JP2006063442A - H-shaped steel excellent in fire resistance and production method therefor - Google Patents

Abstract

<P>PROBLEM TO BE SOLVED: To provide an H-shaped steel excellent in fire resistance in which a flange portion has a tensile strength of 400 Mpa or more at room temperature, a 0.2% proof stress of 157 MPa or more at 600°C, and a Charpy impact absorbed energy of 100 J or more at 0°C, and to provide a production method therefor. <P>SOLUTION: The steel substantially contains, in mass, 0.03 to 0.15% C, 0.1 to 0.6% Mo, ≤0.06% Nb and 0.002 to 0.012% N, wherein a total mol fraction of precipitate of alloy carbides and alloy carbonitrides at 600°C is 0.2% or more, and the ratio of total mol fraction of precipitate of alloy carbonitrides etc.=(the total mol fraction of precipitate of alloy carbonitrides etc. at 300°C/the total mol fraction of precipitate of alloy carbonitrides etc. at 600°C) is 4.0 or less. The production method therefor comprises the steps of: reheating a cast steel having the above components, to 1,100 to 1,300°C; hot rolling the cast steel to form the shaped steel at 860°C or more as measured on a surface of a flange portion of the shaped steel; and after terminating the hot rolling, naturally cooling or rapidly cooling then naturally cooling the shaped steel. <P>COPYRIGHT: (C)2006,JPO&NCIPI

Description

  The present invention relates to an H-section steel having a low yield point ratio and excellent toughness and fire resistance (hereinafter abbreviated as refractory H-section steel) and a method for producing the same.

  Previously, based on the Ministry of Construction (now Ministry of Land, Infrastructure, Transport and Tourism) Notification No. 332 established in March 1987, high temperature design strength can be secured, and fireproof coating of steel used for building structures can be made unnecessary or reduced. Refractory steel materials for use in the “New Fireproof Design Method” are provided.

  Corresponding to such trends, for H-section steel, a technique for securing strength at 600 ° C. by precipitation of Mo-based carbides, that is, securing strength and yield ratio in a high temperature range based on precipitation strengthening technology, Many prior arts that ensure fire resistance have been disclosed.

  The present inventor, for example, as shown in Patent Document 1, in addition to (1) precipitation strengthening technology, (2) the number of oxides obtained by deoxidizing dissolved oxygen in molten steel with Ti, B, Mg, etc. We have developed a refractory H-section steel with a control technology, that is, an oxide metallurgy technology. The oxide metallurgy technology has the following effects.

  In the manufacturing process of H-section steel, the rolling finish temperature and the cooling rate are different depending on the part due to the non-uniform shape in the cross section of the steel material. Non-uniformity occurs, and variations in cross-sectional mechanical properties occur.

  In particular, among the flange portions, the fillet portion (see FIG. 1) where the flange and the web are coupled has a smaller amount of strain due to rolling compared to other flange portions, and the processing in the high temperature region is strong. It is done.

  In the H section (see FIG. 1), a difference of about 150 ° C. may occur as a finishing temperature difference between the three points of the fillet, the quarter flange, and the web. The difference in mechanical properties between the cross-sectional parts due to this difference in rolling temperature history must be eliminated.

  Here, by dispersing the intragranular transformation nuclei such as Ti oxide and promoting the intragranular transformation, the temperature dependence of the microstructure formation in hot rolling is reduced, and the microstructure is made fine and uniform. Achieve homogenous characteristics. Furthermore, since not only the effect of homogenizing the crystal grains but also the effect of refining is exhibited, the toughness of the base material can be improved.

JP-A-9-104944

  The present inventor has produced H-shaped steel having excellent fire resistance such as that disclosed in Patent Document 1, for example. The following problems are clarified and improvements are made to solve the problems. It was.

  As in Patent Document 1, the technology for imparting fire resistance to the H-shaped steel by applying the oxide metallurgy technology is a general-purpose process such as oxygen concentration control before Ti addition and subsequent Ti addition in the steel slab manufacturing stage. Compared to the above, since complicated processes are required, problems such as a reduction in production capacity and an increase in manufacturing costs have occurred.

Furthermore, precipitation strengthening is effective in ensuring strength and yield ratio in a high temperature range, that is, ensuring fire resistance, but the Mo-based carbide disclosed in Patent Document 1 is mainly Mo ₂ C. In the component range disclosed in Patent Document 1, it is predicted that the solid solution is completely dissolved in the steel in the temperature range of 600 to 650 ° C.

  In such a case, the effect of precipitation strengthening of the alloy carbide and the alloy carbonitride on the steel strength is lost.

  The main control element for precipitation strengthening is the precipitation amount of alloy carbide and alloy carbonitride (hereinafter referred to as “precipitation molar fraction”), which has temperature dependency. The temperature dependence is affected by the thermodynamic characteristics resulting from the carbon concentration in the steel, the type of alloy carbide and alloy carbonitride, and the like.

  The effect of the carbon concentration is that the solid solution carbon in the ferrite accompanying the temperature drop when the concentration is sufficiently higher than the concentration of Mo, Ti, V, Nb, Cr, etc. (alloy elements) that produce carbonitrides. As the concentration decreases, the amount of carbon contributing to precipitation increases, so the total precipitation mole fraction of alloy carbide and alloy carbonitride also increases.

  As a result, even when the temperature drop width was the same, when the carbon concentration was high and the increase in the total precipitation mole fraction of the alloy carbide and alloy carbonitride was large, the strength and yield ratio at room temperature were too large.

  The present inventor believes that alloy carbide and alloy carbonitride should be alloy carbide and alloy carbonitride that are once solutionized in the reheating process during material production and precipitated in the cooling process in the subsequent hot rolling process. For example, since it does not contribute effectively to precipitation strengthening, it is not only desirable that the thermodynamic properties be stable for alloy carbides and alloy carbonitrides, but it is necessary to make the solution solid at heating temperature. I came up with the idea.

  And as a result of examining the above problems, it is necessary to appropriately control the fluctuation range of the precipitation molar fraction of alloy carbide and alloy carbonitride in the high temperature range and the normal temperature range. Found that it was necessary to design the steel components in consideration of the following items.

  (I) In order to suppress an excessive increase in strength and yield ratio when reaching the normal temperature range, the total precipitation mole fraction of alloy carbide and alloy carbonitride in the normal temperature range is suppressed.

  (Ii) In order to ensure strength in the high temperature range, a total precipitation mole fraction of a predetermined amount or more of the alloy carbide and alloy carbonitride in the predetermined high temperature range is ensured.

  Based on this viewpoint, the present inventor has made various designs of alloy carbides and alloy carbonitrides. The amount of precipitation is controlled by the quantitative balance between the alloy elements of Mo, Ti, V, Nb, Cr, etc., and the alloy element group constituting the alloy carbonitride, and C and N. Invented H-section steel which is controlled by mutual quantitative balance among elements.

  In addition, the “alloy carbonitride / precipitate” targeted in the present invention means the total of alloy carbide and alloy carbonitride excluding cementite.

Specifically, particularly in a temperature range of 600 to 650 ° C., in the case where Mo-based carbide (mainly Mo ₂ C), which is highly likely to be completely dissolved in steel, is mainly composed of precipitates, the strength of the steel material is reduced. The contribution due to precipitation strengthening of alloy carbide may be lost.

Therefore, as part of Mo carbides alternative focuses on highly stable MCN type carbonitride in a high temperature range than the M ₂ C type, compared to the M ₂ C-type carbide, MCN type carbonitride It has been found that it is effective to increase the amount of precipitation of the above to solve the above problems.

In order to form alloy carbonitrides other than Mo-based carbides (M ₂ C-based), it is preferable to mainly increase Nb as an alternative to Mo, and there is an appropriate balance of addition amounts of Nb and Mo. It has been found that the production of preferred types of alloy carbides and alloy carbonitrides mainly composed of Nb and Mo can be controlled.

  Note that the normal temperature generally refers to a temperature range of about 0 to 30 ° C., but the actual precipitation behavior is a solid solution of alloy elements, carbon, nitrogen, etc. in steel near the normal normal temperature. It is known that the diffusion of elements is significantly reduced.

  That is, in the temperature range from less than 300 ° C. to a general room temperature, an equilibrium state of 300 ° C. is almost maintained. Therefore, in the description of the present invention, for convenience, the 300 ° C. precipitation state is substituted for the room temperature precipitation state. I let you.

  This invention was made | formed based on the said knowledge, The summary is as follows.

  (1) By mass%, essentially C: 0.03-0.15%, Mo: 0.1-0.6%, Nb ≦ 0.06%, N: 0.002-0.012% (X) alloy carbide and alloy carbonitride at 600 ° C. having a total precipitation molar fraction of 0.2% or more, and (y) alloy carbide and alloy Total precipitation mole fraction ratio of carbonitride = (total precipitation mole fraction of alloy carbide and alloy carbonitride at 300 ° C.) / (Total precipitation mole fraction of alloy carbide and alloy carbonitride at 600 ° C.) Is an H-section steel excellent in fire resistance, characterized by being 4.0 or less.

  (2) Further, by mass, Si: 0.05 to 0.50%, Mn: 0.4 to 1.6%, Nb: 0.02 to 0.06%, Al ≦ 0.01% H-shaped steel with excellent fire resistance, (z1) normal temperature tensile strength of flange part is 400 MPa or more, (z2) 0.2% proof stress at 600 ° C is 157 MPa or more, and (z3) at 0 ° C The H-section steel having excellent fire resistance according to the above (1), wherein Charpy impact absorption energy is 100 J or more.

  (3) Further, in mass%, one of V ≦ 0.20%, Ti ≦ 0.02%, Cr ≦ 0.7%, Ni ≦ 1.0%, Cu ≦ 1.0%, or The H-section steel excellent in fire resistance as described in (1) or (2) above, comprising two or more types.

  (4) After reheating the steel slab having the composition described in any one of (1) to (3) above to 1100 to 1300 ° C, hot rolling is started and the surface temperature of the flange portion is set to 860 ° C. A method for producing an H-section steel excellent in fire resistance, characterized by finishing as described above and then allowing to cool, or cooling after accelerated cooling.

  (5) In the hot rolling step, the surface of the flange portion of the H-shaped steel is water-cooled to 700 ° C. or less, and the water-cooling / rolling cycle of rolling in the reheating process is performed once or more (4 The manufacturing method of H-section steel excellent in fire resistance as described in 1).

  (6) In the accelerated cooling, accelerated cooling is performed such that an average cooling rate up to 600 ° C. is 0.5 to 5.0 ° C./s, and the fire resistance as described in (5) above is excellent. A manufacturing method of H-shaped steel.

  According to the present invention, a steel slab having a predetermined component composition is subjected to a predetermined manufacturing process such as hot rolling at a rolling finishing temperature of 860 ° C. or higher, thereby achieving an appropriate balance of Nb and Mo addition amounts. An alloy carbide and alloy carbonitride mainly composed of Nb and Mo are formed, and a fire-resistant H-section steel having required high-temperature strength and mechanical properties at room temperature and a method for producing the same can be provided. .

  First, the reason why the chemical composition of the steel slab used for rolling is limited in the present invention will be described. In addition,% means the mass%.

  C is added as an effective component for improving the strength of the steel. However, if it is less than 0.03%, the strength required for structural steel cannot be obtained. On the other hand, excessive addition exceeding 0.15% Material toughness, weld crack resistance, weld heat affected zone (HAZ) toughness, etc. are significantly reduced. Therefore, C is preferably 0.03 to 0.15%.

  In addition to functioning as a deoxidizing element, Si is a component necessary for ensuring the strength of the base material, but if it is less than 0.05%, it hardly contributes to strength improvement, while if it exceeds 0.50% , HAZ produces high carbon island martensite, which is a hardened structure, and significantly deteriorates toughness. Therefore, Si is preferably 0.05 to 0.50%.

  Mn needs to be added in an amount of 0.4% or more in order to ensure the strength and toughness of the base material, but if it exceeds 1.6%, the HAZ toughness and crack resistance are impaired. Therefore, Mn is preferably 0.4 to 1.6%.

  Mo is an alloy element that generates carbides. Since the present invention secures the normal temperature and high temperature strength by precipitation of the carbide, it is an effective component for securing these strengths. If it is less than 0.1%, sufficient high-temperature strength cannot be ensured. On the other hand, if it exceeds 0.6%, the hardenability is excessively increased and the base metal toughness and the HAZ toughness are impaired. Therefore, Mo is preferably 0.1 to 0.6%.

  Furthermore, if it limits, 0.25 to 0.4% is more preferable, and the most preferable is 0.25 to 0.3%.

  Nb, like V and Ti, is an alloy element that constitutes an M (C, N) type carbonitride and contributes to precipitation strengthening. However, when Nb exceeds 0.06%, the amount of carbonitride that is not solutionized increases even at a heating temperature of 1100 to 1300 ° C. before hot rolling, and does not contribute to precipitation strengthening. Further, Nb is more preferably 0.03 to 0.04%.

  Moreover, when Nb is less than 0.02%, the strength improvement by precipitation strengthening is insufficient. Therefore, Nb is preferably 0.02 to 0.06%.

  N is an important component constituting carbonitride, and if it is less than 0.002%, the amount of precipitation is insufficient, while if it exceeds 0.012%, the base metal toughness is remarkably lowered. Therefore, N is preferably 0.002 to 0.012%.

  Al is a strong deoxidizing element, but if it is contained in an amount exceeding 0.01%, it combines with N to precipitate AlN, thereby reducing the amount of precipitation of carbonitride that is a feature of the present invention. Therefore, Al is preferably 0.01% or less.

  Next, the reason for limitation related to the concentration range of the alloy element selectively added in the present invention will be described.

  V, like Nb and Ti, constitutes an M (C, N) type carbonitride and is an alloy element that contributes to precipitation strengthening. V is solid-solved in M (C, N) type carbonitride produced by Nb addition or Nb and Ti combined addition in the present invention, and (Nb, V) (C, N) or (Nb , V, Ti) (C, N) to change the thermal stability of the carbonitride.

Specifically, the addition of V can shift the thermodynamic characteristics of the M (C, N) type carbonitride to the low temperature region side and lower the solution temperature. In addition to this, V forms a solid solution (Mo, V) ₂ C in an M ₂ C type carbide mainly composed of Mo, and shifts the thermodynamic characteristics of the M ₂ C type carbide to a high temperature side.

  However, if V exceeds 0.20%, the amount of carbonitride that is not solutionized increases even at a heating temperature of 1100 to 1300 ° C. before hot rolling due to the combined effect with Nb, and does not contribute to precipitation strengthening. Therefore, V is preferably 0.20% or less.

  Ti, like Nb and V, is an alloy element that constitutes M (C, N) carbonitride and contributes to precipitation strengthening. In the present invention, it is dissolved in M (C, N) type carbonitride produced by V addition or combined addition of V and Nb, and (V, Ti) (C, N) or (V, Ti, Nb) (C, N) is configured to change the thermal stability of the carbonitride.

  Specifically, the thermal stability of the M (C, N) type carbonitride is expanded to a high temperature range by adding Ti. However, if Ti exceeds 0.02%, the amount of carbonitride that is not solutionized increases even at a heating temperature of 1100 to 1300 ° C. before hot rolling, and does not contribute to precipitation strengthening. Therefore, Ti is preferably 0.02% or less.

  Cr is not only an effective component for increasing the normal temperature strength and high temperature strength of the base metal by improving hardenability and precipitation hardening, but it also suppresses grain boundary oxidation on the steel surface and improves surface properties (smoothness). Is also a functional alloying element. However, addition of over 0.7% adversely affects the base metal toughness and the HAZ toughness. Therefore, Cr is preferably 0.7% or less.

  Ni is an alloy element effective for increasing the toughness of the base material. However, addition of over 1.0% significantly increases the component cost, so Ni is preferably 1.0% or less.

  Cu is an alloy element effective for strengthening the base material, but is also an alloy element that simultaneously increases the hardenability and impairs the base material toughness and the HAZ toughness. Therefore, Cu is preferably 1.0% or less.

  The composition of alloy carbides and alloy carbonitrides and the molar fraction of precipitation can actually be measured by microscopic observation and analysis at the electron microscope level, but as a relatively simple determination method, thermodynamic equilibrium You may employ | adopt the calculation using a calculation program.

  The thermodynamic equilibrium calculation program employed in the present invention is the commercially available software “Thermo-Calc”, and the database is “SSOL”. Not as long.

In the present invention, the precipitation molar fraction of alloy carbide and alloy carbonitride is MCN type Face Centered Cubic type and M ₂ C type Hexagonal Close-Packed type 2 The total value of the precipitation mole fractions of various types of alloy carbides and alloy carbonitrides was used. Under these calculation conditions, the components and temperature were changed, and the precipitation mole fraction was evaluated.

  It is preferable to increase Nb as a partial replacement for the Mo-based carbide that is highly likely to be completely dissolved in steel in the temperature range of 600 to 650 ° C. and not contribute to precipitation strengthening.

  In the component range of the present invention containing Nb in a range of 0.06% or less, pay attention to Mo and Nb in the alloy element group constituting the alloy carbide and the alloy carbonitride, and balance with the amounts of C and N From the above, the types of alloy carbides and alloy carbonitrides and the resulting precipitation strengthening molar fraction at each temperature were varied, and the mechanical properties at each temperature and room temperature were examined.

Regarding the precipitation mole fraction, the alloy carbide containing Mo as a main component and partly dissolved in V and Nb is M ₂ C type, and the alloy carbonitride containing V and Nb as the main components and partly dissolved in Mo. Is preferably MCN type and is predicted by thermodynamic equilibrium calculation. In the case of an actual process performed by continuous cooling, it may be corrected because it is slightly different from the calculated thermodynamic equilibrium.

  First, in order to ensure fire resistance, mechanical properties at 600 ° C., particularly 0.2% proof stress are important. 157 MPa or more is required.

  The balance of the alloy element group, C and N was changed, and the total precipitation mole fraction of the alloy carbide and the alloy carbonitride was changed in the range of 0 to 1.0%, and as a result, it was found to be 0.2% or more. I knew it was good.

  Next, in order to solve the problem that the strength and yield ratio at room temperature are too large, the ratio of the total precipitation mole fraction of alloy carbide and alloy carbonitride, that is, the value at 300 ° C. and the value at 600 ° C. The ratio between the values and the suitability of the mechanical properties at each temperature and room temperature were investigated.

  That is, in order to ensure fire resistance while suppressing excessive strength at room temperature and to achieve both, (1) Tensile strength of the flange portion is 400 MPa or more and (2) Charpy at 0 ° C. As for fire resistance, it was found that (2) the 0.2% proof stress at 600 ° C. should be 157 MPa or more while ensuring the impact absorption energy of 100 J or more.

  In particular, Nb is contained as an alternative to Mo, and the balance between the amount of the alloy element group and C and N is changed, so that the total precipitation molar fraction of the alloy carbide and the alloy carbonitride is 0 to 10. As a result of changing within a range of 0%, it was found that it should be 4.0 or less.

  As described above, as a result of changing the balance between the amount of the alloy element group and C and N after containing Nb as an alternative to Mo, for example, as shown in FIGS. 2 (a) and 2 (b) New alloy carbides and alloy carbonitrides can be newly designed.

Here, the total amount of the M ₂ C type, which is an alloy carbide in which V and Nb are also partly dissolved, and the MCN type, which is an alloy carbonitride in which V and Nb are the main components and Mo is also partly dissolved. Is indicated by a broken line as “total”.

  An example containing 0.02% Nb is shown in FIG. The total precipitation mole fraction of alloy carbide and alloy carbonitride at 600 ° C. is 0.54%. In addition, the total precipitation mole fraction of alloy carbide and alloy carbonitride when the temperature drops from 600 ° C. to 300 ° C. only increases from about 0.54% to about 1.28%, Is 2.37, and it can be seen that the requirement specified in the present invention is 4.0 or less.

  An example containing 0.06% Nb is shown in FIG. The total precipitation mole fraction of alloy carbide and alloy carbonitride at 600 ° C. is about 0.41%. On the other hand, the total precipitation mole fraction of alloy carbide and alloy carbonitride at 300 ° C. is about 1.14%. As a result, the ratio between the two is suppressed to 2.78, and it is understood that the requirement defined by the present invention is 4.0 or less.

  As shown in these figures, as a result, it is a component design technique that can suppress the excessive increase in strength at room temperature while maintaining the fire resistance at 600 ° C. with respect to the precipitation strengthening effect of the precipitate.

  The prior art is shown in FIGS. In both FIG. 3A and FIG. 3B, the component ranges are included in the component ranges defined in the present invention.

  In the prior art shown in FIG. 3 (a), the total precipitation mole fraction of alloy carbide and alloy carbonitride at 600 ° C. is about 0.16%, which is 0.20% or more of the requirement defined in the present invention. Not satisfied. Furthermore, the total precipitation mole fraction ratio at 600 ° C. and 300 ° C. is 4.6, which does not satisfy the requirement of 4.0 or less, and as a result, the strength at normal temperature is increased. It has passed.

  In the prior art shown in FIG. 3B, the total precipitation mole fraction of the alloy carbide and the alloy carbonitride at 600 ° C. is about 0.204%, which is 0.2% or more of the requirement defined in the present invention. Is satisfied.

  However, the ratio of the total precipitation mole fraction at 600 ° C. and 300 ° C. is that the total precipitation mole fraction of the alloy carbide and alloy carbonitride at 300 ° C. is about 0.84%. .2, which deviates from the requirement of 4.0 or less defined in the present invention. Thereby, it turns out that the intensity | strength in normal temperature is strengthened too much.

  Since the size of precipitates such as alloy carbide and alloy carbonitride is also affected, it is preferable to add a device such as reducing the size to 10 to 1000 nm according to the target strength. A process of forming a solution before hot rolling and precipitating in the cooling process from hot rolling, and a process of precipitating by holding at 600 ° C. are preferable.

In the present invention, the precipitation amount of MCN type carbonitride showing a stable precipitation amount with little temperature dependency in the range of 300 to 600 ° C. and M ₂ C type carbide having a relatively high temperature dependency in the same temperature range. Although the ratio is not particularly defined, the ratio of MCN-type carbonitride precipitation / M ₂ C-type carbide precipitation (hereinafter referred to as MCN / M ₂ C ratio) is increased as compared with the prior art. .

For example, when 600 ° C. is taken up as an example of the high temperature region, the effect becomes remarkable by setting the MCN / M ₂ C ratio to 0.1 or more. However, when the total amount of M ₂ C is reduced, it is difficult to determine only by the MCN / M ₂ C ratio. Therefore, even if it is less than 0.1, it is not excluded from the range defined in the present invention.

  Next, the reason why the process in the hot rolling process is limited will be described.

  First, a steel piece is reheated to 1100-1300 degreeC. The reheating temperature is limited to 1100 to 1300 ° C., in the hot rolling of H-section steel, while securing a sufficient temperature to perform processing in the austenite region, alloy carbide and alloy carbonitride are once solutionized. This is for fully expressing precipitation strengthening.

  After reheating, the steel slab is hot-rolled. The process is basically a breakdown process by punching, an intermediate rolling process by an intermediate universal rolling mill group consisting of an edger rolling mill and a universal rolling mill. And a finish rolling process using a universal rolling mill. In addition, the said process also includes the skew roll rolling process which controls the web height of H-section steel.

  In the rolling step, in the breakdown step, a rolling roll having a protrusion at the center of the hole bottom and a plurality of hole molds having different hole bottom widths is subjected to rolling in the width direction of the steel slab. Ensure flange width and web height.

  Subsequently, in the intermediate rolling step, an appropriate flange width is secured with an edger rolling mill, and an appropriate web thickness and flange thickness are secured with a universal rolling mill. Further, in the finish rolling process, the surface temperature of the flange portion is set to 860 ° C. or higher to form a predetermined H-section steel size.

  In the present invention, for example, a large H-section steel having a web thickness of 40 mm, a flange thickness of 60 mm, a web height of 500 mm, and a flange width of 500 mm is selected from an H-section steel having a web thickness of 9 mm, a flange thickness of 12 mm, a web height of 500 mm, and a flange width of 200 mm. It is preferable to be a target.

  In the hot rolling after reheating, it is preferable to perform one or more water-cooling / rolling cycles in which the surface of the flange portion of the H-shaped steel is water-cooled to 700 ° C. or lower during rolling and rolled in the reheating process.

  As described above, the fillet and the flange are hotter than the web due to the shape of the H-section steel. Therefore, in order to reduce this temperature non-uniformity, the water cooling / rolling cycle is performed once in the rolling process. Do it above. The water-cooling / rolling cycle is preferably carried out one or more times in accordance with the size of the H-section steel and the number of rolling passes.

  In the present invention, it is preferable to cool after completion of hot rolling, or to cool after accelerated cooling. In this cooling process, the microstructure can be refined to increase the normal temperature strength, toughness and high temperature strength of the H-section steel.

  When performing accelerated cooling before standing to cool, it is preferable to perform accelerated cooling with an average cooling rate up to 600 ° C. of 0.5 to 5.0 ° C./s in terms of further miniaturizing the microstructure.

  In the present invention, after the cooling step, mechanical properties of the normal temperature tensile strength of the flange portion of 400 MPa class, 0.2% proof stress at 600 ° C. of 157 MPa or more, and 0 ° C. Charpy impact absorption energy of 100 J or more. H-shaped steel having excellent fire resistance, and a machine having a flange portion having a normal temperature tensile strength of 490 MPa class, a 0.2% proof stress at 600 ° C. of 217 MPa or more, and a 0 ° C. Charpy impact absorption energy of 100 J or more H-shaped steel having excellent properties and fire resistance can be produced.

  Next, examples of the present invention will be described. The conditions of the examples are one example of conditions adopted for confirming the feasibility and effects of the present invention, and the present invention is limited to this one example of conditions. Is not to be done. The present invention can adopt various conditions as long as the object of the present invention is achieved without departing from the gist of the present invention.

  The steels (invention steel and comparative steel) that had the component compositions shown in Table 1 and were obtained by respectively determining the precipitation mole fractions of alloy carbide and alloy carbonitride at 600 ° C. and 300 ° C. by equilibrium calculation were converted. The slab steel pieces having a thickness of 240 to 300 mm were cast by melting in a furnace and continuously casting.

  In the table, “tr” means a trace (= not detectable even if analyzed, or unavoidable impurity level).

  For the trial steels a to e used as comparative steels, a is that the total precipitation molar fraction of alloy carbide and alloy carbonitride at 600 ° C., which is a requirement of the present invention, does not satisfy 0.2% or more. , A to e have a ratio obtained by (total precipitation mole fraction of alloy carbide and alloy carbonitride at 300 ° C.) / (Total precipitation mole fraction of alloy carbide and alloy carbonitride at 600 ° C.). It does not meet the requirement of 0.0 or less.

  After the steel slab is reheated to 1100-1300 ° C., a breakdown process by hole rolling, an intermediate rolling process by an intermediate universal rolling mill group consisting of an edger rolling mill and a universal rolling mill, and finishing by a universal rolling mill It used for the hot rolling process comprised by a rolling process, and manufactured the H-shaped steel of a predetermined size.

  In the hot rolling process, the web height of the H-section steel was appropriately controlled by the skew roll rolling process.

  The H-shaped steel was manufactured in a size range from a web thickness of 9 mm, a flange thickness of 12 mm, a web height of 500 mm, and a flange width of 200 mm to a web thickness of 40 mm, a flange thickness of 60 mm, a web height of 500 mm, and a flange width of 500 mm.

As shown in FIG. 1, the mechanical properties of the manufactured H-section steel are ¼ and ½ width of the flange width overall length (B) at the center portion (½ t ₂ ) of the plate thickness t ₂ in the flange. The test pieces collected from three locations (1 / 4B and 1 / 2B, respectively) and 1 / 2H of the web height at the center of the thickness of the web were obtained by performing various tests.

  The mechanical properties of the flange 1/4 part (1 / 4B) can represent the mechanical characteristics of the flange part of the H-shaped steel, but this time there is a problem that the mechanical characteristics of the web at room temperature are excessively strengthened. In order to confirm that it was eliminated, the mechanical properties of the web and the average value of the three locations were investigated. The ratio of the mechanical properties of the web divided by the average value of the mechanical properties at three locations was investigated. The measurement was performed on the C cross section.

  Table 2 shows the room temperature yield strength, room temperature tensile strength, room temperature yield ratio, 0 ° C impact absorption energy value (3-point average value) in Charpy test, and 0.2% yield strength at 600 ° C obtained as a result of the above test. Indicates. In addition, the Charpy test employ | adopted the measured value in the flange 1 / 2B part (fillet) used as the lowest value in a H-section steel cross-section part.

  For the 0.2% proof stress at 600 ° C., the measured value at the flange 1 / 4B portion was adopted as a portion representing the strength of the H-shaped steel. There are two types of strength classes required for steel.

  One is a 400 MPa class normal temperature tensile strength defined as SN 400 class (this time, as an example, 400 to 520 MPa level), and the other is SN 490 class (as an example, 500 to 611 MPa level). The room temperature tensile strength is of the 490 MPa class, and these are shown separately. Moreover, the ratio which remove | divided the mechanical characteristic of the web by the average value of the mechanical characteristic of three places was written together.

  The invention steel satisfies the requirements regarding the component composition and the precipitation mole fraction of alloy carbides and alloy carbonitrides defined in the present invention, and, depending on the production conditions, yield strength, tensile strength, 0 ° C. impact absorption energy, etc. Satisfying the target mechanical properties and mechanical properties at high temperature (600 ° C) with respect to the strength ratio obtained by (0.2% yield strength at 600 ° C) / (yield strength at normal temperature) and the yield ratio at normal temperature is doing.

  Among the inventive steels, No. 3 and no. FIG. 4 (a) and FIG. 4 (b) show the temperature transition of the total amount of precipitation mole fraction of alloy carbide and alloy carbonitride for No. 6.

  As a method of suppressing the ratio of the total precipitation mole fraction of alloy carbide and alloy carbonitride to 300 ° C. and 600 ° C. to 4.0 or less, the total precipitation mole fraction of alloy carbide and alloy carbonitride at 300 ° C. An approach to suppress and an approach to increase the total precipitation mole fraction of alloy carbide and alloy carbonitride at 600 ° C. are shown.

  On the other hand, although the comparative steel satisfies the component conditions specified in the present invention, it does not meet the requirements for the precipitation molar fraction of alloy carbide and alloy carbonitride, so it has the target normal temperature mechanical properties and high temperature mechanical properties. It turns out that one or more of them are not satisfied.

  For comparative steels a, c, and d, 0.2% proof stress at a high temperature (here, 600 ° C.) of SN400 class a is less than 157 MPa even though the target is 157 MPa or more. Although steels c and d are targeted at 217 MPa, they do not reach 206 MPa and 209 MPa.

  Further, the comparative steel b has an SN400 class target tensile strength of 510 MPa or less, which is 544 MPa, which is too large at room temperature.

  The comparative steel e has exceeded the SN490 grade target value of 610 MPa. Furthermore, 0 ° C. shock absorption energy targeting 100 J or more is insufficient.

  As described above, according to the present invention, an alloy carbide and an alloy carbonitride mainly composed of Nb and Mo are formed in an H-section steel manufacturing process under an appropriate balance of the amount of Nb and Mo added. It is possible to provide an H-shaped steel having a required high temperature strength and mechanical properties at room temperature and excellent in fire resistance.

  Therefore, the above H-section steel is useful as a structural material for steel structures, and the present invention has great industrial applicability.

In the H-section steel, the position at which a specimen for obtaining the microstructure and mechanical properties is taken (1/4 and 1/1 of the flange width overall length (B) at the center (1 / 2t2) of the thickness t2 of the flange. It is a figure which shows 2 width (1 / 4B, 1 / 2B, respectively) and 1 / 2H of the web height in the plate | board thickness center part in a web. It is a figure which shows an example of this invention containing 0.02% of Nb. The total precipitation mole fraction of alloy carbide and alloy carbonitride at 600 ° C. is approximately 0.54%, and the total precipitation mole fraction of alloy carbide and alloy carbonitride at 600 ° C. and 300 ° C. The ratio is about 2.37. It is a figure which shows an example of this invention containing 0.06% of Nb. The total precipitation mole fraction of alloy carbide and alloy carbonitride at 600 ° C. is 0.41%, and the total precipitation mole fraction ratio of alloy carbide and alloy carbonitride at 600 ° C. and 300 ° C. is It is about 2.78. The temperature transition (%) of the precipitation mole fraction of M ₂ C type, which is an alloy carbide in which V and Nb are also partly dissolved, and alloy carbonitriding in which V and Nb are the main components and Mo is also partly dissolved. Three types of temperature transition (%) of the precipitation mole fraction of the MCN type, which is a product, and the temperature transition (%) of the “total” precipitation mole fraction, which is the sum of these two, are illustrated by broken lines. It is a figure which shows a prior art. Three types of M ₂ C type, MCN type, and “total” obtained by adding them are shown by broken lines. An example is shown in which the temperature dependency of MCN is lower than that of M ₂ C. The total precipitation mole fraction of alloy carbide and alloy carbonitride at 600 ° C. and the total precipitation mole fraction ratio of alloy carbide and alloy carbonitride at 600 ° C. and 300 ° C. are outside the scope of the present invention. ing. It is a figure which shows a prior art. Three types of M ₂ C type, MCN type, and “total” obtained by adding them are shown by broken lines. The total precipitation mole fraction of alloy carbide and alloy carbonitride at 600 ° C. is within the scope of the present invention, but the total precipitation mole fraction ratio of alloy carbide and alloy carbonitride at 600 ° C. and 300 ° C. Is outside the scope of the present invention. Invention Example No. of the embodiment of the present invention. FIG. Invention Example No. of the embodiment of the present invention. FIG.

Explanation of symbols

1 H-section steel 2 Flange 3 Web 4 Fillet part B Flange width H Web height t1 Web thickness t2 Flange thickness

Claims (6)

In mass%, essentially
C: 0.03-0.15%,
Mo: 0.1 to 0.6%,
Nb ≦ 0.06%,
N: 0.002 to 0.012%
H-shaped steel with excellent fire resistance, including
(X) the total precipitation mole fraction of alloy carbide and alloy carbonitride at 600 ° C. is 0.2% or more, and
(Y) Total precipitation mole fraction ratio of alloy carbide and alloy carbonitride = (total precipitation mole fraction of alloy carbide and alloy carbonitride at 300 ° C.) / (Alloy carbide and alloy carbonitride at 600 ° C. H-shaped steel excellent in fire resistance, characterized in that the total precipitation molar fraction of) is 4.0 or less. Furthermore, in mass%,
Si: 0.05 to 0.50%,
Mn: 0.4 to 1.6%,
Nb: 0.02 to 0.06%,
Al ≦ 0.01%
H-shaped steel with excellent fire resistance, including
(Z1) The normal temperature tensile strength of the flange portion is 400 MPa or more,
(Z2) 0.2% yield strength at 600 ° C. is 157 MPa or more, and
(Z3) The H-section steel excellent in fire resistance according to claim 1, wherein Charpy impact absorption energy at 0 ° C is 100 J or more. Furthermore, in mass%,
V ≦ 0.20%,
Ti ≦ 0.02%,
Cr ≦ 0.7%,
Ni ≦ 1.0%,
Cu ≦ 1.0%,
1 type or 2 types or more of these are included, The H-section steel excellent in fire resistance of Claim 1 or 2 characterized by the above-mentioned.   After the steel slab having the component composition according to any one of claims 1 to 3 is reheated to 1100 to 1300 ° C, hot rolling is started, and the surface temperature of the flange portion is finished at 860 ° C or higher, Next, the method for producing an H-shaped steel excellent in fire resistance, wherein the method is allowed to cool or is cooled after accelerated cooling.   The said hot rolling process WHEREIN: The surface of the flange part of H-section steel is water-cooled to 700 degrees C or less, and the water-cooling and rolling cycle which rolls in a recuperation process is performed 1 time or more. A method for producing H-shaped steel with excellent fire resistance.   The H-shaped steel excellent in fire resistance according to claim 5, wherein the accelerated cooling is performed so that an average cooling rate up to 600 ° C is 0.5 to 5.0 ° C / s. Manufacturing method.

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