JPS6132375B2 - - Google Patents

Info

Publication number
JPS6132375B2
JPS6132375B2 JP20617381A JP20617381A JPS6132375B2 JP S6132375 B2 JPS6132375 B2 JP S6132375B2 JP 20617381 A JP20617381 A JP 20617381A JP 20617381 A JP20617381 A JP 20617381A JP S6132375 B2 JPS6132375 B2 JP S6132375B2
Authority
JP
Japan
Prior art keywords
steel
amount
added
temperature
killed
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
Application number
JP20617381A
Other languages
Japanese (ja)
Other versions
JPS58107414A (en
Inventor
Yoshikuni Tokunaga
Masato Yamada
Current Assignee (The listed assignees may be inaccurate. Google has not performed a legal analysis and makes no representation or warranty as to the accuracy of the list.)
Nippon Steel Corp
Original Assignee
Nippon Steel Corp
Priority date (The priority date 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 date listed.)
Filing date
Publication date
Application filed by Nippon Steel Corp filed Critical Nippon Steel Corp
Priority to JP20617381A priority Critical patent/JPS58107414A/en
Publication of JPS58107414A publication Critical patent/JPS58107414A/en
Publication of JPS6132375B2 publication Critical patent/JPS6132375B2/ja
Granted legal-status Critical Current

Links

Classifications

    • CCHEMISTRY; METALLURGY
    • C21METALLURGY OF IRON
    • C21DMODIFYING 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
    • C21D8/00Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/04Modifying 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 Sheet Steel (AREA)
  • Heat Treatment Of Steel (AREA)

Description

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

本発明は超深絞り用鋼板の製造方法に関するも
のである。 超深絞り用鋼板としては、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。例えばNbキルド鋼板は前記特許出願時
に比べて技術的進歩が著しく、上記特許公報開示
技術では経済的でないため現実性に乏しい。C<
500ppmの範囲にCが容易に低減できるようにな
つたことから、最近では上記特許公報の範囲より
も低い炭素量、Nb量で製造できることが、学術
文献として報告されている例もある。これらの鋼
板は極低炭素にすることを前提として、炭窒化物
形成能の強いTiあるいはNbを添加して、侵入型
元素C、Nのほとんどない鋼板にしているところ
から、連続焼鈍によつても箱焼鈍と同レベルの製
品を製造できる利点がある。 しかし、実際にこれらの鋼板を連続焼鈍で製造
してみると次のような欠点を持つている。Tiキ
ルド鋼板については、二次加工割れが発生し易い
ことである。特に材質のよいものを狙い、Tj添
加量がC、Nに対して当量以上の範囲である場合
には、P量が多くなると二次加工割れが発生し易
くなる。また、P添加した場合r値がかなり劣化
することも欠点である。さらに、連続焼鈍の一種
であるゼンジマー式連続溶融亜鉛めつきラインに
て、合金化亜鉛めつき鋼板を製造する際には、パ
ウダリング(合金化が過度に進行して、プレス成
形時にめつき層が剥脱する現象)が発生し易い欠
点もある。しかし、この鋼種は熱延巻取温度が、
通常の600〜650℃でも連続焼鈍で安定した材質の
ものが製造できる利点がある。 それに対して、Nbキルド鋼は熱延で高温巻取
(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(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と
すだけでなく、コイル端部は冷却速度が速いため
に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。 本発明者等は、Tiキルド鋼、Nbキルド鋼の持
つこれらの長所、短所を詳細に検討した結果、ま
ず上記の挙動の違いについて以下のように考え
た。Tiは極めて強い窒化物形成元素であり、熱
延の加熱炉中ですでにTiNが形成されている炭化
物は窒化物と比較すると析出温度は低いものの、
窒化物を析出サイトにして600〜650℃の巻取中に
もかなりのものが析出してくると考えられる。従
つて低い巻取温度でもかなりの析出物が、ホツト
コイル中で大きな析出物として析出しており、冷
延後の連続焼鈍時に析出してくるものはわずかで
ある。従つて、再結晶温度は極端には高くならな
いですんでおり、かつ材質はかなり均質になると
考えられる。しかし、炭窒化物としてC、Nがほ
とんど析出してしまつて粒界は清浄になり、そこ
に粒界脆化を起こすP等の不純物元素が偏析を起
こすことにより、二次加工割れが起こると考えら
れる。さらにTiキルド鋼では、合金化亜鉛めつ
き鋼板を製造する場合、鋼中Tiにより地鉄と溶
融亜鉛の合金化反応が促進されるために、過合金
化が進行しやすく、パウダリングが起こりやすい
と考えられる。 それに対してNbは、窒化物形成能がTiと比較
してかなり劣る。実際Nb添加鋼では、Nbは炭化
物を形成するが、窒素はAlNとして析出してい
る。AlNは低温巻取では形成され難く、巻取温度
を700℃以上にしないと熱延板中では形成され
ず、冷延後の連続焼鈍時に微細に析出し、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、熱延コイルの前後端部は
高温巻取をしても、冷却速度が速いために実際に
は低温巻取に近い材質にとどまる。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AlNの形成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と考えられる。 しかし、Nbは炭窒化物形成能がTiに比較して
劣るため、数ppmの炭素は粒界に偏析してお
り、これが粒界の結合エネルギーを高めるため
に、P含有量がかなり高くても二次加工割れは心
配ないものと考えられる。またNbは合金化亜鉛
めつき鋼板製造時には、地鉄と溶融亜鉛の合金化
反応をTiほど促進する傾向はみられないため、
Nbキルド鋼では、パウダリングはTiキルド鋼に
比較して起こり難い。 こうした考えに基づいて、コイル内材質の均質
性が良好でかつ二次加工割れの心配がなく、また
合金化亜鉛めつき鋼板製造時にはパウダリングの
発生し難い超深絞り性鋼板製造の考え方は以下の
ようなものになつた。即ち、NはAlNではなくTi
によつてTiNとして仕上げ熱延前に析出させ、ま
たCは〔Ti、Nb〕Cの複合炭化物として析出さ
せた鋼板が本発明の基本的な考え方である。 本発明鋼は後述するようにTiキルド鋼と比較
して、P添加をして高強度化した場合にr値が劣
化せず、かつ二次加工割れが極めて起こり難い点
で優れている。さらに合金化亜鉛めつき鋼板を製
造する場合には、パウダリングが発生し難いとい
う長所を有する。次にNbキルド鋼に比較して、
NをAlNではなくTiNとして固定しているため
に、高温巻取を実施せずとも低温巻取でもほぼ同
じレベルの材質を得ることができ、かつコイル長
方向及びコイル幅方向の材質が非常に均質である
点で優れている。これらの長所は従来のTiキル
ド鋼、Nbキルド鋼と比較して、それぞれの鋼板
が持つていない種々の優れた材質特性を有するこ
とを明確に示しており、本発明鋼は極めて有利な
ものである。 更にTi、Nbを複数添加した本発明鋼が、従来
のTiキルド鋼、Nbキルド鋼の性質からは全く予
想もできない独特の性質を持つているのは、r値
の異方性が極めて小さいことである。一般に、
Tiキルド鋼やNbキルド溝のr値は、圧延方向
(L方向)又は45゜方向が最も劣り、圧延直角方
向(C方向)が最も優れている。しかし本発明鋼
は、冷延率の大小によらず、L、C、45゜方向が
ほとんど等しいか、又は45゜方向がいくらか大き
いr値を持つている。こうした異方性を持つた高
r値鋼板は今までに全く知られておらず、学問的
にも興味深いが、実用上も大きな利点がある。特
に四角筒絞り(角部が45゜方向)の場合に、極め
て優れた成形性を示すことは容易に考えられる。
しかし円筒深絞り成形の場合でも、破断限界は最
小のr値で決まる場合が多い。また、例えば乾電
池の外筒のように、絞り成形後の板厚の均一性が
重要なものや、耳の発生を極力抑えたい用途に
は、異方性が少ないことは多大な利点がある。従
つて単にr値を3方向に平均値だけで見るだけで
なく、本発明鋼のような新鋼板の出現で、r値の
異方性についても巷間の関心が高まると考えられ
る。 このように本発明鋼は、Ti又はNbの単独添加
鋼と比較してあらゆる点で優れ、かつ全く予想も
できなかつた性質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との関係で決まり、少くともNを
固定するに足る量を添加すべきである。 即ち48/14〔N(%)−0.002%(20ppm)〕<Ti
である。 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、炭化物になつ
て二次加工性を劣化させたり、P添加時の材質を
悪くするので、多くともTi添加量はCとNとの
当量以下にすべきである。即ち、Ti(%)<〔48/
12C(%)+48/14N(%)〕になる。 一方NbはCとの関係で決まり、原子比でC量
の0.3倍以上、即ち Nb(%)>0.3×93/12C(%)=2.33(%)で、
かつ0.003%以上で0.025%未満の量添加すること
が望ましい。 Nb(%)/C(%)<2.33では〔Ti、Nb〕Cの
複合炭化物を形成しないで固溶Cが残存し、非時
効性の鋼板が得られないという問題があり、また
Nb0.025%では材質特性値は、従来のNbキルド
鋼に類似の挙動を示し、再結晶温度が上昇し、コ
イル前後端部の材質劣化が大きくなり、本発明の
原理からはずれることになる。 第1図、第2図はTi、Nb量の関係から本発明
の範囲を示したものである。 第1図はNb量を一定量(0.022%)に固定し、
Ti量を変化させた場合の材質特性値である。
C;0.005、Si;0.01、Mn;0.25、P;0.02、
S;0.01、Sol.Al;0.06、N;0.005(各%)の試
料について熱間圧延で720℃巻取を行ない、a;
コイル長方向中心部、b;コイル前後端部を示し
た。Ti量がNを固定するために不十分な量の場
合、即ち48/14〔N(%)−0.002%〕=0.010%>
Tiの場合には、コイル前後端部の材質劣化が大
きく、またr値の異方性は、Nbを微量添加した
極低炭素鋼のr値の特性を示しており、Ti、Nb
複合添加の効果が小さい。逆にTiを、C、Nに
対して当量以上添加した場合には、コイル前後端
部の材質劣化は非常に小さいが、r値の異方性は
Tiキルドのそれとほとんど等しいものとなる。
即ちTiを適当量添加した場合にのみ、Ti、Nb複
合添加による集合組織改善効果が現われ、極めて
異方性の小さい優れたr値特性を示すものであ
る。このような優れたr値の等方性は、Ti、Nb
の複合添加を必須とするもので、Ti、Nb単独添
加鋼では得られない集合組織によるものである。 又、Tiの添加量は第1図に示す如く、下限値
はTi>48/14〔N(%)−0.002〕から、0.010%と
なり、上限値はTi<〔4.00C(%)+3.43N(%)〕
から0.037%となる。 Tiが0.010%未満であるとコイル前後端部にお
いて、TiがNをTiNとして固定させる比率が少な
く、逆にAlNとして析出するNの比率が多くな
り、AlNの析出物サイズが極めて小さいので、コ
イル前後端部の材質を劣化させ、好ましくないも
のである。 Tiの量が0.037以上では鋼中のCとNの全量が
Tiによつて固定される。この場合には、Ti単独
添加鋼と同じ材質特性となり、一次加工性が劣化
したり、パウダリングが発生しやすくなり、好ま
しくないものである。 第2図はTi量をN固定するに十分な量だけ一
定量(0.02%)添加し、Nb添加量を変化させた
場合の材質特性を示したものである。試料の化学
成分は第1図とほぼ同一に水準のものである。
Nb量がC量に対して低すぎる場合には、第1図
と同様単なる極低炭素鋼に類似の特徴を示し、45
゜方向のr値が極めて小さく、異方性が非常に大
きい。さらにコイル前後端部の材質劣化が大き
く、非時効性が得られないのは第1図の場合と同
様である。Nb量が0.025%を超えると、r値の異
方性は同じ傾向であるが、rcの値が低くなる欠
点が現われ、コイル長前後端部の材質劣化が極め
て大きいというNbキルド鋼特有の特性に類似す
る。第2図からもTi、Nbの適当な複合添加によ
つてのみ、r値の異方性が極めて小さく、コイル
長方向の材質が均一な優れた特性が得られる。こ
のような特性は、Ti、Nb単独添加鋼では得られ
ず、Ti、Nbの複合添加が非須であることを示す
ものである。このようにNb単独添加に比べて巻
取温度に依存せず、高い安定した材質が出るこ
と、及びTi単独添加に比べて少ない合金元素添
加量でも高い材質が得られるのは次のような理由
による。 本発明範囲のように、TiとNbの添加量をバラ
ンスよく含んだ場合には〔Ti、Nb〕Cの複合析
出物が出来るが、これはTiC、NbCに比べて析出
開始温度が高く、大きな析出物として析出するの
で巻取温度が低くとも良好な再結晶挙動を示すも
のと思われる。これが結果的に等方的なr値につ
ながつていると考えられる。それに対してNb添
加鋼、又はNb0.025%添加した場合は生成する
のにNbCであり、巻取温度によつて析出状態が大
きく変化して、低温巻取の場合は、微細なNbCが
連続焼鈍時の再結晶温度を上げるために材質は大
きく劣化する。またTiの単独添加では、TiとC
+Nの原子比を1以上にしなければ材質は大きく
劣化するが、これはTi量を多くしなければTiCの
析出が熱延板中で十分に起らずに、連続焼鈍時に
微細に析出するために硬質でかつ延性が劣化する
ものと思われる。 以上詳述したようにNb<0.025%にして、Nbと
Tiを複合添加した場合に、延性に優れてかつ深
絞り性に優れた鋼板が得られるのである。 以上述べた如く、本発明鋼の成分範囲は、Ti
添加量は鋼中のN量に依存し、Ti(%)>48/14
〔N(%)−0.002%)かつTi(%)<〔48/12C
(%)+48/14N(%)〕を満たす量含有し、Nb添
加量は鋼中C量で決まり、Nb(%)>0.3×93/12
C(%)=2.33C(%)かつ0.003%≦Nb(%)<
0.025%を満たす量含有するというものである。 Ti、Nb以外の成分範囲はC:0.007%以下、
Si:0.8%以下、Mn:1.0%以下、P:0.1%以
下、Al:0.01〜0.1%、N:80ppm以下及び他の
不可避的不純物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、〔Ti、Nb〕Cの生
成量が増えるため、結晶粒の成長を阻害し、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は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ため
0.8%以下とする。特に合金化処理を実施しない
場合には0.3%以下が望ましい。 Mnは多量に加えるとr値の劣化が著しいため
1.0%を上限するが、高r値の観点から低い方が
望ましい。 本発明鋼は、Ti、Nbの複合添加により二次加
工割れを起こし難いが、Pが多量に含まれると、
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0.1%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Ti量が増加し好ましくない。このた
めNは80ppm以下とする。 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。
The present invention relates to a method of manufacturing a steel plate for ultra-deep drawing. Two types of steel plates for ultra-deep drawing are known: Ti killed steel plates (Japanese Patent Publication No. 18066/1982) and Nb killed steel plates (Japanese Patent Publication No. 1245/1983). For example, the technology of Nb killed steel sheets has significantly improved compared to when the patent application was filed, and the technology disclosed in the patent publication is not economical and therefore impractical. C<
Since it has become possible to easily reduce C to a range of 500 ppm, there have recently been reports in academic literature that it is possible to produce with lower amounts of carbon and Nb than the ranges in the above-mentioned patent publications. These steel sheets are designed to have extremely low carbon, and are made by adding Ti or Nb, which has a strong carbonitride-forming ability, to create steel sheets with almost no interstitial elements C and N. It also has the advantage of being able to produce products on the same level as box annealing. However, when these steel plates are actually produced by continuous annealing, they have the following drawbacks. Ti-killed steel sheets are susceptible to secondary processing cracks. In particular, if a material with good quality is targeted and the amount of Tj added is in a range equivalent to or higher than the amount of C and N, secondary processing cracks are likely to occur as the amount of P increases. Another drawback is that the r value deteriorates considerably when P is added. Furthermore, when manufacturing alloyed galvanized steel sheets on a Sendzimer-type continuous hot-dip galvanizing line, which is a type of continuous annealing, powdering (alloying progresses excessively and the plating layer is coated during press forming) is produced. It also has the disadvantage of being prone to the phenomenon of peeling off. However, for this steel type, the hot rolling winding temperature is
It has the advantage of being able to produce stable materials by continuous annealing even at the usual 600-650℃. On the other hand, Nb killed steel requires hot rolling and high temperature coiling (coiling temperature 700°C). At normal coiling temperatures, the complete recrystallization temperature becomes extremely high, and unrecrystallized areas may remain in the continuous annealing furnace's possible temperature range (approximately 850°C or less), and depending on the amount of Nb. There are large variations in material. 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 ends of the hot-rolled coil. . However, high-temperature winding not only makes the scale thicker and drastically reduces pickling efficiency, but also makes it impossible to obtain sufficient material at the ends of the coil due to the fast cooling rate, resulting in a decrease in yield.
Some Nb-killed steels are particularly large. As a result of a detailed study of the advantages and disadvantages of Ti-killed steel and Nb-killed steel, the present inventors first considered the above-mentioned difference in behavior as follows. Ti is an extremely strong nitride-forming element, and although the carbide in which TiN has already been formed in the hot-rolling furnace has a lower precipitation temperature than the nitride,
It is thought that a considerable amount of nitrides will precipitate even during winding at 600 to 650°C using nitrides as precipitation sites. Therefore, even at a low coiling temperature, a considerable amount of precipitates are precipitated as large precipitates in the hot coil, and only a small amount of precipitates are precipitated during continuous annealing after cold rolling. Therefore, the recrystallization temperature does not need to become extremely high, and the material is considered to be fairly homogeneous. However, most of the C and N precipitate as carbonitrides, and the grain boundaries become clean, and impurity elements such as P, which cause grain boundary embrittlement, segregate there, causing secondary work cracking. Conceivable. Furthermore, when manufacturing alloyed galvanized steel sheets with Ti-killed steel, the Ti in the steel promotes the alloying reaction between the base iron and molten zinc, which tends to cause overalloying and powdering. it is conceivable that. On the other hand, Nb has a considerably inferior nitride formation ability compared to Ti. In fact, in Nb-added steel, Nb forms carbides, but nitrogen precipitates as AlN. AlN is difficult to form during low-temperature coiling, and is not formed in hot-rolled sheets unless the coiling temperature is 700°C or higher.AlN is finely precipitated during continuous annealing after cold rolling, and is used to increase yield strength and reduce elongation. This causes material deterioration such as deterioration. Therefore, even if the front and rear ends of the hot-rolled coil are coiled at a high temperature, the material remains close to that used for low-temperature coiling because the cooling rate is fast. In addition, due to slight variations in hot-rolling winding temperature and differences in the degree of AlN formation at the center and ends of the coil,
It is thought that deterioration of the material at the front and rear ends of the hot-rolled coil and variations in the material inside the coil occur. However, since Nb has inferior carbonitride formation ability compared to Ti, several ppm of carbon is segregated at grain boundaries, and this increases the binding energy of grain boundaries, even if the P content is quite high. It is considered that secondary processing cracking is not a cause for concern. In addition, during the production of alloyed galvanized steel sheets, Nb does not tend to accelerate the alloying reaction between base iron and molten zinc as much as Ti.
Powdering is less likely to occur in Nb-killed steel than in Ti-killed steel. Based on this idea, the following is the concept of manufacturing ultra-deep drawable steel sheets that have good homogeneity of the material inside the coil and no worries about secondary processing cracks, and that powdering is difficult to occur when manufacturing alloyed galvanized steel sheets. It became something like That is, N is not AlN but Ti
The basic concept of the present invention is a steel sheet in which TiN is precipitated as a composite carbide of [Ti, Nb]C before finish hot rolling. As will be described later, the steel of the present invention is superior to Ti-killed steel in that the r value does not deteriorate even when P is added to increase the strength, and secondary processing cracks are extremely unlikely to occur. Furthermore, when producing an alloyed galvanized steel sheet, it has the advantage that powdering is less likely to occur. Next, compared to Nb killed steel,
Because N is fixed as TiN instead of AlN, almost the same level of material can be obtained even with low-temperature winding without high-temperature winding, and the material quality in the coil length direction and coil width direction is extremely high. It is superior in that it is homogeneous. These advantages clearly indicate that the steel of the present invention has various superior material properties that are not possessed by the conventional Ti-killed steel and Nb-killed steel, and the steel of the present invention is extremely advantageous. be. Furthermore, the steel of the present invention with multiple additions of Ti and Nb has unique properties that cannot be predicted from the properties of conventional Ti-killed steel and Nb-killed steel, because the anisotropy of the r value is extremely small. It is. in general,
The r value of Ti-killed steel and Nb-killed grooves is poorest in the rolling direction (L direction) or 45° direction, and best in the direction perpendicular to rolling (C direction). However, the steel of the present invention has an r value that is almost equal in L, C, and 45° directions, or is somewhat larger in the 45° direction, regardless of the cold rolling rate. A high r-value steel sheet with such anisotropy has not been known at all so far, and it is academically interesting, but it also has great practical advantages. In particular, it is easy to imagine that extremely excellent formability is exhibited in the case of a rectangular cylindrical drawing (corners angled at 45°).
However, even in the case of cylindrical deep drawing, the fracture limit is often determined by the minimum r value. In addition, low anisotropy is a great advantage for applications where uniformity of plate thickness after drawing is important, such as outer cylinders for dry cell batteries, and for applications where it is desired to minimize the occurrence of selvage. Therefore, it is thought that the appearance of new steel sheets such as the steel of the present invention will increase public interest in the anisotropy of the r value, rather than simply looking at the average value in three directions. As described above, the steel of the present invention is a completely new steel plate that is superior in all respects to steels with only Ti or Nb added, and has completely unexpected properties, making it extremely advantageous. Next, we will discuss the component range. As mentioned above,
The amount of Ti added is determined by the relationship with N, and should be added at least in an amount sufficient to fix N. That is, 48/14 [N (%) - 0.002% (20 ppm)] <Ti
It is. If Ti is added in an amount exceeding the equivalent amount to N (i.e. 48/14N (%)), it will form sulfides or become carbides, deteriorating secondary workability, and worsening the quality of the material when P is added. At most, the amount of Ti added should be equal to or less than the equivalent amount of C and N. That is, Ti (%) < [48/
12C (%) + 48/14N (%)]. On the other hand, Nb is determined by the relationship with C, and is 0.3 times or more the amount of C in terms of atomic ratio, that is, Nb (%) > 0.3 × 93/12C (%) = 2.33 (%),
And it is desirable to add it in an amount of 0.003% or more and less than 0.025%. When Nb (%) / C (%) < 2.33, there is a problem that a composite carbide of [Ti, Nb] C is not formed and solid solution C remains, making it impossible to obtain a non-aging steel sheet.
At 0.025% Nb, the material property values show behavior similar to conventional Nb killed steel, the recrystallization temperature increases, and material deterioration at the front and rear ends of the coil increases, which deviates from the principle of the present invention. FIGS. 1 and 2 show the scope of the present invention from the relationship between the amounts of Ti and Nb. Figure 1 shows the amount of Nb fixed at a constant amount (0.022%).
These are material property values when changing the amount of Ti.
C; 0.005, Si; 0.01, Mn; 0.25, P; 0.02,
A sample of S; 0.01, Sol.Al; 0.06, N; 0.005 (each %) was hot-rolled at 720°C, and a;
Center part in the coil length direction, b; front and rear ends of the coil are shown. When the amount of Ti is insufficient to fix N, i.e. 48/14 [N (%) - 0.002%] = 0.010%>
In the case of Ti, the material deterioration at the front and rear ends of the coil is large, and the anisotropy of the r value shows the characteristics of the r value of ultra-low carbon steel with a trace amount of Nb added.
The effect of compound addition is small. On the other hand, when Ti is added in an amount equivalent to or more than C and N, the material deterioration at the front and rear ends of the coil is very small, but the anisotropy of the r value is
It is almost the same as that of Ti Killed.
That is, only when an appropriate amount of Ti is added, the effect of improving the texture due to the combined addition of Ti and Nb appears, and the material exhibits excellent r-value characteristics with extremely low anisotropy. Such excellent r-value isotropy is due to Ti, Nb
This is due to the texture that cannot be obtained with steel with only Ti and Nb added. Furthermore, as shown in Figure 1, the amount of Ti added is 0.010% because the lower limit is Ti > 48/14 [N (%) - 0.002], and the upper limit is Ti < [4.00C (%) + 3.43N]. (%)〕
0.037%. If Ti is less than 0.010%, the ratio of Ti fixing N as TiN at the front and rear ends of the coil will be small, and conversely, the ratio of N precipitated as AlN will increase, and the size of AlN precipitates will be extremely small, so the coil This is undesirable because it deteriorates the material of the front and rear ends. When the amount of Ti is 0.037 or more, the total amount of C and N in the steel is
Fixed by Ti. In this case, the material properties are the same as those of steel with only Ti added, and primary workability is deteriorated and powdering is more likely to occur, which is undesirable. Figure 2 shows the material properties when a constant amount (0.02%) sufficient to fix the amount of Ti to N is added and the amount of Nb added is varied. The chemical composition of the sample is at almost the same level as in Figure 1.
If the Nb content is too low relative to the C content, it will exhibit characteristics similar to those of simple ultra-low carbon steel as shown in Figure 1, and 45
The r value in the ° direction is extremely small, and the anisotropy is extremely large. Furthermore, the material deterioration at the front and rear ends of the coil is large, and non-aging properties cannot be obtained, as in the case of FIG. 1. When the amount of Nb exceeds 0.025%, the anisotropy of the r value follows the same trend, but there is a drawback that the value of r c becomes low, and the material deterioration at the front and rear ends of the coil length is extremely large, which is unique to Nb killed steel. Similar to characteristics. As can be seen from FIG. 2, only by adding a suitable combination of Ti and Nb can the anisotropy of the r value be extremely small and the excellent properties of uniform material in the longitudinal direction of the coil can be obtained. Such properties cannot be obtained with steels to which only Ti and Nb are added, indicating that the combined addition of Ti and Nb is non-essential. The following are the reasons why, compared to the addition of Nb alone, a highly stable material is produced that does not depend on the coiling temperature, and compared to the addition of Ti alone, a high quality material can be obtained even with a small amount of alloying element added. by. When the addition amount of Ti and Nb is well-balanced as in the range of the present invention, a composite precipitate of [Ti, Nb]C is formed, but this has a higher precipitation initiation temperature than TiC and NbC, and a large Since it precipitates as a precipitate, it seems to exhibit good recrystallization behavior even at a low coiling temperature. It is thought that this results in an isotropic r value. On the other hand, in Nb-added steel or when 0.025% Nb is added, NbC is produced, and the precipitation state changes greatly depending on the coiling temperature, and in the case of low-temperature coiling, fine NbC is continuous. The material quality deteriorates significantly due to the increased recrystallization temperature during annealing. In addition, when Ti is added alone, Ti and C
If the atomic ratio of +N is not greater than 1, the material will deteriorate significantly. This is because unless the amount of Ti is increased, TiC will not precipitate sufficiently in the hot rolled sheet, and will precipitate finely during continuous annealing. It is thought that it becomes hard and has poor ductility. As detailed above, Nb<0.025% and Nb
When Ti is added in combination, a steel plate with excellent ductility and deep drawability can be obtained. As stated above, the composition range of the steel of the present invention is Ti
The amount added depends on the amount of N in the steel, Ti (%) > 48/14
[N (%) - 0.002%) and Ti (%) < [48/12C
(%) + 48/14N (%)], and the amount of Nb added is determined by the amount of C in the steel, Nb (%) > 0.3 × 93/12
C (%) = 2.33C (%) and 0.003%≦Nb (%)<
The content is 0.025%. The range of components other than Ti and Nb is C: 0.007% or less,
It consists of Si: 0.8% or less, Mn: 1.0% or less, P: 0.1% or less, Al: 0.01 to 0.1%, N: 80ppm or less, and other unavoidable impurities. When the amount of C is large, Nb is inevitably added to fix C.
A correspondingly larger amount is required, and the amount of [Ti, Nb]C produced increases, which inhibits the growth of crystal grains, leading to a decrease in r value, increase in yield strength, and decrease in elongation. Therefore, from the perspective of producing ultra-deep drawable steel sheets, C
shall be 0.007% or less. When producing hot-dip galvanized steel sheets, Si tends to reduce the adhesion of the plating layer film.
Should be 0.8% or less. Particularly when alloying treatment is not performed, a content of 0.3% or less is desirable. If Mn is added in large amounts, the r value will deteriorate significantly.
The upper limit is 1.0%, but a lower value is preferable from the viewpoint of high r value. The steel of the present invention is difficult to cause secondary processing cracks due to the combined addition of Ti and Nb, but if a large amount of P is included,
The upper limit of P is set to 0.1% because the amount of grain boundary segregation increases, embrittles the grain boundaries, and promotes secondary work cracking. 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. becomes significant. On the other hand, if it is added in a large amount, Al 2 O 3 inclusions will increase, which is not preferable. For the above reasons, Al is set to 0.01 to 0.1%. N is fixed to Ti as TiN, but if the N content is large, the required amount of Ti increases, which is not preferable. Therefore, N should be 80 ppm or less. 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.

【表】 上記の供試鋼を熱間仕上げ温度910℃、巻取温
度720℃、620℃の2水準で、板厚4.0mmに熱間圧
延し、0.8mmまで冷間圧延した後、第3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
た。即ち800〜850℃×30sec保定し、約400℃まで
の冷却速度5〜100℃/secとした。 このようにして得た冷延鋼板の材質試験結果を
第2表に示す。第2表は(1)は巻取温度720℃、第
2表(2)は巻取温度620℃の例である。
[Table] The above test steel was hot-rolled to a thickness of 4.0 mm at two levels: a hot finishing temperature of 910°C, a coiling temperature of 720°C, and 620°C, and then cold-rolled to 0.8 mm. Annealing was performed on a continuous annealing line using the annealing cycle shown in the figure. That is, the temperature was maintained at 800 to 850°C for 30 seconds, and the cooling rate to about 400°C was 5 to 100°C/sec. Table 2 shows the material test results of the cold-rolled steel sheets obtained in this way. In Table 2, (1) is an example of a winding temperature of 720°C, and Table 2 (2) is an example of a winding temperature of 620°C.

【表】【table】

【表】 第4表は供試鋼のコイル長方向の機械的性質の
分布の概略を示したものである。図中A;Ti、
Nb添加鋼(供試鋼2)B;Tiキルド鋼(6)C;Nb
キルド鋼(4)を示し、a;巻取温度720℃、b;巻
取温度620℃の例である。 第2表および第2図から本発明のTi、Nb複合
添加鋼が従来のTiキルド鋼、Nbキルド鋼に比較
して極めて優れた材質特性を有することが明らか
である。 Nbキルド鋼は通常の620℃の巻取温度では再結
晶温度が非常に高くなり、結果として降伏強度が
高く、伸びが低い。巻取温度720℃の高温巻取材
においても、コイル端部では冷却速度が大きいた
め、通常巻取材に近い材質となつており、極めて
歩留が低いという結果を得た。これに対して、
Tiキルド鋼はC、Nを析出させるのに十分なTi
を添加した場合(6)には、コイル長方向に優れた材
質の均一性を有す。しかし、Ti添加量がC、N
を析出させる量よりも不足する場合、即ちTi/
C+N(原子比)<1の場合(7)には著しく材質が
劣化する。これに対し、Ti、Nb複合添加鋼は、
十分なTi量を添加したTiキルド鋼とほぼ同様の
優れた材質の均一性を示した。 次に二次加工割れの試験(絞り比3.0)を行な
つた結果、第5図に示すごとく、Tiキルド鋼は
割れ発生温度がNbキルド鋼、Ti、Nb添加鋼に比
較して約30℃高いという欠点を有することが明ら
かになつた。これに対し、Ti、Nb添加鋼はNbキ
ルド鋼と同等の良好なレベルである。しかしなが
ら、箱焼鈍材のように焼鈍後の冷却速度が遅い場
合には、冷却中にPの粒界偏析が起こつて脆性発
生温度が高くなるので、本発明鋼は連続焼鈍で製
造することが必要である。 さらに特筆すべき点はr値の異方性に関してで
ある。第4図に示す如く、通常温度巻取材ではど
の鋼種もΔrは比較的小さいが、高温巻取材では
Tiキルド鋼、Nbキルド鋼ではΔrが非常に大き
い。第6図に各鋼種の値のr値の面内異方性の
代表値を示したが、Tiキルド鋼、Nbキルド鋼の
45゜は特に高温巻取材で非常に低く、絞り成形
時に問題となる可能性が高い。それに対して
Ti、Nb複合添加鋼は、Nbキルド鋼のように低温
巻取材の値が極端に低いこともなくかつ異方性
が極めて小さく、rL、rCに比較してr45゜はほ
とんど等しいかあるいはわずかに大きい値を示
し、四角筒絞り成形の場合などに、特に優れた成
形性を発揮する。 第7図、冷間圧延率を変化させた場合のr値の
挙動を示したものである。図中a;巻取温度720
℃、b;巻取温度620℃の例である。Ti、Nb複合
添加鋼はr値の異方性がTiキルド鋼、Nbキルド
鋼に比べて著しく低いのは既に述べた通りである
が、この特性は冷延率の大小によらず認められる
ものである。さらにTi、Nb添加鋼はTiキルド
鋼、Nbキルド鋼と比較して、低い冷延率でも比
較的高いr値を有し、実操業の面からも優れた鋼
種であると言える。 また第2表に示す如く、Ti、Nb複合添加鋼は
優れた加工硬化係数n値を有し、Tiキルド鋼、
Nbキルド鋼と同様に非時効性を示している。 実施例 2 第3表は本発明鋼及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] Table 4 shows an outline of the distribution of mechanical properties in the coil length direction of the test steel. A in the figure; Ti,
Nb-added steel (sample steel 2) B; Ti-killed steel (6) C; Nb
Killed steel (4) is shown, a: coiling temperature is 720°C, b: coiling temperature is 620°C. It is clear from Table 2 and FIG. 2 that the Ti and Nb composite steel of the present invention has extremely superior material properties compared to conventional Ti-killed steel and Nb-killed steel. Nb killed steel has a very high recrystallization temperature at the normal coiling temperature of 620°C, resulting in high yield strength and low elongation. Even in high-temperature web material with a winding temperature of 720°C, the cooling rate is high at the ends of the coil, so the material is close to that of normal web material, resulting in extremely low yields. On the contrary,
Ti-killed steel has enough Ti to precipitate C and N.
When (6) is added, the material has excellent uniformity in the length direction of the coil. However, the amount of Ti added is
If the amount is insufficient than the amount to precipitate Ti/
When C+N (atomic ratio)<1 (7), the material quality deteriorates significantly. On the other hand, Ti and Nb composite added steel
It showed excellent material uniformity, almost the same as Ti-killed steel with a sufficient amount of Ti added. Next, we conducted a secondary work cracking test (drawing ratio 3.0) and found that the crack initiation temperature in Ti-killed steel was approximately 30°C compared to Nb-killed steel, Ti, and Nb-added steel, as shown in Figure 5. It has become clear that it has the disadvantage of being expensive. On the other hand, steel with Ti and Nb additions has a good level equivalent to that of Nb killed steel. However, when the cooling rate after annealing is slow, such as in box-annealed materials, grain boundary segregation of P occurs during cooling, raising the temperature at which brittleness occurs, so the steel of the present invention must be manufactured by continuous annealing. It is. A further noteworthy point is the anisotropy of the r value. As shown in Figure 4, Δr is relatively small for all steel types in normal temperature rolled material, but in high temperature rolled material, Δr is relatively small.
Δr is very large in Ti-killed steel and Nb-killed steel. Figure 6 shows typical values of the in-plane anisotropy of the r value for each steel type. It is highly likely that On the other hand
Unlike Nb-killed steel, Ti and Nb composite additive steel does not have an extremely low value for low-temperature web material and has extremely small anisotropy, and compared to r L and r C , r 45 ° is almost equal. Alternatively, it exhibits a slightly larger value, and exhibits particularly excellent formability in the case of rectangular tube drawing. FIG. 7 shows the behavior of the r value when the cold rolling rate is changed. Figure a: Winding temperature 720
°C, b: This is an example of a winding temperature of 620 °C. As already mentioned, the anisotropy of the r value of Ti and Nb composite steel is significantly lower than that of Ti-killed steel and Nb-killed steel, but this property is observed regardless of the cold rolling rate. It is. Furthermore, compared to Ti-killed steel and Nb-killed steel, Ti and Nb-added steel have a relatively high r value even at a low cold rolling rate, and can be said to be an excellent steel type from the standpoint of actual operation. In addition, as shown in Table 2, Ti and Nb composite added steel has an excellent work hardening coefficient n value, Ti killed steel,
Like Nb killed steel, it exhibits non-aging properties. Example 2 Table 3 shows the compositions of the steel of the present invention and the test steel used for comparison.

【表】 上記の供試鋼は、本願発明のTi、Nb複合添加
鋼、従来のTiキルド鋼、Nbキルド鋼に合金元素
(主としてP)を添加して高強度化したものであ
る。これらの鋼を熱間仕上げ温度910℃、巻取温
度720℃で板厚4.0mmに熱間圧延し、0.8mmまで冷
間圧延した後、第3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た。 このようにして得た冷延鋼板の材質試験結果を
第4表に示す。
[Table] The above test steels are made by adding alloying elements (mainly P) to the Ti and Nb composite addition steel of the present invention, conventional Ti killed steel, and Nb killed steel to increase their strength. These steels were hot rolled to a thickness of 4.0 mm at a hot finishing temperature of 910°C and a coiling temperature of 720°C, then cold rolled to 0.8 mm, and then transferred to a continuous annealing line using the annealing cycle shown in Figure 3. and annealed. Table 4 shows the material test results of the cold-rolled steel sheets obtained in this way.

【表】【table】

【表】 第8図は各供試鋼のコイル長方向の材質特性値
の分布を示したものである。図中A;Ti、Nb複
合添加鋼(供試鋼8、9)、B;Tiキルド鋼(11)、
C;Nbキルド鋼(10)について示した。第4表、第
8図から次のことが明らかである。 Tiキルド鋼にPを添加した場合には、Ti、Nb
添加鋼、Nbキルド鋼に比べてコイル中心部でr
値が約0.2劣るという欠点を有する。またTiキル
ド鋼にP添加した場合には、第5図に示した二次
加工割れ発生温度がさらに上昇する傾向を示し
た。Nbキルド鋼で、コイル端部の材質劣化が著
しい。これらの従来鋼に比較して、Ti、Nbb複合
添加鋼は、r値のレベルは、コイル長方向中心部
でNbキルド鋼と同等に高く、コイル長方向の材
質分布は、Tiキルド鋼と同様に極めて均一であ
る。さらにr値の異方性が極めて小さいという、
Tiキルド鋼、Nbキルド鋼にない優れた特性を有
している。このように合金元素を添加して高強度
化した場合にも、本発明鋼の優位性が明らかとな
つた。 実施例 3 第1表、第3表に示し供試鋼のうち2、3、
5、6、8、10、11について、実施例2の場合と
同一条件にて冷間圧延まで行なつた後、第9図に
示すサイクルを用いて溶融亜鉛めつき鋼板を製造
した。図中800〜850℃×30sec保定し、(a)、約450
℃まで冷却速度3〜100℃/secで冷却し(b)、亜鉛
浴450〜500℃(c)で処理し、(イ)、(ロ)は約500〜560℃
で合金化処理(d)を示す。(イ)のサイクルは合金化処
理を行なわない場合に相当し、(ロ)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。 合金化処理の有無により、機械的性質はほとん
ど影響を受けなかつたが、第5表に(ロ)の合金化処
理を行なつた場合の材質特性値を示す。
[Table] Figure 8 shows the distribution of material property values in the coil length direction for each sample steel. In the figure, A: Ti, Nb composite addition steel (sample steel 8, 9), B: Ti killed steel (11),
C: Shown for Nb killed steel (10). The following is clear from Table 4 and Figure 8. When P is added to Ti-killed steel, Ti, Nb
r at the center of the coil compared to additive steel and Nb-killed steel.
The disadvantage is that the value is about 0.2 lower. Furthermore, when P was added to Ti-killed steel, the secondary processing crack initiation temperature shown in FIG. 5 showed a tendency to further increase. With Nb killed steel, there is significant material deterioration at the end of the coil. Compared to these conventional steels, the Ti and Nbb composite addition steel has an r value as high as that of the Nb killed steel at the center in the longitudinal direction of the coil, and the material distribution in the longitudinal direction of the coil is similar to that of the Ti killed steel. It is extremely uniform. Furthermore, the anisotropy of r value is extremely small.
It has excellent properties not found in Ti-killed steel and Nb-killed steel. The superiority of the steel of the present invention became clear even when the strength was increased by adding alloying elements in this manner. Example 3 Among the test steels shown in Tables 1 and 3, 2, 3,
Examples 5, 6, 8, 10, and 11 were subjected to cold rolling under the same conditions as in Example 2, and then hot-dip galvanized steel sheets were manufactured using the cycle shown in FIG. In the figure, 800 to 850℃ x 30sec, (a), approx. 450℃
℃ at a cooling rate of 3 to 100℃/sec (b), treated in a zinc bath at 450 to 500℃ (c), and (a) and (b) to about 500 to 560℃
shows the alloying treatment (d). Cycle (a) corresponds to the case where alloying treatment is not performed, and cycle (b) corresponds to the case where alloying treatment is performed to produce an alloyed galvanized steel sheet. Although the mechanical properties were hardly affected by the presence or absence of alloying treatment, Table 5 shows the material property values when alloying treatment (b) was performed.

【表】 各供試鋼の材質特性値は、実施例1、実施例2
で得られたものとほぼ同等の傾向を示し、溶融亜
鉛めつき鋼板としても、本発明鋼が極めて優れて
いることを表わすものである。 合金化亜鉛めつき鋼板の場合は、合金化が過度
に進行すると、脆弱な合金層が成長してプレス成
形時にパウダリングを引き起こす危険性が存在す
る。 第6表は、各鋼種とも製造した10コイルにつ
き、各10ケ所、計100ケ所からサンプルを切り出
し、パウダリング試験を行なつた結果を示したも
のである。
[Table] Material property values of each test steel are shown in Example 1 and Example 2.
This indicates that the steel of the present invention is extremely superior as a hot-dip galvanized steel sheet. In the case of alloyed galvanized steel sheets, if alloying progresses excessively, there is a risk that a brittle alloy layer will grow and cause powdering during press forming. Table 6 shows the results of a powdering test performed on samples cut out from 10 locations (10 locations in total) for 10 coils manufactured for each type of steel.

【表】 Tiキルド鋼は、Tiが地鉄と溶融亜鉛の合金化
を促進して過合金化を助長するためにパウダリン
グ発生率が非常に高い。本発明のTi、Nb複合添
加鋼ではNbキルド鋼とほぼ同等のレベルにあ
り、耐パウダリング性に優れており、この点から
も良好な合金化亜鉛めつき鋼板素材として最適で
ある。 以上のようにTi、Nbを複合添加することによ
りTi、Nbをそれぞれ単独に添加した材料では得
られない種々の優れた特性が得られることにな
り、本発明の新規性が示された。
[Table] Ti-killed steel has a very high powdering rate because Ti promotes alloying of base iron and molten zinc and promotes overalloying. The Ti and Nb composite additive steel of the present invention has excellent powdering resistance, which is almost at the same level as Nb killed steel, and from this point of view as well, it is optimal as a good alloyed galvanized steel sheet material. As described above, by adding Ti and Nb in combination, various excellent properties that cannot be obtained with a material in which Ti and Nb are added individually are obtained, demonstrating the novelty of the present invention.

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

第1図はTi、Nb複合添加鋼の材質特性に及ぼ
すTi量の影響を示す説明図、第2図はTi、Nb複
合添加鋼の材質特性に及ぼすNb量の影響を示す
説明図、第3図は焼鈍サイクルを示す説明図、第
4図はコイル長方向の材質試験値の分布を示す説
明図、第5図は二次加工割れ発生温度を示す説明
図、第6図はr値およびr値の異方性を示す説明
図、第7図はr値の冷間圧延率依存性を示す説明
図、第8図はコイル長方向の材質試験値の分布を
示す説明図、第9図は焼鈍サイクルを示す説明
図、である。
Figure 1 is an explanatory diagram showing the influence of the amount of Ti on the material properties of steel with Ti and Nb composite additions. Figure 2 is an explanatory diagram showing the influence of the amount of Nb on the material properties of steel with Ti and Nb composite additions. Figure 4 is an explanatory diagram showing the annealing cycle, Figure 4 is an explanatory diagram showing the distribution of material test values in the longitudinal direction of the coil, Figure 5 is an explanatory diagram showing the temperature at which secondary processing cracks occur, and Figure 6 is an explanatory diagram showing the r value and r Figure 7 is an explanatory diagram showing the anisotropy of values, Figure 7 is an explanatory diagram showing the dependence of r value on cold rolling rate, Figure 8 is an explanatory diagram showing the distribution of material test values in the coil length direction, and Figure 9 is an explanatory diagram showing the distribution of material test values in the coil length direction. It is an explanatory view showing an annealing cycle.

Claims (1)

【特許請求の範囲】 1 C:0.007%以下、Si:0.8%以下、Mn:1.0
%以下、P:0.1%以下、Al:0.01〜0.1%、N:
80ppm以下及び他の不可避的不純物から成り、
かつTiとNbを添加し、Tiは48/14〔N(%)−
0.002%〕<Ti、かつTi(%)<〔4.00C(%)+
3.43N(%)〕を満たし、Tiは0.010%〜0.037%の
範囲内で含有し、NbはNb(%)>2.33C(%)
で、かつ0.003%以上で0.025%未満の量添加した
成分の鋼を、熱間圧延および冷間圧延後、700℃
以上Ac3変態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: 0.1% or less, Al: 0.01-0.1%, N:
Consisting of less than 80ppm and other unavoidable impurities,
And Ti and Nb are added, Ti is 48/14 [N (%) -
0.002%〕<Ti, and Ti(%)<[4.00C(%)+
3.43N (%)], Ti is contained within the range of 0.010% to 0.037%, and Nb is Nb (%) > 2.33C (%).
The steel containing the added ingredients in an amount of 0.003% or more and less than 0.025% is heated to 700℃ after hot rolling and cold rolling.
A method for producing a steel plate for ultra-deep drawing, characterized by continuous annealing at a temperature below the Ac 3 transformation point.
JP20617381A 1981-12-22 1981-12-22 Manufacture of super deep drawing steel sheet Granted JPS58107414A (en)

Priority Applications (1)

Application Number Priority Date Filing Date Title
JP20617381A JPS58107414A (en) 1981-12-22 1981-12-22 Manufacture of super deep drawing steel sheet

Applications Claiming Priority (1)

Application Number Priority Date Filing Date Title
JP20617381A JPS58107414A (en) 1981-12-22 1981-12-22 Manufacture of super deep drawing steel sheet

Publications (2)

Publication Number Publication Date
JPS58107414A JPS58107414A (en) 1983-06-27
JPS6132375B2 true JPS6132375B2 (en) 1986-07-26

Family

ID=16519011

Family Applications (1)

Application Number Title Priority Date Filing Date
JP20617381A Granted JPS58107414A (en) 1981-12-22 1981-12-22 Manufacture of super deep drawing steel sheet

Country Status (1)

Country Link
JP (1) JPS58107414A (en)

Cited By (3)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
WO2001098552A1 (en) 2000-06-20 2001-12-27 Nkk Corporation Thin steel sheet and method for production thereof
WO2013054464A1 (en) 2011-10-13 2013-04-18 Jfeスチール株式会社 High-strength cold-rolled steel plate having excellent deep drawability and in-coil material uniformity, and method for manufacturing same
WO2013114850A1 (en) 2012-01-31 2013-08-08 Jfeスチール株式会社 Hot-dip galvanized steel sheet and production method therefor

Families Citing this family (13)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
JPS6082615A (en) * 1983-10-11 1985-05-10 Kawasaki Steel Corp Production of steel sheet having high drawability
JPS6126756A (en) * 1984-07-17 1986-02-06 Kawasaki Steel Corp Dead soft steel sheet having high suitability to chemical conversion treatment
JPS6160872A (en) * 1984-08-30 1986-03-28 Kawasaki Steel Corp Hot dip zn-al alloy coating steel sheet superior in press formability and its manufacture
JPS6167721A (en) * 1984-09-10 1986-04-07 Kawasaki Steel Corp Manufacture of non-aging cold rolled steel plate by continuous annealing
JPS6176621A (en) * 1984-09-25 1986-04-19 Kawasaki Steel Corp Manufacture of ultralow carbon cold rolled steel sheet superior in phosphate treatability and formability
JPS61113725A (en) * 1984-11-08 1986-05-31 Nippon Steel Corp Manufacture of cold rolled steel sheet extremely superior in press formability
JPS61113724A (en) * 1984-11-08 1986-05-31 Nippon Steel Corp Manufacture of cold rolled steel sheet extremely superior in press formability
JPS61276962A (en) * 1985-05-31 1986-12-06 Kawasaki Steel Corp Alloyed and galvanized steel sheet for deep drawing having excellent baking hardenability and powdering resistance
JPS61276931A (en) * 1985-05-31 1986-12-06 Kawasaki Steel Corp Production of cold rolled steel sheet having extra-deep drawing having baking hardenability
JPS61276961A (en) * 1985-05-31 1986-12-06 Kawasaki Steel Corp Alloyed and galvanized steel sheet for extra-deep drawing and its production
US4973367A (en) * 1988-12-28 1990-11-27 Kawasaki Steel Corporation Method of manufacturing steel sheet having excellent deep-drawability
JP2514298B2 (en) * 1992-11-25 1996-07-10 株式会社神戸製鋼所 Method for producing galvannealed steel sheet with excellent press formability
KR20120040758A (en) * 2006-12-20 2012-04-27 제이에프이 스틸 가부시키가이샤 Cold-rolled steel sheet and process for producing the same

Cited By (8)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
WO2001098552A1 (en) 2000-06-20 2001-12-27 Nkk Corporation Thin steel sheet and method for production thereof
US6743306B2 (en) 2000-06-20 2004-06-01 Nkk Corporation Steel sheet and method for manufacturing the same
US7252722B2 (en) 2000-06-20 2007-08-07 Nkk Corporation Steel sheet
EP2312009A1 (en) 2000-06-20 2011-04-20 JFE Steel Corporation Steel sheet and method for manufacturing the same
EP2312010A1 (en) 2000-06-20 2011-04-20 JFE Steel Corporation Steel sheet and method for manufacturing the same
WO2013054464A1 (en) 2011-10-13 2013-04-18 Jfeスチール株式会社 High-strength cold-rolled steel plate having excellent deep drawability and in-coil material uniformity, and method for manufacturing same
WO2013114850A1 (en) 2012-01-31 2013-08-08 Jfeスチール株式会社 Hot-dip galvanized steel sheet and production method therefor
US9322091B2 (en) 2012-01-31 2016-04-26 Jfe Steel Corporation Galvanized steel sheet

Also Published As

Publication number Publication date
JPS58107414A (en) 1983-06-27

Similar Documents

Publication Publication Date Title
US4504326A (en) Method for the production of cold rolled steel sheet having super deep drawability
EP3696292A1 (en) A high tensile strength galvanized steel sheet with excellent formability and anti-crush properties and method of manufacturing the same
JPS6132375B2 (en)
JP2007530783A (en) High strength bake hardening type cold rolled steel sheet, hot dip plated steel sheet and method for producing the same
KR950007472B1 (en) High strength cold rolled steel sheet having excellent non-aging property at room temperature and suitable for drawing and method of producing the same
JPS5974232A (en) Production of bake hardenable galvanized steel sheet for ultradeep drawing having extremely outstanding secondary processability
JP4065579B2 (en) Ferritic stainless steel sheet with small in-plane anisotropy and excellent ridging resistance and method for producing the same
JPH0128817B2 (en)
JPS5967319A (en) Manufacture of steel plate for extremely deep drawing
JP2787366B2 (en) Manufacturing method of hot-dip galvanized high-tensile cold-rolled steel sheet
JPH03277741A (en) Dual-phase cold roller steel sheet excellent in workability, cold nonaging properties and baking hardenability and its manufacture
JPS59190332A (en) Production of galvanized steel plate for ultradeep drawing having extremely good secondary processability
JP2004143470A (en) Steel sheet excellent in paint bake hardenability and retarded natural aging hardenability and its manufacturing process
WO2021020439A1 (en) High-strength steel sheet, high-strength member, and methods respectively for producing these products
JPS59197526A (en) Preparation of deep drawing cold rolled steel plate having excellent quality uniformity
JP2582894B2 (en) Hot-rolled steel sheet for deep drawing and its manufacturing method
JP4114521B2 (en) Ultra-high strength cold-rolled steel sheet having excellent formability and method for producing the same
JPH0665645A (en) Production of high ductility hot rolled high tensile strength steel sheet
JP2001011538A (en) Production of high tension hot dip galvanized steel sheet
JP3613149B2 (en) Hot-dip galvanized steel sheet
RU2788613C1 (en) Cold-rolled coated steel sheet and method for production thereof
JPS6152218B2 (en)
JPH02156043A (en) Al killed steel sheet for porcelain enameling and its production
JPH02145747A (en) Hot rolled steel sheet for deep drawing and its manufacture
JPS6114219B2 (en)