JPS6111294B2 - - Google Patents
Info
- Publication number
- JPS6111294B2 JPS6111294B2 JP56116489A JP11648981A JPS6111294B2 JP S6111294 B2 JPS6111294 B2 JP S6111294B2 JP 56116489 A JP56116489 A JP 56116489A JP 11648981 A JP11648981 A JP 11648981A JP S6111294 B2 JPS6111294 B2 JP S6111294B2
- Authority
- JP
- Japan
- Prior art keywords
- steel
- less
- temperature
- present
- rolling
- Prior art date
- Legal status (The legal status is an assumption and is not a legal conclusion. Google has not performed a legal analysis and makes no representation as to the accuracy of the status listed.)
- Expired
Links
- 229910000831 Steel Inorganic materials 0.000 claims description 61
- 239000010959 steel Substances 0.000 claims description 61
- 238000000137 annealing Methods 0.000 claims description 21
- 238000005098 hot rolling Methods 0.000 claims description 18
- 238000005097 cold rolling Methods 0.000 claims description 9
- 238000004519 manufacturing process Methods 0.000 claims description 9
- 239000010960 cold rolled steel Substances 0.000 claims description 8
- 229910052757 nitrogen Inorganic materials 0.000 claims description 8
- 229910052698 phosphorus Inorganic materials 0.000 claims description 7
- 229910052710 silicon Inorganic materials 0.000 claims description 7
- 230000009466 transformation Effects 0.000 claims description 7
- 238000000034 method Methods 0.000 claims description 6
- 229910052758 niobium Inorganic materials 0.000 claims description 5
- 229910052719 titanium Inorganic materials 0.000 claims description 5
- 229910052720 vanadium Inorganic materials 0.000 claims description 5
- 229910052726 zirconium Inorganic materials 0.000 claims description 5
- 238000001953 recrystallisation Methods 0.000 claims description 4
- 238000005096 rolling process Methods 0.000 claims description 4
- 239000012535 impurity Substances 0.000 claims description 3
- 229910052748 manganese Inorganic materials 0.000 claims description 3
- 239000006104 solid solution Substances 0.000 description 18
- 230000000694 effects Effects 0.000 description 11
- 230000032683 aging Effects 0.000 description 10
- 229910000859 α-Fe Inorganic materials 0.000 description 9
- 239000000463 material Substances 0.000 description 6
- 239000000047 product Substances 0.000 description 6
- 229910052799 carbon Inorganic materials 0.000 description 5
- 230000006866 deterioration Effects 0.000 description 4
- 238000010438 heat treatment Methods 0.000 description 4
- 239000000654 additive Substances 0.000 description 3
- 230000000996 additive effect Effects 0.000 description 3
- 238000007872 degassing Methods 0.000 description 3
- 239000000203 mixture Substances 0.000 description 3
- 229910001566 austenite Inorganic materials 0.000 description 2
- 238000001816 cooling Methods 0.000 description 2
- 230000006911 nucleation Effects 0.000 description 2
- 238000010899 nucleation Methods 0.000 description 2
- 239000002244 precipitate Substances 0.000 description 2
- 238000005728 strengthening Methods 0.000 description 2
- 239000000126 substance Substances 0.000 description 2
- 238000004804 winding Methods 0.000 description 2
- 229910000975 Carbon steel Inorganic materials 0.000 description 1
- 230000003679 aging effect Effects 0.000 description 1
- 230000003712 anti-aging effect Effects 0.000 description 1
- 230000000052 comparative effect Effects 0.000 description 1
- 239000002131 composite material Substances 0.000 description 1
- 238000009749 continuous casting Methods 0.000 description 1
- 238000007796 conventional method Methods 0.000 description 1
- 238000005336 cracking Methods 0.000 description 1
- 230000007423 decrease Effects 0.000 description 1
- 230000002542 deteriorative effect Effects 0.000 description 1
- 238000005265 energy consumption Methods 0.000 description 1
- 238000005516 engineering process Methods 0.000 description 1
- 230000009422 growth inhibiting effect Effects 0.000 description 1
- 229910052738 indium Inorganic materials 0.000 description 1
- 238000002844 melting Methods 0.000 description 1
- 230000008018 melting Effects 0.000 description 1
- 150000001247 metal acetylides Chemical class 0.000 description 1
- 150000004767 nitrides Chemical class 0.000 description 1
- 238000005554 pickling Methods 0.000 description 1
- 238000011084 recovery Methods 0.000 description 1
- 230000003014 reinforcing effect Effects 0.000 description 1
- 239000000243 solution Substances 0.000 description 1
- 238000011426 transformation method 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
本発明は連続焼鈍によるプレス成形性に優れた
加工用高強度冷延鋼板の製造方法に関する。
従来、連続焼鈍法を利用した高強度冷延鋼板の
製造方法として種々の提案がなされており、その
一例としては、Ｐ添加鋼に代表されるような固溶
体強化を利用する製造法、或は変態組織を利用し
て所謂複合組織鋼とする製造方法等がある。しか
しながら、このような従来法によつて得られる高
強度冷延鋼板は総じてプレス成形性、とりわけ深
絞り性に劣るという難点がある。もつとも、この
ような加工性に関しては、例えばＰ添加鋼を箱焼
鈍して製造した鋼板は高強度にもかかわらず、絞
り用軟質冷延鋼板に近いプレス成形性を得ること
ができるが、ここで利用される箱焼鈍は能率面や
消費エネルギー等の面で連続焼鈍に較べはるかに
劣つており、高水準の生産性を期待することはで
きない。連続焼鈍材の深絞り性改善については、
熱延段階における高温巻取りが有効であることは
知られているが、反面、熱延高温巻取材は粗大化
した炭化物の影響により延性・張り出し性が劣化
するという問題があり、従来この点に関する十分
な解決は与えられていない。なお、特殊な例とし
て、脱ガス極低Ｃ―Ti添加鋼などの所謂I.F.鋼を
ベースとして、これにSi，Mn，Ｐ等の固溶強化
元素を添加する方法もあるが、この方法では固溶
〓〓〓〓
Ｃ，Ｎを完全に固定するためTi等の特殊炭窒化
物形成元素を多量（固溶Ｃ，Ｎを完全に固定する
のに必要な量の数倍程度）に添加する必要があ
り、このためコスト高になるという問題があり、
同時にかかる技術では前記したように固溶Ｃ，Ｎ
を完全に固定してしまうため、焼付硬化性が全く
期待できないという難点がある。
本発明は以上のような従来の問題点を解消すべ
く創案されたもので、TS＝35〜45Kg/mm2、El37
％及び1.4程度のプレス成形性に優れた高強
度冷延鋼板を高能率に、しかも低コストで製造す
ることができる方法を提供せんとするものであ
り、その特徴は、Ｃ：0.010％以下、Mn：0.05〜
0.30％、Si：0.01〜0.70％、Ｐ：0.01〜0.15％、
Ｓ：0.020％以下、SolAl：0.070％以下、Ｎ：
0.00050％以下、Ｂ：0.0010〜0.0050％、残部Fe
及び不可避的不純物からなる鋼を、熱間圧延段階
で仕上温度850〜900℃で圧延した後巻取り、次い
で冷間圧延後再結晶温度以上A3変態点以下で連
続焼鈍することにある。また他の特徴は上記成分
系に加え、Ti，Nb，Zr，Ｖのうち１種又は２種
以上を合計で0.001〜0.100％含有せしめることに
ある。以上により特殊元素などを多量に添加する
ことなく、高強度でしかも深絞り性等のプレス成
形性に優れた冷延鋼板の高能率且つ低コストでの
製造が容易に可能である。
本発明による化学成分は次の如き範囲において
調整される。
Ｃ：0.010％以下
Mn：0.05〜0.30％
Si：0.01〜0.70％
Ｐ：0.01〜0.15％
Ｓ：0.020％以下
SolAl：0.070％以下
Ｎ：0.0050％以下
Ｂ：0.0010〜0.0050％
また本発明のものは、上記のような基本成分に
対し、更にTi，Nb，Zr，Ｖのうち１種又は２種
以上を合計で0.001〜0.100％添加することができ
る。
本発明において、上記のように成分範囲を限定
した理由について説明すると以下の通りである。
Ｃは、脱ガス処理にて0.010％以下とする。Ｃ
は低い程好ましく深絞り性及び耐時効性が向上す
る。Ｃは0.005％未満が好ましい範囲ではある
が、現状の脱ガス設備能力から0.010％以下と規
定した。
Mnは、0.30％を超えると深絞り性の劣化が著
しい。Mnは深絞り性改善の見地からは低い程好
ましいが、表面性状や熱間脆性の問題を考慮して
下限を0.05％とする。
SiとＰは、鋼板の強度レベル調整を目的とした
強化元素として適量添加するものである。しかし
Siが0.70％を超えると表面性状の劣化を招くので
好ましくなく、またＰも0.15％を超えるとスポツ
ト溶接性の低下を招くので好ましくない。なお、
Si，Ｐともに0.01％以下とすることは製鋼上難し
く、またコストも上昇することから下限をこのよ
うに規定した。
Ｓは、延性を劣化させるので低い方が望まし
く、0.020％をその上限とする。
SolAlは、脱酸を図るため、またＢ／Ｎ＜１の
場合にＢで固定しきれない固溶ＮをAlNとして固
定するために添加するもので、0.070％を上限と
する。これ以上の添加はコスト高となり好ましく
ない。
Ｎは、必然的に混入するものであるが、低い程
好ましく0.0050％を上限とする。Ｎ量は後述する
Ｂ添加量とも密接に関係するが、Ｎが0.0050％を
超えると多量の窒化物が生成し、これにより焼鈍
時のフエライト粒の成長が阻害されて加工性が劣
化する。またＮが多いと固溶Ｎ固定のための添加
元素量が増すためコスト的にも不利となる。
Ｂは、本発明の最も重要な添加元素である。Ｂ
はＮとの親和力の強い元素であり、Ｎと結合して
BNを形成し、耐時効性を改善する効果があるこ
とは既に知られている。本発明者等は、このＢの
微量添加鋼と特定の熱間圧延仕上温度とを組み合
せることにより、連続焼鈍の如き急速加熱焼鈍に
おいても、深絞り性の優れた鋼板が得られること
を知見したものである。Ｂ添加及び特定の熱延仕
上温度の２条件の組み合せにより得られる上記効
果についての詳細は後述するが、熱延のオーステ
ナイト粒径の調整作用に基づくものと考えられ
る。本発明で規定するＢ量は0.0010〜0.0050％で
ある。Ｎ固定を目的としてＢを単独添加する場合
は、Ｂ／Ｎ１の条件を満たす必要があるが、逆
にこの比が大きくなり過ぎると固溶Ｂが残存して
〓〓〓〓
製品のプレス成形性に悪影響を及ぼす。本発明で
は熱延の粒調整効果をひいては冷延焼鈍後の深絞
り性改善効果を得るためＢを添加するものであ
り、Ｂ／Ｎが当量である必要はない。固溶Ｎの固
定に関してはSolAl，Ti等の添加でその目的が十
分達せられる。ただ、固溶Ｂによるプレス成形性
の劣化は避ける必要があり、この意味でＢの上限
をＮの上限との関係で0.0050％とした。またＢが
0.0050％を超えるとスラブのエツジ割れを生じ易
いという問題もある。Ｂの下限は0.0010％であ
り、これ以下ではＢ添加の効果が得られない。
Ti，Nb，Zr，Ｖについては、これらのものは
炭窒化物形成元素であり、これらを単独或は複合
添加して固溶Ｎの完全固定と固溶Ｃの一部又は全
量固定を図る。これらの元素を十分添加すれば、
固溶Ｃ，Ｎが完全に固定された所謂I.F.鋼とな
り、焼鈍後の製品は非時効性となる。しかし、こ
の非時効性を目的とした過去の例では、添加元素
量が固溶Ｃ，Ｎを完全に固定するのに必要な量の
数倍程度必要となり、コスト面で非常に不利とな
ることは前述した通りである。そこで本発明はコ
スト面に主眼を置き、固溶Ｎは完全固定するが、
固溶Ｃは一部ないし全量固定するのに必要な最低
限度の量を添加することとし、上限を単独又は複
合添加の合計で0.100％とした。また、合計で
0.001％を下回ると、その添加効果が得られず、
このため下限をこのように規定した。
このような成分からなる鋼の製造条件として、
本発明では熱延段階において、仕上温度850〜900
℃の範囲で仕上圧延を行うもので、これが本発明
の大きな特徴の１つであり、この熱処理と上記成
分系、特にＢとの組み合せにより好適なプレス成
形性が得られる。これは前述した如く熱延のオー
ステナイト粒径の調整作用に基づくものと思われ
る。第１図にＢ添加材（Ｂ：0.0035％、巻取温
度：660℃）の熱延仕上温度と熱延板フエライト
粒径との関係を示す。この第１図に示されるよう
に、熱延仕上温度が900℃超となると急激に熱延
板のフエライト粒径が大きくなる。この原因は明
確ではないが、固溶Ｂが変態時の核発生頻度を低
下させること、及び極低Ｃ鋼のためにフエライト
粒径自体が大きいこと等が原因であると考えられ
る。仕上温度が850〜900℃（図中斜線部分）では
適正なフエライト粒径を呈するが、これは温度の
低下によりBNが析出して固溶Ｂが減少し、上記
したような変態時の核発生頻度低下作用がなくな
ること、及び析出したBN等の析出物により粒成
長阻害作用を生ずることが原因であると推定され
る。また仕上温度が850℃未満では低温仕上層が
現出しはじめるために平均フエライト粒径は大き
くなる。第２図は第１図と同様Ｂ添加材（Ｂ：
0.0035％、巻取温度：660℃、連続焼鈍加熱温
度：850℃）の熱延仕上温度と冷延連続焼鈍後の
深絞り性（値）との関係を示すものであるが、
これからも判るように熱延板のフエライト粒径が
大きくなり過ぎると、冷延焼鈍後の製品の深絞り
性は適正粒径（図中斜線部分）のものに較べ、か
なり劣つたものとなつており、これは仕上温度が
900℃超では、冷延前粒径が極めて大きいこと
が、また仕上温度が850℃未満では低温仕上組織
が出現してしまうことが原因であると考えられ
る。
以上のような仕上温度で仕上圧延された鋼板
は、コイルに巻取られるが、この巻取温度に関し
ては特に限定はされない。即ち、550〜720℃程度
の低温ないし高温巻取りが行われる。次いで鋼板
は脱スケール処理後、圧延率60〜90％の通常の冷
間圧延が行われ、さらに連続焼鈍される。この連
続焼鈍温度は再結晶温度以上、A3変態点以下と
する。この温度範囲内では高温側ほどフエライト
粒成長が進み深絞り性に優れた製品が得られる利
点がある。しかし、A3変態点を超えると集合組
織がランダム化して深絞り性が劣化する。本発明
は極低Ｃ鋼をその対象としているため、その効果
は焼鈍後の冷却条件には依存せず、従つて冷却条
件は特に規定しない。
また本発明は以上の成分範囲及び熱処理で十分
その目的とする性質、つまりTS＝35〜45Kg/mm2，
El37％，1.4程度を得ることができ、従つ
て低炭素鋼で通常必要とされる過時効処理が必要
でなく、これが本発明の大きな特徴でもある。
実施例
第１表に示す化学成分の鋼を溶製し、連続鋳造
でスラブとした。表中、Ｂ鋼〜Ｆ鋼、Ｊ鋼〜Ｍ鋼
が本発明鋼、Ａ鋼及びＧ鋼〜Ｉ鋼が比較鋼であ
る。
〓〓〓〓
The present invention relates to a method for manufacturing a high-strength cold-rolled steel sheet for processing with excellent press formability by continuous annealing. Conventionally, various proposals have been made as methods for manufacturing high-strength cold-rolled steel sheets using continuous annealing. Examples include manufacturing methods that utilize solid solution strengthening, as typified by P-added steel, or transformation methods. There are manufacturing methods that utilize the structure to produce so-called composite structure steel. However, high-strength cold-rolled steel sheets obtained by such conventional methods generally have poor press formability, particularly poor deep drawability. However, regarding such workability, for example, a steel plate manufactured by box annealing P-added steel can obtain press formability close to that of a soft cold-rolled steel plate for drawing, despite its high strength. The box annealing used is far inferior to continuous annealing in terms of efficiency and energy consumption, and a high level of productivity cannot be expected. Regarding the improvement of deep drawability of continuously annealed materials,
It is known that high-temperature coiling in the hot-rolling stage is effective, but on the other hand, hot-rolled high-temperature coiled material has the problem of deteriorating ductility and stretchability due to the influence of coarse carbides. No adequate solution has been given. As a special example, there is a method in which solid solution strengthening elements such as Si, Mn, and P are added to a so-called IF steel such as degassing ultra-low C-Ti added steel, but this method Melting〓〓〓〓
In order to completely fix C and N, it is necessary to add a large amount of special carbonitride forming elements such as Ti (several times the amount required to completely fix solute C and N). There is a problem of high cost,
At the same time, in this technology, as mentioned above, solid solution C, N
Since it is completely fixed, there is a drawback that no bake hardenability can be expected. The present invention was devised to solve the above-mentioned conventional problems.TS=35~45Kg/ mm2 , El37
The purpose of the present invention is to provide a method for manufacturing high-strength cold-rolled steel sheets with excellent press formability of approximately 1.4% and 1.4% with high efficiency and at low cost. Mn: 0.05~
0.30%, Si: 0.01-0.70%, P: 0.01-0.15%,
S: 0.020% or less, SolAl: 0.070% or less, N:
0.00050% or less, B: 0.0010 to 0.0050%, balance Fe
and unavoidable impurities, the steel is rolled at a finishing temperature of 850 to 900°C in the hot rolling stage, then coiled, and then continuously annealed at a temperature above the recrystallization temperature and below the A3 transformation point after cold rolling. Another feature is that in addition to the above-mentioned component system, one or more of Ti, Nb, Zr, and V are contained in a total of 0.001 to 0.100%. As described above, it is possible to easily produce a cold-rolled steel sheet with high strength and excellent press formability such as deep drawability at high efficiency and at low cost without adding large amounts of special elements. The chemical components according to the present invention are adjusted within the following ranges. C: 0.010% or less Mn: 0.05-0.30% Si: 0.01-0.70% P: 0.01-0.15% S: 0.020% or less SolAl: 0.070% or less N: 0.0050% or less B: 0.0010-0.0050% In addition to the above basic components, one or more of Ti, Nb, Zr, and V can be added in a total amount of 0.001 to 0.100%. In the present invention, the reason why the component range is limited as described above is as follows. C is reduced to 0.010% or less by degassing treatment. C
The lower the value, the better the deep drawability and aging resistance. Although C is preferably less than 0.005%, it is specified to be 0.010% or less based on the current degassing equipment capacity. When Mn exceeds 0.30%, deep drawability deteriorates significantly. The lower the Mn content, the better from the standpoint of improving deep drawability, but the lower limit is set at 0.05% in consideration of surface texture and hot embrittlement problems. Si and P are added in appropriate amounts as reinforcing elements for the purpose of adjusting the strength level of the steel plate. but
If Si exceeds 0.70%, this is undesirable, as it causes deterioration of the surface properties, and if P exceeds 0.15%, it also causes deterioration in spot weldability, which is undesirable. In addition,
Since it is difficult to make both Si and P 0.01% or less in terms of steel manufacturing, and the cost also increases, the lower limit was defined in this way. Since S deteriorates ductility, it is desirable to have a low content, and the upper limit is set at 0.020%. SolAl is added to deoxidize and to fix solid solution N that cannot be fixed by B when B/N<1 as AlN, and has an upper limit of 0.070%. Addition of more than this amount increases the cost and is not preferable. Although N is inevitably mixed, the lower the content, the more preferable the upper limit is 0.0050%. The amount of N is closely related to the amount of B added, which will be described later, but when N exceeds 0.0050%, a large amount of nitrides are generated, which inhibits the growth of ferrite grains during annealing and deteriorates workability. Furthermore, if the amount of N is large, the amount of added elements for fixing N in solid solution increases, which is disadvantageous in terms of cost. B is the most important additive element of the present invention. B
is an element that has a strong affinity with N, and when combined with N,
It is already known that it forms BN and has the effect of improving aging resistance. The present inventors have discovered that by combining this steel with a small amount of B added and a specific hot rolling finishing temperature, a steel plate with excellent deep drawability can be obtained even in rapid heating annealing such as continuous annealing. This is what I did. The above effect obtained by the combination of the two conditions of B addition and a specific hot rolling finishing temperature will be described in detail later, but it is thought to be based on the effect of adjusting the austenite grain size of the hot rolling. The amount of B specified in the present invention is 0.0010 to 0.0050%. When adding B alone for the purpose of N fixation, it is necessary to satisfy the condition of B/N1, but conversely, if this ratio becomes too large, solid solution B will remain.
It has a negative effect on the press formability of the product. In the present invention, B is added in order to obtain the effect of grain adjustment during hot rolling and the effect of improving deep drawability after cold rolling annealing, and B/N does not need to be equivalent. Regarding the fixation of solid solution N, the purpose can be sufficiently achieved by adding SolAl, Ti, etc. However, it is necessary to avoid deterioration of press formability due to solid solution B, and in this sense, the upper limit of B was set at 0.0050% in relation to the upper limit of N. Also B
If it exceeds 0.0050%, there is also the problem that edge cracking of the slab is likely to occur. The lower limit of B is 0.0010%, and below this the effect of B addition cannot be obtained. Ti, Nb, Zr, and V are carbonitride-forming elements, and these are added singly or in combination to completely fix the solid solution N and to fix part or all of the solid solution C. If enough of these elements are added,
It becomes a so-called IF steel in which solid solution C and N are completely fixed, and the product after annealing becomes non-aging. However, in past examples aimed at achieving non-aging properties, the amount of added elements was required to be several times the amount required to completely fix the solid solution C and N, which was extremely disadvantageous in terms of cost. is as described above. Therefore, the present invention focuses on cost, and completely fixes the solid solution N.
Solid solution C was added in the minimum amount necessary to fix part or all of it, and the upper limit was set at 0.100% for the total amount of C added alone or in combination. Also, in total
If it is less than 0.001%, the effect of the addition cannot be obtained,
For this reason, the lower limit was defined as follows. The manufacturing conditions for steel made of these components are as follows:
In the present invention, in the hot rolling stage, the finishing temperature is 850 to 900.
Finish rolling is carried out in the temperature range of 0.degree. C., which is one of the major features of the present invention, and the combination of this heat treatment and the above-mentioned component system, especially B, provides suitable press formability. This seems to be due to the effect of adjusting the austenite grain size during hot rolling as described above. FIG. 1 shows the relationship between the hot-rolling finishing temperature of the B additive material (B: 0.0035%, coiling temperature: 660°C) and the ferrite grain size of the hot-rolled sheet. As shown in FIG. 1, when the hot-rolling finishing temperature exceeds 900°C, the ferrite grain size of the hot-rolled sheet suddenly increases. Although the cause of this is not clear, it is thought that the solid solution B reduces the frequency of nucleation during transformation, and that the ferrite grain size itself is large due to the extremely low C steel. When the finishing temperature is 850 to 900℃ (the shaded area in the figure), the ferrite grain size exhibits an appropriate size, but this is because BN precipitates due to the temperature drop and solid solution B decreases, resulting in nucleation during transformation as described above. This is presumed to be due to the absence of the frequency-reducing effect and the fact that precipitates such as BN cause a grain growth inhibiting effect. Furthermore, when the finishing temperature is lower than 850°C, the average ferrite grain size becomes large because a low-temperature finishing layer begins to appear. Figure 2 shows the B additive material (B:
This shows the relationship between the hot rolling finishing temperature (0.0035%, coiling temperature: 660°C, continuous annealing heating temperature: 850°C) and deep drawability (value) after cold rolling continuous annealing.
As can be seen from this, when the ferrite grain size of the hot-rolled sheet becomes too large, the deep drawability of the product after cold-rolling and annealing becomes considerably inferior to that of the product with the appropriate grain size (the shaded area in the figure). This means that the finishing temperature is
This is thought to be due to the extremely large grain size before cold rolling at temperatures above 900°C, and the appearance of a low-temperature finished structure at finishing temperatures below 850°C. The steel plate finish-rolled at the finishing temperature as described above is wound into a coil, but there are no particular limitations on this winding temperature. That is, low-temperature to high-temperature winding of about 550 to 720° C. is performed. Next, the steel plate is descaled, then subjected to normal cold rolling at a rolling reduction of 60 to 90%, and then continuously annealed. The continuous annealing temperature is set to be higher than the recrystallization temperature and lower than the A3 transformation point. Within this temperature range, there is an advantage that the higher the temperature, the more ferrite grains grow and a product with excellent deep drawability can be obtained. However, when the A3 transformation point is exceeded, the texture becomes random and deep drawability deteriorates. Since the present invention targets ultra-low C steel, its effects do not depend on the cooling conditions after annealing, and therefore the cooling conditions are not particularly defined. In addition, the above component range and heat treatment are sufficient for the present invention to obtain the desired properties, that is, TS = 35 to 45 Kg/mm 2 ,
El37%, about 1.4 can be obtained, and therefore there is no need for over-aging treatment, which is normally required for low carbon steels, which is also a major feature of the present invention. Example Steel having the chemical composition shown in Table 1 was melted and made into a slab by continuous casting. In the table, B steel to F steel and J steel to M steel are the invention steels, and A steel and G steel to I steel are comparative steels. 〓〓〓〓
【表】
上記スラブは第２表に示す種々の熱延条件で板
厚2.8mmに熱間圧延した。次いで酸洗で脱スケー
ル後、板厚0.8mm（圧延率71.4％）に冷間圧延
し、その後連続焼鈍を行つた。その機械的性質を
第２表に合せて示す。なお、第２表中、Ｂ(1)鋼〜
Ｂ(5)鋼はＢ鋼を種々の熱延条件及び焼鈍条件で処
理したものであつて、このうち、Ｂ(2)鋼、Ｂ(4)
鋼、Ｂ(5)鋼が本発明鋼である。[Table] The above slabs were hot rolled to a thickness of 2.8 mm under various hot rolling conditions shown in Table 2. Next, after descaling by pickling, it was cold rolled to a plate thickness of 0.8 mm (rolling ratio 71.4%), and then continuously annealed. Its mechanical properties are also shown in Table 2. In addition, in Table 2, B(1) steel ~
B(5) steel is obtained by processing B steel under various hot rolling conditions and annealing conditions, and among these, B(2) steel, B(4) steel
Steel, B(5) steel is the steel of the present invention.
【表】
〓〓〓〓
[Table] 〓〓〓〓
【表】
第２表において、Ａ鋼はＢが添加されておら
ず、Ｇ鋼はＢが添加されていないのに加えＣの含
有量が高く、Ｈ鋼も同じくＣの含有量が高く、ま
たＩ鋼はMnの含有量が高く、いずれも本発明の
範囲外であつて、このため値が低い。
一方、Ｃ鋼、Ｄ鋼、Ｅ鋼はいずれも本発明鋼で
あつて、TS，El，値ともに良好な性質を得ら
れている。またＢ鋼のうち、Ｂ(2)鋼、Ｂ(4)鋼及び
Ｂ(5)鋼は本発明鋼であり、良好な性質を示してい
るが、Ｂ(1)鋼及びＢ(3)鋼は熱延段階での仕上温度
が本発明の範囲外（900℃、830℃）にあり、この
ため値が低い。Ｆ鋼、Ｊ鋼、Ｋ鋼、Ｌ鋼、Ｍ鋼
は、本願第２の発明に係る実施例であつて、炭窒
化物形成元素たるTi，Nb，Zr，Ｖをそれぞれ単
独または複合で適量添加したものであり、これに
よればTS，El，値ともに良好な性質を示して
いるのに加え、特に時効指数に関し他の本発明鋼
が3.0〜3.5Kg/mm2の値を示しているのに対し、0.8
〜2.3Kg/mm2と比較的低い値となつており、他の本
発明鋼に較べて適度な非時効性を得ていることが
判る。ただ、本発明鋼を全体としてみれば、耐時
効性を評価する時効指数は１〜3.5Kg/mm2程度のレ
ベルにあり、この状態は少量の固溶Ｃが存在し完
全非時効ではないが、時効による降伏点伸びの回
復及び材質の時効劣化量は実用上ほとんど問題と
ならない程度で遅時効性であると言える。また逆
に、このように少量の固溶Ｃが存在することは焼
付硬化性を有することを意味し、自動車部品等の
製品化後の高い降伏強度が期待できる。
以上述べたように本発明によれば、高強度であ
りながら深絞り性等のプレス成形性に優れた鋼板
を連続焼鈍により高能率に、しかも多量の特殊元
素を添加することなく低コストで製造することが
でき、実用的価値の極めて高い発明であると言う
ことができる。[Table] In Table 2, steel A has no B added, steel G has no B added but has a high C content, steel H also has a high C content, and I steel has a high Mn content, both of which are outside the scope of the present invention, and therefore the value is low. On the other hand, C steel, D steel, and E steel are all steels of the present invention, and have good properties in terms of TS, El, and values. Among B steels, B(2) steel, B(4) steel and B(5) steel are steels of the present invention and exhibit good properties, but B(1) steel and B(3) steel The finishing temperature in the hot rolling stage is outside the range of the present invention (900°C, 830°C), and therefore the value is low. F steel, J steel, K steel, L steel, and M steel are examples according to the second invention of the present application, in which appropriate amounts of carbonitride forming elements Ti, Nb, Zr, and V are added individually or in combination. According to this, in addition to showing good properties in both TS and El values, in particular, regarding the aging index, other steels of the present invention show values of 3.0 to 3.5 Kg/ mm2 . against 0.8
The value is relatively low at ~2.3 Kg/mm 2 , and it can be seen that a suitable anti-aging property has been obtained compared to other steels of the present invention. However, when looking at the steel of the present invention as a whole, the aging index for evaluating aging resistance is at a level of about 1 to 3.5 Kg/ mm2 , and although this state is not completely non-aging due to the presence of a small amount of solid solution C, It can be said that the recovery of yield point elongation due to aging and the amount of aging deterioration of the material are of a slow aging nature to the extent that there is almost no problem in practical use. Conversely, the presence of such a small amount of solid solution C means that the material has bake hardenability, and high yield strength can be expected after it is manufactured into products such as automobile parts. As described above, according to the present invention, a steel plate with high strength and excellent press formability such as deep drawability can be produced with high efficiency by continuous annealing and at low cost without adding large amounts of special elements. This invention can be said to have extremely high practical value.
第１図は本発明の成分範囲による鋼板の熱延仕
上温度と熱延板フエライト粒度との関係を示すも
のである。第２図は同じく本発明の成分範囲によ
る鋼板の熱延仕上温度と冷延連続焼鈍後のとの
関係を示すものである。
〓〓〓〓
FIG. 1 shows the relationship between the hot-rolling finishing temperature of a steel sheet and the ferrite grain size of the hot-rolled sheet according to the composition range of the present invention. FIG. 2 similarly shows the relationship between the hot rolling finishing temperature of a steel sheet according to the composition range of the present invention and that after cold rolling and continuous annealing. 〓〓〓〓
Claims (1)
0.01〜0.70％、Ｐ：0.01〜0.15％、Ｓ：0.020％以
下、SolAl：0.070％以下、Ｎ：0.0050％以下、
Ｂ：0.0010〜0.0050％、残部Fe及び不可避的不純
物からなる鋼を、熱間圧延段階で仕上温度850〜
900℃で圧延した後巻取り、次いで冷間圧延後再
結晶温度以上A3変態点以下で連続焼鈍すること
を特徴とする連続焼鈍による加工用高強度冷延鋼
板の製造方法。 ２ Ｃ：0.010％以下、Mn：0.05〜0.30％、Si：
0.01〜0.70％、Ｐ：0.01〜0.15％、Ｓ：0.020％以
下、SolAl：0.070％以下、Ｎ：0.0050％以下、
Ｂ：0.0010〜0.0050％、Ti、Nb、Zr、Ｖのうち１
種又は２種以上が合計で0.001〜0.100％、残部Fe
及び不可避的不純物からなる鋼を、熱間圧延段階
で仕上温度850〜900℃で圧延した後巻取り、次い
で冷間圧延後再結晶温度以上A3変態点以下で連
続焼鈍することを特徴とする連続焼鈍による加工
用高強度冷延鋼板の製造方法。[Claims] 1 C: 0.010% or less, Mn: 0.05 to 0.30%, Si:
0.01-0.70%, P: 0.01-0.15%, S: 0.020% or less, SolAl: 0.070% or less, N: 0.0050% or less,
B: Steel consisting of 0.0010 to 0.0050%, balance Fe and unavoidable impurities is heated to a finishing temperature of 850 to 850 in the hot rolling stage.
A method for producing a high-strength cold-rolled steel sheet for processing by continuous annealing, which comprises rolling at 900°C, then coiling, and then continuous annealing at a temperature above the recrystallization temperature and below the A3 transformation point after cold rolling. 2 C: 0.010% or less, Mn: 0.05-0.30%, Si:
0.01-0.70%, P: 0.01-0.15%, S: 0.020% or less, SolAl: 0.070% or less, N: 0.0050% or less,
B: 0.0010-0.0050%, 1 of Ti, Nb, Zr, V
Species or two or more species total 0.001-0.100%, balance Fe
and unavoidable impurities, is rolled at a finishing temperature of 850 to 900°C in the hot rolling stage, then coiled, and then continuously annealed at a temperature above the recrystallization temperature and below the A3 transformation point after cold rolling. A method for manufacturing high-strength cold-rolled steel sheets for processing by continuous annealing.
Priority Applications (1)
| Application Number | Priority Date | Filing Date | Title |
|---|---|---|---|
| JP11648981A JPS5819442A (en) | 1981-07-27 | 1981-07-27 | Manufacture of high strength cold rolled steel plate for working by continuous annealing |
Applications Claiming Priority (1)
| Application Number | Priority Date | Filing Date | Title |
|---|---|---|---|
| JP11648981A JPS5819442A (en) | 1981-07-27 | 1981-07-27 | Manufacture of high strength cold rolled steel plate for working by continuous annealing |
Publications (2)
| Publication Number | Publication Date |
|---|---|
| JPS5819442A JPS5819442A (en) | 1983-02-04 |
| JPS6111294B2 true JPS6111294B2 (en) | 1986-04-02 |
Family
ID=14688383
Family Applications (1)
| Application Number | Title | Priority Date | Filing Date |
|---|---|---|---|
| JP11648981A Granted JPS5819442A (en) | 1981-07-27 | 1981-07-27 | Manufacture of high strength cold rolled steel plate for working by continuous annealing |
Country Status (1)
| Country | Link |
|---|---|
| JP (1) | JPS5819442A (en) |
Families Citing this family (7)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| JPH0768634B2 (en) * | 1985-07-03 | 1995-07-26 | 新日本製鐵株式会社 | Zinc-based plated steel sheet with excellent corrosion resistance, coating performance and workability |
| JPS6280251A (en) * | 1985-10-04 | 1987-04-13 | Kawasaki Steel Corp | Low-carbon steel sheet for working excellent in ridging resistance |
| JPH0830245B2 (en) * | 1987-03-23 | 1996-03-27 | 住友金属工業株式会社 | High-strength cold-rolled steel sheet for processing and its manufacturing method |
| JPS6425945A (en) * | 1987-07-20 | 1989-01-27 | Sumitomo Metal Ind | Cold rolled steel plate for drawing having excellent elongation and its production |
| 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 |
| KR20020040433A (en) * | 2000-11-24 | 2002-05-30 | 이구택 | method of manufacturing a cold-rolled steel with good formability |
| JP6052145B2 (en) | 2013-11-28 | 2016-12-27 | Ｊｆｅスチール株式会社 | Bake-hardening hot-dip galvanized steel sheet |
Citations (1)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| JPS616133A (en) * | 1984-06-16 | 1986-01-11 | Ishizuka Glass Ltd | Forehearth for melting crystal glass |
-
1981
- 1981-07-27 JP JP11648981A patent/JPS5819442A/en active Granted
Patent Citations (1)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| JPS616133A (en) * | 1984-06-16 | 1986-01-11 | Ishizuka Glass Ltd | Forehearth for melting crystal glass |
Also Published As
| Publication number | Publication date |
|---|---|
| JPS5819442A (en) | 1983-02-04 |
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