High-nitrogen ferritic heatresisting steel with high niobium content and method of production thereof

Abstract: A high-nitrogen ferritic heat-resisting steel with high niobium content comprises, in weight per cent, 0.01 - 0.30% C, 0.02 - 0.80% Si, 0.20 - 1.00% Mn, 8.00 - 13.00% Cr, 0.005 - 1.00% Mo, 0.20 - 1.50% W, 0.05 - 1.00% V, over 0.12 up to 2.00% Nb and 0.10 - 0.50% N, the balance being Fe and unavoidable impurities. A method of producing the steel comprises melting and equilibrating the steel components in an atmosphere of a mixed gas of a prescribed nitrogen partial pressure or nitrogen gas and thereafter casting or solidifying the resulting melt in an atmosphere controlled to have a total pressure of not less than 2.5 bar and a nitrogen partial pressure of not less than 1.0 bar, with the relationship between the nitrogen partial pressure p and the total pressure P being P > 2.5p.

Description: BACKGROUND OF THE INVENTION Field of the Invention This invention relates to a ferritic heat-resisting steel, more particularly to a high-nitrogen ferritic heat-resisting steel containing chromium and appropriate for use in a high-temperature, high-pressure environment, and to a method of producing the same. Description of the Prior Art Recent years have seen a marked increase in the temperatures and pressures under which thermal power plant boilers are required to operate. Some plans already call for operation at 566°C and 314 bar and it is expected that operation under conditions of 650°C and 355 bar will be implemented in the future. These are extremely severe conditions from the viewpoint of the boiler materials used. At operating temperatures exceeding 550°C, it has, from the points of oxidation resistance and high-temperature strength, been necessary to switch from ferritic 2 · 1/4 Cr - 1 Mo steel to high-grade austenitic steels such as 18 - 8 stainless steel. In other words, it has been necessary to adopt expensive materials with properties exceeding what is required. Decades have been spent in search of steels for filling in the gap between 2 · 1/4 Cr - 1 Mo steel and austenitic stainless steel. Medium Cr (e.g. 9 Cr and 12 Cr) steel boiler pipes are made of heat-resisting steels that were developed against this backdrop. They achieve high-temperature strength and creep rupture strength on a par with austenitic steels by use of a base metal composition which includes various alloying elements for precipitation hardening and solution hardening. The creep rupture strength of a heat-resisting steel is governed by solution hardening in the case of short-term aging and by precipitation hardening in the case of prolonged aging. This is because the solution-hardening elements initially present in solid solution in the steel for the most pert precipitate as stable carbides such as M 23 C 6 during aging, and then when the aging is prolonged these precipitates coagulate and enlarge, with a resulting decrease in creep rupture strength. Thus, with the aim of maintaining the creep rupture strength of heat-resisting steels at a high level, a considerable amount of research has been done for discovering ways for avoiding the precipitation of the solution hardening elements and maintaining them in solid solution for as long as possible. For example, Japanese Patent Public Disclosures No. Sho 63-89644, Sho 61-231139 and Sho 62-297435 teach ferritic steels that achieve dramatically higher creep rupture strength than conventional Mo-containing ferritic heat-resisting steels by the use of W as a solution hardening element. While the solution hardening by W in these steels may be more effective than by Mo, the precipitates are still fundamentally carbides of the N 23 C 6 type, so that it is not possible to avoid reduction of the creep rupture strength with prolonged aging. Moreover, the use of ferritic heat-resisting steels at up to 650°C has been considered difficult because of their inferior high-temperature oxidation resistance as compared with austenitic heat-resisting steels. A particular problem with these steels is the pronounced degradation of high-temperature oxidation resistance that results from the precipitation of Cr in the form of coarse N 23 C 6 type precipitates at the grain boundaries. FR-A-1 140 573 discloses ferritic steels comprising especially higher levels of chromium (9-20%) and 0.001-0.2% nitrogen. Although the document teaches to improve hardness by precipitation it does not mention the specific nature of these precipitates. Further, the document does not disclose the importance of nitrides informing precipitates or even the importance of a high nitrogen content at all. The highest temperature limit for use of ferritic heat-resisting steel has therefore been considered to be 600°C. The need for heat-resisting steels capable of standing up under extremely severe conditions has grown more acute not only because of the increasingly severe operating conditions mentioned earlier but also because of plans to reduce operating costs by extending the period of continuous power plant operation from the current 100 thousand hours up to around 150 thousand hours. Although ferritic heat-resisting steels are somewhat inferior to austenitic steels in high-temperature strength and anticorrosion property, they have a cost advantage. Furthermore, for reasons related to the difference in thermal expansion coefficient, among the various steam oxidation resistance properties they are particularly superior in scale defoliation resistance. For these reasons, they are attracting attention as a boiler material. For the reasons set out above, however, it is clearly not possible with the currently available technology to develop ferritic heat-resisting steels that are capable of standing up for 150 thousand hours under operating conditions of 650°C and 355 bar, that are low in price and that exhibit good steam oxidation resistance. Based on the foregoing knowledge and as described in Japanese Patent Application No. Hei 2-37895, the inventors earlier disclosed that a high-nitrogen ferritic heat-resisting steel estimated by linear extrapolation to exhibit a creep rupture strength of not less than 147 MPa under operating conditions of 650°C and 355 bar for 150 thousand hours can be obtained by using a pressurized atmosphere to add nitrogen exceeding the solution limit and thus inducing precipitation of the excess nitrogen in the form of fine nitrides and carbo-nitrides. The gist of their disclosure was a ferritic heat-resisting steel characterized in comprising, in weight per cent, 0.01 - 0.30% C, 0.02 - 0.80% Si, 0.20 - 1.00% Mn, 8.00 - 13.00% Cr, 0.50 - 3.00% W, 0.005 - 1.00% Mo, 0.05 - 0.50% V, 0.02 -0.12% Nb and 0.20 - 0.50% N and being controlled to include not more than 0.050% P, not more than 0.010% S and not more than 0.020% O, and optionally comprising (A) one or both of 0.01 - 1.00% Ta and 0.01 - 1.00% Hf and/or (B) one or both of 0.0005 - 0.10% Zr and 0.01 - 0.10% Ti, the balance being Fe and unavoidable impurities and a method of producing the steel wherein the steel components are melted and equilibrated in an atmosphere of a mixed gas of a prescribed nitrogen partial pressure or nitrogen gas and the resulting melt is thereafter cast or solidified in an atmosphere controlled to have a nitrogen partial pressure of not less than 1.0 bar and a total pressure of not less than 4.0 bar, with the relationship between the partial pressure p and the total pressure P being 10 p < P 0.37 + log 10 6 thereby obtaining good quality ingot free of blowholes. Based on the results of tests for determining the creep rupture strength of the steel taught by Japanese Patent Application No. Hei 2-37895 up to 50 thousand hours, the inventors discovered that the creep rupture strength of the steel at 150 thousand hours, as estimated by linear extrapolation, is no more than 176 MPa and, in particular, that the steel experiences a marked decrease in creep rupture strength between 30 and 50 thousand hours. Further studies showed that the reason for the decrease in creep rupture strength was that during the creep test large Fe 2 W grains measuring 1 µm or more in diameter precipitated in large amounts, principally at the grain boundaries, leading to large-scale loss of W as a solid solution element from the steel. Moreover, they further discovered that by limiting the W content to not more than 1.5% so as to prevent precipitation of W as Fe 2 W and by adding Nb in excess of 0.12% so that NbN and (Nb, V)N, the most stable of all nitrides, become the principal precipitation hardening factors, it is possible to obtain a ferritic heat-resisting steel exhibiting a creep rupture strength at 650°C, 355 bar and 150 thousand hours of not less than 200 MPa, as estimated by linear extrapolation. In addition they discovered that owing to the increase in the N solution limit resulting from the addition of a large amount of Nb the pressurized atmosphere conditions required for casting of sound ingot become a total pressure of not less than 2.5 bar and a nitrogen partial pressure of not less than 1.0 bar, with the relationship between the total pressure P and the nitrogen partial pressure p being P > 2.5p. There have been few papers published on research into high-nitrogen ferritic heat-resisting steels and the only known published report in this field is Ergebnisse der Werkstoff-Forschung, Band I, Varlag Schweizerische Akademie der Werkstoffwissenschaften "Thubal-Kain", Zurich, 1987, 161 - 180. However, the research described in this report is limited to that in connection with ordinary heat-resisting steel and there is no mention of materials which can be used under such severe conditions as 650°C, 355 bar and 150 thousand hours continuous operation. SUMMARY OF THE INVENTION An object of this invention is to provide a high-nitrogen ferritic heat-resisting steel which overcomes the shortcomings of the conventional heat-resisting steels and particularly to provide such a steel exhibiting outstanding creep rupture strength and capable of being used under severe operating conditions, wherein the decrease in creep rupture strength following prolonged aging and the degradation of high-temperature oxidation resistance caused by precipitation of carbides are mitigated by adding nitrogen to supersaturation so as to precipitate fine nitrides and/or carbo-nitrides which suppress the formation of carbides such as the M 23 C 6 precipitates seen in conventional steels. This invention was accomplished in the light of the aforesaid knowledge and, in one aspect, pertains substantially to a high-nitrogen ferritic heat-resisting steel with high niobium content comprising, in weight per cent, 0.01 - 0.30% C, 0.02 - 0.80% Si, 0.20 - 1.00% Mn, 8.00 - 13.00% Cr, 0.005 - 1.00% Mo, 0.20 - 1.50% W, 0.05 - 1.00% V, over 0.12 up to 2.00% Nb and more than 0.20 - 0.50% N and being controlled to include not more than 0.050% P, not more than 0.010% S and not more than 0.020% O, and optionally comprising (A) one or both of 0.01 - 1.00% Ta and 0.01 - 1.00% Hf and/or (B) one or both of 0.0005 - 0.10% Zr and 0.01 - 0.10% Ti, the balance being Fe and unavoidable impurities, wherein the N content of the steel of said chemical composition exceeds the N concentration in the steel equilibrated with 1 bar N gas during casting. Another aspect of the invention pertains to a method of producing such a high-nitrogen ferritic heat-resisting steel with high niobium content wherein the steel components are melted and equilibrated in an atmosphere of a mixed gas of a prescribed nitrogen partial pressure or nitrogen gas and the resulting melt is thereafter cast or solidified in an atmosphere controlled to have a total pressure of not less than 2.5 bar and a nitrogen partial pressure of not less than 1.0 bar, with the relationship between the nitrogen partial pressure p and the total pressure P being P > 2.5p. The above and other features of the present invention will become apparent from the following description made with reference to the drawings. BRIEF DESCRIPTION OF THE DRAWINGS Figure 1 is a perspective view of an ingot and the manner in which it is to be cut. Figure 2 is a graph showing the relationship between the steel nitrogen content and the weight percentage of the total of M 23 C 6 + M 6 C + NbC + Cr 2 N + NbN among the precipitates in the steel accounted for by M 23 C 6 + M 6 C + NbC and the relationship between the steel nitrogen content and the weight percentage of the total of M 23 C 6 + M 6 C + NbC + Cr 2 N + NbN among the precipitates in the steel accounted for by Cr 2 N + NbN. Figure 3 is a graph showing conditions under which blowholes occur in the ingot in terms of the relationship between the total pressure and nitrogen partial pressure of the atmosphere during casting. Figure 4 is a schematic view showing the manner in which creep test pieces are taken from a pipe specimen and a rolled plate specimen. Figure 5 is a graph showing the relationship between steel nitrogen content and estimated creep rupture strength at 650°C, 150 thousand hours. Figure 6 is a graph showing the relationship between steel Nb content and estimated creep rupture strength at 650°C, 150 thousand hours. Figure 7 is a graph showing the relationship between steel W content and estimated creep rupture strength at 650°C, 150 thousand hours. Figure 8 is a graph showing an example of creep test results in terms of stress vs rupture time. Figure 9 is a graph showing the relationship between steel nitrogen content and Charpy impact absorption energy at 0°C following aging at 700°C for 10 thousand hours. Figure 10 is a graph showing the relationship between steel nitrogen content and the thickness of the oxidation scale formed on the surface of a test piece after oxidation at 650°C for 10 thousand hours. DETAILED DESCRIPTION OF THE PREFERRED EMBODIMENT The reasons for the limits placed on the components of the high-nitrogen ferritic heat-resisting steel with high Nb content according to this invention will now be explained. C is required for achieving strength. Adequate strength cannot be achieved at a C content of less than 0.01%, while at a C content exceeding 0.30% the steel is strongly affected by welding heat and undergoes hardening which becomes a cause for low-temperature cracking. The C content range is therefore set at 0.01 - 0.30%. Si is important for achieving oxidation resistance and is also required as a deoxidizing agent. It is insufficient for these purposes at a content of less than 0.02%, whereas a content exceeding 0.80% reduces the creep rupture strength. The Si content range is therefore set at 0.02 - 0.80%. Mn is required for deoxidation and also for achieving strength. It has to be added at least 0.20% for adequately exhibiting its effect. When it exceeds 1.00% it may in some cases reduce creep rupture strength. The Mn content range is therefore set at 0.20 - 1.00%. Cr is indispensable to oxidation resistance. It also contributes to increasing creep resistance by combining with N and finely precipitating in the base metal matrix in the form of Cr 2 N, Cr 2 (C, N) and the like. Its lower limit is set at 8.00% from the viewpoint of oxidation resistance. Its upper limit is set at 13.00% for maintaining the Cr equivalent value at a low level so as to realize a martensite phase texture. W produces a marked increase in creep rupture strength by solution hardening. Its effect toward increasing creep rupture strength over long periods at high temperatures of 550°C and higher is particularly pronounced. Its upper limit is set at 1.50% because at contents higher than this level it precipitates in large quantities in the form of carbide and intermetallic compounds which sharply reduce the toughness of the base metal. The lower limit is set at 0.20% because it does not exhibit adequate solution hardening effect at lower levels. Mo increases high-temperature strength through solution hardening. It does not exhibit adequate effect at a content of less than 0.005% and at a content higher than 1.00% it may, when added together with W, cause heavy precipitation of Mo 2 C type oxides which markedly reduce base metal toughness. The Mo content range is therefore set at 0.005 - 1.00%. V produces a marked increase in the high-temperature strength of the steel regardless of whether it forms precipitates or, like W, enters solid solution in the matrix. When it precipitates, the resulting VN and (Nb, V)N serve as precipitation nuclei for Cr 2 N and Cr 2 (C, N), which has a pronounced effect toward promoting fine dispersion of the precipitates. It has no effect at below 0.05% and reduces toughness at higher than 1.00%. The V content range is therefore set at 0.05 - 1.00%. Nb is an element which increases high-temperature strength by precipitating as NbN, (Nb, V)N, Nb(C, N) and (Nb, V)(C, N). Also, similarly to V, it promotes fine precipitate dispersion by forming precipitation nuclei for Cr 2 N, Cr 2 (C, N) and the like. For it to disperse in the steel as the primary precipitation hardening factor it has to be added in excess of 0.12%. However, its upper limit is set at 2.00% because when present at higher levels it reduces strength by causing precipitate coagulation and enlargement. N dissolves in the matrix and also forms nitride and carbo-nitride precipitates. As the form of the precipitates is mainly Cr 2 N and Cr 2 (C, N), there is less precipitate-induced consumption of Cr and W than in the case of the M 23 C 6 , M 6 C and other such precipitates observed in conventional steels. N thus increases oxidation resistance and creep rupture strength. At least 0.10% is required for precipitation of nitrides and carbo-nitrides and suppressing precipitation of M 23 C 6 and M 6 C. The upper limit is set at 0.50% for preventing coagulation and enlargement of nitride and carbo-nitride precipitates by the presence of excessive nitrogen. Steels according to the invention are defined to comprise from >0.20 to 0.50% N. P, S and O are present in the steel according to this invention as impurities. P and S hinder the achievement of the purpose of the invention by lowering strength, while 0 has the adverse effect of forming oxides which reduce toughness. The upper limits on these elements is therefore set at 0.050%, 0.010% and 0.020%, respectively. The basic components of the steel according to this invention (aside from Fe) are as set out above. Depending on the purpose to which the steel is to be put, however, it may additionally contain (A) one or both of 0.01 - 1.00% Ta and 0.01 - 1.00% Hf and/or (B) one or both of 0.0005 - 0.10% Zr and 0.01 - 0.10% Ti. At low concentrations Ta and Hf act as deoxidizing agents. At high concentrations they form fine high melting point nitrides and carbo-nitrides and, as such, increase toughness by decreasing the austenite grain size. In addition, they also reduce the degree to which Cr and W dissolve in precipitates and by this effect enhance the effect of supersaturation with nitrogen. Neither element exhibits any effect at less than 0.01%. When either is present at greater than 1.00%, it reduces toughness by causing enlargement of nitride and carbo-nitride precipitates. The content range of each of these elements is therefore set at 0.01 - 1.00%. Acting to govern the deoxidation equilibrium in the steel, Zr suppresses the formation of oxides by markedly reducing the amount of oxygen activity. In addition, its strong affinity for N promotes precipitation of fine nitrides and carbo-nitrides which increase creep rupture strength and high-temperature oxidation resistance. When present at less than 0.0005% it does not provide an adequate effect of governing the deoxidation equilibrium and when present at greater than 0.10% it results in heavy precipitation of coarse ZrN and ZrC which markedly reduce the toughness of the base metal. The Zr content range is therefore set at 0.0005 - 0.10%. Ti raises the effect of excess nitrogen by precipitating in the form of nitrides and carbo-nitrides. At a content of less than 0.01% it has no effect while a Ti content of over 0.10% results in precipitation of coarse nitrides and carbo-nitrides which reduce toughness. The Ti content range is therefore set at 0.01 - 0.10%. The aforesaid alloying components can be added individually or in combinations. The object of this invention is to provide a ferritic heat-resisting steel that is superior in creep rupture strength and high-temperature oxidation resistance. Depending on the purpose of use it can be produced by various methods and be subjected to various types of heat treatment. These methods and treatments in no way diminish the effect of the invention. However, in view of the need to supersaturate the steel with nitrogen, it is necessary during casting to raise the total pressure of the atmosphere to not less than 2.5 bar and to control the relationship between the total pressure P and the nitrogen partial pressure p to satisfy the inequation P > 2.5p. As an auxiliary gas to be mixed with the nitrogen gas it is appropriate to use an inert gas such as Ar, Ne, Xe or Kr. These casting conditions were determined by the following experiment. Steel of a chemical composition, aside from nitrogen, as indicated in claims 1 - 4 was melted in an induction heating furnace installed in a chamber that could be pressurized up to 150 bar. A mixed gas of argon and nitrogen having a nitrogen partial pressure adequate for achieving the target nitrogen content was introduced into the furnace and maintained at a pressure which was varied from test to test. After the nitrogen and molten metal had reached chemical equilibrium, the molten metal was cast into a mold that had been installed in the chamber beforehand, whereby there was obtained a 5-ton ingot. The ingot was cut vertically as shown in Figure 1 and the ingot 1 was visually examined for the presence of blowholes. Following this examination, a part of the ingot was placed in a furnace and maintained at 1180°C for 1 hour and then forged into a plate measuring 50 mm in thickness, 750 mm in width and 4,000 mm in length. This plate was subjected to solution treatment at 1200°C for 1 hour and to tempering at 800°C for 3 hours. The steel was then chemically analyzed and the dispersion state and morphology of the nitrides and carbo-nitrides were investigated by observation with an optical microscope, an electron microscope, X-ray diffraction and electron beam diffraction, whereby the chemical structure was determined. Among the precipitates present within the as-heat-treated steel, Figure 2 shows how the proportion of the precipitates in the steel accounted for by M 23 C 6 type carbides and M 6 C or NbC type carbides and the proportion thereof accounted for by Cr 2 N type nitrides and NbN type nitrides vary with nitrogen concentration. At a nitrogen concentration of 0.10%, nitrides account for the majority of the precipitates in the steel of the invention, while at a nitrogen concentration of 0.15%, substantially 100% of the precipitates are nitrides with virtually no carbides present whatsoever. Thus for the effect of this invention to be adequately manifested it is necessary for the nitrogen concentration of the steel to be more than 0.2%. The graph of Figure 3 shows how the state of blowhole occurrence varies depending on the relationship between the total and nitrogen partial pressures of the atmosphere. For achieving a nitrogen concentration of 0.10% or higher it is necessary to use a total pressure of not less than 2.5 bar. Equilibrium calculation based on Sievert's law shows that in this case the nitrogen partial pressure in the steel of this invention is not less than 1.0 bar. Moreover, where for controlling the amount of nitride and carbo-nitride precipitation the nitrogen partial pressure is maintained at 1.0 - 6.0 bar (nitrogen concentration within the steel of approximately 0.5 mass%), it becomes necessary to vary the total pressure between 2.5 and about 15 bar, the actual value selected depending on the nitrogen partial pressure. Namely, it is necessary to use a total pressure falling above the broken line representing the boundary pressure in Figure 3. When the boundary line of Figure 3 is determined experimentally it is found to lie at P = 2.5p meaning that the steel according to this invention can be obtained by selecting an atmosphere of a pressure and composition meeting the condition of the inequality P > 2.5p. It is therefore necessary to use furnace equipment enabling pressure and atmosphere control. Without such equipment, it is difficult to produce the steel of the present invention. There are no limitations whatever on the melting method. Based on the chemical composition of the steel and cost considerations, it suffices to select from among processes using a converter, an induction heating furnace, an arc melting furnace or an electric furnace. The situation regarding refining is similar. Insofar as the atmosphere is controlled to a total pressure of not less than 2.5 bar and a nitrogen partial pressure of not less than 1.0 bar, it is both possible and effective to use a ladle furnace, an electro-slag remelting furnace or a zone melting furnace. After casting under a pressurized atmosphere of a total pressure of not less than 2.5 bar and a nitrogen partial pressure of not less than 1.0 bar, it is possible to process the steel into billet, bloom or plate by forging or hot rolling. Since the steel of this invention includes finely dispersed nitrides and carbo-nitrides, it is superior to conventional ferritic heat-resisting steels in hot-workability. This is also one reason for employing nitrides and carbo-nitrides obtained by adding nitrogen to beyond the solution limit. For processing the steel into products, it is possible to first process it into round or rectangular billet and then form it into seamless pipe or tube by hot extrusion or any of various seamless rolling methods. Otherwise it can be formed into sheet by hot and cold rolling and then made into welded tube by electric resistance welding. Alternatively, it can be processed into welded pipe or tube by use of TIG, MIG, SAW, LASER and EB welding, individually or in combination. Moreover, it is possible to expand the size range of products to which the present invention can be applied by following any of the aforesaid processes by hot or warm stretch reduction or sizing. The steel according to the invention can also be provided in the form of plate or sheet. The plate or sheet can, in its hot-rolled state or after whatever heat treatment is found necessary, be provided as a heat-resisting material in various shapes, without any influence on the effects provided by the invention. The pipe, tube, plate, sheet and variously shaped heat-resisting materials referred to above can, in accordance with their purpose and application, be subjected to various heat treatments, and it is important for them to be so treated for realizing the full effect of the invention. While the production process ordinarily involves normalizing (solution heat treatment) + tempering, it is also possible and useful additionally to carry out one or a combination of two or more of quenching, tempering and normalizing. It is also possible, without influencing the effects of the present invention in any way, to repeatedly carry out one or more of the aforesaid processes to whatever degree is necessary for adequately bringing out the steel properties. The aforesaid processes can be appropriately selected and applied to the manufacture of the steel according to the invention. Working examples: The steels indicated in Tables 1 - 14, each having a composition according to one of claims 1 - 4, were separately melted in amounts of 5 tons each in an induction heating furnace provided with pressurizing equipment. The resulting melt was cleaned by ladle furnace processing (under bubbling with a gas of the same composition as the atmosphere) for reducing its impurity content, whereafter the atmosphere was regulated using a mixed gas of nitrogen and argon so as to satisfy the conditions of the inequality shown in claim 5. The melt was then cast into a mold and processed into a round billet, part of which was hot extruded to obtain a tube 60 mm in outside diameter and 10 mm in wall thickness and the remainder of which was subjected to seamless rolling to obtain a pipe 380 mm in outside diameter and 50 mm in wall thickness. The tube and pipe were subjected to a single normalization at 1200°C for 1 hour and were then tempered at 800°C for 3 hours. In addition, a 5 ton ingot was cast and forged into a slab which was hot rolled into 25 mm and 50 mm thick plates. As shown in Figure 4, creep test pieces 6 measuring 6 mm in diameter were taken along the axial direction 4 of the pipe or tube 3 and along the rolling direction 5 of the plates and subjected to creep test measurement at 650°C. Based on the data obtained, a linear extrapolation was made for estimating the creep rupture strength at 150 thousand hours. A creep rupture strength of 150 MPa was used as the creep rupture strength evaluation reference value. The creep rupture strength at 650°C, 150 thousand hours is hereinafter defined as the linearly extrapolated value at 150 thousand hours on the creep rupture strength vs rupture time graph. Toughness was evaluated through an accelerated evaluation test in which aging was carried out at 700°C for 10 thousand hours. JIS No. 4 tension test pieces were cut from the aged steel and evaluated for impact absorption energy. Assuming a water pressure test at 0°C, the toughness evaluation reference value was set at 10 J. High-temperature oxidation resistance was evaluated by suspending a 25 mm x 25 mm x 5 mm test piece cut from the steel in 650°C atmospheric air in a furnace for 10 thousand hours and then cutting the test piece parallel to the direction of growth of the scale and measuring the oxidation scale thickness. The 650°C, 150 thousand hour creep rupture strength, the Charpy impact absorption energy at 0°C after aging at 700°C for 10 thousand hours and the oxidation scale thickness after oxidation at 650°C for 10 thousand hours are shown in Tables 2, 4, 6, 8, 10, 12, and 14. For comparison, steels of compositions not falling within any of the claims 1 to 4 were melted, processed and tested in the same way as described above. Their chemical compositions and the evaluation results are shown in Tables 15 and 16. Figure 5 shows the relationship between the nitrogen content of the steels and the estimated creep rupture strength at 650°C, 150 thousand hours. It will be noted that the creep rupture strength attains high values exceeding 150 MPa at a steel nitrogen content of 0.1% or higher but falls below 150 MPa and fails to satisfy the evaluation reference value that was set at a steel nitrogen content of less than 0.1%. Figure 6 shows the relationship between the Nb content of the steels and the estimated creep rupture strength at 650°C, 150 thousand hours. It will be noted that the creep rupture strength attains values exceeding 150 MPa at a steel Nb content exceeding 0.12% but at a Nb content of 2.0% or higher the creep rupture strength is instead lowered owing to the precipitation of coarse NbN and Fe 2 Nb type Laves phase at the melting stage. Figure 7 shows the relationship between the W content of the steels and the estimated creep rupture strength at 650°C, 150 thousand hours. The creep rupture strength is below 150 MPa at a W content of less than 0.2% and is 150 MPa or higher in a content range of 0.2 - 1.5%. When the W is present in excess of 1.5%, the creep rupture strength falls below 150 MPa owing to coarse Fe 2 W precipitating at the grain boundaries. Figure 8 shows the results of the creep test in terms of stress vs rupture time. A good linear relationship can be noted between stress and rupture time at a steel nitrogen content of not less than 0.1%. Moreover, the creep rupture strength is high. On the other hand, when the steel nitrogen content falls below 0.1%, the relationship between stress and rupture time exhibits a pronounced decline in creep rupture strength with increasing time lapse. Either the linearity is not maintained, or the slope of the creep rupture curve is steep, with the short-term side creep rupture strength being high but the long-term creep rupture strength being low, or the creep rupture strength is low throughout. This is because W and the other solution hardening elements precipitate as carbides whose coagulation and enlargement degrades the creep rupture strength property of the base metal. In contrast, at a nitrogen content of 0.1% or higher, fine nitrides are preferentially precipitated so that the formation of carbides is greatly delayed. Therefore, since the dissolution of the solution hardening elements into carbides was suppressed and also because the finely precipitated nitrides remained present in a stable state without coagulating and enlarging during the long-term high-temperature creep test, a high creep rupture strength was maintained in the long-term creep test. Figure 9 shows the relationship between Charpy impact absorption energy at 0 °C following aging at 700°C for 10 thousand hours and steel nitrogen content. When the steel nitrogen content falls within the range of 0.1 - 0.5%, the impact absorption energy exceeds 10 J. In contrast, when it falls below 0.1%, there is little or no suppression of grain growth by residual high melting point nitrides during solution treatment and, as a result, the impact absorption energy decreases, and when it exceeds 0.5%, the impact absorption energy is reduced by heavy nitride precipitation. Figure 10 shows the relationship between the thickness of the oxidation scale formed on the surface of a test piece after oxidation at 650°C for 10 thousand hours and the steel nitrogen content. Although the oxidation scale thickness is between 400 and 900 µm when the steel nitrogen content falls below 0.1%, it decreases to 50 µm or less when the steel nitrogen content is 0.1% or higher. Reference is now made to the comparison steels shown in Table 15. Nos. 161 and 162 are examples in which insufficient steel nitrogen content resulted in a low estimated creep rupture strength at 650°C, 150 thousand hours and also to poor high-temperature oxidation resistance. Nos. 163 and 164 are examples in which excessive steel nitrogen content caused heavy precipitation of coarse nitrides and carbo-nitrides, resulting in a Charpy impact absorption energy at 0°C after aging at 700°C for 10 thousand hours of not more than 10 J. No. 165 is an example in which a low W concentration resulted in a low creep rupture strength at 650°C, 150 thousand hours owing to insufficient solution hardening notwithstanding that the steel nitrogen content fell within the range of the invention. No. 166 is an example in which a high W concentration led to low rupture strength and toughness owing to precipitation of coarse Fe 2 W type Laves phase at the grain boundaries during creep. No. 167 is an example in which a low Nb content resulted in a low estimated creep rupture strength at 650°C, 150 thousand hours. No. 168 is an example in which a high Nb content caused profuse precipitation of coarse Fe 2 Nb type Laves phase during creep, which in turn lowered both the estimated creep rupture strength at 650°C, 150 thousand hours and the Charpy impact absorption energy at 0°C after aging at 700°C for 10 thousand hours. No. 169 is an example in which heavy precipitation of coarse ZrN caused by a Zr concentration in excess of 0.1% resulted in a Charpy impact absorption energy at 0°C after aging at 700°C for 10 thousand hours of less than 10 J. Nos. 170, 171 and 172 are examples similar to the case of No. 169 except that the elements present in excess were Ta, Hf and Ti, respectively. As a result, heavy precipitation of coarse TaN, HfN and TiN resulted in a Charpy impact absorption energy at 0°C after aging at 700°C for 10 thousand hours of less than 10 J. No. 173 is an example in which, notwithstanding that the steel composition satisfied the conditions of claims 1 to 4, since the nitrogen partial pressure was 2.2 bar and the total pressure was 2.5 bar, values not satisfying the inequality of claim 5, many large blowholes formed in the ingot, making it impossible to obtained either a sound ingot or a plate and leading to a reduction in both the estimated creep rupture strength at 650°C, 150 thousand hours and the Charpy impact absorption energy at 0°C after aging at 700°C for 10 thousand hours. (mass%) Invention steels No. C Si Mn P S Nb V Cr Mo 1 0.195 0.251 0.801 0.048 0.008 0.666 0.094 9.52 0.586 2 0.070 0.475 0.632 0.013 0.005 0.132 0.370 10.38 0.038 3 0.192 0.709 0.261 0.042 0.009 1.185 0.174 8.70 0.707 4 0.219 0.220 0.299 0.018 0.004 1.064 0.496 11.18 0.427 5 0.096 0.190 0.620 0.022 0.003 0.608 0.066 9.82 0.865 6 0.270 0.515 0.216 0.014 0.001 0.680 0.462 9.07 0.843 7 0.098 0.369 0.583 0.039 0.003 0.927 0.105 11.51 0.748 8 0.259 0.477 0.995 0.039 0.002 0.582 0.089 10.40 0.179 9 0.083 0.594 0.404 0.021 0.007 1.824 0.050 11.70 0.820 10 0.206 0.478 0.961 0.035 0.004 0.336 0.073 8.84 0.366 11 0.100 0.663 0.812 0.026 0.003 1.163 0.469 12.74 0.504 12 0.127 0.600 0.467 0.024 0.002 0.392 0.454 9.62 0.843 13 0.092 0.722 0.203 0.022 0.006 1.747 0.246 12.49 0.704 14 0.243 0.320 0.395 0.036 0.007 1.583 0.165 9.38 0.613 15 0.060 0.713 0.388 0.013 0.001 0.194 0.285 11.20 0.732 16 0.101 0.227 0.520 0.041 0.008 0.897 0.200 8.70 0.117 (Continued from Table 1) (mass%) Invention steels No. W Zr Ta Hf Ti N O CS MPa VE J TO µm 1 1.154 - - - - 0.375 0.002 231 19.4 46 2 0.716 - - - - 0.242 0.010 214 17.7 46 3 1.169 - - - - 0.422 0.019 230 55.3 38 4 0.497 - - - - 0.210 0.007 183 54.3 47 5 1.187 - - - - 0.376 0.002 209 77.9 19 6 0.608 - - - - 0.217 0.012 153 64.1 36 7 0.473 - - - - 0.440 0.006 157 16.9 36 8 0.329 - - - - 0.301 0.017 226 31.2 20 9 0.420 0.050 - - - 0.270 0.009 226 14.1 16 10 0.999 0.092 - - - 0.427 0.006 178 49.6 11 11 1.006 0.031 - - - 0.445 0.019 201 45.7 33 12 0.686 0.015 - - - 0.218 0.012 175 32.4 35 13 1.464 0.026 - - - 0.228 0.017 213 35.1 38 14 0.323 0.028 - - - 0.276 0.002 166 64.9 46 15 1.363 0.40 - - - 0.438 0.005 202 28.6 26 16 1.086 - 0.275 - - 0.293 0.013 195 13.6 34 CS: Creep rupture strength at 650 °C, 150 thousand hours; VE: Charpy impact absorption energy at 0°C after aging at 700°C for 10 thousand hours; TO: Oxidation scale thickness after 650°C, 10 thousand hour high-temperature oxidation (mass%) Invention steels No. C Si Mn P S Nb V Cr Mo 17 0.027 0.796 0.445 0.045 0.005 1.306 0.157 12.42 0.859 18 0.273 0.684 0.616 0.036 0.006 1.054 0.209 12.23 0.954 19 0.093 0.727 0.740 0.018 0.007 1.824 0.100 10.41 0.459 20 0.120 0.290 0.555 0.017 0.004 1.164 0.363 10.34 0.055 21 0.248 0.694 0.487 0.019 0.007 0.572 0.391 10.20 0.241 22 0.055 0.208 0.602 0.046 0.001 1.432 0.286 9.19 0.080 23 0.198 0.752 0.361 0.048 0.002 1.631 0.296 12.13 0.939 24 0.292 0.519 0.709 0.013 0.005 1.583 0.056 9.91 0.011 25 0.106 0.449 0.217 0.049 0.003 1.617 0.334 11.12 0.494 26 0.127 0.685 0.219 0.023 0.002 0.190 0.304 11.57 0.563 27 0.118 0.576 0.459 0.022 0.006 0.492 0.278 9.73 0.406 28 0.111 0.602 0.658 0.043 0.006 1.759 0.272 11.70 0.878 29 0.063 0.416 0.740 0.013 0.007 1.035 0.252 10.85 0.331 30 0.220 0.528 0.523 0.040 0.009 0.723 0.380 9.33 0.379 31 0.211 0.280 0.956 0.039 0.006 0.957 0.055 12.74 0.470 32 0.158 0.227 0.728 0.025 0.005 0.334 0.227 9.64 0.408 33 0.241 0.279 0.798 0.020 0.004 1.557 0.210 12.12 0.195 34 0.078 0.197 0.915 0.030 0.008 0.262 0.310 8.54 0.524 (Continued from Table 3) (mass%) Invention steels No. W Zr Ta Hf Ti N O CS MPa VE J TO µm 17 1.374 - 0.611 - - 0.377 0.001 238 15.8 45 18 0.756 - 0.221 - - 0.329 0.016 160 30.8 49 19 1.272 - 0.295 - - 0.468 0.007 232 63.1 13 20 1.225 - 0.134 - - 0.266 0.015 225 55.9 23 21 1.356 0.007 0.337 - - 0.379 0.019 159 32.3 29 22 1.460 0.012 0.494 - - 0.378 0.019 221 68.5 21 23 0.246 0.003 0.098 - - 0.411 0.016 194 59.9 37 24 1.227 0.017 0.325 - - 0.353 0.005 200 21.0 26 25 1.401 0.001 0.118 - - 0.284 0.017 158 69.4 13 26 0.337 0.084 0.681 - - 0.476 0.009 186 77.7 47 27 0.899 0.031 0.079 - - 0.399 0.002 152 24.2 15 28 1.463 - - 0.509 - 0.301 0.019 172 67.8 14 29 1.266 - - 0.36 - 0.423 0.014 165 26.7 46 30 0.495 - - 0.954 - 0.202 0.007 226 50.2 43 31 1.456 - - 0.215 - 0.439 0.012 235 66.8 49 32 0.429 - - 0.239 - 0.254 0.013 221 16.1 33 33 1.116 - - 0.775 - 0.229 0.007 152 48.8 13 34 1.143 - - 0.511 - 0.241 0.008 195 71.6 18 CS: Creep rupture strength at 650 °C, 150 thousand hours; VE: Charpy impact absorption energy at 0°C after aging at 700°C for 10 thousand hours; TO: Oxidation scale thickness after 650°C, 10 thousand hour high-temperature oxidation (mass%) Invention steels No. C Si Mn P S Nb V Cr Mo 35 0.013 0.514 0.430 0.043 0.001 0.562 0.385 9.86 0.024 36 0.244 0.202 0.686 0.027 0.008 1.922 0.074 9.51 0.451 37 0.251 0.594 0.652 0.029 0.003 0.881 0.135 10.99 0.969 38 0.229 0.256 0.857 0.036 0.002 0.636 0.203 10.17 0.274 39 0.091 0.135 0.271 0.016 0.001 0.590 0.138 8.04 0.047 40 0.032 0.745 0.624 0.017 0.007 1.097 0.471 12.99 0.466 41 0.252 0.293 0.474 0.026 0.007 1.554 0.160 8.45 0.225 42 0.081 0.484 0.769 0.030 0.009 1.403 0.416 12.86 0.661 43 0.011 0.299 0.870 0.018 0.008 0.988 0.430 12.23 0.017 44 0.093 0.653 0.510 0.015 0.004 0.628 0.051 9.89 0.914 45 0.236 0.532 0.650 0.031 0.006 1.657 0.267 11.83 0.408 46 0.060 0.620 0.630 0.026 0.007 1.736 0.134 9.91 0.926 47 0.138 0.327 0.757 0.021 0.005 0.544 0.396 12.80 0.318 48 0.061 0.155 0.791 0.024 0.004 1.198 0.245 11.28 0.057 49 0.095 0.311 0.345 0.023 0.005 1.922 0.377 11.03 0.788 50 0.169 0.165 0.971 0.015 0.007 0.839 0.384 9.23 0.673 51 0.018 0.714 0.898 0.045 0.002 0.192 0.314 9.23 0.154 52 0.271 0.143 0.664 0.020 0.004 0.309 0.471 9.42 0.402 53 0.119 0.553 0.823 0.029 0.003 0.583 0.261 8.20 0.605 54 0.033 0.126 0.712 0.036 0.003 0.506 0.447 8.99 0.944 55 0.240 0.227 0.929 0.013 0.003 0.416 0.228 10.50 0.658 56 0.054 0.575 0.388 0.045 0.009 1.583 0.129 10.93 0.898 (Continued from Table 5) (mass%) Invention steels No. W Zr Ta Hf Ti N O CS MPa VE J TO µm 35 0.750 0.065 - 0.132 - 0.363 0.010 171 28.4 13 36 0.200 0.047 - 0.829 - 0.469 0.007 249 38.4 49 37 0.378 0.073 - 0.297 - 0.243 0.011 247 78.5 31 38 0.882 0.066 - 0.709 - 0.348 0.009 234 57.7 48 39 1.307 0.043 - 0.767 - 0.271 0.002 178 23.3 26 40 0.583 0.022 - 0.647 - 0.277 0.004 199 23.7 14 41 0.908 0.072 - 0.033 - 0.242 0.010 213 36.2 18 42 0.573 0.064 - 0.619 - 0.444 0.003 243 40.8 47 43 1.144 0.064 - 0.603 - 0.359 0.016 205 14.2 27 44 0.883 0.042 - 0.150 - 0.257 0.015 217 39.2 46 45 0.699 - 0.207 0.955 - 0.218 0.012 200 25.1 33 46 0.258 - 0.341 0.135 - 0.227 0.007 201 35.9 16 47 0.468 - 0.360 0.355 - 0.220 0.002 207 51.8 27 48 0.441 - 0.396 0.595 - 0.377 0.009 157 10.9 15 49 1.254 - 0.097 0.849 - 0.272 0.010 167 54.5 12 50 0.646 - 0.016 0.264 - 0.363 0.001 166 61.3 16 51 1.421 - 0.143 0.434 - 0.318 0.019 184 28.0 26 52 1.206 - 0.178 0.542 - 0.404 0.001 191 73.5 11 53 0.881 0.045 0.012 0.273 - 0.301 0.017 202 32.9 15 54 1.365 0.019 0.047 0.337 - 0.272 0.004 242 29.6 11 55 0.980 0.015 0.162 0.924 - 0.223 0.013 227 74.0 43 56 0.592 0.051 0.010 0.719 - 0.404 0.019 227 75.1 20 CS: Creep rupture strength at 650°C, 150 thousand hours; VE: Charpy impact absorption energy at 0°C after aging at 700°C for 10 thousand hours; TO: Oxidation scale thickness after 650°C, 10 thousand hour high-temperature oxidation (mass%) Invention steels No. C Si Mn P S Nb V Cr Mo 57 0.168 0.707 0.952 0.030 0.009 1.092 0.253 10.45 0.161 59 0.253 0.712 0.448 0.022 0.010 0.980 0.296 12.97 0.574 60 0.027 0.196 0.545 0.045 0.003 1.114 0.203 11.93 0.423 61 0.228 0.180 0.543 0.036 0.006 1.228 0.370 11.44 0.614 62 0.184 0.028 0.988 0.022 0.005 0.696 0.428 12.41 0.834 63 0.092 0.657 0.819 0.012 0.006 1.776 0.235 8.22 0.755 64 0.275 0.590 0.894 0.049 0.006 1.727 0.127 10.93 0.521 65 0.201 0.362 0.750 0.049 0.006 1.062 0.487 9.47 0.886 66 0.161 0.761 0.800 0.011 0.002 0.717 0.195 11.38 0.866 67 0.254 0.099 0.223 0.031 0.003 0.383 0.187 12.86 0.066 68 0.299 0.227 0.243 0.015 0.002 0.718 0.155 11.73 0.969 69 0.063 0.509 0.608 0.049 0.001 0.188 0.184 8.92 0.777 70 0.239 0.643 0.497 0.048 0.006 0.746 0.137 12.49 0.502 71 0.142 0.665 0.549 0.019 0.008 1.266 0.357 9.91 0.796 72 0.260 0.786 0.889 0.031 0.008 0.605 0.300 12.25 0.083 73 0.298 0.405 0.687 0.049 0.009 0.167 0.197 11.60 0.461 74 0.198 0.566 0.429 0.017 0.001 1.715 0.383 11.77 0.241 (Continued from Table 7) (mass%) Invention steels No. W Zr Ta Hf Ti N O CS MPa VE J TO µm 57 1.183 0.039 0.104 0.666 - 0.321 0.002 195 14.8 30 59 1.250 0.042 0.071 0.233 - 0.487 0.013 235 53.2 49 60 0.762 0.065 0.019 0.660 - 0.464 0.003 212 31.2 29 61 0.676 - - - 0.082 0.204 0.014 183 57.2 26 62 0.581 - - - 0.081 0.380 0.013 167 38.9 46 63 0.716 - - - 0.090 0.264 0.004 184 59.2 38 64 1.407 - - - 0.076 0.268 0.001 241 61.1 29 65 1.145 - - - 0.082 0.415 0.017 152 61.7 16 66 0.758 - - - 0.065 0.309 0.012 168 26.2 40 67 0.956 - - - 0.039 0.473 0.004 238 58.4 24 68 1.400 0.002 - - 0.069 0.349 0.006 190 60.0 41 69 1.017 0.024 - - 0.037 0.210 0.002 243 79.0 16 70 1.317 0.083 - - 0.052 0.466 0.016 198 55.2 34 71 1.405 0.051 - - 0.075 0.218 0.010 194 64.5 41 72 0.412 0.066 - - 0.023 0.482 0.009 237 28.1 14 73 0.479 0.004 - - 0.027 0.347 0.003 223 45.0 34 74 0.394 0.069 - - 0.016 0.396 0.001 162 79.7 14 CS: Creep rupture strength at 650°C, 150 thousand hours; VE: Charpy impact absorption energy at 0°C after aging at 700°C for 10 thousand hours; TO: Oxidation scale thickness after 650°C, 10 thousand hour high-temperature oxidation (mass%) Invention steels No. C Si Mn P S Nb V Cr Mo 75 0.019 0.256 0.274 0.023 0.004 0.876 0.266 11.27 0.616 76 0.061 0.732 0.450 0.038 0.008 1.892 0.357 11.79 0.383 77 0.282 0.275 0.824 0.036 0.007 1.357 0.470 10.74 0.763 78 0.212 0.148 0.595 0.044 0.007 0.453 0.098 8.14 0.897 79 0.099 0.043 0.977 0.015 0.006 0.226 0.380 8.02 0.757 80 0.237 0.048 0.922 0.017 0.004 1.291 0.312 11.17 0.856 81 0.167 0.070 0.793 0.043 0.004 0.482 0.135 11.37 0.489 82 0.123 0.680 0.262 0.021 0.003 1.753 0.495 12.72 0.084 83 0.156 0.085 0.278 0.020 0.009 1.696 0.356 11.17 0.827 84 0.038 0.770 0.994 0.047 0.002 0.663 0.475 10.15 0.416 85 0.108 0.533 0.552 0.029 0.009 1.814 0.485 10.58 0.526 86 0.255 0.467 0.614 0.024 0.003 0.392 0.272 10.10 0.910 87 0.177 0.373 0.265 0.023 0.004 1.667 0.476 8.30 0.017 88 0.156 0.164 0.520 0.029 0.001 1.393 0.118 9.17 0.815 89 0.030 0.730 0.614 0.030 0.005 1.490 0.395 11.15 0.877 90 0.239 0.382 0.759 0.046 0.009 1.912 0.116 9.62 0.474 91 0.111 0.044 0.701 0.026 0.007 1.124 0.489 10.50 0.281 92 0.283 0.645 0.990 0.036 0.001 0.417 0.250 12.20 0.493 93 0.226 0.762 0.575 0.024 0.005 0.131 0.157 10.70 0.298 (Continued from Table 9) (mass%) Invention steels No. W Zr Ta Hf Ti N O CS MPa VE J TO µm 75 0.381 - 0.089 - 0.094 0.310 0.017 170 39.6 40 76 0.263 - 0.646 - 0.046 0.218 0.012 221 68.1 43 77 1.425 - 0.213 - 0.019 0.273 0.005 227 25.4 23 78 0.904 - 0.941 - 0.034 0.213 0.019 216 19.2 15 79 1.275 - 0.589 - 0.060 0.268 0.006 189 17.4 20 80 1.104 - 0.440 - 0.031 0.210 0.015 155 77.8 34 81 1.060 - 0.545 - 0.056 0.450 0.016 236 62.2 38 82 0.271 0.039 0.030 - 0.024 0.298 0.004 225 67.8 45 83 0.296 0.052 0.119 - 0.043 0.234 0.010 192 38.7 20 84 0.729 0.085 0.011 - 0.084 0.334 0.019 202 57.0 13 85 1.148 0.009 0.066 - 0.073 0.466 0.016 176 31.0 29 86 0.921 0.007 0.023 - 0.086 0.432 0.009 219 32.3 20 87 1.171 0.032 0.149 - 0.086 0.262 0.011 187 67.8 28 88 0.364 0.049 0.199 - 0.024 0.232 0.010 213 76.1 39 89 1.085 - - 0.518 0.085 0.321 0.002 158 40.8 35 90 0.507 - - 0.911 0.096 0.245 0.010 168 31.1 16 91 0.813 - - 0.693 0.072 0.482 0.010 161 61.2 49 92 0.958 - - 0.058 0.032 0.307 0.015 204 73.3 40 93 1.176 - - 0.171 0.035 0.389 0.018 217 45.3 39 CS: Creep rupture strength at 650°C, 150 thousand hours; VE: Charpy impact absorption energy at 0°C after aging at 700°C for 10 thousand hours; TO: Oxidation scale thickness after 650°C, 10 thousand hour high-temperature oxidation (mass%) Invention steels No. C Si Mn P S Nb V Cr Mo 94 0.283 0.545 0.806 0.037 0.009 0.633 0.067 11.89 0.372 95 0.079 0.145 0.236 0.044 0.003 0.416 0.070 11.77 0.918 96 0.112 0.546 0.487 0.046 0.001 1.964 0.230 11.21 0.157 97 0.168 0.078 0.806 0.016 0.007 0.672 0.102 10.14 0.713 98 0.274 0.484 0.754 0.035 0.005 1.318 0.165 12.68 0.954 99 0.180 0.543 0.766 0.042 0.006 1.076 0.457 11.88 0.494 100 0.258 0.024 0.373 0.023 0.007 1.989 0.391 9.73 0.029 101 0.130 0.645 0.536 0.044 0.006 0.611 0.278 12.22 0.208 102 0.277 0.191 0.985 0.033 0.009 0.139 0.145 10.17 0.214 103 0.181 0.384 0.681 0.017 0.004 1.865 0.224 12.09 0.745 104 0.044 0.160 0.419 0.022 0.003 0.142 0.082 11.61 0.652 105 0.116 0.041 0.761 0.018 0.003 1.168 0.072 11.08 0.397 106 0.015 0.763 0.554 0.019 0.001 1.115 0.076 8.29 0.928 107 0.231 0.128 0.741 0.033 0.004 1.269 0.393 9.13 0.422 108 0.237 0.626 0.679 0.028 0.005 1.969 0.434 9.59 0.072 109 0.135 0.148 0.803 0.030 0.007 1.742 0.254 8.08 0.083 110 0.279 0.092 0.512 0.044 0.007 0.434 0.209 8.66 0.454 111 0.237 0.466 0.610 0.049 0.009 1.793 0.296 11.37 0.184 (Continued from Table 11) (mass%) Invention steels No. W Zr Ta Hf Ti N O CS MPa VE J TO µm 94 1.059 - - 0.733 0.010 0.207 0.004 219 60.8 11 95 0.872 - - 0.945 0.019 0.237 0.014 229 63.5 46 96 1.198 - - 0.523 0.070 0.432 0.018 246 10.9 50 97 0.892 - - 0.752 0.055 0.346 0.001 190 51.7 34 98 0.490 0.011 - 0.559 0.099 0.453 0.013 193 77.2 49 99 1.317 0.029 - 0.965 0.049 0.263 0.011 230 73.7 26 100 1.413 0.081 - 0.038 0.034 0.482 0.005 175 57.5 24 101 1.023 0.035 - 0.412 0.068 0.392 0.007 195 75.0 32 102 0.370 0.012 - 0.798 0.080 0.257 0.012 211 70.9 16 103 1.018 0.078 - 0.887 0.040 0.360 0.015 240 22.6 25 104 1.318 0.055 - 0.912 0.010 0.373 0.018 214 55.7 24 105 1.461 - 0.893 0.227 0.099 0.433 0.005 197 59.3 30 106 1.291 - 0.389 0.493 0.025 0.233 0.005 219 36.6 31 107 1.031 - 0.277 0.404 0.029 0.442 0.010 199 70.7 14 108 0.928 - 0.128 0.139 0.039 0.339 0.013 155 69.5 45 109 0.875 - 0.348 0.329 0.098 0.297 0.017 221 49.5 27 110 1.142 - 0.247 0.075 0.020 0.494 0.004 217 10.9 15 111 0.401 - 0.372 0.357 0.076 0.337 0.005 242 58.3 10 CS: Creep rupture strength at 650°C, 150 thousand hours; VE: Charpy impact absorption energy at 0°C after aging at 700°C for 10 thousand hours; TO: Oxidation scale thickness after 650°C, 10 thousand hour high-temperature oxidation (mass%) Invention steels No. C Si Mn P S Nb V Cr Mo 112 0.157 0.497 0.978 0.037 0.010 0.234 0.321 12.72 0.652 113 0.154 0.087 0.687 0.026 0.008 1.123 0.289 8.62 0.764 114 0.144 0.189 0.303 0.023 0.001 1.021 0.072 10.12 0.231 115 0.176 0.143 0.360 0.015 0.009 0.282 0.193 8.02 0.213 116 0.159 0.104 0.608 0.010 0.002 1.169 0.288 10.81 0.401 117 0.159 0.680 0.631 0.032 0.008 0.675 0.143 10.13 0.649 118 0.152 0.556 0.529 0.035 0.004 1.745 0.275 9.10 0.748 (Continued from Table 13) (mass%) Invention steels No. W Zr Ta Hf Ti N O CS MPa VE J TO µm 112 0.429 0.091 0.090 0.113 0.055 0.338 0.002 206 37.3 16 113 0.610 0.090 0.715 0.686 0.085 0.421 0.017 197 51.9 12 114 0.554 0.012 0.131 0.772 0.091 0.269 0.017 183 43.9 24 115 0.478 0.071 0.497 0.062 0.047 0.350 0.014 231 26.6 40 116 1.283 0.091 0.492 0.726 0.013 0.340 0.003 212 13.8 29 117 0.833 0.091 0.089 0.591 0.015 0.339 0.009 192 24.2 31 118 0.668 0.038 0.034 0.026 0.020 0.218 0.008 218 54.3 10 CS: Creep rupture strength at 650°C, 150 thousand hours; VE: Charpy impact absorption energy at 0°C after aging at 700°C for 10 thousand hours; TO: Oxidation scale thickness after 650°C, 10 thousand hour high-temperature oxidation Comparison steels (mass%) No. C Si Mn P S Nb V Cr Mo W 161 0.02 0.06 0.45 0.014 0.008 0.051 0.20 8.53 0.592 0.87 162 0.07 0.07 0.57 0.012 0.004 0.066 0.22 8.99 0.424 0.88 163 0.23 0.10 0.52 0.016 0.004 0.074 0.18 9.05 0.560 0.86 164 0.15 0.13 0.45 0.011 0.002 0.033 0.19 9.23 0.550 0.77 165 0.29 0.12 0.98 0.008 0.001 0.127 0.08 12.1 0.320 0.11 166 0.15 0.09 0.24 0.009 0.001 0.290 0.45 12.7 0.007 1.80 167 0.18 0.09 0.44 0.009 0.001 0.014 0.31 11.4 0.989 0.99 168 0.11 0.12 0.74 0.004 0.001 2.880 0.22 11.8 0.679 0.67 169 0.09 0.21 0.81 0.044 0.002 0.321 0.20 10.5 0.814 0.23 170 0.10 0.23 0.99 0.024 0.001 0.545 0.19 10.0 0.333 0.77 171 0.22 0.21 0.11 0.012 0.001 1.227 0.40 9.76 0.545 1.21 172 0.07 0.27 0.07 0.013 0.009 1.621 0.49 9.00 0.512 1.09 173 0.23 0.23 0.09 0.002 0.010 1.998 0.07 8.86 0.533 1.15 (Continued from Table 15) Comparison steels (mass%) No. Zr Ta Hf Ti N O CS MPa VE J TO (µm) 161 0.007 0.65 - 0.034 0.072 0.007 120 70 760 162 0.008 0.77 - 0.031 0.081 0.009 107 28 660 163 0.002 0.78 - 0.044 0.872 0.007 205 6 50 164 0.012 0.71 0.66 0.100 0.525 0.002 185 3 35 165 0.011 0.76 0.87 0.010 0.164 0.002 65 95 35 166 - - 0.81 - 0.128 0.002 90 2 20 167 - - 0.59 - 0.154 0.001 70 60 25 168 - - - 0.060 0.332 0.002 116 4 5 169 0.145 0.89 - 0.071 0.425 0.002 187 4 25 170 - 1.21 - 0.032 0.202 0.002 153 7 40 171 0.011 0.32 1.13 - 0.191 0.008 220 7 45 172 0.540 0.05 0.22 0.29 0.103 0.012 210 8 15 173 0.880 - 0.10 - 0.200 0.006 24 2 30 CS: Creep rupture strength at 650°C, 150 thousand hours; VE: Charpy impact absorption energy at 0 °C after aging at 700°C for 10 thousand hours; TO: Oxidation scale thickness after 650 °C, 10 thousand hour high-temperature oxidation

Claims: A high-nitrogen ferritic heat-resisting steel with high Nb content comprising, in weight per cent 0.01 - 0.30% C, 0.02 - 0.80% Si, 0.20 - 1.00% Mn, 8.00 - 13.00% Cr, 0.005 - 1.00% Mo, 0.20 - 1.50% W, 0.05 - 1.00% V, over 0.12 up to 2.00% Nb, and > 0.20 - 0.50% N, and being controlled to include not more than 0.050% P, not more than 0.010% S, and not more than 0.020% O, the remainder being Fe and unavoidable inpurities, wherein the N content of the steel of said chemical composition exceeds the N concentration in the steel equilibrated with 1 bar N gas during casting. A high-nitrogen ferritic heat-resisting steel with high Nb content comprising, in weight per cent 0.01 - 0.30% C, 0.02 - 0.80% Si, 0.20 - 1.00% Mn, 8.00 - 13.00% Cr, 0.005 - 1.00% Mo, 0.20 - 1.50% W, 0.05 - 1.00% V, over 0.12 up to 2.00% Nb, and > 0.20 - 0.50% N, and one or both of 0.01 - 1.00% Ta and 0.01 - 1.00% Hf and being controlled to include not more than 0.050% P, not more than 0.010% S, and not more than 0.020% O, the remainder being Fe and unavoidable inpurities, wherein the N content of the steel of said chemical composition exceeds the N concentration in the steel equilibrated with 1 bar N gas during casting. A high-nitrogen ferritic heat-resisting steel with high Nb content comprising, in weight per cent 0.01 - 0.30% C, 0.02 - 0.80% Si, 0.20 - 1.00% Mn, 8.00 - 13.00% Cr, 0.005 - 1.00% Mo, 0.20 - 1.50% W, 0.05 - 1.00% V, over 0.12 up to 2.00% Nb, and > 0.20 - 0.50% N, and one or both of 0.0005 - 0.10% Zr and 0.01 - 0.10% Ti and being controlled to include not more than 0.050% P, not more than 0.010% S, and not more than 0.020% O, the remainder being Fe and unavoidable inpurities, wherein the N content of the steel of said chemical composition exceeds the N concentration in the steel equilibrated with 1 bar N gas during casting. A high-nitrogen ferritic heat-resisting steel with high Nb content comprising, in weight per cent 0.01 - 0.30% C, 0.02 - 0.80% Si, 0.20 - 1.00% Mn, 8.00 - 13.00% Cr, 0.005 - 1.00% Mo, 0.20 - 1.50% W, 0.05 - 1.00% V, over 0.12 up to 2.00% Nb, and > 0.20 - 0.50% N, one or both of 0.01 - 1.00% Ta and 0.01 - 1.00% Hf and one or both of 0.0005 - 0.10% Zr and 0.01 - 0.10% Ti and being controlled to include not more than 0.050% P, not more than 0.010% S, and not more than 0.020% O, the remainder being Fe and unavoidable inpurities, wherein the N content of the steel of said chemical composition exceeds the N concentration in the steel equilibrated with 1 bar N gas during casting. A method of producing a high-nitrogen ferritic heat-resisting steel with high Nb content having a composition according to any of claims 1 - 4, wherein the steel is melted and equilibrated in an atmosphere of a mixed gas of a prescribed nitrogen partial pressure or nitrogen gas and is thereafter cast or solidified in an atmosphere controlled to have a total pressure of not less than 2.5 bar and a nitrogen partial pressure of not less than 1.0 bar, with the relationship between the nitrogen partial pressure p and the total pressure P being P > 2.5p.

Inventors: HASEGAWA YASUSHI, MASUYAMA FUJIMITSU, NAOI HISASHI, OHGAMI MASAHIRO, HASEGAWA, YASUSHI, MASUYAMA, FUJIMITSU, NAOI, HISASHI, OHGAMI, MASAHIRO, Hasegawa, Yasushi, c/o R &amp; D Laboratories, Masuyama, Fujimitsu, c/o Nagasaki R. &amp; Dev. Center, Naoi, Hisashi, c/o R &amp; D Laboratories, Ohgami, Masahiro, c/o R &amp; D Laboratories
Applicants: MITSUBISHI HEAVY IND LTD, NIPPON STEEL CORP, MITSUBISHI JUKOGYO KABUSHIKI KAISHA, NIPPON STEEL CORPORATION
Application year: 1992
Main classification: C22C38
IPC codes: C21C7, B22D23, C22C38