US20090214376A1 - Creep-resistant steel - Google Patents

Creep-resistant steel Download PDF

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Publication number: US20090214376A1
Authority: US; United States
Prior art keywords: creep; resistant steel; ppm; weight; maximum
Prior art date: 2008-02-25
Legal status (The legal status is an assumption and is not a legal conclusion. Google has not performed a legal analysis and makes no representation as to the accuracy of the status listed.): Abandoned

Application number

US12/390,740

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English (en)

Inventor

Mohamed Nazmy

Andreas Kuenzler

Markus Staubli

Current Assignee (The listed assignees may be inaccurate. Google has not performed a legal analysis and makes no representation or warranty as to the accuracy of the list.)

GE Vernova GmbH

Original Assignee

Alstom Technology AG

Priority date (The priority date is an assumption and is not a legal conclusion. Google has not performed a legal analysis and makes no representation as to the accuracy of the date listed.)

2008-02-25

Filing date

2009-02-23

Publication date

2009-08-27

2009-02-23 Application filed by Alstom Technology AG filed Critical Alstom Technology AG

2009-02-23 Assigned to ALSTOM TECHNOLOGY LTD reassignment ALSTOM TECHNOLOGY LTD ASSIGNMENT OF ASSIGNORS INTEREST (SEE DOCUMENT FOR DETAILS). Assignors: NAZMY, MOHAMED, STAUBLI, MARKUS, KUENZLER, ANDREAS

2009-08-27 Publication of US20090214376A1 publication Critical patent/US20090214376A1/en

Status Abandoned legal-status Critical Current

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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

LCF Low Cycle Fatigue
Martensitically hardenable steel materials based on 9-12% chromium are materials in widespread use in power station technology. They were developed for application in steam power stations at operating temperatures of above 600° C. and steam pressures of above 250 bar, in order to increase the efficiency of the power stations. Under these operating conditions, the creep resistance and the oxidation resistance of the material play a particular part.
chromium in the abovementioned range not only allows high resistance to atmospheric corrosion, but also makes it possible to have the full hardenability of thick-walled forgings, such as are employed, for example, as monobloc rotors or as rotor disks in gas and steam turbines.
Proven alloys of this type can contain about 0.08 to 0.2% carbon which, in solution, makes it possible to set a hard martensitic structure.
a good combination of heat resistance and ductility of martensitic steels is made possible by an annealing treatment in which a particle-stabilized subgranular structure is formed as a result of the precipitation of carbon in the form of carbides, while at the same time the dislocation substructure is recovered.
the annealing behavior and the properties resulting from this can be influenced effectively by the choice and the co-ordination, in terms of quantity fractions, of special carbide formers, such as, for example, Mo, W, V, Nb and Ta.
the contents of Cr, Mo and W were modified, taking into account N, Nb and/or B, in order to improve the creep and rupture strengths for applications at 600° C.
the carbides such as, for example, M 23 C 6 , are to be stabilized by the addition of boron.
the Ni contents were restricted to values of lower than 0.25%.
a disadvantage is that the fracture toughness values are low, and, although this does not play an important part in steam turbine applications and may therefore be ignored, it should be avoided in gas turbine applications.
European patent application EP 0 931 845 A1 describes a nickel-containing 12% chromium steel which is similar in constitution to the German steel X12CrNiMo12 and in which the element, molybdenum, was reduced, as compared with the known steel X12CrNiMo12, but an increased content of tungsten was added for alloying.
DE 198 32 430 A1 discloses a modification of a steel which is of the same type as X12CrNiMo12 and in which the tendency to embrittlement in the temperature range between 425 and 500° C. is limited by the addition of rare earth elements or boron.
a disadvantage is that the strength, in particular the heat resistance at temperatures of between 300 and 600° C., could not be improved in any of the abovementioned developed steels at a high ductility level comparable to that of the steel X12CrNiMo12.
EP 0 866 145 A2 describes a class of martensitic chromium steels with nitrogen contents in the range between 0.12 and 0.25%, and, in EP 1 158 067 A1, with nitrogen contents of 0.12 to 0.18%, the weight ratio V/N lying in the range between 3.5 and 4.2.
the entire structural generation is controlled by the formation of special nitrides, in particular of vanadium nitrides, which can be distributed in many different ways by forging treatment, by austenitization, by controlled cooling treatment or by annealing treatment.
an aim is to set a high ductility by the distribution and morphology of the nitrides, but, above all, by limiting the granular coarsening during forging and during solution heat treatment.
a heat-resistant steel with good toughness properties for use as a turbine rotor is known from EP 0867 522 A2 and has the following chemical composition (% by weight): 0.05-0.30 C, 0.20 or less Si, 0-1.0 Mn, 8-14 Cr, 0.5-3.0 Mo, 0.10-0.50 V, 1.5-5.0 Ni, 0.01-0.5 Nb, 0.01-0.08 N, 0.001-0.020 B, the rest iron and unavoidable impurities.
Microalloying with boron can lead to precipitations at the grain boundaries and increase the long-term stability of the carbonitrides at high temperatures, although higher contents of B reduce the toughness of the steel.
Disadvantages of this proposed composition are also the relatively high permitted Si values in the amount of 0.2%.
Si advantageously serves as a deoxidant at the melting time point, on the other hand, parts of this remain as oxides in the steel, and this is reflected adversely in a reduced toughness.
the steels also have the following elements (values in % by weight): 0.08-0.15 C, at least one element from the group of noble metals, such as Ru, Rh, Os, Pt, Pd, Ir, in the range of 0.01-2.00, 0.01-0.1 Si, at least one element from the group of W and Mo in the range of 0.50-4.00, at least one austenite stabilizer (such as Ni, Co, Mn Cu) in the range of 0.001-6.00, 0.25-0.40 V, 0.001-0.025 Al, max. 0.01 P, max. 0.004 S, max. 0.060 N, max. 2 ppm H, max. 50 ppm 0, max. 0.006 As, max. 0.003 Sb, max.
noble metals such as Ru, Rh, Os, Pt, Pd, Ir
W and Mo in the range of 0.50-4.00
austenite stabilizer such as Ni, Co, Mn Cu
the steel may additionally contain up to 0.50% by weight Nb.
Nb 0.50% by weight
the austenite stabilizers it is described, with regard to the austenite stabilizers, that the steel is to contain as much Co as possible, while at the same time the Ni content is to be minimized. According to the authors' statement, this balance between the Ni and Co content is used to suppress undesirable embrittlement phenomena and at the same time to ensure the desired toughness of the steel.
good properties in high-temperature applications are to be achieved, that is to say balanced mechanical and oxidation properties.
a steel for high-temperature turbine components is thereby to be made available, which has high resistance to embrittlement, oxidation and creep.
a creep-resistant steel comprising: a chemical composition (values in % by weight) of: about 0.10 to 0.15 C, 8 to 13 Cr, 0.1 to 0.5 Mn, 2 to 3 Ni; at least one or both elements from the group Mo, W in each case in a range of about 0.5 to 2.0 or, if both elements are present, a maximum total of about 3.0; about 0.02 to 0.2 Nb, 0.05 to 2 Ta, 0.1 to 0.4 V, 0.005 to 2 Pd, 0.02 to 0.08 N, 0.03 to 0.15 Si; about 80 to 120 ppm B, maximum about 100 ppm Al, maximum about 150 ppm P, maximum about 250 ppm As, maximum 120 ppm Sn, maximum 30 about ppm Sb, maximum about 50 ppm S, a remainder of the composition being iron and impurities.
a chemical composition values in % by weight of: about 0.10 to 0.15 C, 8 to 13 Cr, 0.1 to 0.5 Mn, 2 to 3 Ni; at least one or both elements from
a creep-resistant steel comprising: a chemical composition (values in % by weight) of: iron; 8 to 13 Cr; designated percentages by weight of Mn, Ni, Nb, Ta, V, N and Si; about 0.005 to 2 Pd; about 80 to 120 ppm B; and at least one of Mo and W.
FIG. 1 shows a graphical illustration in which stresses of selected alloys (VL 1 according to the prior art and an exemplary L 1 according to the present disclosure) are plotted against time, at a temperature of 550° C., up to the fracture of the material, two different heat treatment methods (with two different annealing temperatures) having been used for alloy L 1 ;
FIG. 2 shows a graphical illustration in which an exemplary elongation amplitude is plotted against the number of load cycles up to incipient cracking at 575° C. for the alloy L 1 according to the disclosure and at 500° C. for the comparative alloy VL 1 , and
FIG. 3 shows a graphical illustration in which the fracture toughness (left part image) and the notched bar impact work (right part image) at room temperature are compared for the two alloys L 1 and VL 1 after heat treatment and additional age hardening of 3000 hours at 480° C.
Exemplary embodiments disclosed herein can provide an 8-13% Cr steel which can possess an increased creep resistance at temperatures of 550° C. and above, and also possess improved LCF properties and comparatively high toughness relative to known steel. Exemplary embodiments can be used for rotors of thermal turbomachines, so that the efficiency and output of these machines can be increased, as compared with known steel.
an exemplary steel can include (e.g., consist of) a chemical composition (values in % by weight) of: about 0.10 to 0.15 C, 8 to 13 Cr, 0.1 to 0.5 Mn, 2 to 3 Ni; at least one or both elements from the group Mo, W in each case in a range of about 0.5 to 2.0 or, if both elements are present, a maximum total of about 3.0; about 0.02 to 0.2 Nb, 0.05 to 2 Ta, 0.1 to 0.4 V, 0.005 to 2 Pd, 0.02 to 0.08 N, 0.03 to 0.15 Si; and about 80 to 120 ppm B, maximum about 100 ppm Al, maximum about 150 ppm P, maximum about 250 ppm As, maximum about 120 ppm Sn, maximum about 30 ppm Sb, maximum about 50 ppm S, a remainder of the composition being iron and impurities (e.g., unavoidable impurities).
the term “about” refers to an amount (e.g.
Exemplary preferred ranges for individual alloying elements of the composition according to the disclosure include chemical composition (values in % by weight) of: 0.12 C, 11.5 Cr, 0.2 Mn, 2.5 Ni, 1.7 Mo, 0.25 V, 0.03 Nb, 0.06 Ta, 50 ppm Pd, 100 ppm B, 0.04 N, ⁇ 0.01 Al, ⁇ 0.01 P, ⁇ 0.005 S, ⁇ 0.05 Si, ⁇ 0.012 Sn, ⁇ 0.025 As, ⁇ 0.0025 Sb, a remainder of the composition being iron and impurities (e.g., unavoidable impurities).
chemical composition values in % by weight of: 0.12 C, 11.5 Cr, 0.2 Mn, 2.5 Ni, 1.7 Mo, 0.25 V, 0.03 Nb, 0.06 Ta, 50 ppm Pd, 100 ppm B, 0.04 N, ⁇ 0.01 Al, ⁇ 0.01 P, ⁇ 0.005 S, ⁇ 0.05 Si, ⁇ 0.012 Sn, ⁇ 0.025 As, ⁇ 0.0025 Sb
An exemplary advantage of the alloy according to the disclosure is creep properties at temperatures of 550° C. and above, in comparison with known alloys of similar composition without a B addition or without a Pd addition. Improved toughness properties and higher fatigue strength (LCF) can also be achieved in accordance with exemplary embodiments.
a tempered structure can be set which is distinguished by a tough basic matrix and by a presence of nitrides, borides and carbides which afford heat resistance.
the toughness of the basic matrix can be set by the presence of substitution elements, such as by nickel.
substitution elements such as by nickel.
the contents of these substitution elements can be determined such that they allow an optimal development both of martensitic hardening and of particle hardening due to the precipitation of special nitrides, for example vanadium nitrides or niobium nitrides, for the purpose of setting the highest possible heat resistances.
a ductility minimum is in this case characteristically observed in the region of secondary hardening. This ductility minimum need not be caused solely by the actual precipitation hardening mechanism. A certain contribution to embrittlement may also be made by the segregation of impurities up to the grain boundaries or, possibly, also by near-order settings of dissolved alloy atoms.
a rise in the annealing temperature over the secondary hardening range leads to complete precipitation, with a marked growth of carbides.
the strength decreases and the ductility increases.
Ductility can increase to a great extent due to the simultaneous recovery of the dislocation substructure and particle coarsening, so that the combination of strength and ductility, overall, can be improved. This improvement can be attributed to the formation of a particle-stabilized subgranular structure. It can be assumed, in this case, that both the ductility and the strength of particle-stabilized subgranular structures can be reduced by nonuniformities in the topology of the particle subgranular structure.
Precipitations at subgranular boundaries can be subject to accelerated coarsening and can tend to coagulate with adjacent precipitations. Coarse and coagulated phases can generate fracture-triggering stress peaks which lower the ductility. Above all, however, the hardening mechanism which is the most effective at high temperatures, to be precise the particle hardening, can be seriously limited by the nonuniform distribution of the precipitations.
the strength is to be increased by a lowering of the annealing temperature to below 700° C.
a partial conversion of ferrite into austenite can be reckoned on during annealing. This is associated with a certain ductility-promoting grain reforming.
the carbide precipitation takes place only incompletely above the Ac1 temperature, since the solubility of the austenite-stabilizing element, carbon, is higher in austenite than in ferrite.
the austenite which is formed is not sufficiently stabilized, and therefore a larger volume fraction of the reformed austenite can be subjected to further martensitic conversion during recooling after annealing.
Manganese lies on the left side next to the element iron, in the periodic system of elements. It is an electron-leaner element, and therefore its action in a solid solution should be markedly different from that of nickel. Nonetheless, it is an austenite-stabilizing element which can greatly lower the Ac1 temperature, and does not have an especially positive, but, instead, a somewhat adverse effect on the ductility. With regard to carbon-containing 12% chromium steels, manganese is understood to be a contaminating element which appreciably promotes annealing embrittlement. The manganese content can therefore be limited to very small quantities.
a weight fraction of 8-13% chromium allows a good full hardenability of thick-walled components and ensures sufficient oxidation resistance up to a temperature of 550° C.
a weight fraction of below 8% can be detrimental to full hardening.
Contents above 13% can lead to the accelerated formation of hexagonal chromium nitrides during the annealing operation, which, in addition to nitrogen, can also bind vanadium, and, consequently, lower the effectiveness of hardening by vanadium nitrides.
An exemplary optimal chromium content is about 11 to 12%.
the range to be specified can lie, for manganese, in the range between about 0.1 and 0.5% by weight, preferably between about 0.1 and 0.25%, in particular at 0.2% by weight, and, for silicon, at about 0.03-0.15, preferably at ⁇ 0.05% by weight.
Nickel can be used as an austenite-stabilizing element for the suppression of delta ferrite. Furthermore, as a dissolved element in the ferritic matrix, it can improve ductility. Nickel contents of 2 to about 3% by weight can be expedient. Nickel contents above 4% by weight can intensify the austenite stability in such a way that, after solution heat treatment and annealing, an increased fraction of residual austenite or annealing austenite may be present in the hardened martensite. The nickel content lies at, for example, about 2.3 to 2.7, in particular at 2.5% by weight.
Molybdenum and tungsten can improve the creep resistance by solid solution hardening as partially dissolved elements and by precipitation hardening during long-term stress.
an excessively high fraction of these elements can lead to embrittlement during long-term age hardening, which can be due to the precipitation and coarsening of the Laves phase (W, Mo) and sigma phase (Mo).
An exemplary desired range for Mo and W is in each case about 0.5 to 2% by weight, preferably about 1.6 to 1.8% by weight, in particular 1.7% by weight. If both elements are present, the overall fraction can amount to a maximum of about 3% by weight.
V/N ratio sometimes can also increase the stability of the vanadium nitride with respect to the chromium nitride.
the actual content of nitrogen and vanadium nitrides can depend on the optimal volume fraction of the vanadium nitrides which can remain as insoluble primary nitrides during the solution heat treatment. The larger the overall fraction of vanadium and nitrogen is, the larger that fraction of the vanadium nitrides which is no longer dissolved is, and the higher the grain-refining action.
An exemplary preferred nitrogen content lies in the range of about 0.02 to 0.08% by weight, preferably about 0.025 to 0.055% by weight, particularly preferably at 0.04% by weight N, and the vanadium content lies in the range of between about 0.1 and 0.4% by weight, preferably about 0.2 to 0.3% by weight, and, in particular, at 0.25% by weight.
Niobium is a strong nitride former which promotes the grain-refining action. In order to keep the volume fraction of primary nitrides low, its overall fraction can be limited. Niobium dissolves in small quantities in vanadium nitride and can consequently improve the stability of the vanadium nitride. Niobium can be added for alloying in the range of between about 0.02 and 0.2% by weight, preferably about 0.02 to 0.04% by weight, and, in particular, at 0.03% by weight.
These elements can intensify annealing embrittlement during long-term age hardening in the range between about 350 and 500° C. These elements can therefore be limited to maximum acceptable fractions (150 ppm P, 120 ppm Sn).
Ta can positively influence the creep resistance. Alloying with about 0.05 to 2% by weight Ta has the effect that, because of the greater tendency of tantalum to form carbides than chromium, on the one hand, the precipitation of undesirable chromium carbides at the grain boundaries can be diminished and, on the other hand, the undesirable depletion of the chromium mixed crystal can also be reduced.
An exemplary preferred range for Ta is about 0.05 to 0.1% by weight, and, in particular, a Ta content of about 0.06% by weight should be set.
Carbon during annealing, forms chromium carbides which are conducive to improved creep resistance. At carbon contents which are too high, however, the increased volume fraction of carbides which results from this can lead to a ductility reduction which, in particular, can be reflected by carbide coarsening during long-term age hardening.
the carbon content can therefore have an upper limit of about 0.15% by weight.
An exemplary disadvantage can be that carbon intensifies hardening during welding.
An exemplary preferred carbon content lies in the range between about 0.10 and 0.14% by weight, preferably at 0.12% by weight.
Boron stabilizes the M 23 C 6 precipitations, and hence can improve the creep resistance of the steel and reduce the annealing embrittlement, although the formation of boron nitrides at the expense of the vanadium carbonitrides can be prevented.
the austenitization temperature can be increased in order to obtain homogeneous boron in the matrix, but this, in turn can lead to an increase in the grain size and consequently to poorer properties of the material.
the boron content can therefore be limited to about 80 to 120 ppm. Also, an exemplary B content of about 100 ppm can preferably be set.
Pd forms, with the iron of the steel, an ordered intermetallic Fe—Pd L 1 0 phase, the ⁇ ′′ phase.
This stable ⁇ ′′ phase can increase the rupture strength at high temperatures by the stabilization of the grain boundary precipitations, such as, for example, M 23 C 6 , and therefore has a positive effect on the creep properties.
palladium can have the exemplary disadvantage of high costs.
the Pd content of the proposed steel should lie in the range of about 0.005 to 2, preferably of about 0.005 to 0.01% by weight, a content of 0.005% by weight, that is to say 50 ppm, of Pd being particularly suitable.
the investigated alloy L 1 according to the disclosure had the following chemical composition (values in % by weight): 0.12 C, 11.5 Cr, 0.2 Mn, 2.5 Ni, 1.7 Mo, 0.25 V, 0.03 Nb, 0.06 Ta, 0.04 N, 0.005 Pd, 0.01 B, ⁇ 0.01 Al, ⁇ 0.01 P, ⁇ 0.005 S, ⁇ 0.05 Si, ⁇ 0.012 Sn, ⁇ 0.025 As, ⁇ 0.0025 Sb, the remainder of the composition being iron and unavoidable impurities.
the comparative alloy VL 1 used was a commercial steel of the type X12CrNiMoV11-2-2 which is known from the prior art and which includes (e.g., consists of) chemical composition (values in % by weight) of: 0.10-0.14 C, 11.0-12.0 Cr, 0.25 Mn, 2.0-2.6 Ni, 1.3-1.8 Mo, 0.2-0.35 V, 0.02-0.05 N, 0.15 Si, 0.026 P, 0.015 S, the remainder of the composition being Fe and impurities (e.g., unavoidable impurities).
the two alloys therefore have an approximately comparable composition, the difference being that the alloy L 1 according to the disclosure is additionally microalloyed with Nb, B, Ta and Pd.
the alloy L 1 according to the disclosure was subjected to the following exemplary two-stage heat treatment processes:
the exemplary comparative alloy VL 1 was solution heat-treated at 1065° C. and was subsequently subjected to annealing treatment at 640° C.
FIG. 1 shows the properties during creeping, that is to say the rupture strength, at 550° C. for the two alloys VL 1 and L 1 .
This graph thus illustrates the average times up to fracture as a function of the stress at 550° C.
the exemplary alloy L 1 according to the disclosure advantageously involves, both after heat treatment “A” and after heat treatment “B”, longer times under the action of the same stress up to fracture than the comparative alloy VL 1 .
the sample, given an arrow in FIG. 1 of the alloy L 1 has not yet even been fractured.
a marked shift toward longer times can be seen, and this can be particularly advantageous for the planned use as a gas turbine rotor or steam turbine rotor.
an exemplary elongation amplitude is plotted against the number of load cycles up to incipient cracking at 575° C., with a holding time of 10 minutes in the tensile range, for the alloy L 1 according to the disclosure.
the fracture toughness and the notched bar impact work at room temperature are compared for the two alloys investigated, after the above-described heat treatment state with subsequent age hardening (3000 h at 480° C.).
the exemplary alloy according to the disclosure there is scarcely any impairment in the fracture toughness and the notched bar impact work is slightly increased.
the exemplary alloy L 1 according to the disclosure therefore has no greater tendency to embrittlement than the comparative alloy VL 1 .
This very good property combination (very high creep resistance at temperatures of 550° C. and above, good toughness properties after long-term age hardening at high temperatures, and, moreover, very high fatigue strength at these high temperatures) can be achieved, as compared with the prior art, by the alloying elements as a whole, for example, by the combination of B, Ta and Pd in the exemplary ranges specified.
the exemplary alloy according to the disclosure can be distinguished by a very high creep resistance and high resistance to low-cycle fatigue at temperatures of 550° C. and above and can be consequently superior to the known 12% Cr steels. This is attributable at least in part to the influence of boron, tantalum and palladium which are added for alloying in the specified range. Boron, tantalum and palladium can stabilize the M 23 C 6 precipitations which can play a substantial consolidating part during creeping, Pd additionally forming a stable intermettalic phase with the iron, this also contributing to increasing the creep resistance. In addition, the dislocation density up to fracture can be maintained and therefore the strength of the steel can be improved. On the other hand, exemplary alloys according to the disclosure can have improved resistance to embrittlement during long-term ageing and comparatively high toughness and also high resistance to fatigue.
Exemplary alloys according to the disclosure can therefore advantageously be used particularly for rotors in gas and steam turbines which are exposed to high inlet temperatures of, for example, 550° C. and above.

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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)

US12/390,740 2008-02-25 2009-02-23 Creep-resistant steel Abandoned US20090214376A1 (en)

Applications Claiming Priority (2)

Application Number	Priority Date	Filing Date	Title
CH00270/08		2008-02-25
CH2702008		2008-02-25

Publications (1)

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US12/390,740 Abandoned US20090214376A1 (en)	2008-02-25	2009-02-23	Creep-resistant steel

Country Status (5)

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

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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

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KR100532877B1 (ko) *	2002-04-17	2005-12-01	스미토모 긴조쿠 고교 가부시키가이샤	고온강도와 내식성이 우수한 오스테나이트계 스테인레스강및 상기 강으로부터 이루어지는 내열 내압부재와 그제조방법
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2009
- 2009-01-29 AT AT09151605T patent/ATE492661T1/de active
- 2009-01-29 EP EP09151605A patent/EP2116626B1/de not_active Not-in-force
- 2009-02-23 US US12/390,740 patent/US20090214376A1/en not_active Abandoned
- 2009-02-24 JP JP2009040079A patent/JP2009280901A/ja active Pending
- 2009-02-25 CN CN200910004415.5A patent/CN101519757B/zh not_active Expired - Fee Related

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Publication number	Publication date
JP2009280901A (ja)	2009-12-03
CN101519757A (zh)	2009-09-02
EP2116626A1 (de)	2009-11-11
EP2116626B1 (de)	2010-12-22
ATE492661T1 (de)	2011-01-15
CN101519757B (zh)	2013-07-17

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