JP5157072B2 - Manufacturing method of high strength and high toughness thick steel plate with excellent tensile strength of 900 MPa and excellent in cutting crack resistance - Google Patents
Manufacturing method of high strength and high toughness thick steel plate with excellent tensile strength of 900 MPa and excellent in cutting crack resistance Download PDFInfo
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本発明は，天然ガスや原油の輸送用として用いられる引張強度が９００ＭＰａ以上の高強度・高靱性のラインパイプ用厚鋼板の製造方法に関し，特に，せん断加工での切断の際，切断面での耐切断割れ性に優れるものに関する。 The present invention relates to a process for the preparation of natural gas and crude oil tensile strength to be used for the transport 900MPa or more high strength and high toughness thick steel plate for a line pipe, in particular, upon cleavage at shearing, cutting plane It relates to a material having excellent cutting crack resistance.
近年，天然ガスや原油の輸送用として使用されるラインパイプは，高圧化による輸送効率の向上や薄肉化による現地溶接施工能率の向上のため，年々高強度化され、既にＡＰＩ規格でＸ１００グレードのラインパイプが実用化されているが，さらに，引張強度９００ＭＰａを超えるＸ１２０グレードに対する要求が具体化されつつある。 In recent years, line pipes used for transportation of natural gas and crude oil have been strengthened year by year in order to improve transport efficiency by increasing pressure and to improve local welding efficiency by reducing wall thickness, and already have X100 grade of API standard. Although the line pipe has been put into practical use, the demand for the X120 grade exceeding the tensile strength of 900 MPa is being realized.
このような高強度ラインパイプ用溶接鋼管用の厚鋼板の製造方法に関し，例えば特許文献１においては，熱間圧延後２段冷却を行い，２段目の冷却停止温度を３００℃以下とすることで，高強度化を達成する技術が開示されている。また，特許文献２においては，Ｃｕ析出強化を利用した高強度化のための加速冷却＋時効熱処理条件に関する技術が開示されている。
しかしながら，特許文献１記載のように、冷却停止温度を低くして，低温変態生成する硬質なベイナイトあるいはマルテンサイト組織を導入することで高強度化を達成した場合，冷却ままの鋼板を必要なサイズにせん断加工で切断する際，鋼中に残存する拡散性水素が原因で板面に平行な割れ（以降、切断割れと称する）が発生する。 However, as described in Patent Document 1, when high strength is achieved by introducing a hard bainite or martensite structure that generates a low-temperature transformation by lowering the cooling stop temperature, a steel plate that has been cooled is the required size. When cutting by shearing, cracks parallel to the plate surface (hereinafter referred to as cutting cracks) occur due to diffusible hydrogen remaining in the steel.
一方，特許文献２のように，加速冷却後に熱処理を行った場合，鋼中の水素は十分拡散させられるので，切断割れは抑制できるものの，熱処理過程でミクロ組織中にセメンタイトが析出・粗大化し，靱性低下，特に脆性亀裂伝播停止特性の評価を行うＤＷＴＴ（ＤｒｏｐＷｅｉｇｈｔ Ｔｅａｒ Ｔｅｓｔ）特性が劣化する。 On the other hand, as in Patent Document 2, when heat treatment is performed after accelerated cooling, hydrogen in the steel is sufficiently diffused, so that cracking can be suppressed, but cementite precipitates and coarsens in the microstructure during the heat treatment process. DWTT (Drop Weight Tear Test) characteristics for evaluating toughness reduction, particularly brittle crack propagation stop characteristics, deteriorate.
また，天然ガス輸送用ラインパイプにおいては，外因性の事故により発生した延性破壊の亀裂伝播停止を目的として，高い母材シャルピー吸収エネルギー値や、脆性亀裂発生防止の観点から優れたＣＴＯＤ（Ｃｒａｃｋ Ｔｉｐ Ｏｐｅｎｉｎｇ Ｄｉｓｐｌａｃｅｍｅｎｔ）値といった様々な靭性要求を満足しなくてはならない。 In addition, natural gas transport line pipes have a high CTOD (Cracking Tip) value from the standpoint of high base metal Charpy absorbed energy value and prevention of brittle cracking for the purpose of stopping crack propagation of ductile fracture caused by exogenous accidents. Various toughness requirements such as Opening Displacement values must be satisfied.
本発明は，種々の靭性要求を満足し、且つ耐切断割れ性を改善した高強度厚鋼板の製造方法を提供することを目的とする。 An object of the present invention is to provide a method for producing a high-strength thick steel plate that satisfies various toughness requirements and has improved cut cracking resistance.
発明者らは，加速冷却ままの高強度厚鋼板の切断割れについて，鋭意研究を重ねた結果，切断割れについては以下の１）、２）、セメンタイトの粗大化挙動については３）、ＤＷＴＴ特性については４）〜６）の知見を得た。
１）鋼中の拡散性水素が，各トラップサイトにトラップされることを阻止するために，少なくとも３００℃以上での脱水素の熱処理が必要である。
As a result of extensive research on cutting cracks in high-strength thick steel plates with accelerated cooling, the inventors have conducted the following 1), 2) for cutting cracks, 3) for coarsening behavior of cementite, and DWTT characteristics. Obtained the findings of 4) to 6).
1) In order to prevent diffusible hydrogen in steel from being trapped at each trap site, a heat treatment for dehydrogenation at least at 300 ° C. or higher is required.
２）加速冷却停止後，ただちに再加熱を開始し，鋼板温度を３００℃以上に昇温することで，水素の拡散が促進される結果，鋼中に残留する水素が割れ発生限界量を下回る。 2) Immediately after the accelerated cooling is stopped, reheating is started and the steel sheet temperature is raised to 300 ° C. or higher, so that hydrogen diffusion is promoted.
３）再加熱時の加熱速度を速くすることで，３００〜５００℃といった温度域に加熱してもセメンタイトの粗大化が抑制され，ＤＷＴＴ特性が劣化しなくなる。 3) By increasing the heating rate during reheating, coarsening of cementite is suppressed even when heated to a temperature range of 300 to 500 ° C., and the DWTT characteristics do not deteriorate.
４）熱間圧延時におけるオーステナイト未再結晶域圧延の強化が有効であるが、大圧下を行うとシャルピー吸収エネルギーおよびＣＴＯＤ値の低下が生じる。 4) Strengthening of austenite non-recrystallized region rolling at the time of hot rolling is effective, but when large reduction is performed, Charpy absorbed energy and CTOD value are lowered.
５）オーステナイト未再結晶域圧延を強化した材料は、シャルピー試験あるいはＣＴＯＤ試験時にセパレーションと呼ばれる板面と平行な面が剥離するような亀裂発生が生じ、その結果シャルピー試験でのシャルピー吸収エネルギーおよびＣＴＯＤ値が低下する。 5) In the material strengthened with austenite non-recrystallized zone rolling, cracking occurs in the Charpy test or CTOD test where the plane parallel to the plate surface peels off, resulting in Charpy absorbed energy and CTOD in the Charpy test. The value drops.
６）セパレーションを抑制しつつ、優れたＤＷＴＴ特性を得るためにはオーステナイト未再結晶域における圧延温度、特に圧延終了温度を厳格に制御することが有効である。 6) In order to obtain excellent DWTT characteristics while suppressing separation, it is effective to strictly control the rolling temperature in the austenite non-recrystallized region, particularly the rolling end temperature.
本発明は得られた知見を基に更に検討を加えてなされたもので、すなわち、本発明は、
１．質量％で、
Ｃ：０．０３〜０．１２％
Ｓｉ：０．０１〜０．５％
Ｍｎ：１．５〜３．０％
Ａｌ：０．０１〜０．０８％
Ｎｂ：０．０１〜０．０８％
Ｔｉ：０．００５〜０．０２５％
Ｎ：０．００１〜０．０１％
Ｂ：０．０００５〜０．００３％以下
更に
Ｃｕ：０．０１〜２％
Ｎｉ：０．０１〜３％
Ｃｒ：０．０１〜１％
Ｍｏ：０．０１〜１％
Ｖ：０．０１〜０．１％
の一種または二種以上を含有し、残部Ｆｅ及び不可避的不純物からなる鋼を，
１０００〜１２００℃に加熱後，熱間圧延を開始し，
９５０℃以下の温度域での累積圧下量≧７０％となるよう圧延を行い，
８００℃以上で圧延を終了し，その後
冷却開始温度７００℃以上，８００℃未満より，
冷却速度２０〜８０℃／ｓの加速冷却を開始し，
２５０℃未満の温度域で冷却停止し，
冷却停止後ただちに５℃／ｓ以上の昇温速度で
３００℃以上５００℃以下の温度に再加熱することを特徴とする，耐切断割れ性に優れた
引張強度９００ＭＰａ以上の高強度・高靱性厚鋼板の製造方法。
The present invention was made by further study based on the obtained knowledge, that is, the present invention is
1 . % By mass
C: 0.03-0.12%
Si: 0.01 to 0.5%
Mn: 1.5 to 3.0%
Al: 0.01 to 0.08%
Nb: 0.01 to 0.08%
Ti: 0.005-0.025%
N: 0.001 to 0.01%
B: 0.0005 to 0.003% or less Cu: 0.01 to 2%
Ni: 0.01 to 3%
Cr: 0.01 to 1%
Mo: 0.01 to 1%
V: 0.01 to 0.1%
A steel containing one or more of the following, the balance Fe and unavoidable impurities,
After heating to 1000-1200 ° C, hot rolling is started,
Rolling is performed so that the cumulative reduction amount ≧ 70% in the temperature range of 950 ° C. or lower,
Rolling is completed at 800 ° C or higher, and then the cooling start temperature is 700 ° C or higher and lower than 800 ° C.
Accelerated cooling is started at a cooling rate of 20 to 80 ° C./s,
Stop cooling in the temperature range below 250 ℃,
Immediately after stopping cooling, it is reheated to a temperature not lower than 300 ° C and not higher than 500 ° C at a rate of temperature increase of 5 ° C / s or more.
A method for producing a high-strength, high-toughness thick steel plate having a tensile strength of 900 MPa or more .
２．１記載の成分に加え，質量％で、
Ｃａ：０．０００５〜０．０１％
ＲＥＭ：０．０００５〜０．０２％
Ｚｒ：０．０００５〜０．０３％
Ｍｇ：０．０００５〜０．０１％
の一種あるいは二種以上を含有し
残部Ｆｅ及び不可避的不純物からなる鋼を，
１０００〜１２００℃に加熱後，熱間圧延を開始し，
９５０℃以下の温度域での累積圧下量≧７０％となるよう圧延を行い，
８００℃以上，９００℃以下の温度で圧延を終了し，ひきつづき
７００℃以上，８００℃未満の温度から，
冷却速度２０〜８０℃／ｓの加速冷却を開始し，
２５０℃未満の温度域で冷却停止し，
冷却停止後ただちに５℃／ｓ以上の昇温速度で
３００℃以上５００℃以下の温度に再加熱することを特徴とする，耐切断割れ性に優れた引張強度９００ＭＰａ以上の高強度・高靱性厚鋼板の製造方法。
2 . In addition to the ingredients described in 1,
Ca: 0.0005 to 0.01%
REM: 0.0005 to 0.02%
Zr: 0.0005 to 0.03%
Mg: 0.0005 to 0.01%
A steel containing one or more of the following, the balance Fe and unavoidable impurities,
After heating to 1000-1200 ° C, hot rolling is started,
Rolling is performed so that the cumulative reduction amount ≧ 70% in the temperature range of 950 ° C. or lower,
Rolling is finished at a temperature of 800 ° C. or higher and 900 ° C. or lower, and subsequently from a temperature of 700 ° C. or higher and lower than 800 ° C.,
Accelerated cooling is started at a cooling rate of 20 to 80 ° C./s,
Stop cooling in the temperature range below 250 ℃,
Immediately after the cooling is stopped, it is reheated to a temperature of 300 ° C. or more and 500 ° C. or less at a temperature rising rate of 5 ° C./s or more. A method of manufacturing a steel sheet.
本発明により，切断割れの発生防止に優れる，引張強度９００ＭＰａ以上の高強度・高靱性を備えた厚鋼板の製造が可能で、産業上極めて有用である。 According to the present invention, it is possible to produce a thick steel plate having a high strength and a high toughness with a tensile strength of 900 MPa or more, which is excellent in preventing the occurrence of cutting cracks, which is extremely useful in industry.
本発明では成分組成、製造条件を規定する。鋼板の成分限定理由を説明する。尚、本発明で、高強度とは引張強度９００ＭＰａ以上、高靱性とは、ＤＷＴＴ特性（試験温度−３０℃での脆性破面率７５％以上）、試験温度−３０℃でのシャルピー吸収エネルギ−２００Ｊ以上、試験温度−２０℃での限界開口変位量δｃ：０．１５ｍｍ以上、厚鋼板とは板厚６ｍｍ以上とする。 In this invention, a component composition and manufacturing conditions are prescribed | regulated. The reason for limiting the components of the steel sheet will be described. In the present invention, high strength means a tensile strength of 900 MPa or more, and high toughness means a DWTT characteristic (a brittle fracture surface ratio of 75% or more at a test temperature of −30 ° C.), a Charpy absorbed energy at a test temperature of −30 ° C. The limit opening displacement amount δc at a test temperature of −20 ° C. is 200 J or more and 0.15 mm or more, and the thick steel plate is 6 mm or more.
［成分組成］％は質量％とする。
Ｃ：０．０３〜０．１２％
Ｃは低温変態組織においては過飽和固溶することで強度上昇に寄与する。これらの効果を得るためには０．０３％以上の添加が必要であるが，０．１２％を超えて添加すると，
パイプの円周溶接部の硬度上昇が著しくなり，溶接低温割れが発生しやすくなるため，上限を０．１２％とする。
[Component composition]% is mass%.
C: 0.03-0.12%
C contributes to an increase in strength by being supersaturated in a low temperature transformation structure. In order to obtain these effects, addition of 0.03% or more is necessary, but when adding over 0.12%,
The upper limit is set to 0.12% because the increase in the hardness of the circumferential welded part of the pipe becomes significant and the cold cracking of the weld tends to occur.
Ｓｉ：０．０１〜０．５％
Ｓｉは脱酸材として作用し，さらに固溶強化により鋼材の強度を増加させる元素であるが、０．０１％以下ではその効果がなく，０．５％を超えて添加すると靱性が著しく低下
するため上限を０．５％とする。
Si: 0.01 to 0.5%
Si is an element that acts as a deoxidizer and increases the strength of the steel by solid solution strengthening. However, it is not effective at 0.01% or less, and if added over 0.5%, the toughness is significantly reduced. Therefore, the upper limit is made 0.5%.
Ｍｎ：１．５〜３．０％
Ｍｎは焼入性向上元素として作用する。１．５％以上の添加によりその効果は得られるが，連続鋳造プロセスでは中心偏析部の濃度上昇が著しく，３．０％を超える添加を行うと，偏析部での遅れ破壊の原因となるため，上限を３．０％とする。
Mn: 1.5 to 3.0%
Mn acts as a hardenability improving element. The effect can be obtained by addition of 1.5% or more, but the concentration in the central segregation part is remarkably increased in the continuous casting process, and if addition over 3.0% causes delayed fracture in the segregation part. The upper limit is set to 3.0%.
Ａｌ：０．０１〜０．０８％
Ａｌは脱酸元素として作用する。０．０１％以上の添加で十分な脱酸効果が得られるが，０．０８％を超えて添加すると鋼中の清浄度が低下し，靱性劣化の原因となるため，上限を０．０８％とする。
Al: 0.01 to 0.08%
Al acts as a deoxidizing element. Sufficient deoxidation effect can be obtained with addition of 0.01% or more, but if added over 0.08%, the cleanliness in the steel is lowered and the toughness is deteriorated, so the upper limit is 0.08%. And
Ｎｂ：０．０１〜０．０８％
Ｎｂは熱間圧延時のオーステナイト未再結晶領域を拡大する効果があり，特に９５０℃以下を未再結晶領域とするため、０．０１％以上添加する。一方，０．０８％を超えて添加すると，ＨＡＺの靱性を著しく損ねることから上限を０．０８％とする。
Nb: 0.01 to 0.08%
Nb has an effect of expanding the austenite non-recrystallized region at the time of hot rolling. In particular, Nb is added in an amount of 0.01% or more to make the non-recrystallized region at 950 ° C. or less. On the other hand, if added over 0.08%, the toughness of the HAZ is remarkably impaired, so the upper limit is made 0.08%.
Ｔｉ：０．００５〜０．０２５％
Ｔｉは窒化物を形成し，鋼中の固溶Ｎ量低減に有効であるほか，析出したＴｉＮがピンニング効果でオーステナイト粒の粗大化抑制防止をすることで，母材，ＨＡＺの靱性向上に寄与する。
Ti: 0.005-0.025%
Ti forms nitrides and is effective in reducing the amount of solute N in the steel. Precipitated TiN prevents the austenite grains from becoming coarse by the pinning effect, contributing to improved toughness of the base metal and HAZ. To do.
必要なピンニング効果を得るためには０．００５％以上の添加が必要であるが，０．０２５％を超えて添加すると炭化物を形成するようになり，その析出硬化で靱性が著しく劣化するため，上限を０．０２５％とする。 Addition of 0.005% or more is necessary to obtain the required pinning effect, but if added over 0.025%, carbides are formed, and the toughness deteriorates significantly due to precipitation hardening. The upper limit is 0.025%.
Ｎ：０．００１〜０．０１％
Ｎは通常鋼中の不可避不純物として存在するが，前述の通りＴｉ添加を行うことで，オーステナイト粒の粗大化を抑制するＴｉＮを形成する。
N: 0.001 to 0.01%
N is usually present as an unavoidable impurity in steel, but TiN is added to form TiN that suppresses coarsening of austenite grains as described above.
必要とするピンニング効果を得るためには０．００１％以上鋼中に存在することが必要であるが，０．０１％を超える場合，溶接部，特に溶融線近傍で１４５０℃以上に加熱されたＨＡＺでＴｉＮが分解し，固溶Ｎの悪影響が著しいため，上限を０．０１％とする。 In order to obtain the required pinning effect, 0.001% or more must be present in the steel. However, if it exceeds 0.01%, it was heated to 1450 ° C or more near the weld, particularly in the vicinity of the melting line. The upper limit is set to 0.01% because TiN decomposes in HAZ and the negative effect of solute N is significant.
Ｂ：０．０００５〜０．００３％
Ｂはオーステナイト粒界に偏析し、特に０．０００５％以上の添加でフェライト変態が抑制され、フェライト生成により２相組織となりシャルピー吸収エネルギーが低下することを防止する。一方、０．００３％を超えて添加しても効果が飽和するため、上限を０．００３％とする。
B: 0.0005 to 0.003%
B segregates at the austenite grain boundary, and when ferrite is added in an amount of 0.0005% or more, the ferrite transformation is suppressed and a two-phase structure is formed by the formation of ferrite, thereby preventing a decrease in Charpy absorbed energy. On the other hand, since the effect is saturated even if added over 0.003%, the upper limit is made 0.003%.
Ｃｕ，Ｎｉ，Ｃｒ，Ｍｏ，Ｖの一種または二種以上
Ｃｕ，Ｎｉ，Ｃｒ，Ｍｏ，Ｖはいずれも焼入性向上元素として作用するため、高強度化を目的に、これらの元素の一種，または二種以上を添加する。
One or more of Cu, Ni, Cr, Mo, and V Cu, Ni, Cr, Mo, and V all act as a hardenability improving element. Therefore, for the purpose of increasing the strength, Or two or more of them are added.
Ｃｕ：０．０１〜２％
Ｃｕは，０．０１％以上添加することで鋼の焼入性向上に寄与する。しかし，２％以上の添加を行うと，靱性劣化が生じるため，上限を２％とする。尚，０．８％以上添加した場合，時効熱処理による析出強化が著しく溶接熱影響部の軟化防止にも寄与するため、好ましくは０．８〜２％とする。
Cu: 0.01-2%
Cu contributes to the improvement of hardenability of steel by adding 0.01% or more. However, if 2% or more is added, toughness deterioration occurs, so the upper limit is made 2%. If 0.8% or more is added, precipitation strengthening due to aging heat treatment remarkably contributes to prevention of softening of the heat affected zone, so 0.8 to 2% is preferable.
Ｎｉ：０．０１〜３％
Ｎｉは０．０１％以上添加することで鋼の焼入性向上に寄与する。特に，多量に添加しても靱性劣化を生じないため，強靱化に有効であるが，高価な元素であり，かつ３％を超えて添加しても強度上昇が飽和するため，上限を３％とする。
Ni: 0.01 to 3%
Ni contributes to improving the hardenability of steel by adding 0.01% or more. In particular, even if added in a large amount, it does not cause toughness deterioration, so it is effective for toughening, but it is an expensive element, and even if added over 3%, the increase in strength is saturated, so the upper limit is 3% And
Ｃｒ：０．０１〜１％
Ｃｒもまた０．０１％以上添加することで鋼の焼入性向上に寄与する。一方，１％を超えて添加すると，靱性が劣化するため，上限を１％とする。
Cr: 0.01 to 1%
Cr also contributes to improving the hardenability of steel by adding 0.01% or more. On the other hand, if added over 1%, the toughness deteriorates, so the upper limit is made 1%.
Ｍｏ：０．０１〜１％
Ｍｏもまた０．０１％以上添加することで鋼の焼入性向上に寄与する。一方，１％を超えて添加すると，靱性が劣化するため，上限を１％とする。
Mo: 0.01 to 1%
Mo also contributes to improving the hardenability of steel by adding 0.01% or more. On the other hand, if added over 1%, the toughness deteriorates, so the upper limit is made 1%.
Ｖ：０．０１〜０．１％
Ｖは炭窒化物を形成することで析出強化し，特に溶接熱影響部の軟化防止に寄与する。０．０１％以上の添加によりこの効果が得られるが，０．１％を超えて添加すると，析出強化が著しく靱性が低下するため，上限を０．１％とする。
V: 0.01 to 0.1%
V strengthens precipitation by forming carbonitrides, and contributes particularly to the prevention of softening of the heat affected zone. This effect can be obtained by addition of 0.01% or more, but if added over 0.1%, precipitation strengthening remarkably reduces toughness, so the upper limit is made 0.1%.
以上が本発明に係る鋼の基本成分組成であるが、溶接部の靭性を更に向上させる場合、Ｃａ，ＲＥＭ，Ｚｒ，Ｍｇの一種または二種以上を添加する。 The above is the basic component composition of the steel according to the present invention. When the toughness of the weld is further improved, one or more of Ca, REM, Zr, and Mg are added.
Ｃａ，ＲＥＭ，Ｚｒ，Ｍｇ
Ｃａ，ＲＥＭ，Ｚｒ，Ｍｇは鋼中の非金属介在物であるＭｎＳの形態制御、あるいは酸化物、窒化物を形成し，主に溶接熱影響部におけるオーステナイト粒粗大化をピンニング効果で抑制し，靱性を向上させる目的で添加する。
Ca, REM, Zr, Mg
Ca, REM, Zr, and Mg form MnS, which is a non-metallic inclusion in steel, or form oxides and nitrides, mainly suppressing the austenite grain coarsening in the heat affected zone by the pinning effect. Add for the purpose of improving toughness.
Ｃａ：０．０００５〜０．０１％
Ｃａは鋼中の硫化物の形態制御に有効な元素であり、０．０００５％以上添加することで靭性に有害なＭｎＳの形成を抑制する。一方、０．０１％を超えた場合、ＣａＯーＣａＳのクラスタを形成し、母材を含めて靱性が低下するうえに、取鍋のノズル閉塞の原因となり、生産性を阻害するため、上限は０．０１％とし、添加する場合は、０．０００５〜０．０１％とする。
Ca: 0.0005 to 0.01%
Ca is an element effective for controlling the form of sulfide in steel, and by adding 0.0005% or more, formation of MnS harmful to toughness is suppressed. On the other hand, when the content exceeds 0.01%, a CaO-CaS cluster is formed, and the toughness including the base material is reduced. Further, it causes nozzle clogging of the ladle and inhibits the productivity. 0.01%, and when added, 0.0005 to 0.01%.
ＲＥＭ：０．０００５〜０．０２％
ＲＥＭは鋼中で酸硫化物を形成し，０．０００５％以上添加することで溶接熱影響部の粗大化を防止するピンニング効果をもたらす。しかし，高価な元素であり，かつ０．０２％を超えて添加しても効果が飽和するため，上限を０．０２％とし、添加する場合は、０．０００５〜０．０２％とする。
REM: 0.0005 to 0.02%
REM forms an oxysulfide in steel, and adding 0.0005% or more provides a pinning effect that prevents the weld heat affected zone from becoming coarse. However, since it is an expensive element and the effect is saturated even if added over 0.02%, the upper limit is made 0.02%, and when added, it is made 0.0005 to 0.02%.
Ｚｒ：０．０００５〜０．０３％
Ｚｒは鋼中で炭窒化物を形成し，とくに溶接熱影響部においてオーステナイト粒の粗大化を抑制するピンニング効果をもたらす。十分なピンニング効果を得るためには，０．０００５％以上の添加が必要であるが，０．０３％を超えて添加すると，鋼中の清浄度が著しく低下し，靱性が低下するようになるため，上限を０．０３％とし、添加する場合は、０．０００５〜０．０３％とする。
Zr: 0.0005 to 0.03%
Zr forms carbonitrides in steel and brings about a pinning effect that suppresses the coarsening of austenite grains, particularly in the weld heat affected zone. In order to obtain a sufficient pinning effect, addition of 0.0005% or more is necessary, but if added over 0.03%, the cleanliness in the steel is remarkably lowered and the toughness is lowered. Therefore, the upper limit is made 0.03%, and when added, the content is made 0.0005 to 0.03%.
Ｍｇ：０．０００５〜０．０１％
Ｍｇは製鋼過程で鋼中に微細な酸化物として生成し，特に，溶接熱影響部においてオーステナイト粒の粗大化を抑制するピンニング効果をもたらす。
Mg: 0.0005 to 0.01%
Mg is produced as fine oxides in the steel during the steel making process, and has a pinning effect that suppresses austenite grain coarsening, particularly in the weld heat affected zone.
十分なピンニング効果を得るためには，０．０００５％以上の添加が必要であるが，０．０１％を超えて添加すると，鋼中の清浄度が低下し，靱性が低下するようになるため，上限を０．０１％とし、添加する場合は、０．０００５〜０．０１％とする。 In order to obtain a sufficient pinning effect, addition of 0.0005% or more is necessary. However, if it exceeds 0.01%, the cleanliness in the steel decreases and the toughness decreases. The upper limit is 0.01%, and when added, the content is 0.0005 to 0.01%.
［製造条件］
加熱温度：１０００〜１２００℃
熱間圧延する際、鋼片全体をオーステナイト化するため、１０００℃以上に加熱する。一方，１２００℃を超える温度まで鋼片を加熱すると，ＴｉＮピンニングによっても、オーステナイト粒成長が著しく，母材靱性が劣化するため，１２００℃以下とする。
[Production conditions]
Heating temperature: 1000-1200 ° C
When hot rolling, in order to austenite the entire steel slab, it is heated to 1000 ° C. or higher. On the other hand, when the steel slab is heated to a temperature exceeding 1200 ° C., the austenite grain growth is remarkably caused by TiN pinning, and the base material toughness deteriorates.
９５０℃以下での累積圧下量≧７０％
前述の通り，Ｎｂ添加によって９５０℃以下はオーステナイト未再結晶域である。この温度域にて累積で大圧下を行うことにより，オーステナイト粒が伸展し特に板厚方向では細粒となり、この状態で加速冷却して得られる鋼のＤＷＴＴ特性は良好となる。
Cumulative reduction at 950 ° C or lower ≥ 70%
As described above, 950 ° C. or lower is an austenite non-recrystallized region due to Nb addition. By accumulating large pressures in this temperature range, austenite grains are expanded and become fine grains particularly in the thickness direction, and the DWTT characteristics of steel obtained by accelerated cooling in this state are improved.
圧下量が７０％未満では，細粒化効果は不十分でＤＷＴＴ特性向上が得られないため，累積圧下量の下限を７０％とする。尚、著しく靱性向上を狙うための好適範囲は７５％以上である。 If the reduction amount is less than 70%, the effect of refining is insufficient and the DWTT characteristics cannot be improved. Therefore, the lower limit of the cumulative reduction amount is set to 70%. In addition, the suitable range for aiming at a remarkable toughness improvement is 75% or more.
圧延終了温度：８００℃以上
オーステナイト未再結晶域での累積大圧下は、ＤＷＴＴ特性の向上に有効であるが、８００℃未満の低温域での圧延はオーステナイト粒に集合組織が形成され、加速冷却後の変態組織にも受け継がれてセパレーションを生じやすくし、シャルピー吸収エネルギーやＣＴＯＤ値を低下させるため、圧延終了温度は８００℃以上とする。
Rolling end temperature: 800 ° C or higher Cumulative large pressure in the austenite non-recrystallized region is effective for improving the DWTT characteristics, but rolling in the low temperature region below 800 ° C forms a texture in the austenite grains and accelerates cooling. The rolling end temperature is set to 800 ° C. or higher in order to easily inherit separation in the later transformation structure and reduce Charpy absorbed energy and CTOD value.
圧延終了温度の上限は特に定めないが、圧延中の温度上昇により再結晶と未再結晶の境界温度である９５０℃以上となることを防止するため、好ましくは９００℃以下とする。 The upper limit of the rolling end temperature is not particularly defined, but is preferably set to 900 ° C. or lower in order to prevent the boundary temperature between recrystallization and non-recrystallization from becoming 950 ° C. or higher due to temperature rise during rolling.
加速冷却の冷却開始温度：７００℃以上８００℃未満
加速冷却開始温度は７００℃以上とする。加速冷却開始温度が低いと、熱間圧延後、加速冷却開始までの空冷過程においてオーステナイト粒界から初析フェライトが生成し、母材強度が低下するようになるので、７００℃以上とする。
Accelerated cooling start temperature: 700 ° C. or higher and lower than 800 ° C. The accelerated cooling start temperature is 700 ° C. or higher. If the accelerated cooling start temperature is low, proeutectoid ferrite is generated from the austenite grain boundaries in the air cooling process after hot rolling until the accelerated cooling starts, and the base material strength is lowered.
一方、８００℃以上の温度から加速冷却を行うと、オーステナイト未再結晶域での強圧下で形成されたオーステナイトの集合組織が回復する前に、ベイナイト変態、マルテンサイト変態が生じ、セパレーションが生じやすく、シャルピー吸収エネルギーおよびＣＴＯＤ値が低下するため、８００℃未満とし、好ましくは７５０℃以下とする。 On the other hand, when accelerated cooling is performed from a temperature of 800 ° C. or higher, bainite transformation and martensite transformation occur before the austenite aggregate structure formed under strong pressure in the austenite non-recrystallized region is recovered, and separation is likely to occur. Since Charpy absorbed energy and CTOD value are lowered, the temperature is set to less than 800 ° C., preferably 750 ° C. or less.
加速冷却の冷却速度：２０〜８０℃／ｓ
加速冷却の冷却速度は２０℃／ｓ以上とする。加速冷却の冷却速度が遅い場合、フェライト変態が生じ、鋼板の強度が低下する。一方，８０℃／ｓを超える冷却速度としたとき，特に鋼板表面近傍ではマルテンサイト変態が生じ，鋼板強度は上昇するものの，靱性劣化，特にシャルピー吸収エネルギー低下が著しいため，冷却速度の上限を８０℃／ｓとする。
Accelerated cooling rate: 20-80 ° C / s
The cooling rate for accelerated cooling is 20 ° C./s or more. When the cooling rate of accelerated cooling is slow, ferrite transformation occurs and the strength of the steel sheet is reduced. On the other hand, when the cooling rate exceeds 80 ° C./s, martensite transformation occurs particularly near the steel plate surface, and the strength of the steel plate increases. However, the toughness deterioration, particularly the Charpy absorption energy decrease, is significant, so the upper limit of the cooling rate is 80 It is set to ° C / s.
加速冷却の冷却停止温度：≦２５０℃
鋼板の高強度化のため，加速冷却の停止温度を下げて，低温で変態するベイナイトやマルテンサイト組織を生成させる。冷却停止温度が２５０℃を超える温度の場合，靱性が低い上部ベイナイト組織となるため，冷却停止温度は２５０℃以下とする。
Cooling stop temperature for accelerated cooling: ≤250 ° C
In order to increase the strength of the steel sheet, the stop temperature of accelerated cooling is lowered to produce bainite and martensite structures that transform at low temperatures. When the cooling stop temperature is higher than 250 ° C, the upper bainite structure is low in toughness, so the cooling stop temperature is 250 ° C or lower.
再加熱処理
加速冷却で低温変態させて高強度化させた鋼板は，加速冷却後、空冷させても鋼中の拡散性水素が残留し，切断割れが生じることがある。そこで，冷却停止後、速やかに再加熱を行う。再加熱方法は，炉加熱，誘導加熱いずれでもかまわない。
Reheating treatment Steel plates that have undergone low-temperature transformation by accelerated cooling to increase strength may cause diffusible hydrogen in the steel to remain and cause cracking even if air-cooled after accelerated cooling. Therefore, reheating is performed immediately after cooling is stopped. The reheating method may be either furnace heating or induction heating.
再加熱温度：３００℃〜５００℃
再加熱温度が３００℃未満の場合，十分水素が拡散せず，切断割れを防止することができないため，再加熱温度は３００℃以上とする。一方，５００℃を超える温度まで加熱すると，焼き戻しによる軟化で強度低下が著しいことから，上限を５００℃とする。
Reheating temperature: 300 ° C to 500 ° C
When the reheating temperature is less than 300 ° C., hydrogen does not diffuse sufficiently and cutting cracks cannot be prevented. Therefore, the reheating temperature is set to 300 ° C. or higher. On the other hand, when heated to a temperature exceeding 500 ° C., the upper limit is set to 500 ° C. because the strength is significantly reduced due to softening by tempering.
再加熱時の昇温速度：≧５℃／ｓ
再加熱時の昇温速度が５℃／ｓ未満の場合，特に３００℃を超えるような温度まで加熱する途中でセメンタイトが生成，粗大化するため，ＤＷＴＴ特性の劣化が著しい。再加熱時の昇温速度を５℃／ｓ以上とし、セメンタイトの粗大化を抑制する。
Heating rate during reheating: ≧ 5 ° C / s
When the heating rate at the time of reheating is less than 5 ° C./s, cementite is generated and coarsened particularly during heating to a temperature exceeding 300 ° C., so that the DWTT characteristics are remarkably deteriorated. The heating rate during reheating is set to 5 ° C./s or more to suppress cementite coarsening.
再加熱までの時間が長いと，その間の空冷過程での温度低下によって水素が拡散しにくくなるため，加熱開始は好ましくは３００秒以内、更に好ましくは１００秒以内とする。 If the time until reheating is long, hydrogen becomes difficult to diffuse due to a temperature drop during the air cooling process, so the heating start is preferably within 300 seconds, more preferably within 100 seconds.
［ミクロ組織］
引張強度≧９００ＭＰａの高強度を達成するためには低温変態して転位密度の高い組織とする必要がある．また，このような硬質なミクロ組織中に少量のフェライトやパーライトが混入すると母材ＣＴＯＤ値がポップインの発生によりばらつきやすくなることから，本発明における鋼のミクロ組織は，ベイナイト，マルテンサイトあるいはベイナイト＋マルテンサイト２相組織に限定する．
本発明において鋼の製鋼方法については特に限定しないが，経済性の観点から，転炉法による製鋼プロセスと，連続鋳造プロセスによる鋼片の鋳造を行うことが望ましい。
[Microstructure]
In order to achieve a high strength of tensile strength ≧ 900 MPa, it is necessary to transform to a low temperature to obtain a structure with a high dislocation density. In addition, when a small amount of ferrite or pearlite is mixed in such a hard microstructure, the base material CTOD value is likely to vary due to the occurrence of pop-in. Therefore, the microstructure of the steel in the present invention is bainite, martensite or bainite. + Limited to martensite two-phase structure.
In the present invention, the steel making method is not particularly limited. However, from the viewpoint of economy, it is desirable to cast a steel slab by a steelmaking process by a converter method and a continuous casting process.
表１に示す化学組成の鋼を用い，表２に示す熱間圧延・加速冷却，再加熱条件で鋼板Ａ
〜Ｋを作製した．なお，再加熱には，加速冷却設備と同一ライン上に設置した誘導加熱型の加熱装置を用いて行った。
Using steel with the chemical composition shown in Table 1, steel plate A was used under the hot rolling / accelerated cooling and reheating conditions shown in Table 2.
-K was produced. Reheating was performed using an induction heating type heating device installed on the same line as the accelerated cooling equipment.
得られた鋼板をせん断機により２０箇所切断し，その後，鋼板切断面を磁粉探傷により調査し，切断割れが認められた切断端面の数を求めた。ここで，１つの端面内に複数の割れが確認できた場合でも，端面としては１つなので，切断割れの発生数は１とした。全ての切断箇所において，切断割れが認められない場合，（切断割れ発生数０）を良好とした。 The obtained steel plate was cut at 20 points with a shearing machine, and then the cut surface of the steel plate was examined by magnetic particle inspection to determine the number of cut end faces where cut cracks were observed. Here, even when a plurality of cracks could be confirmed in one end face, the number of cut cracks was set to 1 because there was only one end face. In the case where no cut cracks were observed in all the cut portions, (the number of cut crack occurrences 0) was considered good.
次に，得られた鋼板の強度と靱性を評価するため，ＡＰＩ−５Ｌに準拠した全厚引張試験片およびＤＷＴＴ試験片およびＢＳ７４４８に準拠したＢ（板厚）×２Ｂサイズの３点曲げＣＴＯＤを採取した。 Next, in order to evaluate the strength and toughness of the obtained steel sheet, full thickness tensile test piece and DWTT test piece according to API-5L and B (plate thickness) × 2B size 3 point bending CTOD according to BS7448 are used. Collected.
板厚中央位置からＪＩＳ Ｚ２２０２（１９８０）のＶノッチシャルピー衝撃試験片を
採取して，鋼板の引張試験，ＤＷＴＴ試験，ＣＴＯＤ試験およびシャルピー衝撃試験を実施した。
A V-notch Charpy impact test piece of JIS Z2202 (1980) was taken from the center of the plate thickness, and a tensile test, a DWTT test, a CTOD test, and a Charpy impact test were performed on the steel sheet.
また，板圧延方向断面に平行にミクロ組織観察用サンプルを採取し，鏡面研磨後，硝酸アルコールエッチング処理を行ってから光学顕微鏡にて組織観察を行い，鋼のミクロ組織の種類を調査した．鋼板のせん断加工試験結果，母材の強度・靱性試験結果，ミクロ組織観察結果をまとめて表３に示す。 In addition, a sample for microstructural observation was taken parallel to the cross section in the plate rolling direction. After mirror polishing, nitric alcohol etching was performed, and the microstructure was observed with an optical microscope to investigate the type of microstructure of the steel. Table 3 summarizes the results of the steel plate shearing test, the base metal strength and toughness test results, and the microstructure observation results.
尚、表３においてＣＴＯＤ−２０（ｍｍ）はー２０℃での限界開口変位量δｃ（ｍｍ）を示す。 In Table 3, CTOD-20 (mm) represents the limit opening displacement δc (mm) at −20 ° C.
本発明範囲は、せん断加工試験では割れ無し、母材の強度は引張強度９００ＭＰａ以上、降伏強度７５０ＭＰａ以上、母材靭性ｖＥ−３０２３０Ｊ以上、ＤＷＴＴ ＳＡ−３０ ７５％以上、ＣＴＯＤ−２００．１５以上を本発明範囲とした。 The scope of the present invention is that there is no cracking in the shearing test, the strength of the base material is tensile strength 900 MPa or more, yield strength 750 MPa or more, base material toughness vE- 30 230J or more, DWTT SA- 30 75% or more, CTOD- 20 0.15. The above is the scope of the present invention.
化学組成および圧延・冷却・再加熱条件が本発明の範囲内である，発明例１〜８は切断
割れが発生することなく，かつ高強度・高靱性を示した。
Inventive Examples 1 to 8, whose chemical composition and rolling / cooling / reheating conditions are within the scope of the present invention, showed high strength and high toughness without occurrence of cutting cracks.
一方，圧延終了温度が本発明の下限を下回った比較例Ｎｏ．９，Ｎｏ．１０はＣＴＯＤ
値が低下した。
On the other hand, Comparative Example No. in which the rolling end temperature was lower than the lower limit of the present invention. 9, no. 10 is CTOD
The value decreased.
冷却停止温度が本発明の上限を上回った比較例Ｎｏ．１１は，一部上部ベイナイト組織
が生成したため強度が低下した。
Comparative Example No. in which the cooling stop temperature exceeded the upper limit of the present invention. No. 11 had a lower strength because part of the upper bainite structure was formed.
再加熱昇温速度が本発明の下限を下回った比較例Ｎｏ．１２は，セメンタイトの粗大化が生じたため、シャルピー吸収エネルギーおよびＤＷＴＴ特性が低下した。 Comparative Example No. in which the reheating temperature rising rate was lower than the lower limit of the present invention. In No. 12, cementite coarsening occurred, and thus Charpy absorbed energy and DWTT characteristics decreased.
再加熱温度が本発明の下限を下回った比較例Ｎｏ．１３は，加熱温度が低すぎて充分な脱水素が起こらなかったため，切断割れが多数発生した。一方、再加熱温度が本発明の上限を上回った比較例Ｎｏ．１４は，加熱温度が高すぎたため，強度が低下した。 Comparative Example No. in which the reheating temperature was below the lower limit of the present invention. In No. 13, since the heating temperature was too low and sufficient dehydrogenation did not occur, many cut cracks occurred. On the other hand, the comparative example No. in which the reheating temperature exceeded the upper limit of the present invention. In No. 14, the strength decreased because the heating temperature was too high.
鋼板のＣ添加量が本発明の上限を上回る鋼種Ｇを用いた比較例Ｎｏ．１５は，高い強度を示したもののセメンタイトの密度が高くなりすぎて切断割れを起こした。また，シャルピー吸収エネルギーおよびＤＷＴＴ特性、ＣＴＯＤ値が低下した。 Comparative Example No. using steel type G in which the C addition amount of the steel sheet exceeds the upper limit of the present invention. No. 15 showed high strength, but the density of cementite became too high, causing cut cracks. In addition, Charpy absorbed energy, DWTT characteristics, and CTOD value decreased.
鋼板のＭｎ添加量が本発明の下限を下回る鋼種Ｈを用いた比較例Ｎｏ．１６は，強度が低下した。 Comparative Example No. using steel type H in which the Mn addition amount of the steel sheet is below the lower limit of the present invention. No. 16 decreased in strength.
鋼板のＮｂ添加量が本発明の上限を上回る鋼種Ｊを用いた比較例Ｎｏ．１７は，Ｎｂ系
炭化物が増加しすぎたために，シャルピー吸収エネルギーおよびＤＷＴＴ特性およびＣＴＯＤ値が低下した。
Comparative Example No. using steel type J in which the Nb addition amount of the steel sheet exceeds the upper limit of the present invention. For No. 17, Charpy absorbed energy, DWTT characteristics, and CTOD value decreased because Nb-based carbides increased too much.
鋼板のＢ添加量が本発明の下限を下回る鋼種Ｋを用いた比較例Ｎｏ．１８は，加速冷却を開始する前の空冷過程で初析フェライトが生成し，ミクロ組織がフェライト+ベイナイトとなったことから，シャルピー吸収エネルギーおよびＣＴＯＤ値が低下した。 Comparative Example No. using steel type K in which the B addition amount of the steel sheet is below the lower limit of the present invention. In No. 18, proeutectoid ferrite was formed in the air cooling process before the start of accelerated cooling, and the microstructure became ferrite + bainite. Therefore, the Charpy absorbed energy and the CTOD value decreased.
Claims (2)
Ｃ：０．０３〜０．１２％
Ｓｉ：０．０１〜０．５％
Ｍｎ：１．５〜３．０％
Ａｌ：０．０１〜０．０８％
Ｎｂ：０．０１〜０．０８％
Ｔｉ：０．００５〜０．０２５％
Ｎ：０．００１〜０．０１％
Ｂ：０．０００５〜０．００３％以下
更に
Ｃｕ：０．０１〜２％
Ｎｉ：０．０１〜３％
Ｃｒ：０．０１〜１％
Ｍｏ：０．０１〜１％
Ｖ：０．０１〜０．１％
の一種または二種以上を含有し、残部Ｆｅ及び不可避的不純物からなる鋼を，
１０００〜１２００℃に加熱後，熱間圧延を開始し，
９５０℃以下の温度域での累積圧下量≧７０％となるよう圧延を行い，
８００℃以上で圧延を終了し，その後
冷却開始温度７００℃以上，８００℃未満より，
冷却速度２０〜８０℃／ｓの加速冷却を開始し，
２５０℃未満の温度域で冷却停止し，
冷却停止後ただちに５℃／ｓ以上の昇温速度で
３００℃以上５００℃以下の温度に再加熱することを特徴とする，耐切断割れ性に優れた
引張強度９００ＭＰａ以上の高強度・高靱性厚鋼板の製造方法。 % By mass
C: 0.03-0.12%
Si: 0.01 to 0.5%
Mn: 1.5 to 3.0%
Al: 0.01 to 0.08%
Nb: 0.01 to 0.08%
Ti: 0.005-0.025%
N: 0.001 to 0.01%
B: 0.0005 to 0.003% or less Cu: 0.01 to 2%
Ni: 0.01 to 3%
Cr: 0.01 to 1%
Mo: 0.01 to 1%
V: 0.01 to 0.1%
A steel containing one or more of the following, the balance Fe and unavoidable impurities,
After heating to 1000-1200 ° C, hot rolling is started,
Rolling is performed so that the cumulative reduction amount ≧ 70% in the temperature range of 950 ° C. or lower,
Rolling is completed at 800 ° C or higher, and then the cooling start temperature is 700 ° C or higher and lower than 800 ° C.
Accelerated cooling is started at a cooling rate of 20 to 80 ° C./s,
Stop cooling in the temperature range below 250 ℃,
Immediately after the cooling is stopped, it is reheated to a temperature of 300 ° C. or more and 500 ° C. or less at a temperature rising rate of 5 ° C./s or more. A method of manufacturing a steel sheet.
Ｃａ：０．０００５〜０．０１％
ＲＥＭ：０．０００５〜０．０２％
Ｚｒ：０．０００５〜０．０３％
Ｍｇ：０．０００５〜０．０１％
の一種あるいは二種以上を含有し
残部Ｆｅ及び不可避的不純物からなる鋼を，
１０００〜１２００℃に加熱後，熱間圧延を開始し，
９５０℃以下の温度域での累積圧下量≧７０％となるよう圧延を行い，
８００℃以上，９００℃以下の温度で圧延を終了し，ひきつづき
７００℃以上，８００℃未満の温度から，
冷却速度２０〜８０℃／ｓの加速冷却を開始し，
２５０℃未満の温度域で冷却停止し，
冷却停止後ただちに５℃／ｓ以上の昇温速度で
３００℃以上５００℃以下の温度に再加熱することを特徴とする，耐切断割れ性に優れた引張強度９００ＭＰａ以上の高強度・高靱性厚鋼板の製造方法。 In addition to the ingredients according to claim 1,
Ca: 0.0005 to 0.01%
REM: 0.0005 to 0.02%
Zr: 0.0005 to 0.03%
Mg: 0.0005 to 0.01%
A steel containing one or more of the following, the balance Fe and unavoidable impurities,
After heating to 1000-1200 ° C, hot rolling is started,
Rolling is performed so that the cumulative reduction amount ≧ 70% in the temperature range of 950 ° C. or lower,
Rolling is finished at a temperature of 800 ° C. or higher and 900 ° C. or lower, and subsequently from a temperature of 700 ° C. or higher and lower than 800 ° C.,
Accelerated cooling is started at a cooling rate of 20 to 80 ° C./s,
Stop cooling in the temperature range below 250 ℃,
Immediately after the cooling is stopped, it is reheated to a temperature of 300 ° C. or more and 500 ° C. or less at a temperature rising rate of 5 ° C./s or more. A method of manufacturing a steel sheet.
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| KR101167388B1 (en) | 2010-02-26 | 2012-07-19 | 현대제철 주식회사 | Steel Marine, and method for producing the same |
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| JP2002115024A (en) * | 2000-10-06 | 2002-04-19 | Nkk Corp | Wear resistant steel having excellent toughness and delayed-fracture resistance and its production method |
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| JP5176271B2 (en) * | 2005-03-22 | 2013-04-03 | 新日鐵住金株式会社 | Method for producing high-strength steel sheet for line pipe with tensile strength of 760 MPa or higher with suppressed increase in yield strength after heating by coating treatment, and method for producing high-strength steel pipe for line pipe using the same |
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