JPH0128817B2 - - Google Patents
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- Publication number
- JPH0128817B2 JPH0128817B2 JP58065218A JP6521883A JPH0128817B2 JP H0128817 B2 JPH0128817 B2 JP H0128817B2 JP 58065218 A JP58065218 A JP 58065218A JP 6521883 A JP6521883 A JP 6521883A JP H0128817 B2 JPH0128817 B2 JP H0128817B2
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- 229910000831 Steel Inorganic materials 0.000 claims description 74
- 239000010959 steel Substances 0.000 claims description 74
- 238000000137 annealing Methods 0.000 claims description 22
- 238000001816 cooling Methods 0.000 claims description 17
- 238000005098 hot rolling Methods 0.000 claims description 12
- 238000010438 heat treatment Methods 0.000 claims description 10
- 238000004519 manufacturing process Methods 0.000 claims description 10
- 229910052757 nitrogen Inorganic materials 0.000 claims description 10
- 238000005097 cold rolling Methods 0.000 claims description 8
- 238000001953 recrystallisation Methods 0.000 claims description 8
- 229910052698 phosphorus Inorganic materials 0.000 claims description 6
- 239000010960 cold rolled steel Substances 0.000 claims description 5
- 239000000203 mixture Substances 0.000 claims description 5
- 239000012535 impurity Substances 0.000 claims description 2
- 238000005096 rolling process Methods 0.000 claims description 2
- 239000000463 material Substances 0.000 description 39
- 230000000694 effects Effects 0.000 description 27
- 239000002244 precipitate Substances 0.000 description 22
- 239000006104 solid solution Substances 0.000 description 20
- 239000002131 composite material Substances 0.000 description 14
- 230000006866 deterioration Effects 0.000 description 11
- ATJFFYVFTNAWJD-UHFFFAOYSA-N Tin Chemical compound [Sn] ATJFFYVFTNAWJD-UHFFFAOYSA-N 0.000 description 9
- 239000000126 substance Substances 0.000 description 9
- 238000004804 winding Methods 0.000 description 9
- 238000006243 chemical reaction Methods 0.000 description 8
- 229910052799 carbon Inorganic materials 0.000 description 7
- 238000005336 cracking Methods 0.000 description 5
- 238000001556 precipitation Methods 0.000 description 5
- 229910001209 Low-carbon steel Inorganic materials 0.000 description 4
- 150000004767 nitrides Chemical class 0.000 description 4
- 238000005554 pickling Methods 0.000 description 4
- 230000002829 reductive effect Effects 0.000 description 4
- 238000012360 testing method Methods 0.000 description 4
- 229910019142 PO4 Inorganic materials 0.000 description 3
- 230000032683 aging Effects 0.000 description 3
- 230000007423 decrease Effects 0.000 description 3
- NBIIXXVUZAFLBC-UHFFFAOYSA-K phosphate Chemical compound [O-]P([O-])([O-])=O NBIIXXVUZAFLBC-UHFFFAOYSA-K 0.000 description 3
- 239000010452 phosphate Substances 0.000 description 3
- 230000001376 precipitating effect Effects 0.000 description 3
- 238000005728 strengthening Methods 0.000 description 3
- 239000000654 additive Substances 0.000 description 2
- 230000002411 adverse Effects 0.000 description 2
- 230000002776 aggregation Effects 0.000 description 2
- 230000015572 biosynthetic process Effects 0.000 description 2
- 238000009792 diffusion process Methods 0.000 description 2
- 230000001976 improved effect Effects 0.000 description 2
- 230000000670 limiting effect Effects 0.000 description 2
- 238000000034 method Methods 0.000 description 2
- 238000010422 painting Methods 0.000 description 2
- 230000008092 positive effect Effects 0.000 description 2
- 238000005204 segregation Methods 0.000 description 2
- 229910052717 sulfur Inorganic materials 0.000 description 2
- 229910000975 Carbon steel Inorganic materials 0.000 description 1
- 229910001021 Ferroalloy Inorganic materials 0.000 description 1
- 229910000655 Killed steel Inorganic materials 0.000 description 1
- 239000002253 acid Substances 0.000 description 1
- 150000007513 acids Chemical class 0.000 description 1
- 230000000996 additive effect Effects 0.000 description 1
- 238000005054 agglomeration Methods 0.000 description 1
- 238000004220 aggregation Methods 0.000 description 1
- 230000003679 aging effect Effects 0.000 description 1
- 229910001566 austenite Inorganic materials 0.000 description 1
- 230000008901 benefit Effects 0.000 description 1
- 230000008859 change Effects 0.000 description 1
- 230000000052 comparative effect Effects 0.000 description 1
- 239000013078 crystal Substances 0.000 description 1
- 230000002542 deteriorative effect Effects 0.000 description 1
- 238000010586 diagram Methods 0.000 description 1
- 230000009931 harmful effect Effects 0.000 description 1
- 230000006872 improvement Effects 0.000 description 1
- 230000002401 inhibitory effect Effects 0.000 description 1
- 231100000989 no adverse effect Toxicity 0.000 description 1
- 229910052760 oxygen Inorganic materials 0.000 description 1
- 230000008569 process Effects 0.000 description 1
- 239000000047 product Substances 0.000 description 1
- 230000001737 promoting effect Effects 0.000 description 1
- 230000003014 reinforcing effect Effects 0.000 description 1
- 230000002040 relaxant effect Effects 0.000 description 1
- 238000010583 slow cooling Methods 0.000 description 1
- 239000002436 steel type Substances 0.000 description 1
- 238000009864 tensile test Methods 0.000 description 1
- 238000004154 testing of material Methods 0.000 description 1
Classifications
-
- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
- C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
- C21D8/04—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips to produce plates or strips for deep-drawing
Landscapes
- Chemical & Material Sciences (AREA)
- Engineering & Computer Science (AREA)
- Physics & Mathematics (AREA)
- Thermal Sciences (AREA)
- Crystallography & Structural Chemistry (AREA)
- Mechanical Engineering (AREA)
- Materials Engineering (AREA)
- Metallurgy (AREA)
- Organic Chemistry (AREA)
- Heat Treatment Of Steel (AREA)
- Heat Treatment Of Sheet Steel (AREA)
Description
本発明は極めて優れた二次加工性を有する超深
絞り用鋼板の製造方法に関するものである。
従来、連続焼鈍用の深絞り性鋼板として、炭窒
化物形成元素を添加した極低炭素鋼が開発されて
いるが、かかる鋼板は苛酷な深絞り加工後に二次
加工を受けると脆性的に破壊する傾向を有してお
り、特にＰ，Si，Mn等を添加して高強度鋼板を
製造する場合には、Ｐ，Siは鋼板を脆化させる性
質が強いため、上記二次加工脆性は極めて発生し
易くなる。ＢはＣと同様に結晶粒界を強化する働
きがあるとされるが、本発明者等は実際に調査検
討した結果、Ｂ添加による鋼板材質への影響は鋼
種、製造条件によつて様々に異なるという以下の
新規な知見を得、これに基づき本発明を完成した
ものである。
Ti添加極低炭素鋼にＢを添加した場合には、
二次加工性は改善される傾向を示すが、その改善
効果は比較的小さく、またＢ未添加材と比較して
深絞り性（ｒ値）、延性（El）の劣化が極めて大
きいものがある。Ti添加鋼ではTiが鋼中のＯ，
Ｎ，Ｓ，Ｃとの析出物形成傾向が極めて強いため
に粒界が極めて清浄であり、粒界強度は非常に弱
い。Ｂを添加した場合にも、脆弱な粒界の性質は
残存するため、二次加工性の改善効果は比較的小
さいのである。
Nb添加極低炭素鋼にＢを添加した場合には、
添加するＢ量が微量の場合には二次加工性の改善
効果は小さく、逆に二次加工性を改善する効果が
現れる程度に、Ｂ添加量を増加した場合には、前
記Ti単独添加鋼と同様ｒ値、Elの劣化が極めて
大きい。
かかる現象の原因は、Nb添加鋼の場合には、
窒化物形成傾向がNb，Alと比較してＢの方が大
きいために、添加したＢはBNを形成し、二次加
工性を改善する効果を有する固溶Ｂの状態で存在
するものが少ないために、微量のＢ添加時には効
果が小さいものである。
固溶Ｂとして存在するＢを確保するためにはＮ
とBNを形成する量以上のＢ量を添加する必要が
ある。しかしながらBNはｒ値、Elを劣化させる
傾向が強いために、材質劣化が大きく深絞り用鋼
板として好ましくない。
更に、BNとなるＢ量は鋼巾Ｎ量によつて決ま
るために、実機製造時のＮ含有量の変動を考慮す
れば、添加Ｂ量は安全性を考えて多くする必要が
ある。
固溶Ｂとして存在する場合においても、Ｂは材
質を劣化させる傾向があることから、Ｂ添加量を
多くする必要のあるNb添加鋼では材質劣化、材
質変動が極めて大きい欠点を有するのである。
本発明者等はＢ添加に起因する上記の問題点以
外に、従来の極低炭素系深絞り用鋼板は以下の欠
点があるとの新規知見を得た。
Ti単独添加鋼は、Ti添加量をＣとＮの当量以
下にした場合には、炭化物（TiC）が微細に析出
するために延性、降伏強度、深絞り性、時効性等
の材質が著しく劣化する傾向がある。
従つて深絞り性に優れた材質を得るには、Ti
添加量をＣとＮの当量以上にする必要があり、こ
の場合には、固溶Ｃが鋼板中にほとんど存在せず
に二次加工性は極めて劣化し、更にTi添加量が
多くなるために、塗装下地処理として施されるリ
ン酸塩処理性の劣化が大きい。
Nb添加鋼では、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。
これはAlN，NbCの生成に関係しており、こ
れら析出物が熱延板中にて十分な大きさを持つた
析出物になつていないために、再結晶を抑制する
ためと考えられる。
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。
これはAlN，NbCの生成に関係し、高温巻取
では熱延板中にこれら析出物が、大きな寸法の析
出物として生成するためである。
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とすだけでな
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、歩留りの低下はNb添加鋼では
特に大きいものがある。
第２は冷延後の焼鈍温度と焼鈍後の冷却速度の
問題である。
特開昭55−141526号、特開昭55−141555号公報
にある如く、高温（約900℃以上）で焼鈍すると
AlN，NbCが再溶解するために固溶Ｃ，Ｎが出
来て、焼鈍後徐冷をしなければ遅時効性にはなら
ない。従つて操業性、経済性の面から問題とな
る。
本発明は、これら従来のTi添加鋼及びNb添加
鋼の持つ欠点をなくした鋼の製造に関するもので
もある。
即ち、優れた深絞り性と苛酷な深絞り加工を受
けた場合にも、二次加工割れの発生しにくい性能
を有し、熱延巻取条件に鈍感な鋼板の開発を目的
として行なわれたものであり、その骨子は、Ｃ：
0.005％以下、Si：0.8％以下、Mn：1.0％以下、
Ｐ：0.1％以下、Al：0.01〜0.1％、Ｎ：0.005％以
下及び他の不可避的不純物から成り、かつＢ，
Ti，Nbを複合添加することを必須条件とし、Ｂ
は2ppm以上10ppm未満の範囲内で添加し、Tiは
48／14〔Ｎ（％）−0.003％〕≦Ti（％）≦48／12C
（％）＋48／14N（％）の条件を満たす範囲内で含
有し、Nbは7C（％）＞Nb（％）＞2C（％）でかつ
0.003％≦Nb＜0.035％を満たす範囲内の含有量
で、かつ0.017％＜Ti＋Nb（％）＜0.06％を満たす
成分の鋼を加熱温度1300℃以下の条件で熱間圧延
した後、脱スケール処理、延間圧延後、再結晶温
度以上AC3点以下の温度で連続焼鈍を施した後、
650℃から450℃の間を平均50℃／sec超の冷却速
度で冷却することを特徴とする極めて優れた二次
加工性を有する超深絞り用冷延鋼板の製造方法で
ある。
本発明鋼の基本原理を以下に述べる。
本発明鋼は、鋼板中に存在する固溶Ｂおよび固
溶Ｃの粒界濃化による粒界強度を著しく高め、極
めて優れた二次加工性を付与することを発明の根
本思想とする。
さらに、鋼中に添加したＢを固溶Ｂとして上記
効果を発揮せしめるために、Tiを複合添加する。
複合添加するTiの効果は、鋼中のＮをTiNとし
て析出固定することにより、添加したＢがBNを
形成するのを妨げ、固溶Ｂとなすものである。
従つて、添加Ｂ量は微量で有効であり、Ｂ添加
による延性（El）、深絞り性（ｒ値）の劣化を抑
制できる。更に、複合添加するNbの効果は、鋼
中のＣの一部をNbCとして析出固定し、固溶Ｃ
量を実質的に非時効となる如く低減することを目
的とするものである。
本発明鋼が従来のTi単独添加鋼、Nb単独添加
鋼と比較して優れた深絞り性と、二次加工性を共
に兼備しているのは、微量のＢ，Ti，Nbを複合
添加することによるものである。
即ち、複合添加した微量のTiによつて、鋼中
のＮはTiNとして既に熱延加熱炉中で析出固定
されている。
TiNは窒化物として極めて安定であるので熱
延、冷延、再結晶焼鈍の各工程において何ら変化
するものではなく、従つて熱延の巻取温度や連続
焼鈍温度やその後の冷却速度によつてその析出形
態は変わらない。
鋼中に添加したＢは、窒化物形成傾向がTiに
比べて小さいため、固溶Ｂとして存在し、微量の
添加量で粒界強度を高める効果を有するのであ
る。これに対して、Nb単独添加鋼にＢを添加し
た場合、窒化物形成傾向はNb，Alに比べてＢの
方が大きいため、添加したＢはBNを形成する。
従つてＢ添加量が少ない場合は、二次加工性を
改善する効果を有する固溶Ｂが存在しない（Ｂは
BNとして存在）ために、二次加工性改善効果は
ない。
固溶Ｂとして存在するＢを確保するには、Ｎと
BNを形成する量以上のＢを添加する必要がある
が、BN及び固溶Ｂはｒ値、Elを劣化させる傾向
が強いために、Ｂ添加量を多くすることは材質劣
化を招き、深絞り用鋼板として好ましくない。
また、Ti単独添加鋼にＢを添加した場合は、
公知の如く、Ｂ未添加材と比較してｒ値、Elの劣
化が極めて大きい。更に、Tiは鋼中のＯ，Ｎ，
Ｓ，Ｃとの析出物形成傾向が極めて強いために、
粒界が極めて清浄で粒界強度は非常に弱い。
従つてＢを添加して脆弱な粒界の性質を改善す
るためには、Ｂ添加量を多くする必要があり、材
質の観点から好ましくない。
これに対して本発明鋼におけるTiの添加は、
ＮをTiNとして析出固定するための役割をなす
ものであり、上記Ti単独添加鋼にみられる欠点
を引き起こすものではない。
添加Ｂ量を種々の添加量で一定値に固定し、本
発明鋼と上記Ti，Nb単独添加鋼の材質を繰り返
し比較調査した結果においても、本発明鋼は最も
延性、深絞り性が優れており、二次加工性の点か
らも、明確な優位性を示した。
本発明鋼はＢ，Tiと共にNbを複合添加するも
のであるが、Ti，Nbの共存により（Ti，Nb）
Ｃの如き複合析出物が熱間圧延時の仕上前（即ち
オーステナイト温度域）から形成されて析出を始
めるために、巻方温度が低目でもかなり良好な材
質を得ることができる。
本発明鋼が従来のNb単独添加鋼と比較して優
れた材質特性を有するのは、ＮをTiNとして
析出固定することにより、微量Ｂ添加により安定
して二次加工性を著しく向上できる点Ｎを
TiNとして熱延加熱炉中で、既に析出させてＮ
に起因する巻取温度の材質への変動要因をなくし
たことTi，Nbの複合添加により、（Ti，Nb）
Ｃの如き複合析出物を仕上前から形成して巻取温
度が低目でもかなり良好な材質を得ることができ
る点にある。
次に成分範囲について述べる。
まずＢ添加量については、2ppm以上10ppm未
満の範囲内で添加する必要がある。本発明鋼にお
けるＢの添加は、二次加工性の向上効果にあり、
その効果は固溶状態で存在するＢによるものであ
る。
本発明鋼では、Tiの複合添加によりＮをTiN
として析出せしめているため、添加したＢは固溶
Ｂとなり、添加量は微量で十分に有効である。
Ｂ添加量が増加するとｒ値、Elが若干劣化する
傾向にあり、超深絞り用鋼板という本発明鋼の特
性から、上限を2ppm以上10ppm未満とする。
TiはＮを固定してその害をなくすために添加
するものであり、48／14〔Ｎ（％）−0.003％〕以上
の添加を必要とする。即ちTi添加量の下限は、
計算上Tiで析出固定できないＮ量が30ppm以下
である。
通常のアルミキルド鋼では、30ppmのＮは悪影
響を及ぼす量であるが、Tiを複合添加すると、
TiNを析出核としてAlNを析出した（Ti，Al）
Ｎの複合析出物が形成され、極めて高温から安定
析出物となるため、実質上全Ｎ量をTiNとして
析出させたのと同様の効果を有するとの知見を得
た。
上記効果を十分顕著ならしめるには、0.002％
以上のTi添加量が望ましい。またＣとＮの和の
当量を越えて添加すると、Ti添加鋼と同様の性
質が強くなり、二次加工割れが発生し易くなるた
め、上限を48／12C（％）＋48／14N（％）未満と
する。延性、降伏強度および経済的観念からは
Ti添加量はTiCを生成しない48／14N（％）以下
で、0.015％以下が最も好ましい。
Nbの添加量は、複合析出物を形成するために
は、2C（％）以上の添加を必要とし、かつ0.003％
未満ではその効果は小さく、またNb添加量が7C
（％）以上の場合は、NbCの組成に近い析出物に
なり、Nb単独添加鋼の持つ欠点が如実に現れる
ことになり望ましくない。最も好ましくはNb＜
0.025％の添加量である。
又本発明の特徴は固溶Ｃを残存させることにあ
り、NbをＣとの化学当量すなわち93／12C（％）
〔＝7.75C（％）〕以上入れると固溶Ｃがなくなり、
本発明の特徴が出せなくなる。
このためNbは7.75C（％）以下にしなければな
らないが、優れた二次加工性を得るためには更に
Nbの量を下げ、最低限の固溶Ｃ量を確保する必
要があり、このためNbの上限は7.0C（％）とす
る。
また冷延鋼板は、塗装下地処理としてリン酸塩
処理（ボンデ処理）を施されるが、いわゆるボン
デ性にも優れたものである必要がある。しかし、
極低炭素鋼では、NbやTiを添加するとボンデ性
が大きく劣化する性質がある。
特に溶接部をグラインダー手入れして新生面の
露出した場所についても、良好な化成処理性を保
障するには、Ti，Nb添加量をTi（％）＋Nb（％）
＜0.06％に制限することが必要である。最も望ま
しくはTi（％）＋Nb（％）＜0.05％の範囲である。
（Ti，Nb）Ｃの如き複合析出物を生成させる
ためには、TiとNbの量は実施例に示す如くTi
（％）＋Nb（％）＞0.017％が必要である。
次にＢ，Ti，Nb以外の元素の範囲について記
す。
Ｃは量が多いと、必然的にＣを固定するための
Nb量が多くなり、製造コストが高くなり、また
複合析出物の生成量が増えるため、析出強化要素
が大きくなり材質の低下を招く。このため0.005
％以下とする。
Siは高強度鋼板にする場合添加することがある
が、脆性を助長する元素であり、かつ化成処理性
を阻害する元素でもあり、0.8％以下にすべきで
ある。
Mnも高強度化するに際して、使用することが
できる。しかしｒ値を劣化させる性質があるこ
と、合金鉄のコストが高いことから1.0％以下に
する。
Ｐは、最も強化能の大きな元素であり、高強度
化する場合添加されるが、多量に含まれると粒界
偏析量が多くなつて脆化、即ち二次加工割れをひ
き起こすので上限は0.1％とする。
Ｎは、（Ti，Al）Ｎとして実質的に全Ｎ量が固
定されるが、Ｎ含有量が多いと、Ti添加量も多
く必要になるので0.005％以下とする。Ｃ，Ｎを
50ppm以下の極低量範囲に制限することにより、
析出物量が減少し、延性が良好で降伏強度が低く
なり、Ti，Nb添加量が増加した場合の悪影響は
軽減される傾向を示す。
次に製造条件について述べる。
本発明鋼は、ＮをTiによつて析出固定するこ
とにより無害化しており、またTi，Nbの複合添
加により（Ti，Nb）Ｃ複合析出物を高温から析
出させているが、熱延加熱温度を1300℃以下とす
ることにより、これら析出物あるいは析出核が加
熱炉中で十分存在することになる。
この結果、微量のTi添加量で実質上全Ｎ量を
（Ti，Al）Ｎとして析出させることが可能となつ
たものであり、また、（Ti，Nb）Ｃ複合析出物
が、仕上前の高温域から析出し始めることにな
る。
従つて、低目の巻取温度でも、熱延板の状態で
析出物がかなり凝集し、巻取温度に鈍感な材質挙
動を示すとの新規知見を得たのである。加熱温度
を1300℃以下に制限することにより、析出物の凝
集度がよくなり、その悪影響が低下することか
ら、Ti添加量、Nb添加量の上限も若干緩和され
る。
また、材質特にｒ値が向上することから、二次
加工性に対しても好影響を与え、Ｂ添加効果は顕
著に現われ2ppm以上の添加量で十分有効である。
析出物の粗大凝集を促進することは、化成処理
性に対しても好影響を及ぼし、Ti，Nb添加量総
和の上限を緩和する。即ち（Ti，Nb）Ｃ、（Ti，
Al）Ｎ等の析出物は、Fe3Cに比べて酸に溶解し
にくいため、リン酸塩結晶が析出しにくく、化成
処理性に悪影響を及ぼすものであるが、凝集させ
ることで、かかる析出物密度が減少し、化成処理
性が改善されるのである。
本発明鋼では、他の熱間圧延条件は特に規定す
る必要はない。ただし熱延仕上温度が低下するに
伴い、ｒ値、Elが低下する傾向があることから、
850℃以上の仕上温度が好ましい。巻取温度に関
しても前記理由により特に規定する必要はない。
冷間圧延条件についても特に規定する必要はな
い。冷延率を増加するに伴い、深絞り性は向上す
る傾向があり、二次加工脆性割れは鋼板のｒ値が
高い程発生し難いことから、本発明鋼の特性を更
に優位づけるためは、50％以上の冷延率が最も好
ましい。
本発明鋼はTi，Nb添加量が微量でよいため、
再結晶温度は低いが、冷延率を増加することは、
更に再結晶温度を低下させ焼鈍温度を下げること
に対しても有効である。
焼鈍条件については、再結晶温度以上AC3点以
下の温度で連続焼鈍することとする。箱型焼鈍は
冷却速度が極めて遅いため、冷却中にＰの粒界へ
の拡散が起こり望ましくない。
冷延鋼板を製造する場合には、焼なまし処理後
の冷却速度の制限が必要である。本発明鋼は二次
加工性には極めて優れた材料であるが、あまり遅
い冷却速度では、Ｐ等の粒界偏析により二次加工
性は発生し易くなる傾向はある。
Ｐの粒界への拡散を考えると、650℃から450℃
の間の冷却速度が問題で、その冷却速度を50℃／
sec超にすべきである。
以下、実施例について述べる。
実施例 １
第１表に示す成分の鋼スラブを溶製し、第１表
に示す熱延条件により熱間圧延をした。
仕上温度はいずれも890〜910℃である。熱延板
厚さは3.8mmであり、酸洗後0.8mmに冷間圧延した
後、連続焼鈍炉にて焼鈍した。焼鈍サイクルは約
10℃／secで、780〜820℃まで加熱して、該温度
範囲に40秒保持した後、室温まで平均冷速50〜
100℃／secで冷却した。
第１図は焼鈍サイクルを示す。
スキンパスを0.8％かけた後材質試験に供して、
その結果を化成処理性、二次加工割れ試験の結果
と共に第２表に示す。巻取温度の高い一部の材料
については熱延コイル長手方向中心部（上段）、
長手方向端部（下段）相当位置の材質を示した。
The present invention relates to a method for manufacturing a steel plate for ultra-deep drawing that has extremely excellent secondary workability. Conventionally, ultra-low carbon steels with added carbonitride-forming elements have been developed as deep-drawable steel sheets for continuous annealing, but such steel sheets become brittle and fracture when subjected to secondary processing after severe deep drawing. In particular, when producing high-strength steel sheets by adding P, Si, Mn, etc., the above-mentioned secondary work embrittlement is extremely difficult because P and Si have a strong tendency to embrittle steel sheets. It is more likely to occur. B, like C, is said to have the effect of strengthening grain boundaries, but as a result of actual research and study by the present inventors, we found that the effect of B addition on the steel sheet material varies depending on the steel type and manufacturing conditions. The present invention has been completed based on the following new findings that the invention is different. When B is added to Ti-added ultra-low carbon steel,
Although the secondary workability tends to improve, the improvement effect is relatively small, and there are cases where the deterioration of deep drawability (r value) and ductility (El) is extremely large compared to materials without B additives. . In Ti-added steel, Ti is O in the steel,
Since the tendency to form precipitates with N, S, and C is extremely strong, the grain boundaries are extremely clean and the grain boundary strength is extremely weak. Even when B is added, the brittle grain boundary properties remain, so the effect of improving secondary workability is relatively small. When B is added to Nb-added ultra-low carbon steel,
When the amount of B added is small, the effect of improving secondary workability is small, and on the contrary, when the amount of B added is increased to the extent that the effect of improving secondary workability appears, the above-mentioned Ti-only added steel Similarly, the deterioration of r value and El is extremely large. The cause of this phenomenon is that in the case of Nb-added steel,
Since B has a greater tendency to form nitrides than Nb and Al, the added B forms BN and very little exists in the form of solid solution B, which has the effect of improving secondary workability. Therefore, when a trace amount of B is added, the effect is small. In order to secure B existing as solid solution B, N
It is necessary to add an amount of B greater than the amount required to form BN. However, since BN has a strong tendency to deteriorate the r value and El, the material deteriorates significantly and is not preferred as a steel sheet for deep drawing. Furthermore, since the amount of B that becomes BN is determined by the amount of N in the steel width, the amount of B added needs to be increased in consideration of safety, taking into account the fluctuations in the N content during actual manufacturing. Even when B exists as a solid solution, B tends to deteriorate the material, so Nb-added steel, which requires a large amount of B, has the disadvantage of extremely large material deterioration and material fluctuation. In addition to the above-mentioned problems caused by the addition of B, the present inventors have obtained new knowledge that conventional ultra-low carbon steel sheets for deep drawing have the following drawbacks. In steel with only Ti added, if the amount of Ti added is less than the equivalent amount of C and N, carbide (TiC) will precipitate finely, resulting in significant deterioration of material properties such as ductility, yield strength, deep drawability, and aging resistance. There is a tendency to Therefore, in order to obtain a material with excellent deep drawability, Ti
It is necessary to make the addition amount equal to or higher than the equivalent amount of C and N. In this case, there is almost no solid solution C in the steel sheet, resulting in extremely poor secondary workability, and the addition amount of Ti increases. , the deterioration of the phosphate treatment applied as a base treatment for painting is significant. For Nb-added steel, 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 the possible temperature range of continuous annealing furnaces (usually about
(below 850℃), unrecrystallized areas remain, and the material quality varies greatly depending on the amount of Nb. This is related to the formation of AlN and NbC, and is thought to be because these precipitates do not have sufficient size in the hot-rolled sheet to suppress recrystallization. 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 about 800 to 850°C, except for the longitudinal ends of the hot-rolled coil. It is. This is related to the formation of AlN and NbC, and is because these precipitates form in the hot-rolled sheet as large-sized precipitates during high-temperature coiling. 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 the normal winding temperature. Since the material cannot be obtained, the decrease in yield is particularly large in Nb-added steel. The second problem is the annealing temperature after cold rolling and the cooling rate after annealing. As stated in JP-A-55-141526 and JP-A-55-141555, annealing at high temperatures (approximately 900°C or higher)
Since AlN and NbC are redissolved, solid solution C and N are formed, and slow aging cannot be achieved unless slow cooling is performed after annealing. Therefore, it becomes a problem from the viewpoint of operability and economy. The present invention also relates to the production of steel that eliminates the drawbacks of these conventional Ti-added steels and Nb-added steels. In other words, the objective was to develop a steel sheet that has excellent deep drawability and is resistant to secondary processing cracking even when subjected to severe deep drawing processing, and is insensitive to hot rolling winding conditions. The gist of it is C:
0.005% or less, Si: 0.8% or less, Mn: 1.0% or less,
Consisting of P: 0.1% or less, Al: 0.01 to 0.1%, N: 0.005% or less and other unavoidable impurities, and B,
The essential condition is to add Ti and Nb in combination, and B
is added within the range of 2ppm or more and less than 10ppm, and Ti is added within the range of 2ppm or more and less than 10ppm.
48/14 [N (%) - 0.003%] ≦Ti (%) ≦48/12C
(%) + 48/14N (%), and Nb is 7C (%) > Nb (%) > 2C (%) and
Steel with a content within the range of 0.003%≦Nb<0.035% and a composition satisfying 0.017%<Ti+Nb (%)<0.06% is hot rolled at a heating temperature of 1300℃ or less, and then descaled. , After rolling, continuous annealing is performed at a temperature above the recrystallization temperature and below AC 3 points,
This is a method for producing a cold-rolled steel sheet for ultra-deep drawing, which has extremely excellent secondary workability, and is characterized by cooling between 650°C and 450°C at an average cooling rate of over 50°C/sec. The basic principle of the steel of the present invention will be described below. The basic idea of the present invention is to significantly increase grain boundary strength through grain boundary concentration of solid solution B and solid solution C present in the steel sheet, and to provide extremely excellent secondary workability. Furthermore, Ti is added in combination in order to exhibit the above effects by turning the B added into the steel into a solid solution B.
The effect of the composite addition of Ti is to precipitate and fix N in the steel as TiN, thereby preventing the added B from forming BN and forming solid solution B. Therefore, the amount of B added is effective even in a small amount, and the deterioration of ductility (El) and deep drawability (r value) due to B addition can be suppressed. Furthermore, the effect of composite addition of Nb is that some of the C in the steel is precipitated and fixed as NbC, and solid solution C is
The objective is to reduce the amount so that it is substantially non-aging. The reason why the steel of the present invention has superior deep drawability and secondary workability compared to conventional steels with only Ti added and steels with only Nb added is because of the combined addition of trace amounts of B, Ti, and Nb. This is due to a number of things. That is, due to the small amount of Ti added in a composite manner, N in the steel is already precipitated and fixed as TiN in the hot rolling furnace. Since TiN is extremely stable as a nitride, it does not change in any way during the hot rolling, cold rolling, and recrystallization annealing processes. The precipitation form remains unchanged. Since B added to steel has a smaller tendency to form nitrides than Ti, it exists as solid solution B, and has the effect of increasing grain boundary strength with a small amount of addition. On the other hand, when B is added to steel with only Nb added, B has a greater tendency to form nitrides than Nb and Al, so the added B forms BN. Therefore, when the amount of B added is small, there is no solid solution B that has the effect of improving secondary processability (B is
(exists as BN), so it has no effect on improving secondary workability. To ensure that B exists as a solid solution, N and
It is necessary to add B in an amount greater than the amount required to form BN, but since BN and solid solution B have a strong tendency to deteriorate r value and El, increasing the amount of B added will lead to material deterioration, resulting in deep drawing. It is not preferred as a steel plate for industrial use. In addition, when B is added to steel with only Ti added,
As is well known, the deterioration of the r value and El is extremely large compared to the material to which B is not added. Furthermore, Ti is O, N,
Because it has an extremely strong tendency to form precipitates with S and C,
The grain boundaries are extremely clean and the grain boundary strength is very weak. Therefore, in order to improve the properties of brittle grain boundaries by adding B, it is necessary to increase the amount of B added, which is not preferable from the viewpoint of material quality. On the other hand, the addition of Ti in the steel of the present invention
This serves to precipitate and fix N as TiN, and does not cause the drawbacks seen in the steel with only Ti added above. The results of repeatedly comparing the materials of the steel of the present invention and the above-mentioned steel with only Ti and Nb added by fixing the amount of added B to a constant value at various additive amounts also show that the steel of the present invention has the best ductility and deep drawability. It also showed clear superiority in terms of secondary processability. The steel of the present invention has a composite addition of Nb along with B and Ti, but due to the coexistence of Ti and Nb, (Ti, Nb)
Since composite precipitates such as C are formed and begin to precipitate before finishing during hot rolling (ie, in the austenite temperature range), a fairly good material can be obtained even at a low winding temperature. The reason why the steel of the present invention has superior material properties compared to conventional steels with only Nb added is that by precipitating and fixing N as TiN, the addition of a small amount of B can stably improve secondary workability.N of
TiN has already been precipitated in the hot rolling heating furnace.
The factor of variation in the winding temperature due to the material has been eliminated. By the combined addition of Ti and Nb,
The advantage is that composite precipitates such as C can be formed before finishing and a fairly good quality material can be obtained even at a low winding temperature. Next, we will discuss the component range. First, the amount of B added must be within the range of 2 ppm or more and less than 10 ppm. The addition of B in the steel of the present invention has the effect of improving secondary workability,
This effect is due to B existing in solid solution. In the steel of the present invention, N is replaced with TiN by the composite addition of Ti.
Since B is precipitated as B, the added B becomes solid solution B, and even a small amount is sufficiently effective. As the amount of B added increases, the r value and El tend to deteriorate slightly. Considering the characteristics of the steel of the present invention, which is a steel plate for ultra-deep drawing, the upper limit is set to 2 ppm or more and less than 10 ppm. Ti is added to fix N and eliminate its harmful effects, and needs to be added in an amount of 48/14 [N (%) - 0.003%] or more. In other words, the lower limit of the amount of Ti added is
Calculations show that the amount of N that cannot be precipitated and fixed by Ti is 30 ppm or less. In normal aluminium-killed steel, 30 ppm of N has an adverse effect, but when Ti is added in combination,
AlN was precipitated using TiN as a precipitation nucleus (Ti, Al)
It was found that a composite precipitate of N is formed and becomes a stable precipitate even at extremely high temperatures, so that it has the same effect as precipitating substantially the entire amount of N as TiN. In order to make the above effect sufficiently noticeable, 0.002%
It is desirable to add Ti in the above amount. Furthermore, if more than the equivalent of the sum of C and N is added, the properties similar to Ti-added steel will become stronger and secondary processing cracks will occur more easily, so the upper limit should be set at 48/12C (%) + 48/14N (%). less than From ductility, yield strength and economic considerations
The amount of Ti added is 48/14N (%) or less, which does not generate TiC, and is most preferably 0.015% or less. The amount of Nb added requires 2C (%) or more to form a composite precipitate, and 0.003%
The effect is small when the amount of Nb added is less than 7C.
(%) or more, the precipitates will have a composition close to that of NbC, and the drawbacks of steel with only Nb added will clearly appear, which is undesirable. Most preferably Nb<
The amount added is 0.025%. Moreover, the feature of the present invention is that solid solution C remains, and the chemical equivalent of Nb with C, that is, 93/12C (%)
If you add more than [=7.75C (%)], there will be no solid solution C,
The features of the present invention cannot be brought out. For this reason, Nb must be kept below 7.75C (%), but in order to obtain excellent secondary workability, it must be further reduced.
It is necessary to reduce the amount of Nb and secure the minimum amount of solid solution C, and therefore the upper limit of Nb is set at 7.0C (%). Furthermore, cold-rolled steel sheets are subjected to phosphate treatment (bonding treatment) as a base treatment for painting, but they also need to have excellent bonding properties. but,
Ultra-low carbon steel has the property that the addition of Nb or Ti significantly deteriorates bonding properties. In order to ensure good chemical conversion properties, especially for areas where welded parts are cleaned with a grinder and new surfaces are exposed, the amount of Ti and Nb added must be Ti (%) + Nb (%).
It is necessary to limit it to <0.06%. The most desirable range is Ti (%) + Nb (%) < 0.05%. In order to form a composite precipitate such as (Ti, Nb)C, the amounts of Ti and Nb should be adjusted as shown in the example.
(%) + Nb (%) > 0.017% is required. Next, the range of elements other than B, Ti, and Nb will be described. When the amount of C is large, it is necessary to fix C.
As the amount of Nb increases, manufacturing costs increase, and the amount of composite precipitates produced increases, precipitation strengthening factors become large, leading to deterioration of material quality. For this reason 0.005
% or less. Si is sometimes added when making high-strength steel sheets, but it is an element that promotes brittleness and also inhibits chemical conversion treatment properties, so it should be kept at 0.8% or less. Mn can also be used to increase the strength. However, since it has the property of deteriorating the r value and the cost of ferroalloy is high, it is limited to 1.0% or less. P is an element with the greatest strengthening ability and is added to increase strength, but if it is included in a large amount, the amount of grain boundary segregation increases and causes embrittlement, that is, secondary work cracking, so the upper limit is 0.1 %. Substantially the total amount of N is fixed as (Ti, Al)N, but if the N content is large, a large amount of Ti is required to be added, so it is set to 0.005% or less. C, N
By limiting the amount to an extremely low amount of 50ppm or less,
The amount of precipitates is reduced, ductility is good, yield strength is low, and the negative effects of increasing the amount of Ti and Nb added tend to be alleviated. Next, the manufacturing conditions will be described. In the steel of the present invention, N is rendered harmless by precipitating and fixing it with Ti, and (Ti, Nb)C composite precipitates are precipitated at high temperatures by the combined addition of Ti and Nb. By setting the temperature to 1300° C. or lower, these precipitates or precipitation nuclei are sufficiently present in the heating furnace. As a result, it has become possible to precipitate virtually the entire amount of N as (Ti, Al)N by adding a small amount of Ti, and (Ti, Nb)C composite precipitates can be precipitated before finishing. Precipitation begins in the high temperature range. Therefore, we have obtained the new knowledge that even at a low coiling temperature, the precipitates are considerably agglomerated in the hot rolled sheet state, and the material behavior is insensitive to the coiling temperature. By limiting the heating temperature to 1300° C. or less, the degree of aggregation of precipitates is improved and the adverse effects thereof are reduced, so the upper limits of the amounts of Ti and Nb added are also slightly relaxed. In addition, since it improves the quality of the material, especially the r value, it has a positive effect on the secondary workability, and the effect of B addition is noticeable and is sufficiently effective when added in an amount of 2 ppm or more. Promoting coarse agglomeration of precipitates also has a positive effect on chemical conversion treatment properties, relaxing the upper limit on the total amount of Ti and Nb added. That is, (Ti, Nb)C, (Ti,
Precipitates such as Al)N are less soluble in acids than Fe 3 C, making it difficult for phosphate crystals to precipitate and having a negative impact on chemical conversion properties. The material density is reduced and chemical conversion treatment properties are improved. In the steel of the present invention, there is no need to specify other hot rolling conditions. However, as the hot rolling finishing temperature decreases, the r value and El tend to decrease.
A finishing temperature of 850°C or higher is preferred. There is no need to particularly specify the winding temperature for the above reasons. There is no need to particularly specify the cold rolling conditions. As the cold rolling rate increases, deep drawability tends to improve, and secondary work brittle cracking is less likely to occur as the r value of the steel sheet increases. Therefore, in order to further improve the characteristics of the steel of the present invention, A cold rolling rate of 50% or more is most preferred. Since the steel of the present invention requires only a small amount of Ti and Nb to be added,
Although the recrystallization temperature is low, increasing the cold rolling rate
Furthermore, it is effective in lowering the recrystallization temperature and lowering the annealing temperature. Regarding the annealing conditions, continuous annealing is performed at a temperature above the recrystallization temperature and below AC 3 points. Since the cooling rate of box-shaped annealing is extremely slow, diffusion of P to grain boundaries occurs during cooling, which is undesirable. When manufacturing cold-rolled steel sheets, it is necessary to limit the cooling rate after annealing. Although the steel of the present invention is a material with extremely excellent secondary workability, if the cooling rate is too slow, secondary workability tends to occur easily due to grain boundary segregation of P and the like. Considering the diffusion of P into the grain boundaries, the temperature ranges from 650℃ to 450℃.
The problem is the cooling rate between
It should be over sec. Examples will be described below. Example 1 Steel slabs having the components shown in Table 1 were melted and hot rolled under the hot rolling conditions shown in Table 1. The finishing temperature is 890-910°C in all cases. The hot-rolled sheet thickness was 3.8 mm, and after pickling, it was cold rolled to 0.8 mm and then annealed in a continuous annealing furnace. The annealing cycle is approx.
After heating to 780-820℃ at 10℃/sec and holding in the temperature range for 40 seconds, average cooling rate 50-820℃ to room temperature.
It was cooled at 100°C/sec. FIG. 1 shows an annealing cycle. After applying a skin pass of 0.8%, it was subjected to a material test.
The results are shown in Table 2 along with the results of chemical conversion treatment and secondary processing cracking tests. For some materials with high coiling temperatures, the longitudinal center of the hot rolled coil (upper stage),
The material of the position corresponding to the longitudinal end (lower stage) is shown.
【表】【table】
【表】【table】
【表】
本発明品（供試鋼No.１〜４）はいずれも良好な
結果を示している。
供試鋼No.５はTi，Nb添加量が多い（Ti（％）＋
Nb（％）＞0.06％）ために化成処理性が劣る。No.
６は熱延加熱温度が高いために、Tiの複合添加
効果が小さく、No.１と比較して材質、二次加工性
が劣る。
No.７はＢを添加していないために二次加工割れ
が発生し易く、逆にNo.８はＢ添加量が多過ぎて
YP，El，ｒ値が良くない。
No.９はTi添加量が多いために、Ti添加鋼に近
い性質となり、二次加工性、化成処理性が劣る。
No.10はNb量が少ないため固溶Ｃが多くなり、時
効性が大きく材質も劣る。
No.11はNb量が多すぎてNb添加鋼に近い材質と
なり、700℃以下の巻取温度では良好な材質が得
られない。
No.12〜14はTiを添加しない材料で、この場合
はＢはＮとBNを形成するために、Ｂによる二次
加工性改善効果がない（No.12）。
また巻取温度の低い場合（No.13）に材質劣化が
大きい。No.14の如くＢ添加量を増やすと二次加工
性は改善されるが、材質が劣る。
No.15，16はTi添加鋼にNbを添加せずにＢだけ
を添加した場合であるが、この場合はＢ添加によ
る材質劣化が大きく、二次加工性自体の改善効果
が小さく、更に化成処理性が劣る。
比較例 １
次に冷却速度が低い場合について述べる。
第１表に示すNo.２，４の成分の鋼スラブを用い
て、加熱温度1200℃、仕上温度900℃、巻取温度
各々700℃，650℃で3.8mm厚のコイルに巻取つた。
酸洗後、冷間圧延をして0.8mmのコイルにして
から第２図、第３図のイ〜ホのサイクルで、連続
焼鈍後0.8％のスキンパスをかけた。
第３表は連続焼鈍後の冷却条件を示し、第４表
はかかる条件下で得られた冷延鋼板の材質結果を
示す。[Table] All of the products of the present invention (test steel Nos. 1 to 4) showed good results. Test steel No. 5 has a large amount of Ti and Nb added (Ti (%) +
Nb (%)>0.06%), chemical conversion treatment properties are poor. No.
No. 6 has a high hot rolling heating temperature, so the composite addition effect of Ti is small, and compared to No. 1, the material quality and secondary workability are inferior. No. 7 does not contain B, so secondary processing cracks tend to occur, while No. 8, on the other hand, has too much B added.
YP, El, r values are not good. Since No. 9 has a large amount of Ti added, its properties are similar to those of Ti-added steel, and its secondary workability and chemical conversion treatment properties are poor.
No. 10 has a small amount of Nb, so there is a lot of solid solution C, and the aging property is high and the material quality is also inferior. No. 11 has too much Nb, making the material similar to Nb-added steel, and a good material cannot be obtained at a coiling temperature of 700°C or lower. Nos. 12 to 14 are materials to which Ti is not added, and in this case, B forms BN with N, so B has no effect of improving secondary workability (No. 12). In addition, material deterioration is large when the winding temperature is low (No. 13). As shown in No. 14, increasing the amount of B added improves secondary workability, but the material quality is inferior. Nos. 15 and 16 are cases in which only B is added without adding Nb to Ti-added steel, but in this case, the addition of B causes significant material deterioration, the effect of improving secondary workability itself is small, and furthermore, Processability is poor. Comparative Example 1 Next, a case where the cooling rate is low will be described. Using steel slabs having compositions No. 2 and 4 shown in Table 1, they were wound into a 3.8 mm thick coil at a heating temperature of 1200°C, a finishing temperature of 900°C, and a winding temperature of 700°C and 650°C, respectively. After pickling, it was cold rolled into a 0.8 mm coil, and then subjected to continuous annealing and a 0.8% skin pass using cycles I to E in Figures 2 and 3. Table 3 shows the cooling conditions after continuous annealing, and Table 4 shows the material properties of cold rolled steel sheets obtained under these conditions.
【表】【table】
【表】
は割れの発生しない最大絞り比を示した
。 引張試験値はホツトコイル長手方向中
心部に相当する位置のものを示した。
650℃から450℃の間の平均冷却速度が30℃／
sec，10℃／sec，３℃／sec，２℃／secと遅いサ
イクルでは、二次加工性は高いレベルにあるもの
の実施例１と比較すると低下する傾向を示す。
二次加工性の著しく優れた本発明鋼の特性を最
大限に発揮せしめるためには、冷却速度を50℃／
sec超にすべきである。
実施例 ２
第５表に示す成分の鋼スラブを用いて加熱温度
1180℃、仕上温度890℃、巻取温度680℃にて熱間
圧延し、3.8mmのコイルとした。酸洗、冷間圧延
を行なつて0.8mmのコイルとした後、第１図に示
すサイクルで連続焼鈍し、スキンパスを0.8％か
けた後、材質試験に供した。その結果を第６表に
示す。[Table] shows the maximum drawing ratio without cracking. The tensile test value is in the longitudinal direction of the hot coil.
The position corresponding to the heart is shown.
The average cooling rate between 650℃ and 450℃ is 30℃/
sec, 10° C./sec, 3° C./sec, and 2° C./sec, the secondary workability is at a high level but tends to be lower compared to Example 1. In order to maximize the properties of the steel of the present invention, which has extremely excellent secondary workability, the cooling rate must be set to 50℃/
Should be over sec. Example 2 Using a steel slab with the composition shown in Table 5, the heating temperature was
It was hot rolled at 1180°C, finishing temperature 890°C, and coiling temperature 680°C to form a 3.8 mm coil. After pickling and cold rolling to form a 0.8 mm coil, it was continuously annealed in the cycle shown in Figure 1, subjected to a skin pass of 0.8%, and then subjected to material testing. The results are shown in Table 6.
【表】【table】
【表】
従来、高ｒ値を有する高強度鋼板は、TS＝40
Kgｆ／mm2級が限界であつた。これは更に強度を付
与するためにはＰ，Si等の強化元素を添加する必
要があるが、これらの元素は著しく脆化を促進す
るために、二次加工割れを起こし易いことが阻害
要因であつた。Ｂを添加して、二次加工性を改善
することを試みれば、材質が著しく劣化するとの
欠点も同時に有していたものである。
第６表に示す如く、従来のTi，Nb単独添加鋼
にＢを添加すると、材質が著しく劣化すると共
に、微量のＢでは二次加工性改善効果も非常に小
さい。本発明鋼は微量のＢ添加量で、二次加工性
は著しく優れたものとなり、材質の観点でも、Ｂ
添加、Ｐ，Si，Mnの添加の悪影響がない。
従つて本発明鋼は、強度の高い高強度鋼板や、
二次加工性を起こし易い厚手鋼板の製造に関して
も極めて有利なものである。[Table] Conventionally, high-strength steel sheets with high r-values have TS=40
Kgf/mm class 2 was the limit. In order to provide further strength, it is necessary to add reinforcing elements such as P and Si, but these elements significantly promote embrittlement and are likely to cause secondary processing cracks, which is an inhibiting factor. It was hot. At the same time, if an attempt was made to improve the secondary workability by adding B, the material quality would deteriorate significantly. As shown in Table 6, when B is added to the conventional steel with only Ti and Nb added, the material quality deteriorates significantly, and the effect of improving secondary workability is very small with a small amount of B. With the addition of a small amount of B, the steel of the present invention has extremely excellent secondary workability, and from the viewpoint of material quality, B
There is no adverse effect of addition of P, Si, or Mn. Therefore, the steel of the present invention is a high-strength steel plate with high strength,
It is also extremely advantageous for the production of thick steel plates that are susceptible to secondary workability.
第１，２，３図は本発明の実施例に於ける熱処
理サイクルを示す説明図である。
FIGS. 1, 2, and 3 are explanatory diagrams showing a heat treatment cycle in an embodiment of the present invention.
Claims (1)
を複合添加することを必須条件とし、Ｂは2ppm
以上10ppm未満の範囲内で添加し、Tiは48／14
〔Ｎ（％）−0.003％〕≦Ti（％）≦48／12C（％）＋
48／14N（％）の条件を満たす範囲内で含有し、
Nbは7C（％）＞Nb＞2C（％）で、かつ0.003％≦
Nb（％）＜0.035％を満たす範囲内の含有量で、か
つ0.017％＜Ti（％）＋Nb（％）＜0.06％を満たす成
分の鋼を、熱延加熱温度1300℃以下の条件で熱間
圧延した後、脱スケール処理、冷間圧延後、再結
晶温度以上AC3点以下の温度で連続焼鈍を施した
後、650℃から450℃の間を平均50℃／sec超の冷
却速度で冷却することを特徴とする極めて優れた
二次加工性を有する超深絞り用冷延鋼板の製造方
法。[Claims] 1. C: 0.005% or less, Si: 0.8% or less, Mn: 1.0% or less, P: 0.1% or less, Al: 0.01 to 0.1%, N: 0.005% or less, and unavoidable impurities. and B, Ti, Nb
The essential condition is that B is added in combination, and B is 2ppm.
Added within the range of 10 ppm or more, and Ti is 48/14
[N (%) - 0.003%] ≦Ti (%) ≦48/12C (%) +
Contains within the range that satisfies the condition of 48/14N (%),
Nb is 7C (%) > Nb > 2C (%) and 0.003%≦
Steel with a content that satisfies Nb (%) < 0.035% and a composition that satisfies 0.017% < Ti (%) + Nb (%) < 0.06% is hot-rolled at a hot rolling heating temperature of 1300℃ or less. After rolling, descaling treatment, cold rolling, continuous annealing at a temperature above the recrystallization temperature and below AC 3 points, and then cooling between 650℃ and 450℃ at an average cooling rate of over 50℃/sec. A method for producing a cold-rolled steel sheet for ultra-deep drawing, which has extremely excellent secondary workability.
Priority Applications (1)
| Application Number | Priority Date | Filing Date | Title |
|---|---|---|---|
| JP6521883A JPS59193221A (en) | 1983-04-15 | 1983-04-15 | Rreparation of cold rolled steel plate used in ultra-deep drawing having extremely excellent secondary processability |
Applications Claiming Priority (1)
| Application Number | Priority Date | Filing Date | Title |
|---|---|---|---|
| JP6521883A JPS59193221A (en) | 1983-04-15 | 1983-04-15 | Rreparation of cold rolled steel plate used in ultra-deep drawing having extremely excellent secondary processability |
Publications (2)
| Publication Number | Publication Date |
|---|---|
| JPS59193221A JPS59193221A (en) | 1984-11-01 |
| JPH0128817B2 true JPH0128817B2 (en) | 1989-06-06 |
Family
ID=13280547
Family Applications (1)
| Application Number | Title | Priority Date | Filing Date |
|---|---|---|---|
| JP6521883A Granted JPS59193221A (en) | 1983-04-15 | 1983-04-15 | Rreparation of cold rolled steel plate used in ultra-deep drawing having extremely excellent secondary processability |
Country Status (1)
| Country | Link |
|---|---|
| JP (1) | JPS59193221A (en) |
Families Citing this family (15)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| JPS61246344A (en) * | 1985-04-22 | 1986-11-01 | Kawasaki Steel Corp | Cold rolled steel sheet for super drawing excelling in resistance to secondary operation brittleness |
| JPS61276927A (en) * | 1985-05-31 | 1986-12-06 | Kawasaki Steel Corp | Production of cold rolled steel sheet having good deep drawability |
| JPH0647706B2 (en) * | 1986-08-04 | 1994-06-22 | 日新製鋼株式会社 | Cold-rolled steel sheet for deep drawing with excellent secondary work cracking resistance and method for producing the same |
| JPS63121623A (en) * | 1986-11-11 | 1988-05-25 | Kawasaki Steel Corp | Production of cold rolled steel sheet for deep drawing having excellent ridging resistance and chemical convertibility |
| US4889566A (en) * | 1987-06-18 | 1989-12-26 | Kawasaki Steel Corporation | Method for producing cold rolled steel sheets having improved spot weldability |
| JPH0699779B2 (en) * | 1987-09-14 | 1994-12-07 | 川崎製鉄株式会社 | Hot-rolled steel sheet for ultra deep drawing with good resistance to secondary processing brittleness |
| JPH0668129B2 (en) * | 1988-07-13 | 1994-08-31 | 川崎製鉄株式会社 | Method for producing hot rolled steel sheet with excellent deep drawability |
| JPH0670255B2 (en) * | 1988-11-21 | 1994-09-07 | 川崎製鉄株式会社 | Method for producing hot-rolled steel sheet for deep drawing with excellent surface properties |
| JPH02173242A (en) * | 1988-12-26 | 1990-07-04 | Kawasaki Steel Corp | High tensile cold rolled steel sheet for working and its production |
| JPH0832952B2 (en) * | 1989-12-28 | 1996-03-29 | 川崎製鉄株式会社 | Manufacturing method of cold-rolled steel sheet for press work with excellent chemical conversion treatability, weldability, punchability and slidability |
| US5279683A (en) * | 1990-06-20 | 1994-01-18 | Kawasaki Steel Corporation | Method of producing high-strength cold-rolled steel sheet suitable for working |
| JPH0756051B2 (en) * | 1990-06-20 | 1995-06-14 | 川崎製鉄株式会社 | Manufacturing method of high strength cold rolled steel sheet for processing |
| JP3111462B2 (en) * | 1990-07-19 | 2000-11-20 | 住友金属工業株式会社 | Manufacturing method of high-strength bake hardenable steel sheet |
| JPH0670267B2 (en) * | 1991-05-29 | 1994-09-07 | 株式会社神戸製鋼所 | Cold-rolled steel sheet for processing with excellent strength characteristics of welded parts |
| JP5050459B2 (en) * | 2006-09-14 | 2012-10-17 | Ｊｆｅスチール株式会社 | Cold rolled steel sheet for wound cores for automotive alternators with excellent magnetic properties and burr resistance |
Citations (4)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| JPS5819465A (en) * | 1981-07-27 | 1983-02-04 | Nippon Kokan Kk <Nkk> | Manufacture of galvanized steel plate with superior press formability |
| JPS5825436A (en) * | 1981-08-10 | 1983-02-15 | Kawasaki Steel Corp | Manufacture of deep drawing cold rolling steel plate having slow aging property and small anisotropy |
| JPS58110659A (en) * | 1981-12-25 | 1983-07-01 | Nippon Kokan Kk <Nkk> | Galvanized steel plate for deep drawing and its manufacture |
| JPS5974232A (en) * | 1982-10-20 | 1984-04-26 | Nippon Steel Corp | Production of bake hardenable galvanized steel sheet for ultradeep drawing having extremely outstanding secondary processability |
-
1983
- 1983-04-15 JP JP6521883A patent/JPS59193221A/en active Granted
Patent Citations (4)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| JPS5819465A (en) * | 1981-07-27 | 1983-02-04 | Nippon Kokan Kk <Nkk> | Manufacture of galvanized steel plate with superior press formability |
| JPS5825436A (en) * | 1981-08-10 | 1983-02-15 | Kawasaki Steel Corp | Manufacture of deep drawing cold rolling steel plate having slow aging property and small anisotropy |
| JPS58110659A (en) * | 1981-12-25 | 1983-07-01 | Nippon Kokan Kk <Nkk> | Galvanized steel plate for deep drawing and its manufacture |
| JPS5974232A (en) * | 1982-10-20 | 1984-04-26 | Nippon Steel Corp | Production of bake hardenable galvanized steel sheet for ultradeep drawing having extremely outstanding secondary processability |
Also Published As
| Publication number | Publication date |
|---|---|
| JPS59193221A (en) | 1984-11-01 |
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