EP2240619B1 - Creep resistant steel - Google Patents

Description

Technical area

The invention relates to steels based on 9-12% chromium, which are used for rotors in the power plant sector. It concerns the choice and the proportionate tuning of special alloying elements which allow the setting of an exceptionally good creep resistance at temperatures of 550 ° C and above in this material. The steel according to the invention should also have a high toughness after long-term aging, so that it can be used both in gas and steam turbines.

State of the art

Maragingitic-hardening steels based on 9-12% chromium are widely used materials in power plant technology. They were developed for use in steam power plants at operating temperatures above 600 ° C and steam pressures above 250 bar to increase the efficiency of power plants. Under these operating conditions, the creep resistance and the oxidation resistance of the material play a special role.

It is known that the addition of chromium in the abovementioned range not only provides good resistance to atmospheric corrosion but also complete through-hardenability of thick-walled forgings, for example as monobloc rotors or as rotor disks in gas and steam turbines. Proven alloys of this type usually contain about 0.08 to 0.2% carbon, which in solution allows the setting of a hard martensitic structure. A good combination of heat resistance and ductility of martensitic steels is made possible by a tempering treatment in which the precipitation of carbon in the form of carbides with simultaneous recovery of the dislocation substructure forms a particle-stabilized subgrain structure. The tempering behavior and the resulting properties can be effectively influenced by the choice and proportionate tuning of specific carbide formers such as Mo, W, V, Nb and Ta.

A typical representative who has found widespread use in steam power plants, in particular as rotor steel, is the German steel X20CrMoV12.1 known under DIN.

It is also known that the ductility at a strength level of 850 MPa can be significantly improved by the alloying of nickel. Such alloys are therefore widely used where significantly higher demands are placed on both strength and ductility, typically as disk materials for gas turbine rotors. A typical representative of such alloys, which has found wide use in gas turbine technology, in particular as a material for rotor disks, is the German steel X12CrNiMo12 known under DIN. However, nickel tends to reduce hot strength at high temperatures. This is related to reduced carbide stability in nickel-containing steels.

In the past, various efforts have been made to improve specific properties of the known 9-12% Cr steels. For example, in the publication by Kern et al .: High Temperature Forged Components for Advanced Steam Power Plants, Materials for Advanced Power Engineering 1998, Proceedings of the 6th Liege Conference, ed. by J. Lecomte-Becker et. al., the development of novel rotor steels for steam turbine application en described.

In such alloys, the contents of Cr, Mo, W were optimized taking into account N, Nb and / or B to improve creep and creep rupture strengths for 600 ° C applications. By adding boron, the carbides, such as M ₂₃ C ₆ , are stabilized. Because of the harmful effect of nickel on the long-term properties, the Ni contents were limited to values of less than 0.25% in these steels. In the case of these alloys, the fracture toughness values are disadvantageous, which does not play a major role in steam turbine applications and can therefore be neglected, but must be avoided in gas turbine applications.

In later publications ( F. Kauffmann et. al.: "Microstructural Investigation of Boron containing TAF Steel and the Correlation to the Creep Strength", 31st MPA Seminar in conjunction with the symposium "Material and Component Behavior in Energy and Plant Engineering", 13./14.10.2005, Stuttgart ), the Ni contents were even limited to values <0.002% for a B addition of 0.03% to a 10.5% Cr steel.

Specifically for gas turbine applications, efforts have been made to either improve the creep ruptures in the range of 450 to 500 ° C to high ductility level in 9-12% Cr steels or to reduce the embrittlement tendency at temperatures between 425 and 500 ° C. This is how the European patent application describes EP 0 931 845 A1 a nickel-containing 12% chromium steel similar in structure to the German steel X12CrNiMo12, in which reduces the element molybdenum compared to the known steel X12 CrNiMo12, but an increased content of tungsten was alloyed. In DE 198 32 430 A1 is a further optimization of the X12CrNiMo12 similar steel with the name M152 disclosed in which by the addition of rare earth elements, the embrittlement tendency in the temperature range between 425 and 500 ° C is limited.

The disadvantage is that in none of the abovementioned developments could the strength, in particular the heat resistance at temperatures between 300 and 600 ° C., be improved on a level of ductility comparable to the steel X12CrNiMo12.

One possible approach to improving the high temperature strength and high ductility has been proposed with the development of steels with increased nitrogen contents. In EP 0 866 145 A2 describes a new class of martensitic chromium steels with nitrogen contents in the range of 0.12 to 0.25% and in EP 1 158 067 A1 with nitrogen contents of 0.12 to 0.18%, wherein the weight ratio V / N is in the range between 3.5 and 4.2. In these steels, the entire structure of the structure is controlled by the formation of special nitrides, in particular vanadium nitrides, which can be distributed in a variety of ways by forging, austenitizing, controlled cooling or annealing. While strength is achieved through the nitriding's curing effect, the aim is to achieve high ductility through the distribution and morphology of the nitrides, but above all by limiting grain coarsening during forging and during solution heat treatment.

US2002020473 discloses a high-chromium, heat-resistant, ferritic steel having the following composition (in% by weight): 0.02 to 0.15% C, 0.05 to 1.5% Mn, ≤0.03% P, ≤0.015% S, 8 to 13% Cr, 1.5 to 4% W, 2 to 6% Co, 0.1 to 0.5% V, 0.01 to 0.15% Ta, 0.01 to 0.15% Nb, 0.001 to 0.2% Nd, ≤0.02% N, 0.0005 to 0.02% B, 0.001 to 0.05 % Al, at least one of Ca (≤0.02%), Y (≤0.2%), La (≤0.2%) and Hf (≤0.2%), remainder Fe.

Out EP 0867 522 A2 For example, a heat resistant steel with good toughness properties is known for use as a turbine rotor, which has the following chemical composition (% by weight): 0.05-0.30 C, 0.20 or less Si, 0-1.0 Mn, 8-14 Cr, 0.5-3.0 Mo, 0.10-0.50 V, 1.5-5.0 Ni, 0.01-0.5 Nb, 0.01-0.08 N, 0.001-0.020 B, balance iron and unavoidable impurities. Boron microalloying results in precipitates at the grain boundaries and increases the time stability of the carbonitrides at high temperatures, but higher levels of B reduce the toughness of the steel. Disadvantages of this proposed composition are also the relatively high permitted Si values of 0.2%. Although Si serves advantageously as a deoxidizer at the time of melting, parts of it remain as oxides in the steel, which is disadvantageous in a reduced toughness.

Finally, in US 5820817 proposed stainless steels with 8-13 wt .-% Cr, which include boron and / or rare earth metals in their composition, in order to increase the resistance to embrittlement on long-term aging. According to this document, the maximum content of rare earth should be 0.5% by weight, an optimum proportion being 0.1% by weight.

Presentation of the invention

The invention has for its object to provide a 9-12% Cr steel, which is characterized over the prior art by increased creep strength at temperatures of 550 ° C and above and which improved resistance to embrittlement in long-term aging and a has relatively high toughness, so that it can be used especially in gas turbine, but also in steam turbine power plants. It should preferably find application for rotors of turbomachinery, so that the efficiency and the output can be increased over the known prior art.

The core of the invention is a steel having the following chemical composition (in% by weight): 0.08 to 0.16 C, 9.0 to 12.0 Cr, 0.1 to 0.5 Mn, 2.3 to 3 Ni, 1.5 to 2.0 Mo, 0.1 to 0.4 V, 0.01 to 0.06 Nb, 0.02 to 0.08 N, 0.001 to 2 Ta, 0.001 to 0.5 La, 0.0001 to 1 Pd, 0.004 to 0.012 B, maximum 0.005 P, maximum 0.005 S, maximum 0.05 Si, maximum 0.005 Sn, balance iron and unavoidable impurities ,

Preferred ranges for the individual alloying elements of the composition according to the invention are contained in the subclaims, wherein the steel according to the invention particularly preferably has the following chemical composition (in% by weight): 0.12 C, 11.5 Cr, 0.2 Mn, 2.5 Ni, 1.7 Mo, 0.25 V, 0.03 Nb, 0.04 N, 0.01 Ta, 0.05 La, 0.001 Pd, 0.007 B, 0.005 P, 0.005 S, 0.05 Si, 0.005 Sn, balance iron and unavoidable impurities.

The advantage of the invention is that the inventive alloy compared to the known from the prior art alloys of similar composition, but without B addition or without La and Pd addition, with the same heat treatment improved creep properties at temperatures of 550 ° C and above, while also good toughness properties (elongation, impact work) and improved resistance to embrittlement during long-term aging can be achieved.

It is set a starting structure, which is characterized by a tough matrix and the presence of heat-resistant nitrides, borides and carbides. The toughness of the base matrix is adjusted by the presence of substitution elements, preferably nickel. The contents of these substitution elements are determined to provide optimal unfolding of both martensite hardening and particle hardening by precipitation of special nitrides, e.g. As vanadium nitrides or niobium nitrides, to set the highest heat resistance possible.

Basically, both hardening mechanisms lower the ductility. Characteristically, a minimum ductility is observed in the area of secondary hardening. This minimum ductility need not be caused exclusively by the actual precipitation hardening mechanism. A certain embrittlement contribution may also be provided by segregation of impurities to the grain boundaries or possibly also by near-order adjustments of dissolved alloy atoms.

Increasing the tempering temperature over the secondary hardening area results in complete precipitation with significant growth of carbides. As a result, the strength decreases and the ductility increases. Importantly, the simultaneous recovery of dislocation substructure and particle coarsening increases ductility so as to improve the overall combination of strength and ductility. This improvement is attributable to the formation of a particle-stabilized subgrain structure. It can be assumed that both the ductility and the strength of particle-stabilized subgrain structures are reduced by unevenness in the topology of the particle sub-grain structure. Precipitose precipitates are subject to accelerated coarsening and tend to coagulate with adjacent precipitates. Coarse and coagulated phases generate fracture stress peaks, which reduce ductility. Above all, however, the non-uniform distribution of the precipitates also severely limits the hardening mechanism which is most effective at high temperatures, namely particle hardening.

One measure for increasing the ductility in conventional martensitic-hardening steels is the alloying of nickel. However, the reasons for this are not known in all respects and should depend heavily on the nickel content. Thus, small amounts of nickel can be very Ductititätsfördernd, if it can be completely suppressed, for example, the formation of delta ferrite.

By contrast, nickel contents above 2% by weight are expected to lower the Ac1 temperature (which is the temperature at which ferrite begins to convert to austenite during heating) to temperatures below 700 ° C. So if the strength is increased by lowering the tempering temperature below 700 ° C, then in the presence of increased nickel contents during tempering with a partial conversion of ferrite into austenite can be expected. This is associated with a certain ductility-promoting grain regeneration. However, it should be noted that the carbide precipitation above the Ac1 temperature is only incomplete, since the solubility of the austenite-stabilizing element carbon in austenite is greater than in the ferrite. Furthermore, the austenite which forms is not sufficiently stabilized, so that a larger volume fraction of the reformed austenite undergoes further martensitic transformation in the post-anneal cooling. In addition to the two above-mentioned contributions of nickel to the ductility increase, a certain ductility contribution of nickel can come into solid solution as a substitution element. This can be explained electron-theoretically in such a way that the element nickel feeds additional, free electrons into the iron grid and thereby makes the iron alloys even more "metallic".

Basically, conventional martensitic-hardenable steels which are alloyed with nickel have no particular advantages in heat resistance compared to nickel-poor alloys. This is true at least for test temperatures above 500 ° C and could be related to the above-mentioned reactivation on tempering at elevated nickel levels. It is also known that the alloying of nickel into such steels significantly aggravates the structural instability under long-term aging conditions at elevated temperatures. This long-term structural instability is related to an accelerated coarsening of the carbides.

Manganese is on the left side next to the element iron in the periodic system of elements. It is an electron-poorer element, so its action in solid solution should be distinctly different from that of nickel. Nonetheless, it is an austenite stabilizing element which greatly lowers the Ac1 temperature, but leaves no particularly positive but rather unfavorable effect on ductility. On the carbonaceous 12% chromium steel side, manganese is understood to be an impurity element which promotes temper embrittlement substantially. Therefore, the content of manganese is usually limited to very small amounts.

The preferred amounts in weight percents for each element and the reasons for the selected alloying ranges according to the invention are shown below in connection with the heat treatment possibilities resulting therefrom.

Chrome:

A weight proportion of 9-12% chromium allows good through-hardenability of thick-walled components and ensures sufficient oxidation resistance up to a temperature of 550 ° C. A proportion by weight of less than 9% impairs the through-hardenability. Contents above 12% lead to the accelerated formation of hexagonal chromium nitrides during the tempering process, which, in addition to nitrogen, also cures vanadium, thus reducing the effectiveness of vanadium nitride curing. The optimum chromium content is 10.5 to 11.5%.

Manganese and silicon:

These elements promote temper embrittlement and must therefore be limited to the smallest amounts. The range to be specified should take into account the metallurgical possibilities in the range between 0.1 and 0.5% by weight, preferably between 0.1 and 0.25%, in particular at 0.2% for manganese and at max. 0.05% by weight for silicon.

Nickel:

Nickel is used as an austenite stabilizing element to suppress delta ferrite. In addition, it is said to improve ductility as a dissolved element in the ferritic matrix. Nickel contents of 2.3 to about 3% by weight make sense. Nickel contents above 4% by weight increase the austenite stability such that after solution heat treatment and tempering an increased proportion of retained austenite or tempering austenite can be present in the tempered martensite. The nickel content is preferably 2.3 to 2.8, in particular 2.5% by weight.

Molybdenum:

Molybdenum improves creep strength by solid solution hardening as a partially dissolved element and precipitation hardening during a long-term stress. An excessively high proportion of this element, however, leads to embrittlement during long-term aging, which is due to the precipitation and coarsening of Laves phase (W, Mo) and Sigma phase (Mo). The range for Mo is 1.5 to 2% by weight, preferably 1.6 to 1.8% by weight, in particular 1.7% by weight.

Vanadium and nitrogen:

Together, these two elements significantly control grain size formation and precipitation hardening. A slightly more than stoichiometric V / N ratio sometimes also increases the stability of the vanadium nitride over the chromium nitride. The specific content of nitrogen and vanadium nitrides depends on the optimum volume fraction of the vanadium nitrides, which are to remain as insoluble primary nitrides during the solution annealing. The greater the total content of vanadium and nitrogen, the greater is the proportion of vanadium nitrides, which no longer goes into solution and the greater the grain refining effect. However, the positive effect of grain refining on ductility is limited because with increasing volume fraction of primary nitrides, the primary nitrides themselves limit ductility. The preferred content of nitrogen is in the range from 0.02 to 0.08% by weight, preferably 0.025 to 0.055% by weight, particularly preferably 0.04% by weight N, and that of vanadium is in the range between 0.1 and 0.4% by weight. , preferably 0.2 to 0.3% by weight, and especially at 0.25% by weight.

Niobium:

Niobium is a strong nitride former that aids the grain refining effect. In order to keep the volume fraction of the primary nitrides small, their total proportion must be limited to 0.1% by weight. Niobium dissolves in vanadium nitride in small amounts and can thus improve the stability of the vanadium nitride. Niobium is added in the range between 0.01 and 0.06% by weight, preferably 0.02 to 0.04% by weight, and in particular at 0.03% by weight.

Phosphorus, sulfur, tin:

Together with silicon and manganese, these elements increase the embrittlement of long-term aging in the range between 350 and 500 ° C. These elements should therefore be limited to maximum tolerable levels (0.005 wt%).

tantalum:

Ta influences the creep resistance positively. Addition of 0.001 to 2% by weight of Ta has the effect, on the one hand, that due to the greater tendency of tantalum to form carbides as chromium, the precipitation of undesired chromium carbides at the grain boundaries and, on the other hand, the undesirable depletion of the mixed crystal in chromium are reduced. The preferred range for Ta is 0.005 to 0.1% by weight, in particular a Ta content of 0.01% by weight should be set.

Carbon:

Carbon forms chromium carbides on tempering, which are conducive to improved creep resistance. At too high carbon contents, however, the resulting increased volume fraction of carbides leads to a ductility reduction, which comes into play in particular by the carbide coarsening during long-term storage. The carbon content should therefore be limited upwards to 0.16% by weight. Another disadvantage is the fact that carbon increases the hardening during welding. The preferred carbon content is in the range between 0.10 and 0.14% by weight, preferably 0.12% by weight.

Boron:

Boron stabilizes the M ₂₃ C ₆ precipitates and thus improves the creep resistance of the steel, but the formation of boron nitrides has to be prevented at the expense of the vandium carbonitrides. In addition, however, it should be noted that the austenitizing temperature must be increased in order to obtain homogeneous boron in the matrix, which in turn leads to an increase in the particle size and thus to poorer properties of the material. Therefore, the boron content should be limited to 40 to 120 ppm. Preferably, a B content of 50 ppm to 120 ppm, more preferably of about 70 ppm to adjust.

lanthanum:

Lanthanum binds the sulfur in the steel through the formation of lanthanum sulfide La ₂ S ₃ . La ₂ S ₃ is much more stable than MnS ₂ . It has a melting point of> 2100 ° C, while MnS ₂ decomposes at high temperatures to release S. However, since certain contents of S are necessary on the one hand because of the steelmaking process and on the other hand as impurities are unavoidable, and sulfur adversely increases the embrittlement tendency of the material, stable sulfide formers in steel such as La are much better than Mn. In addition, the grain size is advantageously reduced by the micro-alloying with La, which also has an advantageous effect if the material is tested non-destructively by ultrasonic methods. For example, in the case of a 12% Cr steel doped with B at an austenitizing temperature of 1100 ° C., a grain size ASTM 6 was determined by the applicant, whereas for a 12% Cr steel microblasted with B and La, the grain size at the same austenitizing temperature only remains ASTM 7 was. In addition, because of the very high stability of the lanthanum sulfides and the positive effect on the obstruction of interdental weld cracks, the weldability of the 12% Cr steels is improved. However, since lanthanum also adversely forms oxides, the content of La should be 0.001 to 0.5% by weight, preferably 0.01 to 0.1% by weight, especially 0.05% by weight.

Palladium:

Pd forms an ordered Fe-Pd L1 ₀ intermetallic phase with the iron of the steel, the α "phase. This stable α" phase increases the high temperature creep strength by stabilizing grain boundary precipitates such as M ₂₃ C ₆ and acts thus have a positive effect on the creep properties. However, palladium has the disadvantage of high costs. The Pd content of the proposed steel should be in the range of 0.0001 to 1, preferably 0.0005 to 0.01 wt%, with a content of 0.001 wt% being particularly suitable.

Short description of the drawing

Fig. 1

eine graphische Darstellung, bei welcher die Spannungen ausgewählter Legierungen (nach dem Stand der Technik VL1 bzw. entsprechend der vorliegenden Erfindung L2) bei einer Temperatur von 550 °C über der mittleren Zeit bis zum Bruch bzw. bis zur 1 % Dehnung des Materials aufgetragen sind;

Fig. 2

eine graphische Darstellung analog zu Fig.1, aber bei einer Temperatur von 450 °C;

Fig. 3

eine graphische Darstellung, bei welcher die Bruchzähigkeit (linkes Teilbild) und die Schlagenergie (rechtes Teilbild) für die beiden Legierungen VL1 und L2 bei Raumtemperatur im wärmebehandelten Zustand (ohne Auslagerung) gegenübergestellt sind und

Fig. 4

eine graphische Darstellung analog zu Fig. 3, bei welcher aber die Proben nach der Wärmebehandlung zusätzlich 3000 Stunden bei 480 °C ausgelagert wurden.

In the drawing, an embodiment of the invention is shown. Show it:

Fig. 1

a graph in which the voltages of selected alloys (according to the prior art VL1 and according to the present invention L2) at a temperature from 550 ° C over the mean time to break or to 1% elongation of the material, respectively;

Fig. 2

a graphic representation analogous to Fig.1 but at a temperature of 450 ° C;

Fig. 3

a graph in which the fracture toughness (left partial image) and the impact energy (right partial image) for the two alloys VL1 and L2 are compared at room temperature in the heat treated state (without outsourcing) and

Fig. 4

a graphic representation analogous to Fig. 3 , in which, however, the samples after the heat treatment were additionally stored for 3,000 hours at 480 ° C.

Ways to carry out the invention

Hereinafter, the invention with reference to an embodiment and the Fig. 1 to 4 explained in more detail.

The investigated inventive alloy L2 had the following chemical composition (in% by weight): 0.12 C, 11.5 Cr, 0.2 Mn, 2.5 Ni, 1.7 Mo, 0.25 V, 0.03 Nb, 0.04 N, 0.01 Ta, 0.05 La, 0.001 Pd, 0.0070 B, 0.05 Si, 0.005 P, 0.005 S, 0.005 Sn, balance iron and unavoidable impurities.

The comparative alloy VL1 used was a prior art commercial steel of the type X12CrNiMoV11-2-2 having the following chemical composition (in% by weight): 0.10-0.14 C, 11.0-12.0 Cr, 0.25 Mn, 2.0-2.6 Ni, 1.3-1.8 Mo, 0.2-0.35 V, 0.02-0.05 N, 0.15 Si, 0.026 P and 0.015 S.

Both alloys thus have a comparable composition with the difference that the inventive alloy L2 is additionally microalloyed with Nb, B, as well as La and Pd and contains Ta.

1. 1. Normalisierung bei 1100 °C / 3h / Ventilatorluft-Abkühlung auf Raumtemperatur
2. 2. Anlassbehandlung bei 640 °C / 5h / Luft-Abkühlung auf Raumtemperatur

The inventive alloy L2 and the comparative alloy VL1 were subjected to the following heat treatment processes:

1. 1. Normalization at 1100 ° C / 3h / fan air cooling to room temperature
2. 2. Tempering treatment at 640 ° C / 5h / air cooling to room temperature

From the materials treated in this way, samples were prepared for determining the mechanical properties. Long-term aging was carried out at 450 ° C., 480 ° C. or 550 ° C. under certain mechanical loads, and the impact energy and the fracture toughness at room temperature were determined. The results are in the FIGS. 1 to 4 shown.

Fig. 1 shows the creep properties for the two alloys VL1 and L2, ie the creep rupture strength and the 1% yield strength at 550 ° C. In this diagram, the mean times to break and to reach a 1% elongation are thus shown, depending on the voltage at 550 ° C.

It turns out that at the stated temperature, the alloy L2 according to the invention advantageously requires significantly longer times when the same voltage is applied until a 1% elongation is reached than the comparison alloy VL1. In the time to break (creep strength), this difference is even more evident, since the in Fig. 1 Alloy L2 samples with an arrow have not yet broken. In the case of the alloy L2 according to the invention, a clear shift to longer is here Recognize times, which is of particular advantage for the planned use as a gas turbine or steam turbine rotor.

In Fig. 2 the same dependencies are shown, but for a lower temperature (450 ° C). Although the differences between the behavior of the comparison layer VL1 and the alloy L2 according to the invention can also be seen here, in comparison to those in Fig. 1 These results are not so serious.

In Fig. 3 the fracture toughnesses and impact energies at room temperature are compared for the two investigated alloys in the above-described heat treatment state (without aging). Despite the significantly better creep properties at high temperatures (see Fig. 1 and 2 ), the toughness of the alloy according to the invention scarcely deteriorates.

The influence of a long-term aging at a temperature of 480 ° C shows Fig. 4 , There, the fracture toughness and impact energy at room temperature for the two investigated alloys L2 and VL1 after aging for 3,000 hours at 480 ° C are shown. Despite the significantly better creep properties at high temperatures (significantly better strength values), the toughness properties of the alloy L2 according to the invention are virtually unchanged compared to VL1.

This very good combination of properties (very high creep resistance at temperatures of 450 ° C. and significantly higher, good toughness properties after long-term aging at high temperatures) is achieved in comparison to the prior art by the combination of the alloying elements B, Ta, La and Pd in the specified ranges ,

In summary, it should be said that the alloy according to the invention is distinguished on the one hand by a very good creep resistance at temperatures of 450 ° C., preferably 550 ° C., and above, and is thus superior to the conventional 12% Cr steels. This is mainly due to the influence of boron, tantalum and palladium, which are alloyed in the specified range. Boron, tantalum and palladium stabilize the M ₂₃ C ₆ precipitates, which play a significant strengthening role during creep, with Pd additionally forming a stable intermetallic phase with the iron, which also contributes to increasing creep resistance. In addition, the dislocation density is maintained until breakage, thus improving creep strength of the steel. A similar effect on increasing creep strength is found in Ta and Pd. On the other hand, the alloy of the present invention has improved resistance to embrittlement upon long-term aging and comparatively high toughness. This is due to the addition of lanthanum in the specified range, because both the grain size is reduced as well as stable lanthanum sulfide La ₂ S _{3 are} formed.

The inventive alloy is thus particularly advantageous for rotors in gas and steam turbines, which are exposed to high inlet temperatures of about 550 ° C, can be used advantageously.

Of course, the invention is not limited to the embodiment described.

Claims (27)

1. Creep-resistant steel, characterized by the following chemical composition (given in % by weight): 0.08 to 0.16 C, 9.0 to 12.0 Cr, 0.1 to 0.5 Mn, 2.3 to 3 Ni, 1.5 to 2.0 Mo, 0.1 to 0.4 V, 0.01 to 0.06 Nb, 0.02 to 0.08 N, 0.001 to 2 Ta, 0.001 to 0.5 La, 0.0001 to 1 Pd, 0.004 to 0.012 B, maximum 0.005 P, maximum 0.005 S, maximum 0.05 Si, maximum 0.005 Sn, and the remainder iron and unavoidable impurities.
2. Creep-resistant steel according to claim 1, characterized by 2.3 to 2.8% Ni.
3. Creep-resistant steel according to claim 2, characterized by 2.5% Ni.
4. Creep-resistant steel according to claim 1, characterized by 10 to 12% Cr.
5. Creep-resistant steel according to claim 4, characterized by 10.5 to 11.5% Cr.
6. Creep-resistant steel according to claim 1, characterized by 0.10 to 0.14% C.
7. Creep-resistant steel according to claim 6, characterized by 0.12% C.
8. Creep-resistant steel according to claim 1, characterized by 0.10 to 0.25% Mn.
9. Creep-resistant steel according to claim 8, characterized by 0.20% Mn.
10. Creep-resistant steel according to claim 1, characterized by 1.6 to 1.8% Mo.
11. Creep-resistant steel according to claim 10, characterized by 1.7% Mo.
12. Creep-resistant steel according to claim 1, characterized by 0.2 to 0.3% V.
13. Creep-resistant steel according to claim 12, characterized by 0.25% V.
14. Creep-resistant steel according to claim 1, characterized by 0.02 to 0.04% Nb.
15. Creep-resistant steel according to claim 14, characterized by 0.03% Nb.
16. Creep-resistant steel according to claim 1, characterized by 0.025 to 0.055% N.
17. Creep-resistant steel according to claim 16, characterized by 0.04% N.
18. Creep-resistant steel according to claim 1, characterized by 0.005 to 0.012% B.
19. Creep-resistant steel according to claim 18, characterized by 0.007% B.
20. Creep-resistant steel according to claim 1, characterized by 0.005 to 0.1% Ta.
21. Creep-resistant steel according to claim 20, characterized by 0.01% Ta.
22. Creep-resistant steel according to claim 1, characterized by 0.01 to 0.1% La.
23. Creep-resistant steel according to claim 22, characterized by 0.05% La.
24. Creep-resistant steel according to claim 1, characterized by 0.0001 to 1% Pd.
25. Creep-resistant steel according to claim 24, characterized by 0.0005 to 0.01% Pd.
26. Creep-resistant steel according to claim 25, characterized by 0.001% Pd.
27. Creep-resistant steel according to any of claims 1 to 26, characterized in that it is used for rotors of thermal turbomachines.

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