JPS6145689B2

JPS6145689B2 -

Info

Publication number: JPS6145689B2
Application number: JP57139897A
Authority: JP
Inventors: Yoshikuni Tokunaga; Noryuki Iida; Masato Yamada
Original assignee: Nippon Steel Corp
Current assignee: Nippon Steel Corp
Priority date: 1982-08-13
Filing date: 1982-08-13
Publication date: 1986-10-09
Also published as: JPS5931827A

Description

【発明の詳細な説明】[Detailed description of the invention]

本発明は超深絞り用焼付硬化性鋼板の製造方法
に関するものである。近年、自動車産業界では車体軽量化による燃費
向上と安全性の追求から高強度鋼板に対する要望
が高まりつつある。一方、自動車の販売性は車体
のスタイリングで左右される風潮にあることか
ら、従来に増して鋼板のプレス成形性が重要視さ
れてきた。このような背景から、プレス成形時に
は良好な成形性を示し、塗装焼付後に、降伏強
度、引張強度の上昇する特性、即ち焼付硬化性を
有する鋼板に対する要求が高まつている。本発明
はかかる要求を満足する超深絞り用焼付硬化性鋼
板の製造方法に関するものである。焼付硬化性を有しない超深絞り用鋼板の製造方
法に関するものとしては、Tiキルド鋼板（特公
昭44−18066号公報）及びNbキルド鋼板（特公昭
54−1245号公報）の２つの系統のものが知られて
いる。しかしながらこれらの鋼板は、鋼板中の
Ｃ、Ｎを完全にTiあるいはNb等の析出物として
固定しているために、プレス成形後の塗装焼付時
に歪時効現象が起こらず、従つて焼付硬化性を有
しないものである。次に、焼付硬化性を付与した超深絞り用鋼板の
製造法としては、Nb添加鋼において鋼中のＣ、
Ｎ、Al含有量に応じてNbを添加し、Nb（at
％）／｛固溶Ｃ（at％）＋固溶Ｎ（at％）｝をある
範囲内に制限することにより、鋼板中の固溶Ｃ、
固溶Ｎ量をコントロールし、さらに焼鈍後の冷却
速度を制御することを特徴とする方法（特公昭55
−141526号、55−141555号公報）がある。しかしながら、実際に調査検討してみると、か
かる製造法には次のような欠点がある。まず熱延
巻取温度、焼鈍温度、焼鈍後の冷却速度に対する
制限である。Nb添加鋼では熱延で高温巻取（巻
取温度≧700℃）を必要とする。通常の巻取温度
では完全再結晶温度が非常に高くなつて、連続焼
鈍炉の可能温度範囲（通常は約850℃以下）では
未再結晶部が残つていたり、またNb量の多少に
よつて材質の変動が大きい。高温巻取を行なつた
場合には、熱延コイルのコイル長手方向端部を除
いては約800〜850℃の焼鈍温度で高いｒ値の鋼板
が得られることは種々報告されている通りであ
る。しかし、高温巻取を行なうということは、ス
ケールが厚くなり酸洗能率を極端に落すだけでな
く、コイル端部は冷却速度が速いために通常の巻
取温度と同じ程度の材質となり、十分な材質が得
られないので歩留の低下はNbキルド鋼では特に
大きいものがある。第２は焼鈍温度と焼鈍後の冷却速度の問題であ
る。前記の公知技術にある如く、高温（約900℃
以上）で焼鈍すると熱延で析出していたNb
（Ｃ，Ｎ）あるいはAINが再溶解するために固溶
Ｎ、固溶Ｃが増加し時効性が大きくなる。また、
焼鈍後の冷却速度を該特許にあるように、400℃
までを50℃／sec以下あるいは、400℃までを50
℃／sec以上かつ400℃から200℃までを10℃／sec
以下という様に、徐冷しなければ、固溶Ｎ、固溶
Ｃの影響により遅時効性とならないのである。
Nb添加鋼の場合に特に問題となるのは、高温で
再溶解するAINであり、常温時効性に対して問題
となるのはこのＮ時効によるものである。第３の欠点は亜鉛めつき特性である。一般に溶
融亜鉛めつきを行なつた場合、鋼中のSi含有量が
約0.3％を越えるとめつき密着性が劣化するとさ
れている。これは鋼板表面に濃化するSiによつて
薄い酸化膜が生成され、地鉄と溶融亜鉛との反応
が良好に進まないためとされる。かかる理由から
溶融亜鉛めつきを行なう材料では、表面濃化Si量
を抑えるため鋼中のSi含有量を約0.3％以下に制
限して製造に供してきた。しかしながら引張強度
40Kg／mm^２以上の鋼板を製造する場合、二次加工
割れの問題からＰ含有量が制限され、その他の合
金元素を添加して高強度化しなければならないに
もかかわらず、かかるめつき密着性の観点からSi
およびMnを添加できないため、実質上40Kg／mm
^２を超える引張強度を有する性能の良好な亜鉛め
つき超深絞り用焼付硬化性鋼板は非常に難しい。
以上のSiの表面濃化現象は、徐冷却時における一
般現象であり、現在まで解決されない問題点であ
つた。本発明者等は、従来技術であるNb添加鋼のも
つこれらの欠点を詳細に検討した結果、本質的な
原因は次の３点であるという結論を得た。 (1)熱延巻取時の温度履歴に大きく影響される
AINの析出。(2)焼鈍温度および焼鈍後の冷却速度
に依存して固溶Ｃ、固溶Ｎ量が変化すること。(3)
焼鈍後の冷却速度に依存する合金元素の粒界ある
いは表面への濃化の３点である。 (1) については、実際Nb添加鋼ではNbは炭化物
を形成するが、窒化物形成傾向は比較的弱く、
窒素はAINとして析出している。AINは通常の
巻取温度では形成され難く、巻取温度を700℃
以上にしないと熱延板中では析出せず、冷延後
の連続焼鈍時に微細析出物として析出し、結晶
粒の成長を抑制し、降伏強度を高くしたり、伸
びを低下させる等の材質劣化を引き起こす。従つて、高温巻取をした場合でも、熱延コイ
ル長手方向中心部はAINが析出するが、冷却速
度の速いコイル長手方向前後端部では巻取時に
AINが析出しないことから通常巻取時と同様、
降伏強度の上昇、伸び、ｒ値の極端な低下をも
たらすものである。また、熱延巻取温度の微妙
なバラツキやコイルの中央部、端部ではAINの
析出の程度に差がでてくるために、熱延コイル
の前、後端部の材質劣化やコイル内材質のバラ
ツキが生ずるものである。 (2) については、高温焼鈍を行なうと、熱延巻取
時に析出していたNb（Ｃ，Ｎ）やAIN等の析
出物が、その温度における溶解度積になる如く
再溶解し、固溶Ｃ，固溶Ｎが増加する。特に
AIN溶解に起因する固溶Ｎは固溶Ｃに比較して
常温時効性が大きいため、遅時効性という目標
材質に反することになる。また、焼鈍時に存在
する固溶Ｃ、固溶Ｎはその後の冷却過程でNb
（Ｃ，Ｎ）あるいはAINとして析出するが、特
にAINの析出速度は約815℃にピークを持ち低
温になる程小さくなるため、高温域での徐冷却
が必要となる。該特許では、400℃までを徐冷
もしくは400℃までを急冷して400℃から200℃
までを徐冷しなければならないとしているが、
実質上400℃から200℃の温度域での析出速度は
極めて小さく、高温域での徐冷却が必要という
ことになる。いずれにしてもかかる製造法では
焼鈍温度、焼鈍後の冷却速度のコントロールが
不可欠であることに変わりはない。(3)について
は既に述べた通りである。以上述べてきた考えに基づいて、熱延巻取温度
に依らずにコイル内材質の均質性が良好であり、
焼鈍温度及び焼鈍後の冷却速度に依らずに焼付硬
化性が付与でき、さらに二次加工脆性の問題がな
く、溶融亜鉛めつき鋼板製造法時にめつき密着性
が良好な超深絞り用鋼板製造方法の考え方は以下
のようなものになつた。即ち、ＮはAINではなくTiによつてTiNとして
仕上げ熱延前に析出させ、以下の熱延巻取時、冷
延、再結晶焼鈍時には再固溶、析出といつた反応
を起こさない、言いかえれば熱延以降の工程はＮ
に関しては材質に無関係ということである。また
Ｃは主としてNbによつてNbCとして析出させる
が、焼付硬化性を熱延巻取温度、焼鈍温度、焼鈍
後の冷却速度に関係なく安定して付与するため
に、Ｃ量の含有範囲を、Nb添加量のみとの関係
できまるExcessC量（Ｃ（％）−１２／９３Nb（％）の
値で定義する）が、10〜50ppmの範囲に制御した
鋼板が本発明の基本的な考え方である。本発明鋼が従来のNb添加鋼と比較して著しく
優れた材質特性を有するのは、TiをNbと共に複
合添加したことに依存するものである。即ち鋼中
のＮはTiによつてTiNとして既に熱延加熱炉中で
析出している。公知の如くTiNは窒化物として極
めて高温から安定析出物となるために、熱延、冷
延、再結晶焼鈍の各製造工程において何ら変化す
るものではなく、従つて、かかる製造工程の影響
によつて材質が何ら影響を受けるものではない。
Nbを単独に添加した鋼板では、熱延巻取温度に
よつてAINとしてのＮの析出量が著しく変動する
ために通常巻取では、巻取段階において存在する
固溶Ｎ量が多く、再結晶焼鈍時に微細析出する
AINあるいはNbNの影響で降伏強度が高く、伸
び、ｒ値が低く目標とする材質特性が得られな
い。また高温巻取の場合においても、コイル長手
方向前後端部では巻取時の冷却速度が速いため、
通常巻取相当の材質となることは既に述べた通り
である。本発明は鋼中のＮをTiNとして熱延加熱炉中で
すでに析出させており、TiNは極めて高温から安
定な析出物であるから上記の如く、熱延巻取温度
によつて析出量が変化するものではない。また、
TiNは周知の如くサイズの大きい析出物であるか
ら再結晶焼鈍時に結晶粒の成長を妨げ材質を劣化
させるものではない。以上の如く、本発明鋼はTiの添加により、熱
延巻取温度に関わらず、通常巻取でも良好な材質
を得る製造法となる。また、高温巻取を行なつた
場合においても、コイル内材質は極めて均質なも
のとなり、製造歩留の低下は全く考慮する必要が
ない。次に、焼鈍温度及び冷却速度の影響であるが、
本発明鋼はＮをTiNとして固定しており、焼鈍温
度に比較して極めて高温の温度域においても安定
な析出物であるために、高温焼鈍でも再固溶する
ことはなく固溶Ｎ量にはほとんど全く影響しな
い。従つて焼付硬化性は固溶Ｃ量のみに依つて決
定されるものである。冷却速度に関してもかかる
理由から考慮する必要はなく任意の冷却速度でよ
い。また、室温まで急速冷却したとしても何ら常
温時効性に対して影響を及ぼすものではない。仮りに高温焼鈍を行なつて熱延巻取時に析出し
ていたNbCの一部が再固溶したとして、固溶Ｃ量
がわずかに変動したとしても急速冷速をすること
により、Ｃをはじめとする元素の粒界への偏析が
極端に低くなるため、論文等に記載の如く、粒界
が高純であることの効果により時効処理後、降伏
点伸びが出現することはない。 Tiを複合添加した本発明鋼により、以上のよ
うに焼鈍温度、冷却速度に依存せずに純粋に鋼中
のＣ、Nb量のみを制御することにより、実質上
非時効である超深絞り用焼付硬化性鋼板が製造で
きるものである。更に本発明鋼では、Tiの複合添加により前述
の如く、焼鈍後の冷却速度に依存せずに製造可能
なため、Si，Mn等を添加して40Kg／mm^２を超え
る引張強度の高い高強度溶融亜鉛めつき鋼板製造
時に、めつき密着性の極めて良好な、従来にない
優れた鋼板を製造できる。溶融亜鉛めつき鋼板のめつき密着性は、既に述
べたように、焼鈍冷却中に鋼板表面に濃化する主
としてSi等の合金元素の影響により劣化するた
め、Si含有量は約0.3％以下に制限されていた。
このため特に引張強度40Kg／mm^２以上の性能の良
好な超深絞り用亜鉛めつき鋼板は、これまで実質
上製造されていないと言える。この欠点も、本発
明鋼ではTiの複合添加により材質は冷却速度に
ほとんど全く依存せず、急速冷却も可能となり、
表面に濃化するSi等の合金元素を著しく低減でき
るため、かかるSi含有量の高い超深絞り用焼付硬
化性高強度鋼板の製造が可能となるものである。以上述べてきた如く、本発明鋼はTiの複合添
加により、従来の鋼板にない種々の材質特性を兼
ね備えた全く新しい性格の鋼板であり、極めて有
利なものである。次に成分範囲について述べる。前述の如くTi
添加量はＮとの関係で決まる。Tiを添加するの
は、ＮをAINあるいはNbNとして析出させると、
その析出量が巻取温度によつて大きく左右された
り、焼鈍温度、焼鈍後の冷却速度によつて、最終
的に鋼板中の固溶Ｎ量が影響されるために、コイ
ル長手方向端部の材質が劣化したり、コイル内材
質のバラツキが大きくなつたり、さらには材質の
巻取温度、焼鈍温度、冷却速度に依存して大きく
変化するのを防ぐためである。調査してみると、このようなＮの効果は、含有
量20ppm以下では大きくないため、Ti添加量の
下限はTiで固定できないＮ量が20ppm以下とな
るように決まる。即ち48／14〔Ｎ（％）−0.002％
（20ppm）〕＜Ti（％）である。この場合に、Ｎは
〔Ti、Al〕Ｎとして比較的高温でも安定な析出物
になつているために、実質上全Ｎ量をTiNとして
析出させたのと同様な効果を有することが実験の
結果確められている。またTi添加量の上限は、
TiをＮとの当量（即ち48／14N（％））以上に添
加すると硫化物を形成したり、炭化物となつて延
性及び二次加工性を劣化させたり、焼付硬化性の
制御を困難にする、さらにはＰ添加時にリン化物
を形成して材質を悪くするのでTiはＮとの当量
以上は添加しないことが望ましい。以上よりTi
添加量は48／14〔Ｎ（％）−0.002％〕＜Ti（％）≦
48／14N（％）となる。一方NbはＣとの関係で決まり、ExcessC量
（Ｃ（％）−12／93Nb（％）が10〜50（ppm）の
範囲内で、かつNb量が0.02％以下であることが
望ましい。従つてNb添加量はNb（％）≦0.02％か
つ93／12〔Ｃ（％）−0.005％（50ppm）〕≦Nb
（％）≦93／12〔Ｃ（％）−0.001％（10ppm）〕と
なる。 Nb量の下限は、実質上非時効となる限界から
決まる。Nbに対する当量を超えるＣ量が50ppm
以上になると鋼板内の固溶Ｃが多くなり過ぎ、焼
付硬化性は大きくなるが、時効後、降伏点伸びが
現れ、外板用素材として適さない。Nb量の上限
は固溶Ｃが適当量残存するように決定される。
ExcessC量が10ppm以下になると、最終的な固
溶Ｃ量が少くなり過ぎ、急速冷却の場合でも、焼
付硬化性を十分付与できない。ExcessC量が10〜
50ppmの範囲の場合に、焼鈍温度冷却速度に関
わらず安定して焼付硬化性を付与でき、かつ実質
上非時効となる（第５図に示す）。更にNb添加量が0.02％を超えると、微細NbC
析出量が増加するために、再結晶温度が著しく上
昇したり、NbCによる析出強化要素が大きくな
り、YPの上昇、EI、ｒ値の低下の原因となるの
で本発明鋼の目標材質からはずれるものである。第１図、第２図は本願発明鋼のTi、Nb含有量
範囲をＮ、Ｃ量との関係で図示したものである。
第３図、第４図はTi、Nb量と材質の関係から本
発明の範囲を示したものである。第３図は、Nb量を一定量（0.008％）に固定
し、Ti量を変化させた場合の材質特性値であ
る。他の合金成分はＣ：0.0035、Si：0.01、Mn：
0.25、Ｐ：0.020、Ｓ：0.010、sol.Al：0.040、
Ｎ：0.0040（各wt％）および残部は実質的にFe
からなる試料について、ａ：コイル長手方向中心
部、ｂ：コイル長手方向前後端部を示した。熱延
巻取温度は700℃の高温巻取、焼鈍条件は800℃×
30秒、冷却速度は、室温までの冷却速度が100
℃／secである。また、焼鈍後の調質圧延率は0.8％である。Ti
量がＮをAINとして析出させるかわりにTiNとし
て固定するために不十分な量の場合、即ち48／14
〔Ｎ（％）−0.002％〕＞Tiの場合には、コイル前後
端部の材質劣化（YPの上昇、El、ｒ値の低下、
BH量（第７図参照）が過大）が大きい。コイル
前後端部は巻取時の冷却速度が速いため、通常巻
取に相当する材質を表わすことになり、かかる
Ti量の領域では低温巻取時に良好な材質を確保
できないことになる。逆にTiをＮに対する当量
（48／14N（％））以上添加した場合には、コイル
前後端部、通常巻取時の材質劣化は小さいが、
TiNを形成するために必要な量を超えるTiが含有
されているために、硫化物、炭化物を形成して延
性、ｒ値を劣化させる結果となつている。また、この領域ではTiがTiCとして析出物を形
成するために、固溶Ｃ量が少くなり、適当な焼付
硬化性が得られない。上記の範囲にTi量をコン
トロールすれば、固溶Ｃ量はExcessC量で決ま
り、焼鈍温度、冷却速度に依らずに焼付硬化性を
付与できる。第４図は、Ti量を、Ｎを固定するに足る量だ
け一定量（0.013％）添加し、Nb量を変化させた
場合の材質特性値である。他の合金成分はＣ：
0.0030、Si：0.01、Mn：0.25、Ｐ：0.020、Ｓ：
0.010、sol.Al：0.040、Ｎ：0.0040（各wt％）お
よび残部は実質的にFeよりなる試料について、
ａ：コイル長手方向中心部、ｂ：コイル長手方向
前後端部の材質を示した。他の製造条件は第３図
の場合と同様である。 Nb量を従来のNb添加鋼で使用されている量よ
りも著しく低い量、93／12〔Ｃ（％）−0.005％〕
≦Nb（％）≦93／12〔Ｃ（％）−0.001％〕かつNb
≦0.02％を満たすように添加すれば、適当な焼付
硬化性を付与できることを示している。ここに適当な焼付硬化性とはBH＝３〜５Kg／
mm^２である。Nb量が上記範囲を超えて含有され
ると、コイル前後端部あるいは通常巻取時に微細
NbCの影響でNbCによる析出強化要素が大きくな
り、YPの上昇、El、ｒ値の劣化等材質が劣化
し、また再結晶温度も著しく上昇する。さらに焼
付硬化性も十分付与できなくなり、本願発明の主
旨をはずれることになる。第５図はExcessC量（Ｃ（％）−12／93Nb
（％））と焼付硬化性の関係を示したものである。
焼鈍後の冷却速度は室温まで100℃／sec一定とし
た。ExcessC量のうちどれだけの割合が最終的に
固溶Ｃ量として残るかは別としても、固溶Ｃ量は
ExcessC量と比例して残存する。従つて、固溶Ｃ
量と対応する焼付硬化性はExcessC量と対応する
ことになる。第６図は連続焼鈍における冷却サイクルを示し
たものであり、は冷延鋼板、溶融亜鉛めつき鋼
板（合金化処理なし）用サイクル、は合金化溶
融亜鉛めつき鋼板用サイクルである。焼鈍後の冷却速度は、Tiを複合添加した本発
明鋼では、既に述べてきたように室温まで任意で
よいが、二次加工脆性、溶融亜鉛めつき性の観点
から、室温までの急速冷却（270℃／sec）が望ま
しい。次にTi、Nb以外の合金成分範囲は、Ｃ：0.007
％以下、Si：0.8％以下、Mn：1.0％以下、Ｐ：
0.15％以下、Al：0.01〜0.1％、Ｎ：0.01％以下及
び他の不可避的不純物、残部実質的にFeから成
るものである。Ｃ量が多いと必然的にＣを固定するためのNb
量がそれだけ多く必要となり、製造コストが高く
なり、またNbCの生成量が増えるため、析出強化
要素が大きくなり、結晶粒の成長が阻害され、ｒ
値の低下、降伏強度の上昇、伸びの低下を導く。
このため超深絞り性鋼板の製造という観点からＣ
は0.007％以下とする。 Siは溶融亜鉛めつき鋼板を製造する場合、めつ
き層皮膜の密着性を低下させる傾向を有するが、
本発明では室温までの急速冷却により表面濃化の
影響を抑制できるため、0.8％以下とする。本発明鋼はTi添加量が低く、固溶Ｃが残存す
る鋼板であり、更に急速冷却により粒界偏析量を
著しく低減できるために、二次加工割れを起こし
難いが、Ｐがさらに多量に含まれると、粒界偏析
量が多くなり、粒界を脆化させ、二次加工割れを
助長するため、Ｐの上限を0.15％とする。 AlはTi、Nb添加前の溶鋼脱酸剤として加える
が、少量すぎる場合には、Alによる脱酸が十分
に行なわれず、Ti、Nbが脱酸剤として働くた
め、Ti、Nbの歩留低下が著しくなり、焼付硬化
性の制御も困難になる。逆に多量に加えると
Al₂O₃介在物が増加して好ましくない。以上の理
由によりAlは0.01〜0.1％とする。ＮはTiNとしてTiに固定されるが、Ｎ含有量が
多いと必要なTi量が増加し好ましくない。この
ためＮは0.01％以下とする。 Ti、Nbを複合添加した鋼板としては、特公昭
54−12883号公報、特開昭54−131536号公報、特
開昭56−166331号公報、特開昭57−35673号公報
等に開示されているが、これらはいずれもTi、
Nbを複合添加することを発明の基本思想とする
ものではなく、Ti、NbあるいはZr、Ｖ、Cr等を
単にＣ、Ｎを析出せしめる添加元素として任意に
選択されるものである。該特許においては、TiはＮを析出せしめるに
止まらずＣをTiCとして析出せしめ、Ｃ、Ｎによ
る時効性を抑制するため、いずれもTiおよびNb
の総添加量とＣ、Ｎ量との関係で目標材質を得る
ものである。また、該特許は非時効深絞り鋼板を
発明の基本思想とするものであり、鋼中に固溶Ｃ
を適当量残存せしめ、焼付硬化性を付与するもの
ではない。本発明鋼は既に述べた通り、極低炭素鋼に微量
のTi、Nbを添加することを必須条件とし、極め
て高い深絞り性と焼付硬化性を付与することを根
本原理とする。Tiの添加は材質に悪影響を及ぼ
すＮを無害化するためであり、上記発明の如く、
TiCをも生成せしめるものではない。更に、Nb
の添加は固溶Ｃ量を常温時効性に対して有害とな
らない範囲内に低減せしめ、固溶Ｃの残存により
焼付硬化性を付与する目的にある。従つて本発明鋼は、Ti、Nbを複合添加するこ
とを基本思想とするが、Tiの添加はＮの無害
化、Nbの添加はＣによる常温時効性の無害化と
焼付硬化性を付与することにあり、いずれもＮ
量、Ｃ量に対して当量以下の微量添加である。ま
た、極低炭素鋼を基本成分とする本発明鋼は、
Ac₃点以上の焼鈍を行なうとランダム方位結晶粒
を生成せしめ、ｒ値の劣化を招くため、焼鈍温度
は再結晶温度以上Ac₃点以下となる。さらに、焼
鈍後の冷却速度は任意でよく、生産性、めつき特
性、鋼板の二次加工脆性の観点からは70℃／sec
以上が望ましく過時効処理も必要としない。以上の理由から本発明鋼は、根本原理及び得ら
れる鋼板の材質とも該特許とは本質的に異なる。本発明鋼は、Ti、Nb含有量を従来の鋼板に比
較して著しく低減せしめているため、再結晶温度
は熱間圧延条件に関わらず低い。熱間圧延時に高
温巻取を行なうことにより、析出物の凝集が促進
され、再結晶温度は更に低下する。従つて本発明
鋼は低温焼鈍でも高いｒ値が得られ、ブリキの如
き極薄鋼板製造に対しても超加工用鋼板を提供す
るものである。本発明においては、TiN，NbCを粗大析出させ
ることにより、発明の範囲内で熱延条件に左右さ
れずに高い材質を得るものである。TiNの析出開
始温度、NbCの析出開始温度は、それぞれ鋼中の
Ｎ，Ti量，Ｃ，Nb量に依存し、かかる量が少な
い場合は、析出開始が低温となつて析出物が粗大
化しにくい。かかる観点から、本発明の主旨に沿つて実施例
に示した範囲から、TiとNbの総和の下限を0.014
％、上限は同様に実施例に示した範囲に基づき
0.027％とする。これはTiとNbの総和が0.027％を
超えると、逆に炭化物（NbC）が微細になつて降
伏強度が上昇したり、再結晶温度が上昇したりす
る弊害が生じるためである。又、本発明での熱延仕上げ温度は、Ar₃未満で
は圧延中にフエライト相が混入して深絞り性を劣
化させるのでAr₃以上とする。次に巻取り温度について述べると、本発明にお
いては、炭化物、窒化物を粗大凝集させることに
より、深絞り性を高めている。かかる効果は、巻
取り温度が高いほど顕著となる。かかる観点から
実施例に示す範囲に基づき、巻取り温度を600℃
以上と限定するものである。 The present invention relates to a method for manufacturing a bake-hardenable steel plate for ultra-deep drawing. In recent years, demand for high-strength steel sheets has been increasing in the automobile industry in pursuit of improved fuel efficiency and safety through lighter vehicle bodies. On the other hand, as the salability of automobiles is increasingly influenced by the styling of the car body, more emphasis has been placed on the press formability of steel sheets than ever before. Against this background, there is an increasing demand for a steel sheet that exhibits good formability during press forming and has the property of increasing yield strength and tensile strength after painting and baking, that is, has bake hardenability. The present invention relates to a method for manufacturing a bake-hardenable steel sheet for ultra-deep drawing that satisfies such requirements. Regarding the manufacturing method of ultra-deep drawing steel sheets that do not have bake hardenability, Ti killed steel sheets (Japanese Patent Publication No. 18066/1986) and Nb killed steel sheets (Japanese Patent Publication No. 18066/1986) are available.
54-1245) are known. However, in these steel sheets, since the C and N in the steel sheets are completely fixed as precipitates such as Ti or Nb, strain aging does not occur during paint baking after press forming, and therefore the bake hardenability is reduced. It is something that we do not have. Next, as a manufacturing method for ultra-deep drawing steel sheets with bake hardenability, C in Nb-added steel,
Nb (at
%)/{solid solute C (at%) + solid solute N (at%)} within a certain range, the solid solute C in the steel plate,
A method characterized by controlling the amount of solid solute N and further controlling the cooling rate after annealing (Special Publication No. 55
-141526, 55-141555). However, upon actual investigation and consideration, this manufacturing method has the following drawbacks. First, there are restrictions on hot rolling coiling temperature, annealing temperature, and cooling rate after annealing. Nb-added steel requires hot rolling and high temperature coiling (coiling temperature ≧700℃). At normal coiling temperatures, the complete recrystallization temperature becomes very high, and in the temperature range possible with continuous annealing furnaces (usually about 850°C or less), unrecrystallized areas may remain, and depending on the amount of Nb. There are large variations in the material used. It has been variously reported that when high-temperature coiling is performed, a steel plate with a high r value can be obtained at an annealing temperature of approximately 800 to 850°C, except for the longitudinal ends of the hot-rolled coil. be. However, high-temperature winding not only makes the scale thicker and drastically reduces pickling efficiency, but also because the cooling rate at the end of the coil is fast, the material remains at the same temperature as normal winding. The decrease in yield is particularly large in Nb-killed steel because the material cannot be obtained. The second problem is the annealing temperature and the cooling rate after annealing. As in the above-mentioned known technology, high temperature (approximately 900℃)
Nb that was precipitated during hot rolling when annealed with
Since (C, N) or AIN is redissolved, solute N and C increase, resulting in increased aging properties. Also,
The cooling rate after annealing was set to 400℃ as stated in the patent.
up to 50°C/sec or less, or up to 400°C up to 50°C
℃/sec or more and 10℃/sec from 400℃ to 200℃
As described below, unless slow cooling is performed, slow aging properties cannot be achieved due to the effects of solid solute N and solid solute C.
In the case of Nb-added steel, a particular problem is AIN, which remelts at high temperatures, and it is this N aging that poses a problem with room temperature aging. A third drawback is the galvanizing properties. Generally, when hot-dip galvanizing is performed, plating adhesion is said to deteriorate if the Si content in the steel exceeds approximately 0.3%. This is thought to be because a thin oxide film is formed due to Si concentrating on the surface of the steel sheet, and the reaction between the base iron and molten zinc does not proceed well. For this reason, materials subjected to hot-dip galvanizing have been manufactured with the Si content in the steel limited to about 0.3% or less in order to suppress the amount of Si concentrated on the surface. However, the tensile strength
When manufacturing steel sheets with a weight of 40Kg/mm2 or ^more , the P content is limited due to the problem of secondary processing cracking, and other alloying elements must be added to increase the strength. Si in terms of
and Mn cannot be added, so it is practically 40Kg/mm
It is very difficult to find a galvanized ultra-deep drawing bake-hardenable steel sheet with good performance and a tensile strength of more than ² .
The above-mentioned surface concentration phenomenon of Si is a general phenomenon during slow cooling, and has been an unsolved problem until now. As a result of a detailed study of these drawbacks of the conventional Nb-added steel, the present inventors came to the conclusion that the essential causes are the following three points. (1) Significantly affected by temperature history during hot rolling and winding
Precipitation of AIN. (2) The amount of solid solute C and solid solute N changes depending on the annealing temperature and the cooling rate after annealing. (3)
There are three points: the concentration of alloying elements at grain boundaries or on the surface, which depends on the cooling rate after annealing. Regarding (1), although Nb actually forms carbides in Nb-added steel, the tendency to form nitrides is relatively weak;
Nitrogen is precipitated as AIN. AIN is difficult to form at normal winding temperatures, so the winding temperature is set at 700℃.
Otherwise, it will not precipitate in the hot-rolled sheet, but will precipitate as fine precipitates during continuous annealing after cold rolling, suppressing the growth of crystal grains, increasing yield strength, decreasing elongation, etc., resulting in material deterioration. cause. Therefore, even when high-temperature winding is performed, AIN precipitates at the longitudinal center of the hot-rolled coil, but AIN precipitates at the longitudinal center of the hot-rolled coil during winding.
As AIN does not precipitate, it is the same as normal winding.
This results in an increase in yield strength, elongation, and an extreme decrease in r value. In addition, due to slight variations in hot-rolling winding temperature and differences in the degree of AIN precipitation at the center and ends of the coil, material deterioration at the front and rear ends of the hot-rolled coil and material inside the coil may deteriorate. This causes variations in the Regarding (2), when high-temperature annealing is performed, precipitates such as Nb (C, N) and AIN that were precipitated during hot-rolling and coiling are re-dissolved to the solubility product at that temperature, and solid solution C is formed. , solute N increases. especially
Since solid solution N resulting from AIN dissolution has a higher room temperature aging property than solid solution C, this goes against the target material of slow aging property. In addition, solid solution C and solid solution N present during annealing become Nb in the subsequent cooling process.
It precipitates as (C,N) or AIN, but the precipitation rate of AIN peaks at about 815°C and decreases as the temperature decreases, so slow cooling in a high temperature range is required. In this patent, cooling from 400℃ to 200℃ by slow cooling to 400℃ or rapid cooling to 400℃
However, it is necessary to slowly cool down to
Substantially, the precipitation rate in the temperature range of 400°C to 200°C is extremely low, which means that gradual cooling is required in the high temperature range. In any case, in this manufacturing method, it is still essential to control the annealing temperature and the cooling rate after annealing. Regarding (3), it has already been stated. Based on the ideas described above, the homogeneity of the material inside the coil is good regardless of the hot rolling coiling temperature,
Production of ultra-deep drawing steel sheets that can impart bake hardenability regardless of the annealing temperature and post-annealing cooling rate, are free from the problem of secondary processing brittleness, and have good plating adhesion when manufacturing hot-dip galvanized steel sheets. The idea behind the method was as follows. In other words, N is precipitated as TiN by Ti rather than AIN before finishing hot rolling, and does not cause reactions such as redissolution or precipitation during hot rolling, cold rolling, and recrystallization annealing. In other words, the process after hot rolling is N.
This means that it has nothing to do with the material. In addition, C is mainly precipitated as NbC by Nb, but in order to stably impart bake hardenability regardless of hot rolling coiling temperature, annealing temperature, and cooling rate after annealing, the content range of C amount is set as follows: The basic idea of the present invention is to provide a steel sheet in which the amount of ExcessC (defined as the value of C (%) - 12/93Nb (%)), which is determined by the relationship only with the amount of Nb added, is controlled within the range of 10 to 50 ppm. . The fact that the steel of the present invention has significantly superior material properties compared to conventional Nb-added steels is due to the composite addition of Ti and Nb. That is, N in the steel has already been precipitated as TiN by Ti in the hot rolling furnace. As is well known, since TiN becomes a stable precipitate as a nitride at extremely high temperatures, it does not change in any way during the manufacturing processes of hot rolling, cold rolling, and recrystallization annealing. The material is not affected in any way.
In steel sheets to which Nb has been added alone, the amount of N precipitated as AIN varies markedly depending on the hot-rolling and coiling temperature. Fine precipitation occurs during annealing
Due to the influence of AIN or NbN, the yield strength is high and the elongation and r value are low, making it impossible to obtain the desired material properties. In addition, even in the case of high-temperature winding, the cooling rate at the front and rear ends of the coil during winding is fast.
As already mentioned, the material is equivalent to that used for normal winding. In the present invention, the N in the steel is already precipitated in the hot rolling heating furnace as TiN, and since TiN is a stable precipitate even at extremely high temperatures, the amount of precipitation varies depending on the hot rolling coiling temperature as described above. It's not something that changes. Also,
As is well known, TiN is a large-sized precipitate, so it does not hinder the growth of crystal grains during recrystallization annealing and does not deteriorate the material quality. As described above, by adding Ti, the steel of the present invention can be produced by a method of producing good material even when it is normally coiled, regardless of the hot rolling coiling temperature. Further, even when high-temperature winding is performed, the material inside the coil becomes extremely homogeneous, and there is no need to consider a decrease in manufacturing yield at all. Next, as for the effects of annealing temperature and cooling rate,
In the steel of the present invention, N is fixed as TiN, which is a stable precipitate even in an extremely high temperature range compared to the annealing temperature, so it does not re-dissolve even during high-temperature annealing, and the amount of solid solute N decreases. has almost no effect. Therefore, bake hardenability is determined only by the amount of solid solute C. For this reason, there is no need to consider the cooling rate, and any cooling rate may be used. In addition, even if it is rapidly cooled to room temperature, it does not affect the room temperature aging property in any way. Even if some of the NbC that had precipitated during hot-rolling and coiling was re-dissolved during high-temperature annealing, even if the amount of solute C slightly fluctuated, rapid cooling would reduce the amount of NbC and other components. Since the segregation of the elements to the grain boundaries is extremely low, as described in papers, etc., yield point elongation does not appear after aging treatment due to the effect of the high purity of the grain boundaries. By controlling only the C and Nb contents in the steel without depending on the annealing temperature and cooling rate, the steel of the present invention with composite addition of Ti can be used for ultra-deep drawing, which is virtually non-aging. Bake-hardenable steel sheets can be produced. Furthermore, the steel of the present invention can be manufactured without depending on the cooling rate after annealing as mentioned above due to the composite addition of Ti, so by adding Si, Mn, etc., high tensile strength exceeding 40 Kg/ ^mm2 can be achieved. When manufacturing hot-dip galvanized steel sheets, it is possible to manufacture unprecedentedly superior steel sheets with extremely good plating adhesion. As mentioned above, the plating adhesion of hot-dip galvanized steel sheets deteriorates mainly due to the influence of alloying elements such as Si, which concentrate on the steel sheet surface during annealing and cooling. It was restricted.
For this reason, it can be said that a galvanized steel sheet for ultra-deep drawing with good performance, particularly a tensile strength of 40 Kg/mm ² or more, has not been substantially manufactured to date. This drawback can also be overcome by the composite addition of Ti in the steel of the present invention, which makes the material almost completely independent of the cooling rate and allows for rapid cooling.
Since alloying elements such as Si that concentrate on the surface can be significantly reduced, it becomes possible to produce a bake-hardenable high-strength steel sheet for ultra-deep drawing with a high Si content. As described above, the steel of the present invention is a completely new steel sheet that has various material properties not found in conventional steel sheets due to the composite addition of Ti, and is extremely advantageous. Next, we will discuss the component range. As mentioned above, Ti
The amount added is determined by the relationship with N. Ti is added when N is precipitated as AIN or NbN.
The amount of N precipitated is greatly affected by the coiling temperature, and the amount of solid solute N in the steel sheet is ultimately affected by the annealing temperature and cooling rate after annealing. This is to prevent the material from deteriorating, the variation in the material inside the coil from increasing, and further from changing greatly depending on the winding temperature, annealing temperature, and cooling rate of the material. Investigation revealed that such an effect of N is not large when the content is less than 20 ppm, so the lower limit of the amount of Ti added is determined so that the amount of N that cannot be fixed by Ti is 20 ppm or less. That is, 48/14 [N (%) - 0.002%
(20ppm)]<Ti (%). In this case, because N is precipitated as [Ti, Al]N, which is stable even at relatively high temperatures, experiments have shown that it has the same effect as precipitating substantially all the N as TiN. The results have been confirmed. In addition, the upper limit of the amount of Ti added is
If Ti is added in an amount exceeding the equivalent amount to N (i.e. 48/14N (%)), it will form sulfides or carbides, which will deteriorate ductility and secondary workability, and make it difficult to control bake hardenability. Moreover, when P is added, phosphides are formed and the quality of the material deteriorates, so it is desirable not to add Ti in an amount exceeding the equivalent amount of N. From the above, Ti
Addition amount is 48/14 [N (%) - 0.002%] <Ti (%) ≦
It becomes 48/14N (%). On the other hand, Nb is determined by the relationship with C, and it is desirable that the amount of ExcessC (C (%) - 12/93Nb (%)) is within the range of 10 to 50 (ppm) and the amount of Nb is 0.02% or less. Therefore, the amount of Nb added is Nb (%) ≦ 0.02% and 93/12 [C (%) - 0.005% (50 ppm)] ≦ Nb
(%)≦93/12 [C (%) - 0.001% (10 ppm)]. The lower limit of the amount of Nb is determined from the limit at which it becomes virtually non-aging. The amount of C that exceeds the equivalent amount to Nb is 50ppm
If the temperature exceeds this value, the solid solution C in the steel plate becomes too large, and the bake hardenability increases, but yield point elongation appears after aging, making it unsuitable as a material for outer panels. The upper limit of the amount of Nb is determined so that an appropriate amount of solid solution C remains.
When the amount of Excess C is less than 10 ppm, the final amount of solid solute C becomes too small, and even in the case of rapid cooling, sufficient bake hardenability cannot be imparted. ExcessC amount is 10~
In the case of a range of 50 ppm, bake hardenability can be stably imparted regardless of the annealing temperature and cooling rate, and substantially no aging occurs (as shown in FIG. 5). Furthermore, when the amount of Nb added exceeds 0.02%, fine NbC
As the amount of precipitation increases, the recrystallization temperature increases significantly and the precipitation strengthening factor due to NbC increases, causing an increase in YP and a decrease in EI and r value, which deviates from the target material of the steel of the present invention. It is. FIGS. 1 and 2 illustrate the Ti and Nb content ranges of the steel according to the present invention in relation to the N and C contents.
FIGS. 3 and 4 show the scope of the present invention from the relationship between the amounts of Ti and Nb and the materials. Figure 3 shows the material property values when the Nb content is fixed at a constant level (0.008%) and the Ti content is varied. Other alloy components are C: 0.0035, Si: 0.01, Mn:
0.25, P: 0.020, S: 0.010, sol.Al: 0.040,
N: 0.0040 (each wt%) and the remainder is substantially Fe
For a sample consisting of the following, a: the center portion in the longitudinal direction of the coil, and b: the front and rear ends of the coil in the longitudinal direction are shown. Hot rolling coiling temperature is 700℃, annealing condition is 800℃×
30 seconds, cooling rate is 100 seconds to room temperature
°C/sec. In addition, the skin pass rolling ratio after annealing is 0.8%. Ti
If the amount is insufficient to fix N as TiN instead of precipitating as AIN, i.e. 48/14
[N (%) - 0.002%] > Ti, material deterioration at the front and rear ends of the coil (increase in YP, decrease in El and r values,
BH amount (excessive amount (see Figure 7)) is large. The front and rear ends of the coil cool down quickly during winding, so they represent the material that is normally used for winding.
In the range of Ti content, good material quality cannot be ensured during low temperature winding. Conversely, when Ti is added in an amount equal to or more than N (48/14N (%)), material deterioration at the front and rear ends of the coil and during normal winding is small, but
Since the content of Ti exceeds the amount necessary to form TiN, sulfides and carbides are formed, resulting in deterioration of ductility and r value. Further, in this region, since Ti forms precipitates as TiC, the amount of solid solution C decreases, making it impossible to obtain appropriate bake hardenability. If the amount of Ti is controlled within the above range, the amount of solute C is determined by the amount of Excess C, and bake hardenability can be imparted regardless of the annealing temperature and cooling rate. FIG. 4 shows the material property values when a constant amount (0.013%) of Ti is added, sufficient to fix N, and the amount of Nb is varied. Other alloy components are C:
0.0030, Si: 0.01, Mn: 0.25, P: 0.020, S:
0.010, sol.Al: 0.040, N: 0.0040 (each wt%), and the remainder substantially consists of Fe,
a: The material of the center part in the longitudinal direction of the coil, and b: The material of the front and rear ends of the coil in the longitudinal direction. Other manufacturing conditions are the same as in the case of FIG. The amount of Nb is significantly lower than that used in conventional Nb-added steel, 93/12 [C (%) - 0.005%]
≦Nb (%) ≦93/12 [C (%) - 0.001%] and Nb
This shows that if added in an amount of ≦0.02%, appropriate bake hardenability can be imparted. Appropriate bake hardenability here is BH = 3~5Kg/
^mm2 . If the Nb content exceeds the above range, fine particles may appear at the front and rear ends of the coil or during normal winding.
Due to the influence of NbC, the precipitation strengthening factor due to NbC increases, resulting in material deterioration such as an increase in YP and deterioration of El and r values, and the recrystallization temperature also increases significantly. Furthermore, it becomes impossible to impart sufficient bake hardenability, which deviates from the gist of the present invention. Figure 5 shows ExcessC amount (C (%) - 12/93Nb
(%)) and bake hardenability.
The cooling rate after annealing was kept constant at 100°C/sec until room temperature. Regardless of how much of the ExcessC amount ultimately remains as solid solute C amount, the solid solute C amount is
It remains in proportion to the amount of ExcessC. Therefore, solid solution C
The bake hardenability corresponding to the amount corresponds to the amount of ExcessC. FIG. 6 shows cooling cycles in continuous annealing, where 1 is a cycle for cold-rolled steel sheets and hot-dip galvanized steel sheets (without alloying treatment), and 2 is a cycle for alloyed hot-dip galvanized steel sheets. As mentioned above, the cooling rate after annealing can be any rate up to room temperature for the steel of the present invention with composite addition of Ti, but from the viewpoint of secondary processing brittleness and hot-dip galvanizing properties, rapid cooling to room temperature ( 270℃/sec) is desirable. Next, the range of alloy components other than Ti and Nb is C: 0.007
% or less, Si: 0.8% or less, Mn: 1.0% or less, P:
0.15% or less, Al: 0.01 to 0.1%, N: 0.01% or less and other unavoidable impurities, and the remainder substantially consists of Fe. When the amount of C is large, Nb is inevitably added to fix C.
Since a larger amount is required, the production cost becomes higher, and the amount of NbC produced increases, the precipitation strengthening factor becomes larger, inhibiting the growth of crystal grains, and r
leads to a decrease in value, an increase in yield strength, and a decrease in elongation.
Therefore, from the perspective of producing ultra-deep drawable steel sheets, C
shall be 0.007% or less. When manufacturing hot-dip galvanized steel sheets, Si tends to reduce the adhesion of the plating layer film, but
In the present invention, the effect of surface concentration can be suppressed by rapid cooling to room temperature, so the content is set to 0.8% or less. The steel of the present invention is a steel plate with a low Ti addition and residual solid solution C, and the amount of grain boundary segregation can be significantly reduced by rapid cooling, so secondary processing cracking is less likely to occur, but it also contains a large amount of P. If this occurs, the amount of grain boundary segregation increases, which embrittles the grain boundaries and promotes secondary work cracking. Therefore, the upper limit of P is set at 0.15%. Al is added as a deoxidizer for molten steel before adding Ti and Nb, but if it is too small, deoxidation by Al will not be sufficient and Ti and Nb will act as deoxidizers, resulting in a decrease in the yield of Ti and Nb. This makes it difficult to control bake hardenability. On the other hand, if you add a large amount
This is not preferable because Al ₂ O ₃ inclusions increase. For the above reasons, Al is set to 0.01 to 0.1%. N is fixed on Ti as TiN, but if the N content is high, the required amount of Ti will increase, which is not preferable. Therefore, N should be 0.01% or less. As a steel sheet with composite addition of Ti and Nb,
These are disclosed in JP-A-54-12883, JP-A-54-131536, JP-A-56-166331, JP-A-57-35673, etc., all of which contain Ti,
The basic idea of the invention is not to add Nb in combination, but to simply select Ti, Nb, Zr, V, Cr, etc. as additional elements that precipitate C and N. In this patent, Ti not only precipitates N but also precipitates C as TiC and suppresses the aging properties due to C and N, so both Ti and Nb
The target material quality is obtained based on the relationship between the total addition amount and the amounts of C and N. In addition, the basic idea of the invention is a non-aging deep drawn steel plate, and the patent is based on a non-aging deep drawn steel plate, with solid solution carbon in the steel.
remains in an appropriate amount, and does not impart bake hardenability. As already mentioned, the basic principle of the steel of the present invention is to add a small amount of Ti and Nb to ultra-low carbon steel as an essential condition, and to impart extremely high deep drawability and bake hardenability. The purpose of adding Ti is to render harmless N, which has a negative effect on the material, and as in the above invention,
It does not generate TiC either. Furthermore, Nb
The purpose of the addition is to reduce the amount of solid solute C within a range that is not harmful to room temperature aging properties, and to impart bake hardenability due to the residual solid solute C. Therefore, the basic idea of the steel of the present invention is to add Ti and Nb in combination, but the addition of Ti makes N harmless, and the addition of Nb makes the room temperature aging property harmless and imparts bake hardenability due to C. In particular, both N
It is added in a trace amount, equal to or less than the amount of C. In addition, the steel of the present invention, which has ultra-low carbon steel as its basic component,
If annealing is performed at _three Ac points or more, randomly oriented crystal grains will be generated and the r value will deteriorate, so the annealing temperature will be higher than the recrystallization temperature and lower than Ac ₃ points. Furthermore, the cooling rate after annealing may be arbitrary, and from the viewpoint of productivity, plating characteristics, and secondary processing brittleness of the steel sheet, it is 70℃/sec.
The above is desirable and no over-aging treatment is required. For the reasons mentioned above, the basic principle and the material of the obtained steel sheet of the steel of the present invention are essentially different from those of the patent. Since the steel of the present invention has significantly reduced Ti and Nb contents compared to conventional steel sheets, the recrystallization temperature is low regardless of hot rolling conditions. By performing high-temperature winding during hot rolling, agglomeration of precipitates is promoted and the recrystallization temperature is further lowered. Therefore, the steel of the present invention can obtain a high r value even when annealed at low temperatures, and provides a steel plate for super-processing even in the production of ultra-thin steel plates such as tinplate. In the present invention, by coarsely precipitating TiN and NbC, high quality material can be obtained without being influenced by hot rolling conditions within the scope of the invention. The precipitation start temperature of TiN and the precipitation start temperature of NbC depend on the amounts of N, Ti, C, and Nb in the steel, respectively. If these amounts are small, precipitation starts at a low temperature and the precipitates are difficult to coarsen. . From this point of view, the lower limit of the sum of Ti and Nb was set at 0.014 from the range shown in the examples in accordance with the gist of the present invention.
%, the upper limit is also based on the range shown in the example.
It shall be 0.027%. This is because if the total amount of Ti and Nb exceeds 0.027%, carbides (NbC) become finer, resulting in higher yield strength and higher recrystallization temperature. Further, the hot rolling finishing temperature in the present invention is set to be Ar ₃ or more, since if it is less than Ar ₃ , a ferrite phase will be mixed during rolling and the deep drawability will deteriorate. Next, regarding the winding temperature, in the present invention, deep drawability is improved by coarsely aggregating carbides and nitrides. This effect becomes more pronounced as the winding temperature increases. From this point of view, the winding temperature was set to 600°C based on the range shown in the example.
This is limited to the above.

【表】【table】

【表】以下、実施例について述べる。実施例１第１表は本発明鋼および比較のために用いた供
試鋼の化学成分を示したものである。上記の供試鋼を熱間仕上温度910℃、巻取温度
は700℃の高温巻取および、600℃の通常巻取で板
厚40mmに熱間圧延し、0.8mmまで冷間圧延した
後、第６図に示す焼鈍サイクルを用いて連続焼
鈍ラインにて焼鈍し冷延鋼板を製造した。但し焼
鈍温度は800℃、保定時間30秒、冷却速度は室温
まで100℃／sec一定とした。その後調査圧延を１
％の圧延率で加えた。その材質試験を第２表(i)，
(ii)に示す。[Table] Examples will be described below. Example 1 Table 1 shows the chemical composition of the steel of the present invention and the test steel used for comparison. The above test steel was hot-rolled to a thickness of 40 mm with a hot finishing temperature of 910℃, a coiling temperature of 700℃, and a normal coiling of 600℃, and then cold-rolled to a thickness of 0.8mm. A cold rolled steel plate was manufactured by annealing on a continuous annealing line using the annealing cycle shown in FIG. However, the annealing temperature was 800°C, the holding time was 30 seconds, and the cooling rate was constant at 100°C/sec until room temperature. Then survey rolling 1
It was added at a rolling rate of %. The material test is shown in Table 2 (i).
Shown in (ii).

【表】【table】

【表】第２表(i)，(ii)より本願発明によるNo.１〜No.５の
鋼板は、コイル内材質が均質でバラツキが極めて
小さく、巻取温度によらず通常巻取相当でも良好
な材質が得られることが明らかである。また二次
加工割れの心配がなく化成処理性も良好である。
さらに成分変化にかかわらず、焼付硬化性が精度
よく制御できる優位性を示している。比較材No.６、No.７、No.８はTiが本発明範囲か
らはずれている。No.６、No.８はTi量が低すぎる
ためNbキルド鋼に近い材質となり、通常巻取の
場合及び高温巻取時のコイル端部の材質劣化が大
きく、またAlNの溶解による固溶Ｎの影響によ
り、時効処理後降伏点伸びが出現している。 No.７はTi量が多過ぎるためTiキルド鋼に近い
材質となり二次加工割れが起こり易く、また化成
処理性、焼付硬化性が劣る。No.９はNb量が本発
明範囲をはずれ、Nbキルド鋼に近い材質とな
り、通常巻取の場合及び高温巻取時のコイル端部
の材質劣化が極めて大きく、またNb量が多いた
め化成処理性が劣る。また比較鋼No.６〜８はTi
量、Nb量が本発明範囲をはずれることから焼付
硬化性を適当な値にコントロールすることができ
ない。本発明鋼No.１〜No.５は冷却速度を変化させた場
合にも第２表と同様、常温で実質上非時効であ
り、かつ３〜５Kg／mm^２の焼付硬化性が得られる
ことを確認した。実施例２第３表は本発明鋼および比較のために用いた供
試鋼の成分組成を示したものである。[Table] From Table 2 (i) and (ii), steel plates No. 1 to No. 5 according to the present invention have homogeneous material inside the coil with extremely small variations, and can be used even when the coil is equivalent to normal coiling regardless of the coiling temperature. It is clear that a good material can be obtained. Furthermore, there is no fear of cracking during secondary processing, and the chemical conversion treatment property is also good.
Furthermore, it shows the advantage of being able to accurately control bake hardenability regardless of changes in components. In comparative materials No. 6, No. 7, and No. 8, Ti is outside the range of the present invention. No. 6 and No. 8 have a material similar to Nb-killed steel because the Ti content is too low, and there is significant material deterioration at the end of the coil during normal winding and high-temperature winding, and solid solution N due to dissolution of AlN. Due to the influence of aging, yield point elongation appears after aging treatment. No. 7 has an excessively large amount of Ti, making it a material similar to Ti-killed steel, which is susceptible to secondary processing cracks and has poor chemical conversion treatment properties and bake hardenability. In No. 9, the amount of Nb is outside the range of the present invention, and the material is close to Nb-killed steel, and the material deterioration at the end of the coil during normal winding and high-temperature winding is extremely large, and because the amount of Nb is large, it is treated with chemical conversion treatment. inferior in sex. In addition, comparative steels No. 6 to 8 are Ti
Since the amount of Nb and the amount of Nb are out of the range of the present invention, it is not possible to control the bake hardenability to an appropriate value. Inventive steels No. 1 to No. 5 are substantially non-aging at room temperature even when the cooling rate is changed, as shown in Table 2, and bake hardenability of 3 to 5 kg/mm ² can be obtained. It was confirmed. Example 2 Table 3 shows the composition of the steel of the present invention and the sample steel used for comparison.

【表】【table】

【表】上記供試鋼は本発明範囲の鋼および比較材に合
金元素を添加して高強度化したものである。No.１
〜６までの鋼を実施例１の場合と同一の条件によ
り冷間圧延まで行ない、第６図に示すサイクル
を用いて連続焼鈍し、冷延鋼板を製造した。但し
焼鈍温度は800℃、保定時間30秒、冷却速度は10
℃／sec、50℃／sec、70℃／sec、100℃／sec
（室温まで一定）とした。供試鋼No.７は熱延仕上温度880℃、巻取温度540
℃で板厚4.0mmに熱間圧延し、0.8mmまで冷間圧延
した後、焼鈍温度720℃で箱型焼鈍に供した。 No.１〜６の材料について冷却速度100℃／sec
（室温まで一定）の場合の材質を第４表に示し、
同時にNo.７の材質結果も示した。[Table] The above test steels were made by adding alloying elements to the steels within the scope of the present invention and comparative materials to increase their strength. No.1
The steels of Examples 1 to 6 were cold-rolled under the same conditions as in Example 1, and then continuously annealed using the cycle shown in FIG. 6 to produce cold-rolled steel sheets. However, the annealing temperature is 800℃, the holding time is 30 seconds, and the cooling rate is 10
℃/sec, 50℃/sec, 70℃/sec, 100℃/sec
(constant up to room temperature). Test steel No. 7 has a hot rolling finishing temperature of 880℃ and a coiling temperature of 540℃.
After hot rolling to a plate thickness of 4.0 mm at ℃ and cold rolling to 0.8 mm, it was subjected to box annealing at an annealing temperature of 720 ℃. Cooling rate 100℃/sec for materials No. 1 to 6
(Constant up to room temperature) materials are shown in Table 4.
At the same time, the material results for No. 7 are also shown.

【表】【table】

【表】第４表により、本発明鋼は高強度化した場合に
もｒ値が極めて高く、延性も優れている。軟鋼板
の場合と同様にコイル内材質も均質で化成処理性
も良好であり、また、焼付硬化性も目標値通り得
られている。さらにＰ等合金元素の添加量が多い
にもかかわらず二次加工割れが起こり難いのは固
溶Ｃが存在して粒界を強化しているためと、室温
までの急速冷却により粒界偏析Ｐ量が著しく低減
されているためである。第４表(2)に冷却速度を変化させた場合の二次加
工脆性の発生限界を示す。試験温度は−50℃であ
り、試験片はシヤーエツジであり、円筒成形した
場合の外径を30mmφ一定とした。（試料は全て高
温巻取材のコイル長手方向中心部）[Table] According to Table 4, the steel of the present invention has an extremely high r value and excellent ductility even when it is strengthened. As in the case of the mild steel plate, the material inside the coil is homogeneous and has good chemical conversion treatment properties, and the bake hardenability has also been achieved as per the target value. Furthermore, the reason why secondary work cracking is difficult to occur despite the addition of a large amount of alloying elements such as P is due to the presence of solid solution C, which strengthens the grain boundaries, and the rapid cooling to room temperature, which causes grain boundary segregation P. This is because the amount is significantly reduced. Table 4 (2) shows the limits for the occurrence of secondary work brittleness when the cooling rate is varied. The test temperature was -50°C, the test piece was a shear edge, and the outer diameter when molded into a cylinder was constant at 30 mmφ. (All samples are at the longitudinal center of the coil of high-temperature rolled material)

【表】【table】

【表】第４表(2)の結果より、Ｐ添加を行なつて高強度
鋼板化した場合には二次加工割れの発生限界は軟
鋼板の場合と比較して劣化するが、本発明鋼は
Tiの複合添加によつてもNb添加鋼とほぼ同じレ
ベルであり、二次加工割れを発生し易いTiキル
ド鋼（No.５の材料はTiキルド鋼に近い材質であ
る。）よりも優れる。Nb添加鋼と同一レベルの二
次加工割れ特性を得るには冷却速度は任意で良い
が、箱焼鈍で製造したAl・Ｋ鋼並みの限界絞り
比を得るには70℃／sec以上の冷却速度で冷却す
るのが望ましい。本発明鋼では、Tiの複合添加
により急速冷却も可能であり、Ｐの粒界偏析量を
著しく低減できるためかかる効果が生まれてくる
ものであり、本発明鋼の優位性を示すものであ
る。比較材No.４はNb添加量が多いために通常巻取
の場合及び高温巻取時のコイル端部の材質劣化が
極めて大きく、化成処理性が劣る。また、Nb量
が多過ぎるために急速冷却の場合でも、焼付硬化
性を付与できない。No.５の材料はTi含有量が多
過ぎて二次加工割れを起こし易く、化成処理性が
劣るという欠点を有し、さらに焼付硬化性を付与
できない。 No.６の材料はTi量がＮをTiNとして固定するに
必要な量に足らないため、Nb添加鋼に近い材質
となり、通常巻取時および高温巻取時のコイル前
後端部の材質劣化が極めて大きく、焼付硬化性の
バラツキが大きいという欠点を有する。（通常巻
取材の材質は実施例１の場合と同様いずれも高温
巻取材のコイル前後端部とほとんど同じ材質を示
したので特に記述しない。）実施例３第１表、第３表に示す供試鋼の全てについて、
実施例１の場合と同一条件にて冷間圧延まで行な
つた後、第６図，で示す焼鈍サイクルを用い
て、溶融亜鉛めつき鋼板を製造した。焼鈍温度は
800℃、焼鈍時間は30秒であり、冷却速度は室温
まで一定の100℃／secの冷速で冷却した。第６図は合金化処理を行なわない場合に相当
し、は合金化処理を行なつて合金化亜鉛めつき
鋼板を製造する場合である。これらの場合、冷却
速度のコントロールは鋼板が亜鉛めつき浴に入る
までの冷却速度、めつき浴を出てから室温になる
までの冷却速度を、ともに制御した。上記の製造結果は以下の通りである。１機械試験値は第２表、第４表に示した値とほ
とんど同じであり亜鉛めつきを行なつたこと
は、本発明の主旨に何ら反するものはなかつ
た。合金化処理を行なうことは530℃程度で約
10秒保持されるが、焼付硬化性に実施例１、２
と差がないことは、大部分のＣの析出が500℃
に比較し、より高温域で起こるためと理解され
る。かかる理由により材質特性値は特に示さな
い。但し、亜鉛めつきを行なつた場合のめつき
密着性を第５表に示す。[Table] From the results in Table 4 (2), it is clear that when P is added to make a high-strength steel sheet, the limit of occurrence of secondary work cracking is lower than that of a mild steel sheet, but the present invention steel teeth
Even with the combined addition of Ti, it is at almost the same level as Nb-added steel, and is superior to Ti-killed steel (material No. 5 is similar to Ti-killed steel), which is prone to secondary processing cracks. To obtain the same level of secondary processing cracking properties as Nb-added steel, the cooling rate can be set at any desired rate, but to obtain the same critical drawing ratio as Al/K steel manufactured by box annealing, the cooling rate must be 70°C/sec or higher. It is preferable to cool it down. In the steel of the present invention, rapid cooling is also possible due to the combined addition of Ti, and the amount of grain boundary segregation of P can be significantly reduced, resulting in this effect, which shows the superiority of the steel of the present invention. Comparative material No. 4 has a large amount of Nb added, so the material deterioration at the end of the coil during normal winding and high-temperature winding is extremely large, and the chemical conversion treatment property is poor. Furthermore, since the amount of Nb is too large, bake hardenability cannot be imparted even in the case of rapid cooling. Material No. 5 has the drawbacks of having an excessively high Ti content, easily causing secondary processing cracks, poor chemical conversion treatment properties, and cannot be imparted with bake hardenability. Since the amount of Ti in material No. 6 is not sufficient to fix N as TiN, the material is similar to Nb-added steel, and material deterioration at the front and rear ends of the coil during normal winding and high-temperature winding is avoided. It has the drawback of being extremely large and having large variations in bake hardenability. (As in Example 1, the material of the normal rolled material is almost the same as the front and rear ends of the coil of the high temperature rolled material, so it will not be particularly described.) Example 3 All about test steel,
After cold rolling was performed under the same conditions as in Example 1, a hot-dip galvanized steel sheet was manufactured using the annealing cycle shown in FIG. The annealing temperature is
The temperature was 800°C, the annealing time was 30 seconds, and the cooling rate was a constant 100°C/sec to room temperature. FIG. 6 corresponds to a case where no alloying treatment is performed, and FIG. 6 corresponds to a case where an alloyed galvanized steel sheet is manufactured by performing an alloying treatment. In these cases, the cooling rate was controlled both before the steel plate entered the galvanizing bath and after leaving the galvanizing bath until it reached room temperature. The above manufacturing results are as follows. 1. The mechanical test values were almost the same as those shown in Tables 2 and 4, and the fact that galvanizing was performed did not contradict the gist of the present invention. Alloying treatment is carried out at approximately 530℃.
Although it was held for 10 seconds, the bake hardenability of Examples 1 and 2 was
The fact that there is no difference between the two and
This is understood to be because it occurs in a higher temperature range compared to . For this reason, material property values are not particularly shown. However, Table 5 shows the plating adhesion when galvanizing was performed.

【表】上記結果より、本発明鋼は亜鉛めつき特性も極
めて良好である。特に第３表の試料はSi含有量が
高いにもかかわらず密着性が良好であるのは、急
速冷却により表面濃化Si量を著しく低減できてい
るためである。第１表−No.７、第３表−No.５の試料はTi含有
量が多いために、地鉄と溶融亜鉛の合金化反応が
促進されて、過合金化が進みメツキ層中に硬くて
脆い合金層が形成されたために、密着性が劣化し
たものと考えられる。第１表−No.９、第３表−No.
４の材料も本発明鋼に比べ密着性が劣るのはNb
添加量が多いためと推定される。この観点から
も、合金元素としてTi、Nb添加量の低い本発明
鋼は優位性を有する。実施例４第１表に示す供試鋼No.１，２，８を熱延仕上温
度910℃、巻取温度700℃で板厚2.3mmに熱間圧延
し、0.2mmまで冷間圧延した後、第６図に示す
焼鈍サイクルにより連続焼鈍し、ブリキを製造し
た。焼鈍温度ST＝650℃、均熱時間20秒、冷却速
度10℃／sec一定である。その後調質圧延を0.8％
の圧下率で加えた。その材質結果を第６表に示
す。[Table] From the above results, the steel of the present invention has extremely good galvanizing properties. In particular, the reason why the samples in Table 3 have good adhesion despite their high Si content is that the amount of Si concentrated on the surface can be significantly reduced by rapid cooling. Samples No. 7 in Table 1 and No. 5 in Table 3 have a high Ti content, so the alloying reaction between the base iron and molten zinc is promoted, leading to overalloying and hardening in the plating layer. It is thought that the adhesion deteriorated due to the formation of a brittle alloy layer. Table 1 - No. 9, Table 3 - No.
The material No. 4 also has inferior adhesion compared to the steel of the present invention is Nb.
This is presumed to be due to the large amount added. From this point of view as well, the steel of the present invention with low amounts of Ti and Nb added as alloying elements has an advantage. Example 4 Test steel Nos. 1, 2, and 8 shown in Table 1 were hot rolled to a thickness of 2.3 mm at a finishing temperature of 910°C and a coiling temperature of 700°C, and then cold rolled to 0.2 mm. , Continuous annealing was performed using the annealing cycle shown in FIG. 6 to produce tinplate. Annealing temperature ST = 650°C, soaking time 20 seconds, cooling rate constant 10°C/sec. Then temper rolling to 0.8%
It was added at a reduction rate of . The material results are shown in Table 6.

【表】上記結果より、本発明鋼はTi、Nb添加量を著
しく低減せしめているため、再結晶温度が著しく
低く、低温焼鈍によつて極薄鋼板の製造が可能で
あり、超加工用極薄鋼板を提供するものである。
比較鋼No.８はNb添加量が高く低温焼鈍では未結
晶である。以上のように、Ti、Nbを複合添加することを
必須条件とし、TiはＮを固定する範囲内で添加
し、NbはExcessC量が10〜50ppmとなる如く添
加することにより、熱延巻取温度、焼鈍温度、冷
却温度に依らずに、従来にない極めて優れた種々
の特性を有する超深絞り用焼付硬化性鋼板が得ら
れることになり、本発明の新規性が示された。[Table] From the above results, the steel of the present invention has significantly reduced amounts of Ti and Nb added, so the recrystallization temperature is extremely low, and it is possible to manufacture ultra-thin steel sheets by low-temperature annealing, making it suitable for super-processing. It provides thin steel sheets.
Comparative steel No. 8 has a high Nb content and is non-crystallized during low-temperature annealing. As mentioned above, it is essential to add Ti and Nb in combination, Ti is added within the range that fixes N, and Nb is added so that the Excess C amount is 10 to 50 ppm. The novelty of the present invention was demonstrated by the fact that a bake-hardenable steel sheet for ultra-deep drawing having various extremely excellent properties not found in the prior art was obtained regardless of temperature, annealing temperature, and cooling temperature.

【図面の簡単な説明】[Brief explanation of the drawing]

第１図はTi添加量とＮとの図表、第２図はNb
添加量とＣとの図表、第３図、第４図はTi、Nb
複合添加鋼の材質に及ぼすTi量およびNb量の影
響を示す図表、第５図はExcessC量（Ｃ（％）−
１２／９３Nb（％））と焼付硬化性の関係を示す図表、
第６図は焼鈍サイクルを示す説明図、第７図はBH
の説明図である。 Figure 1 is a graph of Ti addition amount and N, Figure 2 is Nb
Diagrams of addition amount and C, Figures 3 and 4 are for Ti and Nb
Figure 5 shows the effect of Ti and Nb on the material properties of composite additive steel.
A diagram showing the relationship between 12/93Nb (%)) and bake hardenability,
Figure 6 is an explanatory diagram showing the annealing cycle, Figure 7 is BH
FIG.

Claims

【特許請求の範囲】１Ｃ：0.007％以下，Si：0.8％以下， Mn：1.0％以下，Ｐ：0.15％以下， Al：0.01〜0.1％、Ｎ：0.01％以下及び他の不可避的不純物から成り、かつTiとNbを複合添加することを必須条件と
し、Tiは48／14〔Ｎ（％）−0.002％〕＜Ti（％）
かつTi（％）≦48／14N（％）を満たす範囲内で含有し、 Nbは93／12〔Ｃ（％）−0.005％〕≦Nb％≦93／
12〔Ｃ（％）−0.001％〕かつNb≦0.020％を満たし、 TiとNbの合計量が0.014％〜0.027％を満たす範
囲内で、添加した成分の鋼を、Ar₃変態点以上の
仕上温度で熱延した後、600℃以上の巻取り温度
で巻取り、かかる後、冷間圧延後、再結晶温度以
上A₃点以下の温度で連続焼鈍することを特徴と
する超深絞り用焼付硬化性鋼板の製造方法。[Claims] 1 C: 0.007% or less, Si: 0.8% or less, Mn: 1.0% or less, P: 0.15% or less, Al: 0.01 to 0.1%, N: 0.01% or less, and other unavoidable impurities. The essential condition is to add Ti and Nb in combination, and Ti is 48/14 [N (%) - 0.002%] < Ti (%)
and Ti (%)≦48/14N (%) is contained within the range, and Nb is 93/12 [C (%) - 0.005%]≦Nb%≦93/
12 [C (%) - 0.001%] and satisfies Nb≦0.020%, and within the range where the total amount of Ti and Nb satisfies 0.014% to 0.027%, the steel with the added components has a finish of Ar ₃ transformation point or higher Baking for ultra-deep drawing, characterized by hot rolling at a temperature of 600°C or higher, followed by cold rolling, followed by continuous annealing at a temperature above the recrystallization temperature and below _A3 points. Method for manufacturing hardenable steel plate.