenum, max. 2.0% tungsten, remainder 50% nickel and impurities, the component containing semi-finished products of wrought alloy, after welding, only the weld seams and surrounding heat-affected zones undergo annealing between greater than 980 and 1250 C. for 0.05 minutes-24 hours, then cooling in inert protective atmosphere, moving protective gas or air, where: Cr+Al28 and Fp39.9 with Fp=Cr+0.272*Fe+2.36*Al+2.22*Si+2.48*Ti+0.374*Mo+0.538*W11.8*C, Cr, Fe, Al, Si, Ti, Mo, W and C being element wt. % concentrations.

Claims

1: A method for the manufacture of a component with one or more welded seams and/or for installation of a component in a plant with one or more welded seams, which comprise a nickel-chromium-aluminum alloy, containing (in mass-%) more than 18 to 33% chromium, 1.8 to 4.0% aluminum, 0.01 to 7.0% iron, 0.001 to 0.50% silicon, 0.001 to 2.0% manganese, 0.00 to 0.60% titanium, respectively 0.0 to 0.05% magnesium and/or calcium, 0.005 to 0.12% carbon, 0.0005 to 0.050% nitrogen, 0.0001-0.020% oxygen, 0.001 to 0.030% phosphorus, max. 0.010% sulfur, max. 2.0% molybdenum, max. 2.0% tungsten, the rest nickel, greater than or equal to 50% and the common process-related impurities, wherein the component is partly or completely made up of semifinished products of this nickel-chromium-aluminum wrought alloy and, after the welding, only the welded seams of this nickel-chromium-aluminum wrought alloy and the heat-affected zones surrounding the welded seams are subjected, for homogenization of the welded seams and/or for reduction of stresses, to an annealing between 98 and 1250 C. for times of 0.05 minutes to 24 hours, followed by a cooling in stationary shield gas or air, moving (blown) shield gas or air, with the consequence that the creep strength and the creep ductility of the welded seams are improved by this annealing, wherein the following relationships must be satisfied: $\begin{array}{cc}\mathrm{Cr}+\mathrm{Al}28\mathrm{and}& \left(1a\right)\end{array}$ $\begin{array}{cc}\mathrm{Fp}39.9\mathrm{with}& \left(2a\right)\end{array}$ $\begin{array}{cc}\mathrm{Fp}=\mathrm{Cr}+0.272*\mathrm{Fe}+2.36*\mathrm{Al}+2.22*\mathrm{Si}+2.48*\mathrm{Ti}+0.374*\mathrm{Mo}+0.538*W-11.8*C& \left(3a\right)\end{array}$ wherein Cr, Fe, Al, Si, Ti, Mo, W and C are the concentrations of the elements in question in mass-%.

2: The method according to claim 1, wherein the component contains welded seams and, after the welding, the entire component containing the welded seams is subjected, for homogenization of the welded seam and/or for reduction of stresses, to a further annealing between higher than 980 and 1250 C. for times from 0.05 minutes up to 24 hours, followed by a cooling in stationary shield gas or air, moving (blown) shield gas or air or in water, with the consequence that the creep strength and the creep ductility of the welded seams are improved by this.

3: The method according to claim 1, wherein, after a process of grinding of the welded seam and of the heat-affected zone, it is advantageous when roughness values Ra of 0.01 to 15 m are attained, since this improves the corrosion resistance and especially the metal dusting resistance and raises them almost to the value of the parent material.

4: The method according to claim 1, wherein the semifinished product has a grain size of 30 to 600 m.

5: The method according to claim 1, with a chromium content of 20 to 33%.

6: The method according to claim 1, with an aluminum content of 1.8 to 3.2%.

7: The method according to claim 1, with an iron content of 0.01 to 4.0%.

8: The method according to claim 1, if necessary with a content of niobium of 0.0 to 1.1%, wherein the formula (4a) is supplemented by a term for Nb: $\begin{array}{cc}\mathrm{Fp}=\mathrm{Cr}+0.272*\mathrm{Fe}+2.36*\mathrm{Al}+2.22*\mathrm{Si}+2.48*\mathrm{Ti}+1.26*\mathrm{Nb}+0.374*\mathrm{Mo}+0.538*W-11.8*C& \left(3b\right)\end{array}$ and Cr, Fe, Al, Si, Ti, Nb, Mo, W and C are the concentrations of the elements in question in mass-%.

9: The method according to claim 1, optionally with a content of zirconium of 0.0 to 0.20%.

10: The method according to claim 1, optionally with an yttrium content of 0.001 to 0.20%.

11: The method according to claim 1, optionally with a lanthanum content of 0.001 to 0.20%.

12: The method according to claim 1, optionally with a cerium content of 0.001 to 0.20%.

13: The method according to claim 1, optionally with a content of cerium mixed metal of 0.001 to 0.20%.

14: The method according to claim 1, optionally with a content of hafnium of 0.001 to 0.20%.

15: The method according to claim 1, optionally with a content of tantalum of 0.001 to 0.60%.

16: The method according to claim 1, optionally with a content of boron of 0.0001 to 0.008%.

17: The method according to claim 1, further optionally containing 0.0 to 5.0% cobalt.

18: The method according to claim 1, further optionally containing at most 0.5% copper, wherein the formula (4a) is supplemented by a term for Cu: $\begin{array}{cc}\mathrm{Fp}=\mathrm{Cr}+0.272*\mathrm{Fe}+2.36*\mathrm{Al}+2.22*\mathrm{Si}+2.48*\mathrm{Ti}+0.477*\mathrm{Cu}+0.374*\mathrm{Mo}+0.538*W-11.8*C& \left(3c\right)\end{array}$ and Cr, Fe, Al, Si, Ti, Cu, Mo, W and C are the concentrations of the elements in question in mass-%.

19: The method according to claim 1, further optionally containing at most 0.5% vanadium.

20: The method according to claim 1, wherein the impurities are adjusted to contents of max. 0.002% Pb, max. 0.002% Zn, max. 0.002% Sn.

21. (canceled)

Description

EXAMPLES

Manufacture of a Component from Semifinished Product

[0223] FIG. 1a left shows sketches of the semifinished products in the form of sheet (1), strip (2), bar (3), tube (4), wire and weld filler in wire form (5) in top view and cross section. FIG. 1a right shows, for example, how a product or precursor product (7) is formed, in that two components (1a, 1b), for example, are cut out of the semifinished product sheet and chamfered (ii) and then joined (iv) by fusion welding (iii) with a weld filler in the form of wire (5) by means of a V-groove seam (6a). FIG. 1b left shows a further example for the formation of a product or precursor product (7), in which two tubular components (4a, 4b), for example, are cut out of the semifinished product tube and chamfered (ii) and then joined (iv) by fusion welding (iii) with a weld filler in the form of wire (5) by means of a V-groove seam (6b). FIG. 1b middle shows yet another example for the formation of a product or precursor product (7), in which a hole is milled (ii) through a sheet semifinished product or a sheet component (1c), a tube semifinished product or a tube component (3c) is inserted therein and joined (iv) by fusion welding (iii) with a weld filler in the form of wire (5) by means of a fillet weld (6c). FIG. 1b right shows a further example for the formation of a product or of a precursor product (7), in which two strip portions (2a, 2b) are sheared out from the semifinished product strip (ii), machined to make the edges fit and then joined (iv) abuttingly (9) by fusion welding (iii) with one part of the edges as weld filler (6d).

Tests Performed

[0224] The phases occurring in equilibrium were calculated for the various alloy variants with the JMatPro program of Thermotech. The TTNI7 database of Thermotech for nickel alloys was used as the database for the calculations.

[0225] The creep strength is determined in an uninterrupted uniaxial creep test with elongation measurement under tensile load according to DIN EN ISO 204. For this purpose, the specimen is mounted in a creep-testing machine and subjected to a constant test force. In the process, the time to fracture t.sub.u and the creep elongation at break A.sub.u.sup.b are determined. The time to rupture is a measure for the creep strength and the creep elongation at break is a measure for the creep ductility. The tests were performed on round specimens with a diameter of 10 mm in the measurement region and an initial reference length L.sub.r0 of 50 mm. The sampling took place in a manner transverse relative to the direction of forming of the semifinished product.

Description of the Properties

[0226] The nickel-chromium-aluminum alloy named Alloy NiCrAlH used in this invention has, besides an excellent corrosion resistance under highly corrosive conditions, in this case, for example, an excellent metal dusting resistance, a good phase stability and creep strength.

Phase Stability

[0227] In the nickel-chromium-aluminum-iron system having additions of titanium and/or niobium, various embrittling TCP (topologically close packed) phases can be formed, depending on alloying contents, such as, for example, the Laves phase, sigma phase or -phase or even the embrittling -phase or -phase. The calculation of the equilibrium phase proportions as a function of temperature for N06690, Batch 111389, for example (see Table 2 for the compositions used here), theoretically shows the formation of -chromium (BCC phase in FIG. 3) below 720 C. (T.sub.s BCC) in high quantitative proportions. The formation of this phase is generally made difficult, since it is analytically very different from the parent material. However, if the formation temperature T.sub.s BCC of this phase is very high, it can definitely occur, as described, for example, in E. Slevolden, J. Z. Albertsen. U. Fink, Tjeldbergodden Methanol Plant: Metal Dusting Investigations, Corrosion/2011. paper no. 11144 (Houston, TX: NACE 2011), p. 15 for a variant of Alloy 693 (UNS 06693). FIG. 4 and FIG. 5 show the phase diagrams of the Alloy 693 variants (from U.S. Pat. No. 4,882,125 Table 1) for Alloy 3 and Alloy 10 from Table 2. This phase is brittle and leads to an undesired embrittlement of the material. Alloy 3 has a formation temperature T.sub.s BCC of 1079 C. and Alloy 10 of 939 C. In E. Slevolden, J. Z. Albertsen. U. Fink, Tjeldbergodden Methanol Plant: Metal Dusting Investigations, Corrosion/2011. paper no. 11144 (Houston, TX: NACE 2011), p. 15 the exact analysis of the alloy in which -chromium (BCC) forms is not described. It is to be assumed, however, that -chromium (BCC phase) can be formed among the examples cited in Table 2 for Alloy 693 or analyses that theoretically have the highest formation temperatures Is BCC (such as, for example, Alloy 10). In a corrected analysis (with reduced formation temperature T.sub.s BCC), -chromium was subsequently still observed only close to the surface in E. Slevolden, J. Z. Albertsen. U. Fink, Tjeldbergodden Methanol Plant: Metal Dusting Investigations, Corrosion/2011. paper no. 11144 (Houston, TX: NACE 2011), p. 15. In order to avoid the occurrence of such an embrittling phase, the formation temperature T.sub.s BCC in the nickel-chromium-aluminum alloy (Alloy NiCrAlH) used in this invention should be lower than or equal to 939 C.the lowest formation temperature T.sub.s BCC among the examples for Alloy 693 in Table 2 (from U.S. Pat. No. 4,882,125 Table 1).

[0228] This is the case in particular when the following formula is satisfied:

[00010]$\begin{array}{cc}\mathrm{Fp}39.9\mathrm{with}& \left(2a\right)\end{array}$$\begin{array}{cc}\mathrm{Fp}=\mathrm{Cr}+0.272*\mathrm{Fe}+2.36*\mathrm{Al}+2.22*\mathrm{Si}+2.48*\mathrm{Ti}+1.26*\mathrm{Nb}+0.477*\mathrm{Cu}+0.374*\mathrm{Mo}+0.538*W-11.8*C& \left(3d\right)\end{array}$

wherein Cr, Al, Fe, Si, Ti, Nb, Cu, Mo, W and C are the concentrations of the elements in question in mass-%. Table 2 containing the alloys shows that Fp is greater than 39.9 for Alloy 8, Alloy 3 and Alloy 2 and is exactly 39.9 for Alloy 10. For all other alloys, T.sub.s BCC is lower than 939 C. and so Fp is <39.9.
Example for the Manufacture of Components with Welded Seams and their Properties

[0229] Tables 3a and 3b show the analyses of industrially smelted batches of Alloy NiCrAlH alloys, used in this invention, from which sheets and welding rods (weld filler in the form of wire) were made. Formula (2a) Al+Cr28 is satisfied for these batches, as is therefore the requirement that was imposed on the metal dusting resistance. For the compositions in Tables 3a and 3b, the value for Fp was also calculated according to formula (3a). As required, Fp is lower than 39.9.

[0230] These batches were openly smelted in quantities of 16 metric tons, followed by a treatment in a VOD unit. Then electrodes were cast and subjected to ESR. Thereafter the alloy was annealed at temperatures between 900 C. and 1270 C. for 0.1 to 70 hours and hot-formed with intermediate annealings between 900 C. and 1270 C. for 0.1 hours to 70 hours. Hot-rolled sheets with the thickness of 25 and 16 mm were manufactured from Batch 319144. Hot-rolled wire was made from Batch 318385. After the hot rolling, the sheets were solution-annealed at 1100 C. for 40 minutes and then cooled in air. Then they were abrasive blasted, pickled and ground for removal of the oxide layer. The 25-mm-thick sheets had a grain size of approximately 89 m, the 16-mm-thick sheets a grain size of approximately 82 m. The 25-mm sheet and the 16-mm sheet therefore have a comparable grain size. The rolled wire was likewise abrasive blasted, pickled and ground and then cold-drawn to final thickness with intermediate annealings between 80 and 1250 C. for 0.05 minutes to 70 hours. Then the wire is solution-annealed under hydrogen in the temperature range of 800 to 1250 C. for 0.05 minutes to 70 hours and processed to welding rods of 2.0 and 2.4 mm diameter.

[0231] Sheet portions were cut out of the 25-mm-thick solution-annealed semifinished product sheet and annealed at 980 C. for 3 hours with subsequent air cooling. For creep tests, specimens transverse to the rolling direction were made from the sheets or sheet portions that were only solution-annealed and from those that were additionally annealed. Table 5 shows the results from the creep tests according to DIN EN ISO 204.

[0232] Sheet portions measuring 150500 mm were cut out of the 16-mm-thick semifinished product sheet. Respectively 2 pieces were welded with a 70 V-groove seam by means of manual TIG welding under pure argon using the 2.0-mm-thick and 2.4-mm-thick welding rods from Batch 318385 as weld filler together with the welding parameters indicated in Table 4. The welded seam and the heat-affected zone were brushed immediately after the welding. Some partial pieces of the welded sheets manufactured in this way were annealed at 980 C. for 3 hours with subsequent air cooling, others at 1100 C. for 40 minutes with subsequent air cooling and yet others at 1100 C. for 3 hours with subsequent air cooling. Some were not subjected to any annealing. Seams or seam portions were also manufactured that were ground or not subsequently treated at all. For creep tests, specimens transverse to the welded seam were made from the welded sheets and from the welded and annealed sheets or sheet portions. Table 5 shows the results from the creep tests according to DIN EN ISO 204.

[0233] From Table 5, it is evident that, in creep tests at 600 C., an additional annealing of a solution-annealed sheet at 980 C. for 3 hours followed by an air cooling (specimen 19 45B and 19 46B) resulted in a shortening of the time to fracture t.sub.u in comparison with the sheet that was only solution-annealed (specimen 19 23B and 19 7B). A creep test (specimen 302W and 303W) transverse to the welded seam without a further annealing (prior art T) likewise exhibits a significantly shortened time to fracture t.sub.u in comparison with a creep test on the sheet that was only solution-annealed. An additional annealing, at 980 C. for 3 hours followed by an air cooling (specimen 247W and 249W), of the welded portion from which the creep specimens are made, has no marked impact on the time to fracture t.sub.u, but it markedly reduces the creep elongation at break A.sub.u.sup.b. In contrast, an additional inventive annealing, at 1100 C. for 40 minutes followed by an air cooling (specimen 250W and 503W), of the welded portion from which the creep specimens are made, respectively causes a significant prolongation of the time to fracture t.sub.u by a factor of approximately 3 as well as an increase of the creep elongation at break A.sub.u.sup.b, in some cases above the value of the specimens that were not annealed after the welding (specimen 302W and 303W). In particular, an additional inventive annealing, at 1100 C. for 3 hours followed by an air cooling (specimen 511W and 506W), of the welded portion from which the creep specimens are made, respectively causes a further prolongation of the time to fracture t.sub.u as well as an increase of the creep elongation at break A.sub.u.sup.b, almost up to the value of the creep test on the sheet that was only solution-annealed but not welded (specimen 19 23B and 19 7B).

[0234] At 700 C., a creep test (specimen 306W) transverse to the welded seam without a further annealing (prior art T) likewise exhibits, just as at 600 C., a shortened time to fracture t.sub.u in comparison with the creep test on the sheet that was only solution-annealed (specimen 30 34B). An additional annealing, at 980 C. for 3 hours followed by an air cooling (specimen 248W), of the welded portion from which the creep specimens are made, again causes a slight shortening of the time to fracture t.sub.u in comparison with the specimen 306W that was not annealed after the welding. In contrast, an additional inventive annealing, at 1100 C. for 40 minutes followed by an air cooling (specimen 251W), of the welded portion from which the creep specimens are made, causes a prolongation of the time to fracture t.sub.u by the factor of approximately 2 as well as an increase of the creep elongation at break A.sub.u.sup.b, significantly above the value of the specimen (306W) that was were not annealed after the welding. In particular, an additional inventive annealing, at 1100 C. for 3 hours followed by an air cooling (specimen 253W), of the welded portion from which the creep specimens are made, causes a further prolongation of the time to fracture t.sub.u beyond the time to fracture of the creep test on the sheet that was only solution-annealed.

[0235] At 800 C., an additional annealing, at 980 C. for 3 hours followed by an air cooling (specimen 19 49B), of a solution-annealed sheet resulted in a similar time to fracture t.sub.u in comparison with the sheet that was only solution-annealed (specimen 19 22B). A creep test (specimen 309W) transverse to the welded seam without a further annealing (prior art T) also has a similar time to fracture t.sub.u in comparison with the creep test on the sheet that was only solution-annealed (specimen 19 22B). An additional annealing, at 1100 C. for 40 minutes followed by an air cooling (specimen 519W), of the welded portion from which the creep specimens are made, likewise causes a similar time to fracture t.sub.u in comparison with the creep test on the sheet (specimen 19 22B) that was only solution-annealed.

[0236] This means that a heat treatment of the welded seam after the welding at 1100 C. for at least 40 minutes significantly improves, according to the invention, the time to fracture and the creep elongation at break of a creep specimen transverse to the welded seam at temperatures of 600 and 700 C. and thus in the range of -formation. At higher temperatures above the -solvus temperature, the non-annealed welded seam has a similar to better time to fracture t.sub.u in comparison with the sheet that was only solution-annealed. An annealing of the welded seam influences only negligibly the time to fracture at service temperatures above the -solvus temperature.

[0237] The claimed limits for the method and the nickel-chromium-aluminum alloy named Alloy NiCrAlH used in the invention can therefore be justified in detail as follows:

[0238] Too low chromium contents mean that the chromium concentration underneath the oxide layer during use of the alloy in a corrosive atmosphere decreases very rapidly below the critical limit, so that a closed chromium oxide layer can no longer be formed. Therefore 18% chromium is the lower limit for chromium.

[0239] Too high chromium contents worsen the phase stability of the alloy, especially at the high aluminum contents of 1.8%. Therefore 33% chromium is to be regarded as the upper limit.

[0240] The formation of an aluminum oxide layer underneath the chromium oxide layer reduces the oxidation rate. Below 1.8% aluminum, the aluminum oxide layer is too incomplete to develop its effect fully. Too high aluminum contents impair the processability of the alloy. Therefore an aluminum content of 4.0% forms the upper limit.

[0241] The costs for the alloy increase with the reduction of the iron content. Below 0.01%, the costs rise disproportionally, since special precursor material must be used. For cost reasons, therefore, 0.01% iron is to be regarded as the lower limit. With increase of the iron content, the phase stability is reduced (formation of embrittling phases), especially at high chromium and aluminum contents. Therefore 7% Fe is a practical upper limit in order to ensure the phase stability of the alloy according to the invention.

[0242] Silicon is needed for the manufacture of the alloy. A minimum content of 0.001% is therefore necessary. Too high contents in turn impair the processability and the phase stability, especially at high aluminum and chromium contents. The silicon content is therefore limited to 0.50%.

[0243] A minimum content of 0.001% manganese is necessary for improvement of the processability. Manganese is limited to 2.0%, since this element reduces the oxidation resistance.

[0244] Titanium increases the high-temperature strength. Above 0.60%, the oxidation behavior may be impaired, which is why 0.60% is the maximum value.

[0245] Even very low magnesium contents and/or calcium contents improve the processing by the binding of sulfur, whereby the occurrence of low-melting NiS eutectics is avoided. At too high contents, intermetallic NiMg phases or NiCa phases may occur, which again greatly worsen the processability. The magnesium content and/or calcium content is therefore limited to at most 0.05%.

[0246] A minimum content of 0.005% carbon is necessary for a good creep resistance. Carbon is limited to at most 0.12%, since above such a content this element reduces the processability by the excessive formation of primary carbides.

[0247] A minimum content of 0.0005% nitrogen is necessary, whereby the processability of the material is improved. Nitrogen is limited to at most 0.05%, since this element reduces the processability due to the formation of coarse carbonitrides.

[0248] The oxygen content must be 0.020%, in order to ensure the manufacturability of the alloy. A too low oxygen content increases the costs. The oxygen content is therefore 0.0001%.

[0249] The content of phosphorus should be smaller than or equal to 0.030%, since this surface-active element impairs the oxidation resistance. A too low phosphorus content increases the costs. The phosphorus content is therefore 0.001%.

[0250] The contents of sulfur should be adjusted as low as possible, since this surface-active element impairs the oxidation resistance. Therefore at most 0.010% sulfur is specified.

[0251] Molybdenum is limited to at most 2.0%, since this element reduces the oxidation resistance.

[0252] Tungsten is limited to at most 2.0%, since this element likewise reduces the oxidation resistance.

[0253] Nickel is the element comprising the rest. A too low nickel content reduces the phase stability, especially at high chromium contents. Nickel must therefore be higher than or equal to 50%.

[0254] For highly corrosive conditions, but especially for a good metal dusting resistance, it is advantageous when the following relationship is satisfied between Cr and Al:

[00011]$\begin{array}{cc}\mathrm{Cr}+\mathrm{Al}28& \left(1a\right)\end{array}$

wherein Cr and Al are the concentrations of the elements in question in mass-%. Only then is the content of oxide-forming elements high enough to ensure a sufficient metal dusting resistance.

[0255] Beyond this, the following relationship must be satisfied in order that an adequate phase stability is ensured:

[00012]$\begin{array}{cc}\mathrm{Fp}39.9\mathrm{with}& \left(2a\right)\end{array}$$\begin{array}{cc}\mathrm{Fp}=\mathrm{Cr}+0.272*\mathrm{Fe}+2.36*\mathrm{Al}+2.22*\mathrm{Si}+2.48*\mathrm{Ti}+0.374*\mathrm{Mo}+0.538*W-11.8*C& \left(3a\right)\end{array}$

wherein Cr, Fe, Al, Si, Ti, Mo, W and C are the concentrations of the elements in question in mass-%. The limits for Fp and the possible incorporation of further elements have been justified in detail in the foregoing text.

[0256] If necessary, the oxidation resistance may be further improved with additions of oxygen-affine elements, such as, for example, yttrium, lanthanum, cerium, cerium mixed metal, zirconium, hafnium. They do this by being incorporated in the oxide layer, where they block the paths of diffusion of the oxygen to the grain boundaries.

[0257] Yttrium increases the oxidation resistance. For cost reasons, the upper limit is set to 0.20%.

[0258] Lanthanum increases the oxidation resistance. For cost reasons, the upper limit is set to 0.20%.

[0259] Cerium increases the oxidation resistance. For cost reasons, the upper limit is set to 0.20%.

[0260] Cerium mixed metal increases the oxidation resistance. For cost reasons, the upper limit is set to 0.20%.

[0261] If necessary, niobium may be added, since niobium also increases the high-temperature strength. Higher contents very greatly increase the costs. The upper limit is therefore set at 1.10%.

[0262] If necessary, the alloy may also contain tantalum, since tantalum also increases the high-temperature strength and the oxidation resistance. Higher contents very greatly increase the costs. The upper limit is therefore set at 0.60%. A minimum content of 0.001% is necessary in order to achieve an effect.

[0263] If necessary, the alloy may also contain zirconium. Zirconium increases the high-temperature strength and the oxidation resistance. For cost reasons, the upper limit is set to 0.20% zirconium.

[0264] If necessary, the alloy may also contain hafnium. Hafnium increases the high-temperature strength and the oxidation resistance. For cost reasons, the upper limit is set to 0.20% hafnium.

[0265] If necessary, boron may be added to the alloy, since boron improves the creep resistance. Therefore a content of at least 0.0001% should be present. At the same time, this surface-active element worsens the oxidation resistance. Therefore at most 0.008% boron is specified.

[0266] Cobalt up to 5.0% may be contained in this alloy. Higher contents markedly reduce the oxidation resistance.

[0267] Copper is limited to at most 0.5%, since this element reduces the oxidation resistance.

[0268] Vanadium is limited to at most 0.5%, since this element reduces the oxidation resistance.

[0269] Lead is limited to at most 0.002%, since this element reduces the oxidation resistance. The same is true for zinc and tin.

[0270] Too small grain sizes of smaller than 30 m lead to a poor creep strength at higher temperatures. Too large grain sizes of larger than 600 m lead to a very low creep ductility at temperatures in the range of -formation.

[0271] A homogenization of the welded seams and/or for reduction of stresses by an annealing between higher than 980 and 1250 C. for times from 0.05 minutes to 24 hours, followed by a cooling in stationary shield gas or air, moving (blown) shield gas or air, improves the creep strength and the creep ductility of the welded seams in the range of -formation. At a too low temperature lower than or equal to 980 C., the temperature is too low, so that homogenization cannot be carried out economically because of the long times required. At a temperature well above the solution annealing, marked grain growth takes place, which reduces the hot strength at low temperatures as well as the elongation in the creep test in the range of -formation. Short times of less than 0.05 minutes are not adequate even at very high temperatures. Times longer than 24 hours are uneconomical, especially for larger components.

[0272] An annealing under shield gas reduces the oxidation of the material during the annealing and thus a material loss.

DESCRIPTION OF THE FIGURES

[0273] FIG. 1a: Left: Sketches in top view and cross section of the semifinished product forms sheet (, strip 2, bar 3 tube 4 and wire or weld filler in wire form 5. Right: Examples of manufacture of a component 7 by chamfering of two sheets 1a, 1b and joining by fusion welding with a weld filler in wire form 5 by means of a V-groove seam 6a.

[0274] FIG. 1b: Examples of manufacture of components 7. Left: Chamfering of two tubes at one end 4a, 4b and joining by fusion welding with a weld filler in wire form 5 by means of a V-groove seam 6b. Middle: Milling of a hole in a sheet 1c and inserting and joining of a tube by fusion welding with a weld filler in wire form 5 by means of a fillet weld 6c. Right: Joining of two strip portions 2a, 2b machined to make the edges fit abuttingly by fusion welding 6d with one part of the edges as weld filler 9.

[0275] FIG. 2: Loss of metal due to metal dusting as a function of the aluminum and chromium content in a strongly carburizing gas containing 37% CO, 9% H.sub.2O, 7% CO.sub.2, 46% H.sub.2 as well as the following activities a.sub.c=163 and p(O.sub.2)=2.510.sup.27. (from (Hermse, C. G. M. and van Wortel, J. C.: Metal dusting: relationship between alloy composition and degradation rate. Corrosion Engineering, Science and Technology 44 (2009), pp. 182-185).

[0276] FIG. 3: Quantitative proportions of the phases in thermodynamic equilibrium as a function of the temperature of Alloy 690 (N06690) for the example of Batch 111389.

[0277] FIG. 4: Quantitative proportions of the phases in thermodynamic equilibrium as a function of the temperature of Alloy 693 (N06693) for the example of Alloy 3 from Table 2.

[0278] FIG. 5: Quantitative proportions of the phases in thermodynamic equilibrium as a function of the temperature of Alloy 693 (N06693) for the example of Alloy 10 from Table 2.

TABLE-US-00006 TABLE 1 Some alloys according to ASTM B 168-11. All values in mass-%. Alloy Ni Cr Co Mo Nb Fe Mn Al C Cu Si S Ti P Zr Y B N Ce Alloy 72.0 14.0- 6.0- 1.0 0.15 0.5 0.5 0.015 600 - min 17.0 10.0 max. max. max. max. max. N06600 Alloy 58.0- 21.0- Rest 1.0 1.0- 0.10 0.5 0.5 0.015 601 - 63.0 25.0 max. 1.7 max. max. max. max. N06601 Alloy 44.5 20.0- 10.0- 8.0- 3.0 1.0 0.8- 0.05- 1.0 0.5 0.015 0.6 0.006 617 - min 24.0 15.0 10.0 max. max. 1.5 0.15 max. max. max. max. max. N06617 Alloy 58.0 27.0- 7.0- 0.5 0.05 0.5 1.0 0.015 690 - min 31.0 11.0 max. max. max. max. max. N06690 Alloy Rest 27.0- 0.5- 2.5- 1.0 2.5- 0.15 0.5 0.5 0.01 1.0 693 - 31.0 2.5 6.0 max. 4.0 max. max. max. max. max. N06693 Alloy Rest 24.0- 8.0- 0.15 1.8- 0.15- 0.1 0.5 0.010 0.1- 0.020 0.01- 0.05- 602CA - 26.0 11.0 max. 2.4 0.25 max. max. max. 0.2 max. 0.10 0.12 N06025 Alloy 45 26.0- 21.0- 1.0 0.05- 0.3 2.5- 0.010 0.020 0.03- 45 - min 29.0 25.0 max. 0.12 max. 3.0 max. max. 0.09 N06045 Alloy Rest 24.0- 8.0- 0.15 2.4- 0.20- 0.50 0.5 0.010 0.01- 0.020 0.01- 0.01- 603 - 26.0 11.0 max. 3.0 0.40 max. max. max. 0.25 max. 0.10 0.15 N06603 Alloy Rest 28.0- 1.0- 2.0- 1.0 0.15 1.5- 1.0- 0.010 1.0 696 - 32.0 3.0 6.0 max. max. 3.0 2.5 max. max. N06696

TABLE-US-00007 TABLE 2 Compositions of some alloys according to ASTM B 168-11. All values in mass-%. Alloy Batch C S Cr Ni Mn Si Mo Ti Nb Alloy 164310 0.07 0.002 15.75 73.77 0.28 0.32 0.2 600 N06600 Alloy 156656 0.053 0.0016 22.95 59.58 0.72 0.24 0.47 601 N06601 Alloy 111389 0.022 0.002 28.45 61.95 0.12 0.32 0.29 690 N06690 Alloy Alloy 10 *) 0.015 0.01 29.42 60.55 0.014 0.075 0.02 1.04 693 N06693 Alloy Alloy 8 *) 0.007 0.01 30.00 60.34 0.11 0.38 0.23 1.13 693 N06693 Alloy Alloy 3 *) 0.009 0.01 30.02 57.79 0.01 0.14 0.02 2.04 693 N06693 Alloy Alloy 2 *) 0.006 0.01 30.01 60.01 0.12 0.14 0.01 0.54 693 N06693 Alloy 163968 0.170 0.01 25.39 62.12 0.07 0.07 0.13 602 N06025 Alloy 52475 0.225 0.002 25.20 61.6 0.09 0.03 0.16 0.01 603 N06603 Alloy UNS Mitte 0.080 0.01 30.00 61.20 0.1 1.5 2 0.1 696 N06696 T.sub.s BCC Cr + Alloy Cu Fe P Al Zr Y B in C. Al Fp Alloy 0.01 9.42 0.009 0.16 0.001 15.9 19.1 600 N06600 Alloy 0.04 14.4 0.008 1.34 0.015 0 0.001 669 24.3 31.2 601 N06601 Alloy 0.01 8.45 0.005 0.31 0 0 720 28.8 32.7 690 N06690 Alloy 0.03 5.57 3.2 0.002 939 32.6 39.9 693 N06693 Alloy 0.03 4.63 3.08 0.002 979 33.1 41.3 693 N06693 Alloy 0.03 5.57 4.3 0.002 1079 34.3 44.5 693 N06693 Alloy 0.03 5.80 3.27 0.002 948 33.3 40.3 693 N06693 Alloy 0.01 9.47 0.008 2.25 0.08 0.08 0.005 690 27.6 31.8 602 N06025 Alloy 0.01 9.6 0.007 2.78 0.07 0.08 0.003 707 28.0 32.2 603 N06603 Alloy 2 3 792 30.0 35.1 696 N06696 *) Alloy composition from U.S. Pat. No. 4.88.125 Table 1.

TABLE-US-00008 TABLE 3a Composition of the industrially smelted batches (G) of the nickel-chromium-aluminum alloy named Alloy NiCrAlH used in this invention, Part 1. All values in mass-%. T.sub.s BCC Cr + Name Batch C N Cr Ni Mn Si Mo Ti Nb Cu Fe Al W in C. Al Fp H G 25 mm 319144 0.020 0.017 29.47 67.96 <0.01 0.05 <0.01 0.01 0.14 <0.01 0.11 2.11 <0.01 778 31.58 34.5 sheet H G 16 mm 319144 0.020 0.018 29.47 67.89 0.01 0.05 <0.01 0.01 0.14 <0.01 0.12 2.16 <0.01 783 31.63 34.7 sheet H G Welding 318385 0.022 0.026 28.92 68.20 0.01 0.06 <0.01 0.02 0.14 0.01 0.47 2.03 <0.01 757 30.92 33.9 rods (H: Examples of the nickel-chromium-aluminum alloy named Alloy NiCrAlH used in this invention, G: industrially smelted).

TABLE-US-00009 TABLE 3b Composition of the industrially smelted batches (G) of the nickel-chromium-aluminum alloy named Alloy NiCrAlH used in this invention, Part 2. All values in mass-%. Name Batch S P Mg Ca V Zr Co Y La B Hf Ta Ce O H G 25 mm 319144 <0.002 0.002 0.005 <0.001 <0.01 0.03 <0.01 <0.01 <0.01 0.002 <0.01 <0.01 <0.01 0.001 sheet H G 16 mm 319144 <0.002 0.002 0.006 <0.001 <0.01 0.03 <0.01 <0.01 <0.01 0.002 <0.01 <0.01 <0.01 0.001 sheet H G Welding 318385 <0.002 0.002 0.009 <0.001 <0.01 0.02 <0.01 <0.01 <0.01 0.003 <0.01 <0.01 <0.01 0.001 rods (Pb: max. 0.002%, Zn: max. 0.002%, Sn: max. 0.002% are applicable for all alloys; meaning of H, G: see Table 3a).

TABLE-US-00010 TABLE 4 Welding parameters for the welding of the 16-mm-thick sheets (Batch 319144) with welding rods from Batch 318385, of the nickel-chromium-aluminum alloy named Alloy NiCrAlH used in this invention. Energy per Weld filler Welding speed unit length diam. in mm Filler and in cm/min in kJ/cm Thickness Welding Filler and Root pass top passes Filler and Filler and Shield in mm technique Root top passes I in A U in V I in A U in V Root top passes top passes gas 16 m-TIG 2.0 2.0-2.4 120 15 190 17 5 8 20-25 Ar 4.6/ pure Ar

TABLE-US-00011 TABLE 5 Results of the creep tests according to DIN EN ISO 204 on i) 25-mm-thick solution-annealed sheets (1100 C./40 min/air cooling, grain size 89 m) of Batch 319144 (BM) and on ii) 16-mm-thick solution-annealed sheets (1100 C./40 min/air cooling, grain size 82 m) of Batch 319144 (S), welded with welding rods from Batch 318385. Annealing after Temperature Initial stress Time to fracture Creep elongation Specimen Material welding T in C. .sub.0 in MPa t.sub.u in h at break A.sub.u.sup.b in % 19 23B BM none 600 315 4832 7.3 19 45B BM 980 C./3 h/AC 600 315 1175 8.1 T 302W S none 600 305 626 3.5 247W S 980 C./3 h/AC 600 305 593 1.5 E 250W S 1100 C./40 min/AC 600 305 1927 3.4 E 511W S 1100 C./3 h/AC 600 305 2013 7.4 19 7B BM none 600 270 15744 4.5 19 46B BM 980 C./3 h/AC 600 270 4978 5.1 T 303W S none 600 270 1222 2.7 249W S 980 C./3 h/AC 600 270 1224 1.6 E 503W S 1100 C./40 min/AC 600 270 3816 4.1 E 506W S 1100 C./3 h/AC 600 270 5256**) **) 30 34B BM none 700 90 3665 3.8 T 306W S none 700 90 2450 0.7 248W S 980 C./3 h/AC 700 90 1727 0.8 E 251W S 1100 C./40 min/AC 700 90 3812 2.6 E 253W S 1100 C./3 h/AC 700 90 4987 *) 19 22B BM none 800 39 1878 16.9 19 49B BM 980 C./3 h/AC 800 39 1656 16.6 T 309W S none 800 39 2936 4.8 519W S 1100 C./40 min/AC 800 39 1896**) **) AC = air cooling *) Fracture at the shoulder, not measurable, **)Specimen still in testing. E = According to the invention; T = Prior art