JPH0541683B2

JPH0541683B2 -

Info

Publication number: JPH0541683B2
Application number: JP62168576A
Authority: JP
Inventors: Rikio Chijiiwa; Yoshio Terada; Hiroshi Tamehiro; Masazumi Hirai; Seiji Isoda
Original assignee: Nippon Steel Corp
Current assignee: Nippon Steel Corp
Priority date: 1987-07-08
Filing date: 1987-07-08
Publication date: 1993-06-24
Also published as: JPS6415320A

Description

[Detailed description of the invention]

〔産業上の利用分野〕本発明はとくに小入熱溶接から大入熱溶接に至
るまで熱影響部（HAZ）の低温靱性が優れ、か
つ板厚方向の特性が良好な低温用高張力鋼の製造
法に関する。鉄鋼業においては厚板ミルに適用す
ることが最も好ましいが、ホツトコイル、形鋼な
どにも適用可能である。また、この方法で製造し
た鋼は圧力容器、ラインパイプ（耐サワーガスラ
インパイプを含む）など苛酷な環境下で使用され
る溶接鋼構造物に用いることができる。〔従来の技術〕低合金鋼のHAZ靱性は、(1)結晶粒のサイズ、
(2)高炭素島状マルテンサイト（M′）、上部ベイナ
イト（Bu）などの硬化相の分散状態、(3)粒界脆
化の有無、(4)元素のミクロ偏析など種々の治金学
的要因に支配される。なかでもHAZの結晶粒の
サイズは低温靱性に大きな影響を与えることが知
られており、HAZ組織を微細化する数多くの技
術が開発実用化されている。TiNなど高温でも
比較的に安定な窒化物を鋼中に微細分散させ、こ
れによつてHAZのオーステナイト（γ）粒の粗
大化を抑制する技術はとくに有名である。しかしHAZの1400℃以上に加熱される領域で
は、TiNは粗大化もしくは溶解し、γ粒の粗大
化抑制能力は消失する。このため溶融線近傍での
靱性劣化が大きく、HAZ全域で安定して高靱性
を得ることができない。これに対してTi酸化物（主としてTi₂O₃）を微
細分散させた鋼（特開昭61−79745号公報）は溶
融線近傍でもHAZ組織を小さくすることができ、
TiN鋼に比較して優れた低温靱性が得られる。
しかし、この鋼の大入熱溶接HAZの靱性はシヤ
ルピー遷移温度−15〜−35℃程度であり、十分と
は言えない。さらに海洋構造物や耐サワーガスラインパイプ
などでは板厚方向の特性も問題となり、前者で
は、板厚方向に応力が作用するため耐ラメラーテ
ア性が、後者では耐水素誘起割れ性が要求され
る。しかしながら現在のところ溶融線近傍まで
HAZ組織を安定して微細化し、板厚方向の特性
を改善する技術は存在しない。〔発明が解決しようとする問題点〕本発明はHAZ靱性および板厚方向の特性が優
れた鋼を安価に製造する技術を提供するものであ
る。本発明法で製造した鋼は溶融線近傍から
HAZまで全域で組織が微細化し、優れた低温靱
性を示すとともに良好な耐ラメラーテア性、耐水
素誘起割れ性を有する。〔問題点を解決するための手段〕本発明の要旨は、重量％で、Ｃ；0.01〜0.15
％、Si；0.5％以下、Mn；0.5〜2.0％、Ｐ；0.025
％以下、Ｓ；0.002％以下、Al；0.004％以下、
Ca；0.001〜0.005％、Ｎ；0.0010〜0.0040％、
Ｏ；0.0010〜0.0050％を含有し、且つ、 −0.010≦〔Ca〕−2.5 〔Ｏ〕−1.25〔Ｓ〕≦０を満足する残部が鉄および不可避的不純物からな
る鋼を連続鋳造法によつてスラブとし、これを
1250℃以下の温度で再加熱後、鋼を製造すること
を特徴とする溶接部靱性の優れた低温用高張力鋼
の製造法である。更に本発明は必要によりNb；0.005〜0.060％、
Ｖ；0.005〜0.060％、Ni；0.05〜2.00％、Cu；
0.05〜1.00％、Cr；0.05〜1.00％、Mo；0.05〜
0.40％、Ｂ；0.0003〜0.0020％の一種または二種
以上を更に加え、更にZr；0.005〜0.025％、Ti；
0.005〜0.025％の一種または二種を加え、更に
REM；0.005〜0.050％を加えることができる。〔作用〕本発明者らの研究によれば、HAZ靱性は、１）
鋼の化学成分、２）組織（結晶粒の大きさと硬化
相の分布状態）に大きく依存し、鋼成分の適正化
とこれによる結晶粒の微細化が高靱性化に不可欠
であると考えられた。また板厚方向の特性を改善
するには、鋼のＳ量の低減とCaによる硫化物系
介在物の形態制御が不可欠である。そこで極低Ｓ鋼にCaを添加し、微細なCa酸化
物（主としてCaO）を形成、分散させ、これによ
つて組織を微細化、かつ残存するCaでＳを固定
する新しい方法を発明した。 Ca酸化物はγ粒の粗大化抑制能力は小さいが、
γ−α変態時にγ粒内に存在するCaOを核とし
て、放射状に微細なアシキユラーフエライト
（AF）が生成するので、HAZ組織は著しく微細
化する。CaOは溶融線近傍の1400℃以上に加熱さ
れる領域でも安定であり、この領域でも組織の微
細化に効果を発揮する。その結果、溶接部は全域
にわたつて微細化し、極めて優れた低温靱性が得
られる。酸化物によつてHAZ靱性を改善する方法には、
特開昭61−79745号公報のようにTi酸化物を利用
するものもあるが、Ca酸化物とTi酸化物でAF生
成能力に大きな差はない。CaOはTi₂O₃よりも生
成温度が高く、凝固冷却速度の影響を受け難いの
で、スラブ全厚にわたつて均一微細分散が可能な
点で優れている。さらに酸素と結合して残つた
Caは、Ｓと結合して圧延によつて伸びないCaS
を形成して、耐ラメラーテア性、耐水素誘起割れ
性をも向上させる。通常の製鋼法において鋼中にCaOを微細分散さ
せ、CaSを形成させるためには、とくにCa、Ｏ
およびＳ量とそのバランスの適正化が必須であ
る。このためCa、Ｏ、Ｓ量をそれぞれCa：0.01
−0.005％、Ｏ：0.0010−0.0050％、Ｓ：0.002％
以下に限定し、かつそのバランスを−0.010％≦
［Ca］−2.5［Ｏ］−1.25［Ｓ］≦０％とする必要があ
る。 Ca、Ｏ量の下限は、HAZにおいてCaOを生成
するための必要最少量である。Caの上限はCa系
の大型介在物やクラスターの生成による鋼の内質
劣化を防止するためである。Ｏの上限は非金属介
在物の生成による鋼の清浄度、靱性の劣化を防止
するためである。またＳ量の上限0.002％は板厚方向の特性を改
善するためである。Ｓ量が多過ぎるとCa添加量
を増加させねばならず、鋼の清浄度が悪化する。
Ｓ量は低いほど良く、0.0010％以下が望ましい。しかし単に個々の元素量を限定するだけでは微
細なCaOを安定して得ることはできないし、Ca
添加による硫化物系介在物の形態制御もできない
ので、Ca、Ｏ、Ｓ量のバランスを−0.010％≦
［Ca］−2.5［Ｏ］−1.25［Ｓ］≦０％の狭い範囲に限
定した。ここで、Tiを添加した場合、CaはTiより酸素
との結合力が強く、CaOは溶鋼段階ではすでに生
成する。Tiは溶鋼の凝固過程ですでに生成して
いるCaOを核に生成し、Ca−Ti−Ｏの複合酸化
物を生成する。この複合酸化物はCaOと同等の効
果を持ち、酸化物の数もCa単独添加の場合とか
わらない。従つて、Tiを添加した場合でも、Caの規定式
には、考慮しないでも良い。 Ca、Ｏ、Ｓ量が、この範囲にあるとHAZ靱性
および耐ラメラーテア性、耐水素誘起割れ性は飛
躍的に向上する。下限はCa量の不足によるCaO
の生成量の不足を防ぐためであり、上限は過剰の
Caによる清浄度の悪化を防止するためである。しかし、たとえCaOが鋼中に微細分散し、硫化
物の形態制御が完璧でも基本成分が適当でない
と、優れたHAZ靱性および耐ラメラーテア性、
耐水素誘起割れ性は得られない。以下、この点について説明する。Ｃの下限0.01％は、母材および溶接部の強度の
確保ならびにNb、Ｖなどの添加時に、これらの
効果を発揮させるための最小量である。しかしＣ
量が多過ぎると、HAZ靱性に悪影響をおよぼす
だけでなく母材靱性、溶接性をも劣化させるの
で、上限を0.15％とした。Ｃ量が多いとHAZに
マルテンサイト（M′）、凝似パーライト（P′）が
生成して低温靱性を著しく劣化させる。 Siは脱酸上、鋼に含まれる元素であるが、多く
添加すると溶接性、接合部の靱性が劣化するた
め、上限を0.5％に限定した。 Mnは強度、靱性を確保する上で不可欠な元素
であり、その下限は0.5％である。HAZ靱性を改
善するには、γ粒界に生成する粗大な初析フエラ
イトを防止する必要があるが、Mn添加は、これ
を抑制する効果がある。しかしMn量が多過ぎると焼入性が増加して溶
接性、HAZ靱性を劣化させるだけでなく、スラ
ブの中心偏析を助長して板厚方向の特性を劣化さ
せるので上限を2.0％とした。本発明鋼において不純物であるＰを0.025％以
下とした。これは、母材、HAZの低温靱性をよ
り一層向上させ、スラブの中心偏析を軽減するた
めである。Ｐ量の低減は、HAZにおける粒界破
壊傾向を減少させる傾向がある。最も好ましいＰ
量は0.010％以下である。 Alは、一般に脱酸上鋼に含まれる元素である
が、本発明鋼では好ましくない元素であり、その
上限を0.004％とした。これはAlが鋼中に含まれ
ているとＯと結合してCaOができないためであ
る。脱酸はCaだけでも可能であり、本発明にお
いてAl量は少ないほど良く、0.002以下が望まし
い。つぎに選択元素の限定理由について説明する。 Nbは高強度と良好な靱性を得るために有効な
元素であり、本発明鋼においても同様である。こ
の効果を得るためには最低0.005％のNb量が必要
である。しかしながら、Nb量が多すぎると溶接
性を劣化させるので、その上限を0.060％とした。 Tiは本発明鋼に添加するとTiNを形成して
HAZの1400℃以下に加熱された領域の組織を微
細化するので、大入熱溶接用鋼に添加することが
望ましい。下限の0.005％は、この効果を得るた
めの最小量であり、また上限の0.025％はTiC形
成によるHAZ靱性の劣化を防止するためである。 ZrはほぼTiと同様の効果をもつ元素である。
その上下限は、それぞれ0.005％、0.025％であ
る。つぎにＶ、Ni、Cu、Cr、Mo、Ｂ、REMを添
加する理由について説明する。基本となる成分にさらに、これらの元素を添加
する主たる目的は、本発明鋼の優れた特徴を損な
うことなく、強度、靱性など特性の向上をはかる
ためである。したがつて、その添加量は自ら制限
されるべき性質のものである。ＶはNbとほぼ同じ効果を持つ元素であるが、
0.005％以下では効果が少なく、上限は0.060％ま
で許容できる。 Niは溶接性、HAZ靱性に悪影響をおよぼすこ
となく、母材の強度、靱性を向上させるが、2.0
％を超えると溶接性に好ましくないため上限を
2.0％とした。 CuはNiとほぼ同様の効果とともに耐食性、耐
水素誘起割れ性などにも効果があるが、1.0％を
超えると熱間圧延時にCu−クラツクが発生し、
製造困難となる。このため上限を1.0％とした。 Crは母材、溶接部の強度を高めるが、多過ぎ
ると溶接性やHAZ靱性を劣化させる。その上限
は1.0％である。 Moは母材の強度、靱性をともに向上させる元
素であるが、多過ぎるとCrと同様に母材、HAZ
の靱性、溶接性の劣化を招き好ましくない。その
上限は0.40％である。なおCr、Moの添加量の下限は、材質上での効
果が得られるための最小量とすべきで、いずれも
0.05％である。Ｂは鋼の焼入性を増大させ強度を増加させる元
素である。HAZのγ粒界に偏析した固溶Ｂはフ
エライトの生成を抑制し、CaOからの微細なAF
の生成を助ける。またＮと結合したBNはフエラ
イト発生核としての作用をもちHAZ組織を微細
化する。このようなＢの効果を得るためには、最低
0.0003％のＢ量が必要である。しかしＢ量が多過
ぎるとFe₂₃（CB）₆などの粗大な析出物がγ粒界に
析出して低温靱性を劣化させる。このためＢ量の
上限を0.002％に制限する必要がある。 REMはCaと同様に硫化物（MnS）の形態を制
御し、低温靱性を向上（シヤルピー吸収エネルギ
ーを増加）させるほか、耐水素誘起割れ性の改善
にも効果を発揮する。しかし0.005％以下では実
用上効果がなく、また0.05％を超えて添加すると
REM−Ｏ、REM−Ｓが多量に生成して大型介在
物となり、鋼の靱性のみならず清浄度も害し、ま
た溶接性にも悪影響を与える。このため添加量の
範囲を0.005〜0.05％に制限した。鋼の成分を上記のように限定しても、製造法が
適切でなければ溶接前の鋼中に微細なCaOを分散
させることはできない。このため製造条件につい
ても限定する必要がある。まず、この鋼は工業的
には連続鋳造法で製造することが必須である。こ
の理由は連続鋳造法では溶鋼の凝固冷却速度が速
くスラブ中に微細なCaOが多量に得られるためで
ある。大型鋼塊による造塊−分塊法では、CaOを
スラブ中に微細分散させることは難しい。連続鋳造法の場合、スラブ厚によつて冷却速度
が異なるが、HAZ靱性の観点からその厚みは350
mm以下が望ましい。さらにスラブの再加熱温度を
1250℃以下とする必要がある。これ以上の温度で
再加熱するとγ粒が粗大化して、母材の低温靱性
が劣化する。なお本発明においては、スラブの再
加熱は必ずしも実施する必要はなく、ホツトチヤ
ージ圧延^*やダイレクト圧延^**を行つても全く問
題はない。ホツトチヤージ圧延^* 加熱炉挿入時のスラブ温度：700〜850℃ 加熱炉温度：1100〜1250℃ （圧延以降の条件は目標材質で決定する）ダイレクト圧延^** 圧延開始温度：1050〜930℃ （圧延以降の条件は目標材質で決定する）本発明ではスラブ再加熱後の圧延法などについ
ては、とくに限定しないが、いわゆる加工熱処理
や圧延後の焼入焼戻、焼ならし処理が強度、靱性
を確保する上で適切である。これは、たとえ優れ
たHAZ靱性が得られても母材の靱性が劣つてい
ると鋼材としては不十分なためである。母材の低
温靱性を優れたものとするには鋼の結晶粒を微細
化する必要がある。加工熱処理の方法としては、１）制御圧延、
２）制御圧延−加速冷却、３）圧延直接焼入−焼
戻などが挙られるが、最も好ましいのは制御圧延
と加速冷却の組み合わせである。なお、この鋼を
製造後、脱水素などの目的でAC₁変態点以下の温
度に再加熱しても本発明の特徴を損なうものでは
ない。〔実施例〕転炉−連続鋳造−厚板工程で種々の鋼成分の鋼
板（厚み30mm）を製造し、溶接熱サイクル再現装
置を使用してHAZ靱性を２mmＶノツチシヤルピ
ー試験によつて調査した。再現熱サイクル試験
は、1/4ｔ位置から採取し試験片を用い、ピーク
温度（最高到達温度）1400℃、800〜500℃の冷却
時間を192、54秒の２条件で行なつた。これらの条件はそれぞれ溶接入熱200、50KJ／
cmに相当する溶融線近傍の熱サイクルを模したも
のである。さらに板厚方向の特性を直径10mmの板厚方向引
張試験片によつて調べた。表１に実施例を示す。本発明法で製造した鋼板（本発明鋼）は全て良
好な母材特性、板厚方向特性およびHAZ靱性を
有するのに対して、本発明法によらない比較鋼は
いずれかの特性が劣り、厳しい環境下で使用され
る鋼板として適切でない。比較鋼において鋼15はＣ量が多過ぎるために
HAZ靱性が劣る。鋼16は低Ｃ量にもかかわらず、
Mn量が多過ぎるためにやはりHAZ靱性が悪い。
鋼17はＳ量が高く、かつCa量がすくないために、
CaOの生成量が不足、さらに硫化物の形態制御も
不完全になり、HAZ靱性、板厚方向特性が劣る。鋼18は鋼成分は本発明の条件を満足している
が、大型鋼塊−分塊法で製造したために微細な
CaOの量が不足し、HAZ靱性が十分でない。鋼
19はAl量が0.008％と多く、やはりCaOの量が不
足したために、HAZ靱性が劣る。また鋼20は鋼
成分において酸素が高過ぎるために母材靱性が不
十分である。 [Industrial Application Field] The present invention is particularly directed to the use of high-strength steel for low-temperature use, which has excellent low-temperature toughness in the heat-affected zone (HAZ) and good properties in the thickness direction, from low heat input welding to high heat input welding. Regarding manufacturing methods. In the steel industry, it is most preferably applied to plate mills, but it can also be applied to hot coils, shaped steel, etc. Furthermore, the steel produced by this method can be used in welded steel structures used in harsh environments such as pressure vessels and line pipes (including sour gas resistant line pipes). [Conventional technology] The HAZ toughness of low alloy steel is determined by (1) grain size;
(2) Dispersion state of hardened phases such as high-carbon island martensite (M′) and upper bainite (Bu), (3) presence or absence of grain boundary embrittlement, and (4) various metallurgy such as micro-segregation of elements. controlled by factors. In particular, it is known that the grain size of HAZ has a large effect on low-temperature toughness, and a number of techniques have been developed and put into practical use to refine the HAZ structure. A particularly well-known technique is to finely disperse nitrides such as TiN, which are relatively stable even at high temperatures, into steel, thereby suppressing the coarsening of austenite (γ) grains in the HAZ. However, in the region of the HAZ heated to 1400°C or higher, TiN coarsens or dissolves, and the ability to suppress the coarsening of γ grains disappears. For this reason, the toughness deteriorates significantly near the fusion line, making it impossible to stably obtain high toughness throughout the HAZ. On the other hand, steel in which Ti oxide (mainly Ti ₂ O ₃ ) is finely dispersed (Japanese Unexamined Patent Publication No. 61-79745) can reduce the HAZ structure even near the fusion line,
Superior low-temperature toughness can be obtained compared to TiN steel.
However, the toughness of the high heat input welding HAZ of this steel is at a Charpy transition temperature of -15 to -35°C, which cannot be said to be sufficient. Furthermore, properties in the thickness direction are also an issue for offshore structures and sour gas line pipes, etc. The former requires lamellar tear resistance because stress acts in the thickness direction, and the latter requires hydrogen-induced cracking resistance. However, at present it is close to the melting line.
There is no technology that stably refines the HAZ structure and improves the properties in the thickness direction. [Problems to be Solved by the Invention] The present invention provides a technology for inexpensively producing steel with excellent HAZ toughness and properties in the thickness direction. The steel produced by the method of the present invention starts from the vicinity of the fusion line.
The structure is refined throughout the entire area up to the HAZ, and it exhibits excellent low-temperature toughness, as well as good lamellar tear resistance and hydrogen-induced cracking resistance. [Means for solving the problems] The gist of the present invention is that C: 0.01 to 0.15 in weight%
%, Si; 0.5% or less, Mn; 0.5-2.0%, P; 0.025
% or less, S; 0.002% or less, Al; 0.004% or less,
Ca; 0.001-0.005%, N; 0.0010-0.0040%,
A steel containing O; 0.0010 to 0.0050% and satisfying -0.010≦[Ca]-2.5 [O]-1.25[S]≦0, the balance consisting of iron and unavoidable impurities, is produced by a continuous casting method. Slab this
This is a method for producing high-strength steel for low temperature use with excellent weld toughness, which is characterized by producing steel after reheating at a temperature of 1250°C or less. Furthermore, the present invention may contain Nb; 0.005 to 0.060%, if necessary.
V; 0.005-0.060%, Ni; 0.05-2.00%, Cu;
0.05~1.00%, Cr; 0.05~1.00%, Mo; 0.05~
0.40%, B; 0.0003 to 0.0020%, one or more of them further added, Zr; 0.005 to 0.025%, Ti;
Add 0.005 to 0.025% of one or two kinds, and
REM: 0.005-0.050% can be added. [Effect] According to the research conducted by the present inventors, HAZ toughness is 1)
High toughness largely depends on the chemical composition of the steel, and 2) its structure (crystal grain size and hardened phase distribution), and optimization of the steel composition and the resulting refinement of the grain size were considered essential for achieving high toughness. . Furthermore, in order to improve properties in the thickness direction, it is essential to reduce the amount of S in steel and control the morphology of sulfide-based inclusions using Ca. Therefore, we invented a new method of adding Ca to ultra-low S steel, forming and dispersing fine Ca oxides (mainly CaO), thereby refining the structure, and fixing S with the remaining Ca. Although Ca oxide has a small ability to suppress coarsening of γ grains,
During the γ-α transformation, fine axial ferrite (AF) is generated in a radial manner using CaO present in the γ grains as nuclei, so the HAZ structure becomes significantly finer. CaO is stable even in the region heated to over 1400°C near the melting line, and is effective in refining the structure even in this region. As a result, the welded area becomes finer over the entire area, resulting in extremely excellent low-temperature toughness. Methods for improving HAZ toughness with oxides include:
Although there are some methods that use Ti oxide, such as in Japanese Patent Application Laid-open No. 61-79745, there is no big difference in AF generation ability between Ca oxide and Ti oxide. CaO has a higher formation temperature than Ti ₂ O ₃ and is less affected by the solidification cooling rate, so it is superior in that it can be uniformly and finely dispersed over the entire thickness of the slab. Furthermore, it combines with oxygen and remains
Ca combines with S and does not elongate by rolling.
This also improves lamellar tear resistance and hydrogen-induced cracking resistance. In order to finely disperse CaO in steel and form CaS in the normal steel manufacturing method, especially Ca, O
It is also essential to optimize the amount of S and its balance. For this reason, the amounts of Ca, O, and S are each Ca: 0.01
-0.005%, O: 0.0010-0.0050%, S: 0.002%
Limited to the following, and the balance is -0.010%≦
[Ca]-2.5[O]-1.25[S]≦0%. The lower limit of the amount of Ca and O is the minimum amount necessary to generate CaO in the HAZ. The upper limit of Ca is set to prevent deterioration of the internal quality of steel due to the formation of large Ca-based inclusions and clusters. The upper limit of O is set to prevent deterioration of the cleanliness and toughness of the steel due to the formation of nonmetallic inclusions. The upper limit of the S content is 0.002% to improve the properties in the thickness direction. If the amount of S is too large, the amount of Ca added must be increased, and the cleanliness of the steel will deteriorate.
The lower the S content, the better, and desirably 0.0010% or less. However, it is not possible to stably obtain fine CaO by simply limiting the amount of each element, and
Since it is not possible to control the morphology of sulfide-based inclusions by addition, the balance of Ca, O, and S amounts should be maintained at -0.010%≦
It was limited to a narrow range of [Ca]-2.5[O]-1.25[S]≦0%. Here, when Ti is added, Ca has a stronger binding force with oxygen than Ti, and CaO is already generated at the molten steel stage. Ti is generated by nucleating CaO, which has already been generated during the solidification process of molten steel, and forms a complex oxide of Ca-Ti-O. This composite oxide has the same effect as CaO, and the number of oxides is the same as when Ca is added alone. Therefore, even if Ti is added, it may not be considered in the formula for Ca. When the amounts of Ca, O, and S are within this range, HAZ toughness, lamellar tear resistance, and hydrogen-induced cracking resistance are dramatically improved. The lower limit is CaO due to insufficient Ca amount
This is to prevent insufficient production of
This is to prevent deterioration of cleanliness due to Ca. However, even if CaO is finely dispersed in the steel and the sulfide morphology is perfectly controlled, if the basic components are not appropriate, excellent HAZ toughness and lamellar tear resistance may occur.
Hydrogen-induced cracking resistance cannot be obtained. This point will be explained below. The lower limit of 0.01% of C is the minimum amount to ensure the strength of the base metal and welded part and to exhibit these effects when adding Nb, V, etc. But C
If the amount is too large, it not only adversely affects the HAZ toughness but also deteriorates the base metal toughness and weldability, so the upper limit was set at 0.15%. When the amount of C is large, martensite (M') and aggregated pearlite (P') are generated in the HAZ, which significantly deteriorates the low-temperature toughness. Si is an element contained in steel for deoxidation purposes, but if too much is added, weldability and toughness of joints deteriorate, so the upper limit was limited to 0.5%. Mn is an essential element for ensuring strength and toughness, and its lower limit is 0.5%. In order to improve HAZ toughness, it is necessary to prevent coarse pro-eutectoid ferrite from forming at the γ grain boundaries, and the addition of Mn has the effect of suppressing this. However, if the amount of Mn is too large, it not only increases hardenability and deteriorates weldability and HAZ toughness, but also promotes segregation at the center of the slab and deteriorates properties in the thickness direction, so the upper limit was set at 2.0%. In the steel of the present invention, P, which is an impurity, is set to 0.025% or less. This is to further improve the low-temperature toughness of the base metal, HAZ, and reduce center segregation of the slab. Reducing the amount of P tends to reduce the tendency for intergranular fracture in the HAZ. Most preferred P
The amount is 0.010% or less. Al is an element that is generally included in deoxidized steel, but it is an undesirable element in the steel of the present invention, and its upper limit was set at 0.004%. This is because when Al is contained in steel, it combines with O to form CaO. Deoxidation is possible with Ca alone, and in the present invention, the smaller the amount of Al, the better, and preferably 0.002 or less. Next, the reason for limiting the selected elements will be explained. Nb is an effective element for obtaining high strength and good toughness, and the same applies to the steel of the present invention. To obtain this effect, a minimum Nb content of 0.005% is required. However, if the amount of Nb is too large, weldability deteriorates, so the upper limit was set at 0.060%. When Ti is added to the steel of the present invention, it forms TiN.
It is desirable to add it to steel for high heat input welding because it refines the structure of the region of the HAZ heated to 1400°C or less. The lower limit of 0.005% is the minimum amount to obtain this effect, and the upper limit of 0.025% is to prevent deterioration of HAZ toughness due to TiC formation. Zr is an element that has almost the same effect as Ti.
The upper and lower limits are 0.005% and 0.025%, respectively. Next, the reason for adding V, Ni, Cu, Cr, Mo, B, and REM will be explained. The main purpose of adding these elements to the basic components is to improve properties such as strength and toughness without impairing the excellent characteristics of the steel of the present invention. Therefore, the amount added must be limited. V is an element that has almost the same effect as Nb, but
The effect is small below 0.005%, and an upper limit of 0.060% is acceptable. Ni improves the strength and toughness of the base metal without adversely affecting weldability and HAZ toughness, but 2.0
If it exceeds %, it is not favorable for weldability, so set the upper limit.
It was set at 2.0%. Cu has almost the same effects as Ni, as well as corrosion resistance and hydrogen-induced cracking resistance, but if it exceeds 1.0%, Cu cracks will occur during hot rolling.
It becomes difficult to manufacture. For this reason, the upper limit was set at 1.0%. Cr increases the strength of the base metal and weld zone, but too much Cr deteriorates weldability and HAZ toughness. Its upper limit is 1.0%. Mo is an element that improves both the strength and toughness of the base metal, but if it is too much, it will cause damage to the base metal and HAZ like Cr.
This is undesirable because it causes deterioration in toughness and weldability. Its upper limit is 0.40%. The lower limit of the amount of Cr and Mo added should be the minimum amount to achieve the desired effect on the material.
It is 0.05%. B is an element that increases the hardenability and strength of steel. The solid solution B segregated at the γ grain boundaries of the HAZ suppresses the formation of ferrite, and the fine AF from CaO
helps generate. Furthermore, BN combined with N acts as a ferrite generation nucleus and refines the HAZ structure. In order to obtain this effect of B, the minimum
A B content of 0.0003% is required. However, if the amount of B is too large, coarse precipitates such as Fe ₂₃ (CB) ₆ will precipitate at the γ grain boundaries, deteriorating the low-temperature toughness. Therefore, it is necessary to limit the upper limit of the amount of B to 0.002%. Like Ca, REM controls the morphology of sulfide (MnS), improves low-temperature toughness (increases Charpy absorbed energy), and is also effective in improving hydrogen-induced cracking resistance. However, if it is less than 0.005%, it has no practical effect, and if it is added in excess of 0.05%,
REM-O and REM-S are generated in large quantities and become large inclusions, which impair not only the toughness but also the cleanliness of the steel, and also have an adverse effect on the weldability. For this reason, the range of addition amount was limited to 0.005 to 0.05%. Even if the steel components are limited as described above, fine CaO cannot be dispersed in the steel before welding unless the manufacturing method is appropriate. For this reason, it is also necessary to limit the manufacturing conditions. First, industrially, it is essential to manufacture this steel using a continuous casting method. The reason for this is that in the continuous casting method, the solidification and cooling rate of molten steel is fast and a large amount of fine CaO is obtained in the slab. In the ingot-blending method using large steel ingots, it is difficult to finely disperse CaO in the slab. In the case of continuous casting, the cooling rate varies depending on the slab thickness, but from the perspective of HAZ toughness, the thickness is 350 mm.
Desirably less than mm. Furthermore, the reheating temperature of the slab
Must be below 1250℃. If reheated at a temperature higher than this, the γ grains will become coarser and the low-temperature toughness of the base material will deteriorate. Note that in the present invention, it is not necessary to reheat the slab, and there is no problem at all even if hot charge rolling ^* or direct rolling ^** is performed. Hot charge rolling ^* Slab temperature when inserted into heating furnace: 700-850℃ Heating furnace temperature: 1100-1250℃ (Conditions after rolling are determined by target material) Direct rolling ^** Rolling start temperature: 1050-930℃ (After rolling In the present invention, the rolling method after reheating the slab is not particularly limited, but so-called processing heat treatment, quenching and tempering after rolling, and normalizing treatment ensure strength and toughness. It is appropriate to do so. This is because even if excellent HAZ toughness is obtained, if the toughness of the base material is poor, the steel material will not be sufficient. In order to improve the low-temperature toughness of the base material, it is necessary to refine the grains of steel. Methods of processing heat treatment include 1) controlled rolling;
Examples include 2) controlled rolling-accelerated cooling, and 3) rolling direct quenching-tempering, but the most preferred is a combination of controlled rolling and accelerated cooling. Note that even if this steel is reheated to a temperature below the AC ₁ transformation point for the purpose of dehydrogenation or the like after production, the features of the present invention will not be impaired. [Example] Steel plates (thickness: 30 mm) of various steel compositions were manufactured using a converter-continuous casting-thick plate process, and their HAZ toughness was investigated by a 2 mm V notch pea test using a welding thermal cycle reproduction device. The reproducible thermal cycle test was conducted using a test piece taken from the 1/4 t position under two conditions: a peak temperature (maximum temperature) of 1400°C and a cooling time of 800 to 500°C of 192 seconds and 54 seconds. These conditions are welding heat input of 200 and 50KJ/
This model simulates the thermal cycle near the melting line, which corresponds to cm. Furthermore, the properties in the thickness direction were investigated using a tensile test piece in the thickness direction with a diameter of 10 mm. Examples are shown in Table 1. The steel sheets manufactured by the method of the present invention (inventive steel) all have good base material properties, properties in the thickness direction, and HAZ toughness, whereas comparative steels that are not made by the method of the present invention are inferior in any of the properties. Not suitable for steel plates used in harsh environments. Among comparative steels, Steel 15 has too much C content.
HAZ toughness is poor. Despite the low C content of steel 16,
The HAZ toughness is still poor due to the excessive amount of Mn.
Steel 17 has a high S content and a low Ca content, so
The amount of CaO produced is insufficient, and the morphological control of sulfides is also incomplete, resulting in poor HAZ toughness and properties in the thickness direction. Steel 18 satisfies the conditions of the present invention in terms of steel composition, but because it was produced by the large steel ingot-blooming method, it contains fine particles.
The amount of CaO is insufficient and the HAZ toughness is insufficient. steel
In No. 19, the amount of Al was as high as 0.008%, and the amount of CaO was insufficient, so the HAZ toughness was poor. Further, Steel 20 has insufficient base metal toughness because the oxygen content in the steel composition is too high.

【表】【table】

〔Effect of the invention〕

本発明により、母材はもとよりHAZ全域にお
いて優れた低温靱性を有し、かつ耐ラメラーテア
性、耐水素誘起割れ性の良好な鋼を大量、且つ安
価に製造することが可能になつた。その結果、溶
接構造物の施工能率が著しく向上するとともにそ
の安全を大きく向上させることができた。 The present invention has made it possible to manufacture steel in large quantities and at low cost, which has excellent low-temperature toughness not only in the base material but also throughout the HAZ, and has good lamellar tear resistance and hydrogen-induced cracking resistance. As a result, we were able to significantly improve the construction efficiency of welded structures, as well as their safety.

Claims

[Claims] 1% by weight: C; 0.01-0.15% Si; 0.5% or less Mn; 0.5-2.0% P; 0.025% or less S; 0.002% or less Al; 0.004% or less Ca; 0.001-0.005% N ; 0.0010 to 0.0040% O; 0.0010 to 0.0050%, and the balance satisfies -0.010≦[Ca]-2.5 [O]-1.25[S]≦0, with the balance consisting of iron and unavoidable impurities. A slab is made by casting method, and this is
A method for manufacturing high-strength steel for low temperature use with excellent weld toughness, which is characterized by manufacturing the steel after reheating at a temperature of 1250°C or lower. 2 In weight%, C; 0.01-0.15% Si; 0.5% or less Mn; 0.5-2.0% P; 0.025% or less S; 0.002% or less Al; 0.004% or less Ca; 0.001-0.005% N; 0.0010-0.0040% O 0.0010~0.0050% Nb; 0.005~0.060% V; 0.005~0.060% Ni; 0.05~2.00% Cu; 0.05~1.00% Cr; 0.05~1.00% Mo; 0.05~0.40% B; 0.0003~0.0020 %, and the balance is iron and unavoidable impurities, and the steel satisfies -0.010≦[Ca]-2.5 [O]-1.25[S]≦0 by a continuous casting method. and make it into a slab.
A method for manufacturing high-strength steel for low temperature use with excellent weld toughness, which is characterized by manufacturing the steel after reheating at a temperature of 1250°C or lower. 3. Claim 1 or Claim 2, characterized in that the composition of the steel is, in weight%, one or two of the following: Zr; 0.005 to 0.025%; Ti; 0.005 to 0.025%.
A method for manufacturing a high-strength steel for low temperature use with excellent weld toughness as described in 2. 4. The welding according to claim 1, claim 2, or claim 3, characterized in that the steel composition is expressed in weight% and REM; 0.005 to 0.050% is further added. A manufacturing method for low-temperature high-strength steel with excellent partial toughness.