EP1158067B1 - Martensitic hardenable heat treatable steel with improved thermal resistance and ductility - Google Patents

Description

Technical area

The invention relates to martensitic-hardenable steels with increased nitrogen contents. It relates both to the choice and the proportionate tuning of special alloying elements, which allow the setting of an exceptionally good combination of heat resistance and ductility, as well as a method for heat treatment of the inventive alloy.

State of the art

Martensite-hardenable steels based on 9-12% chromium are widely used materials in power plant technology. 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 is the setting of a hard martensitic structure allows. 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.

Strengths above 850 MPa for 9-12% chromium steels can be adjusted by keeping the tempering temperature low, typically in the range 600 to 650 ° C. However, the use of low tempering temperatures leads to high transition temperatures from the brittle to the ductile state (above 0 ° C), with which the material exhibits brittle fracture behavior at room temperature. Significantly improved ductilities can be achieved if the tempered strength is lowered below 700 MPa. This is achieved by raising the tempering temperature to over 700 ° C. The use of increased tempering temperatures has the advantage that the set structural states are stable for longer times at elevated temperatures. 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 alloying nickel. For example, it is known that by alloying about 2 to 3% nickel, even after a tempering treatment at temperatures of 600 to 650 ° C, the transition temperature from brittle to ductile state is still below 0 ° C, resulting in a significantly improved overall combination of strength and ductility. Such alloys therefore find a wide there Use where significantly higher demands are placed on both strength and ductility, typically as disc 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.

In recent years, various efforts have been made to improve specific properties of these steels. For example, in the publication of 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 , describing the development of novel rotor steels for steam turbine applications. In such alloys, the contents of Cr, Mo, W were further optimized, taking into account about 0.03 to 0.07% N, 0.03 to 0.07% Nb and / or 50 to 100 ppm B, to creep and creep strengths for 600 ° C applications to improve.

On the other hand, especially for gas turbine applications, efforts have been made either to improve the creep ruptures in the range of 450 to 500 ° C at high ductility level or to reduce the embrittlement tendency at temperatures between 425 and 500 ° C. That's how it describes European patent application EP 0 931 845 A1 a nickel-containing 12% chromium steel similar in structure to the German steel X12CrNiMo12, in which the element molybdenum is reduced compared to the known steel X12 CrNiMo12, but an increased content of tungsten was added. 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 0.12 to 0.25%. In this class of steel, all structural formation is controlled by the formation of special nitrides, in particular vanadium nitrides, which can be distributed in a variety of ways through forging, austenitization, controlled cooling or annealing. While strength is achieved by the nitriding curing effect, the high ductility setting in this patent application is aimed at by the distribution and morphology of the nitrides, but most importantly, by limiting grain coarsening during forging and solution heat treatment. This is achieved in the cited document by an increased volume fraction as well as by a high particle coarsening resistance of sparingly soluble nitrides, so that a dense dispersion of nitrides can still effectively limit grain growth even at austenitizing temperatures of 1150 to 1200 ° C. The essential benefit of in EP 0 866 145 A2 The alloys listed above have the possibility of optimally influencing the combination of strength and ductility solely by the formation of nitrides with regard to distribution and morphology by a suitable definition of the heat treatment.

However, an optimized formation state of nitrides is only one factor for achieving maximum ductility. Another factor is the Effect of dissolved substitution elements such as nickel, cobalt and manganese expected. Manganese is known from carbon steels to be embrittling rather than promoting ductility. In particular, it causes embrittlement when the alloy is subjected to long-term annealing at temperatures in the range of 350 to 500 ° C. It is also known that nickel in carbon steels improves ductility but tends to lower high temperature hot strength as well. This is related to reduced carbide stability in nickel-containing steels. In contrast, the effect of cobalt on the combination of heat resistance and ductility is very little known even in carbonaceous 9-12% chromium steels.

Presentation of the invention

The invention has for its object to provide a martensitic-hardenable tempering steel with high ductility, which is characterized over the prior art, in particular the steel X12CrNiMo12, by an increased thermal stability at temperatures of 300 to 600 ° C. On the one hand, a suitable steel composition and, on the other hand, a heat treatment process for materials of this composition should be specified, which enables the formation of a ductile and, at the same time, heat-resistant martensitic tempering structure.

Core of the invention is a martensitic-hardenable tempering steel having the following composition (in wt .-%): 9 to 13% Cr, 0.001 to 0.25% Mn, 2 to 7% Ni, 0.001 to 8% Co, at least one of W and Mo in the sum between 0.5 and 4%, 0.5 to 0.8% V, at least one of Nb, Ta, Zr and Hf in the sum of 0.001 to 0.1%, 0.001 to 0.05% Ti, 0.001 to 0.15% Si, 0.01 to 0.1 % C, 0.12 to 0.18% N, maximum 0.025% P, maximum 0.015% S, maximum 0.01% Al, maximum 0.0012% Sb, maximum 0.007% Sn, maximum 0.012 % As, the remainder iron and common impurities, and the proviso that the weight ratio of vanadium to nitrogen V / N ranges between 3.5 and 4.2.

Preferred ranges for the individual alloying elements of the inventive composition are contained in the subclaims.

• Lösungsglühen bei 1050 bis 1250 °C,
• danach Abkühlung auf eine Temperatur unter 300°C,
• Anlassbehandlung, partielle oder vollständige Reaustenitisierung bei 600 bis 900 °C,
• Abkühlung auf eine Temperatur unter 300°C,
• Glühung bei einer Temperatur von 550 bis 650 °C.

The heat treatment process for the alloy according to the invention is characterized by the following steps:

• Solution annealing at 1050 to 1250 ° C,
• then cooled to a temperature below 300 ° C,
• Tempering treatment, partial or complete re-esterification at 600 to 900 ° C,
• Cooling to a temperature below 300 ° C,
• Annealing at a temperature of 550 to 650 ° C.

The advantage of the invention is that in the said alloy, a starting structure is set, which is characterized by a tough matrix and the presence of heat-resistant nitrides. The toughness of the base matrix is adjusted by the presence of substitution elements, preferably nickel and, secondarily, cobalt. The levels of these substitution elements are determined to permit optimum unfolding of both martensite setting and particle hardening by means of speederitrides, preferably vanadium nitrides, to set the highest levels of heat resistance.

In the known martensitic-hardenable 9-12% chromium steels, a good combination of heat resistance and ductility can be adjusted by a tempering treatment which includes austenitizing treatment, quenching treatment and tempering treatment. The achievable Strength is hereby significantly limited by the basic hardness of the quenched martensite and the potential particle hardening effect of precipitation phases which are formed in tempering. An increase in strength beyond the basic hardness of the quenched martensite is possible by a tempering treatment in the so-called secondary hardening range. For the 12% chromium steels, which are well known in power plant engineering, this temperature is between 450 and 530 ° C.

Basically, both hardening mechanisms, both martensitic hardening and precipitation hardening, reduce 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 settings 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. Ductile structures are formed in the low-nickel 9-13% chromium steels by a tempering treatment above 700 ° C. 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. Precipitates at sub-grain boundaries 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.

A limited in their effect ductility increase is possible by reducing the particle size. However, this is difficult to implement for the alloys known in the art in large components forging and leaves little effect. A slightly more important measure to increase ductility in conventional, martensitic-hardenable steels is the alloying of nickel. However, the causes of this are not known in all respects and are likely to 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. On the other hand, with nickel contents above 2% by weight, the temperature of Ac1 (that is, the temperature at which ferrite begins to convert to austenite during heating) is lowered to temperatures below 700 ° C. So if the strength is to be 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 is to 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 to the effect of nickel to increase the ductility, a certain ductility contribution of nickel in its effect as a substitution element can also be found come in solid solution. 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 alloying nickel into such steels significantly aggravates structural instability under long-term aging conditions at elevated temperatures. This long-term structural instability is related to an accelerated coarsening of the carbides.

Cobalt is an austenite stabilizing element similar to nickel. In solid solution, therefore, a nickel-like effect in terms of ductility is to be expected. From the chemical point of view, however, there is an important difference in that cobalt promotes ferromagnetism, which increases Curie temperature. Since the self-diffusion within the iron matrix increases abruptly when the Curie temperature is exceeded, all the diffusion-controlled recovery and coarsening processes are accelerated when the Curie temperature is exceeded. By increasing the Curie temperature, an improvement in the tempering resistance can thus also be expected. Cobalt-alloyed structures should then soften delayed during tempering, thus allowing adjustment of increased strength. Also of importance is the fact that cobalt significantly lowers the Ac1 temperature per weight percent alloy addition than nickel.

Unlike nickel and cobalt, manganese is on the left side of 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 nickel and cobalt. Nonetheless, it is an austenite stabilizing element which greatly lowers the Ac1 temperature, but leaves no particularly positive but rather unfavorable effect on ductility.

1. 1) Die erfindungsgemässe Legierung soll ein wirksames Kornneubildungsverhalten aufweisen, so dass durch Schmieden und Normalisierung (Austenitisierung) verfeinerte Korn- und Blockstrukturen erzeugt werden können. Diese Kornfeinung soll durch eine Dispersion schwerlöslicher Nitride erzeugt werden, welche durch das Zulegieren von Stickstoff zu starken Nitridbildnern wie Vanadium, Niob, Tantal, Titan, Hafnium und Zirkon erzeugt wird. Die Kornfeinung selbst kann einen Beitrag zur Warmfestigkeit liefern, sofern die eingestellte Korn- und Blockstruktur durch die Nitride gegen Vergröberung gut stabilisiert ist. Entscheidend ist jedoch, dass die günstige Wirkung der Kornfeinung auf die Duktilität gegenüber der negativen Wirkung der gröberen Primärphasen auf die Duktilität überwiegt.
2. 2) Eine verbesserte Warmfestigkeit wird durch eine thermisch stabile Ausscheidungsphase gewährleistet, welche in einem kleinen Volumenanteil, das heisst bei geringer Löslichkeit und damit hohem Widerstand gegen Vergröberung, eine maximale Teilchenhärtungswirkung pro Molvolumen bei erhöhten Prüftemperaturen ermöglicht. Die bevorzugte Ausscheidungsphase soll von demselben Typ sein wie diejenigen Nitride, welche die Korngrössenbegrenzung in Punkt 1 ermöglichen.
3. 3) Die Ausscheidungsreaktionen sollen gleichmässig verlaufen, so dass die Kombination von Duktilität und Warmfestigkeit nicht durch grobe, filmartige Ausscheidungen und deren ungleichmässige Verteilung auf den Korn- und Subkorngrenzen beeinträchtigt wird.
4. 4) Es soll Nickel zulegiert werden, um seine duktilitätsfördernde Wirkung als Substitutionselement zu nutzen und um die Ac1- und Ac3-Temperatur zu senken. Dieses Absenken der Umwandlungstemperatur ermöglicht die Ausscheidung von Nitriden bei tiefen Austenitisierungstemperaturen, nachdem das Material von Raumtemperatur auf die Aushärtungstemperatur aufgeheizt wird.
5. 5) Die Wirkung von Kobalt soll als komplementäres Substitutionselement zu Nickel zur Duktilitätssteigerung genutzt werden. Ungleich des Nickels soll es ausschliesslich als ein gelöstes Element und nicht zur Phasenumwandlungsbeeinflussung (Ferrit/Austenit) genutzt werden. Als Substitutionselement soll es auch die Anlassbeständigkeit verbessern.
6. 6) Die gewünschte Legierung, insbesondere das Einbringen von Stickstoff, soll unter der Rahmenbedingung hergestellt werden können, dass die Gehalte an versprödungsbringenden Elementen wie Silizium und Mangan den Anforderungen entsprechend gering gehalten werden können.

Based on these findings and hypotheses, the following alloy design is proposed for improvement from the combination of heat resistance and ductility:

1. 1) The alloy according to the invention should have an effective grain regeneration behavior, so that refined grain and block structures can be produced by forging and normalization (austenitization). This grain refining is to be produced by a dispersion of sparingly soluble nitrides, which is produced by the alloying of nitrogen to form strong nitride formers such as vanadium, niobium, tantalum, titanium, hafnium and zirconium. The grain refining itself can contribute to the heat resistance provided that the set grain and block structure is well stabilized by the nitrides against coarsening. Crucial, however, is that the favorable effect of the grain refining on the ductility over the negative effect of the coarser primary phases outweighs the ductility.
2. 2) An improved heat resistance is ensured by a thermally stable precipitation phase, which in a small volume fraction, that is, with low solubility and thus high resistance to coarsening, a maximum particle hardening effect per mol volume at elevated test temperatures. The preferred excretion phase should be of the same type as those nitrides which allow the grain size limitation in point 1.
3. 3) The precipitation reactions should be uniform, so that the combination of ductility and heat resistance is not affected by coarse, film-like precipitates and their uneven distribution on the grain and sub-grain boundaries.
4. 4) Nickel should be alloyed in order to use its ductility-promoting effect as a substitution element and to lower the Ac1 and Ac3 temperature. This lowering of the transition temperature allows the precipitation of nitrides at low austenitizing temperatures after the material is heated from room temperature to the curing temperature.
5. 5) The effect of cobalt should be used as a complementary substitution element to nickel for ductility increase. Unlike nickel, it should be used exclusively as a solute element and not for phase change (ferrite / austenite). As a substitution element, it should also improve tempering resistance.
6. 6) The desired alloy, in particular the introduction of nitrogen, should be able to be produced under the condition that the contents of brittleness-bearing elements such as silicon and manganese can be kept correspondingly low according to requirements.

Out EP 0 866 145 A2 It is known that the controlled alloying of nitrogen, vanadium and other special nitride formers, such as niobium, titanium, tantalum, zirconium and hafnium, satisfies the first three points and therefore provides a better opportunity for steel alloy development Properties can be used. A key role is played by vanadium nitride, which can be used effectively for grain refining as well as precipitation hardening. It is crucial that a tempering treatment of such steels at temperatures between 600 and 650 ° C is able to significantly increase the thermal stability compared to similar tempered, but conventional alloys. This is due to the onset of precipitation hardening by vanadium nitrides in this temperature range, which at 700 ° C for the first time by Göcmen, A. et al.: Precipitation Behavior and Stability of Nitrides in High Nitrogen Martensitic 9% and 12% Chromium Steels, ISIJ Int., 1996, 36, p. 769, was observed. Of importance is that in this case fine and dense precipitation states with high coherence of the vanadium nitrides to the iron grid were found. It can be concluded that secondary hardening with vanadium nitrides at 600 to 650 ° C offers no special ductility advantages compared to classical secondary hardening at 450 to 530 ° C.

Another important role in the pamphlet EP 0 866 145 A2 listed alloys according to the invention is played by the element manganese. The element manganese, especially in high-nitrogen steels, has an important role to play because it increases the solubility for nitrogen in the melt as well as in the austenite matrix. Manganese also has the property of shifting the austenite-ferrite conversion nose to longer times. These properties of manganese give favorable conditions for the vanadium nitrides to be separated again after a solution annealing treatment, before the martensitic transformation in the area of the metastable austenite. On the side of the carbon-containing 12% chromium steels, on the other hand, manganese is understood as an impurity element which promotes temper embrittlement considerably. Therefore, the manganese content is usually limited to very small amounts, especially with regard to applications in the temperature range between 350 and 500 ° C.

Substitution of manganese by nickel in high-nitrogen containing 9-13% chromium steels creates new advantages and possibilities. It can be assumed that nickel as a substitution element in solid solution improves the ductility of the crystal matrix. Alloying with nickel further lowers the Ac1 and Ac3 temperatures. This provides the advantage in N- and V-alloyed systems that vanadium nitrides can be precipitated at low austenitizing temperatures, ie in an austenitic matrix. The key advantage, however, is that the vanadium nitrides, which are inherently difficult to germinate, can easily be germinated in the martensitic, dislocation-rich matrix before being converted into the austenitic matrix. Thus, if the intention is to finely disperse the vanadium nitrides in the austenitic matrix, then it is no longer necessary to carry out the removal directly after the solution annealing treatment in the "metastable" austenite, as described in US Pat EP 0 866 145 A2 has been described. Since the nuclei for the vanadium nitrides can now be easily formed in the martensitic matrix, the aging time for the maturation of the nitrides in austenite can be significantly shortened. By alloying with nickel, a new possibility is thus offered to precipitate the vanadium nitrides rapidly and effectively in the transformable austenitic matrix. Since the austenite hardening can now be carried out effectively without manganese, the stability of the martensitic matrix against temper embrittlement can be further improved by a strong limitation of manganese.

It is further to be considered that the alloy provides sufficiently high solubility for the preferred levels of nitrogen. It is known that manganese increases the solubility for nitrogen and nickel lowers the solubility for nitrogen. The particular advantage of the desired alloy design lies in the fact that the required solubility for nitrogen is already provided by the element vanadium, which is used to form an optimal microstructure in an almost stoichiometric ratio to nitrogen is added. The dominant effect of vanadium on the solubility of nitrogen makes it possible for the increased and preferred levels of nitrogen to be introduced by vanadium and nearly stoichiometrically to vanadium without application of overpressure and, therefore, only inferiorly hampered by the presence of nickel and cobalt.

The element cobalt also offers the possibility of retarding the aging of the nitrides and the recovery of the dislocations during tempering without provoking an increased austenitic reversion during tempering.

The preferred amounts in weight percent for each element and the reasons for the selected alloy regions according to the invention in connection with the resulting possibilities of heat treatments are shown below.

Chrome:

A weight proportion of 9-13% 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 13% 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 together with silicon promote temper embrittlement and must therefore be limited to very small contents. The area to be specified should take into account the possibilities of ladle metallurgy ranging between 0.001 and 0.25% for manganese and between 0.001 and 0.15% 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 up to about 3.5% by weight remain homogeneously dissolved in the matrix if the tempering temperature or the stress relief annealing temperature does not exceed 600 ° C. to complete the entire tempering treatment. For alloys which are to be annealed at low temperatures, ie at 600 to 640 ° C, a preferred nickel content is from 3 to 4 wt%. Nickel contents above 4% by weight enhance 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. In the presence of stoichiometric amounts of nitrogen and vanadium, however, a special heat treatment is suitable for the high nickel steels. If such an alloy is solution-annealed at high temperatures, for example at 1150 to 1200 ° C., then an increased residual austenite content after tempering is due to the effect of the high concentration of nitrogen and vanadium in solution due to the resulting increase in the martensite start temperature. However, re-reactitizing at temperatures between 700 and 850 ° C allows re-precipitation of vanadium nitrides, which can again raise the martensite start temperature such that complete re-conversion into the martensite by quenching becomes possible again. Such low readenitization temperatures prevent premature over-aging of the vanadium nitrides, so that they still provide a substantial particle hardening contribution. This process can result in a martensite that is well stabilized with vanadium nitrides, which, due to the previous grain remodeling process, sets a particularly high level of martensite Ductility allows. A further tempering treatment at about 600 ° C leads to the formation of small austenite islands, which is sufficiently stabilized against a back transformation into the martensite. The volume fraction of this austenite is below 5%, provided that the nickel content does not exceed 7%. Even higher volume fractions increase the risk of embrittlement during long-term aging at elevated temperatures. This type of heat treatment is suitable for alloys with 2 to 7% nickel. A particularly good combination of heat resistance and ductility is achieved in consideration of this special heat treatment technology with nickel contents in the range between 4.5 and 6.5%.

Cobalt:

This element is used as a substitution element for iron in solid solution for the final fine tuning of ductility and heat resistance. An amount by weight of up to 10% cobalt can be alloyed without expectation of austenite conversion at tempering temperatures in the range of 600 to 650 ° C. The optimum content of cobalt depends on the proportion of molybdenum and tungsten. Above about 8% by weight, the alloying of cobalt proves uneconomical. A preferred range of alloys accounting for the high alloying cost of cobalt is 3.5 to 4.5 weight percent.

Molybdenum and tungsten:

Both elements improve creep resistance by solid solution hardening as partially dissolved elements and precipitation hardening during long-term stress. However, an excessively high proportion of these elements leads to embrittlement during long-term aging, which is due to the precipitation and coarsening of Laves phase (W, Mo) and Sigma phase (Mo). For this reason, the total Mo + W must be limited to 4%. An ideal range for W + Mo is between 1 and 4%.

Molybdenum is preferred because of its higher solubility to tungsten. A preferred range is given by a molybdenum content in the range of 1 to 2% and a tungsten content of less than 1%. Better is a molybdenum content of 1 to 2.5% and a tungsten content of less than 0.5%. A particularly preferred range is given by a negligibly small tungsten content, but molybdenum contents of 1 to 3%.

Vanadium and nitrogen:

Together, these two elements significantly control grain size formation and precipitation hardening. The microstructural forms are optimal when the elements vanadium and nitrogen are alloyed in a nearly stoichiometric ratio. The ideal weight ratio V / N is 3.6. Since the nitrogen solubility is improved by vanadium, a slightly more than stoichiometric V / N ratio is desirable. A slightly more than stoichiometric ratio sometimes also increases the stability of the vanadium nitride over the chromium nitride. Overall, a V / N ratio in the range between 3.5 to 4.2 is preferred. A particularly preferred range is 3.8 to 4.2. 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 amount of vanadium and nitrogen, the greater the proportion of vanadium nitrides which no longer dissolve 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 0.13 to 0.18 wt .-% and that of vanadium is in the range between 0.5 and 0.8 wt .-%.

Titanium:

Titanium nitride is a sparingly soluble nitride that aids grain refinement. In contrast to vanadium nitride, however, it can form in the melting phase and in particular in the solidification phase, with the result that the solidification takes place more calmly and more finely. Too high proportions by weight, however, lead to very large primary nitrides, which deteriorate the ductility. Therefore, the upper titanium content must be limited to 0.05%.

Niobium, tantalum, zirconium and hafnium:

These are all strong nitride formers that support 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%. A particularly preferred nitride former is niobium, since niobium dissolves in vanadium nitride in small amounts and can thus improve the stability of the vanadium nitride. Preferably, niobium is added in the range between 0.01 and 0.07%.

Phosphorus, sulfur, arsenic, antimony and 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 the minimum tolerable levels.

Aluminum:

This element is a strong nitride former, which already sets nitrogen in the melt and thus strongly affects the effectiveness of the added nitrogen. The aluminum nitrides formed in the melt are very coarse and reduce the ductility. Aluminum must therefore be limited to a weight fraction of 0.01%.

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 0.1%. Another disadvantage is the fact that carbon increases the hardening during welding. The particularly preferred carbon content is in the range between 0.02 and 0.07% by weight.

Short description of the drawing

Fig. 1

eine graphische Darstellung, bei der die Streckgrenze ausgewählter Legierungen bei Raumtemperatur in Abhängigkeit der Übergangstemperatur des Bruchaussehens (FATT) aufgetragen ist und die Wirkung des Nickels und der Wärmebehandlungstemperaturen auf die Streckgrenze und auf FATT entnehmbar ist;

Fig. 2

eine graphische Darstellung, bei der die Streckgrenze ausgewählter Legierungen bei Raumtemperatur in Abhängigkeit der Übergangstemperatur des Bruchaussehens (FATT) aufgetragen ist und die Wirkung des Nickels und des Kobalts auf die Streckgrenze und auf FATT entnehmbar ist;

Fig. 3

eine graphische Darstellung, bei der die Streckgrenze ausgewählter Legierungen bei einer Prüftemperatur von 550 °C gegenüber der Übergangstemperatur des Bruchaussehens (FATT) aufgetragen ist und die Wirkung des Nickels und des Kobalts auf die Streckgrenze bei 550 °C und auf FATT entnehmbar ist;

Fig. 4

eine graphische Darstellung, bei der die Streckgrenze der erfindungsgemässen Legierung "alloy D" aus verschiedenen Wärmebehandlungen resultierend zusammen mit den Vergleichslegierungen X12CrNiMo12 (martensitisch-härtbarer Stahl) und IN706 (ausscheidungshärtbare Ni-Fe Legierung) und den dazugehörigen Kerbschlagarbeiten Av gegenüber der Prüftemperatur aufgetragen ist

Fig. 5

eine graphische Darstellung, bei der die Streckgrenze der erfindungsgemässen Legierung "alloy E" aus verschiedenen Wärmebehandlungen resultierend zusammen mit den Vergleichslegierungen X12CrNiMo12 (martensitisch-härtbarer Stahl) und IN706 (ausscheidungshärtbare Ni-Fe Legierung) und den dazugehörigen Kerbschlagarbeiten Av gegenüber der Prüftemperatur aufgetragen ist

In the drawing, several embodiments of the invention are shown. Show it:

Fig. 1

a graph in which the yield strength of selected alloys at room temperature as a function of the transition temperature of the fracture appearance (FATT) is plotted and the effect of the nickel and the heat treatment temperatures on the yield strength and on FATT is removable;

Fig. 2

a graph in which the yield strength of selected alloys at room temperature as a function of the transition temperature of the fracture appearance (FATT) is plotted and the effect of the nickel and the cobalt on the yield strength and FATT is removable;

Fig. 3

a graph in which the yield strength of selected alloys at a test temperature of 550 ° C against the transition temperature of the fracture appearance (FATT) is plotted and the effect of nickel and cobalt on the yield strength at 550 ° C and on FATT is removable;

Fig. 4

a graph in which the yield strength of the inventive alloy "alloy D" from various heat treatments resulting together with the comparative alloys X12CrNiMo12 (martensitic-hardenable steel) and IN706 (precipitation-hardenable Ni-Fe alloy) and the associated impact tests Av against the test temperature is plotted

Fig. 5

a graph in which the yield strength of the inventive alloy "alloy E" from various heat treatments resulting together with the reference alloys X12CrNiMo12 (martensitic-hardenable steel) and IN706 (precipitation-hardenable Ni-Fe alloy) and the associated impact tests Av against the test temperature is plotted

Ways to carry out the invention

The invention will be explained in more detail with reference to embodiments and FIGS. 1 to 5.

Table 1 represents a series of alloys according to the invention.

With the exception of the alloys AP35 and AP38, which were melted as 10 kg melt in an induction furnace, all other alloys were produced in the form of 60-80 kg electrodes by the electro-slag remelting process. Further, with the exception of alloys AP28M, no overpressure was used in setting the specified nitrogen content during melting or during the remelting process. These alloys were therefore melted at 0.9 bar (atmospheric pressure), or remelted. The resulting nitrogen analyzes (Table 1) show that the preferred nitrogen contents can be introduced during production even at high nickel contents (up to 5.5%) without overpressure.

The following heat treatments provide a framework for solution annealing and tempering (reaustenitizing temperature) within which the heat treatments were carried out:

W2:

Solution annealing at 1080 ° C / 2h / Air cooling to room temperature Tempering treatment at 640 ° C / 2h / Air cooling to room temperature Low stress annealing at 600 ° C / 1h

W4:

Solution annealing at 1180 ° C / 2h / Air cooling to room temperature Tempering treatment at 640 ° C / 2h / Air cooling to room temperature Low stress annealing at 600 ° C / 1h

T2C:

Solution annealing at 1180 ° C / 2h / air cooling to room temperature Reaustenitisierung at 750 ° C / 2h / air cooling to room temperature stress relief annealing at 600 ° C / 1h

In all other heat treatments, solution annealing and tempering temperatures were at best changed to still be between those of W2, W4, and T2C. For the heat treatment forged blocks were used, which had cross-sectional dimensions of 7x7 cm ² .

FIG. 1 shows for 3 different alloys according to the invention AP28M, "alloy D" and "alloy E", all of which are alloyed with 4% by weight of cobalt and otherwise differ mainly in nickel content, the adjustable combination of yield strength at room temperature and transition temperature of Breaking appearance (FATT). The results are compared with those of a commercial X12 CrNiMo12 alloy which has been solution annealed at 1060 ° C, annealed at 640 ° C and low stress annealed at 600 ° C. In principle, it can be seen that improved yield strength values and / or ductility can be achieved by the selected alloys. Crucial here is the finding that the achievable combination of yield strength and ductility can be sensitively influenced both by the content of nickel and by the solution annealing temperature. From Figure 1 it is clear that by increasing the solution annealing temperature, the yield strength can be improved effectively, but at the cost of equally reduced ductility. Furthermore, it is clear from Figure 1 that by increasing the nickel content, the ductility can be improved effectively, but at the expense of strength. The combination of both findings opens up new possibilities for adjusting alloys with improved strength and ductility properties by optimizing the nickel content on the one hand and optimizing the heat treatment on the other. Based on a simple heat treatment (W2: 1080 ° C / 640 ° C / 600 ° C) an optimal combination of heat resistance and ductility with nickel contents in the range between 3 and 3.5% is achieved. An exceptional combination of yield strength and ductility can be adjusted to 1180 and 750 ° C with two - stage austenitizing treatments, especially with high - nickel alloys ("alloy E"). This favorable combination of properties is made possible by the low Ac3 temperature of the high-nickel alloy. For example, in alloys with nickel contents of 5.5% by weight, the matrix at 750 ° C. is almost completely austenitic. This means that a large volume fraction of vanadium nitrides, which was dissolved at 1180 ° C, can be re-separated in the subsequent annealing at 750 ° C in austenite. The alloying with nickel appears to be well used to post-anneal and quench to room temperature at lower austenitizing temperatures before continuing the heat treatment with the classic quench and temper treatment.

FIG. 2 shows, by means of various alloys, the effect of cobalt on the combination of yield strength at room temperature and the transition temperature of the fracture appearance FATT. In this example, various heat treatments were tested with solution annealing temperatures between 1080 and 1200 ° C, tempering temperatures between 640 and 750 ° C and a final Stress relief annealing at 600 ° C. It can be clearly seen that all alloys with low cobalt contents in the lower quadrant are at low yield strength values and high FATT values, ie are significantly inferior to similar alloys with increased cobalt contents in terms of the adjustable combination of yield strength and ductility.

In FIG. 3, similar to the results shown in FIG. 2, the test results for the same alloys (same heat treatment) are shown at room temperature, taking into account the hot yielding strength at 550.degree. Alloys with cobalt offer, z. B. AP28M, a much better combination of the hot yield point at 550 ° C and the transition temperature the fracture appearance as alloys without cobalt, z. Eg "alloy A". Best properties are achieved with a two-stage solution annealing treatment (1180 ° C, 750 ° C).

FIGS. 4 and 5 show graphs in which the yield strength of the alloys "alloy D" and "alloy E" according to the invention is plotted against the test temperature as a function of the heat treatment carried out. Further, the yield strengths of the comparative alloys X12CrNiMo12 (matensitic-hardenable steel) and IN706 (precipitation-hardenable alloy), as well as their notch impact work Av are plotted in comparison to the inventive alloys. It can be seen that, for the alloys "alloy D" and "alloy E" according to the invention, an improvement in the heat resistance up to high test temperatures is maintained independently of the heat treatment and independently of the nickel content. The novel alloys according to the invention exhibit an extraordinarily good combination of notch impact work at room temperature and heat resistance at 550 ° C. in comparison with austenitic high-temperature alloys (IN706) which are designed on a nickel-iron base.

Of course, the invention is not limited to the described embodiments. Table 1 X12CrNiMo12 AP28M AP35 AP38 alloyA alloyB alloyC Alloyd alloyE Weight 80 kg 10 kg 10 kg 60 kg 60 kg 60 kg 60 kg 60 kg production ESU DESU ESU ESU ESU ESU ESU In% by weight Cr 11.5 11.9 10.2 10.1 10.7 10.8 10.4 11.2 11.2 Mn <0.25 0.04 0.18 0.21 0.1 0.07 0.07 0.05 0.04 Ni 2.3 2.6 4.4 3.4 2.5 1.8 2.4 3.1 5.6 Co 4 3.9 4.1 <0.02 1.0 1.0 4.0 4.0 Not a word 1.5 1.49 1.5 1.6 1.5 1.4 1.8 1.8 1.8 V 0.25 0.66 0.65 0.64 0.51 0.53 0.52 0.61 0.58 Nb 0.04 0.04 0.04 0.03 0.03 0.03 0.028 0.031 Ti 0.03 0.03 0.03 <0.02 <0.02 <0.02 <0.02 <0.02 Si <0.15 0.15 <0.05 <0.05 <0.02 <0.02 <0.02 <0.02 <0.02 C 0.12 0,036 0,042 0,039 0.058 0,065 0,065 0,042 0,035 N 0,035 0.18 0.15 0.15 0.16 0.16 0.15 0.156 0,159

Claims (17)

1. Martensitic-hardenable heat-treated steel, characterized by the following composition (data in % by weight): 9 to 13% Cr, 0.001 to 0.25% Mn, 2 to 7% Ni, 0.001 to 8% Co, at least one of W and Mo in total between 0.5 and 4%, 0.5 to 0.8% V, at least one of Nb, Ta, Zr and Hf in total between 0.001 and 0.1%, 0.001 to 0.05% Ti, 0.001 to 0.15% Si, 0.01 to 0.1% C, 0.12 to 0.18% N, at most 0.025% P, at most 0.015% S, at most 0.01% Al, at most 0.0012% Sb, at most 0.007% Sn, at most 0.012% As, remainder iron and customary impurities, and the proviso that the ratio by weight of vanadium to nitrogen V/N lies in the range between 3.5 and 4.2.
2. Martensitic-hardenable heat-treated steel according to Claim 1, characterized by 2 to 4.5% Ni.
3. Martensitic-hardenable heat-treated steel according to Claim 1, characterized by 2.7 to 3.7% Ni.
4. Martensitic-hardenable heat-treated steel according to Claim 1, characterized by 4 to 7% Ni.
5. Martensitic-hardenable heat-treated steel according to Claim 1, characterized by 4.5 to 6.5% Ni.
6. Martensitic-hardenable heat-treated steel according to one of Claims 1 to 5, characterized by 0.5 to 6% Co.
7. Martensitic-hardenable heat-treated steel according to Claim 6, characterized by 2 to 6% Co.
8. Martensitic-hardenable heat-treated steel according to Claim 6, characterized by 3.5 to 4.5% Co.
9. Martensitic-hardenable heat-treated steel according to one of Claims 1 to 8, characterized by 10 to 12% Cr.
10. Martensitic-hardenable heat-treated steel according to Claim 9, characterized by 10.5 to 11.5% Cr.
11. Martensitic-hardenable heat-treated steel according to one of Claims 1 to 10, characterized by 0.02 to 0.07% C.
12. Martensitic-hardenable heat-treated steel according to one of Claims 1 to 11, characterized by 0.5 to 0.7% V and 0.14 to 0.17% N.
13. Martensitic-hardenable heat-treated steel according to one of Claims 1 to 12, characterized by 0.01 to 0.07% Nb.
14. Martensitic-hardenable heat-treated steel according to one of Claims 1 to 13, characterized by a total of Mo and W which lies in the range between 1 and 4%.
15. Martensitic-hardenable heat-treated steel according to Claim 14, characterized by less than 1% W and a total of Mo and W which lies in the range between 1 and 2.5%.
16. Martensitic-hardenable heat-treated steel according to Claim 15, characterized by less than 0.5% W and a total of Mo and W which lies in the range between 1 and 2.5%.
17. Process for the heat treatment of a steel having a composition according to one of Claims 1 to 16, characterized by the following successive process steps:

    - Solution annealing at 1050 to 1250°C,

    - Cooling to a temperature below 300°C,

    - Tempering treatment, partial or complete reaustenitization at 600 to 900°C,

    - Cooling to a temperature below 300°C,

    - Annealing at a temperature of from 550 to 650°C.

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