WO2018076965A1 - 一种抗拉强度在1500MPa以上且成形性优良的冷轧高强钢及其制造方法 - Google Patents

一种抗拉强度在1500MPa以上且成形性优良的冷轧高强钢及其制造方法 Download PDF

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WO2018076965A1
WO2018076965A1 PCT/CN2017/102384 CN2017102384W WO2018076965A1 WO 2018076965 A1 WO2018076965 A1 WO 2018076965A1 CN 2017102384 W CN2017102384 W CN 2017102384W WO 2018076965 A1 WO2018076965 A1 WO 2018076965A1
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cold
strength
rolled high
strength steel
steel
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PCT/CN2017/102384
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English (en)
French (fr)
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周澍
钟勇
王利
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宝山钢铁股份有限公司
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Priority to EP17863565.2A priority Critical patent/EP3533894A4/en
Priority to US16/342,842 priority patent/US11279986B2/en
Priority to JP2019517956A priority patent/JP6770640B2/ja
Priority to KR1020197014024A priority patent/KR20190071755A/ko
Publication of WO2018076965A1 publication Critical patent/WO2018076965A1/zh

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    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/008—Martensite

Definitions

  • the present invention relates to a steel material and a method of manufacturing the same, and more particularly to a cold rolled steel material and a method of manufacturing the same.
  • the publication number is CN102227511A, and the publication date is October 26, 2011.
  • the Chinese patent document entitled "High-strength cold-rolled steel sheet, high-strength hot-dip galvanized steel sheet having excellent formability and a method for producing the same" discloses a kind having 1180 MPa.
  • the patent is set by a continuous annealing process, so that the steel plate has martensite, ferrite and retained austenite phase, the martensite content is above 30%, and the ratio of martensite content to ferrite content is 0.45-1.5.
  • the tensile strength of the steel sheet disclosed in this patent document is 1180 MPa or more, and the strength requirement of 1500 MPa or more cannot be satisfied.
  • the publication number is CN104160055A, and the publication date is November 19, 2014.
  • the Chinese patent document entitled "High-strength cold-rolled steel sheet and its manufacturing method" discloses a high-strength cold-rolled steel sheet, but its tensile strength reaches 1180 MPa or more. The elongation after fracture is 16% or more, and the strength requirement of 1500 MPa or more cannot be satisfied.
  • One of the objects of the present invention is to provide a cold-rolled high-strength steel having a tensile strength of 1500 MPa or more and excellent formability, which is excellent in formability under the condition of obtaining ultra-high strength by reasonable composition design and microstructure control. Its elongation after break is ⁇ 12%.
  • the present invention provides a cold-rolled high-strength steel having a tensile strength of 1500 MPa or more and excellent formability, and the chemical element mass percentage thereof is:
  • the microstructure of the cold rolled high strength steel has 5-20% retained austenite and 70-90% martensite, and the ratio of retained austenite to martensite carbon concentration is greater than 3.5 and less than 15.
  • C is a solid solution strengthening element necessary for ensuring strength in steel, and is an austenite stabilizing element.
  • the mass percentage of C is less than 0.25%, the content of retained austenite is insufficient, and the strength of steel is low; and when the mass percentage of C is more than 0.40%, the weldability of the steel is remarkably deteriorated. Therefore, the mass percentage of carbon in the cold-rolled high-strength steel according to the present invention is controlled to be 0.25-0.40%.
  • the mass percentage of carbon is further controlled between 0.28 and 0.32%.
  • Si In the technical solution described in the present invention, it promotes the enrichment of carbon in the austenite phase, suppresses the formation of carbides, and further stabilizes the retained austenite phase.
  • the mass percentage of Si is less than 1.50%, the effect is unclear; however, when the mass percentage of Si added exceeds 2.50%, the embrittlement of the steel sheet becomes conspicuous, which in turn causes cracks in the edge of the steel sheet during cold rolling, which reduces production. effectiveness. Therefore, the mass percentage of silicon in the cold-rolled high-strength steel according to the present invention is limited to 1.5 to 2.5%.
  • the mass percentage of silicon is further controlled between 1.6 and 2.0%.
  • Mn increases the stability of austenite, and at the same time reduces the critical cooling temperature and the martensite transformation temperature Ms at the time of steel quenching, and improves the hardenability of the steel sheet. Further, Mn is a solid solution strengthening element and is advantageous for increasing the strength of the steel sheet. When the mass percentage of manganese is higher than 3.0%, manganese causes cracks in the billet and affects the weldability of the steel. Therefore, the cold rolled high strength steel according to the present invention controls the mass percentage of Mn to be 2.0 to 3.0%.
  • the mass percentage of manganese is further controlled between 2.6 and 2.9%.
  • Al is added for deoxidation of steel.
  • the mass percentage of Al is less than 0.03%, the purpose of deoxidation cannot be achieved; and when the mass percentage of Al is more than 0.06%, the deoxidation effect is saturated. Therefore, the mass percentage of aluminum in the cold-rolled high-strength steel according to the present invention is controlled to be 0.03-0.06%.
  • Chromium and molybdenum increase the hardenability of the steel and thus the strength of the steel.
  • the present invention controls the mass percentage of Cr and Mo to at least one of 0.1-1.0% Cr and 0.1-0.5% Mo.
  • Phosphorus In the technical solution described in the present invention, P deteriorates the weldability, increases the cold brittleness of the steel, and reduces the plasticity of the steel. Therefore, it is necessary to control the mass percentage of P to be 0.02% or less.
  • S acts as a harmful element, which deteriorates the weldability and lowers the plasticity of the steel. Therefore, it is necessary to control the mass percentage of S to be 0.01% or less.
  • N and Al combine with AlN particles, thereby affecting the ductility and thermoplasticity of the steel sheet. Therefore, it is necessary to control the mass percentage of N to be 0.01% or less.
  • the microstructure of the cold rolled high strength steel has 5-20% of retained austenite and 70-90% of martensite. This is because during the deformation process of cold-rolled high-strength steel, a certain amount of retained austenite transforms into martensite, which produces a phase change-induced plastic effect (that is, a TRIP effect), thereby ensuring the invention.
  • Cold rolled high strength steel has good formability while having a tensile strength of 1500 MPa or more. When the retained austenite content is less than 5%, the TRIP effect is not significant, and the high formability of the steel sheet cannot be ensured.
  • the cold-rolled high-strength steel according to the present invention has a content of martensite of 70% to 90%. Further, in the cold-rolled high-strength steel according to the present invention, the ratio of the carbon concentration of retained austenite to martensite is more than 3.5 and less than 15. This is because the carbon concentration in the retained austenite represents the stability of the retained austenite.
  • the ratio of the carbon concentration of retained austenite to martensite is greater than 3.5 and less than 15.
  • the microstructure of the cold rolled high strength steel is retained austenite + martensite + ferrite or retained austenite + martensite + bainite.
  • the chemical element mass percentage thereof also satisfies: Mn+Cr+Mo ⁇ 3.8%. This is because when the sum of the mass percentages of Cr, Mn, and Mo is higher than 3.8%, the steel sheet is likely to be markedly banded, and the cold rolling resistance becomes large.
  • the chemical element mass percentage also satisfies: C+Si/30+Mn/20+2P+4S ⁇ 0.56%.
  • the inventor of the present invention found through a large number of research experiments that when the mass percentage of C, Si, Mn, P and C satisfies C+Si/30+Mn/20+2P+4S ⁇ 0.56%, the strength of cold-rolled high-strength steel is high, and the welding performance is higher.
  • each chemical element that satisfies the chemical mass ratio of the formula has a better effect on the structural strengthening and solid solution strengthening of steel.
  • the elongation after break is ⁇ 12%.
  • the cold-rolled high-strength steel according to the present invention further contains at least one of Nb: 0.01-0.1%, V: 0.01-0.2%, and Ti: 0.01-0.05%.
  • Nb is added because the Nb element can strengthen the steel by precipitation strengthening, and at the same time prevent the growth of austenite grains and refine the crystal grains, thereby improving the strength and elongation of the steel.
  • the mass percentage of Nb is controlled to 0.01-0.10% because: when the mass percentage of Nb is less than 0.01%, it has no effect; but when the mass percentage of Nb exceeds 0.10%, the precipitation strengthening excessively acts, resulting in cold rolling. The formability of high-strength steel is reduced, and at the same time, the manufacturing cost is increased.
  • V is added because V can form carbides and increase the strength of the steel.
  • the mass percentage of V is controlled at 0.01-0.20% because the precipitation strengthening effect is not significant when the mass percentage of V is less than 0.01%; however, when the mass percentage of V is more than 0.2%, the precipitation strengthening effect excessively acts, resulting in The formability of the steel sheet is lowered.
  • Ti is added because Ti can form precipitates with C, S, and N to effectively increase the strength and toughness of the steel sheet.
  • the mass percentage of Ti is controlled to 0.01-0.05% because when the mass percentage of titanium is less than 0.01%, the effect is not obvious; when the mass percentage of Ti is higher than 0.05%, the content is increased, for steel. The improvement is not significant.
  • Another object of the present invention is to provide a method for producing the above-described cold-rolled high-strength steel, which comprises the steps of:
  • Continuous annealing heating the strip to a soaking temperature of 800-900 ° C, holding for more than 60 s, then cooling to 150-300 ° C at a rate of 30-80 ° C / s, and then heating to 350-440 ° C , keep warm for 30-300s, and finally cool to room temperature.
  • the key is to select the parameters of the continuous annealing process in the step (5), because the continuous annealing process can diffuse the super-saturated carbon in the martensite into the retained austenite.
  • Austenite is rich in carbon and stably remains at room temperature, and martensite transformation occurs during deformation by retained austenite, resulting in a TRIP effect, thereby ensuring high formability of the cold-rolled high-strength steel according to the present invention.
  • a large amount of martensite in the microstructure ensures the high strength of the cold-rolled high-strength steel according to the present invention.
  • the present invention selects various parameters of the annealing process in order to ensure sufficient retained austenite and martensite in the structure.
  • the design principle is as follows:
  • the soaking temperature is controlled at 800-900 °C, and the temperature is kept for more than 60 s. It is because the austenite transformation of cold-rolled high-strength steel is insufficient when the soaking temperature is lower than 800 °C or the holding time is less than 60 s. The body structure is uneven, the ferrite content is too much, and in the subsequent process, a sufficient amount of retained austenite and martensite cannot be formed.
  • the strength of the steel is low and the elongation is insufficient after the break; when the soaking temperature is higher than At 900 °C, the austenite transformation occurs in the microstructure of the steel sheet after soaking, and the austenite stability is reduced, so that the retained austenite content in the steel sheet matrix after annealing is reduced, the strength of the steel is high and the elongation is insufficient after the fracture. .
  • the soaking temperature is controlled at 820-870 °C.
  • Cooling to 150-300 ° C at a rate of 30-80 ° C / s because: in the technical solution of the present invention, the martensite critical cooling rate is 30 ° C / s, therefore, only to occur in the cooling process Martensitic transformation, cooling rate is not less than 30 ° C / s. However, when the cooling rate exceeds 80 ° C / s, the cooling cost required for production will increase significantly. When the cooling termination temperature at this rate is lower than 150 ° C, the austenite undergoes martensite transformation, and the microstructure obtained when the steel is cooled at room temperature is no retained austenite, resulting in the elongation of the steel.
  • the cooling termination temperature is set at 180-270 °C.
  • the reheating temperature is controlled at 350-440 ° C, and the reheating time is 30-300 s. This is because when the reheating temperature is lower than 350 ° C or the reheating time is less than 30 s, the residual austenite stabilization process in the cold-rolled high-strength steel is insufficient, and the retained austenite in the microstructure obtained when the steel is cooled at room temperature. The body content is insufficient; when the reheating temperature is higher than 440 ° C or the reheating time is higher than 300 s, the steel undergoes significant temper softening, resulting in a significant decrease in the material strength of the cold rolled high strength steel.
  • the slab in the heating stage, the slab is heated to 1200-1300 ° C and kept for 0.5-4 h; in the rolling stage, the final rolling temperature is controlled ⁇
  • the coiling temperature is controlled at 850 ° C and is 400-600 ° C.
  • the holding time of the heating stage is controlled to be 0.5-4h. When the holding time exceeds 4h, the grain structure in the slab is coarse and the surface of the slab is decarburized seriously; when the holding time is less than 0.5h, the internal temperature of the slab is not yet Evenly.
  • the heating temperature in the heating stage is further preferably from 1210 to 270 °C.
  • the control of the finish rolling temperature and the coiling time is because the final rolling temperature is too low, which may cause the slab deformation resistance to be too high, so that it is difficult to produce a finished product of suitable quality.
  • the coiling temperature is higher than 600 °C, the scale of the iron oxide on the surface of the steel sheet is too thick to be pickled; when the coiling temperature is lower than 400 °C, the coiling strength is too high, which is difficult to be cold-rolled, which affects production efficiency.
  • the coiling temperature is further preferably from 450 to 550 °C.
  • the pickling speed is controlled to be 80 to 120 m/min. This is because, in the manufacturing method of the present invention, when the pickling speed is higher than 120 m/min, the scale of the steel sheet surface cannot be completely removed, thereby forming a surface defect of the steel sheet; when the pickling speed is lower than 80 m/min, The mill speed is low, which in turn affects the shape control of the steel sheet and reduces the production efficiency.
  • the cold rolling reduction is controlled to be 40 to 60%.
  • Increasing the cold rolling reduction in the technical solution described in the present invention is advantageous for increasing the austenite formation rate in the subsequent continuous annealing process and contributing to the improvement of the uniformity of the steel sheet. Improve the ductility of the steel sheet.
  • the cold rolling reduction is higher than 60%, the deformation resistance of the material is extremely high due to work hardening, which makes the cold-rolled high-strength steel sheet deteriorate.
  • the cold-rolled high-strength steel sheet having a tensile strength of 1500 MPa or more and excellent formability according to the present invention is excellent in formability under high strength by a reasonable composition design and microstructure control, and the elongation after fracture is ⁇ 12. %.
  • the manufacturing method of the present invention has the above advantages, and the cold-rolled high-strength steel sheet of the present invention has excellent strength and high formability by controlling the process parameters, especially the selection of the continuous annealing process parameters, in the automobile structural parts. It has a good application prospect and is particularly suitable for manufacturing vehicle structural parts and safety parts which require high forming performance and strength.
  • Table 1 lists the mass distribution ratios of the chemical elements in the cold-rolled high-strength steel of each example and the conventional steel sheet of the comparative example.
  • Table 2 lists the specific process parameters of the manufacturing methods of the respective examples and comparative examples.
  • the rapid cooling rate in Table 2 refers to the cooling rate when cooled to 150-300 ° C at 30-80 ° C / s as defined in the claims; the chemical composition numbers in Table 2 use the corresponding chemistry listed in Table 1.
  • the mass ratio of each chemical element of the component number for example, the mass ratio of each chemical element of the chemical component A in Table 1 is used in the embodiment 1.
  • the tensile test method was as follows: JIS No. 5 tensile test specimen was used, and the stretching direction was parallel to the rolling direction.
  • Residual austenite ratio V ⁇ test method A 15 ⁇ 15 mm sample was cut from a steel plate, and after grinding and polishing, an XRD quantitative test was performed.
  • Martensite comparison V ⁇ test method A 15 ⁇ 15 mm sample was cut from a steel plate, and after grinding and polishing, EBSD quantitative analysis was performed.
  • Residual austenite carbon concentration Test method Assuming that the Mn and Al concentrations of the constituent phases in the steel sheet structure did not change, the lattice constant a ⁇ was read using the diffraction peak data of the retained austenite in XRD, and the empirical formula was used: Calculation. among them, with Respectively represent the C concentration, Mn concentration and Al concentration of retained austenite.
  • Martensite carbon concentration Test method according to the obtained V ⁇ , V ⁇ and And the C concentration of the design component x C , by the formula: Calculated.
  • Table 3 lists the test results obtained after testing the cold-rolled high-strength steel of each example and the comparative steel plate of the comparative example.
  • Comparative Example 3 uses the chemical composition E ratio, the mass percentage of carbon is less than 0.25%, resulting in the strength of Comparative Example 3 not reaching 1500 MPa; Comparative Example 4 is based on the chemical composition F ratio, The mass percentage of silicon was less than 1.5%, resulting in a lower elongation after fracture of Comparative Example 4.
  • the soaking temperature of Comparative Example 1 is higher than 900 ° C, and the steel sheet undergoes complete austenite transformation after the heat treatment, and the austenite content is less than 5%. Therefore, the tensile strength of Comparative Example 1 Although it is higher than 1500 MPa, its elongation after break is insufficient and the formability is poor.
  • the rapid cooling termination temperature of Comparative Example 2 was higher than 300 ° C, resulting in insufficient martensite transformation, and the content of retained austenite was low during the subsequent reheating, eventually resulting in insufficient elongation after the fracture of Comparative Example 2.

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Abstract

一种抗拉强度在1500MPa以上且成形性优良的冷轧高强钢，其化学元素质量百分配比为：C 0.25-0.40％，Si 1.50-2.50％，Mn 2.0-3.0％，Al 0.03-0.06％，P≤0.02％，S≤0.01％，N≤0.01％以及0.1-1.0％的Cr和0.1-0.5％的Mo的至少其中之一，余量为Fe和其他不可避免的杂质；所述冷轧高强钢的微观组织具有5-20％的残余奥氏体和70-90％的马氏体，并且残余奥氏体与马氏体的碳浓度之比大于3.5且小于15。该冷轧高强度钢板，通过合理的成分设计和微观组织控制，使其强度高，成形性优良。

Description

一种抗拉强度在1500MPa以上且成形性优良的冷轧高强钢及其制造方法 技术领域
本发明涉及一种钢材及其制造方法，尤其涉及一种冷轧钢材及其制造方法。
背景技术
近年来，为实现车身减重，达到节能减排、提高碰撞安全性以及降低制造成本等方面的目的，越来越多抗拉强度在1000MPa及以上的钢板被应用在复杂的车身结构件和安全件的设计上。然而，随着钢板强度级别的提高，其成形性急剧下降，零件冲压开裂问题日益突出。因此，改善超高强度钢板的成形性，避免冲压成形出现开裂便成为了重要课题。
公开号为CN102227511A，公开日为2011年10月26日，名称为“成形性优良的高强度冷轧钢板、高强度热镀锌钢板及它们的制造方法”的中国专利文献公开了一种具有1180MPa以上的TS、且扩孔性和弯曲性等成形性优良的高强度冷轧钢板、高强度热镀锌钢板及它们的制造方法。该专利通过连续退火工艺设定，使得钢板具有马氏体、铁素体和残余奥氏体相，马氏体含量在30％以上，马氏体含量与铁素体含量之比在0.45～1.5之间，然而该专利文献公开的钢板抗拉强度在1180MPa以上，无法满足1500MPa以上的强度需求。
公开号为CN104160055A，公开日为2014年11月19日，名称为“高强度冷轧钢板及其制造方法”的中国专利文献公开了一种高强度冷轧钢板，然而其抗拉强度达到1180MPa以上，断后延伸率16％以上，无法满足1500MPa以上的强度需求。
鉴于此，期望获得一种冷轧高强度钢板，在具备1500MPa以上的超强强度的同时仍具有良好的成形性，因而特别适用于制造复杂的车身结构件和安全件，满足汽车轻量化需求。与此同时，还期望获得该钢板的制造方法，所述制造方法工艺简单，适用性强，可用在现有大多数生产线上。
发明内容
本发明的目的之一在于提供一种抗拉强度在1500MPa以上且成形性优良的冷轧高强钢，通过合理的成分设计和微观组织控制，使其在获得超高强度的情况下成形性优良，其断后延伸率≥12％。
基于上述发明目的，本发明提供了一种抗拉强度在1500MPa以上且成形性优良的冷轧高强钢，其化学元素质量百分比为：
C 0.25-0.40％，Si 1.50-2.50％，Mn 2.0-3.0％，Al 0.03-0.06％，P≤0.02％，S≤0.01％，N≤0.01％以及0.1-1.0％的Cr和0.1-0.5％的Mo的至少其中之一，余量为Fe和其他不可避免的杂质；
所述冷轧高强钢的微观组织具有5-20％的残余奥氏体和70-90％的马氏体，并且残余奥氏体与马氏体的碳浓度之比大于3.5且小于15。
本发明所述的冷轧高强钢的各化学元素的设计原理为：
碳：C是钢中为了确保强度而必需的固溶强化元素，是奥氏体稳定化元素。当C的质量百分比低于0.25％时，残余奥氏体的含量不足，且钢的强度偏低；而当C的质量百分比高于0.40％时，钢材的焊接性能显著恶化。因此，本发明所述的冷轧高强钢中碳的质量百分比控制0.25-0.40％。
在一些优选的实施方式中，碳的质量百分比进一步控制在0.28-0.32％。
硅：Si在本发明所述的技术方案中，起到促进碳在奥氏体相中的富集、抑制碳化物的生成，进而使残余奥氏体相稳定化的作用。当Si的质量百分比低于1.50％，其作用效果不明；然而，当添加Si的质量百分比超过2.50％，钢板脆化变得显著，进而导致在冷轧时易在钢板边部产生裂纹，降低生产效率。因此，本发明所述的冷轧高强钢中对硅的质量百分比限定在1.5-2.5％。
在一些优选的实施方式中，硅的质量百分比进一步控制在1.6-2.0％。
锰：在本发明所述的技术方案中，Mn增加奥氏体的稳定性，同时降低了钢淬火时的临界冷却温度以及马氏体转变温度Ms，提高钢板的淬透性。此外，Mn是固溶强化元素，对提高钢板的强度有利。当锰的质量百分比高于3.0％时，锰会导致钢坯裂纹产生，影响钢的焊接性能。因此，本发明所述的冷轧高强钢控制Mn的质量百分比为2.0-3.0％。
在一些优选的实施方式中，锰的质量百分比进一步控制在2.6-2.9％。
铝：在本发明所述的技术方案中，Al是为了钢的脱氧而添加的。当Al的质量百分比低于0.03％时，无法达到脱氧的目的；而Al的质量百分比高于0.06％时，脱氧效果饱和。因此，本发明所述的冷轧高强钢中对铝的质量百分比控制在0.03-0.06％。
铬和钼：Cr和Mo可提高钢的淬透性，进而提高钢的强度。为保证钢的淬透性，本发明对Cr和Mo的质量百分比控制在0.1-1.0％的Cr和0.1-0.5％的Mo的至少其中之一。
磷：在本发明所述的技术方案中，P使得焊接性能变差，增加钢的冷脆性，降低了钢的塑性，因此，需要控制P的质量百分比在0.02％以下。
硫：在本发明所述的技术方案中，S作为有害元素，使得焊接性能变差，降低钢的塑性，因此，需要控制S的质量百分比在0.01％以下。
氮：在本发明所述的技术方案中，N与Al结合AlN颗粒，进而影响钢板的延展性和热塑性。因此，需要将N的质量百分比控制在0.01％以下。
此外，在本发明所述的冷轧高强钢中所述冷轧高强钢的微观组织具有5-20％的残余奥氏体和70-90％的马氏体。这是因为：在冷轧高强钢变形过程中，一定量的残余奥氏体发生相变转变为马氏体，产生相变诱导塑性效应(也就是TRIP效应)，从而保证了本发明所述的冷轧高强钢在具有1500MPa以上抗拉强度的同时具有良好的成形性。当残余奥氏体含量＜5％，TRIP效应不显著，无法保证钢板的高成形性；而残余奥氏体量＞20％，则钢板中相应的马氏体含量偏少，无法实现钢板的高强度。马氏体作为钢板的硬相，为钢板提供了高强度。为达到1500MPa及以上的高强度，本发明所述的冷轧高强钢对马氏体的含量限定在70％-90％。此外，本发明所述的冷轧高强钢中，残余奥氏体与马氏体的碳浓度之比大于3.5且小于15。这是因为：残余奥氏体中的碳浓度代表了残余奥氏体的稳定性。当残余奥氏体与马氏体的碳浓度之比小于3.5时，残余奥氏体的稳定性不足，残余奥氏体易在变形过程初期发生马氏体相变，导致TRIP效应不足，无法显著提升钢板的成形性；当残余奥氏体与马氏体的碳浓度之比超过15时，残余奥氏体过于稳定，无法在变形过程中发生马氏体相变，也会导致TRIP效应不足，钢板的成形性同样无法提高。因此，在本发明所述的技术方案中，残余奥氏体与马氏体的碳浓度之比大于3.5且小于15。
进一步地，本发明所述的冷轧高强钢中，所述冷轧高强钢的微观组织为残余奥氏体+马氏体+铁素体或残余奥氏体+马氏体+贝氏体。
进一步地，本发明所述的冷轧高强钢中，其化学元素质量百分比还满足：Mn+Cr+Mo≤3.8％。这是由于当Cr、Mn和Mo的质量百分比之和高于3.8％时，钢板容易出现带状明显，冷轧抗力变大。
进一步地，本发明所述的冷轧高强钢中，其化学元素质量百分比还满足：C+Si/30+Mn/20+2P+4S≤0.56％。本案发明人通过大量研究实验发现，当C、Si、Mn、P和的质量百分比满足C+Si/30+Mn/20+2P+4S≤0.56％获得的冷轧高强钢强度高，焊接性能较好，这是由于满足该式的化学质量配比的各化学元素对钢的组织强化和固溶强化效果较好。
进一步地，本发明所述的冷轧高强钢中，其断后延伸率≥12％。
进一步地，本发明所述的冷轧高强钢还含有Nb：0.01-0.1％，V：0.01-0.2％和Ti：0.01-0.05％的至少其中之一。
其中，添加Nb是因为Nb元素可以通过析出强化而对钢进行强化，同时阻止奥氏体晶粒的长大，细化晶粒，从而提高了钢的强度和断后延伸率。将Nb的质量百分比控制在0.01-0.10％是因为：当Nb的质量百分比小于0.01％时，无法起到作用效果；但当Nb的质量百分比超过0.10％时，析出强化过度发挥作用，导致冷轧高强钢的成形性下降，同时，增加制造成本。
添加V是因为V可形成碳化物，提高钢的强度。对V的质量百分比控制在0.01-0.20％是因为：当V的质量百分比小于0.01％时，析出强化效果不显著；然而，当V的质量百分比大于0.2％时，析出强化效果过度发挥作用，导致钢板成形性下降。
添加Ti是因为Ti可以与C、S、N形成析出物而有效地提高钢板的强度和韧性。将Ti的质量百分比控制在0.01-0.05％是因为：当钛的质量百分比低于0.01％时，其作用效果不明显；当Ti的质量百分比高于0.05％时，再增加其含量，对于钢的改善效果并不显著。
此外，本发明的另一目的在于提供一种上述冷轧高强钢的制造方法，其包括步骤：
(1)冶炼和铸造；
(2)热轧；
(3)酸洗；
(4)冷轧；
(5)连续退火：将带钢加热至均热温度800-900℃之间，保温60s以上，然后以30-80℃/s的速度冷却至150-300℃，然后再加热至350-440℃，保温30-300s，最后冷却至室温。
在本发明所述的制造方法中，关键在于步骤(5)中对连续退火工艺各参数的选取，这是由于连续退火工艺能够使马氏体中的过饱和碳扩散至残余奥氏体中，奥氏体富碳并稳定保留在室温下，并通过残余奥氏体在变形过程中发生马氏体相变，产生TRIP效应，从而保证本发明所述的冷轧高强钢的高成形性。同时，微观组织中大量的马氏体保证了本发明所述的冷轧高强钢的高强度。
因此，本发明对于退火工艺各参数的选取，是为了保证组织中具有足够的残余奥氏体和马氏体。其设计原理具体如下所述：
对于连续退火中均热温度控制在800-900℃，保温60s以上，是由于：当均热温度低于800℃或保温时间不足60s，冷轧高强钢的奥氏体相变不充分，奥氏体组织不均匀，铁素体含量偏多，在随后的工艺过程中，无法形成足够量的残余奥氏体和马氏体，钢的强度偏低同时断后延伸率不足；当均热温度高于900℃时，均热处理后的钢板基体组织发生完全奥氏体相变，奥氏体稳定性降低，从而使得退火后钢板基体中残余奥氏体含量减少，钢的强度偏高同时断后延伸率不足。
在一些优选的实施方式中，均热温度控制在820-870℃。
以30-80℃/s的速度冷却至150-300℃，是因为：在本发明所述的技术方案中，马氏体临界冷却速度为30℃/s，因此，为保证冷却过程中仅发生马氏体相变，冷却速度不小于30℃/s。但当冷速超过80℃/s时，生产所需要的冷却成本将大幅上升。而当以该速度冷却的冷却终止温度低于150℃时，则奥氏体全部发生马氏体转变，钢在室温冷却时获得的微观组织为无残余奥氏体生成，导致钢的断后延伸率不足；而当以该速度冷却的冷却终止温度高于300℃时，则马氏体相变发生不充分，在随后的再加热过程中，残余奥氏体的含量和稳定性不足，进而影响钢的强度和成形性。
在一些优选的实施方式中，冷却终止温度设定在180-270℃。
在本发明所述的制造方法，步骤(5)中对再加热温度控制在350-440℃，再加热时间在30-300s。这是因为：当再加热温度低于350℃或再加热时间低于30s时，冷轧高强钢中残余奥氏体稳定化过程不充分，钢在室温冷却时获得的微观组织中的残余奥氏体含量不足；当再加热温度高于440℃或再加热时间高于300s时，钢发生显著的回火软化，导致冷轧高强钢的材料强度明显下降。
进一步地，本发明所述的制造方法中，在所述步骤(2)中，在加热阶段，将板坯加热到1200-1300℃并保温0.5-4h；在轧制阶段，控制终轧温度≥850℃，控制卷取温度为400-600℃。
这是因为：加热阶段时，加热温度高于1300℃时，会造成板坯过烧，板坯内晶粒组织粗大导致其热加工性能降低，并且超高温会引起板坯表面严重脱碳；当加热温度低于1200℃时，板坯经高压水除鳞和初轧后，精轧温度过低，会造成坯料变形抗力过大。此外，加热阶段的保温时间控制在为0.5-4h，当保温时间超过4h，会造成板坯内晶粒组织粗大同时板坯表面脱碳严重；当保温时间低于0.5h，板坯内部温度尚未均匀。
在一些优选的实施方式中，加热阶段的加热温度进一步优选为1210-1270℃。
另外，对终轧温度和卷取时间的控制是因为：终轧温度过低会造成板坯变形抗力过高，从而难以生产出质量合适的成品。当卷取温度高于600℃，钢板表面的氧化铁皮生成过厚，难于酸洗；当卷取温度低于400℃，则卷取强度偏高，难于冷轧，影响生产效率。
在一些优选的实施方式中，卷取温度进一步优选为450-550℃。
进一步地，本发明所述的制造方法中，在所述步骤(3)中，控制酸洗速度为80-120m/min。这是因为：在本发明所述的制造方法中，酸洗速度高于120m/min，则钢板表面的氧化铁皮无法完全去除，从而形成钢板表面缺陷；当酸洗速度低于80m/min，则轧机速度较低，进而影响对钢板板形控制，并且降低了生产效率。
进一步地，本发明所述的制造方法中，在所述步骤(4)中，控制冷轧压下量为40-60％。在本发明所述的技术方案中增加冷轧压下量有利于后续的连续退火过程中提高奥氏体形成速率，并且有助于提高钢板的组织均匀性，从 而提高钢板的延展性。但冷轧压下量高于60％时，因加工硬化导致材料的变形抗力非常高，使得冷轧高强钢板形变差。
本发明所述的抗拉强度在1500MPa以上且成形性优良的冷轧高强度钢板，通过合理的成分设计和微观组织控制，使其在强度高的情况下成形性优良，其断后延伸率≥12％。
本发明所述的制造方法具有上述优点以外，通过对工艺参数的控制，尤其是连续退火工艺参数的选取上，使得本发明所述的冷轧高强度钢板强度高成形性优良，在汽车结构件中具有很好的应用前景，特别适合于制造对成形性能和强度都要求较高的车辆结构件和安全件。
具体实施方式
下面将结合具体的实施例对本发明所述的抗拉强度在1500MPa以上且成形性优良的冷轧高强钢及其制造方法做进一步的解释和说明，然而该解释和说明并不对本发明的技术方案构成不当限定。
实施例1-8和对比例1-4
实施例1-8的冷轧高强钢和对比例1-4的常规钢板采用下述步骤制得：
(1)采用表1所列出的化学元素的质量百分配比冶炼和铸造板坯；
(2)热轧：在加热阶段，将板坯加热到1200-1300℃并保温0.5-4h，在轧制阶段，控制终轧温度≥850℃，控制卷取温度为400-600℃；
(3)酸洗：控制酸洗速80-120m/min；
(4)冷轧：控制冷轧压下量为40-60％；
(5)连续退火：将钢加热至均热温度800-900℃之间，保温60s以上，然后以30-80℃/s的速度冷却至150-300℃，然后再加热至350-440℃，保温30-300s，最后冷却至室温。
表1列出了各实施例的冷轧高强钢和对比例的常规钢板中各化学元素的质量百分配比。
表1.(wt％，余量为Fe和除了S、P、N以外的其他不可避免杂质元素)
Figure PCTCN2017102384-appb-000001
Figure PCTCN2017102384-appb-000002
表2列出了各实施例和对比例的制造方法的具体工艺参数。
表2
Figure PCTCN2017102384-appb-000003
Figure PCTCN2017102384-appb-000004
注：表2中快速冷却速度是指以权利要求中限定的30-80℃/s冷却至150-300℃时的冷却速度；表2中的化学组分序号采用表1所列的相对应化学组分序号的各化学元素质量配比，例如：实施例1采用的是表1中化学组分A的各化学元素质量配比。
对上述各实施例的冷轧高强钢和对比例的常规钢板，进行各项性能测试，将试验测得到的相关结果列于表3中。其中，屈服强度和抗拉强度测试为常规测试，其他测试方法具体如下所述：
拉伸试验方法为：采用JIS5号拉伸试样，拉伸方向平行于轧制方向。
残余奥氏体相比例Vγ测试方法：从钢板上切取15×15mm尺寸的试样，经过研磨和抛光后，进行XRD定量测试。
马氏体相比例Vα测试方法：从钢板上切取15×15mm尺寸的试样，经过研磨和抛光后，进行EBSD定量分析。
残余奥氏体碳浓度
Figure PCTCN2017102384-appb-000005
测试方法：假设钢板组织中各组成相的Mn和Al浓度未发生变化，利用XRD中残余奥氏体的衍射峰数据读取晶格常数aγ，并利用经验公式：
Figure PCTCN2017102384-appb-000006
计算。其中，
Figure PCTCN2017102384-appb-000007
和
Figure PCTCN2017102384-appb-000008
分别代表残余奥氏体的C浓度，Mn浓度和Al浓度。
马氏体碳浓度
Figure PCTCN2017102384-appb-000009
测试方法：根据得到的Vγ，Vα和
Figure PCTCN2017102384-appb-000010
以及设计成分的C浓度xC，通过公式：
Figure PCTCN2017102384-appb-000011
计算得到。
表3列出了各实施例的冷轧高强钢和对比例的常规钢板测试后所获得的测试结果。
表3
Figure PCTCN2017102384-appb-000012
Figure PCTCN2017102384-appb-000013
从表3可以看出，本案各实施例的抗拉强度均在1500MPa以上，断后延伸率在12％以上，说明本案各实施例的冷轧高强钢具有高强度和良好的成形性。
结合表1和表3可知，对比例3采用化学组成分E配比，其碳的质量百分比低于0.25％，导致对比例3强度无法达到1500MPa；对比例4采用化学组成分F配比，其硅的质量百分比低于1.5％，导致对比例4的断后延伸率较低。
结合表2和表3可知，对比例1的均热温度高于900℃，均热处理后钢板发生完全奥氏体相变，奥氏体含量低于5％，因此，对比例1的抗拉强度虽然高于1500MPa，然而其断后延伸率不足，成形性差。对比例2的快速冷却终止温度高于300℃，导致其马氏体相变不充分，在随后的再加热过程中，残余奥氏体的含量偏低，最终导致对比例2断后延伸率不足。
需要注意的是，以上列举的仅为本发明的具体实施例，显然本发明不限于以上实施例，随之有着许多的类似变化。本领域的技术人员如果从本发明公开的内容直接导出或联想到的所有变形，均应属于本发明的保护范围。

Claims (10)

  1. 一种抗拉强度在1500MPa以上且成形性优良的冷轧高强钢，其特征在于：
    其化学元素质量百分配比为：C 0.25-0.40％，Si 1.50-2.50％，Mn 2.0-3.0％，Al 0.03-0.06％，P≤0.02％，S≤0.01％，N≤0.01％以及0.1-1.0％的Cr和0.1-0.5％的Mo的至少其中之一，余量为Fe和其他不可避免的杂质；
    所述冷轧高强钢的微观组织具有5-20％的残余奥氏体和70-90％的马氏体，并且残余奥氏体与马氏体的碳浓度之比大于3.5且小于15。
  2. 如权利要求1所述的冷轧高强钢，其特征在于，所述冷轧高强钢的微观组织为残余奥氏体+马氏体+铁素体或残余奥氏体+马氏体+贝氏体。
  3. 如权利要求1所述的冷轧高强钢，其特征在于，其化学元素质量百分比还满足：Mn+Cr+Mo≤3.8％。
  4. 如权利要求1-3中任意一项所述的冷轧高强钢，其特征在于，其化学元素质量百分比还满足：C+Si/30+Mn/20+2P+4S≤0.56％。
  5. 如权利要求1所述的冷轧高强钢，其特征在于，其还含有Nb：0.01-0.1％，V：0.01-0.2％和Ti：0.01-0.05％的至少其中之一。
  6. 如权利要求1所述的冷轧高强钢，其特征在于，其断后延伸率≥12％。
  7. 如权利要求1-6中任意一项所述的冷轧高强钢的制造方法，其特征在于，包括步骤：
    (1)冶炼和铸造；
    (2)热轧；
    (3)酸洗；
    (4)冷轧；
    (5)连续退火：将带钢加热至均热温度800-900℃之间，保温60s以上，然后以30-80℃/s的速度冷却至150-300℃，然后再加热至350-440℃，保温30-300s，最后冷却至室温。
  8. 如权利要求7所述的制造方法，其特征在于，在所述步骤(2)中，在加热阶段，将板坯加热到1200-1300℃并保温0.5-4h；在轧制阶段，控制终轧温度≥850℃，控制卷取温度为400-600℃。
  9. 如权利要求7所述的制造方法，其特征在于，在所述步骤(3)中，控制酸洗速度80-120m/min。
  10. 如权利要求7所述的制造方法，其特征在于，在所述步骤(4)中，控制冷轧压下量为40-60％。
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CN108018484B (zh) 2020-01-31
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