JP4786556B2 - Steel material excellent in fire resistance and method for producing the same - Google Patents

Description

  The present invention relates to a steel material having a low yield point ratio and excellent toughness and fire resistance, particularly an H-section steel (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 technology for securing strength at 600 ° C. by precipitation strengthening of Mo-based carbides, that is, securing strength and yield ratio in a high temperature region based on precipitation strengthening technology, That is, many prior arts that ensure fire resistance are disclosed.

  So far, for example, Patent Document 1 has proposed a refractory H-shape by the aforementioned Mo-based carbide precipitation strengthening technique. In this Patent Document 1, it is essential to add Mo 0.3% or more in order to ensure high temperature strength. If Mo is less than 0.3%, the precipitation amount of Mo-based carbide is insufficient and sufficient high temperature is obtained. The knowledge that strength cannot be secured has been obtained. On the other hand, Mo is a relatively expensive alloy element, and from the viewpoint of economy, production with an alloy component that stably has a fireproof function without addition of Mo has been required.

JP 05-105947 A

  The inventors of the present invention, for example, manufactured an H-shaped steel having excellent fire resistance disclosed in Patent Document 1, but made the following problems clear and made improvements.

  In order to ensure the strength and yield ratio in the high temperature range by controlling the precipitation of Mo-based carbide as in Patent Document 1, a sufficient amount of Mo is required to precipitate the Mo-based carbide necessary for ensuring the strength. Although it is necessary to add, Mo is a relatively expensive alloy element, and in order to spread refractory steel materials to the market, it has been necessary to find an inexpensive component system in pursuit of economy.

  Further, precipitation strengthening is effective in ensuring strength and yield ratio in a high temperature range, that is, ensuring fire resistance. However, the Mo-based carbide disclosed in Patent Document 1 is mainly Mo2C. In the component range, the solid solution is expected to be 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. Further, Mo is a relatively rare alloy element, and from the viewpoint of saving resources and avoiding high manufacturing costs, it is necessary to exhibit similar characteristics with a component to which Mo is not added as much as possible.

  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.

  As precipitates other than Mo, precipitation strengthening of Ti, V, Nb, Cr and other alloy carbides and alloy carbonitrides, and in addition to these, precipitation strengthening by controlling fine precipitation of sulfides by Mn, Cr, etc. were studied. . Precipitation strengthening due to precipitates is effective for precipitation strengthening unless it is an alloy carbide and alloy carbonitride that are once solutionized in the reheating process during material production and then precipitated in the cooling process in the subsequent hot rolling process. Therefore, it has not been desirable for alloy carbides and alloy carbonitrides to have stable thermodynamic properties, but has led to the idea that they should be dissolved at heating temperature.

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

That is,
(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 the strength in the high temperature region, it is preferable to ensure the total precipitation molar fraction of the alloy carbide and the alloy carbonitride in a predetermined amount or more in the predetermined high temperature region.

  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 that constitute alloy carbides and alloy carbonitrides such as Ti, V, Nb, and Cr, and C and N. The thermodynamic characteristics of the alloy elements are Invented the H-section steel which is controlled by mutual quantitative balance. In addition, the “alloy carbonitride” targeted in the present invention means the total of alloy carbide excluding cementite and alloy carbonitride.

  Specifically, Mo, which has been indispensable for securing the high temperature strength of the refractory H-shaped steel, is not added. Instead, V alone is basically used as an alternative, or V + Nb, V + Ti, It has been found that the combined addition of V + Nb + Ti and the like has high temperature stability and can control the formation of a preferred type of alloy carbonitride, and it has been found effective to solve the above problems. Furthermore, by optimizing the addition amount of Mn and Cr, the tensile strength of the component at a normal temperature strength level from 400 MPa class to 490 MPa class, excellent toughness, and excellent fire resistance at 600 ° C. A combination was found.

This invention was made | formed based on the said knowledge, The summary is as follows.
(1) By mass%, C: 0.03 to 0.10%, Si: 0.05 to 0.50%, Mn: 0.05 to 1.3%, Al ≦ 0.01%, V: 0 .15-0.35%, Cr: 1.5-3%, N: 0.002-0.008%, S: 0.003-0.02%, P: 0.0001-0.1%, hints balance iron and unavoidable refractory excellent steel which is characterized by comprising a non-pure product.
( 2 ) A method for producing a steel material by re-heating a steel slab having the component described in (1 ), followed by hot rolling.
(A) After reheating to 1100-1300 ° C., hot rolling is started,
(B) After rolling,
(C) The method for producing a steel material having excellent fire resistance according to the above (1 ), wherein the steel material is allowed to cool or is cooled after accelerated cooling.
( 3 ) The method for producing a steel material having excellent fire resistance according to ( 2 ) , wherein the surface of the steel material is water-cooled to 700 ° C. or lower and a water-cooling / rolling cycle for rolling in the reheating process is performed once or more. .
( 4 ) Excellent fire resistance as described in ( 2 ) or ( 3 ), wherein accelerated cooling is performed at an average cooling rate of 0.5 to 5.0 ° C./s from the rolling end temperature to 600 ° C. Steel manufacturing method.
( 5 ) A method of manufacturing a steel material by re-heating a steel slab having the component according to (1 ), followed by hot rolling,
(A) After reheating to 1100-1300 ° C., hot rolling is started,
(B) The rolling end temperature is 850 ° C. or more at the steel surface temperature,
(C) After the rolling is completed, the average cooling rate in the range from 800 ° C. to 500 ° C. is accelerated to 3 to 15 ° C./s, the stop temperature is set to the range of 300 to 550 ° C. at the surface temperature of the steel material, and then allowed to cool. A method for producing a steel material having excellent fire resistance as described in (1 ) .

  In the present invention, Mo is not actively added, but it is allowed to contain 0.03% or less of Mo as an impurity.

  According to the present invention, a steel slab having a predetermined component composition is subjected to a predetermined manufacturing process such as hot rolling, so that Mo is not added, and V is mainly contained under an appropriate V addition amount balance. To provide a steel material with excellent fire resistance that has the required high-temperature strength and mechanical properties at room temperature, by forming the alloy carbide and alloy carbonitride that have been used, and by appropriately combining precipitation strengthening with fine sulfides, etc. Further, it is possible to provide a steel material and a method for producing an H-shaped steel having excellent fire resistance.

  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 the structural steel cannot be obtained. On the other hand, excessive addition exceeding 0.10% Material toughness, weld crack resistance, weld heat affected zone (HAZ) toughness, etc. are significantly reduced. The limited range of the C concentration is preferably 0.03 to 0.08%.

  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, the limited range of the Si concentration is set to 0.05 to 0.50%. Preferably it is 0.05 to 0.20%.

  When Mn is added in an amount of 0.4% or less, the Mn-based sulfide (MnS) is refined and the toughness is improved, particularly the toughness in the sheet width direction is remarkably improved, and the precipitation strengthening mechanism is effective. Demonstrated and suppresses strength drop in high temperature range, especially compared to normal temperature strength. Further, this finely precipitated MnS functions as a precipitation site for V-based M (C, N) type carbonitrides and functions effectively for precipitation strengthening. However, if the amount of Mn is less than 0.05%, the generation of sulfide is insufficient and the amount of solute S increases and the grain boundary becomes brittle, which causes steel slab cracking and steel material cracking during hot rolling. Therefore, the limited range of the Mn concentration is set to 0.05 to 0.4%. A preferable lower limit is 0.1% or more.

  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, the limited range of the Al concentration is set to 0.01% or less. On the other hand, since the effect of the present invention can be obtained even when Al is 0%, the lower limit of Al includes 0%.

  V is an alloy element that constitutes carbonitride and contributes to precipitation strengthening. By optimizing the addition amount of V under the Mo-free condition, the precipitation amount and distribution state of the M (C, N) type carbonitride are controlled. Moreover, the precipitation strengthening function is effectively exhibited by preferentially depositing on the finely precipitated MnS. At this time, by adding a large amount of V equivalent to C or N, an increase in the precipitation strengthening contribution accompanying an increase in the precipitation amount in the low temperature region is avoided. In order to reliably suppress excessive precipitation in this low temperature range, it is necessary to contain V by 0.15% or more. On the other hand, if V is added in excess of 0.35%, the amount of precipitation of the carbonitride becomes excessive, and the base metal toughness and the HAZ toughness are impaired. Therefore, the V concentration is set to 0.35% or less. Therefore, the limited range of the V concentration is 0.15 to 0.35%.

  Cr not only generates carbides and sulfides, but also improves hardenability when contained in a large amount, and improves normal temperature strength and high temperature strength. If Cr is added in an amount of less than 1.5%, the strength improving effect cannot be sufficiently exhibited. On the other hand, addition exceeding 3% causes a decrease in toughness. From the above, the limited range of the Cr concentration is 1.5 to 3%. In addition, Preferably it is set as 1.7 to 3% of range.

  N is an important component constituting carbonitride, and if it is less than 0.002%, the amount of precipitation is insufficient. On the other hand, if it exceeds 0.008%, it is continuously formed by N dissolved in the steel in a high temperature range. There is concern about the possibility of steel piece cracking during casting. Therefore, the limited range of the N concentration is set to 0.002 to 0.008%. Further, in order to suppress the possibility of steel piece cracking in continuous casting, preferably, when at least one of Ti and Nb described later is added, 0.004 to 0.008%, both Ti and Nb When not added, the range of 0.002 to 0.006% is preferable.

  S has conventionally been considered as an element present in steel as an inevitable impurity, but in the present invention, by combining with Mn and Cr, sulfide is finely precipitated and functions to improve toughness and high-temperature strength by dispersion strengthening. You can make it. If it is 0.003% or less, the volume fraction of sulfide is low, and the above effect is not sufficient. Further, if the addition exceeds 0.02%, the sulfide becomes coarse, so the above-mentioned effect of improving the mechanical properties due to the refinement of the sulfide is not exhibited. From the above, the limited range of the S concentration is set to 0.003 to 0.02%. 0.005 to 0.015% is preferable for best exhibiting the effect of refining sulfide.

  P is an element present in steel as an unavoidable impurity, and it is desirable to reduce it as much as possible. However, for the reduction to less than 0.0001%, the cost increase required for the dephosphorization treatment becomes significant. On the other hand, if the content exceeds 0.1%, the toughness is remarkably lowered. From the above, the content range of P is set to 0.0001 to 0.1%. A preferred range is 0.001 to 0.03%.

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

  Nb, like V and Ti, is an alloy element that constitutes an M (C, N) type carbonitride and contributes to precipitation strengthening. However, when the Nb addition amount exceeds 0.04%, 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.

  Moreover, when Nb addition amount is less than 0.005%, the strength improvement by precipitation strengthening is inadequate, and it is equivalent to no addition. Therefore, it is preferable that the limited range of the Nb concentration is 0.005 to 0.04%.

  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 region by adding 0.005% or more of Ti. However, if the Ti addition amount 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, it is preferable that the limited range of the Ti concentration is 0.005 to 0.02%.

  Ni is an alloy element effective for increasing the toughness of the base material. However, addition of more than 1.0% remarkably increases the component cost, so the Ni concentration limit range is preferably 1.0% or less. In order to sufficiently obtain the above effects, the lower limit of Ni is preferably set to 0.3%.

  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, the upper limit of the Cu concentration is preferably 1.0% or less. In order to sufficiently obtain the above effect, the lower limit of Cu is preferably set to 0.3%.

  In order to ensure fire resistance, mechanical properties at 600 ° C., particularly 0.2% proof stress are important. 157 MPa or more is required for steel materials having a strength level of 400 to 520 MPa class at ordinary temperature and 217 MPa or more for steel materials having a strength level of 490 to 611 MPa class at ordinary temperature. Moreover, in order to ensure sufficient toughness, Charpy impact absorption energy at 0 ° C. is required to be 100 J or more.

  As a result of investigating the mechanical properties of the rolled steel material by changing the balance of the alloy element group, C and N, it was found that the above-mentioned target can be realized in the above-described component range.

  Next, as a specific example of the steel material, the reason for limiting the manufacturing conditions in the hot rolling process of the H-section steel 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. Furthermore, it is preferable that the surface temperature of the flange portion is maintained at 800 ° C. or higher in the finish rolling process, and is formed into a predetermined H-section steel size.

  The manufacturing method of the present invention described above includes, for example, an H-shaped 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, 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. Suitable for application to large H-section steel. When the steel material is a thick steel plate, the reheating temperature of the steel slab is set to 1100 to 1300 ° C., and after carrying out normal thick plate rolling, it is allowed to cool, or after accelerated cooling, it is allowed to cool. Moreover, it is preferable that finishing temperature shall be 800 degreeC or more with the surface temperature of a steel plate.

  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. It is preferable to perform the 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. Even when the steel material is a thick steel plate, the surface of the steel plate may be water-cooled to 700 ° C. or less and rolled in a double heat process.

  In the present invention, after hot rolling is finished, it may be allowed to cool, or after accelerated cooling, it is preferably allowed to cool. In this cooling process, the microstructure can be refined to increase the normal temperature strength, toughness, and high temperature strength of steel materials, particularly H-shaped steel.

  In the case of performing accelerated cooling before allowing to cool, performing the accelerated cooling from the rolling end temperature to 600 ° C. at an average cooling rate of 0.5 to 5.0 ° C./s makes the microstructure finer. preferable.

  In the present invention, after the cooling step, the strength ratio obtained by (0.2% yield strength at 600 ° C.) / (Yield strength at normal temperature) of the flange portion is 50% or more, and the yield ratio at normal temperature. Is an H-shaped steel excellent in fire resistance having mechanical properties of 80 J or less and Charpy impact absorption energy at 0 ° C. of 100 J or more, and the normal temperature tensile strength of the flange portion is 400 MPa depending on the strength level. Grade, 0.2% proof stress at 600 ° C. is 157 MPa or more, and 0 ° C. Charpy impact absorption energy is 100 J or more. H with excellent fire resistance with mechanical properties of room temperature tensile strength of 490 MPa, 0.2% proof stress at 600 ° C. of 217 MPa or more, and 0 ° C. Charpy impact absorption energy of 100 J or more It is possible to manufacture the steel. In the present invention, JIS No. 2 mmV notch was used as the Charpy test piece, JIS No. 2 was used as the high-temperature tensile test piece, and JIS 13A or JIS 13B was used as the tensile test piece at room temperature.

Furthermore, the manufacturing method of the fireproof steel materials shown in said ( 5 ) is demonstrated.

  First, the steel slab is reheated to 1100 to 1300 ° C. The reason for this is that, similarly to the above-described method for manufacturing the H-shaped steel, a sufficient temperature for processing in the austenite region is secured, and alloy carbides and alloys This is because the carbonitride is once solutionized to sufficiently develop precipitation strengthening.

  Next, the re-heated steel slab is hot-rolled a plurality of times to a predetermined thickness, and the surface temperature of the steel is maintained at 850 ° C. or higher in the final rolling pass, that is, the finish rolling process, and hot-rolled. To complete. If the surface temperature of the steel material is less than 850 ° C. in the finish rolling, the cooling rate between 800 and 500 ° C. is set to 3 ° C./s or more when the start of subsequent accelerated cooling such as water cooling is delayed. This is because it becomes difficult.

  The accelerated cooling after completion of hot rolling may be performed by water cooling, but may be performed by mist cooling or air blowing depending on the cooling equipment. The accelerated cooling may be performed immediately after the end of hot rolling, and the starting temperature is preferably set to 850 ° C. or more at the surface temperature of the steel material.

  Accelerated cooling needs to be performed such that the average cooling rate between 800 ° C. and 500 ° C. is 3 to 15 ° C./s. In this cooling process, precipitation of the alloy carbonitride can be suppressed, and an increase in strength at room temperature can be suppressed. Moreover, the microstructure is also refined by accelerated cooling, and the toughness can be increased.

  The stop temperature of this accelerated cooling is the surface temperature of the steel material and needs to be in the range of 300 to 550 ° C. This is because, when the water cooling stop is less than 300 ° C., the self-tempering effect cannot be obtained and the toughness is lowered when the water is cooled. That is, after the accelerated cooling is stopped, the toughness is improved by the self-tempering effect by cooling. In order to obtain this effect, it is preferable to stop the accelerated cooling at a higher temperature. On the other hand, if the accelerated cooling stop temperature is higher than 550 ° C, it is difficult to set the average cooling rate of 800 ° C to 500 ° C to 3 ° C / s or more. In addition, when accelerating cooling is stopped at 550 ° C., cooling at 550 ° C. or lower is allowed to cool.

The steel material manufactured by such a manufacturing process has excellent mechanical properties such that the yield ratio at room temperature is 80% or less and the Charpy impact absorption energy at 0 ° C. is 100 J or more. Moreover, by satisfying the production condition ( 5 ) above, the 0.2% proof stress at 700 ° C. becomes 2/3 or more of the lower limit of the standard value of the yield strength at room temperature, and the fire resistance at 700 ° C. is improved. Excellent steel material can be obtained. The required strength at 700 ° C. is that the normal temperature tensile strength is 400 MPa class, the 0.2% proof stress at 700 ° C. is 157 MPa or more, the normal temperature tensile strength is 490 MPa class, and the strength at 700 ° C. The 0.2% proof stress is 217 MPa or more.

  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.

  Trial steels (invention steel and comparative steel) having the composition shown in Table 1 were melted in a converter and slab steel pieces having a thickness of 240 to 300 mm were cast by continuous casting. “Tr” means a trace (= not detectable even if analyzed, or unavoidable impurity level). In addition, Mo was not added about all the Examples, and all the analysis values were 0.01% or less.

  For the trial steels a to g used as comparative steels, a to c depart from the component ranges that are requirements of the present invention.

  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 characteristics of the manufactured H-section steel are ¼ and ½ width of the flange width overall length (B) at the center portion (1 / 2t ₂ ) of the plate thickness t ₂ in the flange. Various tests were conducted on test specimens collected from three locations (1 / 2B and 1 / 2B, respectively) and 1 / 2H of the web height at the center of the thickness (1 / 2t ₁ ) of the web. Asked. The mechanical properties of the flange 1/4 part (1 / 4B) can be representative of the mechanical properties of the H-shaped steel flange, but this time the problem of excessively strengthening the mechanical properties of the web at room temperature has been resolved. In order to confirm that this was done, 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. The ratio of 0.2% yield strength at 600 ° C. and yield strength at room temperature is shown. 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 400 to 520 MPa level) normal temperature tensile strength defined as SN 400 class, and the other is SN 490. Class (this time, as an example, 500 to 611 MPa level) and room temperature tensile strength of 490 MPa class, which are shown separately.

  Invented steel satisfies the requirements of the component composition specified in the present invention, and has target mechanical properties for yield strength, tensile strength, mechanical properties at normal temperature such as 0 ° C. impact absorption energy, and 0.2% proof stress at 600 ° C. Is satisfied. The target of SN 400 class, that is, TS 400 MPa class or higher, is that yield strength YP at room temperature is 235 MPa or more, preferably 355 MPa or less, tensile strength TS is 400 to 520 MPa, impact absorption energy at 0 ° C. is 100 J or more, and is 0. The 2% yield strength PS is 157 MPa or more. Further, the target of SN 490 class, that is, TS 490 MPa class or higher, is that YP is 325 MPa or more, preferably 445 MPa or less, TS is 490 to 611 MPa, 0 ° C. impact absorption energy is 100 J or more, and PS is 217 MPa or more.

  On the other hand, it can be seen that the comparative steel does not satisfy one or more of the target normal temperature mechanical properties and high temperature mechanical properties because the component conditions defined in the present invention are not satisfied.

Since the comparative steel g has C, Mn, V and Nb outside the scope of the present invention, the tensile strength at normal temperature exceeds 611 MPa and the 0 ° C. impact absorption energy is insufficient. Since the SN400 grade comparative steel b has C and Cr, and the comparative steel c has Cr outside the scope of the present invention, the 0.2% proof stress at 600 ° C. is less than 157 MPa. Since the comparative steel b has a small amount of C, the tensile strength at room temperature is also reduced to less than 400 MPa.
Since one or more of C, Mn, V, Cr, and Nb are out of the scope of the present invention in each of SN490 grade comparative steels a, d, and f, 0.2% proof stress at 600 ° C. is 158, The target of 207 or 175 MPa is not achieved, and the comparative steel d has a tensile strength at room temperature exceeding 611 MPa. Since SN, 490 grade comparative steel e has C, Mn, V and Cr outside the scope of the present invention, the tensile strength at room temperature exceeds the upper limit of 611 MPa.

Furthermore, although the Example of this invention is described, this condition is also one example of conditions adopted in order to confirm the feasibility and effect of this invention like the above-mentioned Example.
Trial steels (invention steel and comparative steel) having the composition shown in Table 3 were melted in a converter and slab steel pieces having a thickness of 240 to 250 mm were cast by continuous casting. “Tr” means a trace (= not detectable even if analyzed, or unavoidable impurity level). In addition, Mo was not added about all the Examples, and all the analysis values were 0.01% or less. The trial steels aa to ag used as comparative steels deviate from the component range which is a requirement of the present invention, and the underline means outside the scope of the present invention.

The steel slab having the above components was reheated under the conditions shown in Table 4 and then hot rolled to obtain a steel plate, and accelerated cooling by water cooling was performed after the rolling. A sample for investigating normal temperature tensile characteristics, Charpy characteristics, and high temperature tensile characteristics was collected from the central part of the thickness of the steel sheet in the sheet width direction perpendicular to the rolling direction.
For room temperature tensile properties, the number of test samples was 2, and the average value was used as representative data. For the Charpy characteristics, the test temperature was 0 ° C., the number of test samples was 3, and the average value was used as representative data. The high temperature tensile properties were measured in accordance with JIS G 0567, after raising the temperature to 700 ° C. and holding for 10 minutes, then applying a tensile load and measuring 0.2% proof stress and tensile strength.

Table 4 shows the room temperature yield strength, room temperature tensile strength, 0 ° C impact absorption energy value (3-point average value) in Charpy test, and 0.2% proof stress at 700 ° C, obtained as a result of the above test, together with the manufacturing conditions. Show. In addition, the underline of Table 4 means that the target mechanical characteristics are not satisfied.
Here, the target value of the TS400 MPa class is that the yield strength YP at room temperature is 235 MPa or more, preferably 355 MPa or less, the tensile strength TS is 400 to 520 MPa, the impact absorption energy at 0 ° C. is 100 J or more, and 0.2% at 700 ° C. Yield strength PS is 157 MPa or more. Further, target values of TS490MPa or higher class are YP of 325 MPa or more, preferably 445 MPa or less, TS of 490 to 611 MPa, 0 ° C. impact absorption energy of 100 J or more, and PS of 217 MPa or more.

  In Table 4, 101 to 117 are invention steels that satisfy the requirements of the component composition and production conditions defined in the present invention. These materials have good mechanical properties such as yield strength at normal temperature, tensile strength and 0 ° C. impact absorption energy, and 0.2% proof stress at 700 ° C., and satisfy the target mechanical properties. On the other hand, aa to ag, which are comparative steels, have properties that do not satisfy the target, since each component composition is outside the scope of the present invention.

  Since the comparative steel aa has less Mn and V than the range of the present invention, the tensile strength at normal temperature is lowered. Since the comparative steel ab has less C, V and Cr than the range of the present invention, the normal temperature and 700 ° C. strength are insufficient, and since S and N are larger than the range of the present invention, the toughness is also lowered. Since the comparative steel ac has more Mn and less Cr than the scope of the present invention, the normal temperature tensile strength and the 700 ° C. 0.2% proof stress are lowered. Since the comparative steel ad has more Mn and less V and Nb than the range of the present invention, the normal temperature strength and the 700 ° C. 0.2% proof stress are lowered.

  The comparative steel ae is an example in which Cr is less than the range of the present invention, but C, Si, Mn, V and Nb are large, and the normal temperature strength exceeds the upper limit of the standard. The comparative steel af has more C, S, N, and Ti than the range of the present invention, the normal temperature tensile strength and the 700 ° C. 0.2% proof stress do not reach the targets, and the toughness is also lowered. In the comparative steel ag, C, Mn, S, and Nb are excessive, the room temperature tensile strength exceeds the upper limit of specification, and the 0 ° C. Charpy impact absorption energy is less than 100J. In addition, although comparative steel ag has insufficient V, since Nb which contributes to precipitation strengthening is included excessively, 700 degreeC 0.2% yield strength has reached the target.

Table 5 shows the steel No. which is within the range of the component composition defined in the reference example or the present invention. Nos. 108 to 110 are used to show examples in which the production conditions do not satisfy the requirements of the present invention. Component is in the range of the present invention, even 1 10, when the cooling rate of the accelerated cooling after rolling is small, if the water cooling stop temperature is 300 ° C. or less insufficient strength or toughness shortage is generated.

  As described above, according to the present invention, in the process of manufacturing an H-section steel, an alloy carbonitriding mainly composed of V under an appropriate balance of added amounts of V, C, and N in a component system in which Mo is not added. A steel material excellent in fire resistance having required high temperature strength and mechanical properties at room temperature can be provided by forming an article and adding an alloy element as appropriate according to the required strength level. Among the above steel materials, steel materials represented by H-shaped steel are particularly useful as structural materials 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.

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 (5)

% By mass
C: 0.03-0.10%,
Si: 0.05 to 0.50%,
Mn: 0.05 to 0.4%,
Al ≦ 0.01%
V: 0.15-0.35%,
Cr: 1.5-3%,
N: 0.002 to 0.008%,
S: 0.003 to 0.02%,
P: 0.0001 to 0.1%
A steel material excellent in fire resistance, characterized by comprising iron and the inevitable impurities. A method of manufacturing a steel material by hot rolling after reheating a steel slab having the component according to claim 1 ,
(A) After reheating to 1100-1300 ° C., hot rolling is started,
(B) After rolling,
(C) The method for producing a steel material having excellent fire resistance according to claim 1, wherein the steel material is allowed to cool or is cooled after accelerated cooling. The method for producing a steel material having excellent fire resistance according to claim 2, wherein the surface of the steel material is water-cooled to 700 ° C or lower and a water-cooling / rolling cycle for rolling in the reheating process is performed once or more. The method for producing a steel material excellent in fire resistance according to claim 2 or 3 , wherein accelerated cooling is performed from the rolling end temperature to 600 ° C at an average cooling rate of 0.5 to 5.0 ° C / s. A method of manufacturing a steel material by hot rolling after reheating a steel slab having the component according to claim 1 ,
(A) After reheating to 1100-1300 ° C., hot rolling is started,
(B) The rolling end temperature is 850 ° C. or more at the steel surface temperature,
(C) After rolling, the average cooling rate in the range from 800 ° C. to 500 ° C. is accelerated to 3 to 15 ° C./s, the stop temperature is set to the range of 300 to 550 ° C. at the surface temperature of the steel material, and then allowed to cool. The method for producing a steel material having excellent fire resistance according to claim 1 .

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