JP4265582B2 - Hot-rolled steel sheet with excellent impact properties after quenching and method for producing the same - Google Patents
Hot-rolled steel sheet with excellent impact properties after quenching and method for producing the same Download PDFInfo
- Publication number
- JP4265582B2 JP4265582B2 JP2005233094A JP2005233094A JP4265582B2 JP 4265582 B2 JP4265582 B2 JP 4265582B2 JP 2005233094 A JP2005233094 A JP 2005233094A JP 2005233094 A JP2005233094 A JP 2005233094A JP 4265582 B2 JP4265582 B2 JP 4265582B2
- Authority
- JP
- Japan
- Prior art keywords
- quenching
- less
- hot
- present
- steel sheet
- 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 - Fee Related
Links
Images
Landscapes
- Heat Treatment Of Sheet Steel (AREA)
Description
本発明は、自動車の構造部品等に使用される薄鋼板およびその製造方法に関する。 The present invention relates to a thin steel plate used for a structural part of an automobile and a manufacturing method thereof.
現在、ドアインパクトビームやセンターピラー等の自動車構造部品として、軽量かつ高耐久性の観点から980MPa以上の高強度の鋼板が使用されている。しかし、これらの部品は成形性が厳しいため、上記高強度の鋼板を使用した場合、割れや形状不良の問題が多く、また、素材コストも高い。 Currently, high-strength steel sheets of 980 MPa or more are used as automotive structural parts such as door impact beams and center pillars from the viewpoint of light weight and high durability. However, since these parts have severe formability, when the above-described high-strength steel sheet is used, there are many problems of cracks and defective shapes, and the material cost is high.
近年では、このような問題を背景に440MPaレベルの低強度の薄鋼板を用いて成形を行い、高周波焼入れ等により高強度化が図られている。このような例として、非特許文献1では、センターピラーリンフォースメントやフロントクロスメンバー等において、それぞれ、440MPa、390MPaの鋼板を用いて高周波焼入れにより高強度化している。そして、表面が3次元形状をしている部品に対し、高周波焼入れを行うに際して焼入れコイルをロボットに支持させ、これを部品形状に沿って精密に移動させながら焼入れを行う方法を新規に開発している。
In recent years, with such a problem as a background, forming is performed using a thin steel plate having a low strength of 440 MPa, and high strength is achieved by induction hardening or the like. As such an example, in Non-Patent
また、後熱処理により高強度化する技術としては、特許文献1において、レーザー照射による部分強化の方法が開示されている。
As a technique for increasing the strength by post-heat treatment,
特許文献2では、高エネルギー密度ビーム照射により強化する技術が開示されている。
しかしながら、非特許文献1に記載の技術では、焼入れ条件の変動を小さくするため、莫大な設備投資が必要となっている。
However, the technique described in Non-Patent
特許文献1に記載の技術は、レーザー照射部は極く僅かであり、部材の強度上昇には長時間を要する。また、設備投資も莫大となりコスト増を招く。
In the technique described in
特許文献2に記載の技術は、局所的な強化を行うだけであるため、得られる強度レベルも710MPa程度に過ぎない。
Since the technique described in
このように、焼入れ安定性に優れ、かつ焼入れ後の衝撃特性に優れる鋼板は未だ提案されてないのが現状である。 Thus, the present condition is that the steel plate which is excellent in quenching stability and is excellent in the impact characteristic after quenching has not been proposed yet.
よって、本発明は、焼入れ条件による変動が小さく、焼入れ後の衝撃特性に優れる薄鋼板およびその製造方法を提供することを目的とする。 Therefore, an object of the present invention is to provide a thin steel sheet that is less affected by quenching conditions and has excellent impact characteristics after quenching, and a method for producing the same.
本発明者らが上記目的を達成するために、鋭意研究を重ねた結果、以下のことを見出した。 In order to achieve the above object, the present inventors have conducted intensive research and found the following.
1)加熱温度が1000℃以下、特に950℃以下での焼入れ性に対しては、成分組成が大きく影響し、C、Bの添加が必須である。 1) The component composition greatly affects the hardenability when the heating temperature is 1000 ° C. or less, particularly 950 ° C. or less, and the addition of C and B is essential.
2)焼入れ後の衝撃特性に対しては、析出物の粒径、ミクロ組織の影響が大きく、Ti含有鋼において、TiNの形態が加熱時のオーステナイト粒径を大きく変化させ、TiNが微細に析出している場合、著しくオーステナイト粒径が微細化するため、冷却時にフェライトが部分的に生成してしまいフェライトとオーステナイト界面で亀裂が伸展しやすくなり衝撃特性が低下する。 2) The impact size after quenching is greatly affected by the grain size and microstructure of the precipitate, and in Ti-containing steel, the TiN morphology greatly changes the austenite grain size during heating, and TiN precipitates finely. In this case, since the austenite grain size is remarkably refined, ferrite is partially formed during cooling, and cracks are likely to extend at the interface between ferrite and austenite, resulting in reduced impact characteristics.
3)さらに、高周波加熱後の冷却までの時間の変動といった焼入れ条件の変動に対しては、B-(10.8/14)N*の影響が大きく、B-(10.8/14)N*が小さい場合、高周波加熱後の冷却時にフェライトが生成し、オーステナイト粒径が細粒化した場合と同じくフェライトとオーステナイトとの界面で亀裂が伸展しやすくなり衝撃特性が低下する。 3) In addition, the effect of B- (10.8 / 14) N * is large and the value of B- (10.8 / 14) N * is small for fluctuations in quenching conditions such as fluctuations in the time until cooling after induction heating In addition, ferrite is generated during cooling after high-frequency heating, and cracks are likely to extend at the interface between ferrite and austenite, as in the case where the austenite grain size is reduced, and the impact characteristics are reduced.
本発明はかかる知見に基づきなされたもので、鋼成分としてmass%で、C:0.10〜0.37%、Si:1%以下、Mn:1.4%以下、P:0.1%以下、S:0.03%以下、sol.Al:0.01〜0.1％、N:0.0005〜0.0050％、Ti:0.005〜0.05％、B:0.0003〜0.0050％を含有し、
B-(10.8/14)N*≧0.0005％
(N*=N-(14/48)Ti、但し、右辺≦0の場合、N*=0)
を満足し、残部Fe及び不可避不純物からなり、鋼中析出物であるTiNの平均粒径が0.06〜0.30μmであり、かつ焼入れ後の旧オーステナイト粒径が2〜25μmであることを特徴とする焼入れ後の衝撃特性に優れる熱延鋼板である。
The present invention has been made on the basis of such findings, and as a steel component in mass%, C: 0.10 to 0.37%, Si: 1% or less, Mn: 1.4% or less, P: 0.1% or less, S: 0.03% or less, sol.Al: 0.01-0.1%, N: 0.0005-0.0050%, Ti: 0.005-0.05%, B: 0.0003-0.0050%,
B- (10.8 / 14) N * ≧ 0.0005%
(N * = N- (14/48) Ti, where N * = 0 if right side ≤ 0)
And the balance is composed of the remaining Fe and inevitable impurities, the average grain size of TiN being a precipitate in steel is 0.06 to 0.30 μm, and the prior austenite grain size after quenching is 2 to 25 μm It is a hot-rolled steel sheet with excellent impact characteristics after quenching.
この発明において、さらに、鋼成分としてmass%で、Ni、Cr、Moの1種以上を、合計で1%以下含有することを特徴とする請求項1記載の焼入れ後の衝撃特性に優れる熱延鋼板とすることもできる。
In the present invention, the hot rolling excellent in impact characteristics after quenching according to
さらにまた、鋼成分としてmass%で、Nb:0.1%以下を含有することを特徴とする焼入れ後の衝撃特性に優れる熱延鋼板とすることができる。 Furthermore, it can be set as the hot-rolled steel plate which is excellent in the impact property after hardening characterized by mass% as a steel component and containing Nb: 0.1% or less.
上記の発明の熱延鋼板を得ることができる製造方法の発明は、上記の鋼成分を有する鋼を、巻取温度620℃以上720℃以下で熱間圧延することを特徴とする焼入れ後の衝撃特性に優れる薄鋼板の製造方法である。 The invention of the production method capable of obtaining the hot-rolled steel sheet of the above invention is a post-quenching impact characterized by hot-rolling a steel having the above steel components at a coiling temperature of 620 ° C. or higher and 720 ° C. or lower. This is a method for producing a thin steel sheet having excellent characteristics.
本発明によれば、低温短時間での焼入れ性に優れ、焼入れ条件による変動が小さい焼入れ後の衝撃特性に優れる薄鋼板を得ることができる。さらに、上記薄鋼板が安定して低コストで得られるため、高強度部材として工業的に有用な効果をもたらし、例えば、自動車構造部品として最適である。 ADVANTAGE OF THE INVENTION According to this invention, the thin steel plate which is excellent in the hardenability in a low temperature for a short time, and is excellent in the impact characteristic after quenching with the small fluctuation | variation by quenching conditions can be obtained. Furthermore, since the thin steel sheet can be obtained stably and at low cost, it provides an industrially useful effect as a high-strength member, and is optimal, for example, as an automobile structural component.
まず、鋼板の鋼成分について限定理由を説明する。 First, the reasons for limitation of the steel components of the steel sheet will be described.
C:0.10〜0.37%
Cは、焼入れ後の強度を得るための重要な元素であり、980MPa以上を得るには少なくとも0.10%以上が必要である。しかし、0.37%を超えて添加すると強度は得られるものの衝撃特性が著しく低下する。従って、本発明においてCの添加範囲は0.10〜0.37%とする。優れた衝撃特性を得るには0.30%以下が好ましい。
C: 0.10-0.37%
C is an important element for obtaining strength after quenching, and at least 0.10% or more is necessary to obtain 980 MPa or more. However, if added over 0.37%, although the strength is obtained, the impact characteristics are remarkably lowered. Therefore, in the present invention, the addition range of C is set to 0.10 to 0.37%. In order to obtain excellent impact characteristics, 0.30% or less is preferable.
Si:1%以下
Siは焼入れ性を向上させるとともに固溶強化により強度を上昇させる元素である。しかし、1%を超えて添加すると、熱延板において偏析帯であるバンド組織が著しくなるため衝撃特性が劣化する。従って、本発明においてはSiの添加範囲は1%以下とする。また、優れた衝撃特性を得るには0.5%以下が好ましい。
Si: 1% or less
Si is an element that improves the hardenability and increases the strength by solid solution strengthening. However, if it exceeds 1%, the band structure, which is a segregation band, becomes noticeable in the hot-rolled sheet, so that the impact characteristics are deteriorated. Therefore, in the present invention, the Si addition range is 1% or less. Further, 0.5% or less is preferable for obtaining excellent impact characteristics.
Mn:2.5%以下
Mnは焼入れ性を向上させるとともに固溶強化により強度を上昇させる元素である。しかし、2.5%を超える添加は、偏析帯であるマンガンバンドの生成が顕著となり衝撃特性が劣化する。従って、本発明においてMnの添加範囲は2.5%以下とする。また、優れた衝撃特性を得るには1.5%以下が好ましい。
Mn: 2.5% or less
Mn is an element that improves hardenability and increases strength by solid solution strengthening. However, when the content exceeds 2.5%, the formation of a manganese band, which is a segregation band, becomes prominent and the impact characteristics deteriorate. Therefore, in the present invention, the Mn addition range is 2.5% or less. Further, 1.5% or less is preferable for obtaining excellent impact characteristics.
P:0.1%以下
Pは焼入れ性を向上させるとともに固溶強化により強度を上昇させる元素である。しかし、Pは粒界に偏析し衝撃特性を低下させる元素でもある。B添加により粒界偏析は抑制されるが、それでもPの0.1%を超える添加は粒界脆化を招き衝撃特性が劣化する。よって、本発明においてはPの添加範囲は0.1%以下とする。また、優れた衝撃特性を得るには0.05%以下が好ましい。
P: 0.1% or less
P is an element that improves hardenability and increases strength by solid solution strengthening. However, P is an element that segregates at the grain boundary and lowers the impact characteristics. Grain boundary segregation is suppressed by addition of B, but addition exceeding 0.1% of P still causes grain boundary embrittlement and deteriorates impact characteristics. Therefore, in the present invention, the addition range of P is 0.1% or less. Further, 0.05% or less is preferable for obtaining excellent impact characteristics.
S:0.03%以下
Sは、硫化物を形成し衝撃特性を低下させるため、低減しなければならない元素である。含有量が0.03%を超える場合、衝撃特性が著しく劣化するため、0.03%以下に抑制しなければならない。よって、本発明においてSの添加範囲は0.03%以下とする。なお、優れた衝撃特性を得るには0.02%以下が好ましい。
S: 0.03% or less
S is an element that must be reduced in order to form sulfides and reduce impact properties. When the content exceeds 0.03%, the impact characteristics are remarkably deteriorated, so the content must be suppressed to 0.03% or less. Therefore, in the present invention, the S addition range is 0.03% or less. In order to obtain excellent impact characteristics, 0.02% or less is preferable.
sol.Al:0.01〜0.1％
sol.Alは脱酸剤として用い鋼の清浄度を向上させる元素である。0.01%未満の添加は、清浄度が低下し介在物が増大し、衝撃特性を低下させる。一方、0.1％を越える添加はAlNの形成が顕著となり、焼入れ時のオーステナイトが微細化し冷却時にフェライトが生成してしまい衝撃特性が劣化する。よって、本発明においてsol.Alの添加範囲は0.01〜0.1％とする。なお、優れた衝撃特性を得るには0.03〜0.07％が好ましい。
sol.Al:0.01-0.1%
sol.Al is an element used as a deoxidizer to improve the cleanliness of steel. Addition of less than 0.01% reduces cleanliness, increases inclusions, and lowers impact properties. On the other hand, if the addition exceeds 0.1%, the formation of AlN becomes remarkable, the austenite at the time of quenching becomes fine, and ferrite is formed at the time of cooling, so that the impact characteristics deteriorate. Therefore, in the present invention, the addition range of sol.Al is set to 0.01 to 0.1%. In order to obtain excellent impact characteristics, 0.03% to 0.07% is preferable.
N:0.0005〜0.0050％
NはTiNを形成し加熱時のオーステナイトの粒成長を抑制し衝撃特性を向上させる重要な元素であり、少なくとも0.0005％以上が必要である。一方、0.0050％を越える添加はTiNのみならずBN、AlNの形成も顕著となり、焼入れ時のオーステナイトが微細化し冷却時にフェライトが生成してしまい衝撃特性が劣化する。よって、本発明においてNの添加範囲は0.0005〜0.0050％とする。
N: 0.0005-0.0050%
N is an important element that forms TiN, suppresses the grain growth of austenite during heating, and improves the impact characteristics, and at least 0.0005% or more is necessary. On the other hand, the addition exceeding 0.0050% not only forms TiN but also forms BN and AlN, austenite at the time of quenching becomes fine and ferrite is formed at the time of cooling, resulting in deterioration of impact characteristics. Therefore, in the present invention, the addition range of N is 0.0005 to 0.0050%.
Ti:0.005〜0.05％
Tiは、NとTiNを形成し、オーステナイト粒の粗大化を抑制し、衝撃特性を向上させる重要な元素である。しかし、添加量が0.005％未満の場合、十分な効果が得られず、0.05％を超える過剰な添加はTiCの形成が顕著となり、低温短時間焼入れ時のオーステナイト粒成長を著しく抑制し、加熱後の冷却時にフェライトが生成し衝撃特性が劣化する。よって、本発明においてTiの添加範囲は、0.005〜0.05％とする。
Ti: 0.005-0.05%
Ti is an important element that forms N and TiN, suppresses coarsening of austenite grains, and improves impact properties. However, if the amount added is less than 0.005%, sufficient effect cannot be obtained, and if it exceeds 0.05%, TiC formation becomes prominent, austenite grain growth during quenching at a low temperature for a short time is remarkably suppressed, and after heating Ferrite is formed during the cooling of the steel, and the impact characteristics deteriorate. Therefore, in the present invention, the addition range of Ti is set to 0.005 to 0.05%.
B:0.0003〜0.0050％
Bは焼入れ性を高めるとともに、加熱後冷却時のフェライト生成を抑制し衝撃特性を向上させる重要な元素である。しかし、添加量が0.0003％未満の場合、十分な効果が得られない。一方、0.0050％を超える添加は熱間圧延の負荷が高くなり操業性が低下するととともに、加工性が低下する。よって、本発明においてBの添加範囲は、0.0003〜0.0050％とする。なお、極めて優れた効果を得るには0.0005〜0.0020％が好ましい。
B: 0.0003-0.0050%
B is an important element that improves hardenability and suppresses the formation of ferrite during cooling after heating and improves impact properties. However, when the addition amount is less than 0.0003%, a sufficient effect cannot be obtained. On the other hand, addition over 0.0050% increases the hot rolling load, lowers operability, and lowers workability. Therefore, in the present invention, the addition range of B is 0.0003 to 0.0050%. In addition, 0.0005 to 0.0020% is preferable for obtaining an extremely excellent effect.
有効B:B-(10.8/14)N*≧0.0005％
有効Bは、焼入れ条件の変動に対して大きな影響を及ぼす比率である。
ここで、N*=N-(14/48)Ti (但し、右辺≦0の場合、N*=0)
そこで、焼入れ後の衝撃特性に及ぼす有効B:B-(10.8/14)N*の影響について調査した。
Effective B: B- (10.8 / 14) N * ≧ 0.0005%
Effective B is a ratio that has a great influence on fluctuations in quenching conditions.
Where N * = N- (14/48) Ti (however, if right side ≤ 0, N * = 0)
Therefore, the effect of effective B: B- (10.8 / 14) N * on the impact properties after quenching was investigated.
ベース成分として、C:0.15%、Si:0.02%、Mn:0.90%、P:0.020%、S:0.015%、sol.Al:0.035%、Ti:0.01%とし、N:0.0018〜0.0030%、B:0〜0.0031%、B-(10.8/14)N*:0〜0.0017%の化学成分を有する鋼を溶製し、次いで、加熱温度:1200℃、熱延仕上温度:870℃、中間温度:700℃、巻取温度:620℃で熱延し、酸洗後、冷圧率:50%、焼鈍温度:720℃で1.2mmtの冷延板を製造した。 As base components, C: 0.15%, Si: 0.02%, Mn: 0.90%, P: 0.020%, S: 0.015%, sol.Al: 0.035%, Ti: 0.01%, N: 0.0018 to 0.0030%, B : 0 ~ 0.0031%, B- (10.8 / 14) N *: 0 ~ 0.0017% steel with chemical composition is melted, then heating temperature: 1200 ° C, hot rolling finishing temperature: 870 ° C, intermediate temperature: Hot rolled at 700 ° C. and coiling temperature: 620 ° C. After pickling, a cold rolled sheet of 1.2 mmt was manufactured at a cold pressure rate of 50% and an annealing temperature of 720 ° C.
次いで、得られたサンプルについて高周波焼入後の衝撃特性を評価した。高周波焼入れは、平板(幅35mm×長さ300mm)に対し高周波コイルを移動させながら加熱・焼入れを実施した。図1に高周波焼入れの実施態様を示す。この時の加熱温度は、900℃の低温とし、加熱時間は、900℃までの通電時間を4秒とした。 Subsequently, the impact characteristics after induction hardening were evaluated about the obtained sample. In the induction hardening, heating and hardening were performed while moving the high-frequency coil on a flat plate (width 35 mm x length 300 mm). FIG. 1 shows an embodiment of induction hardening. The heating temperature at this time was a low temperature of 900 ° C., and the heating time was 4 seconds for the energization time up to 900 ° C.
冷却開始時間は、通常行われる即冷却とし0.5秒と、焼入れ安定性を評価するために1.5秒、3秒の3パターンを実施した。 The cooling start time was 0.5 seconds for the immediate cooling that is normally performed, and three patterns of 1.5 seconds and 3 seconds were performed to evaluate the quenching stability.
高周波焼入れ後の評価としては、シャルピー衝撃試験を実施した。シャルピー衝撃試験は、図2に示すような試験片形状にて、試験温度:-50℃、n=3で行った。 As evaluation after induction hardening, the Charpy impact test was implemented. The Charpy impact test was performed with a test piece shape as shown in FIG. 2 at a test temperature of −50 ° C. and n = 3.
得られた結果を図3に示す。図3より、B-(10.8/14)N*が0.0005%以上で冷却開始時間が3秒においても安定して高いシャルピー衝撃吸収エネルギーが得られることがわかる。 The obtained results are shown in FIG. FIG. 3 shows that high Charpy impact absorption energy can be stably obtained even when B- (10.8 / 14) N * is 0.0005% or more and the cooling start time is 3 seconds.
また、B-(10.8/14)N*が0.0005%未満の場合、焼入れ加熱時の固溶B量が十分確保されず、加熱後の冷却開始時間の遅れるような場合、フェライトが生成し衝撃特性の劣化を招く。 In addition, when B- (10.8 / 14) N * is less than 0.0005%, the amount of dissolved B during quenching heating is not sufficiently secured, and when the cooling start time after heating is delayed, ferrite is generated and the impact characteristics Cause deterioration.
よって、生産上のバラツキを低減し安定して高い衝撃特性を得るために、本発明において、B-(10.8/14)N*が0.0005%以上とする。ただし、N*=N-(14/48)Tiであり、右辺≦0の場合、N*=0である。 Therefore, in order to reduce production variations and stably obtain high impact characteristics, B- (10.8 / 14) N * is set to 0.0005% or more in the present invention. However, when N * = N− (14/48) Ti and the right side ≦ 0, N * = 0.
Ni、Cr、Mo:添加する場合1種以上を合計1%以下
Ni、Cr、Moは焼入れ性向上元素であり、1種以上を添加しても良い。しかし、過剰な添加はコスト増を招くため、Ni、Cr、Moの1種以上を合計1%以下とする。
Ni, Cr, Mo: 1% or more when added, totaling 1% or less
Ni, Cr, and Mo are hardenability improving elements, and one or more of them may be added. However, excessive addition causes an increase in cost, so at least one of Ni, Cr, and Mo is made 1% or less in total.
なお、本発明において、加熱時のオ−ステナイト粒の粗大化抑制のためにNbを0.1%以下、Vを0.1%以下、および延性向上のためにCaを0.01%以下添加しても良い。また、耐食性向上のためにCuを1%を超えない範囲で添加しても良い。 In the present invention, Nb may be 0.1% or less, V may be 0.1% or less for suppressing the coarsening of austenite grains during heating, and Ca may be added 0.01% or less for improving ductility. In order to improve corrosion resistance, Cu may be added within a range not exceeding 1%.
また、本発明において、上記元素以外は実質的にＦｅであり、本発明の作用効果を無くさない限り、不可避不純物をはじめ、他の微量元素を含有するものが本発明の範囲に含まれ得ることを意味する。 Further, in the present invention, elements other than the above elements are substantially Fe, and those containing other trace elements including inevitable impurities can be included in the scope of the present invention unless the effects of the present invention are lost. Means.
次に、析出物について限定理由を説明する。 Next, the reasons for limiting the precipitate will be described.
TiNの平均粒径:0.06〜0.30μm
TiNは、焼入れ加熱時のオーステナイト粒の粗大化を抑制する析出物である。TiN平均粒径が0.06μm未満の場合、オーステナイト粒が極めて微細となり、加熱後の冷却時にフェライトが生成し、衝撃特性が劣化する。一方、0.30μmを超える粗大な析出物の場合、オーステナイトの粒成長を抑制することができない。よって、本発明においてTiN平均粒径は、0.06〜0.30μmとする。
Average particle size of TiN: 0.06-0.30μm
TiN is a precipitate that suppresses coarsening of austenite grains during quenching heating. When the TiN average particle size is less than 0.06 μm, the austenite grains become extremely fine, and ferrite is generated during cooling after heating, resulting in deterioration of impact characteristics. On the other hand, in the case of coarse precipitates exceeding 0.30 μm, austenite grain growth cannot be suppressed. Therefore, in the present invention, the TiN average particle size is 0.06 to 0.30 μm.
次に、ミクロ組織について限定理由を説明する。 Next, the reason for limiting the microstructure will be described.
焼入れ後の旧オーステナイト粒径:2〜25μm
焼入れ後の旧オーステナイト粒径、即ち焼入れ後に測定される変態前の旧オーステナイト粒径は、衝撃特性に大きな影響を及ぼす。旧オーステナイト粒径が2μm未満の場合、加熱後冷却時に一部フェライトが生成しフェライトとオーステナイト界面の応力集中に起因して衝撃特性が低下する。一方、25μmを越えるような粗大粒の場合、粒界脆化が顕著となり従来のJSC980Y(鉄連規格)より衝撃特性が低下する。よって、本発明において焼入れ後の旧オーステナイト粒径は、2〜25μmとする。
Old austenite grain size after quenching: 2-25μm
The prior austenite grain size after quenching, that is, the prior austenite grain size before transformation measured after quenching, has a great influence on the impact properties. When the prior austenite grain size is less than 2 μm, some ferrite is formed during cooling after heating, and the impact characteristics are degraded due to stress concentration at the ferrite-austenite interface. On the other hand, in the case of coarse grains exceeding 25 μm, grain boundary embrittlement becomes remarkable, and impact characteristics are deteriorated as compared with the conventional JSC980Y (iron standard). Therefore, in the present invention, the prior austenite grain size after quenching is 2 to 25 μm.
次に製造方法の限定理由について説明する。 Next, the reason for limiting the manufacturing method will be described.
巻取温度:720℃以下
熱間圧延での巻取温度については、720℃を超えるとパーライトのラメラ間隔が大きくなり、焼入性が低下するとともに、焼入時にセメンタイトが溶け残り衝撃特性が低下する。よって、本発明において、熱間圧延での巻取温度は720℃以下とする。
Winding temperature: 720 ° C or less When the coiling temperature in hot rolling exceeds 720 ° C, the pearlite lamella spacing increases, hardenability decreases, and cementite melts during quenching and impact characteristics decrease. To do. Therefore, in this invention, the coiling temperature in hot rolling shall be 720 degrees C or less.
熱延後の球状化焼鈍温度:640℃以上Ac1変態点以下
熱延鋼板を酸洗した後、セメンタイトを球状化し、優れた加工性と焼入性を得るため球状化焼鈍を行うことができる。焼鈍温度が640℃未満の場合、セメンタイトの球状化が不十分となり、効果が得られない。一方、焼鈍温度がAc1変態点を超える場合、部分的にオーステナイト化して冷却中に粗大なパーライトを生成し、加工性が低下するとともに、焼入性も低下する。また、焼入れ時にセメンタイトが溶け残り衝撃特性が低下する。よって、熱延後に球状化焼鈍を行う場合は、焼鈍温度を640℃以上Ac1変態点以下とする。
Spherical annealing temperature after hot rolling: 640 ° C or higher and Ac 1 transformation point or lower After pickling hot rolled steel sheet, cementite can be spheroidized and spheroidized annealing can be performed to obtain excellent workability and hardenability . When the annealing temperature is less than 640 ° C., the cementite is insufficiently spheroidized and the effect cannot be obtained. On the other hand, when the annealing temperature exceeds the Ac 1 transformation point, it partially becomes austenite and coarse pearlite is generated during cooling, resulting in a decrease in workability and a decrease in hardenability. In addition, the cementite remains undissolved during quenching and the impact characteristics are reduced. Therefore, when spheroidizing annealing is performed after hot rolling, the annealing temperature is set to 640 ° C. or more and Ac 1 transformation point or less.
冷間圧延時の圧下率:30%以上
冷間圧延を行う場合の圧下率(冷圧率)は、30%未満であると焼鈍後に未再結晶部が残るとともに、セメンタイトの球状化が不十分となり、軟質化が得られず加工性が劣化する。よって、冷間圧延を行う場合の冷圧率は、30%以上とする。冷圧率の上限は、特に規定しないが、圧延機への負荷が大きくならないように、80%以下とするのが好ましい。
Reduction ratio during cold rolling: 30% or more When the rolling reduction (cold reduction ratio) is less than 30%, unrecrystallized parts remain after annealing, and cementite spheroidization is insufficient. Thus, softening cannot be obtained and workability is deteriorated. Therefore, the cold pressure ratio when performing cold rolling is set to 30% or more. The upper limit of the cold pressure ratio is not particularly defined, but is preferably 80% or less so as not to increase the load on the rolling mill.
冷間圧延後の焼鈍温度:640℃以上Ac1変態点以下
冷間圧延後の焼鈍については、熱延後の球状化焼鈍を省略した場合は、ここで球状化焼鈍を行う。冷間圧延後の球状化焼鈍の焼鈍温度は、前述の熱延後の球状化焼鈍と同様、640℃以上Ac1変態点以下とする。なお、熱延後の球状化焼鈍を行った場合は、ここで再結晶焼鈍を行う。冷間圧延後の再結晶焼鈍の焼鈍温度は、600℃未満では未再結晶部が残り加工性が低下する。一方、焼鈍温度がAc1変態点を超える場合、部分的にオーステナイト化して冷却中に粗大なパーライトを生成し、加工性が低下するとともに、焼入性も低下する。また、焼入れ時にセメンタイトが溶け残り衝撃特性が低下する。よって、熱延後の球状化焼鈍を行った場合は、焼鈍温度を600℃以上Ac1変態点以下とする。
Annealing temperature after cold rolling: 640 ° C. or more and Ac 1 transformation point or less For annealing after cold rolling, when spheroidizing annealing after hot rolling is omitted, spheroidizing annealing is performed here. The annealing temperature of spheroidizing annealing after cold rolling is set to 640 ° C. or more and Ac 1 transformation point or less, similarly to the spheroidizing annealing after hot rolling described above. In addition, when spheroidizing annealing after hot rolling is performed, recrystallization annealing is performed here. If the annealing temperature of recrystallization annealing after cold rolling is less than 600 ° C., unrecrystallized portions remain and the workability deteriorates. On the other hand, when the annealing temperature exceeds the Ac 1 transformation point, it partially becomes austenite and coarse pearlite is generated during cooling, resulting in a decrease in workability and a decrease in hardenability. In addition, the cementite remains undissolved during quenching and the impact characteristics are reduced. Therefore, when spheroidizing annealing after hot rolling is performed, the annealing temperature is set to 600 ° C. or more and Ac 1 transformation point or less.
本発明において、対象とする薄鋼板は、熱延鋼板あるいは冷延鋼板のいずれでも良い。本発明鋼板を製造する場合、素材鋼は、例えば転炉、電気炉等により溶製される。鋼片の製造は造塊-分塊圧延法、連続鋳造法、薄スラブ鋳造法、ストリップ鋳造法のいずれでも構わない。 In the present invention, the target thin steel sheet may be either a hot-rolled steel sheet or a cold-rolled steel sheet. When manufacturing this invention steel plate, raw material steel is smelted by a converter, an electric furnace, etc., for example. The slab can be manufactured by any of the ingot-bundling rolling method, continuous casting method, thin slab casting method, and strip casting method.
熱延プロセスはスラブ加熱後圧延する方法、連続鋳造後短時間の加熱処理を施してあるいは前記加熱工程を省略して直ちに圧延する方法のいずれでもよいが、優れた表面品質を付与するためには、一次スケールのみならず熱間圧延時に生成する二次スケールについても十分に除去するのが好ましい。なお、熱間圧延中においては、バーヒーターにより加熱を行ってもよい。 The hot rolling process may be either a method of rolling after slab heating, a method of performing a heat treatment for a short time after continuous casting, or a method of rolling immediately after omitting the heating step, but in order to give excellent surface quality It is preferable to sufficiently remove not only the primary scale but also the secondary scale generated during hot rolling. In addition, you may heat with a bar heater during hot rolling.
仕上圧延終了温度は、組織の均一性からAr3変態点以上とすることが好ましい。また、組織の均一化を目的として、仕上圧延後1秒以内に200℃/秒以上の急速冷却を行ってもよい。巻取温度は材質安定性の観点から500℃以上とするのが好ましく、一方、上限はスケール生成増大による酸洗性の低下から700℃以下が好ましい。 The finish rolling finish temperature is preferably not less than the Ar 3 transformation point in view of the uniformity of the structure. For the purpose of homogenizing the structure, rapid cooling at 200 ° C./second or more may be performed within 1 second after finish rolling. The coiling temperature is preferably 500 ° C. or more from the viewpoint of material stability, while the upper limit is preferably 700 ° C. or less because of a decrease in pickling properties due to increased scale formation.
冷延鋼板を本発明の薄鋼板として用いる場合、冷間圧延時の圧延率(冷圧率)は80%以下とするのが好ましい。冷圧率が80%を超えるような高い冷圧率の場合、圧延負荷が高くなりすぎるため生産性を低下させる。このときの冷間圧延はタンデム圧延、リバース圧延のいずれでも良い。 When a cold-rolled steel sheet is used as the thin steel sheet of the present invention, the rolling rate (cold pressure rate) during cold rolling is preferably 80% or less. In the case of a high cold pressure ratio that exceeds 80%, the rolling load becomes too high and the productivity is lowered. The cold rolling at this time may be either tandem rolling or reverse rolling.
なお、再結晶焼鈍を行う方法としては、連続焼鈍、箱焼鈍、または溶融亜鉛めっき処理に先行する連続熱処理のいずれでもよい。 In addition, as a method of performing recrystallization annealing, any of continuous annealing, box annealing, or continuous heat treatment preceding hot dip galvanizing treatment may be used.
本発明に係る熱延鋼板、冷延鋼板は、適宜、表面処理(化成処理、溶融亜鉛めっき、合金化溶融亜鉛めっき)が施されて使用されてもよい。 The hot-rolled steel sheet and cold-rolled steel sheet according to the present invention may be used after appropriately being subjected to surface treatment (chemical conversion treatment, hot-dip galvanizing, alloyed hot-dip galvanizing).
表１に示す鋼番1から13の化学成分組成を有する鋼を溶製し、次いで表2に示す製造条件に従って熱間圧延-焼鈍を行い、2.4mmtの熱延板を製造した。このようにして製造した熱延板について引張試験(JIS 5号、C方向(圧延方向に垂直))、TiNの平均粒径測定および高周波焼入れ特性を調査した。
Steels having chemical composition compositions of
TiN平均粒径は、レプリカ法によりTiNを抽出し、透過電子顕微鏡により析出物を撮影し、サンプル数：500個をマイクロアナライザーを用いて測定した。 The TiN average particle diameter was measured by extracting TiN by a replica method, photographing the precipitate with a transmission electron microscope, and using 500 microsamples.
高周波焼入れは、平板(幅35mm×長さ300mm)に対し高周波コイルを移動させながら加熱・焼入れを実施した。図1に高周波焼入れの実施態様を示す。この時の加熱温度は900℃の低温とし、加熱時間は900℃までの通電時間を4秒とした。冷却開始時間は、通常行われる即冷却として0.5秒と、焼入れ安定性を評価するために3秒の2パターンを実施した。高周波焼入れ後の評価として、引張試験(JIS 5号、C方向(圧延方向に垂直))、シャルピー衝撃試験、旧オーステナイト粒径測定を実施した。 In the induction hardening, heating and hardening were performed while moving the high-frequency coil on a flat plate (width 35 mm x length 300 mm). FIG. 1 shows an embodiment of induction hardening. The heating temperature at this time was a low temperature of 900 ° C., and the heating time was 4 seconds for the energization time up to 900 ° C. The cooling start time was 0.5 seconds for the immediate cooling that is normally performed, and two patterns of 3 seconds for evaluating the quenching stability. As evaluation after induction hardening, a tensile test (JIS No. 5, C direction (perpendicular to the rolling direction)), Charpy impact test, and prior austenite particle size measurement were performed.
シャルピー衝撃試験は、図2に示すような試験片形状にて、試験温度:-50℃、n=3で行った。また、熱延板の板厚を1.2mmtに研削加工し、後述の冷延板と同一形状とした。なお、シャルピー衝撃試験値は、同一条件で試験を実施したJSC980Yレベルの0.4kgm以上を合格とした。 The Charpy impact test was performed with a test piece shape as shown in FIG. 2 at a test temperature of −50 ° C. and n = 3. In addition, the hot-rolled sheet was ground to a thickness of 1.2 mm to have the same shape as a cold-rolled sheet described later. In addition, the Charpy impact test value passed 0.4 kgm or more of the JSC980Y level tested on the same conditions.
旧オーステナイト粒径は、サンプルの板厚断面を研磨・腐食後、光学顕微鏡にてミクロ組織を撮影し、マイクロアナライザーを用いて平均粒径を測定した。 For the prior austenite particle size, after polishing and corroding the plate thickness section of the sample, the microstructure was photographed with an optical microscope, and the average particle size was measured using a microanalyzer.
上記より得られた結果を表3に示す。表3より、成分、B-(10.8/14)N*、TiN平均粒径、旧オーステナイト粒径が本範囲内であるNo.A、B、C、E、Gは、焼入れ後の特性として980MPa以上の強度を有し、焼入れ後の冷却開始時間にかかわらず安定してJSC980Y以上(0.4kgm以上)のシャルピー衝撃吸収エネルギーが得られ、優れた衝撃特性が得られていることが明らかとなった。特に、C、Si、Mn、P、Sが低く、sol.Alが0.03〜0.07%、Bが0.0005〜0.0020%であるNo.A、Bはシャルピー衝撃吸収エネルギーが0.5kgm以上であり、極めて優れた衝撃特性が得られていることがわかった。 The results obtained from the above are shown in Table 3. From Table 3, No. A, B, C, E, and G whose components, B- (10.8 / 14) N *, TiN average particle size, and prior austenite particle size are within this range are 980 MPa as the properties after quenching. With the above strength, it became clear that Charpy impact absorption energy of JSC980Y or more (0.4 kgm or more) was stably obtained regardless of the cooling start time after quenching, and excellent impact characteristics were obtained. . In particular, No. A and B with low C, Si, Mn, P, S, sol.Al 0.03-0.07%, B 0.0005-0.0020%, Charpy impact absorption energy is 0.5kgm or more It was found that the impact characteristics were obtained.
一方、Cが本発明範囲外で低いNo.Hは強度が低く、Cが本発明範囲外で高いNo.Iと、Si、Pが本発明範囲外で高いNo.Jと、Mn、Sが本発明範囲外で高いNo.Kは、シャルピー衝撃吸収エネルギーが低く、衝撃特性が劣化している。sol.Al、Nが本発明範囲外で高いNo.Lは、旧オーステナイト粒径が本発明範囲外で小さく、冷却開始時間が遅い場合、シャルピー衝撃吸収エネルギーが低く、衝撃特性が劣化している。Bが本発明範囲外で低く、かつ、B-(10.8/14)N*が本発明範囲外であるNo.Mは、冷却開始時間が遅い場合、フェライトが生成し衝撃特性が劣化している。Tiが本発明範囲外で低く、TiN平均粒径が本発明範囲外で小さく、かつ、B-(10.8/14)N*が本発明範囲外であるNo.Nは、TiNの量が少なくオーステナイト粒成長の抑制がなされず、シャルピー衝撃吸収エネルギーが低く衝撃特性が劣化している。Tiが本発明範囲外で高く、かつ、TiN平均粒径が本発明範囲外で大きいNo.Oは、旧オーステナイト粒径が小さく、冷却開始時間が遅い場合、シャルピー衝撃吸収エネルギーが低く衝撃特性が劣化している。巻取温度が本発明範囲外で高いNo.Dは、焼入時にセメンタイトが溶け残り、シャルピー衝撃吸収エネルギーが低く衝撃特性が劣化している。焼鈍温度が本発明範囲外で高いNo.Fは、部分的にパーライトが生成し、シャルピー衝撃吸収エネルギーが低く衝撃特性が劣化している。 On the other hand, No.H having a low C outside the scope of the present invention has low strength, No.I having a high C outside the scope of the present invention, and No.J having a high Si and P outside the scope of the present invention, Mn, and S. No. K high outside the scope of the present invention has low Charpy impact absorption energy and deteriorates impact characteristics. sol.Al, N, which is high outside the scope of the present invention, No. L, when the prior austenite grain size is small outside the scope of the present invention and the cooling start time is slow, the Charpy impact absorption energy is low and the impact characteristics are degraded. . No. M in which B is outside the scope of the present invention and B- (10.8 / 14) N * is outside the scope of the present invention, when the cooling start time is slow, ferrite is generated and the impact characteristics are deteriorated. . No. N in which Ti is outside the scope of the present invention, TiN average particle size is small outside the scope of the present invention, and B- (10.8 / 14) N * is outside the scope of the present invention is low in the amount of TiN and austenite Grain growth is not suppressed, Charpy impact absorption energy is low, and impact characteristics are deteriorated. No.O with high Ti outside the scope of the present invention and large TiN average grain diameter outside the scope of the present invention has low Charpy impact absorption energy and low impact characteristics when the prior austenite grain size is small and the cooling start time is slow. It has deteriorated. In No. D, which has a high coiling temperature outside the range of the present invention, cementite remains undissolved during quenching, the Charpy impact absorption energy is low, and the impact characteristics are deteriorated. In No. F, which has a high annealing temperature outside the range of the present invention, pearlite is partially generated, the Charpy impact absorption energy is low, and the impact characteristics are deteriorated.
表1に示す鋼番1から13の化学成分組成を有する鋼を溶製し、次いで表4に示す製造条件に従って熱間圧延-冷間圧延-焼鈍を行い、1.2mmtの冷延板を製造した。このようにして製造した冷延板について、実施例1と同様に、引張試験、TiNの平均粒径測定、および高周波焼入れ特性を調査した。結果を表5に示す。 Steel having the chemical composition of steel Nos. 1 to 13 shown in Table 1 was melted, and then hot-rolled-cold-rolled-annealed according to the manufacturing conditions shown in Table 4 to produce 1.2 mmt cold-rolled sheets . For the cold-rolled sheet thus produced, the tensile test, the TiN average particle diameter measurement, and the induction hardening characteristics were investigated in the same manner as in Example 1. The results are shown in Table 5.
表5より、熱延鋼板の場合と同様に、成分、B-(10.8/14)N*、TiN平均粒径、旧オーステナイト粒径が本範囲内であるNo.a、c、d、e、hは、焼入れ後の特性として980MPa以上の強度を有し、焼入れ後の冷却開始時間にかかわらず安定してJSC980Y以上(0.4kgm以上)のシャルピー衝撃吸収エネルギーが得られ、優れた衝撃特性が得られていることが明らかとなった。特に、C、Si、Mn、P、Sが低く、sol.Alが0.03〜0.07%、Bが0.0005〜0.0020%であるNo.a、c、dはシャルピー衝撃吸収エネルギーが0.5kgm以上であり、極めて優れた衝撃特性が得られていることがわかった。 From Table 5, as in the case of hot-rolled steel sheet, No. a, c, d, e, where the component, B- (10.8 / 14) N *, TiN average particle diameter, and prior austenite particle diameter are within this range h has a strength of 980 MPa or more as a characteristic after quenching, and stable Charpy impact absorption energy of JSC980Y or more (0.4 kgm or more) is obtained regardless of the cooling start time after quenching, and excellent impact characteristics are obtained. It became clear that In particular, No. a, c, d with low C, Si, Mn, P, S, sol. Al of 0.03-0.07%, B of 0.0005-0.0020% have Charpy impact absorption energy of 0.5 kgm or more, It was found that excellent impact characteristics were obtained.
一方、TiN平均粒径が本発明範囲外で小さいNo.bは旧オーステナイト粒径が小さく、冷却開始時間が遅い場合、シャルピー衝撃吸収エネルギーが低く衝撃特性が劣化している。また、Cが本発明範囲外で低いNo.iは強度が低く、Cが本発明範囲外で高いNo.jと、Si、Pが本発明範囲外で高いNo.kと、Mn、Sが本発明範囲外で高いNo.lは、シャルピー衝撃吸収エネルギーが低く、衝撃特性が劣化している。sol.Al、Nが本発明範囲外で高いNo.mは、旧オーステナイト粒径が本発明範囲外で小さく、冷却開始時間が遅い場合、シャルピー衝撃吸収エネルギーが低く、衝撃特性が劣化している。Bが本発明範囲外で低く、かつ、B-(10.8/14)N*が本発明範囲外であるNo.nは、冷却開始時間が遅い場合、フェライトが生成し衝撃特性が劣化している。Tiが本発明範囲外で低く、TiN平均粒径が本発明範囲外で小さく、かつ、B-(10.8/14)N*が本発明範囲外であるNo.oは、TiNの量が少なくオーステナイト粒成長の抑制がなされず、シャルピー衝撃吸収エネルギーが低く衝撃特性が劣化している。Tiが本発明範囲外で高く、かつ、TiN平均粒径が本発明範囲外で大きいNo.pは、旧オーステナイト粒径が小さく、冷却開始時間が遅い場合、シャルピー衝撃吸収エネルギーが低く衝撃特性が劣化している。巻取温度が本発明範囲外で高いNo.fは、焼入時にセメンタイトが溶け残り、シャルピー衝撃吸収エネルギーが低く衝撃特性が劣化している。冷間圧延後の焼鈍温度が本発明範囲外で高いNo.gは、部分的にパーライトが生成し、シャルピー衝撃吸収エネルギーが低く衝撃特性が劣化している。 On the other hand, No. b whose TiN average particle size is small outside the range of the present invention has low prior austenite particle size, and when the cooling start time is slow, the Charpy impact absorption energy is low and the impact characteristics are deteriorated. No.i having a low C outside the scope of the present invention is low in strength, No.j having a high C outside the scope of the present invention, No.k having a high Si and P outside the scope of the present invention, Mn, S No. 1 high outside the scope of the present invention has low Charpy impact absorption energy and deteriorated impact characteristics. sol.Al, N is No.m high outside the scope of the present invention, when the prior austenite grain size is small outside the scope of the present invention and the cooling start time is slow, the Charpy impact absorption energy is low and the impact characteristics are deteriorated. . No.n, where B is outside the scope of the present invention and B- (10.8 / 14) N * is outside the scope of the present invention, when the cooling start time is slow, ferrite is generated and the impact characteristics are degraded. . No.o in which Ti is outside the scope of the present invention, TiN average particle size is small outside the scope of the present invention, and B- (10.8 / 14) N * is outside the scope of the present invention is low in the amount of TiN and austenite Grain growth is not suppressed, Charpy impact absorption energy is low, and impact characteristics are deteriorated. No.p with high Ti outside the scope of the present invention and large TiN average grain diameter outside the scope of the present invention has a low Charpy impact absorption energy and low impact characteristics when the prior austenite grain size is small and the cooling start time is slow. It has deteriorated. No. f having a high coiling temperature outside the range of the present invention has no cementite melted during quenching, has low Charpy impact absorption energy, and has deteriorated impact characteristics. No.g, which has a high annealing temperature after cold rolling outside the range of the present invention, partially generates pearlite, has low Charpy impact absorption energy, and deteriorates impact characteristics.
Claims (4)
B-(10.8/14)N*≧0.0005％
N*=N-(14/48)Ti、但し、右辺≦0の場合、N*=0
を満足し、残部Fe及び不可避不純物からなり、鋼中析出物であるTiNの平均粒径が0.06〜0.30μmであり、かつ焼入れ後の旧オーステナイト粒径が2〜25μmであることを特徴とする焼入れ後の衝撃特性に優れる熱延鋼板。 Steel: mass%, C: 0.10 to 0.37%, Si: 1% or less, Mn: 1.4% or less, P: 0.1% or less, S: 0.03% or less, sol.Al: 0.01 to 0.1%, N: 0.0005 -0.0050%, Ti: 0.005-0.05%, B: 0.0003-0.0050%,
B- (10.8 / 14) N * ≧ 0.0005%
N * = N- (14/48) Ti, provided that N * = 0 if right side ≤ 0
And the balance is composed of the remaining Fe and inevitable impurities, the average grain size of TiN being a precipitate in steel is 0.06 to 0.30 μm, and the prior austenite grain size after quenching is 2 to 25 μm Hot rolled steel sheet with excellent impact properties after quenching.
Priority Applications (1)
| Application Number | Priority Date | Filing Date | Title |
|---|---|---|---|
| JP2005233094A JP4265582B2 (en) | 2001-02-07 | 2005-08-11 | Hot-rolled steel sheet with excellent impact properties after quenching and method for producing the same |
Applications Claiming Priority (2)
| Application Number | Priority Date | Filing Date | Title |
|---|---|---|---|
| JP2001030355 | 2001-02-07 | ||
| JP2005233094A JP4265582B2 (en) | 2001-02-07 | 2005-08-11 | Hot-rolled steel sheet with excellent impact properties after quenching and method for producing the same |
Related Parent Applications (1)
| Application Number | Title | Priority Date | Filing Date |
|---|---|---|---|
| JP2001268316A Division JP4123748B2 (en) | 2001-02-07 | 2001-09-05 | Thin steel plate with excellent impact properties after quenching and method for producing the same |
Publications (2)
| Publication Number | Publication Date |
|---|---|
| JP2005325454A JP2005325454A (en) | 2005-11-24 |
| JP4265582B2 true JP4265582B2 (en) | 2009-05-20 |
Family
ID=35472008
Family Applications (1)
| Application Number | Title | Priority Date | Filing Date |
|---|---|---|---|
| JP2005233094A Expired - Fee Related JP4265582B2 (en) | 2001-02-07 | 2005-08-11 | Hot-rolled steel sheet with excellent impact properties after quenching and method for producing the same |
Country Status (1)
| Country | Link |
|---|---|
| JP (1) | JP4265582B2 (en) |
Cited By (3)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| WO2013102986A1 (en) | 2012-01-05 | 2013-07-11 | Ｊｆｅスチール株式会社 | High carbon hot-rolled steel sheet and method for producing same |
| WO2013102987A1 (en) | 2012-01-06 | 2013-07-11 | Ｊｆｅスチール株式会社 | High carbon hot-rolled steel sheet and method for producing same |
| WO2013102982A1 (en) | 2012-01-05 | 2013-07-11 | Ｊｆｅスチール株式会社 | High carbon hot-rolled steel sheet with excellent hardenability and minimal in-plane anisotropy, and method for producing same |
Families Citing this family (1)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| JP4867320B2 (en) * | 2005-12-05 | 2012-02-01 | 住友金属工業株式会社 | High strength steel member and manufacturing method thereof |
-
2005
- 2005-08-11 JP JP2005233094A patent/JP4265582B2/en not_active Expired - Fee Related
Cited By (5)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| WO2013102986A1 (en) | 2012-01-05 | 2013-07-11 | Ｊｆｅスチール株式会社 | High carbon hot-rolled steel sheet and method for producing same |
| WO2013102982A1 (en) | 2012-01-05 | 2013-07-11 | Ｊｆｅスチール株式会社 | High carbon hot-rolled steel sheet with excellent hardenability and minimal in-plane anisotropy, and method for producing same |
| KR20140111002A (en) | 2012-01-05 | 2014-09-17 | 제이에프이 스틸 가부시키가이샤 | High carbon hot-rolled steel sheet with excellent hardenability and minimal in-plane anisotropy, and method for producing same |
| US10323293B2 (en) | 2012-01-05 | 2019-06-18 | Jfe Steel Corporation | High-carbon hot rolled steel sheet with excellent hardenability and small in-plane anistropy and method for manufacturing the same |
| WO2013102987A1 (en) | 2012-01-06 | 2013-07-11 | Ｊｆｅスチール株式会社 | High carbon hot-rolled steel sheet and method for producing same |
Also Published As
| Publication number | Publication date |
|---|---|
| JP2005325454A (en) | 2005-11-24 |
Similar Documents
| Publication | Publication Date | Title |
|---|---|---|
| EP1571230B1 (en) | High tensile strength cold rolled steel sheet having excellent strain age hardening characteristics and the production thereof | |
| JP4772927B2 (en) | High-strength steel sheet, hot-dip galvanized steel sheet, alloyed hot-dip galvanized steel sheet having excellent fatigue characteristics and elongation and impact characteristics, and methods for producing them | |
| KR100513991B1 (en) | Method for production of thin steel sheet | |
| US6692584B2 (en) | High tensile cold-rolled steel sheet excellent in ductility and in strain aging hardening properties, and method for producing the same | |
| JP4265545B2 (en) | High tensile cold-rolled steel sheet with excellent strain age hardening characteristics and method for producing the same | |
| WO2015080242A1 (en) | Hot-formed steel sheet member, method for producing same, and steel sheet for hot forming | |
| KR102654714B1 (en) | High-strength member, method of manufacturing high-strength member, and method of manufacturing steel plate for high-strength member | |
| JP2017048412A (en) | Hot-dip galvanized steel sheet, alloyed hot-dip galvanized steel sheet and production methods therefor | |
| CA2941202A1 (en) | Method for producing a high-strength flat steel product | |
| JP2007162078A (en) | High strength steel sheet and production method | |
| MX2011002559A (en) | High-strength steel plate and manufacturing method thereof. | |
| JP2009035814A (en) | High-strength hot-dip galvanized steel sheet and process for producing the same | |
| KR20210091755A (en) | Hot rolled steel strip and manufacturing method thereof | |
| JP2014019928A (en) | High strength cold rolled steel sheet and method for producing high strength cold rolled steel sheet | |
| KR102541248B1 (en) | High-ductility and high-strength electro-galvanized steel sheet and manufacturing method thereof | |
| JP4983082B2 (en) | High-strength steel and manufacturing method thereof | |
| JP4123748B2 (en) | Thin steel plate with excellent impact properties after quenching and method for producing the same | |
| CN115461482B (en) | Steel sheet, component, and method for manufacturing same | |
| JP4265582B2 (en) | Hot-rolled steel sheet with excellent impact properties after quenching and method for producing the same | |
| JPH06145894A (en) | High strength hot rolled steel sheet excellent in ductility and delayed fracture resistance and its production | |
| JP4519373B2 (en) | High-tensile cold-rolled steel sheet excellent in formability, strain age hardening characteristics and room temperature aging resistance, and method for producing the same | |
| JP4333352B2 (en) | Method for producing high-strength cold-rolled steel sheet excellent in ductility and stretch flangeability | |
| JP4265583B2 (en) | Cold-rolled steel sheet having excellent toughness after quenching and method for producing the same | |
| JP2005256044A (en) | High-strength cold rolled steel sheet having excellent workability and post-painting corrosion resistance and manufacturing method therefor | |
| JP3952714B2 (en) | Hot-rolled steel sheet having excellent toughness after quenching and manufacturing method thereof |
Legal Events
| Date | Code | Title | Description |
|---|---|---|---|
| A621 | Written request for application examination |
Free format text: JAPANESE INTERMEDIATE CODE: A621 Effective date: 20050811 |
|
| RD01 | Notification of change of attorney |
Free format text: JAPANESE INTERMEDIATE CODE: A7421 Effective date: 20060921 |
|
| A131 | Notification of reasons for refusal |
Free format text: JAPANESE INTERMEDIATE CODE: A131 Effective date: 20080909 |
|
| A521 | Written amendment |
Free format text: JAPANESE INTERMEDIATE CODE: A523 Effective date: 20081106 |
|
| TRDD | Decision of grant or rejection written | ||
| A01 | Written decision to grant a patent or to grant a registration (utility model) |
Free format text: JAPANESE INTERMEDIATE CODE: A01 Effective date: 20090127 |
|
| A01 | Written decision to grant a patent or to grant a registration (utility model) |
Free format text: JAPANESE INTERMEDIATE CODE: A01 |
|
| A61 | First payment of annual fees (during grant procedure) |
Free format text: JAPANESE INTERMEDIATE CODE: A61 Effective date: 20090209 |
|
| R150 | Certificate of patent or registration of utility model |
Ref document number: 4265582 Country of ref document: JP Free format text: JAPANESE INTERMEDIATE CODE: R150 Free format text: JAPANESE INTERMEDIATE CODE: R150 |
|
| FPAY | Renewal fee payment (event date is renewal date of database) |
Free format text: PAYMENT UNTIL: 20120227 Year of fee payment: 3 |
|
| FPAY | Renewal fee payment (event date is renewal date of database) |
Free format text: PAYMENT UNTIL: 20120227 Year of fee payment: 3 |
|
| FPAY | Renewal fee payment (event date is renewal date of database) |
Free format text: PAYMENT UNTIL: 20130227 Year of fee payment: 4 |
|
| FPAY | Renewal fee payment (event date is renewal date of database) |
Free format text: PAYMENT UNTIL: 20130227 Year of fee payment: 4 |
|
| LAPS | Cancellation because of no payment of annual fees |