High nickel alloy sheet

Abstract

This high nickel alloy sheet comprises, in mass%, C: 0.020-0.100%, Si: 0.05-1.00%, Mn: 0.05-2.00%, P: 0.035% or less, S: 0.0015% or less, Cr: 18.0-25.0%, Ni: 18.0-50.0%, Al: 0.05-1.00%, Ti: 0.15-1.50%, N: 0.020% or less, O: 0.0030% or less, and B: 0.0003-0.0030%, with the remainder being Fe and unavoidable impurities. The high nickel alloy sheet has: an average crystal particle size of 60 μm or more; an average TiC size of 1 μm or more; a TiC density of 4000 pcs / mm2 or less; a Ti solid solution ratio (solid-solution Ti content / total Ti content) (mass ratio) of 0.75 or more; and a sheet thickness of 0.2-4 mm.

Description

High Ni alloy plate

[0001] The present invention relates to a high-Ni alloy plate having excellent creep properties and high-temperature strength.

[0002] Alloy 800H (ASTM N08010, N08011) is a typical commercial alloy of high-Ni alloy steel used for heat-resistant applications. In recent years, demand has expanded in developing countries, necessitating technological development to enable the supply of products with low cost, good surface quality, and excellent usability. High-Ni alloy steel is primarily used in high-temperature reactors at chemical plants, where it is often used at temperatures of 600°C or higher, and also under high pressure to improve the efficiency of chemical reactions. In these applications, the higher the creep strength, the thinner the wall thickness that can be used.

[0003] Patent Documents 1 and 2 disclose inventions that improve creep properties by optimizing the size and number of Nb and Ti nitrides or carbides in a hot-rolled and annealed high-Ni alloy having a predetermined composition. In Patent Document 1, the cooling rate during production of a steel ingot is controlled to control the secondary dendrite arm spacing, thereby optimizing the size and number of Nb and Ti nitrides or carbides. In an example of Patent Document 2, the annealing temperature after hot rolling is set to 1180°C to 1300°C.

[0004] Non-Patent Document 1 investigates the effect of grain size on the high-temperature creep properties of Hastelloy X, a type of high-Ni alloy. Using samples with grain sizes of 37 to 1220 μm, the time to rupture in high-temperature creep tests was evaluated. The results reported indicate that creep properties improve as the grain size increases from approximately 37 μm to 100 μm, reaching an optimum at a grain size of approximately 100 μm, and gradually deteriorating as the grain size increases further. Non-Patent Document 2 describes the creep properties of Alloy 800H.

[0005] Patent Document 3 relates to an Al- and Ti-containing high Ni-based alloy with excellent resistance to weld hot cracking, and solves the problem by specifying the number density of TiC-based precipitates. Patent Document 4 relates to an Fe—Cr—Ni alloy with excellent surface properties, and solves the problem by specifying the number density of TiN inclusions.

[0006] Japanese Patent Application Laid-Open No. 2017-57461 Japanese Patent No. 7174192 Japanese Patent Application Laid-Open No. 2022-163585 Japanese Patent Application Laid-Open No. 2018-59148

[0007] Yoshihiro Kondo et al., "Effect of Grain Size on the High-Temperature Creep Properties of Hastelloy X," Iron and Steel, 67th year (1981), No. 10, pp. 1805-1814 Katsumi Tachibana et al., "Creep Properties of Alloy 800H," Japan Atomic Energy Research Institute, JAERI-Tech 98-010, March 1998

[0008] High Ni alloy steel is not always provided as a thick plate, but thinner plates are sometimes required depending on the application. Hitherto, no high Ni alloy plate has been realized that satisfies both high levels of creep properties and high-temperature strength.

[0009] An object of the present invention is to provide a high Ni alloy plate having a predetermined component composition, such as Alloy 800H, which has excellent creep properties and high-temperature strength.

[0010] That is, the gist of the present invention is as follows: [1] A steel sheet containing, by mass%, C: 0.020 to 0.100%, Si: 0.05 to 1.00%, Mn: 0.05 to 2.00%, P: 0.035% or less, S: 0.0015% or less, Cr: 18.0 to 25.0%, Ni: 18.0 to 50.0%, Al: 0.05 to 1.00%, Ti: 0.15 to 1.50%, N: 0.020% or less, O: 0.0030% or less, B: 0.0003 to 0.0030%, and the balance being Fe and impurities, wherein the average crystal grain size is 60 μm or more, the average size of TiC is 1 μm or more, and the density of TiC is 4000 particles / mm ２[2] The high Ni alloy plate according to [1], further comprising, in place of a portion of the Fe, one or more of the following, by mass: Ca: 0.0003 to 0.0070%, Mg: 0.0060% or less, Mo: 5.00% or less, W: 2.00% or less, Cu: 3.00% or less, Co: 2.00% or less, V: 1.00% or less, Nb: 1.00% or less, Ta: 1.00% or less, Sn: 0.10% or less, and REM: 0.10% or less.

[0011] The high Ni alloy plate of the present invention has a TiC density of 4000 pieces / mm ２ The average grain size after final annealing is set to 60 μm or more, thereby realizing good creep properties. Furthermore, by increasing the proportion of solute Ti in the contained Ti, fine TiC precipitates when the sheet is used in a high-temperature environment, and high high-temperature strength can be achieved.

[0012] The present invention is directed to a high Ni alloy plate having a thickness of 4 mm or less. In the production of the high Ni alloy plate, a method can be used in which a slab produced in a steelmaking process is subjected to hot working such as continuous hot rolling, and the resulting plate is annealed, or further cold rolled and annealed.

[0013] The component composition of the high Ni alloy plate of the present invention will be explained. % means mass %. In the component content range, the expression "A to B %" means "A % or more, B % or less."

[0014] <<C: 0.020 to 0.100%>> To ensure the high-temperature strength targeted by the present invention, C is set to 0.020% or more, preferably 0.030% or more. On the other hand, if the C content exceeds 0.100%, the pinning effect of precipitated carbides inhibits the grain size from increasing, deteriorating creep properties, and the precipitation of Cr carbides also deteriorates high-temperature strength. Therefore, C is set to 0.100% or less, preferably 0.080% or less.

[0015] <Si: 0.05 to 1.00%> Si is added in an amount of 0.05% or more, preferably 0.20% or more, to improve deoxidation and oxidation resistance. However, because Si is also an element that lowers the melting point of steel, adding more than 1.00% of Si reduces hot ductility at temperatures around 1200°C and the solidification cracking susceptibility and liquation cracking susceptibility during welding. In addition, intermetallic compounds are more likely to precipitate, deteriorating high-temperature properties. Therefore, the upper limit is set to 1.00%. A preferred upper limit is 0.80%.

[0016] <Mn: 0.05-2.00%> Mn has the effect of increasing the stability of the austenite phase and improving heat resistance. For this reason, it is preferable to actively add Mn to the alloy of the present invention. To improve heat resistance, 0.05% or more, preferably 0.20% or more, and more preferably 0.30% or more is added. However, adding more than 2.00% of Mn makes intermetallic compounds more likely to precipitate, deteriorating heat resistance and adversely affecting solidification cracking susceptibility. Therefore, the upper limit is set at 2.00%. A preferred upper limit is 1.50%, and a more preferred upper limit is 1.30%.

[0017] <<P: 0.035% or less>> P is an element that is inevitably mixed in from the raw materials, and has the effect of increasing solidification cracking susceptibility, so its content is limited to 0.035% or less, preferably 0.030% or less.

[0018] <<S: 0.0015% or less>> S is an element that is inevitably mixed in from raw materials and deteriorates hot workability and oxidation resistance, so the content is limited to 0.0015% or less, preferably 0.0010% or less. Although the S content can be reduced by refining, an extreme reduction in the S content increases costs. For this reason, it is preferable to set the lower limit of the S content to 0.0001%.

[0019] <<Cr: 18.0 to 25.0%>> Cr is an essential element responsible for the oxidation resistance and high-temperature corrosion resistance of a heat-resistant alloy as a high-temperature material, and is contained in an amount of 18.0% or more. On the other hand, if the content exceeds 25.0%, the high-temperature structural stability decreases even if a large amount of Ni is contained, and intermetallic compounds begin to precipitate, deteriorating the heat resistance characteristics. Therefore, the upper limit is set to 25.0%.

[0020] <Ni: 18.0 to 50.0%> Ni stabilizes the austenitic structure at high temperatures and improves corrosion resistance and toughness against various acids, so the Ni content is set to 18.0% or more, preferably 20.0% or more. However, Ni is an expensive alloy, and in the steel of the present invention, from the viewpoint of cost, the upper limit is set to 50.0% or less, preferably 48.0% or less.

[0021] <Al: 0.05 to 1.00%> Al is a deoxidizing element and also has the effect of increasing high-temperature strength in high-Ni alloys. In the present invention, in order to increase high-temperature strength, a content of 0.05% or more, preferably 0.10% or more, is required. On the other hand, if the Al content exceeds 1.00%, the high-temperature strength decreases due to intermetallic compounds. For this reason, the upper limit of the Al content is set at 1.00%. The preferred upper limit is 0.80%.

[0022] <<Ti: 0.15 to 1.50%>> Ti achieves high-temperature strength through precipitation strengthening, so the content is set to 0.15% or more. On the other hand, if the Ti content is excessive, the pinning effect of precipitates refines the grain size and reduces creep properties, so the upper limit is set to 1.50%.

[0023] <<N: 0.020% or Less>> If the N content exceeds 0.020%, Ti precipitates as TiN, the precipitation strengthening by TiC is not sufficiently obtained, and the high-temperature strength decreases. Therefore, the upper limit of the N content is set to 0.020%, preferably 0.010%.

[0024] <O: 0.0030% or Less> Oxygen (O) forms oxide-based inclusions with Ca, Mg, Al, and Ti in the alloy of the present invention. The oxygen content corresponds to the total amount of oxide-based inclusions and is an important indicator of the deoxidation state of the alloy. If the oxygen content exceeds 0.0030%, the desired deoxidation equilibrium is not satisfied and nozzle clogging during continuous casting is more likely to occur. Therefore, the upper limit of the oxygen content is set to 0.0030%. A preferred upper limit is 0.0025%, and more preferably 0.0020%. On the other hand, reducing the oxygen content reduces oxide-based inclusions and coarse TiC-based inclusions, which is advantageous in suppressing nozzle clogging and weld hot cracking, but it also generates excess Ca and Mg in the alloy, which can cause a decrease in hot workability. For this reason, the oxygen content is preferably 0.0003% or more.

[0025] <B: 0.0003-0.0030%> B improves high-temperature creep strength and is actively added, especially in applications where the steel will be used in high-temperature environments. While the mechanism by which B improves creep strength is unclear, it is said that segregation at grain boundaries increases grain boundary strength. Since the improvement in hot tensile strength due to B content is apparent at 0.0003% or more, the lower limit is set at 0.0003%. However, because B is an element that lowers the melting point of steel and is prone to segregation, excessive addition promotes solidification cracking and liquation cracking, significantly adversely affecting hot workability, particularly in the I-region embrittlement region (near 1200°C). Therefore, the upper limit for B content is set at 0.0030%.

[0026] The composition of the high Ni alloy sheet of the present invention contains the above-mentioned components, with the balance being Fe and impurities. Furthermore, instead of a part of the Fe, it may selectively contain one or more of the following components (mass %):

[0027] <<Ca: 0.0003-0.0070%>> Ca is an important element for improving the hot workability of the alloy. It is added to fix the S in the alloy as CaS, improving the hot workability. The reaction is as follows: Ca combines with oxygen in the alloy to form CaO, CaO-Al ２ O ３After the dissolved oxygen (free oxygen) in the alloy is reduced to almost zero, the remaining Ca reacts with S in the alloy to form CaS. On the other hand, excessive Ca addition reduces ductility at high temperatures around 1100°C. For this reason, the upper limit of the Ca content is set at 0.0070%.

[0028] <<Mg 0.0060% or Less>> Mg is an element that exhibits a desulfurization effect, and therefore, even if added in small amounts, it is an element that can improve the hot workability of the alloy, but if added in excess, it significantly reduces the hot workability at around 900° C. Therefore, in the present invention, when Mg is added, the upper limit of the Mg content is set to 0.0060%.

[0029] <<Mo: 5.00% or Less>> Mo is an element that enhances the high-temperature strength and high-temperature corrosion resistance of the alloy, and can be added as needed to improve these properties. This effect can be achieved when Mo is contained in an amount of 0.20% or more. However, Mo is an expensive element, and in the steel of the present invention, the upper limit of the Mo content is set to 5.00% from the viewpoint of reducing the alloy cost of the steel.

[0030] <W: 2.00% or less> W is an element that increases high-temperature strength and high-temperature corrosion resistance, and in particular significantly improves creep strength. Therefore, W can be added as needed. However, excessive addition of W reduces hot workability and increases alloy costs, so the upper limit is set at 2.00%.

[0031] Cu: 3.00% or less Cu is an element that enhances the alloy's acid corrosion resistance and dew-point corrosion resistance, which is often a problem in high-temperature equipment, and also has the effect of improving high-temperature strength and structural stability, so it can be added as needed. This effect can be achieved when Cu is contained in an amount of 0.10% or more. On the other hand, if the Cu content exceeds 3.00%, embrittlement occurs during solidification, so the upper limit is set to 3.00%, preferably 2.00%.

[0032] <<Co: 2.00% or Less>> Co is an element effective for improving the high-temperature structural stability and corrosion resistance of the alloy, and is contained as needed to improve these properties. This effect can be achieved when Co is contained in an amount of 0.10% or more. Since Co is an expensive element, if it is contained in an amount exceeding 2.00%, the effect commensurate with the cost will not be achieved, so the upper limit is set at 2.00%. It is preferable to set the upper limit at 1.00%.

[0033] <<V: 1.00% or Less>> When added in an amount of 0.03% or more, V has the effect of improving the high-temperature properties of the alloy through solid solution strengthening or precipitation strengthening. On the other hand, addition of more than 1.00% reduces solidification cracking susceptibility, so the upper limit is set at 1.00%.

[0034] <Nb: 1.00% or less> Nb not only has a high solid solution strengthening ability, but also acts to improve the high-temperature strength of steel, particularly high-temperature creep strength, by finely precipitating carbides and nitrides at service temperatures. Therefore, Nb can be added as needed. This effect can be achieved when the Nb content is 0.03% or more. On the other hand, Nb is also an element that is prone to grain boundary segregation, and excessive addition adversely affects hot workability, so the upper limit in this chemical composition is set to 1.00%. The preferred upper limit is 0.80%.

[0035] <<Ta: 1.00% or Less>> Ta has the effect of improving the high-temperature strength of steel, particularly the high-temperature creep strength, by finely precipitating carbides and nitrides at the temperature of use, and can be added as needed. Note that this effect can be achieved when Ta is contained in an amount of 0.03% or more. Since excessive addition of Ta adversely affects hot workability, the upper limit in this chemical composition is set to 1.00%. It is preferable to set the upper limit to 0.80%.

[0036] <<Sn: 0.10% or less>> Sn is an element that improves the corrosion resistance and high-temperature creep strength of steel, and can be added as needed. This effect can be achieved when the Sn content is 0.01% or more. However, since the addition of more than 0.10% reduces hot workability, the upper limit is set at 0.10%. The preferred upper limit is 0.06%.

[0037] REM: 0.10% or less REM is an element that fixes P and S and improves the oxidation resistance and hot workability of steel, so it can be added as needed. This effect can be achieved when the REM content is 0.01% or more. However, adding more than 0.10% in total generates coarse oxides and nitrides, significantly impairing manufacturability by causing nozzle clogging during refining and increasing surface defects. Therefore, the upper limit of the REM content is set at 0.10%. A preferred upper limit is 0.08%. Here, REM is a collective term for 17 elements, including two elements, Sc and Y, and 15 elements (lanthanides) from La to Lu. REM can be added as individual metals or alloys, or as misch metals.

[0038] The mechanism by which the high Ni alloy plate of the present invention achieves both good creep properties and high high-temperature strength, and the crystalline structure of the high Ni alloy plate of the present invention will be described below.

[0039] <<Mechanism for achieving both good creep properties and high high-temperature strength>> As described in Non-Patent Document 1, in high-Ni alloys, the creep properties improve as the crystal grain size increases within a crystal grain size range of approximately 100 μm. In the present invention, it has been found that the desired creep properties can be obtained by setting the average crystal grain size to 60 μm or more. In the manufacturing process of high-Ni alloy sheet, if a cold-rolling annealing step is not performed after the hot rolling step, the crystal grains are grown in the final annealing step after hot rolling, or in the final annealing step after cold rolling, if the hot rolling step and the cold-rolling annealing step are performed. The average crystal grain size of the final product is set to 60 μm or more.

[0040] On the other hand, the high Ni alloy sheet of the present invention contains Ti and C as its component composition to ensure high-temperature strength. For high Ni alloys containing Ti and C, TiC precipitates during the hot rolling process of sheet production, which involves a hot rolling process and a cold rolling annealing process. TiC also precipitates during heating in the final annealing process after cold rolling. In the final annealing process after cold rolling, the cold-rolled sheet has a large amount of cold-rolling strain, which delays coarsening of the grain size. Furthermore, since there are many TiC precipitates precipitated during the hot rolling process, the pinning effect of the precipitates further delays coarsening of the grain size. Therefore, it has been difficult to achieve an average grain size of 60 μm or more during the final annealing after cold rolling. When attempting to increase the grain size by setting the final annealing temperature to an extremely high temperature above 1200°C, the sheet elongation that occurs during annealing makes passing the sheet itself difficult, making it unpractical.

[0041] In the present invention, it has been found that by optimizing the hot rolling conditions and the temperature rise conditions in the final annealing after hot rolling or cold rolling, the precipitate density in the steel at the final annealing stage can be reduced, thereby suppressing the pinning effect caused by the precipitates and promoting crystal growth, and that the crystal grain size after the completion of the final annealing can be made 60 μm or more, thereby realizing good creep properties. Furthermore, it has been found that by increasing the ratio of the amount of solute Ti to the amount of Ti contained in the steel after the final annealing, fine TiC precipitates when the sheet is used in a high-temperature environment, and high high-temperature strength can be achieved.

[0042] <<Average grain size of 60 μm or more>> By setting the average grain size of the high Ni alloy sheet to 60 μm or more, excellent high-temperature creep properties can be realized. The average grain size can be determined by observing the center of the sheet thickness of a cross section cut by a plane parallel to the rolling direction and perpendicular to the rolling width direction of the final annealed steel sheet with an optical microscope, determining the grain size number by the cutting method specified in JIS G0551:2013, and using the conversion table or graph described in JIS.

[0043] The average size of TiC is 1 μm or more, and the density of TiC is 4000 pieces / mm ２ TiC density of high Ni alloy plate is 4000 pieces / mm ２If the density is less than or equal to 100 μm, the average grain size can be made 60 μm or more by final annealing. By reducing the precipitate density during final annealing, the pinning effect of the precipitates can be suppressed, and the average grain size can be made 60 μm or more. To evaluate the TiC density, a cross section of the final annealed steel sheet cut along a plane parallel to the rolling direction and perpendicular to the rolling width direction is etched by the SPEED method, and then observed in five fields of view at a magnification of 2000 times using an electron microscope, and the number of precipitates identified as TiC by EDS is divided by the observed area to calculate the TiC density.

[0044] By making the average size of TiC in the high Ni alloy plate 1 μm or more, the density of TiC can be increased to 4000 pieces / mm ２ Regarding the evaluation of the average size of TiC, a cross section of the final annealed steel sheet cut along a plane parallel to the rolling direction and perpendicular to the rolling width direction is etched by the SPEED method, and then observed under an electron microscope at a magnification of 2000 times in five fields, and the average value of the major axis and minor axis of precipitates identified as TiC by EDS is taken as the size of the precipitates, and the average size of TiC is calculated as the average value of the sizes of the precipitates.

[0045] <<Solution ratio of Ti (Solution amount of Ti / Content of Ti) (mass ratio) is 0.75 or more>> The high Ni alloy sheet of the present invention achieves a predetermined tensile strength (high-temperature strength) at a high temperature of 800°C. If the solid solution ratio of Ti (Solution amount of Ti / Content of Ti) (mass ratio) of the high Ni alloy sheet is 0.75 or more, the solute Ti is sufficiently contained in the thin sheet, and when used in a high-temperature environment of 800°C, the solute Ti precipitates as fine precipitates in the steel, thereby achieving the required high-temperature strength. The solution ratio of Ti is evaluated as follows. Steel sheets cut into pieces approximately 30 mm square are wet-polished with #600 polishing, and then 1 g of the stainless steel base material is dissolved by electrolysis at a constant potential of -100 mV in a methanol solution of 10% maleic anhydride and 2% tetramethylammonium chloride. The remaining undissolved precipitates were captured using a 200 μm mesh filter, washed with pure water, dried, and then dissolved for elemental analysis using ICP. The amount of precipitated Ti was calculated by dividing the obtained amount of Ti by the mass change of the sample due to electrolysis. The amount of dissolved Ti was determined by subtracting the amount of precipitated Ti from the amount of Ti contained in the steel sheet, and the ratio of the amount of dissolved Ti to the amount of contained Ti (mass ratio) was calculated.

[0046] <<Thickness of 0.2 to 4 mm>> The thickness of the high Ni alloy plate is set to 0.2 to 4 mm. The reason for setting the thickness to 0.2 mm or more is to avoid the risk of holes due to oxidation, taking into consideration the use at high temperatures as a heat-resistant material. Since the present invention aims to achieve a predetermined quality in a high Ni alloy plate with a thickness of 4 mm or less, the thickness is limited to 4 mm or less. Furthermore, if the thickness is 4 mm or less, the heating rate required in the final annealing can be achieved.

[0047] <<Properties that High Ni Alloy Sheets Should Have>> Creep properties are evaluated by the 800°C 80 MPa creep life (hr). The creep test is a constant load test in accordance with JIS Z 2271, and plate-shaped test pieces with a parallel portion 10 mm wide and 35 mm long, with the rolling direction as the longitudinal direction, are prepared from the final annealed steel sheet and used for the test. The test conditions are 800°C and an initial stress of 80 MPa, and the time until fracture is measured. An 800°C 80 MPa creep life of 50 hr or more is considered good. A creep life of 70 hr or more is more preferable.

[0048] The high-temperature strength is evaluated by the 800 ° C tensile strength (MPa). A plate-shaped high-temperature tensile test piece (parallel portion width: 10.5 mm, parallel portion length: 35 mm) with the rolling direction as the longitudinal direction is prepared from the final annealed steel sheet, and a tensile test is performed after holding at 800 ° C for 10 minutes. The strain rate is 0.3% / min up to 0.2% proof stress, and 3% / min thereafter, and a high-temperature tensile test is performed to measure the tensile strength (in accordance with JIS G 0567). An 800 ° C tensile strength of 200 MPa or more is considered good. A value of 240 MPa or more is more preferable.

[0049] <<Manufacturing Method of High Ni Alloy Sheet of the Present Invention>> The high Ni alloy sheet is manufactured by using a slab manufactured in a steelmaking process, through the steps of heating before hot rolling - hot rolling - final annealing, or heating before hot rolling - hot rolling - cold rolling - final annealing. Annealing after hot rolling and before cold rolling, or intermediate annealing during cold rolling may also be performed.

[0050] The slab heating temperature before hot rolling is set to 1240°C or higher, which makes it possible to dissolve the coarse TiC precipitates that existed in the slab before heating. The heating temperature is preferably set to 1300°C or lower.

[0051] During hot rolling, precipitation of TiC precipitates progresses if the temperature remains in the range of 700°C to 950°C. By setting the hot rolling finishing temperature to 1000°C or higher and the coiling temperature to 650°C or lower, the progress of TiC precipitates during hot rolling can be suppressed.

[0052] After hot rolling and before cold rolling, annealing is performed after hot rolling or intermediate annealing during cold rolling at an annealing temperature of 1050° C. or higher for an annealing time of 50 seconds or longer to promote solid dissolution of precipitated TiC.

[0053] In the final annealing after hot rolling when cold rolling is not performed, or in the final annealing after cold rolling when cold rolling is performed, the temperature rise rate from room temperature to 1000°C is set to 10°C / sec or more, thereby minimizing the precipitation of TiC and the coarsening of the precipitates during the temperature rise.

[0054] By taking the above measures before final annealing, the precipitation of TiC and the coarsening of the precipitates are minimized at the start of final annealing (after the annealing temperature is reached), and as a result, the pinning effect of the precipitates during final annealing is minimized, so that the average crystal grain size can be made 60 μm or more by adopting conditions of a final annealing temperature of 1050° C. or higher and a final annealing time of 30 seconds or longer. It is preferable to set the upper limit of the final annealing time to 90 seconds.

[0055] Molten steel having the chemical composition shown in Table 1 was melted in an 80-ton electric furnace and cast into slabs with a thickness of 300 mm using a continuous casting facility.

[0056] In the hot rolling step, the slab was heated at the slab heating temperature, finishing temperature, and coiling temperature shown in Table 2, and hot rolling was performed, followed by coiling. Thereafter, for Examples in which "Cold rolling" is marked "Yes" in the "With or Without Cold Rolling" column in Table 2, hot-rolled sheet annealing was performed at the annealing temperature and time shown in Table 2. The sheet thickness after hot rolling was 6 mm. In cold rolling, cold rolling was performed to the sheet thickness shown in Table 2. When the final sheet thickness was 0.3 mm or less, intermediate annealing was performed under the same conditions as the hot-rolled sheet annealing conditions. For Examples in which "With or Without Cold Rolling" is marked "No" in Table 2 (Manufacturing Method No. B6), the sheet thickness after hot rolling was 4 mm, and annealing after hot rolling was designated "final annealing."

[0057] After cold rolling (after hot rolling in Manufacturing Method No. B6), final annealing was performed at the temperature increase rate (from room temperature to 1000° C.), annealing temperature, and annealing time shown in Table 2.

[0058] The results are shown in Table 3. In Tables 1 to 3, values ​​that deviate from the present invention, values ​​that deviate from the preferred manufacturing conditions of the present invention, and values ​​that deviate from the target properties of the present invention are underlined.

[0059]

[0060]

[0061]

[0062] Inventive Examples C1 to C16 in Table 3 contain the component compositions specified in the present invention, are produced by the preferred production method of the present invention, and have the crystalline structure specified in the present invention. As a result, as shown in Table 3, both the creep properties and high-temperature strength achieve the targets of the present invention.

[0063] Comparative Examples c1 to c7 in Table 3 are comparative examples whose component compositions are outside the range of the present invention.

[0064] In Comparative Example c1, the C content was below the upper limit, resulting in excessive precipitation of TiC precipitates, and the pinning effect of the precipitates prevented grain coarsening, resulting in poor creep properties. TiC was precipitated at the time of product production, resulting in insufficient solute Ti, and precipitation of Cr carbides resulted in poor high-temperature strength. In Comparative Example c2, the C content was below the lower limit, resulting in low high-temperature strength.

[0065] Comparative Example c3 had an Al content outside the upper limit, and the high-temperature strength was poor due to the presence of intermetallic compounds.Comparative Example c4 had an Al content below the lower limit, and the high-temperature strength was poor.

[0066] In Comparative Example c5, the Ti content was above the upper limit, so that the precipitation of TiC precipitates was excessive, and the pinning effect of the precipitates prevented the coarsening of crystal grains, resulting in poor creep properties.In Comparative Example c6, the Ti content was below the lower limit, so that precipitation strengthening by TiC precipitation could not be used, resulting in poor high-temperature strength.

[0067] In Comparative Example c7, the N content was outside the upper limit, and Ti precipitated as TiN, resulting in little precipitation strengthening by TiC and poor high-temperature strength.

[0068] Comparative Examples c8 to c15 are comparative examples in which the manufacturing method was outside the preferred range of the present invention, resulting in crystal structures outside the range of the present invention.

[0069] In Comparative Example c8, the slab heating temperature was low, so the coarse precipitates formed during slab heating could not be completely dissolved in the processes after hot rolling, resulting in poor high-temperature strength. In Comparative Example c9, the finishing temperature of hot rolling was low, so precipitation progressed during hot rolling, increasing the average TiC size. The coarse precipitates could not be completely dissolved in the subsequent processes, resulting in an insufficient proportion of solute Ti and poor high-temperature strength. In Comparative Example c10, the coiling temperature in hot rolling was high, so precipitation progressed during coiling, and the pinning effect of the precipitates prevented grain coarsening during final annealing, resulting in poor creep properties.

[0070] In Comparative Example c11, the annealing temperature of the hot-rolled sheet was low, so precipitates remained, and the pinning effect of the precipitates prevented coarsening of the crystal grains, resulting in poor creep properties. In Comparative Example c12, the annealing time of the hot-rolled sheet was short, so precipitates remained, and the pinning effect of the precipitates prevented coarsening of the crystal grains, resulting in poor creep properties. The high-temperature strength was also low.

[0071] In Comparative Example c13, the temperature rise rate in the final annealing was low, so precipitation progressed during the temperature rise, and the pinning effect of the precipitates prevented coarsening of the crystal grains, resulting in poor creep properties. In Comparative Example c14, the final annealing temperature was low, so the precipitates remaining in the steel were not sufficiently dissolved, resulting in a pinning effect. Also, the final annealing temperature was low, so coarsening of the crystal grains did not progress, resulting in poor creep properties, and the solute Ti in the product was insufficient, resulting in poor high-temperature strength. In Comparative Example c15, the final annealing time was short, so the precipitates remaining in the steel were not sufficiently dissolved, resulting in a pinning effect. Also, the final annealing time was short, so coarsening of the crystal grains did not progress, resulting in poor creep properties, and the solute Ti in the product was insufficient, resulting in poor high-temperature strength.

Claims

1. In mass%, it contains C: 0.020 to 0.100%, Si: 0.05 to 1.00%, Mn: 0.05 to 2.00%, P: 0.035% or less, S: 0.0015% or less, Cr: 18.0 to 25.0%, Ni: 18.0 to 50.0%, Al: 0.05 to 1.00%, Ti: 0.15 to 1.50%, N: 0.020% or less, O: 0.0030% or less, B: 0.0003 to 0.0030%, and the balance being Fe and impurities, and the average crystal grain size is 60 μm or more, the average size of TiC is 1 μm or more, and the density of TiC is 4000 pieces / mm ２ or less, a solid solution ratio of Ti (solid solution Ti amount / contained Ti amount) (mass ratio) is 0.75 or more, and the plate thickness is 0.2 to 4 mm.

2. The high Ni alloy plate according to claim 1, further containing, in mass %, one or more of Ca: 0.0003 to 0.0070%, Mg: 0.0060% or less, Mo: 5.00% or less, W: 2.00% or less, Cu: 3.00% or less, Co: 2.00% or less, V: 1.00% or less, Nb: 1.00% or less, Ta: 1.00% or less, Sn: 0.10% or less, and REM: 0.10% or less, in place of a portion of the Fe.

Images