201209179 六、發明說明： 【發明所屬之技術領域】 本發明係關於主要適用於汽車產業之車架構件用及底 盤構件用之成形性優異的高強度冷軋鋼板及其製造方法。 【先前技術】 近年來’基於地球環保的觀點，改善汽車之燃料消耗 率成爲重要的課題。因此’藉由將車體材料高強度化而謀 求薄型化’且將車體本身輕量化的發展變得活絡起來。然 而’鋼板的高強度化會造成延性變差（亦即成形加工性變 差）’因此期望能開發出兼具高強度和高加工性的材料。 針對這種要求’迄今爲止已開發出肥粒鐵、麻田散鐵 雙相鋼（以下稱DP鋼）、利用殘留沃斯田鐵的變態誘發 塑性之TRIP鋼等各種的複合組織型冷軋鋼板。 例如，專利文獻1揭示一種加工性優異的高強度鋼板 之製造方法，是藉由添加多量的Si來確保殘留沃斯田鐵 而達成高延性。 然而，該等的D P鋼和T RIP鋼，雖然延伸特性優異 但存在著擴孔性差的問題。擴孔性，是將加工孔部擴大而 進行凸緣成形時之加工性指標，其和延伸特性都是高強度 鋼板所要求的重要特性。 作爲凸緣成形性優異的冷軋鋼板製造方法，專利文獻 2所揭示的技術，是在退火均熱後進行淬火-回火而形成 肥粒鐵和回火麻田散鐵之複合組織，藉此來提昇擴孔性。 201209179 然而’這種技術雖能獲得高擴孔性，但有延伸性變差的問 如此般’依據習知的技術，並無法獲得兼具優異的延 伸特性及凸緣成形性之冷軋鋼板。 [專利文獻1]日本特開平2-101117號公報 [專利文獻2]日本特開2004-256872號公報 【發明內容】 本發明是著眼於上述問題點而開發完成的，其目的是 爲了提供一種延伸特性及凸緣成形性優異之高強度冷軋鋼 板及其製造方法。 本發明人等爲了達成上述課題而製造出延伸特性及凸 緣成形性優異之高強度冷軋鋼板，從鋼板組成及微組織的 觀點進行深入的硏究。結果了解到，藉由適當地調整合金 元素，在退火過程中從均熱溫度冷卻時強制冷卻至 1 5 0~3 5 0°C的溫度範圍，然後進行再加熱，藉此獲得以面 稂率計含有肥粒鐵20%以上且回火麻田散鐵1 0~60%，以 體積率計含有殘留沃斯田鐵3~1 5 %之組織，而能具備高 延性及凸緣成形性。 一般而言，若存在有殘留沃斯田鐵，利用殘留沃斯田 鐵之TRIP效果能提昇延性。但已知，附加的應變會使殘 留沃斯田鐵變態而產生非常硬質的麻田散鐵，結果其與主 相之肥粒鐵的硬度差變大而造成凸緣成形性變差。 然而在本發明之成分及組織構造下，可兼具高延性和 -6 - 201209179 高凸緣成形性。即使存在有殘留沃斯田鐵仍能具有高凸緣 成形性的詳細理由雖不甚明白’但藉由使殘留沃斯田鐵和 回火麻田散鐵共存，應可減低殘留沃斯田鐵對凸緣成形性 的不良影響。 再者了解到，藉由獲得麻田散鐵'回火麻田散鐵、殘 留沃斯田鐵所構成之低溫變態相的平均結晶粒徑爲3 // m 以下的鋼板組織，除了具有高加工性還能提昇耐衝擊特性 〇 本發明是根據上述認知而開發完成的，其要旨說明如 下。 第一發明之加工性及耐衝擊性優異之高強度冷軋鋼板 ’以質量％計，係含有 C: 0.05~0.3%、Si: 0.3 〜2.5%、 Mn: 0.5-3.5%、P: 0.003〜0.100%、S: 0,02% 以下、A1 :0_010〜0·5%，剩餘部分爲鐵及不可避免的雜質所構成 ’且具有：以面積率計含有肥粒鐵20%以上、回火麻田 散鐵10〜60%、麻田散鐵0〜10%，以體積率計含有殘留沃 斯田鐵3〜1 5 %之組織。 第二發明，是在第一發明所記載之加工性及耐衝擊性 優異之闻強度冷軋鋼板中，具有前述麻田散鐵、回火麻田 散鐵、殘留沃斯田鐵所構成之低溫變態相的平均結晶粒徑 爲3 # m以下的組織。 第二發明’是在第一發明或第二發明所記載之加工性 及耐衝擊性優異之高強度冷軋鋼板中，進一步含有以質量 /6 b十，运自 Cr. 0.005〜2.00%、Mo: 0.005 〜2.00%、V: 201209179 0.005〜2.00%、Ni:〇〇〇5〜2〇()%、cu:0.005~2.00%* 之1種或2種以上的元素。 第四發明’是在第一〜第三發明所記載之加工性及耐 衝擊性優異之高強度冷軋鋼板中，進一步含有以質量％計 ，選自 Ti: 0.01 〜0.20%、Nb: 0.01 〜0.20% 中之 1 種或 2 種的元素。 第五發明’是在第—〜第四發明所記載之加工性及耐 衝擊性優異之高強度冷軋鋼板中，進一步含有以質量％計 ，B ： 0.0002〜0.005 % 〇 第六發明’是在第一〜第五發明所記載之加工性及耐 衝擊性優異之高強度冷軋鋼板中，進一步含有以質量％計 ，選自 Ca : 0.001 〜0.005 %、REM : 0 · 0 0 1 〜0.0 0 5 % 中之 1 種或2種的元素。 第七發明之加工性及耐衝擊性優異之高強度冷軋鋼板 之製造方法，其特徵在於：將具有第--第六發明所記載 的成分之鋼胚實施熱軋及冷軋而製造出冷軋鋼板，在將該 冷軋鋼板實施連續退火時，在750°C以上的溫度保持1〇 秒以上後，從750°C以平均10°C /s以上的冷卻速度冷卻至 1 50~3 5 0 °C的溫度範圍後，加熱至 3 50〜600 °C並保持 1 0〜600秒後，冷卻至室溫。 第八發明，是在第七發明所記載之加工性及耐衝擊性 優異之高強度冷軋鋼板之製造方法中，在50(TC〜ACl變態 點以平均加熱速度1 〇°C /s以上進行昇溫。 依據本發明可獲得加工性優異的高強度冷軋鋼板，可 -8- 201209179 同時謀求汽車的輕量化及撞擊安全性的提昇，而對汽車車 體的高性能化有很大的幫助。 【實施方式】 以下具體地說明本發明。 1.關於成分組成 首先說明將本發明的鋼成分組成限定在上述範圍的理 由。關於成分％的表示，在沒有特別說明的情況是指質量 % 。 C ·· 0.05〜0.3% C是讓沃斯田鐵穩定化的元素，爲了容易生成肥粒鐵 以外的相而使鋼板強度提昇，又爲了將組織複合化而提昇 TS-EL平衡，其是必要的元素。C量未達0.05%時，即使 能謀求製造條件之最佳化仍難以確保肥粒鐵以外的相，TS xEL會降低。另一方面，若C量超過0.3%，熔接部及熱 影響部的硬化程度變大，熔接部之機械特性變差。基於此 觀點，將 C量設定在 0.05〜0.3%的範圍。較佳爲 0.08〜0· 1 5%的範圍。[Technical Field] The present invention relates to a high-strength cold-rolled steel sheet excellent in formability for use in a frame member for a vehicle industry and a chassis member, and a method for producing the same. [Prior Art] In recent years, improving the fuel consumption rate of automobiles has become an important issue based on the viewpoint of global environmental protection. Therefore, the development of the vehicle body itself is reduced by increasing the strength of the vehicle body material, and the development of the vehicle body itself is reduced. However, the increase in strength of the steel sheet causes deterioration in ductility (i.e., deterioration in formability). Therefore, it has been desired to develop a material having both high strength and high workability. In response to such a demand, various composite structure cold-rolled steel sheets such as ferrite iron, 麻田散铁 duplex steel (hereinafter referred to as DP steel), and TRIP steel using metamorphism-induced plasticity of residual Worth iron have been developed. For example, Patent Document 1 discloses a method for producing a high-strength steel sheet excellent in workability, in which a large amount of Si is added to secure residual Worstian iron to achieve high ductility. However, these D P steels and T RIP steels have excellent elongation characteristics but have a problem of poor hole expandability. The hole expandability is an index of workability when the machined hole portion is enlarged to form a flange, and the elongation characteristics are important characteristics required for a high-strength steel sheet. As a method for producing a cold-rolled steel sheet having excellent flange formability, the technique disclosed in Patent Document 2 is a composite structure in which quenching-tempering is performed after annealing soaking to form ferrite iron and tempered granita iron. Improve hole expandability. 201209179 However, although this technique can achieve high hole expandability, there is such a problem that the elongation is deteriorated. According to the conventional technique, a cold-rolled steel sheet having excellent elongation characteristics and flange formability cannot be obtained. [Patent Document 1] Japanese Laid-Open Patent Publication No. Hei No. Hei. No. Hei. No. Hei. No. Hei. No. 2004-256872. [Invention] The present invention has been developed in view of the above problems, and its purpose is to provide an extension. A high-strength cold-rolled steel sheet excellent in characteristics and flange formability and a method for producing the same. In order to achieve the above-mentioned problems, the present inventors have produced a high-strength cold-rolled steel sheet having excellent elongation properties and excellent flange formability, and have intensively studied from the viewpoint of steel sheet composition and microstructure. As a result, it has been found that by appropriately adjusting the alloying elements, it is forcibly cooled to a temperature range of 150 to 350 ° C during cooling from the soaking temperature during annealing, and then reheated, thereby obtaining a facet ratio. It contains 20% or more of ferrite iron and 10 to 60% of tempered granulated iron, and contains 3 to 15% of residual Worthite iron in volume ratio, and has high ductility and flange formability. In general, if there is residual Worthite iron, the TRIP effect of the residual Worthite iron can improve the ductility. However, it is known that the additional strain causes the residual Worth iron to be metamorphosed to produce a very hard granulated iron, and as a result, the difference in hardness from the ferrite of the main phase becomes large, resulting in deterioration of the flange formability. However, under the composition and structure of the present invention, both high ductility and high flange formability of -6 - 201209179 can be achieved. Even though there are detailed reasons for the high flange formability of the residual Worthite iron, it is not clear, but by residing the residual Worthite iron and the tempered granulated iron, the residual Worthite iron flange should be reduced. Adverse effects of formability. Furthermore, it is understood that the steel sheet structure having an average crystal grain size of 3 // m or less in the low-temperature metamorphic phase composed of the Matian loose iron 'returned to the Ma Tian loose iron and the residual Worth iron is not only high in workability but also high in workability. The present invention has been developed based on the above findings, and the gist thereof is as follows. The high-strength cold-rolled steel sheet having excellent workability and impact resistance according to the first invention contains C: 0.05 to 0.3%, Si: 0.3 to 2.5%, Mn: 0.5 to 3.5%, and P: 0.003 in terms of % by mass. 0.100%, S: 0,02% or less, A1: 0_010 to 0.5%, and the remainder is composed of iron and unavoidable impurities' and has 20% or more of ferrite iron in an area ratio, and tempered Ma Tian 10 to 60% of loose iron, 0 to 10% of granulated iron, and 3 to 15% of residual Worthite iron by volume. According to a second aspect of the invention, in the cold-rolled steel sheet having excellent workability and impact resistance as described in the first aspect of the invention, the low-temperature metamorphic phase composed of the above-mentioned 麻田散铁, tempered 麻田散铁, and residual Worth iron The average crystal grain size is 3 # m or less. The second invention is a high-strength cold-rolled steel sheet having excellent workability and impact resistance as described in the first invention or the second invention, further containing a mass of /6 b, and transported from Cr. 0.005 to 2.00%, Mo. : 0.005 to 2.00%, V: 201209179 0.005 to 2.00%, Ni: 〇〇〇5 to 2 〇 ()%, cu: 0.005 to 2.00%* of one or two or more elements. The fourth invention is a high-strength cold-rolled steel sheet having excellent workability and impact resistance as described in the first to third inventions, further comprising, by mass%, selected from the group consisting of Ti: 0.01 to 0.20%, and Nb: 0.01 〜 One or two elements of 0.20%. The fifth invention is a high-strength cold-rolled steel sheet having excellent workability and impact resistance as described in the first to fourth inventions, and further contains, by mass%, B: 0.0002 to 0.005 %. The high-strength cold-rolled steel sheet excellent in workability and impact resistance described in the first to fifth inventions further contains, in mass%, selected from Ca: 0.001 to 0.005%, and REM: 0 · 0 0 1 to 0.0 0 One or two of the 5 % elements. A method for producing a high-strength cold-rolled steel sheet having excellent workability and impact resistance according to the seventh aspect of the invention, characterized in that the steel preform having the components described in the sixth to sixth invention is subjected to hot rolling and cold rolling to produce cold When the cold-rolled steel sheet is continuously annealed, the rolled steel sheet is maintained at a temperature of 750 ° C or higher for 1 sec. or more, and then cooled from 750 ° C to a cooling rate of 10 ° C /s or more to 1 50 to 3 5 After the temperature range of 0 ° C, heat to 3 50 to 600 ° C and hold for 10 to 600 seconds, then cool to room temperature. According to a seventh aspect of the invention, in the method for producing a high-strength cold-rolled steel sheet having excellent workability and impact resistance according to the seventh aspect of the invention, the method is carried out at an average heating rate of 1 〇 ° C /s or more at 50 (TC to ACl transformation point). According to the present invention, a high-strength cold-rolled steel sheet excellent in workability can be obtained, and the weight reduction and impact safety of the automobile can be improved at the same time, and the performance of the automobile body is greatly improved. [Embodiment] The present invention will be specifically described below. 1. Component composition First, the reason why the steel component composition of the present invention is limited to the above range will be described. The component % is referred to as % by mass unless otherwise specified. ·· 0.05~0.3% C is an element that stabilizes the Worthite iron. It is necessary to increase the strength of the steel sheet in order to easily form a phase other than the ferrite iron, and to increase the TS-EL balance in order to recombine the structure. When the amount of C is less than 0.05%, it is difficult to ensure a phase other than the ferrite iron and the TS xEL is reduced even if the production conditions are optimized. On the other hand, if the amount of C exceeds 0.3%, the welded portion and the thermal image Sclerosis portion becomes large, the mechanical properties of the welded portion is deteriorated. Based on this standpoint, the amount of C is set in the range of 0.05~0.3%. Preferably in a range of · 15% of 0.08~0.
Si : 0.3〜2.5%Si : 0.3~2.5%
Si是對鋼的強化有幫助的元素。又其是肥粒鐵生成 元素，能促進沃斯田鐡中C之濃化及抑制碳化物的生成， ** 9 - 201209179 具有促進殘留沃斯田鐵生成的作用。Si量未達0.3%時， 其添加效果差，因此將下限定爲0.3 %。但添加過多會造 成表面特性、熔接性變差’因此Si含量設定爲2.5%以下 。較佳爲0.7〜2.0%的範圍。 Μη : 0.5 〜3.5% Μη是對鋼的強化有幫助的元素，可促進回火麻田散 鐵等的低溫變態相的生成。這種作用在Μη含量0.5%以 上時會出現。然而若Μη添加過多而超過3·1%，起因於 第二相分率之過度增加及固溶強化，造成肥粒鐵的延性明 顯劣化而使成形性變差。因此，Μη量設定在0.5 ~ 3.5 %的 範圍。較佳爲1 .5〜3.0%的範圍。 Ρ ： 0.003 〜0.100% Ρ是對鋼的強化有幫助的元素，此效果在含量0.003 %以上可獲得。但若添加過多而超過0.1 0 0 %，起因於晶 界偏析會引起脆化，造成耐衝擊性變差。因此Ρ量設定在 0.003 % ~0.1 00% 的範圍。 -10- 1 : 0.02% 以下 S會變成Μ n S等的夾雜物，構成耐衝擊性變差和沿著 熔接部的金屬流裂開的原因，其含量越低越好，基於製造 成本的觀點設定成0.02%以下。 201209179 A1 ： 0.010〜0.5% A1具有脫氧劑的作用，是對鋼的清淨度有幫助的元 素，較佳爲在脫氧步驟中添加。在此，若A1量未達0.01 %，其添加效果差，因此將下限定爲0.01%。但若大量添 加，連續鑄造時之鋼片龜裂發生的危險性增高而使製造性 變差。因此A1添加量的上限定爲〇. 5 %。 本發明之高強度冷軋鋼板，是以上述成分組成作爲基 本成分，剩餘部分爲鐵及不可避免的雜質所構成，按照所 期望的特性，可適當地含有以下所敘述的成分。 選自 Cr : 0.005 〜2.00 % 、Mo : 0.005-2.00 % ' V ： 0.005 〜2.00%、Ni : 0.005-2.00 %、Cu ： 0.005 〜2.00% 中 之1種或2種以上 C r、Μ ο、V、N i、C u，在從退火溫度冷卻時可抑制珍 珠狀組織的生成，可促進低溫變態相的生成而有助於鋼的 強化。此效果，藉由讓C r、Μ ο、V、N i、C u之至少1種 的含量爲0.005%以上即可獲得。但若Cr' Mo、V、Ni、 Cu各成分超過 2.00%，其效果達飽和而造成成本上昇。 因此將 C r、Μ ο、V、N i、C u 的含量分別設定在 0.005〜2.00%的範圍。 選自Ti: 0.01〜0.20%、Nb: 0.01〜0.20%中之1種或2種Si is an element that contributes to the strengthening of steel. It is also a ferrite-iron-forming element that promotes the concentration of C in the Wostian sorghum and inhibits the formation of carbides. ** 9 - 201209179 has the effect of promoting the formation of residual Worth iron. When the amount of Si is less than 0.3%, the effect of addition is poor, so the lower limit is made 0.3%. However, excessive addition causes surface characteristics and weldability deteriorates. Therefore, the Si content is set to 2.5% or less. It is preferably in the range of 0.7 to 2.0%. Μη : 0.5 to 3.5% Μη is an element which contributes to the strengthening of steel and promotes the formation of low-temperature metamorphic phases such as tempered granules. This effect occurs when the Μη content is 0.5% or more. However, if Μη is excessively added and exceeds 3.1%, the excessive increase in the second phase fraction and solid solution strengthening result in deterioration of the ductility of the ferrite iron and deterioration of formability. Therefore, the amount of Μη is set in the range of 0.5 to 3.5%. It is preferably in the range of 1.5 to 3.0%. Ρ : 0.003 ～0.100% Ρ is an element that contributes to the strengthening of steel. This effect is obtained at a content of 0.003% or more. However, if it is added too much and exceeds 0.10%, segregation due to grain boundary may cause embrittlement, resulting in deterioration of impact resistance. Therefore, the amount is set in the range of 0.003 % ~ 0.1 00%. -10- 1 : 0.02% or less S becomes an inclusion such as Μ n S, which causes deterioration of impact resistance and cracking of the metal flow along the welded portion. The lower the content, the better, based on the viewpoint of manufacturing cost. Set to 0.02% or less. 201209179 A1 : 0.010 to 0.5% A1 has a function as a deoxidizer and is an element which contributes to the cleanliness of steel, and is preferably added in the deoxidation step. Here, if the amount of A1 is less than 0.01%, the effect of addition is poor, so the lower limit is made 0.01%. However, if it is added in a large amount, the risk of cracking of the steel sheet during continuous casting increases, and the manufacturability deteriorates. Therefore, the upper limit of the amount of A1 added is 〇. 5 %. The high-strength cold-rolled steel sheet according to the present invention comprises the above-mentioned component composition as a basic component, and the remainder is composed of iron and unavoidable impurities, and the components described below can be appropriately contained in accordance with desired characteristics. It is selected from Cr: 0.005 to 2.00%, Mo: 0.005-2.00% 'V: 0.005 to 2.00%, Ni: 0.005-2.00%, Cu: 0.005 to 2.00%, one or more of C r, Μ ο, V, N i , and C u can suppress the formation of a pearly structure when cooled from the annealing temperature, and promote the formation of a low-temperature metamorphic phase to contribute to the strengthening of steel. This effect can be obtained by making the content of at least one of C r, ο ο, V, N i and C u 0.005% or more. However, if the components of Cr' Mo, V, Ni, and Cu exceed 2.00%, the effect is saturated and the cost increases. Therefore, the contents of C r, ο ο, V, N i , and Cu are set in the range of 0.005 to 2.00%, respectively. One or two selected from the group consisting of Ti: 0.01 to 0.20% and Nb: 0.01 to 0.20%
Ti、Nb會形成碳氮化物，藉由析出強化具有將鋼高 強度化的作用。這種效果在含量爲0.01%以上時會出現。 -11 - 201209179 另一方面，Ti、Nb各含量即使超過0.20%，會因過度的 高強度化而造成延性降低。因此Ti、Nb的含量分別設定 在0.01〜0.20%的範圍。 B ： 0.0002-0.005 % B可抑制從沃斯田鐵晶界生成肥粒鐵，具有讓強度上 昇的作用。此效果在含量0.0002 %以上可獲得。但若B Μ超過0.005%，其效果達飽和而造成成本上昇。因此將 Β含量設定在0.0002〜0.005%的範圍。 選自 Ca: 0.001 〜0.005 %、REM: 0.001 〜0.005 % 中之 1 種 或2種Ti and Nb form carbonitrides, and the precipitation strengthening enhances the strength of the steel. This effect occurs when the content is 0.01% or more. -11 - 201209179 On the other hand, even if the content of Ti and Nb exceeds 0.20%, the ductility is lowered due to excessive strength. Therefore, the contents of Ti and Nb are set in the range of 0.01 to 0.20%, respectively. B : 0.0002-0.005 % B inhibits the formation of ferrite iron from the Worthfield iron grain boundary and has the effect of increasing the strength. This effect is obtained at a content of 0.0002% or more. However, if B Μ exceeds 0.005%, the effect is saturated and the cost increases. Therefore, the niobium content is set in the range of 0.0002 to 0.005%. One or two selected from the group consisting of Ca: 0.001 to 0.005 %, REM: 0.001 to 0.005 %
Ca、REM都能控制硫化物的形態而具有改善加工性 的效果，按照需要能以0.001%以上的含量含有Ca、REM 之1種或2種。但過添加過多會對清淨度造成不良影響， 因此分別設定在0.005%以下。 2.關於組織 接著說明鋼的組織。 肥粒鐵的面積率：20%以上 若肥粒鐵的面積率未達20%，TSxEL會降低，因此 設定爲20%以上。又較佳爲50%以上。 -12- 201209179 回火麻田散鐵的面積率：10〜60% 回火麻田散鐵，是將麻田散鐵在A c,變態點以下（較 佳爲比ACl變態點更低的溫度）加熱而獲得之高差排密度 的肥粒鐵和雪明碳鐵（c e m e n t i t e )之複合組織，對鋼的弓隹 化有幫助。此外，將麻田散鐵在超過Ac,變態點的溫度γ 加熱而獲得的組織，是在肥粒鐵中不含雪明碳鐵組織，| 本上是和本發明之回火麻田散鐵不同的。 此外，回火麻田散鐵對擴孔性的不良影響比麻田散鐵 小，是不致明顯降低擴孔性而有助於確保強度之相。回& 麻田散鐵的面積率未達10%時，要確保強度變困難，若 超過60%時TSxEL降低，因此將回火麻田散鐵的面積率 設定在10〜60%。 麻田散鐵的面積率：0〜10% 麻田散鐵有助於鋼的高強度化，若其面積率超過10 %會造成凸緣成形性明顯降低。因此將麻田散鐵的面積率 設定在0〜10%。 殘留沃斯田鐵的體積率：3〜15% 殘留沃斯田鐵不僅有助於鋼的強化，且有助於鋼之 TSxEL的提高。這種效果在體積率3%以上可獲得。此外 ，若殘留沃斯田鐵超過1 5 %，擴孔性會變差。因此將殘 留沃斯田鐵的體積率設定在3 ~ 1 5 %。 -13- 201209179 麻田散鐵、回火麻田 態相的平均結晶粒徑 麻田散鐵、回火 溫變態相，有助於耐 變態相微細地分散， 均結晶粒徑爲3 // m 相的平均結晶粒徑設 此外，作爲麻田 以外的相，雖然可能 足上述的相構造就沒 觀點，珍珠狀組織含 3 .關於製造條件 將調整成上述成 用連續鑄造法等獲得 ，進行連續退火。關 有特別的限定，以下 鑄造條件 所使用之鋼胚’ 造法來製造，但藉由 此外，在製造鋼胚之 除了此習知方法以外 溫而以溫片的狀態插 散鐵、殘留沃斯田鐵所構成之低溫變 :3 " m以下 麻田散鐵、殘留沃斯田鐵所構成之低 衝擊特性的提昇。特別是藉由讓低溫 可提昇耐衝擊特性，低溫變態相的平 以下時其效果顯著。因此將低溫變態 定爲3 " m以下。 散鐵、回火麻田散鐡、殘留沃斯田鐵 含有珍珠狀組織及變韌鐵，但只要滿 有問題。但基於確保延性及擴孔性的 量宜爲3 %以下。 分組成之鋼使用轉爐等進行熔製，使 鋼胚。將該鋼胚材實施熱軋及冷軋後 於鑄造、熱軋、冷軋之製造方法並沒 是說明較佳的製造方法。 爲了防止成分之微偏析宜採用連續鑄 造塊法、薄鋼胚鑄造法來製造亦可。 後，一旦冷卻至室溫然後再度加熱， ，也能毫無問題地採用：不冷卻到室 入加熱爐，或是進行稍微的保溫之後 14- 201209179 馬上實施輥軋之直送輥軋、直接輥軋等之省能源的製程。 熱軋條件 <鋼胚加熱溫度：1 1 〇 〇 °c以上> 鋼胚加熱溫度，從能量的觀點宜爲低溫加熱，當加熱 溫度未達1 1 00 °C時會產生：碳化物無法充分地固溶、輥 軋負荷的增大造成熱軋時發生阻礙的可能性增加等的問題 。此外’基於氧化重量的增加所伴生之氧化鱗皮損失之增 加等，鋼胚加熱溫度宜爲1 3 0 0 t以下。 又基於即使降低鋼胚加熱溫度仍能防止熱軋時的阻礙 之觀點’亦可活用將板片（sheet bar)加熱之所謂板片加 熱器。 <精軋結束溫度：Ar3變態點以上> 若精軋結束溫度未達Ar3變態點，在輥軋中會生成肥 粒鐵和沃斯田鐵，在鋼板容易生成條帶狀組織，該條帶狀 組織即使在冷軋後或退火後仍會殘留，而可能使材料特性 產生異向性’或成爲加工性變差的原因。因此，精軋結束 溫度宜爲Ar3變態點以上。 〈捲取溫度：450〜700°C > 若捲取溫度未達450 °C，捲取溫度的控制變難而容易 產生溫度不均，結果可能發生冷軋性變差等的問題。又捲 取溫度超過700 °C時，可能在鐵表層發生脫碳等的問題。 -15- 201209179 因此捲取溫度宜設定在450〜700°C的範圍。 又在本發明的熱軋步驟，爲了減少熱軋時之輥軋負荷 ，精軋的一部分或全部可實施潤滑輥軋。基於鋼板形狀的 均一化、材料均一化的觀點，進行潤滑輥軋是有效的。又 在潤滑輥軋時的摩擦係數較佳爲0.2 5〜0.10的範圍。又較 佳爲採用：將陸續送來的板片彼此接合而連續地進行精軋 之連續精軋製程。基於熱軋之作業穩定性的觀點，宜採用 連續精軋製程。 接著，較隹爲將熱軋鋼板表面之氧化鱗皮經由酸洗除 去之後，實施冷軋而成爲既定板厚的冷軋鋼板。在此之酸 洗條件及冷軋條件並沒有特別的限制，依據通常的方法即 可。冷軋之軋縮率宜爲40%以上。 < 500°C〜ACl變態點之平均加熱速度：10°C /s以上> 本發明的鋼，將再結晶溫度範圍（500°C〜AC|變態點 )之平均加熱速度設定爲1 〇 °C /s以上，可抑制加熱昇溫 時的再結晶，有助於在ACl變態點以上所生成之沃斯田鐵 的微細化，進而有助於退火冷卻後之組織的微細化，而能 使低溫變態相的平均粒徑成爲3 // m以下。 當平均加熱速度未達l〇°C/s時，在加熱昇溫時會產 生α的再結晶，導入肥粒鐵中的應變會被釋放而無法達成 充分的微細化。因此，將500°C ~AC|變態點之平均加熱速 度設定爲1〇°C/s以上。該平均加熱速度之較佳範圍爲20 l /s以上。 -16- 201209179 <在750 °C以上的溫度保持1〇秒以上> 當加熱溫度未達7 5 0 t：或保持時間未達1 0秒時，退 火時的沃斯田鐵的生成不足’在退火冷卻後無法確保充分 量的低溫變態相。保持溫度及保持時間的上限雖沒有特別 的規定，但保持溫度900 °C以上及保持時間600秒以上時 效果達飽和而造成成本上昇’因此較佳爲保持溫度未達 9 〇〇 °C及保持時間未達600秒。 〈從7 5 0 °C以1 0 °C /s以上的平均冷卻速度冷卻至1 5 0〜3 5 0 °C的溫度範圍> 從7 5 0°c的冷卻速度，未達l〇°C /S時會生成珍珠狀組 織，T S X E L及擴孔性會變差。因此，從7 5 0 °c的冷卻速度 設定爲1 0 °C /s以上。冷卻到達溫度條件是本技術最重要 的條件之一。冷卻停止時沃斯田鐵的一部分會變態成麻田 散鐵，剩餘的成爲未變態的沃斯田鐵。接著經由再加熱、 電鍍及合金化處理後，冷卻至室溫，麻田散鐵成爲回火麻 田散鐵，未變態沃斯田鐵則成爲殘留沃斯田鐵或麻田散鐵 。從退火之冷卻到達溫度越低，冷卻中生成的麻田散鐵量 越多，未變態沃斯田鐵量越少，因此藉由控制冷卻到達溫 度，來決定最後麻田散鐵及殘留沃斯田鐵和回火麻田散鐵 的面積率。 在冷卻到達溫度比35(TC更高的溫度，冷卻停止時麻 田散鐵變態不足且未變態沃斯田鐵量過多，最後會生成過 -17- 201209179 量之麻田散鐵或殘留沃斯田鐵，而使擴孔性變差。此外， 若冷卻到達溫度比1 50°C更低，冷卻中沃斯田鐵大部分變 態成麻田散鐵而使未變態沃斯田鐵量減少，無法獲得3 % 以上的殘留沃斯田鐵。因此將冷卻到達溫度設定在 1 5 0〜3 5 0 °C的範圍。關於冷卻的方法，只要能達成目標的 冷卻速度及冷卻停止溫度即可，可採用氣體噴射冷卻、噴 編冷卻、水冷、金屬淬火等任何的冷卻方法。 〈加熱至350〜600°C並保持10〜600秒〉 在150〜3 5 0°C的溫度範圍冷卻後，在3 5 0~600°C的溫 度範圍保持1 〇秒以上，藉此使前述冷卻時生成的麻田散 鐵回火而成爲回火麻田散鐵以提昇擴孔性，進一步使在前 述冷卻時未變態成麻田散鐵之未變態沃斯田鐵穩定化，而 最後獲得3 %以上的殘留沃斯田鐵以提昇延性。 藉由再加熱保持使未變態沃斯田鐵穩定化的機制之詳 細內容雖不甚明白，但應是從固溶有過飽和的C之麻田散 鐡讓C往未變態沃斯田鐵擴散而使未變態沃斯田鐵中的C 濃化，藉此使沃斯田鐵穩定化。這時，若麻田散鐡中之雪 明碳鐵的析出比C的擴散更早，未變態沃斯田鐵中C的 濃化會變得不足，因此讓雪明碳鐵的析出延後是重要的， 於是必須添加〇 . 3 %以上的S i。 再加熱溫度未達3 50°C時，麻田散鐵的回火及沃斯田 鐵之穩定化不足，擴孔性及延性變差。若再加熱溫度超過 600 °C，冷卻停止時未變態沃斯田鐵會變態成珍珠狀組織 -18- 201209179 ’最後無法獲得3 %以上的殘留沃斯田鐵。因此將加熱溫 度設疋爲35〇~6〇〇。〇。 在保持時間未達1 0秒時，沃斯田鐵的穩定化不足， 又若超過600秒，冷卻停止時的未變態沃斯田鐵會變態成 變韌鐵’最後無法獲得3 %以上的殘留沃斯田鐵。因此將 再加熱溫度設定爲3 50〜600°C的範圍，且將該溫度範圍的 保持時間設定爲1 0〜6 0 0秒。 又對於退火後的鋼板，爲了矯正形狀、調整表面粗度 等，可實施調質輥軋。此外，即使實施樹脂、油脂被覆、 各種塗裝等的處理也沒有任何的問題。 [實施例1] 將具有表1所示的成分組成且剩餘部分爲Fe及不可 避免的雜質構成之鋼，使用轉爐熔製並藉由連續鑄造法製 得鑄片。將所得的鑄片熱軋成板厚3 _0mm。熱軋條件，是 在精軋溫度900°C、輥軋後的冷卻速度10°C /s、捲取溫度 600 °C下進行。接著，將熱軋鋼板酸洗後，冷軋成板厚 1 · 2 m m，製得冷軋鋼板。 接著，對於該等冷軋鋼板，在連續退火線上依表2所 示的條件實施退火處理。 針對所製得的鋼板之截面微組織、拉伸特性及擴孔性 進行調查。結果如表3所示。 -19- 201209179 【1® (％¥M) 5 槲 粼 m s 粼 粼 粼 豳 m s 翻 镟 m 豳 錨 a 鎰 a 瑙 a g o o to g ο ο m ο ο o d - g o S ο a ο 乏 C4 ο > g ο ο 2 寸 ο ό ο Ζ S ο d 兮 ο ο g ο ο s ο ο s ο ο o o d S o o g ο ο s ο g o o o o s o o CM Ο ο < CO CO o o CO CO p GO CNJ ο ο S ο ο S ο o d CM s d ιη ο ο to g o CO s o a o o CO s o σ> S ο CO CO o o s o o in ο ο (Ο ο ο S ο ο CNJ o o o d g ο ο s o o s o o o o o o o C0 ο ο α s o o in S o CO ο ο g ο ο g ο ο g o d o d ο ο CO o o' CM o d CO o o CJ CM o o r- ο ο C CO csi o c\i CO τ· τ·» CM cJ σ> CO CNJ σ> 产 m evi 00 CO 31 31 W CNJ Ο ΙΟ ir> d - ir> σ> ό CsJ o CNJ CO CO o r— Ο o d r- 〇 d 00 5 ιη eg d CO o d CNJ o’ 兮 o ο τ» ο GO O d σ> O o II Γ-* 5 Ο ο i1ml1 m 翳 < CQ Ο Ω at u. 〇 X 一 -« 醒 ssffiHg 件sslle*璩哩 nf?ii -20 - 201209179 [表2]Both Ca and REM can control the form of the sulfide and have an effect of improving the workability, and if necessary, one or two types of Ca and REM can be contained in an amount of 0.001% or more. However, too much addition will adversely affect the cleanliness, so it is set to 0.005% or less. 2. About the organization Next, explain the organization of steel. Area ratio of ferrite iron: 20% or more If the area ratio of ferrite iron is less than 20%, TSxEL will decrease, so it is set to 20% or more. It is preferably 50% or more. -12- 201209179 Area ratio of tempered loose iron in Matian: 10~60% tempered granulated iron, which is heated under the A c, lower than the metamorphic point (preferably lower than the ACl metamorphic point) The composite structure of the ferrite iron and the cementite obtained by the high difference density density is helpful for the bowing of the steel. In addition, the structure obtained by heating the methadrite iron at a temperature γ exceeding the Ac, metamorphic point is a ferritic carbon-iron structure in the ferrite iron, which is different from the tempered granulated iron of the present invention. . In addition, the adverse effect of the tempered granulated iron on the reaming property is smaller than that of the granulated iron, which is a phase that does not significantly reduce the reaming property and contributes to the strength. When the area ratio of the granules of the granules is less than 10%, it is difficult to ensure the strength. If the TSxEL is lower than 60%, the area ratio of the tempered granules is set at 10 to 60%. Area ratio of granulated iron: 0~10% 麻田散铁 contributes to the high strength of steel. If the area ratio exceeds 10%, the formability of the flange is significantly reduced. Therefore, the area ratio of the granulated iron is set to 0 to 10%. The volume fraction of residual Worthite iron: 3~15% Residual Worthite iron not only contributes to the strengthening of steel, but also contributes to the improvement of TSxEL of steel. This effect is obtained at a volume fraction of 3% or more. In addition, if the remaining Worth iron exceeds 15%, the hole expandability will be deteriorated. Therefore, the volume fraction of the residual Worthfield iron is set at 3 to 15%. -13- 201209179 The average crystal grain size of the granulated iron and tempered sesame phase in the field, the tempering temperature and phase of the tempering phase, contributes to the fine dispersion of the resistant phase, and the average crystal grain size is 3 // m. In addition, as a phase other than the field, the phase other than the field may have no viewpoint, and the pearl-like structure contains 3. The production conditions are adjusted to the above-described continuous casting method or the like, and continuous annealing is performed. There is a special limitation, which is produced by the steel blank used in the following casting conditions, but by this, in addition to the conventional method, the steel embryo is inserted into the molten steel in the form of a warm sheet, and the residual Worth is left. The low temperature change formed by Tiantie: 3 " m below the lower impact characteristics of the Ma Tian loose iron and the residual Worth iron. In particular, by making the low temperature improve the impact resistance, the effect is remarkable when the low temperature metamorphic phase is below the flat. Therefore, the low temperature metamorphism is set to be 3 " m or less. The loose iron, the tempered granules, and the residual Worth iron contain pearly structure and toughened iron, but as long as there are problems. However, the amount of ductility and hole expandability is preferably 3% or less. The divided steel is melted using a converter or the like to make a steel blank. The production method of casting, hot rolling, and cold rolling after hot rolling and cold rolling of the steel preform is not a preferred manufacturing method. In order to prevent the microsegregation of the components, it is preferable to use a continuous casting block method or a thin steel blank casting method. After cooling to room temperature and then heating again, it can be used without any problem: after cooling to the furnace or after a slight heat preservation, 14-201209179 immediately performs rolling direct rolling, direct rolling Wait for the energy-saving process. Hot rolling conditions <steel embryo heating temperature: 1 1 〇〇°c or more> The temperature of the steel embryo heating is preferably low temperature heating from the viewpoint of energy. When the heating temperature is less than 1 00 °C, it may occur: carbide cannot be obtained. Adequate solid solution and an increase in the rolling load cause problems such as an increase in the possibility of hindrance during hot rolling. Further, the steel blast heating temperature is preferably 1 to 300 t or less, based on an increase in oxidized scale loss associated with an increase in oxidized weight. Further, it is based on the viewpoint that the steel sheet heating temperature can be prevented from being hindered during hot rolling, and a so-called sheet heater for heating a sheet bar can be used. <Finishing finish temperature: Ar3 transformation point or more> If the finishing rolling temperature does not reach the Ar3 transformation point, ferrite iron and Vostian iron are formed during rolling, and a strip-like structure is easily formed in the steel sheet. The band structure remains even after cold rolling or after annealing, which may cause anisotropy in material properties or cause deterioration in workability. Therefore, the finish rolling temperature is preferably above the Ar3 metamorphic point. <Winding temperature: 450 to 700 ° C > If the coiling temperature is less than 450 ° C, the control of the coiling temperature becomes difficult and temperature unevenness is likely to occur, and as a result, problems such as deterioration of cold rolling properties may occur. When the temperature exceeds 700 °C, the problem of decarburization may occur in the iron surface layer. -15- 201209179 Therefore, the coiling temperature should be set in the range of 450 to 700 °C. Further, in the hot rolling step of the present invention, in order to reduce the rolling load during hot rolling, a part or all of the finish rolling may be subjected to lubrication rolling. It is effective to perform lubrication rolling based on the viewpoint of uniformity of the shape of the steel sheet and uniformity of the material. Further, the coefficient of friction at the time of lubrication rolling is preferably in the range of 0.2 5 to 0.10. It is also preferable to use a continuous fine rolling process in which the successively fed sheets are joined to each other to continuously perform finish rolling. From the standpoint of the stability of hot rolling operations, a continuous finish rolling process is preferred. Next, the oxidized scale on the surface of the hot-rolled steel sheet is removed by pickling, and then cold-rolled to form a cold-rolled steel sheet having a predetermined thickness. The pickling conditions and the cold rolling conditions are not particularly limited, and may be carried out according to a usual method. The cold rolling reduction ratio is preferably 40% or more. <Average heating rate of 500 ° C to ACl metamorphic point: 10 ° C /s or more> The steel of the present invention sets the average heating rate of the recrystallization temperature range (500 ° C to AC | metamorphic point) to 1 〇 When the temperature is above °C /s, the recrystallization at the time of heating and heating can be suppressed, and the Vostian iron which is formed at a point higher than the ACl transformation point can be made finer, thereby contributing to the refinement of the microstructure after annealing and cooling. The average particle diameter of the low temperature metamorphic phase is 3 // m or less. When the average heating rate is less than 10 ° C / s, recrystallization of α occurs when the temperature is raised by heating, and the strain introduced into the ferrite iron is released, and sufficient refinement cannot be achieved. Therefore, the average heating rate of the 500 ° C ~ AC | metamorphic point is set to 1 ° C / s or more. The average heating rate is preferably in the range of 20 l / s or more. -16- 201209179 <Preservation at temperatures above 750 °C for more than 1 sec.> When the heating temperature is less than 750 ton: or the holding time is less than 10 seconds, the formation of Worthite during annealing is insufficient. 'A sufficient amount of low temperature metamorphic phase cannot be ensured after annealing cooling. Although the upper limit of the temperature and the holding time is not particularly specified, the effect is saturated and the cost rises when the temperature is maintained at 900 ° C or higher and the holding time is 600 seconds or longer. Therefore, it is preferable to keep the temperature below 9 〇〇 ° C and maintain The time is less than 600 seconds. <Cooling from 750 °C to an average cooling rate of 10 °C / s to a temperature range of 150 ° to 350 ° C> From 750 ° C cooling rate, not up to l ° ° When C / S, a pearly structure is formed, and TSXEL and hole expandability are deteriorated. Therefore, the cooling rate from 750 °C is set to 10 °C / s or more. Cooling to temperature conditions is one of the most important conditions of the technology. When the cooling stops, a part of the Worth Iron will be transformed into a granulated iron, and the rest will become the untransformed Worth Iron. After reheating, electroplating, and alloying treatment, it is cooled to room temperature, and the granulated iron becomes tempered granulated iron, and the untransformed Worth iron becomes residual volcanic iron or granulated iron. The lower the temperature from the annealing to the lower temperature, the more the amount of granulated iron produced in the cooling, the less the amount of iron in the untransformed Vostian, so by controlling the cooling to reach the temperature, the final granulated iron and the residual Worthite iron and back are determined. The area ratio of scattered iron in Hematian. When the cooling reaches a temperature higher than 35 (TC higher temperature, when the cooling stops, the numb iron is not metamorphosed and the unconformed Worth iron is too much, and finally the -17-201209179 amount of 麻田散铁 or residual Worthite iron is generated, and In addition, if the cooling reaching temperature is lower than 150 °C, most of the Worth iron in the cooling becomes metamorphic into the granulated iron, which reduces the amount of untransformed Worth, and cannot obtain more than 3% of the residue. Vostian Iron. Therefore, the cooling reaching temperature is set in the range of 150 to 350 ° C. Regarding the cooling method, as long as the target cooling rate and cooling stop temperature can be achieved, gas jet cooling and spraying can be used. Cooling method such as cooling, water cooling, metal quenching, etc. <heating to 350~600 °C and keeping 10~600 seconds> After cooling in the temperature range of 150~3 50 °C, at 305~600 °C The temperature range is maintained for 1 sec or more, thereby tempering the granulated iron which is generated during the cooling, and becomes tempered granulated iron to improve the hole expandability, and further the untransformed state of the granulated iron during the cooling. Vostian Iron is stabilized, Finally, more than 3% of the residual Worthite iron is obtained to improve the ductility. The details of the mechanism for stabilizing the undeformed Worthite iron by reheating are not well understood, but should be from the solid solution of supersaturated C. Ma Tian San Yu makes C diffuse into the untransformed Worthite iron and concentrates the C in the untransformed Worthite iron, thereby stabilizing the Worthite iron. At this time, if the precipitation of Xueming carbon iron in the Matian Sanyu is The diffusion of C is earlier, and the concentration of C in the untransformed Worthite iron becomes insufficient. Therefore, it is important to delay the precipitation of ferritic carbon iron, so it is necessary to add S. 3 % or more of S i . When the temperature reaches 3 50 °C, the tempering of the granulated iron of the Matian iron and the stabilization of the Worthite iron are insufficient, and the hole expandability and ductility are deteriorated. If the reheating temperature exceeds 600 °C, the Worthfield Iron Co. Metamorphosis into pearly tissue-18- 201209179 'The last time I can't get more than 3% of the residual Worthite iron. So set the heating temperature to 35〇~6〇〇.〇. When the holding time is less than 10 seconds, Wo The stability of the Situ iron is insufficient, and if it exceeds 600 seconds, the non-metamorphism when the cooling stops The Worthite iron will be transformed into a toughened iron. Finally, more than 3% of the residual Worthite iron cannot be obtained. Therefore, the reheating temperature is set to a range of 3 50 to 600 ° C, and the holding time of the temperature range is set to In the steel sheet after annealing, in order to correct the shape, adjust the surface roughness, etc., it is possible to carry out temper rolling, and there is no treatment such as resin, grease coating, or various coatings. [Example 1] A steel having the composition shown in Table 1 and having the remainder being Fe and unavoidable impurities was melted in a converter and a cast piece was obtained by a continuous casting method. Hot rolled into a plate thickness of 3 _0mm. The hot rolling conditions were carried out at a finishing temperature of 900 ° C, a cooling rate after rolling, 10 ° C / s, and a coiling temperature of 600 °C. Next, the hot rolled steel sheet was pickled, and then cold rolled into a sheet thickness of 1 · 2 m m to obtain a cold rolled steel sheet. Next, for these cold-rolled steel sheets, annealing treatment was carried out on the continuous annealing line under the conditions shown in Table 2. The cross-sectional microstructure, tensile properties and hole expandability of the obtained steel sheets were investigated. The results are shown in Table 3. -19- 201209179 [1® (%¥M) 5 槲粼ms 粼粼粼豳ms 镟m 豳 anchor a 镒a 瑙 agoo to g ο ο m ο ο od - go S ο a ο 缺 C4 ο > g ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο ο S ο od CM sd ιη ο ο to go CO soaoo CO so σ> S ο CO CO oosoo in ο ο (Ο ο ο S ο ο CNJ ooodg ο ο soosooooooo C0 ο ο α soo in S o CO ο ο g ο ο g ο ο godod ο ο CO oo' CM od CO oo CJ CM oo r- ο ο C CO csi oc\i CO τ· τ·» CM cJ σ> CO CNJ σ> Production m evi 00 CO 31 31 W CNJ Ο ΙΟ ir> d - ir>σ> ό CsJ o CNJ CO CO or — od od r- 〇d 00 5 ιη eg d CO od CNJ o' 兮o ο τ» ο GO O d σ> O o II Γ-* 5 Ο ο i1ml1 m 翳< CQ Ο Ω at u. 〇X - «wake ssffiHg member sslle * Qu miles nf ii -20 -? 201209179 [Table 2]
No. 鋼種 Acj 變態點 500°C~Ac,之 平均加熱速度 最髙到達 m. 保持 時間 平均 冷卻 速度 冷卻後之 到達溫度 再加熱 溫度 再加熱後 保持時間 °C °C/s °c 秒 °C/s °c °C 秒 1 A 721 15 830 60 50 200 400 80 發明例 2 A 15 810 60 50 100 420 80 比較例 3 B 740 20 850 90 80 180 430 60 發明例 4 B 20 720 60 80 250 430 60 比較例 5 C 734 5 &20 90 30 160 450 45 發明例 6 C 5 820 5 30 120 450 45 比較例 7 C 5 820 90 30 30 450 45 比較例 8 D 735 30 780 150 70 150 450 60 發明例 9 D 30 780 120 3 210 450 60 比較例 10 D 30 780 120 100 380 450 50 比較例 11 E 708 7 850 75 80 180 400 30 發明例 12 E 7 850 60 80 200 250 60 比較例 13 E 7 830 75 80 200 650 60 比較例 14 E 7 850 75 80 40 400 30 比較例 15 F 723 15 800 240 90 200 400 90 發明例 16 F 15 820 240 90 220 400 0 比較例 17 F 15 800 240 90 240 500 900 比較例 18 G 725 15 850 60 100 200 500 30 發明例 19 H 720 15 840 120 90 180 400 30 發明例 20 I 718 15 830 75 150 220 500 45 發明例 21 J 743 15 800 45 80 180 400 20 發明例 22 K 730 15 800 200 100 210 550 10 比較例 23 L 686 15 820 120 150 220 400 60 比較例 24 M 745 15 840 90 150 160 400 20 比較例 註:加底線表示超出本發明的範圍 -21 - 201209179 【s |發明例| |比較例I |發明例| |比較例I |發明例| |比較例| 1比較例1 1發明例1 I比較例I 丨比較例| |發明例I 丨比較例I I比較例I I比較例I 發明例 比較例 比較例 |發明例I 1發明例1 I發明例I |發明例I 比較例 1比較例 比較例 AE/TS 0.063 0.063 0.064 0.055 0.058 0.052 0.057 0.067 0.054 0.067 0.055 0.057 0.049 0.054 0.064 0.065 0.063 0.071 0.066 0.064 0.065 0.054 0.052 0.055 10%之吸收 能量(AE) MJ/m 5 CO m CO in Γ ΙΟ (Ο in iO <〇 S m <〇 o 5 S (D in s N 〇> 00 CO s CO CO 03 CO 擴孔率 a*： in CO CM σ> CO CO s in in iO eg σ> CO CO CO csj CO σ> CO CO 00 g esj <〇 CM CO CO l〇 σ> IO CO eg σ> o in IO (O TSXEL MPa·% 23400 18300 22620 17535 22770 18700 18620 25350 18400 24250 22100 18898 18929 18040 22680 18165 18308 26390 22500 24564 22152 17108 18550 16860 izi S S (D CM 〇 CM S σ> CO CM S in OJ a S CM S S in CM 00 eg S (0 MPa 900 in « 870 835 s ⑦ \n CO Oi 980 in Oi S O) 970 850 859 823 820 840 865 796 1015 900 1068 923 CO 1325 CM CD IO 其他 相* 0. a Q. m a m 0. 低溫變態^ 平均粒徑 1 ε ΪΙ r- csi CO e>i O) (O c>i m cj P·^ CO <〇 cj <〇 ci r«* ri U9 CO o cJ CNJ 00 esi pj GO evj CO σ> cj 殘留 沃斯田鐵S mm% <0 04) 兮 OI CO Ή —1 eg ^-1 in cal ^-1 卜 Ή o m 卜 IO —1 O) o 回火 麻田散鐵量 面積％ Oi CM to C5 CO CM c〇l σ> CO 〇>l o m CSJ o in s Si 卜 CO fO s s 〇〇I fSI ml 麻田散鐵量 面積％ 〇 〇 〇 o 〇 ο Ο o o 到 ir> o o o o o o Ο o o o 〇 肥粒鐵S 面積％ IT) CO o <〇 o « in in CM <〇 in In m CD CO IO Oi (D o in <〇 jn CVJ o s CO m to to JO CO c〇 鋼種 < < m m 〇 Ο Ο Q 〇 〇 UJ UJ UJ LU U. u. u. a X - ~3 父 -J i - esi O u> (Ο r- CO Oi o 二 CM CO in (O 卜 GO 〇> s OI CM eo CM s -22 - 201209179 鋼板之截面微組織，是使用3%硝太（Nit al )溶液（ 3%硝酸+乙醇）讓組織顯現出，使用掃描式電子顯微鏡觀 察深度方向板厚1 /4的位置，使用所拍攝的組織相片進行 影像解析處理，將肥粒鐵相的分率定量化（影像解析處理 可使用市售的影像處理軟體）。麻田散鐵面積率、回火麻 田散鐵面積率，是按照組織的細度而拍攝1 000~3 000倍之 適當倍率的SEM相片，使用影像處理軟體予以定量化。 低溫變態相的平均粒徑，是將所觀察的視野之低溫變態相 的面積除以低溫變態相的個數而求出平均面積，將其開根 號而獲得平均粒徑。 殘留沃斯田鐵的體積率，是將鋼板硏磨至板厚方向的 1 /4面，依據該板厚1 /4面的繞射X射線強度來求出。入 射X射線是使用ΜοΚ α線，對於殘留沃斯田鐵相的n} 、{200}、{220 }、{311}面和肥粒鐵相之{110}、{200}、 { 2 1 1 }面之峰値的積分強度所有的組合求出強度比，以其 等的平均値作爲殘留沃斯田鐵的體積率。 此外，拉伸特性’是使用JIS 5號試驗片（以拉伸方 向成爲與鋼板的輥軋方向垂直的方式進行取樣），依據 JIS Z224 1進行拉伸試驗，測定TS (拉伸強度）、EL (伸 長率），求出以強度和伸長率的乘積（TSxEL)來表元；$ 強度-伸長率的平衡値。 再者’作爲評價凸緣成形性的指標，是測定擴孔率λ 。擴孔率λ是依據日本鋼鐵聯盟規格j f s Τ 1 0 0 1來進行丨廣子匕 試驗，根據衝壓時之孔的初期孔徑（1 0 m m φ )和進行擴子匕 -23- 201209179 加工後孔緣的龜裂貫穿板厚時之孔徑的比値來求出。 衝擊吸收特性，是使用從與鋼板的輥軋方向垂直的方 向取樣之平行部寬度5mm、長度7mm的試驗片，以應變 速度2000/s進行拉伸試驗，將所採取的應力-真應變曲線 在應變量〇〜10%的範圍進行積分，算出吸收能量並進行 評價（參照：鋼與鐵83 (1997)ρ·748)。 本發明例的鋼板，顯示TSxEL爲22000MPa · %以上 ，λ爲7〇%以上之優異的強度、延性及凸緣成形性。 相對於此，偏離本發明的範圍之比較例的鋼板，TSx EL未達22000MPa . %且（或）λ未達70%，無法獲得像 本發明例的鋼板那麼優異的強度、延性及凸緣成形性。再 者，藉由使低溫變態相的平均粒徑成爲3 ν m以下，能獲 得吸收能量與TS的比値（AE/TS)爲0.063以上之優異 的耐衝擊特性。 本發明’作爲加工性及耐衝擊性優異的高強度冷軋鋼 板，有助於汽車之輕量化及降低燃料消耗率。 -24-No. Steel type Acj metamorphosis point 500 ° C ~ Ac, the average heating rate reaches the maximum m. Hold time average cooling rate After cooling, the temperature reaches the reheating temperature and then reheats the holding time °C °C / s °c seconds °C /s °c ° C sec 1 A 721 15 830 60 50 200 400 80 Inventive Example 2 A 15 810 60 50 100 420 80 Comparative Example 3 B 740 20 850 90 80 180 430 60 Inventive Example 4 B 20 720 60 80 250 430 60 Comparative Example 5 C 734 5 & 20 90 30 160 450 45 Inventive Example 6 C 5 820 5 30 120 450 45 Comparative Example 7 C 5 820 90 30 30 450 45 Comparative Example 8 D 735 30 780 150 70 150 450 60 Invention Example 9 D 30 780 120 3 210 450 60 Comparative Example 10 D 30 780 120 100 380 450 50 Comparative Example 11 E 708 7 850 75 80 180 400 30 Inventive Example 12 E 7 850 60 80 200 250 60 Comparative Example 13 E 7 830 75 80 200 650 60 Comparative Example 14 E 7 850 75 80 40 400 30 Comparative Example 15 F 723 15 800 240 90 200 400 90 Invention Example 16 F 15 820 240 90 220 400 0 Comparative Example 17 F 15 800 240 90 240 500 900 Comparative Example 18 G 725 15 850 60 100 200 500 30 Inventive Example 19 H 720 15 840 120 90 180 400 30 Inventive Example 20 I 718 15 830 75 150 220 500 45 Inventive Example 21 J 743 15 800 45 80 180 400 20 Inventive Example 22 K 730 15 800 200 100 210 550 10 Comparative Example 23 L 686 15 820 120 150 220 400 60 Comparative Example 24 M 745 15 840 90 150 160 400 20 Comparative Example Note: Bottom line indicates that the scope of the present invention is out of the range of 21 - 201209179 [s | Invention Example | | Comparative Example I | Invention Example | | Comparative Example I | Invention Example | | 1 Comparative Example 1 1 Inventive Example 1 I Comparative Example I 丨 Comparative Example | | Inventive Example I 丨 Comparative Example II Comparative Example II Comparative Example I Inventive Example Comparative Example Comparative Example | Inventive Example I Inventive Example 1 Inventive Example I | Invention Example I Comparative Example 1 Comparative Example Comparative Example AE/TS 0.063 0.063 0.064 0.055 0.058 0.052 0.057 0.067 0.054 0.067 0.055 0.057 0.049 0.054 0.064 0.065 0.063 0.071 0.066 0.064 0.065 0.054 0.052 0.055 10% absorbed energy (AE) MJ/m 5 CO m CO in Γ ΙΟ (Ο in iO <〇S m <〇o 5 S (D in s N 〇> 00 CO s CO CO 03 CO hole expansion ratio a*: in CO CM σ> CO CO s in in iO eg σ> CO CO CO csj CO σ> CO CO 00 g esj <〇CM CO CO l〇σ> IO CO eg σ> o in IO (O TSXEL MPa·% 23400 18300 22620 17535 22770 18700 18620 25350 18400 24250 22100 18898 18929 18040 22680 18165 18308 26390 22500 24564 22152 17108 18550 16860 izi SS (D CM 〇CM S σ> CO CM S in OJ a S CM SS in CM 00 eg S (0 MPa 900 in « 870 835 s 7 \n CO Oi 980 in Oi SO) 970 850 859 823 820 840 865 796 1015 900 1068 923 CO 1325 CM CD IO Other phases* 0. a Q . mam 0. low temperature metamorphosis ^ average particle size 1 ε ΪΙ r- csi CO e>i O) (O c>im cj P·^ CO <〇cj <〇ci r«* ri U9 CO o cJ CNJ 00 Esi pj GO evj CO σ> cj residual Vostian iron S mm% <0 04) 兮OI CO Ή —1 eg ^-1 in cal ^-1 Ή Ή OB IO —1 O) o tempering Ma Tiansan Iron area% Oi CM to C5 CO CM c〇l σ> CO 〇>lom CSJ o in s Si Bu CO fO ss 〇〇I fSI ml Ma Tian scattered iron area % 〇〇〇o 〇ο Ο oo to ir> ; oooooo Ο ooo 〇 粒 粒 S area% IT) CO o <〇o « in in CM <〇in In m CD CO IO Oi (D o in <〇 Jn CVJ os CO m to to JO CO c〇 steel type << mm 〇Ο Ο Q 〇〇UJ UJ UJ LU U. uu a X - ~3 parent-J i - esi O u> (Ο r- CO Oi o Two CM CO in (O Bu GO 〇> s OI CM eo CM s -22 - 201209179 The cross-section microstructure of the steel plate is made using 3% Nit al (Nit al + ethanol) to make the tissue appear Scanning electron microscope is used to observe the position of the depth direction plate thickness of 1/4, and the image of the photographed tissue is used for image analysis to quantify the fraction of the ferrite grain phase. (Image analysis processing can use commercially available image processing software. ). The area ratio of the granulated iron in the field and the area of the tempered iron in the tempering field are taken as SEM photographs of appropriate magnifications of 1,000 to 3,000 times according to the fineness of the organization, and quantified using image processing software. The average particle diameter of the low-temperature metamorphic phase is obtained by dividing the area of the low-temperature metamorphic phase of the observed visual field by the number of low-temperature metamorphic phases to obtain an average area, and opening the root number to obtain an average particle diameter. The volume fraction of the residual Worthite iron is obtained by honing the steel sheet to the 1/4 surface in the thickness direction, and determining the intensity of the diffraction X-ray according to the thickness of the plate. Incident X-rays are ΜοΚ α lines, {110}, {200}, { 2 1 1 for the n}, {200}, {220}, {311} faces of the residual Worthfield iron phase and the ferrite phase. } The integral intensity of the peak of the surface is determined by the combination of all the strength ratios, and the average enthalpy of the surface is taken as the volume fraction of the remaining Worthite iron. In addition, the tensile property 'is a JIS No. 5 test piece (sampling in such a manner that the stretching direction is perpendicular to the rolling direction of the steel sheet), and a tensile test is performed in accordance with JIS Z224 1 to measure TS (tensile strength) and EL. (Elongation), the product of the strength and elongation (TSxEL) is obtained; the balance of the strength-elongation 値. Further, as an index for evaluating the formability of the flange, the hole expansion ratio λ is measured. The hole expansion ratio λ is based on the Japanese Iron and Steel Federation specification jfs Τ 1 0 0 1 to carry out the 丨 匕 匕 匕 test, according to the initial hole diameter (10 mm φ ) of the hole during punching and the hole after machining -23-201209179 The ratio of the pore diameter of the crack is determined by the ratio of the pore diameter. The impact absorption characteristic was obtained by using a test piece having a parallel portion width of 5 mm and a length of 7 mm sampled in a direction perpendicular to the rolling direction of the steel sheet, and performing a tensile test at a strain rate of 2000/s, and the stress-true strain curve taken was The range of the strain 〇~10% was integrated, and the absorbed energy was calculated and evaluated (reference: steel and iron 83 (1997) ρ·748). The steel sheet of the example of the present invention exhibits excellent strength, ductility, and flange formability with a TSxEL of 22,000 MPa·% or more and a λ of 7〇% or more. On the other hand, in the steel sheet of the comparative example which deviated from the range of the present invention, the TSx EL was less than 22,000 MPa% and/or λ was less than 70%, and the strength, ductility and flange forming of the steel sheet of the example of the present invention could not be obtained. Sex. In addition, by setting the average particle diameter of the low-temperature metamorphic phase to 3 ν m or less, it is possible to obtain an excellent impact resistance characteristic of the ratio 値 (AE/TS) of the absorbed energy to TS of 0.063 or more. The present invention is a high-strength cold-rolled steel sheet excellent in workability and impact resistance, contributing to weight reduction of the automobile and reduction in fuel consumption rate. -twenty four-