CREEP-RESISTANT STEEL

Abstract

The disclosure relates to a creep-resistant steel which having a chemical composition (values in % by weight) of: about 0.10 to 0.15 C, 8 to 13 Cr, 0.1 to 0.5 Mn, 2 to 3 Ni; at least one or both elements from the group Mo, W in a range in each case of about 0.5 to 2.0 or, if both elements are present, a maximum total of about 3.0; about 0.02 to 0.2 Nb, 0.05 to 2 Ta, 0.1 to 0.4 V, 0.005 to 2 Pd, 0.02 to 0.08 N, 0.03 to 0.15 Si; and about 80 to 120 ppm B, maximum about 100 ppm Al, maximum about 150 ppm P, maximum about 250 ppm As, maximum about 120 ppm Sn, maximum about 30 ppm Sb, maximum 50 ppm S, a remainder of the composition being iron and impurities.

Description

RELATED APPLICATION

This application claims priority under 35 U.S.C. §119 to Swiss Patent Application No. 00270/08 filed in Switzerland on Feb. 25, 2008, the entire content of which is hereby incorporated by reference in its entirety.

TECHNICAL FIELD

The disclosure relates to steels based on 8-13% chromium which are used for rotors in the power station sector. It relates to the selection and the co-ordination, in terms of quantity fractions, of special alloying elements which make it possible to set an unusually high creep resistance, at temperatures of 550° C. and above, in this material. Moreover, the steel according to the disclosure can also have high resistance to fatigue under a low number of load cycles (LCF=Low Cycle Fatigue) and also high toughness after long-term ageing, so that it can be used both in gas turbines and in steam turbines.

BACKGROUND INFORMATION

Martensitically hardenable steel materials based on 9-12% chromium are materials in widespread use in power station technology. They were developed for application in steam power stations at operating temperatures of above 600° C. and steam pressures of above 250 bar, in order to increase the efficiency of the power stations. Under these operating conditions, the creep resistance and the oxidation resistance of the material play a particular part.

It is known that the addition of chromium in the abovementioned range not only allows high resistance to atmospheric corrosion, but also makes it possible to have the full hardenability of thick-walled forgings, such as are employed, for example, as monobloc rotors or as rotor disks in gas and steam turbines. Proven alloys of this type can contain about 0.08 to 0.2% carbon which, in solution, makes it possible to set a hard martensitic structure. A good combination of heat resistance and ductility of martensitic steels is made possible by an annealing treatment in which a particle-stabilized subgranular structure is formed as a result of the precipitation of carbon in the form of carbides, while at the same time the dislocation substructure is recovered. The annealing behavior and the properties resulting from this can be influenced effectively by the choice and the co-ordination, in terms of quantity fractions, of special carbide formers, such as, for example, Mo, W, V, Nb and Ta.

A representative which has been in widespread use in steam power stations, in particular as rotor steel, is the German steel known under DIN as X20CrMoV12.1.

It is known, furthermore, that ductility can be markedly improved at a strength level of 850 MPa by alloying with nickel. Such alloys are therefore in widespread use where both strength and ductility have to meet markedly higher requirements, such as disk materials for gas turbine rotors. A representative of such alloys which has been in widespread use in gas turbine technology, in particular as material for rotor disks, is the German steel known under DIN as X12CrNiMo12. However, the tendency is for nickel to have the adverse effect of lowering the heat resistance at high temperatures. This is related to reduced carbide stability in nickel-containing steels.

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

In such alloys, the contents of Cr, Mo and W were modified, taking into account N, Nb and/or B, in order to improve the creep and rupture strengths for applications at 600° C. The carbides, such as, for example, M₂₃C₆, are to be stabilized by the addition of boron. On account of the harmful effect of nickel on the long-term properties, in these steels the Ni contents were restricted to values of lower than 0.25%. In these alloys, a disadvantage is that the fracture toughness values are low, and, although this does not play an important part in steam turbine applications and may therefore be ignored, it should 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 specialist conference “Werkstoff-und Bauteilverhalten in der Energie-und Analagentechnik” [“Material and component behavior in power and plant technology”], 13./14.10.2005, Stuttgart), the Ni contents were disclosed as being limited even to the values of <0.002% in the case of an addition of B of 0.03% to a 10.5% Cr steel.

Particularly for gas turbine applications, efforts were made, in 9-12% Cr steels, either to improve the rupture strengths in the range of 450 to 500° C. at a high ductility level or to reduce the tendency to embrittlement at temperatures of between 425 and 500° C. European patent application EP 0 931 845 A1 describes a nickel-containing 12% chromium steel which is similar in constitution to the German steel X12CrNiMo12 and in which the element, molybdenum, was reduced, as compared with the known steel X12CrNiMo12, but an increased content of tungsten was added for alloying.

DE 198 32 430 A1 discloses a modification of a steel which is of the same type as X12CrNiMo12 and in which the tendency to embrittlement in the temperature range between 425 and 500° C. is limited by the addition of rare earth elements or boron.

A disadvantage is that the strength, in particular the heat resistance at temperatures of between 300 and 600° C., could not be improved in any of the abovementioned developed steels at a high ductility level comparable to that of the steel X12CrNiMo12.

One possible approach for improving the heat resistance, while at the same time having high ductility, was the development of steels having increased nitrogen contents. EP 0 866 145 A2 describes a class of martensitic chromium steels with nitrogen contents in the range between 0.12 and 0.25%, and, in EP 1 158 067 A1, with nitrogen contents of 0.12 to 0.18%, the weight ratio V/N lying in the range between 3.5 and 4.2. In these steels, the entire structural generation is controlled by the formation of special nitrides, in particular of vanadium nitrides, which can be distributed in many different ways by forging treatment, by austenitization, by controlled cooling treatment or by annealing treatment. While strength is achieved via the hardening action of the nitrides, an aim is to set a high ductility by the distribution and morphology of the nitrides, but, above all, by limiting the granular coarsening during forging and during solution heat treatment.

A heat-resistant steel with good toughness properties for use as a turbine rotor is known from EP 0867 522 A2 and 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, the rest iron and unavoidable impurities. Microalloying with boron can lead to precipitations at the grain boundaries and increase the long-term stability of the carbonitrides at high temperatures, although higher contents of B reduce the toughness of the steel. Disadvantages of this proposed composition are also the relatively high permitted Si values in the amount of 0.2%. Although Si advantageously serves as a deoxidant at the melting time point, on the other hand, parts of this remain as oxides in the steel, and this is reflected adversely in a reduced toughness.

The document U.S. Pat. No. 5,906,791 describes stainless steels with 8-13% by weight Cr, which have inter alia boron and rare earths in their chemical composition in order to increase resistance to embrittlement under long-term ageing. According to this document, the maximum content of rare earths, for example Y, La, Ce, Er, is to amount to 0.5% by weight, an optimal fraction being given as 0.1% by weight. The boron content is given as 0.001-0.04% by weight. Moreover, the steels also have the following elements (values in % by weight): 0.08-0.15 C, at least one element from the group of noble metals, such as Ru, Rh, Os, Pt, Pd, Ir, in the range of 0.01-2.00, 0.01-0.1 Si, at least one element from the group of W and Mo in the range of 0.50-4.00, at least one austenite stabilizer (such as Ni, Co, Mn Cu) in the range of 0.001-6.00, 0.25-0.40 V, 0.001-0.025 Al, max. 0.01 P, max. 0.004 S, max. 0.060 N, max. 2 ppm H, max. 50 ppm 0, max. 0.006 As, max. 0.003 Sb, max. 0.0050 Sn, the rest Fe. In a special embodiment, the steel may additionally contain up to 0.50% by weight Nb. In this case, it is described, with regard to the austenite stabilizers, that the steel is to contain as much Co as possible, while at the same time the Ni content is to be minimized. According to the authors' statement, this balance between the Ni and Co content is used to suppress undesirable embrittlement phenomena and at the same time to ensure the desired toughness of the steel. With these steels, good properties in high-temperature applications are to be achieved, that is to say balanced mechanical and oxidation properties. For example, a steel for high-temperature turbine components is thereby to be made available, which has high resistance to embrittlement, oxidation and creep.

SUMMARY

A creep-resistant steel is disclosed, comprising: a chemical composition (values in % by weight) of: about 0.10 to 0.15 C, 8 to 13 Cr, 0.1 to 0.5 Mn, 2 to 3 Ni; at least one or both elements from the group Mo, W in each case in a range of about 0.5 to 2.0 or, if both elements are present, a maximum total of about 3.0; about 0.02 to 0.2 Nb, 0.05 to 2 Ta, 0.1 to 0.4 V, 0.005 to 2 Pd, 0.02 to 0.08 N, 0.03 to 0.15 Si; about 80 to 120 ppm B, maximum about 100 ppm Al, maximum about 150 ppm P, maximum about 250 ppm As, maximum 120 ppm Sn, maximum 30 about ppm Sb, maximum about 50 ppm S, a remainder of the composition being iron and impurities.

A creep-resistant steel is disclosed, comprising: a chemical composition (values in % by weight) of: iron; 8 to 13 Cr; designated percentages by weight of Mn, Ni, Nb, Ta, V, N and Si; about 0.005 to 2 Pd; about 80 to 120 ppm B; and at least one of Mo and W.

BRIEF DESCRIPTION OF THE DRAWING

An exemplary embodiment of the disclosure is illustrated in the drawings in which:

FIG. 1 shows a graphical illustration in which stresses of selected alloys (VL1 according to the prior art and an exemplary L1 according to the present disclosure) are plotted against time, at a temperature of 550° C., up to the fracture of the material, two different heat treatment methods (with two different annealing temperatures) having been used for alloy L1;

FIG. 2 shows a graphical illustration in which an exemplary elongation amplitude is plotted against the number of load cycles up to incipient cracking at 575° C. for the alloy L1 according to the disclosure and at 500° C. for the comparative alloy VL1, and

FIG. 3 shows a graphical illustration in which the fracture toughness (left part image) and the notched bar impact work (right part image) at room temperature are compared for the two alloys L1 and VL1 after heat treatment and additional age hardening of 3000 hours at 480° C.

DETAILED DESCRIPTION

Exemplary embodiments disclosed herein can provide an 8-13% Cr steel which can possess an increased creep resistance at temperatures of 550° C. and above, and also possess improved LCF properties and comparatively high toughness relative to known steel. Exemplary embodiments can be used for rotors of thermal turbomachines, so that the efficiency and output of these machines can be increased, as compared with known steel.

As disclosed herein, an exemplary steel can include (e.g., consist of) a chemical composition (values in % by weight) of: about 0.10 to 0.15 C, 8 to 13 Cr, 0.1 to 0.5 Mn, 2 to 3 Ni; at least one or both elements from the group Mo, W in each case in a range of about 0.5 to 2.0 or, if both elements are present, a maximum total of about 3.0; about 0.02 to 0.2 Nb, 0.05 to 2 Ta, 0.1 to 0.4 V, 0.005 to 2 Pd, 0.02 to 0.08 N, 0.03 to 0.15 Si; and about 80 to 120 ppm B, maximum about 100 ppm Al, maximum about 150 ppm P, maximum about 250 ppm As, maximum about 120 ppm Sn, maximum about 30 ppm Sb, maximum about 50 ppm S, a remainder of the composition being iron and impurities (e.g., unavoidable impurities). As referenced herein, the term “about” refers to an amount (e.g., percentage) substantially at the designated value so as to achieve the advantages and effects described herein and can, for example, be on the order of ±10% of the specified value.

Exemplary preferred ranges for individual alloying elements of the composition according to the disclosure include chemical composition (values in % by weight) of: 0.12 C, 11.5 Cr, 0.2 Mn, 2.5 Ni, 1.7 Mo, 0.25 V, 0.03 Nb, 0.06 Ta, 50 ppm Pd, 100 ppm B, 0.04 N, <0.01 Al, <0.01 P, <0.005 S, <0.05 Si, <0.012 Sn, <0.025 As, <0.0025 Sb, a remainder of the composition being iron and impurities (e.g., unavoidable impurities).

An exemplary advantage of the alloy according to the disclosure is creep properties at temperatures of 550° C. and above, in comparison with known alloys of similar composition without a B addition or without a Pd addition. Improved toughness properties and higher fatigue strength (LCF) can also be achieved in accordance with exemplary embodiments.

A tempered structure can be set which is distinguished by a tough basic matrix and by a presence of nitrides, borides and carbides which afford heat resistance. The toughness of the basic matrix can be set by the presence of substitution elements, such as by nickel. The contents of these substitution elements can be determined such that they allow an optimal development both of martensitic hardening and of particle hardening due to the precipitation of special nitrides, for example vanadium nitrides or niobium nitrides, for the purpose of setting the highest possible heat resistances.

Both hardening mechanisms can lower the ductility. A ductility minimum is in this case characteristically observed in the region of secondary hardening. This ductility minimum need not be caused solely by the actual precipitation hardening mechanism. A certain contribution to embrittlement may also be made by the segregation of impurities up to the grain boundaries or, possibly, also by near-order settings of dissolved alloy atoms.

A rise in the annealing temperature over the secondary hardening range leads to complete precipitation, with a marked growth of carbides. As a result, the strength decreases and the ductility increases. Ductility can increase to a great extent due to the simultaneous recovery of the dislocation substructure and particle coarsening, so that the combination of strength and ductility, overall, can be improved. This improvement can be attributed to the formation of a particle-stabilized subgranular structure. It can be assumed, in this case, that both the ductility and the strength of particle-stabilized subgranular structures can be reduced by nonuniformities in the topology of the particle subgranular structure.

Precipitations at subgranular boundaries can be subject to accelerated coarsening and can tend to coagulate with adjacent precipitations. Coarse and coagulated phases can generate fracture-triggering stress peaks which lower the ductility. Above all, however, the hardening mechanism which is the most effective at high temperatures, to be precise the particle hardening, can be seriously limited by the nonuniform distribution of the precipitations.

One measure for increasing the ductility in known martensitically hardenable steels is alloying with nickel. The causes of this, however, are not known in all points and would seem to depend greatly on the nickel content. Thus, even small fractions of nickel may be highly conducive to ductility if, for example, the formation of delta ferrite can thereby be suppressed completely. By contrast, with nickel contents of above 2% by weight, it is expected that nickel lowers the Ac1 temperature (this is the temperature at which ferrite begins to be converted into austenite during heating) to temperatures of below 700° C. If, therefore, the strength is to be increased by a lowering of the annealing temperature to below 700° C., then, if increased nickel contents are present, a partial conversion of ferrite into austenite can be reckoned on during annealing. This is associated with a certain ductility-promoting grain reforming. By contrast, however, the carbide precipitation takes place only incompletely above the Ac1 temperature, since the solubility of the austenite-stabilizing element, carbon, is higher in austenite than in ferrite. Further, the austenite which is formed is not sufficiently stabilized, and therefore a larger volume fraction of the reformed austenite can be subjected to further martensitic conversion during recooling after annealing. In addition to the two abovementioned contributions of nickel with the effect of an increase in ductility, a certain contribution to ductility by nickel can be made in its action as a substitution element in solid solution. This can be explained by electron theory in that the element, nickel, feeds additional free electrons into the iron lattice and thereby makes the iron alloys even more “metallic”.

Known martensitically hardenable steels which are alloyed with nickel do not have any special heat resistance advantages, as compared with low-nickel alloys. This applies, at least, to test temperatures of above 500° C. and, in the case of increased nickel contents, could be connected with the abovementioned reaustenitization during annealing. It is known, furthermore, that the alloying of such steels with nickel can markedly intensify the structural instability under long-term age-hardening conditions at increased temperatures. This long-term structural instability can in this case be related to an accelerated coarsening of the carbides.

Manganese lies on the left side next to the element iron, in the periodic system of elements. It is an electron-leaner element, and therefore its action in a solid solution should be markedly different from that of nickel. Nonetheless, it is an austenite-stabilizing element which can greatly lower the Ac1 temperature, and does not have an especially positive, but, instead, a somewhat adverse effect on the ductility. With regard to carbon-containing 12% chromium steels, manganese is understood to be a contaminating element which appreciably promotes annealing embrittlement. The manganese content can therefore be limited to very small quantities.

Exemplary quantities in percentages by weight for each element and the reasons for the selected alloy ranges according to the disclosure, in their relation to the heat treatment possibilities resulting from them, are listed below.

Chromium:

A weight fraction of 8-13% chromium allows a good full hardenability of thick-walled components and ensures sufficient oxidation resistance up to a temperature of 550° C. A weight fraction of below 8% can be detrimental to full hardening. Contents above 13% can lead to the accelerated formation of hexagonal chromium nitrides during the annealing operation, which, in addition to nitrogen, can also bind vanadium, and, consequently, lower the effectiveness of hardening by vanadium nitrides. An exemplary optimal chromium content is about 11 to 12%.

Manganese and Silicon:

These elements are conducive to annealing embrittlement and can therefore be limited to very small contents. Taking into account the metallurgical possibilities, the range to be specified can lie, for manganese, in the range between about 0.1 and 0.5% by weight, preferably between about 0.1 and 0.25%, in particular at 0.2% by weight, and, for silicon, at about 0.03-0.15, preferably at <0.05% by weight.

Nickel:

Nickel can be used as an austenite-stabilizing element for the suppression of delta ferrite. Furthermore, as a dissolved element in the ferritic matrix, it can improve ductility. Nickel contents of 2 to about 3% by weight can be expedient. Nickel contents above 4% by weight can intensify the austenite stability in such a way that, after solution heat treatment and annealing, an increased fraction of residual austenite or annealing austenite may be present in the hardened martensite. The nickel content lies at, for example, about 2.3 to 2.7, in particular at 2.5% by weight.

Molybdenum and Tungsten:

Molybdenum and tungsten can improve the creep resistance by solid solution hardening as partially dissolved elements and by precipitation hardening during long-term stress. However, an excessively high fraction of these elements can lead to embrittlement during long-term age hardening, which can be due to the precipitation and coarsening of the Laves phase (W, Mo) and sigma phase (Mo). An exemplary desired range for Mo and W is in each case about 0.5 to 2% by weight, preferably about 1.6 to 1.8% by weight, in particular 1.7% by weight. If both elements are present, the overall fraction can amount to a maximum of about 3% by weight.

Vanadium and Nitrogen:

These two elements together can control (e.g., signifcantly control) the grain size formation and the precipitation hardening. A slightly above-stoichiometric V/N ratio sometimes can also increase the stability of the vanadium nitride with respect to the chromium nitride. The actual content of nitrogen and vanadium nitrides can depend on the optimal volume fraction of the vanadium nitrides which can remain as insoluble primary nitrides during the solution heat treatment. The larger the overall fraction of vanadium and nitrogen is, the larger that fraction of the vanadium nitrides which is no longer dissolved is, and the higher the grain-refining action. However, the positive influence of grain refinement on ductility can be limited, since, with an increasing volume fraction of primary nitrides, the primary nitrides themselves limit the ductility. An exemplary preferred nitrogen content lies in the range of about 0.02 to 0.08% by weight, preferably about 0.025 to 0.055% by weight, particularly preferably at 0.04% by weight N, and the vanadium content lies in the range of between about 0.1 and 0.4% by weight, preferably about 0.2 to 0.3% by weight, and, in particular, at 0.25% by weight.

Niobium:

Niobium is a strong nitride former which promotes the grain-refining action. In order to keep the volume fraction of primary nitrides low, its overall fraction can be limited. Niobium dissolves in small quantities in vanadium nitride and can consequently improve the stability of the vanadium nitride. Niobium can be added for alloying in the range of between about 0.02 and 0.2% by weight, preferably about 0.02 to 0.04% by weight, and, in particular, at 0.03% by weight.

Phosphorus, Tin:

These elements, together with silicon and manganese, can intensify annealing embrittlement during long-term age hardening in the range between about 350 and 500° C. These elements can therefore be limited to maximum acceptable fractions (150 ppm P, 120 ppm Sn).

Tantalum:

Ta can positively influence the creep resistance. Alloying with about 0.05 to 2% by weight Ta has the effect that, because of the greater tendency of tantalum to form carbides than chromium, on the one hand, the precipitation of undesirable chromium carbides at the grain boundaries can be diminished and, on the other hand, the undesirable depletion of the chromium mixed crystal can also be reduced. An exemplary preferred range for Ta is about 0.05 to 0.1% by weight, and, in particular, a Ta content of about 0.06% by weight should be set.

Carbon:

Carbon, during annealing, forms chromium carbides which are conducive to improved creep resistance. At carbon contents which are too high, however, the increased volume fraction of carbides which results from this can lead to a ductility reduction which, in particular, can be reflected by carbide coarsening during long-term age hardening. The carbon content can therefore have an upper limit of about 0.15% by weight. An exemplary disadvantage can be that carbon intensifies hardening during welding. An exemplary preferred carbon content lies in the range between about 0.10 and 0.14% by weight, preferably at 0.12% by weight.

Boron:

Boron stabilizes the M₂₃C₆ precipitations, and hence can improve the creep resistance of the steel and reduce the annealing embrittlement, although the formation of boron nitrides at the expense of the vanadium carbonitrides can be prevented. However, the austenitization temperature can be increased in order to obtain homogeneous boron in the matrix, but this, in turn can lead to an increase in the grain size and consequently to poorer properties of the material. The boron content can therefore be limited to about 80 to 120 ppm. Also, an exemplary B content of about 100 ppm can preferably be set.

Palladium:

Pd forms, with the iron of the steel, an ordered intermetallic Fe—Pd L1₀ phase, the α″ phase. This stable α″ phase can increase the rupture strength at high temperatures by the stabilization of the grain boundary precipitations, such as, for example, M₂₃C₆, and therefore has a positive effect on the creep properties. However, palladium can have the exemplary disadvantage of high costs. The Pd content of the proposed steel should lie in the range of about 0.005 to 2, preferably of about 0.005 to 0.01% by weight, a content of 0.005% by weight, that is to say 50 ppm, of Pd being particularly suitable.

Aluminum, Antimony, Arsenic, Sulfur:

Small contents of these elements (e.g., maximum 250 ppm As, maximum 30 ppm Sb, maximum 100 ppm Al, maximum 50 ppm S) cause controlled segregation and secondary phase formation, so that a very clean steel of this type can have increased toughness properties.

The disclosure is explained in more detail below with reference to an exemplary embodiment and to FIGS. 1 to 3.

The investigated alloy L1 according to the disclosure had the following chemical composition (values in % by weight): 0.12 C, 11.5 Cr, 0.2 Mn, 2.5 Ni, 1.7 Mo, 0.25 V, 0.03 Nb, 0.06 Ta, 0.04 N, 0.005 Pd, 0.01 B, <0.01 Al, <0.01 P, <0.005 S, <0.05 Si, <0.012 Sn, <0.025 As, <0.0025 Sb, the remainder of the composition being iron and unavoidable impurities.

The comparative alloy VL1 used was a commercial steel of the type X12CrNiMoV11-2-2 which is known from the prior art and which includes (e.g., consists of) chemical composition (values in % by weight) of: 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, 0.015 S, the remainder of the composition being Fe and impurities (e.g., unavoidable impurities).

The two alloys therefore have an approximately comparable composition, the difference being that the alloy L1 according to the disclosure is additionally microalloyed with Nb, B, Ta and Pd.

The alloy L1 according to the disclosure was subjected to the following exemplary two-stage heat treatment processes:

• 1. Solution heat treatment at 1100° C. and, subsequently,
• 2. “A”: annealing treatment at 670° C. or

“B”: annealing treatment at 640° C.

The exemplary comparative alloy VL1 was solution heat-treated at 1065° C. and was subsequently subjected to annealing treatment at 640° C.

Samples for determining mechanical properties were produced from the materials treated in this way. Long-term age hardenings at 550° C. were carried out under specific mechanical loads, and the notched bar impact toughness and fracture toughness at room temperature and the LCF fatigue behavior at 500° C. and 575° C. were determined. The results are illustrated in FIGS. 1 to 3.

FIG. 1 shows the properties during creeping, that is to say the rupture strength, at 550° C. for the two alloys VL1 and L1. This graph thus illustrates the average times up to fracture as a function of the stress at 550° C.

It is shown that, at the given temperature, the exemplary alloy L1 according to the disclosure advantageously involves, both after heat treatment “A” and after heat treatment “B”, longer times under the action of the same stress up to fracture than the comparative alloy VL1. The sample, given an arrow in FIG. 1, of the alloy L1 has not yet even been fractured. Here, therefore, in the case of the alloy L1 according to the disclosure, a marked shift toward longer times can be seen, and this can be particularly advantageous for the planned use as a gas turbine rotor or steam turbine rotor.

In FIG. 2, an exemplary elongation amplitude is plotted against the number of load cycles up to incipient cracking at 575° C., with a holding time of 10 minutes in the tensile range, for the alloy L1 according to the disclosure. These results are compared with an average value for the comparative alloy VL1 at 500° C. and likewise with a holding time of 10 minutes in the tensile range. It can be seen that the experimentally determined values of L1 at 575° C. lie on the curve of VL1 at 500° C. This means that an improved LCF behavior can be achieved with the alloy according to the disclosure, since the same properties are achieved at a temperature which is 75° C. higher. This is a very considerable improvement.

In FIG. 3, the fracture toughness and the notched bar impact work at room temperature are compared for the two alloys investigated, after the above-described heat treatment state with subsequent age hardening (3000 h at 480° C.). In spite of the markedly better creep properties at high temperatures (see FIG. 1), in the exemplary alloy according to the disclosure there is scarcely any impairment in the fracture toughness and the notched bar impact work is slightly increased. The exemplary alloy L1 according to the disclosure therefore has no greater tendency to embrittlement than the comparative alloy VL1.

This very good property combination (very high creep resistance at temperatures of 550° C. and above, good toughness properties after long-term age hardening at high temperatures, and, moreover, very high fatigue strength at these high temperatures) can be achieved, as compared with the prior art, by the alloying elements as a whole, for example, by the combination of B, Ta and Pd in the exemplary ranges specified.

In summary, on the one hand, the exemplary alloy according to the disclosure can be distinguished by a very high creep resistance and high resistance to low-cycle fatigue at temperatures of 550° C. and above and can be consequently superior to the known 12% Cr steels. This is attributable at least in part to the influence of boron, tantalum and palladium which are added for alloying in the specified range. Boron, tantalum and palladium can stabilize the M₂₃C₆ precipitations which can play a substantial consolidating part during creeping, Pd additionally forming a stable intermettalic phase with the iron, this also contributing to increasing the creep resistance. In addition, the dislocation density up to fracture can be maintained and therefore the strength of the steel can be improved. On the other hand, exemplary alloys according to the disclosure can have improved resistance to embrittlement during long-term ageing and comparatively high toughness and also high resistance to fatigue.

Exemplary alloys according to the disclosure can therefore advantageously be used particularly for rotors in gas and steam turbines which are exposed to high inlet temperatures of, for example, 550° C. and above.

The disclosure, of course, is not restricted to the exemplary embodiments described.

It will be appreciated by those skilled in the art that the present invention can be embodied in other specific forms without departing from the spirit or essential characteristics thereof. The presently disclosed embodiments are therefore considered in all respects to be illustrative and not restricted. The scope of the invention is indicated by the appended claims rather than the foregoing description and all changes that come within the meaning and range and equivalence thereof are intended to be embraced therein.

Claims

1. A creep-resistant steel, comprising:

a chemical composition (values in % by weight) of: about 0.10 to 0.15 C, 8 to 13 Cr, 0.1 to 0.5 Mn, 2 to 3 Ni; at least one or both of the elements from the group Mo, W in each case in a range of about 0.5 to 2.0 or, if both elements are present, a maximum total of about 3.0; about 0.02 to 0.2 Nb, 0.05 to 2 Ta, 0.1 to 0.4 V, 0.005 to 2 Pd, 0.02 to 0.08 N, 0.03 to 0.15 Si; about 80 to 120 ppm B, maximum about 100 ppm Al, maximum about 150 ppm P, maximum about 250 ppm As, maximum about 120 ppm Sn, maximum about 30 ppm Sb, maximum about 50 ppm S, a remainder of the composition being iron and unavoidable impurities.

2. The creep-resistant steel as claimed in claim 1, comprising 2.3 to 2.7% Ni.

3. The creep-resistant steel as claimed in claim 2, comprising 2.5% Ni.

4. The creep-resistant steel as claimed in claim 1, comprising 11 to 12% Cr.

5. The creep-resistant steel as claimed in claim 3, comprising 11.5% Cr.

6. The creep-resistant steel as claimed in claim 1, comprising 0.10 to 0.14% C.

7. The creep-resistant steel as claimed in claim 6, comprising 0.12% C.

8. The creep-resistant steel as claimed in claim 1, comprising 0.10 to 0.25% Mn.

9. The creep-resistant steel as claimed in claim 8, comprising 0.20% Mn.

10. The creep-resistant steel as claimed in claim 1, comprising 1.6 to 1.8% Mo or 1.6 to 1.8% W.

11. The creep-resistant steel as claimed in claim 10, comprising 1.7% Mo or 1.7% W.

12. The creep-resistant steel as claimed in claim 1, comprising 0.2 to 0.3% V.

13. The creep-resistant steel as claimed in claim 12, comprising 0.25% V.

14. The creep-resistant steel as claimed in claim 1, comprising 0.02 to 0.04% Nb.

15. The creep-resistant steel as claimed in claim 14, comprising 0.03% Nb.

16. The creep-resistant steel as claimed in claim 1, comprising 0.025 to 0.055% N.

17. The creep-resistant steel as claimed in claim 16, comprising 0.04% N.

18. The creep-resistant steel as claimed in claim 1, comprising 0.01% B.

19. The creep-resistant steel as claimed in claim 1, comprising 0.05 to 0.1% Ta.

20. The creep-resistant steel as claimed in claim 19, comprising 0.06% Ta.

21. The creep-resistant steel as claimed in claim 1, comprising 0.005 to 0.1% Pd.

22. The creep-resistant steel as claimed in claim 21, comprising 0.005 to 0.01% Pd.

23. The creep-resistant steel as claimed in claim 22, comprising 0.005% Pd.

24. The creep-resistant steel as claimed in claim 1, comprising <0.05% Si.

25. The creep-resistant steel as claimed in claim 1, formed, at least in part, as a rotor of a thermal turbomachine.

26. A creep-resistant steel, comprising:

a chemical composition (values in % by weight) of:

iron; 8 to 13 Cr;

designated percentages by weight of Mn, Ni, Nb, Ta, V, N and Si;

about 0.005 to 2 Pd;

about 80 to 120 ppm B; and

at least one of Mo and W.