WO2026002018A1 - 横纵向强度差异低的低成本冷轧低合金高强钢及其制造方法 - Google Patents

横纵向强度差异低的低成本冷轧低合金高强钢及其制造方法

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Publication number
WO2026002018A1
WO2026002018A1 PCT/CN2025/103371 CN2025103371W WO2026002018A1 WO 2026002018 A1 WO2026002018 A1 WO 2026002018A1 CN 2025103371 W CN2025103371 W CN 2025103371W WO 2026002018 A1 WO2026002018 A1 WO 2026002018A1
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Prior art keywords
low
strength
strength steel
alloy high
transverse
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PCT/CN2025/103371
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English (en)
French (fr)
Inventor
张瀚龙
钟勇
谭文
周文强
黄全伟
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Baoshan Iron and Steel Co Ltd
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Baoshan Iron and Steel Co Ltd
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Publication of WO2026002018A1 publication Critical patent/WO2026002018A1/zh
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Classifications

    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D1/00—General methods or devices for heat treatment, e.g. annealing, hardening, quenching or tempering
    • C21D1/26—Methods of annealing
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C33/00—Making ferrous alloys
    • C22C33/04—Making ferrous alloys by melting
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/02—Ferrous alloys, e.g. steel alloys containing silicon
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/04—Ferrous alloys, e.g. steel alloys containing manganese
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/06—Ferrous alloys, e.g. steel alloys containing aluminium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/14—Ferrous alloys, e.g. steel alloys containing titanium or zirconium

Definitions

  • This invention relates to steel plates and their manufacturing methods, and more particularly to a low-alloy ultra-high-strength steel and its manufacturing method.
  • cold-rolled high-strength steel with a yield strength of 700 MPa is either enhanced by adding high alloys or by combining carbide precipitation with large cold rolling deformation and incomplete annealing.
  • One of the objectives of this invention is to provide a low-cost cold-rolled low-alloy high-strength steel with low difference in transverse and longitudinal strength.
  • This low-alloy high-strength steel has both ultra-low alloy cost and low carbon equivalent, while having a small difference in transverse and longitudinal yield strength.
  • the present invention provides a low-cost cold-rolled low-alloy high-strength steel with low difference in transverse and longitudinal strength, which contains Fe and unavoidable impurities, and also contains the following chemical elements in the following mass percentages:
  • Its microstructure contains ferrite, granular bainite, and precipitated carbonitrides.
  • the mass percentage content of each chemical element is as follows:
  • the balance is Fe and unavoidable impurities.
  • This invention achieves a microstructure of ferrite + granular bainite + precipitated carbonitrides through the coordinated design of composition and process, thereby obtaining a low-cost, low-carbon equivalent steel with high strength, high formability, and low difference between transverse and longitudinal strength.
  • the design principles of each chemical element in the low-alloy ultra-high-strength steel described in this invention are as follows:
  • C In the low-alloy ultra-high-strength steel described in this invention, carbon (C) simultaneously affects phase transformation and carbonitride precipitation.
  • the mass percentage content of C When the mass percentage content of C is too low, effective phase transformation strengthening cannot be achieved, and the amount of carbonitrides formed is also very limited.
  • the mass percentage content of C is too high, although phase transformation is more likely to occur, it also causes a significant increase in carbon equivalent and coarse carbonitride size, affecting weldability and formability. Therefore, in the low-alloy ultra-high-strength steel described in this invention, the mass percentage content of C can be controlled between 0.045% and 0.085%.
  • Mn is one of the core elements affecting strength and controlling phase transformation.
  • the mass percentage content of Mn is too low, the transformation of granular bainite will be insufficient, resulting in insufficient strength; when the mass percentage content of Mn is too high, although it is beneficial to improve the strength of the steel, it will lead to an increase in carbon equivalent, affecting weldability, and will also increase manufacturing costs. Therefore, in the low-alloy ultra-high-strength steel described in this invention, the mass percentage content of Mn can be controlled between 1.16% and 1.50%.
  • Al can act as a deoxidizing element and a carbonitride forming element, improving the quality and strength of the steel. Simultaneously, Al is also a ferrite forming and stabilizing element.
  • the mass percentage content of Al can be controlled between 0.02% and 0.05%.
  • Ti In the low-alloy ultra-high-strength steel described in this invention, Ti is the main carbonitride-forming element, used to refine grains and improve strength. When the mass percentage content of Ti is too low, the precipitation of carbonitrides is insufficient, failing to achieve the required strength; when the mass percentage content of Ti is too high, coarse TiN particles and coarse titanium carbonitride are easily formed in the steel, which is detrimental to the formability of the steel. Therefore, in the low-alloy ultra-high-strength steel described in this invention, the mass percentage content of Ti can be controlled between 0.10 and 0.16.
  • Si is an essential element for steelmaking deoxidation and solid solution strengthening, as well as a ferrite formation and stabilizing element.
  • the mass percentage content of Si can be controlled to 0 ⁇ Si ⁇ 0.2%.
  • the low-alloy high-strength steel described in this invention also contains at least one of the following elements:
  • the low-alloy high-strength steel of the present invention contains Fe and unavoidable impurities, and also contains the following chemical elements in the following mass percentages: C: 0.045–0.085%; Mn: 1.16–1.50%; Al: 0.02–0.05%; Ti: 0.10–0.16%; O ⁇ Si ⁇ 0.2%; Nb: ⁇ 0.02%; Ce: ⁇ 0.006%; La: ⁇ 0.006%; V: ⁇ 0.1%; Cr: ⁇ 0.15%.
  • the mass percentage of each chemical element in the low-alloy high-strength steel of the present invention is as follows: C: 0.045–0.085%; Mn: 1.16–1.50%; Al: 0.02–0.05%; Ti: 0.10–0.16%; 0 ⁇ Si ⁇ 0.2%; Nb: ⁇ 0.02%; Ce: ⁇ 0.006%; La: ⁇ 0.006%; V: ⁇ 0.1%; Cr: ⁇ 0.15%; the balance being Fe and unavoidable impurities.
  • the low-alloy high-strength steel of the present invention contains at least three of Nb, Ce, La, V, and Cr.
  • Nb, Ce, La, V, and Cr can be selectively added to the steel as optional elements.
  • La and Ce, as rare earth elements have the effect of improving steel quality, modifying inclusions, and refining grains;
  • Nb and V, as strong carbonitride forming elements can form a large number of carbonitride precipitates, refine grains, and improve strength;
  • Cr can regulate the phase transformation of steel and also improve hardenability, which is conducive to the formation of granular austenite.
  • the mass percentage content of Nb can be controlled to Nb ⁇ 0.02wt%
  • the mass percentage content of Ce can be controlled to Ce ⁇ 0.006wt%
  • the mass percentage content of La can be controlled to La ⁇ 0.006wt%
  • the mass percentage content of V can be controlled to V ⁇ 0.1wt%
  • the mass percentage content of Cr can be controlled to Cr ⁇ 0.15wt%.
  • the mass percentage content of each chemical element satisfies at least one of the following conditions:
  • Ti/(Al+Nb) ⁇ 3.0 can also be controlled to regulate the precipitation of carbonitrides.
  • Al, Ti, and Nb are all carbonitride-forming elements, but in terms of their ability to form carbonitrides, Nb > Ti > Al. If the Al content in the steel is too high and the Ti and Nb content is too low, fewer carbonitrides will form during manufacturing, resulting in insufficient steel strength. Conversely, if the Nb content is too high, although carbonitride formation is easier during manufacturing, it will lead to coarse carbonitride sizes and consume more carbon, resulting in insufficient or no granular bainite formation.
  • Ti/(Al+Nb) ⁇ 3.0 is controlled. In some embodiments, Ti/(Al+Nb) is controlled within the range of 3.0 to 5.5.
  • the alloy element ratio can also be controlled to ensure that 1.1 ⁇ Mn + Cr - Ti - Nb - V ⁇ 1.6, thereby regulating the optimal ratio of carbonitride precipitation and granular bainite formation.
  • Mn and Cr are beneficial for granular bainite formation, while Ti, Nb, and V are strong carbide-forming elements. Therefore, the higher the value of the above formula, the more favorable it is for granular bainite formation; the lower the value, the more favorable it is for carbonitride precipitation.
  • Mn + Cr - Ti - Nb - V is controlled within the range of 1.1 to 1.4.
  • Cev carbon equivalent
  • unavoidable impurities mainly include phosphorus (P), sulfur (S), and nitrogen (N), and it is desirable to minimize their content when technical conditions permit.
  • the volume fraction of ferrite in its microstructure is ⁇ 90%
  • the volume fraction of granular bainite is ⁇ 5%
  • the carbonitride content is ⁇ 5%.
  • the sum of the volume fractions of ferrite and granular bainite in the microstructure of the low-alloy high-strength steel described in this invention is ⁇ 95%, with the remainder being carbonitrides.
  • ⁇ 60% of the ferrite grains have an aspect ratio ⁇ 4. In some embodiments, 60-90% of the ferrite grains have an aspect ratio ⁇ 4.
  • the difference between its transverse and longitudinal yield strength is ⁇ 50MPa. In some embodiments, the difference between the transverse and longitudinal yield strength of the low-alloy high-strength steel described in this invention is ⁇ 40MPa. In some embodiments, the difference between the transverse and longitudinal yield strength of the low-alloy high-strength steel described in this invention is ⁇ 30MPa.
  • the transverse and longitudinal yield strengths are both >700 MPa
  • its transverse and longitudinal elongation at break are both ⁇ 10%
  • its 90° bending radius to plate thickness ratio r/t ⁇ 0.3
  • the transverse yield strength of the low-alloy high-strength steel described in this invention is ⁇ 705 MPa
  • its longitudinal yield strength is ⁇ 720 MPa
  • the transverse yield strength of the low-alloy high-strength steel described in this invention is 705–790 MPa
  • its longitudinal yield strength is 720–830 MPa.
  • both its transverse and longitudinal tensile strengths are ⁇ 740 MPa.
  • the transverse tensile strength of the low-alloy high-strength steel described in this invention is ⁇ 745 MPa, and its longitudinal tensile strength is ⁇ 770 MPa.
  • the transverse tensile strength of the low-alloy high-strength steel described in this invention is 745–830 MPa, and its longitudinal tensile strength is 770–870 MPa.
  • Another objective of this invention is to provide a method for manufacturing low-alloy high-strength steel that is green and efficient.
  • the resulting low-alloy ultra-high-strength steel has the advantages of being green and low-carbon, as well as having low alloy content, and exhibits low difference in strength between the transverse and longitudinal directions.
  • the present invention provides a method for manufacturing low-alloy high-strength steel, comprising the following steps:
  • Annealing Heat the strip steel to 530-680°C and hold it at that temperature, then cool it to 440-520°C at a first cooling rate of 5-20°C/s and hold it at that temperature, then cool it to room temperature at a second cooling rate of 1-10°C.
  • the manufacturing method described in this invention can obtain a small amount of granular bainite in the hot-rolled precursor structure. Then, by optimizing the cold rolling deformation and annealing process, the optimal balance point between phase transformation strengthening and precipitation strengthening can be obtained, thereby obtaining a low-cost, low-carbon equivalent cold-rolled high-strength steel with high strength, high formability, low difference between transverse and longitudinal strength, and small difference between transverse and longitudinal strength.
  • the pickling and cold rolling step of the present invention by controlling the cold rolling reduction rate to 1-20%, it can be ensured that the strip does not produce elongated ferrite grains after annealing, so as to avoid excessive differences in transverse and longitudinal strength.
  • the lower heating and holding temperatures ensure that the granular bainite in the strip does not decompose, and the carbonitrides do not coarsen as a result.
  • the subsequent cooling, holding, and final cooling processes are to obtain submicron-sized carbonitride precipitation while further controlling the carbonitride growth size.
  • the strip steel is heated to 530-680°C and held at that temperature for 90-180 seconds, then cooled to 440-520°C at a first cooling rate of 5-20°C/s and held at that temperature for 5-30 seconds, and then cooled to room temperature at a second cooling rate of 1-10°C.
  • hot-dip galvanizing of the strip surface is completed in a process segment where the strip is cooled to 440-520°C at a first cooling rate of 5-20°C/s and held at that temperature.
  • the slab exit temperature is controlled at 1250–1300°C
  • the finishing rolling temperature is controlled at 890–940°C.
  • the strip is cooled to 440-520°C at a cooling rate of ⁇ 80°C/s (e.g., 80-160°C/s) after rolling and then coiled.
  • a higher slab exit temperature can ensure that the carbonitrides formed in the slab can be fully re-dissolved; a higher final rolling temperature and ultra-fast cooling rate can avoid the bulk precipitation and coarsening of carbonitrides; and low-temperature coiling at 440 to 520°C can form a small amount of granular bainite and trace amounts of fine carbonitrides in the strip.
  • the low-cost cold-rolled low-alloy high-strength steel with low difference in transverse and longitudinal strength described in this invention, and its manufacturing method, have the following advantages and beneficial effects:
  • the cold-rolled low-alloy high-strength steel described in this invention achieves low transverse and longitudinal strength differences while having both low alloy content and low carbon equivalent.
  • cold-rolled high-strength steel with a transverse and longitudinal yield strength of ⁇ 700MPa, a transverse and longitudinal fracture elongation of ⁇ 10%, a 90° bending radius to plate thickness ratio r/t ⁇ 0.3, and a transverse and longitudinal yield strength difference ⁇ 50MPa can be widely used in the preparation of various automotive structural parts.
  • Figure 1 shows the microstructure of Example 1.
  • Tables 1-1 and 1-2 list the mass percentages of each chemical element in the low-cost cold-rolled low-alloy high-strength steels of Examples 1-6 and the comparative steels of Comparative Examples 1-3.
  • Hot rolling The slab exit temperature is controlled at 1250 ⁇ 1300°C, and the finishing rolling temperature is 890 ⁇ 940°C; after rolling, the ultra-fast cooling low temperature coiling method is adopted, and the steel strip is cooled to 440 ⁇ 520°C at a cooling rate of ⁇ 80°C/s before coiling.
  • Annealing Heat the strip steel to 530-680°C and hold for 90-180s, then cool it to 440-520°C at a first cooling rate of 5-20°C/s and hold for 5-30s, then cool it to room temperature at a second cooling rate of 1-10°C.
  • the hot-dip galvanizing of the strip surface can be completed in the process of cooling to 440-520°C at a first cooling rate of 5-20°C/s and holding at that temperature in step (4), thereby obtaining a hot-dip galvanized product.
  • Tables 2-1 and 2-2 list the specific process parameters for the low-cost cold-rolled low-alloy high-strength steels of Examples 1-6 and the comparative steels of Comparative Examples 1-3.
  • Table 3 lists the microstructure observation results of the low-cost cold-rolled low-alloy high-strength steels of Examples 1-6 of the present invention and the comparative steels of Comparative Examples 1-3.
  • the microstructure of the low-cost cold-rolled low-alloy high-strength steel with low transverse and longitudinal strength difference in Examples 1-6 of the present invention contains ferrite + granular bainite + precipitated carbonitrides.
  • the volume ratio of ferrite is greater than or equal to 90%, the proportion of granular bainite is greater than or equal to 5%, and the proportion of carbonitrides is less than or equal to 5%.
  • Figure 1 shows the microstructure of Example 1.
  • Example 1 the microstructure of Example 1 is ferrite + granular bainite + carbonitrides, wherein the proportion of carbonitrides is 5.5%, the proportion of granular bainite is 4%, and the remainder is ferrite.
  • the low-cost cold-rolled low-alloy high-strength steels with low difference in transverse and longitudinal strength in Examples 1-6 of the present invention have a difference in transverse and longitudinal yield strength of less than 50 MPa, a yield strength of greater than 700 MPa in both transverse and longitudinal directions, a fracture elongation of greater than 10% in both transverse and longitudinal directions, and a ratio of 90° bending radius r (unit: mm) to plate thickness t (unit: mm) r/t of less than or equal to 0.3.

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Abstract

本发明公开了一种横纵向强度差异低的低成本冷轧低合金高强钢，其含有Fe及不可避免的杂质以及质量百分含量如下的下述各化学元素：C：0.045～0.085％；Mn：1.16～1.50％；Al：0.02～0.05％；Ti：0.10～0.16％；0＜Si≤0.2％；其微观组织含有铁素体+粒状贝氏体+析出的碳氮化物。本发明还公开了一种低合金高强钢的制造方法，其包括步骤：冶炼和铸造；热轧；酸洗冷轧：控制冷轧压下率为1～20％；退火：将带钢加热至530～680℃并保温，然后以5～20℃/s的冷速冷却至440～520℃并保温，然后以1～10℃的冷速冷却至室温。

Description

横纵向强度差异低的低成本冷轧低合金高强钢及其制造方法 技术领域
本发明涉及钢板及其制造方法，尤其涉及一种低合金超高强钢及其制造方法。
背景技术
伴随着汽车轻量化发展，越来越多的汽车结构件采用高强钢制备，其所要求的屈服强度逐渐达到或超过700MPa。
在现有技术中，屈服强度达到700MPa级别的冷轧高强钢要么通过添加高合金提升产品强度，要么通过碳化物析出结合冷轧大变形和不完全退火的方式，提升产品强度。
例如，公开号为CN105200332A，公开日为2015年12月30日，名称为“700MPa级薄规格高强钢带及其生产方法”的中国专利文献公开了一种屈服强度700MPa级薄规格高强钢带及其生产方法，其仅Mn+Cr两项合金的添加量之和就高达1.8％以上，导致钢材的碳当量较高，提升了制造成本。
又例如：公开号为CN109207843A，公开日为2019年1月15日，名称为“一种高强度冷轧带钢及其制造方法”的中国专利文献公开了一种高强度冷轧带钢,虽然其带钢强度高、塑性好且碳当量低、可焊性良好，但其横纵向强度差异大，尤其是屈服强度，横向比纵向高达50MPa以上。当横纵向强度差异过大时，会导致钢材在复杂冲压条件下的零件回弹异常和尺寸精度不达标。
又例如，公开号为CN115109994A，公开日为2022年9月27日，名称为“一种高强度冷轧热镀锌微合金带钢及其制造方法”的中国专利文献公开了一种高强度冷轧热镀锌微合金带钢，其采用以Nb代Ti的方法，得到了横纵向屈服强度差异＜50MPa的钢材。但是一方面以Nb代Ti会造成制造成本过高，另一方面上述带钢的屈服强度也低于700MPa。
发明内容
本发明的目的之一在于提供一种横纵向强度差异低的低成本冷轧低合金高强钢，该低合金高强钢在兼具超低合金成本和低碳当量的同时，其横、纵向屈服强度差异小。
为了实现上述目的，本发明提供了一种横纵向强度差异低的低成本冷轧低合金高强钢，其含有Fe及不可避免的杂质，此外其还含有质量百分含量如下的下述各化学元素：
C：0.045～0.085％；
Mn：1.16～1.50％；
Al：0.02～0.05％；
Ti：0.10～0.16％；
0＜Si≤0.2％；
其微观组织含有铁素体+粒状贝氏体+析出的碳氮化物。
进一步地，在本发明所述的低合金高强钢中，其各化学元素质量百分含量为：
C：0.045～0.085％；
Mn：1.16～1.50％；
Al：0.02～0.05％；
Ti：0.10～0.16％；
0＜Si≤0.2％；
余量为Fe和不可避免的杂质。
本发明通过成分和工艺的配合设计获得铁素体+粒状贝氏体+析出的碳氮化物的微观组织，从而获得具有高强度、高成形性能和低横纵向强度差的低成本低碳当量钢。具体来说，在本发明所述的低合金超高强钢中，各化学元素的设计原理具体如下所述：
C：在本发明所述的低合金超高强钢中，C元素同时影响相变和碳氮化物析出。当C元素的质量百分比含量过低时，无法形成有效的相变强化，碳氮化物的形成量也十分有限；当C元素的质量百分比含量过高时，虽然更容易发生相变，但也会造成碳当量的大幅上升和碳氮化物的尺寸粗大，影响可焊接性能额成形性能。因此，在本发明所述的低合金超高强钢中，可以将C元素的质量百分比含量控制在0.045～0.085％之间。
Mn：在本发明所述的低合金超高强钢中，Mn元素是影响强度和控制相变的核心元素之一。当Mn元素的质量百分比含量过低时，会使粒状贝氏体转变不足，导致强度不够；当Mn元素的质量百分比含量过高时，虽然有利于提升钢材强度，但一方面会导致碳当量升高，影响可焊接性能，另一方面也会提升制造成本。因此，在本发明所述的低合金超高强钢中，可以将Mn元素的质量百分比含量控制在1.16～1.50％之间。
Al：在本发明所述的低合金超高强钢中，Al元素可以作为脱氧元素和碳氮化物形成元素，改善钢材质量，提升钢材的强度。同时，Al元素也是铁素体形成元素和稳定元素。当Al元素的质量百分比含量过低时，钢水脱氧不充分，钢质纯净度会受到影响；当Al元素的质量百分比含量过高时，不利于形成粒状贝氏体。因此，在本发明所述的低合金超高强钢中，可以将Al元素的质量百分比含量控制在0.02～0.05％之间。
Ti：在本发明所述的低合金超高强钢中，Ti元素是主要的碳氮化物形成元素，用于细化晶粒、提升强度。当Ti元素的质量百分比含量过低时，形成的碳氮化物析出不够，无法达到需要的强度；当Ti元素的质量百分比含量过高时，容易在钢中形成粗大的TiN颗粒和粗化的碳氮化钛，不利于钢材的成型性能。因此，在本发明所述的低合金超高强钢中，可以将Ti元素的质量百分比含量控制在0.10～0.16之间。
Si：在本发明所述的低合金超高强钢中，Si元素是炼钢脱氧的必要元素和固溶强化元素，也是铁素体形成和稳定元素。但当Si元素的质量百分比含量过高时，会影响粒状贝氏体的形成，同时也会恶化可焊机性能，提升钢材制造的合金成本。因此，在本发明所述的低合金超高强钢中，可以将Si元素的质量百分比含量控制为0＜Si≤0.2％。
进一步地，在本发明所述的低合金高强钢中，其还含有下述元素的至少其中之一：
0＜Nb≤0.02wt％；
0＜Ce≤0.006wt％；
0＜La≤0.006wt％；
0＜V≤0.1wt％；
0＜Cr≤0.15wt％。
因此，在一些实施方案中，本发明所述低合金高强钢含有Fe及不可避免的杂质，此外其还含有质量百分含量如下的下述各化学元素：C：0.045～0.085％；Mn：1.16～1.50％；Al：0.02～0.05％；Ti：0.10～0.16％；0＜Si≤0.2％；Nb：≤0.02％；Ce：≤0.006％；La：≤0.006％；V：≤0.1％；Cr：≤0.15％。在一些实施方案中，本发明所述低合金高强钢的各化学元素质量百分含量为：C：0.045～0.085％；Mn：1.16～1.50％；Al：0.02～0.05％；Ti：0.10～0.16％；0＜Si≤0.2％；Nb：≤0.02％；Ce：≤0.006％；La：≤0.006％；V：≤0.1％；Cr：≤0.15％；余量为Fe和不可避免的杂质。在一些实施方案中，本发明所述低合金高强钢含有Nb、Ce、La、V和Cr中至少三种。
在本发明中，Nb、Ce、La、V和Cr可以作为可选元素选择性地加入钢中。其中，La元素和Ce元素作为稀土元素具有改善钢质、改性夹杂，细化晶粒的功效；Nb元素和V元素作为强碳氮化物形成元素，可以形成大量的碳氮化物析出相，细化晶粒，提升强度；Cr元素可以调节钢材相变，此外还可以提升淬透性，有利于粒状奥氏体形成。
但是上述合金元素的添加一方面会增加钢材的制造成本，另一方面也会增大碳当量，劣化钢材的可焊接性能。因此，在本发明所述的低合金高强钢中，可以将Nb元素的质量百分比含量控制为Nb≤0.02wt％，可以将Ce元素的质量百分比含量控制为Ce≤0.006wt％，可以将La元素的质量百分比含量控制为La≤0.006wt％，可以将V元素的质量百分比含量控制为V≤0.1wt％，可以将Cr元素的质量百分比含量控制为Cr≤0.15wt％。
进一步地，在本发明所述的低合金高强钢中，其各化学元素的质量百分含量满足下述各项的至少其中之一：
Ti/(Al+Nb)≥3.0；
1.1％≤Mn+Cr-Ti-Nb-V≤1.6％。
在本发明的化学成分设计中，在控制单一元素含量的同时，还可以控制Ti/(Al+Nb)≥3.0，以调控碳氮化物析出。
在本发明中，需要在Al、Ti、Nb三种元素的形成碳氮化物能力中寻找一定的平衡。Al、Ti、Nb都是碳氮化物形成元素，但是在形成碳氮化物的能力上Nb＞Ti＞Al。若钢中Al元素含量过高而Ti、Nb元素过少，则在制造过程中碳氮化物形成较少，钢材强度不够；而若钢中Nb元素含量过高，虽然在制造过程中更容易形成碳氮化物，但是一方面会造成碳氮化物尺寸粗大，另一方面也会消耗更多的C元素而导致粒状贝氏体形成不足或无法形成；若Ti元素含量过低，而Nb或Al元素含量过高，会使碳氮化物析出与粒状贝氏体相变不平衡而导致产品强度和成型性能劣化。因此，在在本发明所述的低合金高强钢中，还控制Ti/(Al+Nb)≥3.0。在一些实施方案中，控制Ti/(Al+Nb)在3.0～5.5的范围内。
在本发明的化学成分设计中，在控制单一元素含量的同时，还可以控制合金元素配比，使1.1≤Mn+Cr-Ti-Nb-V≤1.6，以调控碳氮化物析出与粒状贝氏体形成处于最优配比。其中Mn和Cr有利于粒状贝氏体形成，Ti、Nb和V则是强碳化物形成元素，因此，上式值越高，越有利于粒状贝氏体形成，上式值越低，则越有利于析出碳氮化物。在一些实施方案中，控制Mn+Cr-Ti-Nb-V在1.1～1.4的范围内。
进一步地，在本发明所述的低合金高强钢中，其碳当量Cev≤0.32，其中Cev＝C+Mn/6+(Cr+Mo+V)/5+(Ni+Cu+Si)/15，其中各化学元素代入其质量百分含量百分号前的数值。
在本发明中，通过控制碳当量Cev≤0.32，可以保证钢材具有优良的可焊接性。在一些实施方案中，Cev为0.30～0.32。
进一步地，在本发明所述的低合金高强钢的不可避免的杂质中，P≤0.02wt％，S≤0.006wt％，N≤0.006wt％。
在本发明中，不可避免的杂质主要包括磷(P)、硫(S)、氮(N)，在技术条件允许的情况下，期望其含量越少越好。
进一步地，在本发明所述的低合金高强钢中，其微观组织中铁素体的体积相比例≥90％，粒状贝氏体体积相比例≥5％，碳氮化物含量≤5％。在一些实施方案中，本发明所述低合金高强钢的微观组织中铁素体和粒状贝氏体的体积相比例之和≥95％，余下为碳氮化物。
进一步地，在本发明所述的低合金高强钢中，所述铁素体中有≥60％的铁素体晶粒长宽比≤4。在一些实施方案中，所述铁素体中有60～90％的铁素体晶粒长宽比≤4。
进一步地，在本发明所述的低合金高强钢中，其横、纵向屈服强度差≤50MPa。在一些实施方案中，本发明所述的低合金高强钢的横、纵向屈服强度差≤40MPa。在一些实施方案中，本发明所述的低合金高强钢的横、纵向屈服强度差≤30MPa。
进一步地，在本发明所述的低合金高强钢中，其横、纵向屈服强度均＞700MPa以上，横、纵向断裂延伸率均≥10％，90°弯曲半径与板厚的比例r/t≤0.3。在一些实施方案中，本发明所述低合金高强钢的横向屈服强度≥705MPa，纵向屈服强度≥720MPa。在一些实施方案中，本发明所述低合金高强钢的横向屈服强度为705～790MPa，纵向屈服强度为720～830MPa。
进一步地，在本发明所述的低合金高强钢中，其横、纵向抗拉强度均≥740MPa。在一些实施方案中，本发明所述低合金高强钢的横向抗拉强度≥745MPa，纵向抗拉强度≥770MPa。在一些实施方案中，本发明所述低合金高强钢的横向抗拉强度为745～830MPa，纵向抗拉强度为770～870MPa。
本发明的另一目的在于提供一种低合金高强钢的制造方法，该制造方法绿色高效，所制得的低合金超高强钢兼具绿色低碳和低合金添加量的优势，且横纵向强度差异低。
为了实现上述目的，本发明提供了一种低合金高强钢的制造方法，其包括步骤：
冶炼和铸造；
热轧；
酸洗冷轧：控制冷轧压下率为1～20％；
退火：将带钢加热至530～680℃并保温，然后以5～20℃/s的第一冷速冷却至440～520℃并保温，然后以1～10℃的第二冷速冷却至室温。
采用本发明所述的制造方法可以在热轧前驱组织中获得少量的粒状贝氏体，再通过优化冷轧变形量和退火工艺，获得相变强化与析出强化最优平衡点，从而获得具有高强度、高成形性能和低横纵向强度差的成本低碳当量且横、纵向差异小的冷轧高强钢。
在本发明的酸洗冷轧步骤中，通过将冷轧压下率控制为1～20％，可以确保后续带钢在退火后，不产生拉长的铁素体晶粒，以避免横纵向强度差异过大。
在本发明的退火步骤中，较低的加热和保温温度可以确保带钢中的粒状贝氏体不发生分解，同时碳氮化物也不会因此而粗化。之后的冷却、保温和终冷工艺是为了进一步控制碳氮化物生长尺寸的前提下获得亚微米尺寸的碳氮化物析出。
进一步地，在本发明所述的制造方法的退火步骤中，将带钢加热至530～680℃并保温90-180s，然后以5～20℃/s的第一冷速冷却至440～520℃并保温5～30s，然后以1～10℃的第二冷速冷却至室温。
进一步地，在本发明所述的制造方法的退火步骤中，在以5～20℃/s的第一冷速冷却至440～520℃并保温的工序段完成带钢表面的热镀锌。
进一步地，在本发明所述的制造方法的热轧步骤中，控制板坯出炉温度为1250～1300℃，精轧终轧温度为890～940℃。
进一步地，在本发明所述的制造方法的热轧步骤中，轧后以≥80℃/s(如80～160℃/s)的冷速将带钢冷却至440～520℃进行卷取。
在本发明的热轧步骤中，较高的板坯出炉温度可以确保板坯中先形成的碳氮化物能充分重溶；较高的终轧温度和超快冷速可以避免碳氮化物批量析出和粗化；通过440～520℃的低温卷取，可以使带钢内形成少量的粒状贝氏体和微量的细小碳氮化物。
本发明所述的横纵向强度差异低的低成本冷轧低合金高强钢及其制造方法具有如下所述的优点以及有益效果：
本发明所述的冷轧低合金高强钢在兼具低合金添加量和低碳当量的同时，还实现了横纵向强度差异低。
在一些实施方式中，其横、纵向屈服强度均≥700MPa，横、纵向断裂延伸率均≥10％，90°弯曲半径与板厚的比例r/t≤0.3的冷轧高强钢，且横纵向屈服强度差≤50MPa，可广泛应用于各类汽车结构件的制备。
附图说明
图1显示了实施例1的微观组织形貌。
具体实施方式
下面将结合具体的实施例对本发明所述的横纵向强度差异低的低成本冷轧低合金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1-6以及对比例1-3
表1-1和表1-2列出了实施例1-6的低成本冷轧低合金高强钢和对比例1-3的对比钢中各化学元素的质量百分比。
表1-1.(余量为Fe和除了P、S、N以外的其他不可避免的杂质)
表1-2.(余量为Fe和除了P、S、N以外的其他不可避免的杂质)
本发明实施例1-6的低成本冷轧低合金高强钢和对比例1-3的对比钢均采用以下步骤制得：
(1)冶炼和连铸；
(2)热轧：控制板坯出炉温度为1250～1300℃，精轧终轧温度为890～940℃；轧后采用超快冷却低温卷取方式，以≥80℃/s的冷速将钢带冷却至440～520℃卷取；
(3)酸洗冷轧：冷轧压下率控制为1～20％；
(4)退火：将带钢加热至530～680℃并保温90-180s，然后以5～20℃/s的第一冷速冷却至440～520℃并保温5～30s，然后以1～10℃的第二冷速冷却至室温。
此外在一些实施方式中，还可以在步骤(4)中以5～20℃/s的第一冷速冷却至440～520℃并保温的工序段完成带钢表面的热镀锌，从而获得热镀锌产品。
需要说明的是，本发明实施例1-6的成分和工艺均符合本发明要求，而对比例1的成分与实施例6相同，其工艺参数不符合本发明，对比例2-3的成分和工艺均有不符合本发明之处。
表2-1和表2-2列出了实施例1-6的低成本冷轧低合金高强钢和对比例1-3的对比钢的具体工艺参数。
表2-1.
表2-2.退火工艺

为了验证本发明的实施效果，对本发明实施例1-6的低成本冷轧低合金高强钢和对比例1-3的对比钢进行采样，并采用4％硝酸酒精溶液对磨抛后的金相试样侵蚀10s左右，清洗吹干后利用光学显微镜或扫描电子显微镜进行微观组织观察，并将观察结果列于表3中。
表3列出了本发明实施例1-6的低成本冷轧低合金高强钢和对比例1-3的对比钢的微观组织观察结果。
表3.
从上述表3中可以看出，本发明实施例1-6的横纵向强度差异低的低成本冷轧低合金高强钢的微观组织含有铁素体+粒状贝氏体+析出的碳氮化物，其铁素体的体积相比例均大于等于90％，粒状贝氏体的比例均大于等于5％，碳氮化物的比例均小于等于5％。
此外，图1显示了实施例1的微观组织形貌。
如图1所述，实施例1的微观组织为铁素体+粒状贝氏体+碳氮化物，其中碳氮化物比例为5.5％，粒状贝氏体比例为4％，剩余为铁素体。
对实施例1-6的低成本冷轧低合金高强钢和对比例1-3的对比钢进行再次取样，并对其进行各项性能测试，并将测试结果列于表4中。其中，各项性能测试包括：
拉伸性能测：采用GB/T228.1-2021《金属材料拉伸试验第1部分:室温试验方法》对实施例1-6的横纵向强度差异低的低成本冷轧低合金高强钢和对比例1-3的对比钢进行测试。
弯曲性能测试：采用GB/T 232-2010《金属材料弯曲试验方法》对实施例1-6的横纵向强度差异低的低成本冷轧低合金高强钢和对比例1-3的对比钢进行测试。
表4.
从上述表4可以看出，本发明实施例1-6的横纵向强度差异低的低成本冷轧低合金高强钢的横、纵向屈服强度差均小于50MPa，其横、纵向屈服强度均大于700MPa以上，横、纵向断裂延伸率均大于10％，90°弯曲半径r(单位：mm)与板厚t(单位：mm)的比例r/t均小于等于0.3。
需要说明的是，本案中各技术特征的组合方式并不限本案权利要求中所记载的组合方式或是具体实施例所记载的组合方式，本案记载的所有技术特征可以以任何方式进行自由组合或结合，除非相互之间产生矛盾。
还需要注意的是，以上所列举的实施例仅为本发明的具体实施例。显然本发明不局限于以上实施例，随之做出的类似变化或变形是本领域技术人员能从本发明公开的内容直接得出或者很容易便联想到的，均应属于本发明的保护范围。

Claims (15)

  1. 一种横纵向强度差异低的低成本冷轧低合金高强钢，其含有Fe及不可避免的杂质，其特征在于，其还含有质量百分含量如下的下述各化学元素：
    C：0.045～0.085％；
    Mn：1.16～1.50％；
    Al：0.02～0.05％；
    Ti：0.10～0.16％；
    0＜Si≤0.2％；
    其微观组织含有铁素体+粒状贝氏体+析出的碳氮化物。
  2. 如权利要求1所述的低合金高强钢，其特征在于，其各化学元素质量百分含量为：
    C：0.045～0.085％；
    Mn：1.16～1.50％；
    Al：0.02～0.05％；
    Ti：0.10～0.16％；
    0＜Si≤0.2％；
    余量为Fe和不可避免的杂质。
  3. 如权利要求1或2所述的低合金高强钢，其特征在于，其还含有下述元素的至少其中之一：
    0＜Nb≤0.02wt％；
    0＜Ce≤0.006wt％；
    0＜La≤0.006wt％；
    0＜V≤0.1wt％；
    0＜Cr≤0.15wt％。
  4. 如权利要求1或2所述的低合金高强钢，其特征在于，其各化学元素的质量百分含量满足下述各项的至少其中之一：
    Ti/(Al+Nb)≥3.0，3.0≤Ti/(Al+Nb)≤5.5优选；
    1.1％≤Mn+Cr-Ti-Nb-V≤1.6％，优选1.1％≤Mn+Cr-Ti-Nb-V≤1.4％。
  5. 如权利要求1或2所述的低合金高强钢，其特征在于，其碳当量Cev≤0.32，优选0.30～0.32；其中Cev＝C+Mn/6+(Cr+Mo+V)/5+(Ni+Cu+Si)/15，其中各化学元素代入其质量百分含量百分号前的数值。
  6. 如权利要求1或2所述的低合金高强钢，其特征在于，在其不可避免的杂质中，P≤0.02wt％，S≤0.006wt％，N≤0.006wt％。
  7. 如权利要求1或2所述的低合金高强钢，其特征在于，其微观组织中铁素体的体积相比例≥90％，粒状贝氏体的体积相比例≥5％，碳氮化物含量≤5％。
  8. 如权利要求1或2所述的低合金高强钢，其特征在于，所述铁素体中有≥60％、优选60～90％的铁素体晶粒长宽比≤4。
  9. 如权利要求1或2所述的低合金高强钢，其特征在于，其横、纵向屈服强度差≤50MPa。
  10. 如权利要求1或2所述的低合金高强钢，其特征在于，其横、纵向屈服强度均＞700MPa以上，横、纵向断裂延伸率均≥10％，90°弯曲半径与板厚的比例r/t≤0.3。
  11. 如权利要求1-10中任意一项所述的低合金高强钢的制造方法，其特征在于，其包括步骤：
    冶炼和铸造；
    热轧；
    酸洗冷轧：控制冷轧压下率为1～20％；
    退火：将带钢加热至530～680℃并保温，然后以5～20℃/s的第一冷速冷却至440～520℃并保温，然后以1～10℃的第二冷速冷却至室温。
  12. 如权利要求11所述的制造方法，其特征在于，在退火步骤中，将带钢加热至530～680℃并保温90-180s，然后以5～20℃/s的第一冷速冷却至440～520℃并保温5～30s，然后以1～10℃的第二冷速冷却至室温。
  13. 如权利要求11所述的制造方法，其特征在于，在退火步骤中，在以5～20℃/s的第一冷速冷却至440～520℃并保温的工序段完成带钢表面的热镀锌。
  14. 如权利要求11所述的制造方法，其特征在于，在热轧步骤中，控制板坯出炉温度为1250～1300℃，精轧终轧温度为890～940℃。
  15. 如权利要求11-14中任意一项所述的制造方法，其特征在于，在热轧步骤中，轧后以≥80℃/s的冷速将带钢冷却至440～520℃进行卷取。
PCT/CN2025/103371 2024-06-26 2025-06-25 横纵向强度差异低的低成本冷轧低合金高强钢及其制造方法 Pending WO2026002018A1 (zh)

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JP2018070920A (ja) * 2016-10-26 2018-05-10 新日鐵住金株式会社 高張力鋼板の製造方法
CN108350550A (zh) * 2015-11-20 2018-07-31 Posco公司 剪切加工性优异的高强度冷轧钢板及其制造方法
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