EP2116626B1 - Kriechfester Stahl - Google Patents

Kriechfester Stahl Download PDF

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Publication number
EP2116626B1
EP2116626B1 EP09151605A EP09151605A EP2116626B1 EP 2116626 B1 EP2116626 B1 EP 2116626B1 EP 09151605 A EP09151605 A EP 09151605A EP 09151605 A EP09151605 A EP 09151605A EP 2116626 B1 EP2116626 B1 EP 2116626B1
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EP
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Prior art keywords
creep
resistant steel
ppm
steel according
maximum
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German (de)
English (en)
French (fr)
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EP2116626A1 (de
Inventor
Mohamed Nazmy
Andreas KÜNZLER
Markus Staubli
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General Electric Technology GmbH
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Alstom Technology 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
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/44—Ferrous alloys, e.g. steel alloys containing chromium with nickel with molybdenum or tungsten

Definitions

  • the invention relates to steels based on 8-13% chromium, which are used for rotors in the power plant sector. It concerns the selection and proportionate tuning of special alloying elements which allow the setting of exceptionally good creep resistance at temperatures of 550 ° C and above in this material.
  • the steel according to the invention should also have a good low cycle fatigue (LCF) resistance and good toughness after long-term aging so that it can be used in both gas and steam turbines.
  • LCF low cycle fatigue
  • Maragingitic-hardening steels based on 9-12% chromium are widely used materials in power plant technology. They were developed for use in steam power plants at operating temperatures above 600 ° C and steam pressures above 250 bar to increase the efficiency of power plants. Under these operating conditions, the creep resistance and the oxidation resistance of the material play a special role.
  • chromium in the abovementioned range not only provides good resistance to atmospheric corrosion but also complete through-hardenability of thick-walled forgings, for example as monobloc rotors or as rotor disks in gas and steam turbines.
  • Proven alloys of this type usually contain about 0.08 to 0.2% carbon, which in solution allows the setting of a hard martensitic structure.
  • a good combination of heat resistance and ductility of martensitic steels is made possible by a tempering treatment in which the precipitation of carbon in the form of carbides with simultaneous recovery of the dislocation substructure forms a particle-stabilized subgrain structure.
  • the tempering behavior and the resulting properties can be effectively influenced by the choice and proportionate tuning of specific carbide formers such as Mo, W, V, Nb and Ta.
  • German steel X20CrMoV12.1 known under DIN.
  • the contents of Cr, Mo, W were optimized taking into account N, Nb and / or B to improve creep and creep rupture strengths for 600 ° C applications.
  • the carbides such as M 23 C 6 .
  • the Ni contents were limited to values of less than 0.25% in these steels.
  • the fracture toughness values are disadvantageous, which does not play a major role in steam turbine applications and can therefore be neglected, but must be avoided in gas turbine applications.
  • 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.
  • DE 198 32 430 A1 is a further optimization of a similar X12CrNiMo12 steel disclosed in which by the addition of rare earth elements or boron embrittlement tendency in the temperature range 425-500 ° C is limited.
  • EP 0 866 145 A2 describes a new class of martensitic chromium steels with nitrogen contents in the range of 0.12 to 0.25% and in EP 1 158 067 A1 with nitrogen contents of 0.12 to 0.18%, wherein the weight ratio V / N is in the range between 3.5 and 4.2.
  • the entire structure of the structure is controlled by the formation of special nitrides, in particular vanadium nitrides, which can be distributed in a variety of ways by forging, austenitizing, controlled cooling or annealing. While strength is achieved through the nitriding's curing effect, the aim is to achieve high ductility through the distribution and morphology of the nitrides, but above all by limiting grain coarsening during forging and during solution heat treatment.
  • EP 237170 discloses a heat-resistant steel for gas turbine of 0.05-0.2 C, below 0.5 Si, 0.1-0.40 Mn, 8-13 Cr, 1.5-3 Mo, 2-3 Ni, 0.05-0.3 V, 0.02-0.2 Nb and / or Ta, 0.02-0.1 N, below 0.5 Co below l W, below 0.01 B, below 0.3 Al, below 0.5 Ti, below 0.1 Zr, below 0.1 Hf, below 0.01 Ca, below 0.01 Mg, below 0.01 Y, below 0.01 SE, below 0.5 Cu at a Mn to Ni ratio below 0.11, balance iron.
  • a heat resistant steel having good toughness properties is known for use as a turbine rotor having the following chemical composition (% by weight): 0.05-0.30 C, 0.20 or less Si, 0-1.0 Mn, 8-14 Cr, 0.5-3.0 Mo , 0.10-0.50 V, 1.5-5.0 Ni, 0.01-0.5 Nb, 0.01-0.08 N, 0.001-0.020 B, balance iron and unavoidable impurities. Boron microalloying results in precipitates at the grain boundaries and increases the time stability of the carbonitrides at high temperatures, but higher levels of B reduce the toughness of the steel. Disadvantages of this proposed composition are also the relatively high permitted Si values of 0.2%. Although Si serves advantageously as a deoxidizer at the time of melting, parts of it remain as oxides in the steel, which is disadvantageous in a reduced toughness.
  • stainless steels containing 8-13 wt.% Cr which include boron and rare earths in their chemical composition, in order to increase the resistance to long-term aging embrittlement.
  • the maximum content of rare earths, for example Y, La, Ce, Er, according to this document should be 0.5% by weight, an optimum proportion is given as 0.1% by weight.
  • the boron content is given as 0.001-0.04% by weight.
  • the steels still have the following elements (in% by weight): 0.08-0.15 C, at least one element from the group of precious 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 O, max. 0.006 ace, max. 0.003 bb, max.
  • precious 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
  • the steel may additionally contain up to 0.50% by weight of Nb.
  • the austenite stabilizers are described as containing as much as possible Co in the steel while at the same time maintaining the Ni content should be minimized. This balance between the Ni and Co content is important, according to the authors, to suppress unwanted embrittlement phenomena while ensuring the desired toughness of the steel.
  • These steels are said to provide good properties in high temperature applications, ie, balanced mechanical and oxidative properties. For example, to provide a steel for high-temperature turbine components which has good resistance to embrittlement, oxidation and creep.
  • the invention has for its object to provide an 8-13% Cr steel, which is characterized over the prior art by increased creep strength at temperatures of 550 ° C and above, as well as improved LCF properties and a comparatively high Has toughness. It should preferably find application for rotors of thermal turbomachinery, so that the efficiency and the output of these machines over the known prior art can be increased.
  • the core of the invention is a steel having the following chemical composition (in% by weight): 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 the range of 0.5 to 2.0 or in the presence of both elements is a maximum of 3.0, 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, 80 to 120 ppm B, maximum 100 ppm Al, maximum 150 ppm P, maximum 250 ppm As, maximum 120 ppm Sn, maximum 30 ppm Sb, maximum 50 ppm S, balance iron and unavoidable impurities.
  • the steel according to the invention has the following chemical composition (in% by weight): 0.12 C, 11.5 Cr, 0.2 Mn, 2.5 Ni, 1.7 Mo, 0.25 V, 0.03 Nb, 0.06 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, balance iron and unavoidable impurities.
  • the advantage of the invention is that the alloy according to the invention has improved creep properties at temperatures of 550 ° C. and above, compared to alloys of similar composition known from the prior art, but without B additive or without Pd additive Toughness properties and a higher fatigue strength (LCF) is achieved.
  • LCF fatigue strength
  • a starting structure which is characterized by a tough matrix and the presence of heat-resistant nitrides, borides and carbides.
  • the toughness of the base matrix is adjusted by the presence of substitution elements, preferably nickel.
  • substitution elements preferably nickel.
  • the contents of these substitution elements are determined to provide optimal unfolding of both martensite hardening and particle hardening by precipitation of special nitrides, e.g. As vanadium nitrides or niobium nitrides, to set the highest heat resistance possible.
  • both hardening mechanisms lower the ductility. Characteristically, a minimum ductility is observed in the area of secondary hardening. This minimum ductility need not be caused exclusively by the actual precipitation hardening mechanism. A certain embrittlement contribution may also be provided by segregation of impurities to the grain boundaries or possibly also by near-order adjustments of dissolved alloy atoms.
  • Manganese is on the left side next to the element iron in the periodic system of elements. It is an electron-poorer element, so its action in solid solution should be distinctly different from that of nickel. Nonetheless, it is an austenite stabilizing element which greatly lowers the Ac1 temperature, but leaves no particularly positive but rather unfavorable effect on ductility.
  • manganese is understood to be an impurity element which promotes temper embrittlement substantially. Therefore, the content of manganese is usually limited to very small amounts.
  • a weight proportion of 8-13% chromium allows good through-hardenability of thick-walled components and ensures sufficient oxidation resistance up to a temperature of 550 ° C.
  • a weight fraction below 8% 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 11 to 12%.
  • the range to be specified should be in the range between 0.1 and 0.5% by weight, preferably between 0.1 and 0.25%, in particular at 0.2% by weight and for silicon at 0.03-0.15, preferably at ⁇ 0.05%. % By weight.
  • Nickel is used as an austenite stabilizing element to suppress delta ferrite. In addition, it is said to improve ductility as a dissolved element in the ferritic matrix. Nickel contents of 2 to about 3 wt .-% are useful. Nickel contents above 4% by weight increase the austenite stability such that after the solution annealing and tempering an increased proportion of retained austenite or tempering austenite in the tempered Martensite may be present.
  • the nickel content is preferably 2.3 to 2.7, in particular 2.5% by weight.
  • Molybdenum and tungsten improve creep strength by solid solution hardening as partially dissolved elements and precipitation hardening during long-term stress.
  • 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).
  • the desired range for Mo and W is in each case 0.5 to 2% by weight, preferably 1.6 to 1.8% by weight, in particular 1.7% by weight. If both elements are present, the total proportion should not exceed 3% by weight.
  • V / N ratio sometimes also increases the stability of the vanadium nitride over the chromium nitride.
  • the specific content of nitrogen and vanadium nitrides depends on the optimum volume fraction of the vanadium nitrides, which are to remain as insoluble primary nitrides during the solution annealing. The larger the total content of vanadium and nitrogen, the greater is the proportion of vanadium nitrides which no longer dissolve and the greater the grain refining effect.
  • the preferred content of nitrogen is in the range from 0.02 to 0.08% by weight, preferably 0.025 to 0.055% by weight, particularly preferably 0.04% by weight N, and that of vanadium is in the range between 0.1 and 0.4% by weight. , preferably 0.2 to 0.3% by weight, and especially at 0.25% by weight.
  • Niobium is a strong nitride former that aids the grain refining effect. In order to keep the volume fraction of the primary nitrides small, their total proportion must be limited. Niobium dissolves in vanadium nitride in small amounts and can thus improve the stability of the vanadium nitride. Niobium is added in the range between 0.02 and 0.2% by weight, preferably 0.02 to 0.04% by weight, and in particular with 0.03% by weight.
  • these elements increase the embrittlement of long-term aging in the range between 350 and 500 ° C. These elements should therefore be limited to maximum tolerable levels (150 ppm P, 120 ppm Sn).
  • Ta influences the creep resistance positively. Addition of 0.05 to 2% by weight of Ta has the effect, on the one hand, that due to the greater tendency of tantalum to form carbide as chromium, the precipitation of undesirable chromium carbides at the grain boundaries and, on the other hand, the undesirable depletion of the mixed crystal in chromium are reduced.
  • the preferred range for Ta is 0.05 to 0.1% by weight, in particular a Ta content of 0.06% by weight should be set.
  • the carbon content should therefore be limited upwards to 0.15% by weight.
  • the disadvantage is also the The fact that carbon reinforces the hardening during welding.
  • the preferred carbon content is in the range between 0.10 and 0.14% by weight, preferably 0.12% by weight.
  • the boron content should be limited to 80 to 120 ppm. It is preferable to adjust a B content of 100 ppm.
  • Pd forms an ordered Fe-Pd L1 0 intermetallic phase with the iron of the steel, the ⁇ "phase.
  • This stable ⁇ " phase increases the creep rupture strength at high temperatures by stabilizing the grain boundary precipitates, such as M 23 C 6 , and acts thus have a positive effect on the creep properties.
  • palladium has the disadvantage of high costs.
  • the Pd content of the proposed steel should be in the range of 0.005 to 2, preferably 0.005 to 0.01 wt%, with a content of 0.005 wt%, ie 50 ppm Pd, being particularly suitable.
  • the investigated inventive alloy L1 had the following chemical composition (in% by weight): 0.12 C, 11.5 Cr, 0.2 Mn, 2.5 Ni, 1.7 Mo, 0.25 V, 0.03 Nb, 0.06 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, balance iron and unavoidable impurities.
  • the comparative alloy VL1 used was a commercial steel of the type X12CrNiMoV11-2-2 known from the prior art, which is characterized by the following chemical composition (in% by weight): 0.10-0.14 C, 11.0-12.0 Cr, 0.25 Mn , 2.0-2.6 Ni, 1.3-1.8 Mo, 0.2-0.35 V, 0.02-0.05 N, 0.15 Si, 0.026 P, 0.015 S, balance Fe and unavoidable impurities.
  • the comparative alloy VL1 was solution-annealed at 1065 ° C and then subjected to a tempering treatment at 640 ° C.
  • Fig. 1 shows for the two alloys VL1 and L1 the creep properties, ie the creep rupture strength at 550 ° C. In this diagram, the mean times to break are thus shown, depending on the voltage at 550 ° C.
  • the inventive alloy L1 both after a heat treatment "A” and after a heat treatment "B” advantageously longer times when exposed to the same stress to break required as the reference alloy VL1.
  • an arrowheaded sample of alloy L1 has not yet broken.
  • a clear shift towards longer times can be seen here, which is of particular advantage for the planned use as a gas turbine or steam turbine rotor.
  • Fig. 2 the strain amplitude is plotted against the number of cycles until the crack at 575 ° C with 10 minutes holding time in the tensile region for the inventive alloy L1.
  • Fig. 3 the fracture toughness and impact energy at room temperature are compared for the two investigated alloys according to the above-described heat treatment state with a subsequent aging (3000 h at 480 ° C).
  • the fracture toughness hardly deteriorates in the alloy according to the invention Impact work is slightly increased.
  • the alloy L1 according to the invention has no stronger embrittlement tendency than the comparative alloy VL1.
  • the alloy according to the invention is characterized on the one hand by a very good creep resistance and a high resistance to low-cycle fatigue at temperatures of 550 ° C. and above, and thus is superior to conventional 12% Cr steels.
  • This is mainly due to the influence of boron, tantalum and palladium, which are alloyed in the specified range. Boron, tantalum and palladium stabilize the M 23 C 6 precipitates, which play a significant strengthening role during creep, with Pd additionally forming a stable intermetallic phase with the iron, which also contributes to increasing creep resistance.
  • the dislocation density is maintained until fracture, thus improving the strength of the steel.
  • the alloy of the present invention has improved resistance to embrittlement upon long-term aging and comparatively high toughness, as well as high resistance to fatigue.
  • the inventive alloy is thus particularly advantageous for rotors in gas and steam turbines, which are exposed to high inlet temperatures of for example 550 ° C and above, can be used advantageously.

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  • Chemical & Material Sciences (AREA)
  • Engineering & Computer Science (AREA)
  • Materials Engineering (AREA)
  • Mechanical Engineering (AREA)
  • Metallurgy (AREA)
  • Organic Chemistry (AREA)
  • Turbine Rotor Nozzle Sealing (AREA)
  • Heat Treatment Of Steel (AREA)
  • Heat Treatment Of Articles (AREA)
EP09151605A 2008-02-25 2009-01-29 Kriechfester Stahl Active EP2116626B1 (de)

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EP2116626B1 true EP2116626B1 (de) 2010-12-22

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US (1) US20090214376A1 (ja)
EP (1) EP2116626B1 (ja)
JP (1) JP2009280901A (ja)
CN (1) CN101519757B (ja)
AT (1) ATE492661T1 (ja)

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* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
JP2011176183A (ja) * 2010-02-25 2011-09-08 Toyota Motor Corp 半導体装置の製造方法
CN104195472A (zh) * 2014-07-29 2014-12-10 锐展(铜陵)科技有限公司 一种钨钒钼合金钢及其制造方法
CN112696204A (zh) * 2021-02-03 2021-04-23 洛阳九久科技股份有限公司 一种菱形结构帽型齿刀圈及其制造工艺

Family Cites Families (14)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
EP0237170B1 (en) * 1986-02-05 1994-05-11 Hitachi, Ltd. Heat resistant steel and gas turbine composed of the same
US5320687A (en) * 1992-08-26 1994-06-14 General Electric Company Embrittlement resistant stainless steel alloy
JPH06306550A (ja) * 1993-04-28 1994-11-01 Toshiba Corp 耐熱鋼及びその熱処理方法
JPH0726351A (ja) * 1993-07-12 1995-01-27 Hitachi Metals Ltd 高温強度の優れたフェライト系耐熱鋼
DE19712020A1 (de) 1997-03-21 1998-09-24 Abb Research Ltd Vollmartensitische Stahllegierung
JPH10265909A (ja) 1997-03-25 1998-10-06 Toshiba Corp 高靭性耐熱鋼、タービンロータ及びその製造方法
US5820817A (en) 1997-07-28 1998-10-13 General Electric Company Steel alloy
US5906791A (en) * 1997-07-28 1999-05-25 General Electric Company Steel alloys
JPH1171641A (ja) * 1997-08-29 1999-03-16 Japan Casting & Forging Corp 高強度耐熱鋼
JPH11209851A (ja) 1998-01-27 1999-08-03 Mitsubishi Heavy Ind Ltd ガスタービンディスク材
JP4221518B2 (ja) * 1998-08-31 2009-02-12 独立行政法人物質・材料研究機構 フェライト系耐熱鋼
DE10025808A1 (de) 2000-05-24 2001-11-29 Alstom Power Nv Martensitisch-härtbarer Vergütungsstahl mit verbesserter Warmfestigkeit und Duktilität
KR100532877B1 (ko) * 2002-04-17 2005-12-01 스미토모 긴조쿠 고교 가부시키가이샤 고온강도와 내식성이 우수한 오스테나이트계 스테인레스강및 상기 강으로부터 이루어지는 내열 내압부재와 그제조방법
DE502005005216D1 (de) * 2004-10-29 2008-10-09 Alstom Technology Ltd Kriechfester martensitisch-härtbarer vergütungsstahl

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US20090214376A1 (en) 2009-08-27
JP2009280901A (ja) 2009-12-03
ATE492661T1 (de) 2011-01-15
CN101519757B (zh) 2013-07-17
EP2116626A1 (de) 2009-11-11
CN101519757A (zh) 2009-09-02

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