EP0866145B1 - Procédé de traitement thermique pour alliage d'acier complètement martensitique - Google Patents

Procédé de traitement thermique pour alliage d'acier complètement martensitique Download PDF

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EP0866145B1
EP0866145B1 EP98810193A EP98810193A EP0866145B1 EP 0866145 B1 EP0866145 B1 EP 0866145B1 EP 98810193 A EP98810193 A EP 98810193A EP 98810193 A EP98810193 A EP 98810193A EP 0866145 B1 EP0866145 B1 EP 0866145B1
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alloy
temperature
heat treatment
weight
temperatures
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EP0866145A2 (fr
EP0866145A3 (fr
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Peter Dr. Ernst
Peter Prof. Dr. Uggowitzer
Markus Prof. Dr. Dr.H.C. Speidel
Alkan Göcmen
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General Electric Switzerland GmbH
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Alstom Schweiz AG
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    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/30—Ferrous alloys, e.g. steel alloys containing chromium with cobalt
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/002—Heat treatment of ferrous alloys containing Cr
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C33/00—Making ferrous alloys
    • C22C33/02—Making ferrous alloys by powder metallurgy
    • C22C33/0257—Making ferrous alloys by powder metallurgy characterised by the range of the alloying elements
    • C22C33/0278—Making ferrous alloys by powder metallurgy characterised by the range of the alloying elements with at least one alloying element having a minimum content above 5%
    • C22C33/0285—Making ferrous alloys by powder metallurgy characterised by the range of the alloying elements with at least one alloying element having a minimum content above 5% with Cr, Co, or Ni having a minimum content higher than 5%
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/001—Ferrous alloys, e.g. steel alloys containing N
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D1/00—General methods or devices for heat treatment, e.g. annealing, hardening, quenching or tempering
    • C21D1/18—Hardening; Quenching with or without subsequent tempering

Definitions

  • the invention relates to a heat treatment process for alloy specifications from the class of fully martensitic 9-15% chrome steels.
  • controlled excretion management in The quenching phase can have excellent properties and combinations of properties can be set for broad applications in the power plant sector.
  • the mechanical properties important for the use are indicated by a so-called remuneration process. It is done by solution heat treatment, a quenching treatment and subsequent tempering treatment in a medium temperature range.
  • the resulting microstructure is characterized by a dense arrangement of slats, which with excretion phases overgrown. These microstructures are at elevated temperatures unstable. They soften depending on the time, the stress and them imposed deformations.
  • the phase reactions taking place during the heat treatment limit the achievable toughness within the required strength.
  • the phase reactions taking place during operation together with the coarsening of the excretions cause an increased susceptibility to embrittlement and reduce the strains to be borne by the components.
  • the current alloys of the class are fully martensitic 9-15% Chromium steel no longer meets the requirements of modern power plant technology. This primarily concerns the combination of strength and toughness, and also also Combinations of high temperature strength, creep resistance, creep resistance, Resistance to relaxation, resistance to creep embrittlement and thermal Fatigue. A constant improvement in properties of this alloy class by demanding full through-payability, especially in thick-walled Components, narrow metallurgical limits.
  • an optimal combination of grain coarsening resistance, hardenability and tempering resistance is achieved through a suitable (empirical) adjustment to vanadium, niobium, carbon and nitrogen.
  • Optimal combinations are achieved when the proportion of carbon in atomic percentages is higher than that of nitrogen.
  • the optimum carbon content is in the range 0.1-0.2% by weight and the optimum nitrogen content is in the range 0.05-0.1% by weight.
  • nitrogen is added in almost stoichiometric proportions to the special nitride formers vanadium or niobium.
  • the optimum content of vanadium is consequently in the range 0.2-0.35% by weight and that of niobium in the range 0.05-0.4% by weight.
  • the prior art is well represented by the older alloys X22CrMoV121 (X22), X20CrMoV121, X12 CrNiMo2, X19CrMoVNbN111 (X19), and by the newer alloys X10CrMoVNbN91 (P / T91), X12CrMoWVNbN1011 (rotor steel E2), X18CrMoVNbNB91 (rotor steel B2) and through the alloy X20CrMoVNbNB10 1 (TAF). From Goecmen A.
  • the invention has for its object a heat treatment process for alloy specifications for training to identify fully martensitic structures in which a controlled dissolution and re-excretion of special nitrides or special carbonitrides together with the martensitic phase transition to the highest properties and property combinations leads without the properties and combinations of properties to be achieved, are limited by the size of the components to be coated.
  • These specifications, which stand out in terms of composition and heat treatment are then not only used in the area of thin-walled components such as tubes, bolts and blades, but also for rotors, rotor disks, various housing components, boiler systems and much more.
  • the microstructure resulting after the tempering treatment is characterized by a very uniform and dense dispersion of special nitrides and or special carbonitrides in a lath structure, which already were eliminated before the martensitic phase transition.
  • the identified Alloy compositions not only offer an optimal combination of grain coarsening resistance, hardenability and temper resistance, but also enable a targeted influence on the martensitic phase transition through elimination phases for the purpose of improved mechanical Properties and increased structural stability in operation.
  • compositions in which these phase reactions for adjustment increased properties and property combinations can be used can essentially contain 8 to 15% Cr, up to 15% Co, up to 4% Mn, up to 4% Ni, up to 8% Mo, up to 6% W, 0.5 to 1.5% V, up to 0.15% Nb, up to 0.04% Ti, up to 0.4% Ta, to 0.02% Zr, up to 0.02% Hf, up to 0.1% C and 0.12-0.25% N, remainder iron and usual melting-related Impurities.
  • the associated heat treatments, which enable a controlled setting of improved property combinations are characterized as follows. Solution heat treatment is preferably carried out between 1150 and 1250 ° C with holding times between 0.5 and 15h.
  • the cooling takes place below a temperature of 900 ° C with cooling rates less than 120 ° C / h and is carried out as required and by an isothermal annealing in the temperature range between 900 and 500 ° C is interrupted. Cooling and isothermal annealing can be done as needed and by one thermomechanical treatment. The tempering treatment after quenching occurs in the temperature range between 600 and 820 ° C and can be between 0.5 to 25h.
  • the invention has a number of advantages.
  • the specifications formulated above the alloy composition and the heat treatment allow the Adjustability of the highest possible combinations of properties of strength, toughness, High temperature strength, relaxation resistance, creep resistance, creep resistance, Creep ductility, resistance to thermal fatigue and so on.
  • the simple controllability of the resulting excretion states enables one economically efficient development and improvement of products for High temperature applications.
  • the structure aging takes place in the company delayed by the uniformity and stability of the excretion states and controls and thus not only enables extended downtimes, but increases them also the reliability of service life forecasts of the components in operation.
  • the Microstructure formation in thick-walled components, such as in rotors can by influencing and controlling the local cooling speeds Be designed to meet demands, be flexible and optimal. This allows one Considerably improved overall service life optimization of such components taking into account the thermal stresses occurring in them be non-uniform Operating conditions.
  • the specifications developed for the use according to the invention contain in essentially 8 to 15% Cr, up to 15% Co, up to 4% Mn, up to 4% Ni, up to 8% Mo, up to 6% W, 0.5 to 1.5% V, up to 0.15% Nb, up to 0.04% Ti, up to 0.4% Ta, up to 0.04% Zr, up to 0.04% Hf, up to 0.1% C and 0.12-0.25% N and can by casting or on powder metallurgical Be made way. Specifications of this kind are dependent the intended use of targeted dissolution and rejection reactions thermodynamically stable special nitrides and special carbonitrides at high temperatures and benefit from the martensitic phase transition. This becomes the overall stability of the structure that ripens during tempering and operation increased and the mechanical properties improved overall.
  • the amount of excretion can be influenced and improve the stability of the particles against coarsening.
  • the one from the Structures resulting from forging treatment are characterized by the stabilizing effect of primary nitrides very resistant to grain coarsening and allow therefore a controlled partial re-dissolution of primary nitrides during the Solution heat treatment.
  • nitride dispersion Particle size 3-50nm and particle distances between 5 and 100nm become. These influence the morphology and the dislocation density of the resulting Martensite.
  • the uncontrolled formation of coarse grain boundary deposits or the formation of grain boundary films are determined by the type and the kinetics of formation suppresses this special nitride. Bainite conversion is used in such nitrogen and vandium-rich systems were not observed.
  • Chromium promotes corrosion resistance and through-temperability Element.
  • its ferrite-stabilizing effect must be reduced by the austenite-stabilizing one Effect of other elements, such as Co, Mn or Ni can be compensated.
  • This lower disadvantageously for the emergence of a fully martensitic remuneration structure both the martensite start temperature and the ferrite stability the tempering treatment or, as in the case of Co, increase the alloying costs.
  • Cr should not exceed 15% by weight.
  • Less than 8% chromium in turn, not only reduce corrosion and oxidation resistance to an intolerable level Level but also affect hardenability in a way that a flexible precipitation of special nitrides before the martensitic phase change is severely impaired.
  • a particularly preferred range is 10 up to 14% chromium, in particular 11 to 13% chromium.
  • Manganese is an element that promotes through-payability and is for one flexible separation of special nitrides before the martensitic phase transition very important. However, 4% by weight is sufficient for these purposes. Furthermore Mn lowers the martensite start temperature and the ferrite stability at Annealing treatment, leading to undesirable structural training in the fully paid Condition leads. Particularly preferred ranges are up to 2.5%, 0.5 to 2.5% and 0.5 to 1.5% manganese.
  • Nickel like Mn, is an element that promotes through-hardening, but its effect in this regard is not as strong as that of manganese. On the other hand, its effect on austenite stability at high solution annealing temperatures is significantly stronger than that of manganese. Furthermore, its lowering effect on the martensite starting temperature and the ferrite stability when starting is not as high as that of the manganese.
  • a substitution of Ni by Mn depends on the flexibility of the precipitation reactions to be carried out before the martensitic phase transition and on the level of the A c1 temperature required for optimal structure formation in the tempered state. However, the nickel content should not exceed 4% by weight, otherwise the A c1 falls to insufficiently low values. Particularly preferred ranges are up to 2.5%, 0.3 to 2.5%, 0.5 to 2.5%, up to 2% and up to 1.5% nickel.
  • Ni + Mn should not be more than 4% by weight.
  • Cobalt is the most important element for the optimization of a high austenite stability at high solution annealing temperatures and a high A c1 temperature. Its proportion depends on the amount of the ferrite-stabilizing elements Mo, W, V, Nb, Ta, Ti, Zr and Hf that are important for strength. Above 15% by weight, the Ac1 temperature drops to intolerably low values for a fully tempered one Structure from. Preferred ranges are 5 to 15% by weight, 3 to 15% by weight, 1 to 10% by weight, 3 to 10% by weight, 1 to 8% by weight, 3 to 7% by weight. % and 1 to 6% by weight.
  • a particularly preferred range is 5-15% by weight cobalt for alloys which have a very high strength potential due to high molybdenum and tungsten contents, and 1-10% by weight cobalt for alloys with a low to medium strength level. Small strength levels are approximately 700 to 850 MPa, medium levels 850 to 1100 MPa and high levels over 1100MPa.
  • Molybdenum can perform many functions that are important for microstructure formation. Chromium and manganese alike have a strong support in terms of through-hardenability Effect. Furthermore, it can be in solution or via excretion reactions contribute significantly to a further increase in strength. High molybdenum content lower, however, through the rapid coarsening of the intermetallics that form them Elimination phases the toughness. Its ideal salary is based on the intended applications and operating temperatures of the corresponding components. However, molybdenum contents above 8% by weight reduce the toughness and the Martensite start temperature to intolerable values. Preferred molybdenum contents are below 5% by weight, in particular below 4 and 3% by weight.
  • Tungsten works in a similar way to molybdenum and the molybdenum content should be kept below 6% by weight. Its ideal content, like molybdenum, depends on the application and the operating temperature of the corresponding components. Preferred tungsten grades are below 4% by weight, in particular below 3% by weight.
  • the Sum of Mo + W should not be more than 8% by weight.
  • Vandium is that when it comes to setting the highest combinations of properties such as strength and toughness, creep rupture strength and creep ductility as well as structural stability most important alloy element. It ensures together with nitrogen high resistance to grain coarsening at high solution annealing temperatures and a strength-enhancing high separation volume of UN special nitrides at lower excretion temperatures. For a sufficiently high Combination of a high coarsening resistance with an effective one Elimination volumes are, however, necessary at least 0.5% by weight. Increased levels of vanadium make increased solution glow temperatures necessary. At Vandium contents above 1.5% by weight, the solution annealing temperature to be applied increases for increased strength on technically no longer realizable Values. A preferred range is 0.5 to 1 wt% vanadium. A special one preferred range is 0.5 to 0.8 wt .-% vanadium.
  • Vandium's nitrogen is the accompanying element for the formation of MN special nitrides. For a sufficiently good combination of a high coarsening resistance with a strength-effective separation volume at least 0.12% by weight is necessary.
  • the amount to be applied rises like vanadium Solution annealing temperature for improved properties at nitrogen levels above 0.25% by weight on values that are no longer technically feasible.
  • a preferred one The range is 0.12 - 0.2% by weight nitrogen.
  • a particularly preferred area is 0.12 - 0.18% by weight nitrogen.
  • carbon can contribute to an increased excretion volume of special carbonitrides without reducing the coarsening resistance. Excess carbon increases the hardness of the quenched martensite. However, it promotes the formation of toughness-reducing phases of precipitation such as M 23 C 6 and M 2 (C, N), as well as the formation of bainite at low cooling rates.
  • the carbon content should therefore not exceed 0.1% by weight.
  • a preferred range is less than 0.05 wt% C.
  • a particularly preferred range is less than 0.03 wt% C.
  • Nb this is 0.15% by weight, for Ta 0.4% by weight, for Ti 0.04% by weight and for the elements Hf and Zr 0.02% by weight each. These elements can be used alone or in Combination with each other effectively contribute to property improvements.
  • the optimal composition depends on the mechanical properties to be set from.
  • niobium is the preferred element among the special nitride formers.
  • Preferred maximum niobium contents are below 0.1% by weight.
  • Very preferred niobium Contents are 0.02 to 0.1% by weight.
  • Boron is an element that promotes through-hardening and therefore for flexible excretion reactions in austenite before the martensitic phase transformation practical. It also increases the coarsening resistance of excretions in tempered martensite. Because it tends to segregate and high Affinity to nitrogen shows, the boron content must be limited to 0.005 wt .-%.
  • Silicon is an important deoxidation element and is therefore always found in steel. In solution, it can contribute to the strength of the steel and, at the same time, that Increase resistance to oxidation. In large quantities, however, it appears embrittled. The proportion by weight of silicon should therefore not exceed 0.3% by weight.
  • the alloy specifications mentioned ensure a fully martensitic Event structure, which is generated by an extended remuneration process becomes.
  • This consists of a solution heat treatment, a controlled rapid one or slow cooling treatment, with or without one of the martensitic phase transformation previous thermomechanical treatment or isothermal Annealing, and one after quenching to room temperature Tempering treatment.
  • Solution heat treatment is carried out at temperatures between 1150 ° C and 1250 ° C with holding times between 0.5 and 15h.
  • the purpose of this solution heat treatment is the partial dissolution of special nitrides and special carbonitrides.
  • a special one delayed cooling or iosthermal annealing with or without thermomechanical Treatment, i.e. Deformation, in the quenching phase occurs at temperatures between 900 and 500 ° C and can take the entire quenching treatment up to 1000h delay.
  • the occasion treatment takes place at temperatures between 600 and 820 ° C and takes place during glow times between 0.5 and 25h. This intends a partial recovery from the martensitic Phase transformation generated internal tensions.
  • the mean grain diameter of the in the steel alloy by the solution annealing treatment developing structure does not grow beyond a value of 50 ⁇ m.
  • the subsequent cooling down to the martensite start temperature the controlled elimination of vanadium-rich special nitrides or special carbonitrides, be it through a thermomechanical Treatment or be it through an artificially delayed cooling.
  • the AP alloys were at a nitrogen partial pressure of 0.9bar at temperatures melted between 1500 and 1600 ° C.
  • the cast blocks were between Forged at 1230 and 1050 ° C.
  • the heat treatments were forged Plates made with a thickness of 15mm.
  • Solution heat treatment was carried out during the heat treatments for the mechanical tests at 1180 ° C and lasted 1h. Subsequently, an oven was checked Cooling is carried out at a cooling rate of 120 ° C / h.
  • Separate Heat treatments are characterized by isothermal austenite aging (Ausageing) out. The sample is heated to a medium temperature after solution annealing cooled, which is clearly above the martensite start temperature; then kept at this temperature for a certain time and then on Cooled to room temperature. Such a heat treatment is shown schematically in FIG 1 reproduced.
  • the heat treatments T2 and T6 differ from the heat treatment T5 is characterized by very high cooling rates in the quenching phase. In the Heat treatment T5 will also result in a longer isothermal annealing before the martensitic Phase conversion carried out.
  • Fig. 1 shows schematically the time-temperature history of the heat treatment T5.
  • Figure 2 shows the grain sizes resulting from the use of different solution annealing temperatures result.
  • the grain size increases with increasing Solution annealing temperature.
  • a solution annealing temperature 1100 ° C a very pronounced grain coarsening on.
  • the accelerated alloys AP1, AP8 and AP14 are used Grain coarsening only above 1200 ° C.
  • Figure 3 shows for the alloy AP11 how an isothermal Annealing after solution annealing and before the martensitic phase transition affects the hardness of the quenched martensite.
  • the individual samples were each with different austenite aging temperatures and times removed from the oven and quenched in water.
  • the hardness at the time zero corresponds to the martensite hardness in the absence of austenite aging So the solution annealed (1200 ° C / 1h) and directly quenched state.
  • Austenite aging changes the quenching hardness depending on the aging temperature and the aging time before the martensitic phase transition.
  • the hardness course cannot be monotonous. Basically achieved higher quenching hardness at low austenite aging temperatures than at high ones Ausageing.
  • Figure 3 shows, however, that austenite aging treatment are adequately controlled in this way for the purpose of new structural states can that no major losses in hardness are expected.
  • FIG. 4 shows the tempering curves of three alloys examined (AP1, AP8, AP14) in comparison with the known alloy X20 CrMoV 121.
  • inventive Alloys at tempering temperatures above 600 ° C higher tempering temperatures achieved, and this with the same molybdenum content in the alloy (Comparison of AP14 with TAF in Table 1). The influence of molybdenum only becomes very high levels significant (AP8).
  • FIG. 5 shows the influence of a previous austenite aging on the temper resistance of the AP11 alloy.
  • Austenite aging relates to structural conditions, which a minor after austenite aging Martensite hardness than the solution annealed and directly quenched state. It can be seen, however, that the differences to those important for the technology Reduce tempering temperatures above 600 ° C. There are even states (aged austenite: 600 ° C / 150h), which has a higher hardness at a tempering temperature of 650 ° C exhibit. Austenite aging can thus be used for setting higher strengths use.
  • Figure 6 shows the influence of austenite aging on the impact energy and the Notch impact energy transition temperature for the alloy AP1.
  • the transition temperature of the impact work decreases with increasing Tempering temperature and therefore allows the setting of higher impact energy.
  • an austenite overaging does not lead to any significant embrittlement.
  • Figure 7 shows the influence of austenite aging on the yield strength at test temperatures between 23 ° C and 600 ° C. Basically, the yield strength increases with decreasing Tempering temperature. This means achieving high strengths accordingly Figure 6 is at the expense of a significantly reduced impact energy. On the other hand leads to an austenite aging of the AP1 alloy a significant increase in the yield strength up to a temperature of approximately 550 ° C without being linked to embrittlement.
  • Figure 8 shows a comparison of the yield strengths between the Alloy AP1 and known (X20CrMoV121, X12 CrNiMo 12) or in technology Newly-introduced alloy (X12 CrMoWVNbN11 11), being indicated at the Comparative values are minimal standard values. The comparison shows that at similar tempering temperatures for the exemplary alloy AP1 clearly higher yield strengths result.
  • Figure 10 shows the influence of austenite aging on the impact energy and transition temperature the impact energy for an alloy AP8. This is characterized by a high molybdenum content (Table 1). This means that extremely high tempering resistance can be achieved even above one Achieve tempering temperature of 600 ° C ( Figure 4). On the other hand, this is with the The disadvantage of pronounced embrittlement is linked. An increase in the tempering temperature from 710 to 740 ° C is not very effective here. However, for this Alloy the transition temperature of the impact work by a previous one Austenite aging even while maintaining a tempering temperature of 710 ° C considerably lower.
  • FIG. 11 shows the influence of an austenite aging on the yield strength between 23 ° C. and 650 ° C. for the same alloy AP8.
  • austenite aging does not increase the yield strength at room temperature, but austenite aging at lower austenite aging temperatures leads to a significant increase in hot yield strength at temperatures above 500 ° C.

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Claims (3)

  1. Procédé de traitement thermique d'alliages d'acier homogène essentiellement constitués de (en % en poids) 8 à 15 % de Cr, jusque 15 % de Co, jusque 4 % de Mn, jusque 4 % de Ni, jusque 8 % de Mo, jusque 6 % de W, de 0,5 à 1,5 % de V, jusque 0,15 % de Nb, jusque 0,04 % de Ti, jusque 0,4 % de Ta, jusque 0,02 % de Zr, jusque 0,02 % de Hf, au plus 50 ppm de B, jusque 0,1 % de C et de 0,12 à 0,25 % de N, la teneur en Mn + Ni étant inférieure à 4 % et la teneur en Mo + W inférieure à 8 %, le solde étant constitué de fer et des impuretés habituelles de fusion,
    caractérisé en ce que l'alliage subit un recuit de solubilisation à des températures comprises entre 1150°C et 1250°C avec des temps de maintien compris entre 0,5 et 15h, en ce qu'après le recuit de solubilisation, l'alliage est refroidi en dessous d'une température de 900°C à des vitesses de refroidissement inférieures à 120°C/h, en ce que l'alliage est refroidi à température ambiante et ensuite subit un revenu à des températures comprises entre 600°C et 820°C pendant 0,5 à 25h.
  2. Procédé de traitement thermique selon la revendication 1, caractérisé en ce que directement après le traitement de recuit de solubilisation, l'alliage subit en dessous d'une température de 900°C un ou plusieurs recuits isothermes à une température ou à des températures différentes, entre 5 et 500h.
  3. Procédé de traitement thermique selon la revendication 1 ou 2, caractérisé en ce que le traitement thermique après le recuit de solubilisation est associé à une déformation.
EP98810193A 1997-03-21 1998-03-09 Procédé de traitement thermique pour alliage d'acier complètement martensitique Expired - Lifetime EP0866145B1 (fr)

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DE19712020 1997-03-21
DE19712020A DE19712020A1 (de) 1997-03-21 1997-03-21 Vollmartensitische Stahllegierung

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EP0866145A3 EP0866145A3 (fr) 1998-12-23
EP0866145B1 true EP0866145B1 (fr) 2003-05-07

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US (1) US6030469A (fr)
EP (1) EP0866145B1 (fr)
JP (1) JP4204089B2 (fr)
AT (1) ATE239804T1 (fr)
DE (2) DE19712020A1 (fr)
ES (1) ES2199414T3 (fr)

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US7985304B2 (en) 2007-04-19 2011-07-26 Ati Properties, Inc. Nickel-base alloys and articles made therefrom
US10563293B2 (en) 2015-12-07 2020-02-18 Ati Properties Llc Methods for processing nickel-base alloys

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US7235212B2 (en) * 2001-02-09 2007-06-26 Ques Tek Innovations, Llc Nanocarbide precipitation strengthened ultrahigh strength, corrosion resistant, structural steels and method of making said steels
NZ516393A (en) 1999-07-12 2003-01-31 Mmfx Steel Corp Of America Low-carbon steels of enhanced mechanical and corrosion properties with heating and cooling to achieve laths of martensite alternating with films of retained austenite, and no carbides
JP3492969B2 (ja) * 2000-03-07 2004-02-03 株式会社日立製作所 蒸気タービン用ロータシャフト
DE10025808A1 (de) * 2000-05-24 2001-11-29 Alstom Power Nv Martensitisch-härtbarer Vergütungsstahl mit verbesserter Warmfestigkeit und Duktilität
CA2361180A1 (fr) * 2000-11-03 2002-05-03 Michael M. Antony Articles en acier inoxydable resistant a la fatigue thermique
DE10063117A1 (de) * 2000-12-18 2003-06-18 Alstom Switzerland Ltd Umwandlungskontrollierter Nitrid-ausscheidungshärtender Vergütungsstahl
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EP0866145A2 (fr) 1998-09-23
JPH10265914A (ja) 1998-10-06
DE59808216D1 (de) 2003-06-12
US6030469A (en) 2000-02-29
DE19712020A1 (de) 1998-09-24
JP4204089B2 (ja) 2009-01-07
ATE239804T1 (de) 2003-05-15
EP0866145A3 (fr) 1998-12-23

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