JP2870830B2 - Method for producing high tensile strength and high toughness steel sheet excellent in HIC resistance - Google Patents
Method for producing high tensile strength and high toughness steel sheet excellent in HIC resistanceInfo
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- JP2870830B2 JP2870830B2 JP19878789A JP19878789A JP2870830B2 JP 2870830 B2 JP2870830 B2 JP 2870830B2 JP 19878789 A JP19878789 A JP 19878789A JP 19878789 A JP19878789 A JP 19878789A JP 2870830 B2 JP2870830 B2 JP 2870830B2
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Description
【発明の詳細な説明】 ［産業上の利用分野］ この発明は、耐HIC特性に優れた高張力高靭性鋼板の
製造方法に関するものである。Description: TECHNICAL FIELD The present invention relates to a method for producing a high tensile strength and high toughness steel sheet having excellent HIC resistance.
［従来の技術］ 石油あるいはガス輸送ラインパイプに用いられる鋼板
に対して要求される性能は、その敷設環境により異なっ
ている。[Prior Art] The performance required for a steel plate used in an oil or gas transportation line pipe differs depending on the installation environment.
輸送ガス中の硫化水素濃度の高い環境では、耐水素割
れ性および耐硫化物応力腐食割れ性の高い鋼材が要求さ
れている。耐HIC（耐水素誘起割れ）特性の向上のため
には、HICの発生点となるMnSの形態制御にCa,REM等が添
加されている。また、HICの伝播を防ぐため、低温変態
生成物の制御および偏析の低減、添加元素（C,Mn,P）の
制限、加速冷却（急冷）により均一な組織を得ること等
の方法が従来から用いられている（特開昭57−85928号
公報）。このとき、圧延および加速冷却条件としては、
α＋γ二相領域での圧延後の加速冷却においては、組織
の不均一あるいは硬化相を現出させるため、圧延終了温
度をAr3点以上の温度が好ましいとされている（特開昭5
4−118325）。In an environment where the concentration of hydrogen sulfide in the transport gas is high, steel materials having high resistance to hydrogen cracking and sulfide stress corrosion cracking are required. In order to improve the resistance to HIC (hydrogen-induced cracking), Ca, REM, and the like are added to control the morphology of MnS, which is the point at which HIC occurs. In order to prevent the propagation of HIC, methods to control low-temperature transformation products and reduce segregation, limit the added elements (C, Mn, P), and obtain a uniform structure by accelerated cooling (quenching) have been used. (JP-A-57-85928). At this time, rolling and accelerated cooling conditions include:
In the accelerated cooling after rolling in the α + γ two-phase region, it is considered that the rolling end temperature is preferably a temperature of 3 points or more of Ar in order to make the structure uneven or a hardened phase appear (Japanese Patent Laid-Open No.
4-118325).
また、低温環境で使用される鋼材には脆性破壊防止の
ため、良好なDWTT特性が要求されている。良好なDWTT特
性を得る方法としては、α＋γ二相領域での圧延を行な
うことにより、加工フェライトの伸長ならびに衝撃破面
のセパレーションを発生させ遷移温度の低下を図る場合
が多い（特開昭55−115924）。In addition, steel materials used in low-temperature environments are required to have good DWTT characteristics in order to prevent brittle fracture. As a method for obtaining good DWTT characteristics, rolling in the α + γ two-phase region is often performed to reduce the transition temperature by elongation of the processed ferrite and separation of the impact fracture surface (Japanese Patent Laid-Open No. 55-55). 115924).
一方、ラインパイプ材の強度は、輸送効率の上昇、使
用鋼材量低減等のため、年々高強度化している。このよ
うな高強度材の製造は、制御圧延および加速冷却を利用
したTMCPにより製造されている。加速冷却方法として
は、Ar3点以上の温度で圧延を終了させた後に、550〜65
0℃まで５〜20℃/secの冷却速度で冷却する方法（特開
昭54−21917号公報）および500℃以下の温度まで２〜25
℃/secの冷却速度で冷却する方法（特開昭58−120727号
公報）等が開示されている。On the other hand, the strength of line pipe materials is increasing year by year in order to increase transportation efficiency and reduce the amount of steel materials used. Such high-strength materials are manufactured by TMCP using controlled rolling and accelerated cooling. As an accelerated cooling method, after ending the rolling at a temperature of Ar 3 points or more, 550 to 65
Cooling to 0 ° C. at a cooling rate of 5 to 20 ° C./sec (JP-A-54-21917);
A method of cooling at a cooling rate of ° C./sec (JP-A-58-120727) and the like are disclosed.
［発明が解決しようとする課題］ このように、ラインパイプ用の鋼板に要求される性能
が単一の場合には、上述した方法によってその要求性能
を満足させることは可能である。しかしながら、近年、
低温で且つ硫化水素濃度の高い環境で使用される、高強
度で且つ低い降伏比を有する鋼の要求が高まるにつれ
て、従来の解決方法では、この要求を満足できなくなり
つつある。即ち、ラインパイプ用の鋼板を製造する場合
において、γ領域での圧延は耐HIC特性を改善するがDWT
T特性の劣化をまねき、一方、逆にα＋γ領域での圧延
はDWTT特性を向上させるが耐HIC特性の劣化および降伏
比の上昇をまねく。従って、上記２つの特性を同時に満
足させることは従来方法では非常に困難であった。[Problems to be Solved by the Invention] As described above, when the performance required for a steel sheet for a line pipe is a single performance, the required performance can be satisfied by the above-described method. However, in recent years,
As the demand for high strength and low yield ratio steels used in low temperature and high hydrogen sulfide environments increases, conventional solutions are no longer able to satisfy this requirement. In other words, when producing steel sheets for line pipes, rolling in the γ region improves HIC resistance, but DWT
On the other hand, rolling in the α + γ region improves the DWTT characteristics, but leads to deterioration of the HIC resistance characteristics and an increase in the yield ratio. Therefore, it has been extremely difficult with the conventional method to simultaneously satisfy the above two characteristics.
さらに、加速冷却方法においても、上述の要求から、
炭素添加量を低減させ、Mn等の添加によって強度の上昇
を図った鋼においては、圧延終了後に起こる変態終了温
度が加速冷却停止温度以下となる場合がある。このよう
な場合、加速冷却停止後の放冷時にフェライト−パーラ
イト変態が起こり、目的とする均一な組織が得られない
場合がある。また、この組織の降伏比は高いものとな
る。さらに、加速冷却停止温度の低下は、鋼種によって
は、高強度が得られるものの著しい靭性の劣化をまねく
場合がある。Furthermore, also in the accelerated cooling method, from the above-mentioned requirements,
In steels in which the amount of carbon added is reduced and the strength is increased by the addition of Mn or the like, the transformation end temperature that occurs after the end of rolling may be equal to or lower than the accelerated cooling stop temperature. In such a case, ferrite-pearlite transformation occurs at the time of cooling after stopping the accelerated cooling, and a desired uniform structure may not be obtained. Also, the yield ratio of this structure is high. Further, a decrease in the accelerated cooling stop temperature may lead to a significant deterioration in toughness although high strength is obtained depending on the type of steel.
従って、この発明の目的は、上述した多様な要求性能
を満足させるためになされたものであって、耐HIC特性
に優れ、高強度、高靭性且つ低降伏比を有する鋼板の製
造方法を提供することにある。Therefore, an object of the present invention is to satisfy the above-described various required performances, and provides a method for producing a steel sheet having excellent HIC resistance, high strength, high toughness and a low yield ratio. It is in.
［課題を解決するための手段］ 発明者等は、上述した問題を解決すべく鋭意努力し
た。その結果、α＋γ二相領域での圧延により生じるセ
パレーションを発生させず、且つ、合金元素（主に炭
素）の添加を低減し、高強度、高靭性を得るためには、
微細且つ均一なベイナイト組織（アシキュラ−フェライ
ト）が有効であること、および、低炭素鋼においては前
記組織とした場合に耐HIC特性が良好となることを知見
した。[Means for Solving the Problems] The inventors have made intensive efforts to solve the above-mentioned problems. As a result, in order to prevent the separation caused by rolling in the α + γ two-phase region from occurring, reduce the addition of alloying elements (mainly carbon), and obtain high strength and high toughness,
It has been found that a fine and uniform bainite structure (acicular-ferrite) is effective, and that the low carbon steel has good HIC resistance when the structure is used.
この発明は上述の知見に基いてなされたものであり、 C :0.02 〜0.06 wt.％、 Si:0.03 〜0.50 wt.％、 Mn:0.5 〜2.5 wt.％、 Nb:0.005 〜0.100wt.％、 Ti:0.005 〜0.100wt.％、 Al:0.005 〜0.100wt.％、 Ca:0.0005〜0.008wt.％、 S :0.004wt.％以下（０を含む）、 下記からなる群のうち少なくとも１つの元素、 Cu:0.5 wt.％以下（０は含まず）、 Ni:0.5 wt.％以下（０は含まず）、 Cr:0.5 wt.％以下（０は含まず）、 Mo:0.5 wt.％以下（０は含まず）、 V :0.15wt.％以下（０は含まず）、および、 残部:Feおよび不可避不純物、 からなるスラブを1000〜1250℃の温度に加熱し、次い
で、加熱された前記スラブに対し、圧延温度：オーステ
ナイト再結晶温度域、の条件で第１段目の圧延を施し、
次いで、圧延温度：オーステナイト未再結晶温度域、累
積圧下率:70％以上、圧延仕上がり温度:Ar3点以上、の
条件で第２段目の圧延を施し、次いで、鋼種により定ま
るパーライト生成の臨界冷却速度以上の冷却速度で、Ar
3点以上の温度からAr3点−250℃以下の温度まで冷却す
ることに特徴を有するものである。The present invention has been made on the basis of the above-mentioned findings. C: 0.02 to 0.06 wt.%, Si: 0.03 to 0.50 wt.%, Mn: 0.5 to 2.5 wt.%, Nb: 0.005 to 0.100 wt.% , Ti: 0.005 to 0.100 wt.%, Al: 0.005 to 0.100 wt.%, Ca: 0.0005 to 0.008 wt.%, S: 0.004 wt.% Or less (including 0), at least one of the following groups: Element, Cu: 0.5 wt.% Or less (excluding 0), Ni: 0.5 wt.% Or less (excluding 0), Cr: 0.5 wt.% Or less (excluding 0), Mo: 0.5 wt.% The slab consisting of the following (not including 0), V: 0.15 wt.% Or less (not including 0), and the balance: Fe and inevitable impurities, was heated to a temperature of 1000 to 1250 ° C., and then heated. The first stage rolling is performed on the slab under the conditions of rolling temperature: austenite recrystallization temperature range,
Then, the second stage of rolling is performed under the following conditions: rolling temperature: austenite non-recrystallization temperature range, cumulative rolling reduction: 70% or more, rolling finish temperature: 3 points or more, and then the criticality of pearlite formation determined by steel type At a cooling rate higher than the cooling rate, Ar
It is characterized in that it is cooled from a temperature of 3 points or more to a temperature of 3 points of Ar -250 ° C or less.
次に、この発明の鋼板の製造条件および鋼の化学成分
組成を上述のように限定した理由を以下に述べる。Next, the reasons for limiting the production conditions of the steel sheet of the present invention and the chemical composition of the steel as described above will be described below.
スラブ加熱温度 我々は、スラブ加熱温度がDWTT特性に及ぼす影響を下
記の試験によって調べた。即ち、第１表に示す、本発明
内の化学成分組成を有する供試鋼（スラブ）を1000〜13
00℃の種々の温度に加熱し、次いで、オーステナイト再
結晶温度域で第１段目の圧延を施し、次いで、圧延温
度：オーステナイト未再結晶温度域、累積圧下率:75
％、圧延終了温度:800℃、の条件で第２段目の圧延を施
し、次いで800℃から500℃まで、15℃/secの冷却速度で
加速冷却（急冷）し、加熱温度とDWTT試験結果との関係
を調べ、第１図にその結果を示した。第１図に示すよう
に、本発明供試鋼においても、加熱温度が1250℃を超え
ると良好なDWTT特性が得られなくなる。一方、設備上お
よび圧延能率の制約から加熱温度は1000℃以上が好まし
い。Slab heating temperature We examined the effect of slab heating temperature on DWTT characteristics by the following tests. That is, the test steel (slab) having the chemical composition within the present invention shown in Table 1 was used for 1000 to 13
Heating to various temperatures of 00 ° C., then performing the first rolling in the austenite recrystallization temperature range, and then rolling temperature: austenite non-recrystallization temperature range, cumulative rolling reduction: 75
%, Rolling end temperature: 800 ° C, the second stage of rolling, then accelerated cooling (rapid cooling) from 800 ° C to 500 ° C at a cooling rate of 15 ° C / sec, heating temperature and DWTT test results Was examined, and the results are shown in FIG. As shown in FIG. 1, even in the test steel of the present invention, if the heating temperature exceeds 1250 ° C., good DWTT characteristics cannot be obtained. On the other hand, the heating temperature is preferably 1000 ° C. or more from the viewpoint of equipment and the restriction of the rolling efficiency.
オーステナイト未再結晶温度域での累積圧下率 我々は、オーステナイト未再結晶温度域での累積圧下
率がDWTT特性に及ぼす影響を下記の試験によって調べ
た。即ち、第１表に示す、本発明内の化学成分組成を有
する供試鋼（スラブ）を、1200℃の温度に加熱し、次い
で、オーステナイト再結晶温度域で第１段目の圧延を施
し、次いで、オーステナイト未再結晶温度域で、圧延終
了温度を800℃とし、累積圧下率を50〜80％の範囲に変
化させて第２段目の圧延を施し、次いで、800℃から500
℃まで15℃/secの冷却速度で加速冷却し、累積圧下率と
DWTT特性との関係を調べ、第２図にその結果を示した。
第２図に示すように、累積圧下率が70％以上となるとDW
TT特性は向上し、α＋γ二相域圧延を行なわずに良好な
DWTT特性が得られる。 Cumulative rolling reduction in austenite non-recrystallization temperature range We examined the effect of the cumulative rolling reduction in austenite non-recrystallization temperature range on DWTT characteristics by the following test. That is, the test steel (slab) having the chemical composition within the present invention shown in Table 1 was heated to a temperature of 1200 ° C., and then subjected to the first stage rolling in the austenite recrystallization temperature range, Next, in the austenite non-recrystallization temperature range, the rolling end temperature is set to 800 ° C., and the cumulative rolling reduction is changed to a range of 50 to 80%, and the second rolling is performed.
Accelerated cooling to 15 ° C at a cooling rate of 15 ° C / sec.
The relationship with the DWTT characteristics was examined, and the results are shown in FIG.
As shown in Fig. 2, when the cumulative rolling reduction becomes 70% or more, DW
TT characteristics are improved and good without α + γ dual phase rolling
DWTT characteristics are obtained.
圧延終了温度 我々は、圧延終了温度（加速冷却開始温度）がHIC特
性［CLR:Cracking length ratio（割れ長さ率）］に及
ぼす影響を下記の試験によって調べた。即ち、第１表に
示す、本発明内の化学成分組成を有する供試鋼（スラ
ブ）を1200℃に加熱し、次いで、オーステナイト再結晶
温度域で第１段目の圧延を施し、次いで、圧延温度：オ
ーステナイト未再結晶温度域、累積圧下率:70％、圧延
終了温度を680〜800℃の範囲に変化させた条件で第２段
目の圧延を施し、次いで、680〜800℃から500℃まで15
℃/secの冷却速度で加速冷却した。そして、得られた供
試鋼板を硫化水素を飽和させた5wt.％食塩−0.5wt.％酢
酸からなる水溶液に96時間浸漬して耐HIC特性を調べ、
その結果を第３図に示した。なお、供試鋼のAr3点
（℃）は、下記（１）式によって求められる。Rolling end temperature We examined the effect of the rolling end temperature (accelerated cooling start temperature) on HIC characteristics [CLR: cracking length ratio (crack length ratio)] by the following test. That is, the test steel (slab) having the chemical composition according to the present invention shown in Table 1 was heated to 1200 ° C., and then subjected to the first-stage rolling in the austenite recrystallization temperature range, and then to rolling. Temperature: austenite non-recrystallization temperature range, cumulative rolling reduction: 70%, rolling at the second stage under the conditions of changing the rolling end temperature to the range of 680 to 800 ° C, and then from 680 to 800 ° C to 500 ° C Up to 15
Accelerated cooling was performed at a cooling rate of ° C / sec. Then, the obtained test steel sheet was immersed in an aqueous solution consisting of 5 wt.% Salt-0.5 wt.% Acetic acid saturated with hydrogen sulfide for 96 hours to examine the HIC resistance.
The result is shown in FIG. The Ar 3 point (° C.) of the test steel is obtained by the following equation (1).
Ar3＝910−310C−80Mn−20Cu−15Cr −55Ni−80Mo …（１）式 第３図に示すように、圧延終了温度がAr3点（（１）
式より760℃）未満では耐HIC特性の劣化が認められる。
従って、圧延終了温度はAr3点（Ar3温度）以上とすべき
である。 Ar 3 = 910-310C-80Mn-20Cu -15Cr -55Ni-80Mo ... (1) As shown in FIG. 3 type, rolling end temperature is Ar 3 point ((1)
If the temperature is lower than 760 ° C. according to the formula), deterioration of the HIC resistance is observed.
Therefore, the rolling end temperature should be equal to or higher than the Ar 3 point (Ar 3 temperature).
冷却速度 我々は、冷却速度が引張特性および靭性に及ぼす影響
を下記の試験によって調べた。即ち、第１表に示す、本
発明内の化学成分組成を有する供試鋼（スラブ）を1200
℃の温度に加熱し、次いで、オーステナイト再結晶温度
域で第１段目の圧延を施し、次いで、圧延温度：オース
テナイト未再結晶温度域、累積圧下率:75％、圧延終了
温度:800℃の条件で第２段目の圧延を施し、次いで、80
0℃から500℃まで、冷却速度を５〜35℃/sec範囲に変化
させて加速冷却した。そして、得られた供試鋼板の引張
特性および靭性を調べ、その結果を第４図に示した。な
お、パーライト生成の臨界冷却速度は、下記（２）式に
よって求められる。Cooling Rate We examined the effect of cooling rate on tensile properties and toughness by the following tests. That is, a test steel (slab) having a chemical composition within the present invention shown in Table 1 was used for 1200 times.
℃, then the first stage of rolling in the austenite recrystallization temperature range, then rolling temperature: austenite non-recrystallization temperature range, cumulative reduction: 75%, rolling end temperature: 800 ℃ The second stage rolling is performed under the conditions, and then 80
From 0 ° C to 500 ° C, accelerated cooling was performed by changing the cooling rate in the range of 5 to 35 ° C / sec. Then, the tensile properties and toughness of the obtained test steel sheet were examined, and the results are shown in FIG. The critical cooling rate for pearlite formation is determined by the following equation (2).
パーライト生成の臨界冷却速度＝300/kp …（２）式 Log Kp＝0.5966C−0.0992Si＋1.395Mn＋0.385Ni ＋1.295Cr＋0.3978Cu＋3.73Mo−0.8688。 Critical cooling rate for pearlite formation = 300 / kp Equation (2) Log Kp = 0.5966C−0.0992Si + 1.395Mn + 0.385Ni + 1.295Cr + 0.3978Cu + 3.73Mo−0.8688.
ただし、 kp:パーライトの生成に要する800℃から500℃までの
冷却時間。Here, kp: cooling time from 800 ° C to 500 ° C required for pearlite generation.
第４図に示すように、冷却速度がパーライト生成の臨
界冷却速度以上となると、引張強度が上昇し、降伏強度
は低下する。このとき、靭性の劣化は認められない。ま
た、得られた供試鋼板の組織観察の結果では、臨界冷却
速度未満の５℃/secで冷却した鋼板にはパーライトの生
成が認められるが、臨界速度以上となる15℃/secで冷却
した鋼板ではパーライトの生成は認められず均一な組織
となっている。このように、パーライト生成の臨界冷却
速度以上で冷却した場合、均一なベイナイト組織により
引張強度の上昇がもたらされ、且つ、靭性は劣化しな
い。従って、冷却速度は、鋼種によって定まるパーライ
ト生成の臨界冷却速度以上とすべきである。As shown in FIG. 4, when the cooling rate is higher than the critical cooling rate for pearlite formation, the tensile strength increases and the yield strength decreases. At this time, no deterioration in toughness is observed. In addition, the results of microstructure observation of the obtained test steel sheet show that pearlite is generated in the steel sheet cooled at 5 ° C./sec lower than the critical cooling rate, but the steel sheet was cooled at 15 ° C./sec at or above the critical speed. The formation of pearlite is not recognized in the steel sheet, and the steel sheet has a uniform structure. As described above, when cooling is performed at a critical cooling rate or higher for pearlite formation, a uniform bainite structure results in an increase in tensile strength and no deterioration in toughness. Therefore, the cooling rate should be equal to or higher than the critical cooling rate for pearlite formation determined by the type of steel.
冷却停止温度 我々は、冷却停止温度が引張特性および靭性に及ぼす
影響を下記の試験によって調べた。即ち、第２表に示
す、本発明内の化学成分組成を有する供試鋼A,B（スラ
ブ）を1200℃の温度に加熱し、次いで、圧延温度：オー
ステナイト再結晶温度域で第１段目の圧延を施し、次い
で、圧延温度：オーステナイト未再結晶温度域、累積圧
下率:75％、圧延終了温度:800℃の条件で第２段目の圧
延を施し、次いで、800℃から15℃/secで室温から650℃
まで冷却停止温度を変化させて加速冷却した。そして、
得られた供試鋼板の引張特性および靭性を調べ、その結
果を第５図に示した。Cool stop temperature We examined the effect of the cool stop temperature on tensile properties and toughness by the following tests. That is, the test steels A and B (slabs) having the chemical composition according to the present invention shown in Table 2 were heated to a temperature of 1200 ° C., and then the first stage was performed at a rolling temperature: austenite recrystallization temperature range. Rolling temperature: rolling temperature: austenite non-recrystallization temperature range, cumulative rolling reduction: 75%, rolling end temperature: 800 ° C., the second stage of rolling, then from 800 ℃ to 15 ℃ / From room temperature to 650 ℃ in sec
Accelerated cooling was performed by changing the cooling stop temperature until the cooling. And
The tensile properties and toughness of the obtained test steel sheet were examined, and the results are shown in FIG.
第５図に示すように、冷却停止温度がAr3点−250℃以
下では、引張強度の上昇ならびに降伏比の低下が認めら
れる。この場合、靭性の劣化も認められない。これに対
して、Ar3点−250℃を超える温度で加速冷却を停止した
場合、冷却速度が臨界冷却速度以上であっても、加速冷
却後の放冷中にフェライト−パーライト変態が進行し均
一なベイナイト組織が得られない。一方、冷却停止温度
をAr3点−250℃以下とすると変態は完了し、その後の放
冷時に組織変化は起こらず均一なベイナイト組織が得ら
れる。従って、冷却停止温度はAr3点−250℃以下とすべ
きである。As shown in FIG. 5, when the cooling stop temperature is lower than the Ar 3 point −250 ° C., an increase in tensile strength and a decrease in yield ratio are observed. In this case, no deterioration in toughness is observed. In contrast, when accelerated cooling is stopped at a temperature exceeding the Ar 3 point -250 ° C, even when the cooling rate is equal to or higher than the critical cooling rate, the ferrite-pearlite transformation progresses during cooling after accelerated cooling, resulting in uniformity. No bainite structure can be obtained. On the other hand, when the cooling stop temperature is set to an Ar 3 point of −250 ° C. or less, the transformation is completed, and a uniform bainite structure is obtained without any change in structure during the subsequent cooling. Therefore, the cooling stop temperature should be lower than or equal to the Ar 3 point −250 ° C.
なお、Ar3点−250℃以下の温度で冷却を停止した後
は、上述したように、室温まで放冷する。After the cooling is stopped at a temperature lower than or equal to the Ar 3 point of −250 ° C., it is allowed to cool to room temperature as described above.
鋼の化学成分組織 （１）Ｃ（炭素） 化学成分組成に関する本発明のポイントは、Ｃ含有量
にある。我々は、Ｃ含有量が引張強度および靭性に及ぼ
す影響を下記の試験によって調べた。即ち、Ｃ含有量を
0.03〜0.15まで種々変化させ残りの化学成分組成は本発
明範囲内の供試鋼（スラブ）を1200℃に加熱し、次いで
オーステナイト再結晶温度域で第１段目の圧延を施し、
次いで、圧延温度：オーステナイト未再結晶温度域、累
積圧下率:75％、圧延仕上り温度:800℃（Ar3点以上）、
の条件で第２段目の圧延を施し、次いで、パーライト生
成の臨界冷却速度以上の冷却速度で、800℃から500℃ま
たは350℃（両温度とも供試鋼のAr3点−250℃以下）の
温度まで加速冷却した。そして、得られた供試鋼板の引
張特性および靭性を調べ、その結果を第６図に示した。 Chemical composition of steel (1) C (carbon) The point of the present invention regarding the chemical composition is the C content. We investigated the effect of C content on tensile strength and toughness by the following tests. That is, the C content
The test composition (slab) within the range of the present invention is heated to 1200 ° C., and the first stage rolling is performed in the austenite recrystallization temperature range, with the remaining chemical composition varied from 0.03 to 0.15.
Next, rolling temperature: austenite non-recrystallization temperature range, cumulative rolling reduction: 75%, rolling finish temperature: 800 ° C (Ar 3 points or more),
The second stage of rolling is performed under the conditions described above, and then at a cooling rate higher than the critical cooling rate for pearlite formation, from 800 ° C to 500 ° C or 350 ° C (Ar 3 points of the test steel -250 ° C or less at both temperatures). And accelerated cooling to the temperature. Then, the tensile properties and toughness of the obtained test steel sheet were examined, and the results are shown in FIG.
第６図に示すように、Ｃ含有量が0.06wt.％を超える
と、強度が上昇し靭性の劣化が生じる。さらに、Ｃ含有
量の上昇により加速冷却停止温度の変化による強度変化
も大きくなる。Ｃ含有量が0.06wt.％以下であれば靭性
の劣化が小さく且つ加速冷却（水冷）停止温度による強
度変化が少ない。一方、Ｃ含有量が0.02wt.％以上であ
れば、Nb,V,Ti等の析出硬化を有効に利用することがで
きる。従って、Ｃ含有量は0.02〜0.06wt.％の範囲に限
定すべきである。As shown in FIG. 6, when the C content exceeds 0.06 wt.%, The strength increases and the toughness deteriorates. Further, the increase in the C content causes a large change in strength due to a change in the accelerated cooling stop temperature. When the C content is 0.06 wt.% Or less, the deterioration of toughness is small and the strength change due to the accelerated cooling (water cooling) stop temperature is small. On the other hand, when the C content is 0.02 wt.% Or more, precipitation hardening of Nb, V, Ti, and the like can be effectively used. Therefore, the C content should be limited to the range of 0.02 to 0.06 wt.%.
（２）Si（シリコン） Siは脱酸のため添加される。Si含有量が0.03wt.％未
満では上述した作用に所望の効果が得られない。一方、
Si含有量が0.5wt.％を超えると靭性が劣化する。従っ
て、Si含有量は0.03〜0.5wt.％の範囲に限定すべきであ
る。(2) Si (Si) Si is added for deoxidation. If the Si content is less than 0.03 wt.%, Desired effects cannot be obtained in the above-described operation. on the other hand,
If the Si content exceeds 0.5 wt.%, The toughness deteriorates. Therefore, the Si content should be limited to the range of 0.03 to 0.5 wt.%.
（３）Mn（マンガン） Mnは脱酸のため添加される。Mn含有量が0.5wt.％未満
では上述した作用に所望の効果が得られない。一方、Mn
含有量が2.5wt.％を超えると溶接性が劣化する。従っ
て、Mn含有量は、0.5〜2.5wt.％の範囲に限定すべきで
ある。(3) Mn (manganese) Mn is added for deoxidation. If the Mn content is less than 0.5 wt.%, Desired effects cannot be obtained for the above-mentioned effects. On the other hand, Mn
If the content exceeds 2.5 wt.%, The weldability deteriorates. Therefore, the Mn content should be limited to the range of 0.5-2.5 wt.%.
（４）Nb（ニオブ） Nbは強度および靭性を向上させる作用がある。Nb含有
量が0.005wt.％未満では上述した作用に所望の効果が得
られない。一方、Nb含有量が0.1wt.％を超えると母材お
よび溶接部の靭性を劣化する。従って、Nb含有量は0.00
5〜0.1wt.％の範囲に限定すべきである。(4) Nb (niobium) Nb has an effect of improving strength and toughness. If the Nb content is less than 0.005 wt.%, Desired effects cannot be obtained for the above-mentioned effects. On the other hand, if the Nb content exceeds 0.1 wt.%, The toughness of the base metal and the welded portion deteriorates. Therefore, the Nb content is 0.00
It should be limited to the range of 5 to 0.1 wt.%.
（５）Ti（チタン） Tiはスラブ加熱時のオーステナイトの粗大化を防止す
る作用を有する。Ti含有量が0.005wt.％未満では上述の
作用に所望の効果が得られない。一方、Ti含有量が0.1w
t.％を超えると溶接部の靭性を劣化する。従って、Ti含
有量は0.005〜0.1wt.％の範囲に限定すべきである。(5) Ti (Titanium) Ti has an effect of preventing austenite from becoming coarse during slab heating. If the Ti content is less than 0.005 wt.%, Desired effects cannot be obtained in the above-described operation. On the other hand, the Ti content is 0.1w
If the content exceeds t.%, the toughness of the welded portion is deteriorated. Therefore, the Ti content should be limited to the range of 0.005 to 0.1 wt.%.
（６）Al（アルミニウム） Alは脱酸のため添加される。Al含有量が0.005wt.％未
満では上述した作用に所望の効果が得られない。一方、
Al含有量が0.1wt.％を超えると非金属介在物の増加をも
たらしシャルピー吸収エネルギーを低下させる。従っ
て、Al含有量は0.005〜0.1wt.％の範囲に限定すべきで
ある。(6) Al (aluminum) Al is added for deoxidation. If the Al content is less than 0.005 wt.%, Desired effects cannot be obtained in the above-described operation. on the other hand,
If the Al content exceeds 0.1% by weight, non-metallic inclusions increase, and the Charpy absorbed energy decreases. Therefore, the Al content should be limited to the range of 0.005 to 0.1 wt.%.
（７）Ca（カルシウム） Caは介在物の形態を変化させHICの基点を減少させる
作用を有する。Ca含有量が0.0005wt.％未満では上述し
た作用に所望の効果が得られない。一方、Ca含有量が0.
008wt.％を超えると逆に介在物の量の増加をもたらす。
従って、Ca含有量は0.0005〜0.008wt.％の範囲に限定す
べきである。(7) Ca (Calcium) Ca has the effect of changing the form of inclusions and decreasing the base point of HIC. If the Ca content is less than 0.0005 wt.%, The desired effect cannot be obtained in the above-described operation. On the other hand, the Ca content is 0.
If it exceeds 008 wt.%, The amount of inclusions will increase.
Therefore, the Ca content should be limited to the range of 0.0005 to 0.008 wt.%.
（８）Ｓ（硫黄） Ｓは不純物として不可避の元素であるが、0.004wt.％
以下の含有量であれば実害はない。ただし、Ｓ含有量が
0.004wt.％を超えるとHICの基点となる介在物が増加
し、耐HIC特性を劣化させる。従って、Ｓ含有量は0.004
wt.％以下（０を含む）とすべきである。(8) S (sulfur) S is an unavoidable element as an impurity, but 0.004 wt.%
There is no actual harm if the content is below. However, the S content
If the content exceeds 0.004 wt.%, The inclusions serving as the starting point of HIC increase, deteriorating the HIC resistance. Therefore, the S content is 0.004
It should be less than wt.% (including 0).
（９）Cu（銅） Cuは耐食性を向上させる作用を有するが、Cu含有量が
0.5wt.％を超えると溶接部の靭性が劣化する。従って、
Cu含有量は0.5wt.％以下（０は含まず）に限定すべきで
ある。(9) Cu (copper) Cu has the effect of improving corrosion resistance, but the Cu content is
If it exceeds 0.5 wt.%, The toughness of the weld will deteriorate. Therefore,
Cu content should be limited to 0.5 wt.% Or less (excluding 0).
（10）Ni（ニッケル） Niは強度および靭性を向上する作用を有するが、Ni含
有量が0.5wt.％を超えると経済的に不利である。従っ
て、Ni含有量は0.5wt.％以下（０は含まず）に限定すべ
きである。(10) Ni (nickel) Ni has the effect of improving strength and toughness, but is economically disadvantageous if the Ni content exceeds 0.5 wt.%. Therefore, the Ni content should be limited to 0.5 wt.% Or less (excluding 0).
（11）Cr（クロム） Crは耐食性を向上させる作用を有するが、Cr含有量が
0.5wt.％を超えると、溶接部の靭性が劣化する。従っ
て、Cr含有量は0.5wt.％以下（０は含まず）に限定すべ
きである。(11) Cr (chromium) Cr has the effect of improving corrosion resistance, but the Cr content is
If it exceeds 0.5 wt.%, The toughness of the weld will deteriorate. Therefore, the Cr content should be limited to 0.5 wt.% Or less (excluding 0).
（12）Mo（モリブデン） Moは強度および靭性を向上させる作用を有するが、Mo
含有量が0.5wt.％を超えると経済的に不利である。従っ
て、Mo含有量は0.5wt.％以下（０は含まず）に限定すべ
きである。(12) Mo (Molybdenum) Mo has an effect of improving strength and toughness.
If the content exceeds 0.5 wt.%, It is economically disadvantageous. Therefore, the Mo content should be limited to 0.5 wt.% Or less (excluding 0).
（13）Ｖ（バナジウム） Ｖは強度および靭性を向上させる作用を有するが、Ｖ
含有量が0.15wt.％を超えると母材および溶接部の靭性
が劣化する。従って、Ｖ含有量は0.15wt.％以下（０は
含まず）に限定すべきである。(13) V (Vanadium) V has an effect of improving strength and toughness.
If the content exceeds 0.15 wt.%, The toughness of the base metal and the welded portion deteriorates. Therefore, the V content should be limited to 0.15 wt.% Or less (excluding 0).
次に、この発明を実施例により更に詳しく説明する。 Next, the present invention will be described in more detail with reference to examples.
［実施例］ 第３表に示す化学成分組成を有する、本発明鋼（スラ
ブ）A,B,C、および比較鋼（スラブ）Ｄの各々から、本
発明の製造方法または本発明外の製造方法によって、仕
上げ板厚25mmの、本発明鋼板No.1,7,8,10,および、比較
鋼板No.2〜6,9,11,12を調製した。圧延条件、冷却（水
冷）条件は第４表に示した。また、本発明鋼Ａ〜Ｃ、比
較鋼Ｄの未再結晶温度、前述の（１）式から求めたAr3
温度、および、前述の（２）式から求めた臨界冷却速度
を第３表に併せて示した。そして、本発明鋼板および比
較鋼板の各々の降伏強度、引張強さ、降伏比を調べその
結果を第５表に示した。さらに、本発明鋼板および比較
鋼板の各々に対して、シャルピー試験、DWTT試験、HIC
試験を行ないその結果を第５表に併せて示した。[Example] The production method of the present invention or the production method other than the present invention from each of the steels (slabs) A, B, and C of the present invention and the comparative steel (slab) D having the chemical composition shown in Table 3. Thus, steel plates Nos. 1, 7, 8, and 10 of the present invention and comparative steel plates Nos. 2 to 6, 9, 11, and 12 having a finished plate thickness of 25 mm were prepared. Rolling conditions and cooling (water cooling) conditions are shown in Table 4. In addition, the non-recrystallization temperatures of the steels A to C of the present invention and the comparative steel D, and Ar 3 obtained from the above equation (1).
Table 3 also shows the temperature and the critical cooling rate determined from the above-mentioned equation (2). Then, the yield strength, tensile strength and yield ratio of the steel sheet of the present invention and the comparative steel sheet were examined, and the results are shown in Table 5. Furthermore, the Charpy test, the DWTT test, the HIC
The test was performed and the results are shown in Table 5.
第３表〜第５表から明らかなように、本発明鋼板No.1
は、卓越した強度靭性バランスおよび良好な耐HIC特性
を有している。As is clear from Tables 3 to 5, the steel sheet No. 1 of the present invention
Has excellent balance of strength and toughness and good HIC resistance.
比較鋼板No.2〜６は、本発明鋼板Ａを用いているが、
No.2は加熱温度、No.3は仕上げ温度、No.4は未再結晶温
度域での圧下率、No.5は水冷停止温度、No.6は冷却速度
が本発明の製造条件から外れているため、本発明鋼板と
比較し、劣化した材質および耐HIC特性を示した。比較
鋼板No.2,4においてはDWTT特性、比較鋼板No.3において
はHIC特性、比較鋼板No.5,6においては降伏比、DWTT特
性が、本発明鋼板と比較して劣化していた。Comparative steel sheets No. 2 to 6 use the steel sheet A of the present invention,
No. 2 is the heating temperature, No. 3 is the finishing temperature, No. 4 is the rolling reduction in the non-recrystallization temperature range, No. 5 is the water cooling stop temperature, and No. 6 the cooling rate is out of the manufacturing conditions of the present invention. Therefore, compared with the steel sheet of the present invention, the steel sheet showed deteriorated material and HIC resistance. The comparative steel sheets Nos. 2 and 4 had deteriorated DWTT properties, the comparative steel sheet No. 3 had deteriorated HIC properties, and the comparative steel sheets Nos. 5 and 6 had deteriorated yield ratios and DWTT properties compared with the steel sheets of the present invention.
本発明鋼B,Cにおいても、本発明の製造方法によって
製造された本発明鋼板No.7,8,10においては、良好な諸
特性が得られるが、圧延冷却条件が本発明の条件を外れ
た比較鋼板No.9においてはHIC特性が、比較鋼板No.11に
おいては降伏比、DWTT特性が本発明鋼板と比較して劣化
していることが確認された。In the steel sheets B and C of the present invention, in the steel sheet Nos. 7, 8, and 10 of the present invention manufactured by the manufacturing method of the present invention, good properties are obtained, but the rolling and cooling conditions deviate from the conditions of the present invention. It was confirmed that the comparative steel sheet No. 9 had deteriorated HIC characteristics, and the comparative steel sheet No. 11 had deteriorated yield ratio and DWTT characteristics as compared with the steel sheet of the present invention.
Ｃ含有量が本発明外の比較鋼Ｄにおいては、製造条件
が本発明の条件内であっても、靭性、耐HIC特性におい
て良好な特性が得られていない。In Comparative Steel D having a C content outside the present invention, good properties in terms of toughness and HIC resistance were not obtained even when the production conditions were within the conditions of the present invention.
［発明の効果］ 以上説明したように、この発明によれば、耐HIC特性
に優れ、且つ、高強度、高靭性、低降伏比の鋼板が得ら
れる産業上有用な効果がもたらされる。 [Effects of the Invention] As described above, according to the present invention, an industrially useful effect of obtaining a steel sheet having excellent HIC resistance, high strength, high toughness, and a low yield ratio is obtained.
第１図は供試鋼の加熱温度とDWTT試験結果との関係を示
すグラフ、第２図は供試鋼のオーステナイト未再結晶温
度域の累積圧下率とDWTT試験結果との関係を示すグラ
フ、第３図は供試鋼の圧延終了温度とHIC試験結果との
関係を示すグラフ、第４図は供試鋼の加速冷却時の冷却
速度と強度ならびに靭性との関係を示すグラフ、第５図
は供試鋼の加速冷却停止温度と強度ならびに靭性との関
係を示すグラフ、第６図は鋼中のＣ含有量と鋼板の強度
ならびに靭性との関係を示すグラフである。FIG. 1 is a graph showing the relationship between the heating temperature of the test steel and the DWTT test result, and FIG. 2 is a graph showing the relationship between the cumulative rolling reduction in the austenite non-recrystallization temperature region of the test steel and the DWTT test result. FIG. 3 is a graph showing the relationship between the rolling end temperature of the test steel and the result of the HIC test, FIG. 4 is a graph showing the relationship between the cooling rate, the strength and the toughness of the test steel during accelerated cooling, and FIG. Is a graph showing the relationship between the accelerated cooling stop temperature and the strength and toughness of the test steel, and FIG. 6 is a graph showing the relationship between the C content in the steel and the strength and toughness of the steel sheet.
Claims (1)
1250℃の温度に加熱し、次いで、加熱された前記スラブ
に対し、圧延温度：オーステナイト再結晶温度域、の条
件で第１段目の圧延を施し、次いで、圧延温度：オース
テナイト未再結晶温度域、累積圧下率:70％以上、圧延
仕上がり温度:Ar3点以上、の条件で第２段目の圧延を施
し、次いて、鋼種により定まるパーライト生成の臨界冷
却速度以上の冷却速度で、Ar3点以上の温度からAr3点−
250℃以下の温度まで冷却することを特徴とする耐HIC特
性に優れた高張力高靭性鋼板の製造方法。(1) C: 0.02 to 0.06 wt.%, Si: 0.03 to 0.50 wt.%, Mn: 0.5 to 2.5 wt.%, Nb: 0.005 to 0.100 wt.%, Ti: 0.005 to 0.100 wt.%, Al: 0.005 to 0.100 wt.%, Ca: 0.0005 to 0.008 wt.%, S: 0.004 wt.% Or less (including 0), at least one element from the group consisting of: Cu: 0.5 wt.% Or less ( 0: not included), Ni: 0.5 wt.% Or less (not including 0), Cr: 0.5 wt.% Or less (not including 0), Mo: 0.5 wt.% Or less (not including 0), V: 0.15wt.% Or less (0 is not included), and the slab consisting of the balance Fe and unavoidable impurities is 1000-
The slab was heated to a temperature of 1250 ° C., and then subjected to the first-stage rolling under the conditions of rolling temperature: austenite recrystallization temperature range, and then rolling temperature: austenite non-recrystallization temperature range. , The rolling reduction temperature: 70% or more, the rolling finish temperature: Ar 3 points or more, and then the second stage of rolling is performed. Then, at a cooling rate higher than the critical cooling rate of pearlite formation determined by the steel type, Ar 3 Ar 3 points −
A method for producing a high tensile strength and toughness steel sheet having excellent HIC resistance, characterized by cooling to a temperature of 250 ° C or less.
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|---|---|---|---|
| JP19878789A JP2870830B2 (en) | 1989-07-31 | 1989-07-31 | Method for producing high tensile strength and high toughness steel sheet excellent in HIC resistance |
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|---|---|---|---|
| JP19878789A JP2870830B2 (en) | 1989-07-31 | 1989-07-31 | Method for producing high tensile strength and high toughness steel sheet excellent in HIC resistance |
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|---|---|
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| KR20160075900A (en) * | 2014-12-19 | 2016-06-30 | 주식회사 포스코 | Steel having excellent low temperature toughness and hydrogen induced cracking resistance and manufacturing method thereof |
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| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| JP2596868B2 (en) * | 1992-01-23 | 1997-04-02 | 新日本製鐵株式会社 | Welded structure with excellent HIC resistance and SSC resistance |
| JP2647302B2 (en) * | 1992-03-30 | 1997-08-27 | 新日本製鐵株式会社 | Method for producing high-strength steel sheet with excellent resistance to hydrogen-induced cracking |
| US5545269A (en) * | 1994-12-06 | 1996-08-13 | Exxon Research And Engineering Company | Method for producing ultra high strength, secondary hardening steels with superior toughness and weldability |
| ATE321897T1 (en) * | 1997-05-19 | 2006-04-15 | American Cast Iron Pipe Co | SEWER PIPE AND STEEL STRUCTURE PRODUCED BY HIGH SPEED CONTINUOUS CASTING |
| KR20010112876A (en) * | 2000-06-15 | 2001-12-22 | 이흥무 | Compact Disc of Credit Card Type |
| JP5353156B2 (en) * | 2008-09-26 | 2013-11-27 | Ｊｆｅスチール株式会社 | Steel pipe for line pipe and manufacturing method thereof |
-
1989
- 1989-07-31 JP JP19878789A patent/JP2870830B2/en not_active Expired - Fee Related
Cited By (3)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| KR101505262B1 (en) * | 2013-01-31 | 2015-03-24 | 현대제철 주식회사 | High strength steel plate and method for manufacturing the same |
| KR20160075900A (en) * | 2014-12-19 | 2016-06-30 | 주식회사 포스코 | Steel having excellent low temperature toughness and hydrogen induced cracking resistance and manufacturing method thereof |
| KR101639902B1 (en) * | 2014-12-19 | 2016-07-15 | 주식회사 포스코 | Steel having excellent low temperature toughness and hydrogen induced cracking resistance and manufacturing method thereof |
Also Published As
| Publication number | Publication date |
|---|---|
| JPH0364414A (en) | 1991-03-19 |
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