WO2017215478A1 - 一种高强高韧不锈钢及其加工方法 - Google Patents

一种高强高韧不锈钢及其加工方法 Download PDF

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WO2017215478A1
WO2017215478A1 PCT/CN2017/087156 CN2017087156W WO2017215478A1 WO 2017215478 A1 WO2017215478 A1 WO 2017215478A1 CN 2017087156 W CN2017087156 W CN 2017087156W WO 2017215478 A1 WO2017215478 A1 WO 2017215478A1
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stainless steel
martensite
deformation
strength
cooling
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PCT/CN2017/087156
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English (en)
French (fr)
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刘嘉斌
王宏涛
方攸同
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浙江大学
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Priority to JP2018522630A priority Critical patent/JP6605139B2/ja
Priority to US15/778,001 priority patent/US11401566B2/en
Publication of WO2017215478A1 publication Critical patent/WO2017215478A1/zh

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    • 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 by deformation combined with, or followed by, heat treatment
    • C21D8/005—Modifying the physical properties by deformation combined with, or followed by, heat treatment of ferrous alloys
    • 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
    • C21D1/18—Hardening; Quenching with or without subsequent tempering
    • 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
    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/004—Heat treatment of ferrous alloys containing Cr and Ni
    • 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 by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0221—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
    • C21D8/0236—Cold rolling
    • 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
    • C22C38/001—Ferrous alloys, e.g. steel alloys containing N
    • 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
    • C22C38/002—Ferrous alloys, e.g. steel alloys containing In, Mg, or other elements not provided for in one single group C22C38/001 - C22C38/60
    • 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
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • 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
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/58—Ferrous alloys, e.g. steel alloys containing chromium with nickel with more than 1.5% by weight of manganese
    • 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
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/001—Austenite
    • 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
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/008—Martensite

Definitions

  • the invention relates to a high-strength and high-toughness stainless steel material and a processing method thereof, in particular to a stainless steel with high yield strength and high elongation and a processing method thereof.
  • the automobile industry is an important pillar influencing the development of the national economy, technological progress and social modernization. China has clearly proposed to accelerate the development of the automobile industry.
  • the development of the automotive industry is inseparable from the steel materials.
  • the automobile manufacturing industry is the largest user of thin steel sheets. In order to reduce fuel consumption and save energy, automobiles must be developed to be lighter. Therefore, higher and higher requirements are imposed on automotive steel sheets.
  • the German standard is used for hot-dip galvanizing and plating.
  • Fukang car steel plate adopts French standard, mainly hot-dip galvanizing;
  • Xiali car steel plate adopts Japanese standard, mainly hot-dip galvanizing alloy;
  • Cherokee Jeep steel is made of American standard, hot-dip galvanized and hot-dip galvanized alloy.
  • Hot-rolled high-strength steel sheets are mostly used for parts with large stresses such as frames and longitudinal beams of different models. These steel plates are used in large quantities on trucks, accounting for about 60% of the total number of hot-rolled steel sheets for trucks. 70%. Therefore, it is necessary to have higher strength and better formability. Generally, the highest strength of hot-rolled steel sheets for automobiles is 500 MPa. Although the addition of microalloying elements such as Nb and Ti is high in strength, it has an influence on formability, thereby limiting its application. At present, in order to improve the strength level, duplex steel and TRIP steel have been developed.
  • Dual-phase steel (high-strength steel) ferritic-martensitic composite steel sheets are characterized by a distribution of approximately 15% of the hard phase on a fine ferrite matrix and further strengthening by solid solution atoms.
  • the production process is that when the steel sheet is rolled, it stays in the two-phase region of ferrite and austenite for a period of time, a large amount of phase ferrite is precipitated, and the C concentration of the remaining phase austenite is increased, and then the rapid cooling method is adopted.
  • the austenite structure is transformed into a martensite structure.
  • the steel is mainly a ferrite phase, about 80% to 90%, its total elongation is high, and at the same time, due to volume expansion from austenite to martensite, dislocations are formed in the surrounding portion, which lowers the yield strength.
  • the tensile strength of the ferritic-martensitic composite structural steel sheet with good forming properties is between 550 and 650 MPa, and the newly developed martensitic composite hot-rolled steel sheet has a maximum strength of 780 MPa and an elongation of 21%.
  • TRIP steel (high-strength steel) can not reach the strength level of 800MPa.
  • phase change induced plastic steel referred to as TRIP steel
  • TRIP Steel is the star of hope for high-stretch high-strength steel sheets.
  • the composition of TRIP steel is dominated by C-Mn-Si alloy system.
  • the sub-features are low carbon, low alloying and steel purity.
  • the production process is a hot-rolling process in which a double-phase zone critical annealing and a bainite transformation zone heat treatment process have ferrite, bainite and about 10% residual austenite three-phase structure with a residual of 10%.
  • the retained austenite gradually transforms into a martensite structure, and the total elongation is improved by the local deformation caused by the hardening.
  • High strength comes from the combined contribution of solid solution strengthening of martensite, bainite and alloying elements.
  • the performance variation range of TRIP steel is: yield strength 340MPa ⁇ 860MPa, tensile strength 610MPa10 ⁇ 80MPa, elongation 22% ⁇ 37%.
  • the steel sheet is generally required to have a yield strength of 1000 MPa or more and an elongation of not less than 30%.
  • Such high performance indicators are performance levels that are not achievable with both dual phase steels and TRIP. Therefore, it is urgent to develop new advanced high strength and high toughness steel.
  • the object of the present invention is to provide a high-strength and high-toughness stainless steel and a processing method thereof, which solves the contradiction between the strength and the plasticity inherent in the conventional material processing technology.
  • the inventors conducted various studies and found that in the 3XX series stainless steel, in addition to specifying the basic components of the base material, it is necessary to define the organization and processing conditions, and to utilize the dispersed distribution of nano-martens and interface elements.
  • the martensite deactivation effect caused by segregation is used to achieve high strength and toughness comprehensive performance.
  • the present invention provides a stainless steel that includes the following features:
  • the stainless steel contains 0.01% to 0.1% by weight of C, 0.05% to 0.2% of N, not more than 0.03% of P, not more than 0.003% of S, and 0.5 % to 1% of Si, 1.0 % to 2.0% Mn, 15% to 17% Cr, 5% to 7% Ni, the balance being Fe; among the chemical components, P and S are impurities;
  • the stainless steel comprises austenite and strain-induced martensite structure, wherein the martensite is an irregular approximate spindle shape, and the long axis average size is between 50 and 1000 nm, and the short axis average size is 20 to Between 500nm, the volume percentage of martensite in stainless steel is 0.1 % -20%; there is an elemental segregation layer at the interface between martensite and austenite.
  • the thickness of the segregation layer is 1-20nm, Ni in the layer,
  • the content of Mn, N, and Si elements is 1.2 to 3 times the average content of each element in stainless steel, respectively.
  • the invention provides a processing method of the stainless steel, comprising the following steps:
  • the material with the chemical element composition meets the requirements for solution treatment and cooling to obtain a sample;
  • the chemical element composition of the raw material is: C% by weight of 0.01% to 0.1%, 0.05% to 0.2% of N, not P above 0.03%, not higher than 0.003% S, 0.5 % to 1% Si, 1.0 % to 2.0% Mn, 15% to 17% Cr, 5% to 7% Ni, and the rest is Fe ;
  • ⁇ is the cross-sectional shrinkage rate
  • ⁇ is a parameter related to stacking fault energy (SFE), and ⁇ can be obtained by referring to data corresponding to the layer fault energy and ⁇ .
  • SFE stacking fault energy
  • ⁇ is a parameter related to the martensitic phase change kinetic energy of the sample material, and ⁇ can be obtained by referring to the data corresponding to the chemical driving energy and ⁇ .
  • n refers to the pre-factor, usually 2;
  • step (a) in order to obtain a high-strength and expanded austenite region, C and N elements are added to the raw material, but when the C content exceeds 0.1% or the N content exceeds 0.2%, Cr carbides are precipitated at the grain boundaries. Reduce the plasticity of the steel, so the upper limit is set to 0.1% and 0.2%, respectively.
  • Cr and Ni elements are added to the raw material, but too much addition of Cr and Ni elements will result in a material stacking fault energy that is too high to cause martensite transformation at room temperature. Adding too little will cause the material to transform into martensite too early in the cooling process, and therefore the Cr and Ni element contents are limited to 15% to 17% and 5% to 7%, respectively.
  • the temperature of the solution treatment is 1050 ° C to 1150 ° C
  • the holding time is 1 min to 2 h
  • the cooling method is quenching or quenching.
  • the present invention utilizes the strain-induced martensite effect mentioned above, and step (b) deforms the sample to a certain extent at room temperature to convert part of austenite to martensite. If the martensite content is too small, the strengthening effect is not obvious, but if the martensite content is too high, the plasticity will be seriously impaired, so the martensite content is controlled to be 1% to 20%.
  • the strain-induced martensite produced by the multi-pass small deformation is an irregular approximate spindle shape, and the long-axis average size is 50-1000 nm, and the short-axis dimension The average is 20 to 500 nm.
  • the deformation mode is rolling, stamping, forging or drawing.
  • the stacking fault energy and phase change driving energy are determined by the chemical composition of the material.
  • the chemical composition of the material is determined, the corresponding stacking fault energy and phase change driving energy are also determined.
  • the stacking fault energy can be calculated by formula (2):
  • %Ni, %Cr, %Mn, and %Mo represent the respective weight percentages of these elements in the stainless steel.
  • the sample after the multi-pass small deformation in the step (b) is subjected to long-time low-temperature annealing due to Ni, Mn, Si. , N is an austenite magnifying element, it The energy in martensite is higher than that in austenite, and tends to diffuse from martensite to austenite.
  • the heating method is to increase the temperature with the furnace, and the cooling method is cooling with the furnace or air cooling.
  • the method for processing the stainless steel consists of steps (a) to (c).
  • the invention has the beneficial effects that the yield strength of the stainless steel prepared by the room temperature deformation reaches 600 MPa or more and the elongation reaches 30%; the yield strength of the subsequent annealing sample is increased to 1000 MPa or more and the elongation is maintained at 30% or more. Therefore, the stainless steel produced by the invention not only has high strength and high toughness, but also avoids the contradiction between strength and plasticity inherent in the conventional material processing technology.
  • 1 is an engineering stress-strain curve according to Embodiment 1 of the present invention, wherein 1 is a sample obtained by processing deformation at room temperature; and 2 is a sample obtained by annealing at room temperature for a long time after low-temperature annealing.
  • 2 is an X-ray line of the present invention
  • 1 is a sample obtained by processing deformation at room temperature in Example 1
  • 2 is a sample obtained by processing deformation at room temperature in Example 17.
  • Fig. 3 is a photograph of a dark field of a transmission electron microscope center of a sample obtained by deformation at room temperature according to Example 3 of the present invention, in which the white light region is martensite.
  • 4a is a distribution diagram of interface element segregation regions of a three-dimensional atom probe result of a sample obtained after room temperature deformation and low-temperature heat treatment according to Example 1 of the present invention.
  • 4b is a distribution curve of Cr, Ni, Mn, Si, and N elements at the substrate and the interface of the sample obtained by the three-dimensional atom probe obtained after the room temperature deformation and low-temperature heat treatment according to Example 1 of the present invention.
  • the tensile test is carried out according to "GB/T 228.1-2010 tensile test of metallic materials Part 1: room temperature test method", and the yield strength and elongation of the test specimen are tested.
  • the composition around the martensite in the sample was tested using a three-dimensional atom probe.
  • the material composition used is 0.1% C, 0.2% N, 0.03% P, 0.003% S, 0.5% Si, 1.0% Mn, 15% Cr, 5% Ni, and the rest are Fe. The same as in the first embodiment.
  • the material composition used is 0.1% C, 0.05% N, 0.03% P, 0.003% S, 0.5% Si, 1.0% Mn, 15% Cr, 5% Ni, and the rest are Fe. The same as in the first embodiment.
  • the material composition used is 0.05% C, 0.1% N, 0.02% P, 0.001% S, 0.5% Si, 1.0% Mn, 15% Cr, 5% Ni, and the rest are Fe. The same as in the first embodiment.
  • the material composition used is 0.05% C, 0.15% N, 0.02% P, 0.001% S, 1% Si, 2.0% Mn, 15% Cr, 5% Ni, and the rest are Fe. The same as in the first embodiment.
  • the material composition used is 0.05% C, 0.15% N, 0.02% P, 0.001% S, 0.5% Si, 1.0% Mn, 17% Cr, 5% Ni, and the rest are Fe. The same as in the first embodiment.
  • the material composition used is 0.05% C, 0.15% N, 0.02% P, 0.001% S, 0.5% Si, 1.0% Mn, 15% Cr, 7% Ni, and the rest are Fe. The same as in the first embodiment.
  • the material composition used is 0.05% C, 0.15% N, 0.01% P, 0.001% S, 0.8% Si, 1.5% Mn, 16% Cr, 6% Ni, and the rest are Fe. The same as in the first embodiment.
  • the solution temperature used was 1150 ° C, the time was 1 min, and the cooling method was quenching, and the other contents were the same as in Example 1.
  • the solution temperature used was 1100 ° C, the time was 30 min, and the cooling method was quenching, and the other contents were the same as in Example 1.
  • the room temperature deformation mode used was punching instead of rolling, and the other contents were the same as in the first embodiment.
  • the room temperature deformation mode used was forging instead of rolling, and the other contents were the same as those of the first embodiment.
  • the room temperature deformation mode used was drawing rather than rolling, and the other contents were the same as in the first embodiment.
  • the room temperature rolling pass deformation amount used was 0.01, and the other contents were the same as in the first embodiment.
  • the room temperature rolling pass deformation amount was 0.1, and the cumulative deformation amount was 0.3, and the other contents were the same as in the first embodiment.
  • the room temperature rolling pass deformation amount was 0.01, and the cumulative deformation amount was 0.1, and the other contents were the same as in the first embodiment.
  • the material composition used is 0.2% C, 0.25% N, 0.03% P, 0.003% S, 0.5% Si, 1.0% Mn, 15% Cr, 5% Ni, and the rest are Fe. The same as in the first embodiment.
  • the material composition used is 0.05% C, 0.1% N, 0.03% P, 0.003% S, 0.5% Si, 1.0% Mn, 20% Cr, 5% Ni, and the rest are Fe. The same as in the first embodiment.
  • the material composition used is 0.05% C, 0.1% N, 0.03% P, 0.003% S, 0.5% Si, 1.0% Mn, 17% Cr, 9% Ni, and the rest are Fe. The same as in the first embodiment.
  • the material composition used is 0.05% C, 0.1% N, 0.03% P, 0.003% S, 0.5% Si, 1.0% Mn, 13% Cr, 5% Ni, and the rest are Fe. The same as in the first embodiment.
  • the material composition used was 0.05% C, 0.1% N, 0.03% P, 0.003% S, 0.5% Si, 1.0% Mn, 17% of Cr, 3% of Ni, and the balance of Fe were the same as in Example 1.
  • the amount of deformation per pass of the room temperature rolling was 0.2, and the other contents were the same as in Example 1.
  • Table 1 shows the martensite content and size of the sample obtained by long-term low temperature annealing in the above examples.
  • Table 2 shows the yield strength and elongation of the sample obtained by the room temperature deformation in the above embodiment and the sample obtained by long-time low-temperature annealing.
  • Table 3 is the elemental content of the three-dimensional atom probe test of the sample obtained by long-term low temperature annealing in the above embodiment.
  • Examples 1 to 8 are examples in which the influence of the composition on the form, content and size of martensite is examined.
  • spindle martensite was obtained, and the content was between 1% and 20%, the major axis was between 100 and 1000 nm, and the short axis was between 20 and 500 nm.
  • Comparative Example 1 produced a large amount of Cr compounds due to excessive C and N contents; Comparative Example 2 expanded the ferrite region due to the high Cr content, resulting in austenite regions being too small, high martensite content and mutual Connected together to form massive martensite; Comparative Example 3 due to the high Ni content, significantly expanding the austenite region and the austenite is too stable to be strain-induced martensite effect during the room temperature deformation process, no Markov in the tissue Comparative Example 4: Because the Cr content is too low, the Ni content is too high, so that the austenite is too stable to be strain-induced martensite effect in the process of room temperature deformation, no martensite in the structure; Comparative Example 5 due to Ni The content is too low, the austenite is too unstable, and completely transformed into martensite structure during the solution cooling process; the above results show that only the material composition meets the scope disclosed in the present invention, and a reasonable martensite content and size can be obtained.
  • Examples 1, 9, and 10 are examples in which the effect of the solution treatment on the microstructure and properties of the material is examined. Whether it is quenching or quenching, as long as the microstructure of the martensite morphology, content and size can be obtained within the temperature and time range specified by the present invention, the strength is significantly improved and does not decrease after the low-temperature heat treatment. Excellent mechanical properties of elongation.
  • Examples 1, 11 to 13 are examples in which the effect of the room temperature deformation mode on the microstructure and properties of the material was examined. Whether it is rolling, extrusion, forging or drawing, it can obtain the ideal microstructure of martensite morphology, content and size, and exhibits excellent mechanical properties after significantly low heat treatment without increasing the elongation. .
  • Example 1, 14 to 17 are examples of the effects of room temperature deformation and cumulative deformation on the microstructure and properties of the material. child.
  • the microstructure of the desired martensite morphology, content and size can be obtained as long as it is within the range of the ball deformation amount and the cumulative deformation amount defined by the present invention, and exhibits a remarkable increase in strength after low-temperature heat treatment without lowering the elongation.
  • the excellent mechanical properties of the rate It can be seen from Example 1 and Example 14 that the smaller the amount of pass deformation, the smaller the obtained martensite size, and the stronger the reinforcing effect.
  • Examples 1, 18 to 20 are examples of the effects of annealing temperature and time on the microstructure and properties of the material after room temperature deformation.
  • the annealing temperature range and time range defined by the present invention can obtain an ideal microstructure of martensite morphology, content and size, and exhibit excellent strength after low-temperature heat treatment and excellent elongation without lowering the elongation performance.
  • the annealing temperature exceeds the limited range, such as Comparative Example 8, it will cause the martensite to reverse phase back to austenite and produce a Cr compound, which seriously deteriorates the material properties.
  • the annealing time is too short, such as Comparative Example 9, the alloying elements have not yet reached diffusion and enrichment. As shown in Table 3, no significant elemental segregation layer was produced, and the effect of significantly increasing the strength without lowering the elongation could not be achieved.

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Abstract

一种不锈钢及其加工方法，该不锈钢含有重量百分比为0.01％～0.1％的C，0.05％～0.2％的N，不高于0.03％的P，不高于0.003％的S，0.5％～1％的Si，1.0％～2.0％的Mn，15％～17％的Cr，5％～7％的Ni，其余为Fe；包含奥氏体和应变诱发马氏体组织，其中马氏体为不规则的近似纺锤体形状，其长轴平均尺寸在50～1000nm之间，短轴平均尺寸在20～500nm之间，马氏体在不锈钢中的体积百分比为0.1％～20％；马氏体与奥氏体的界面存在一个元素偏聚层，该偏聚层厚度为1～20nm，层内Ni、Mn、N、Si元素的含量分别是各元素在不锈钢中平均含量的1.2～3倍。该不锈钢具备高强高韧性，屈服强度提高至1000MPa以上且延伸率保持在30％以上，避免了传统材料加工技术固有的强度与塑性此消彼长的矛盾。

Description

一种高强高韧不锈钢及其加工方法 技术领域
本发明涉及一种高强高韧不锈钢材料及其加工方法，具体来说是一种具有高屈服强度高延伸率的不锈钢及其加工方法。
技术背景
汽车工业是影响国民经济发展、技术进步和社会现代化的支柱产业具有重要地位，我国明确提出要加速发展汽车工业。汽车工业的发展与钢铁材料是密不可分的，汽车制造行业是薄钢板的最大用户，为了降低油耗，节约能源，汽车要向轻量化发展，因此对汽车用钢板提出了越来越高的要求。
一汽奥迪A6、上海大众B5、上海通用的Buick轿车及一些家用轿车等就大量采用了高强度、镀锌钢板以及激光拼焊板等大众系列轿车用钢板采用德国标准，品种涉及热镀锌、电镀锌和高强度钢板等，最宽达1800mm，最厚达4mm；富康轿车用钢板采用法国标准，以热镀锌为主；夏利轿车用钢板采用日本标准，以热镀锌合金化为主；切诺基吉普车用钢采用美国标准，热镀锌和热镀锌合金化并用。
热轧高强度钢板多用于车架等承受应力较大的零件、不同车型的纵横梁等，此类钢板在载货汽车上用量很大，约占载重车用热轧钢板总量的60％－70％。因此既要有较高的强度，又要有较好的成形性。一般汽车用热轧钢板最高强度为500MPa，虽然添加Nb、Ti等微合金元素高其强度，但对成形性有影响，从而限制了它的应用。目前为了提高强度水平，开发了双相钢和TRIP钢。
双相钢(高强度钢)铁素体-马氏体复合组织钢板的特点是在细小的铁素体基体上分布大约15％的硬相，并通过固溶原子进一步强化。其生产工艺是在轧制钢板时，在铁素体和奥氏体的双相区停留一段时间，有相铁素体大量析出，剩余相奥氏体的C浓度增加，然后采用快速冷却方法使奥氏体组织转变为马氏体组织。由于钢中主要是铁素体相，大约80％～90％，其总延伸率较高，同时由于从奥氏体转变为马氏体时体积膨胀而引起周围部分形成位错,使屈服强度降低并具有良好成形性能铁素体-马氏体复合组织钢板的抗拉强度在550～650MPa之间，新研制的马氏体复合组织热轧钢板最高强度为780MPa，延伸率为21％。
TRIP钢(高强度钢)由于成形性的限制，双相钢无法达到800MPa的强度水平，为满足更高强度的需求，并解决强度与塑性矛盾开发了相变诱导塑性钢，简称TRIP钢，TRIP钢是作为高延伸性高强度钢板的希望之星而登场的。TRIP钢的成分以C-Mn-Si合金系统为主其成 分特点是低碳、低合金化和钢质纯净。其生产工艺是在热轧时，采用双相区临界退火和贝氏体转变区保温的热处理工艺具有铁素体、贝氏体和约10％的残余奥氏体三相组织，具有10％的残余奥氏体的热轧钢板在加工成型时，残余奥氏体逐渐转变为马氏体组织，由于硬化克服了局部变形提高了总延伸率。高强度来自于马氏体、贝氏体和合金元素固溶强化的共同贡献。TRIP钢性能变化范围为：屈服强度340MPa～860MPa，抗拉强度610MPa10～80MPa，延伸率22％～37％。
随着汽车轻量化和安全性能要求的进一步提高，对汽车用高强钢，尤其是作为保险杆用钢提出更高的强韧性要求。通常要求钢板具有1000MPa以上的屈服强度和不低于30％的延伸率。如此高的性能指标是上述双相钢和TRIP均无法达到的性能水平。因此迫切需要开发新型先进高强高韧钢。
发明内容
本发明的目的在于提供一种高强高韧不锈钢及其加工方法，以解决传统材料加工技术固有的强度与塑性此消彼长的矛盾。
为解决上述技术问题，发明人进行了种种研究，结果发现，在3XX系不锈钢中，除了规定母材的基本成分以外，还需限定组织和加工条件，利用弥散分布的纳米马氏体以及界面元素偏聚导致的马氏体失活效应来实现高强高韧综合性能。
本发明采用的技术方案如下：
本发明提供了一种不锈钢，该不锈钢包含如下特征：
(1)所述不锈钢含有重量百分比为0.01％～0.1％的C，0.05％～0.2％的N，不高于0.03％的P，不高于0.003％的S，0.5％～1％的Si，1.0％～2.0％的Mn，15％～17％的Cr，5％～7％的Ni，其余为Fe；各化学组分中，P和S为杂质；
(2)所述不锈钢包含奥氏体和应变诱发马氏体组织，其中马氏体为不规则的近似纺锤体形状，其长轴平均尺寸在50～1000nm之间，短轴平均尺寸在20～500nm之间，马氏体在不锈钢中的体积百分比为0.1％～20％；马氏体与奥氏体的界面存在一个元素偏聚层，该偏聚层厚度为1～20nm，层内Ni、Mn、N、Si元素的含量分别是各元素在不锈钢中平均含量的1.2～3倍。
本发明提供了一种所述不锈钢的加工方法，包括如下步骤：
(a)将化学元素组成符合要求的原料进行固溶处理，冷却得到试样；原料的化学元素组成为：重量百分含量为0.01％～0.1％的C，0.05％～0.2％的N，不高于0.03％的P，不高于0.003％的S，0.5％～1％的Si，1.0％～2.0％的Mn，15％～17％的Cr，5％～7％的Ni，其余为Fe；
(b)将试样在室温下进行一定程度的变形，变形过程采用多道次小变形逐步增加变形量的方式进行，每道次的截面收缩率增量为0.01～0.1，累计总截面收缩率应符合公式(1)
1-exp{-β[1-exp(-αε)]n}<0.3  (1)
ε是截面收缩率，
α是与层错能(SFE)相关的参数，通过查阅相关层错能与α对应的数据可获得α，
β是与该试样材料马氏体相变化学驱动能相关参数，通过查阅化学驱动能与β对应的数据可获知β，
n是指前因子，通常取2；
(c)对经步骤(b)处理的试样进行退火处理，退火温度为50～550℃，退火时间为10min～100h，冷却得到不锈钢。
步骤(a)中，为了获得高强度和扩大奥氏体区，原料中添加了C和N元素，但是C含量超过0.1％或N含量超过0.2％时，在晶界上会析出Cr碳化物，降低钢材塑性，因而将其上限分别定为0.1％和0.2％。为了使该不锈钢在室温具有应变诱发马氏体效应，原料中添加了Cr和Ni元素，但是Cr和Ni元素添加太多将导致材料层错能过高而无法在室温发生马氏体相变，添加太少又将导致材料在冷却过程过早相变为马氏体，因此并将Cr和Ni元素含量分别限定在15％～17％和5％～7％。
进一步，步骤(a)中，固溶处理的温度为1050℃～1150℃，保温时间为1min～2h，冷却方式为淬水或者淬油。
本发明为了获得高强度，利用前述所提的应变诱发马氏体效应，步骤(b)在室温对试样进行一定程度的加工变形，使部分奥氏体转变为马氏体。马氏体含量太少则强化效果不明显，但是马氏体含量太高将严重损害塑性，因此将马氏体含量控制在1％～20％。为了有效发挥马氏体的强化作用，通过多道次小变形的变形方式使所产生的应变诱发马氏体为不规则的近似纺锤体形状，其长轴平均尺寸为50～1000nm，短轴尺寸平均为20～500nm。进一步，步骤(b)中，变形方式为轧制、冲压、锻造或拉拔。其中层错能和相变驱动能均由材料的化学成分确定，当材料的化学成分确定时，其对应的层错能和相变驱动能也确定。其中层错能可通过公式(2)计算：
SFE＝-53+6.2(％Ni)+0.7(％Cr)+3.2(％Mn)+9.3(％Mo)  (2)
式(2)中，％Ni、％Cr、％Mn、％Mo表示这些元素各自在不锈钢中的重量百分含量。
本发明步骤(c)中，为在不降低塑性的条件下还能进一步提高屈服强度，对经步骤(b)多道次小变形后的试样进行长时间低温退火，由于Ni、Mn、Si、N是奥氏体扩大化元素，它 们在马氏体中的能量高于其在奥氏体的能量，倾向于由马氏体向奥氏体扩散，通过本发明的50～550℃退火10min～100h处理，使上述元素发生扩散并在马氏体和奥氏体界面富集产生界面偏聚区，该偏聚区的产生将导致马氏体周围被一层更稳定的奥氏体包围和束缚，无法在后续的变形中继续长大而失活，从而使得试样在变形过程中必须重新通过位错运动、塞积以诱发马氏体形核长大。进一步，步骤(c)中，加热方式为随炉升温，冷却方式为随炉冷却或空冷。
本发明优选所述不锈钢的加工方法由步骤(a)～(c)组成。
本发明的有益效果在于：本发明通过室温变形制备的不锈钢屈服强度达到600MPa以上延伸率达到30％；通过后续退火试样屈服强度提高至1000MPa以上且延伸率保持在30％以上。所以，本发明制得的不锈钢不仅具备高强高韧性，且避免了传统材料加工技术固有的强度与塑性此消彼长的矛盾。
附图说明
图1为本发明实施例1的工程应力应变曲线，1为室温加工变形所获得的试样；2为室温加工之后又长时间低温退火所获得的试样。
图2为本发明的X射线谱线，1为实施例1室温加工变形所获得的试样，2为实施例17室温加工变形所获得的试样。
图3为本发明实施例3室温加工变形所获得的试样的透射电镜中心暗场照片，照片中白亮区域为马氏体。
图4a为本发明实施例1室温变形且低温热处理后所获得试样的三维原子探针结果之界面元素偏聚区分布图。
图4b为本发明实施例1室温变形且低温热处理后所获得试样的三维原子探针结果之基体与界面处Cr、Ni、Mn、Si和N元素的分布曲线。
具体实施方式
下面以具体实施例对本发明的技术方案做进一步说明，但本发明的保护范围不限于此：
实施例1：
将含有重量百分比为0.01％的C，0.2％的N，0.03％的P，0.003％的S，0.5％的Si，1.0％的Mn，15％的Cr，5％的Ni，其余为Fe的钢于1050℃保温2h之后淬水。将所得试样于室温下进行多道次轧制变形，每道次的轧制变形量为0.05，累计总变形量为0.2。之后将所得试样于450℃保温24h，空冷。采用透射电子显微镜观察所得试样内部马氏体的形状和尺寸，采 用X射线衍射测量所得试样马氏体的含量。根据《GB/T 228.1-2010金属材料拉伸试验第1部分：室温试验方法》进行拉伸试验，测试试样屈服强度和延伸率。采用三维原子探针测试试样中马氏体周围的成分。
实施例2：
所用材料成分为0.1％的C，0.2％的N，0.03％的P，0.003％的S，0.5％的Si，1.0％的Mn，15％的Cr，5％的Ni，其余为Fe其他内容均与实施例1相同。
实施例3：
所用材料成分为0.1％的C，0.05％的N，0.03％的P，0.003％的S，0.5％的Si，1.0％的Mn，15％的Cr，5％的Ni，其余为Fe其他内容均与实施例1相同。
实施例4：
所用材料成分为0.05％的C，0.1％的N，0.02％的P，0.001％的S，0.5％的Si，1.0％的Mn，15％的Cr，5％的Ni，其余为Fe其他内容均与实施例1相同。
实施例5：
所用材料成分为0.05％的C，0.15％的N，0.02％的P，0.001％的S，1％的Si，2.0％的Mn，15％的Cr，5％的Ni，其余为Fe其他内容均与实施例1相同。
实施例6：
所用材料成分为0.05％的C，0.15％的N，0.02％的P，0.001％的S，0.5％的Si，1.0％的Mn，17％的Cr，5％的Ni，其余为Fe其他内容均与实施例1相同。
实施例7：
所用材料成分为0.05％的C，0.15％的N，0.02％的P，0.001％的S，0.5％的Si，1.0％的Mn，15％的Cr，7％的Ni，其余为Fe其他内容均与实施例1相同。
实施例8：
所用材料成分为0.05％的C，0.15％的N，0.01％的P，0.001％的S，0.8％的Si，1.5％的Mn，16％的Cr，6％的Ni，其余为Fe其他内容均与实施例1相同。
实施例9：
所用固溶温度为1150℃，时间为1min，冷却方式为淬水，其他内容均与实施例1相同。
实施例10：
所用固溶温度为1100℃，时间为30min，冷却方式为淬油，其他内容均与实施例1相同。
实施例11：
所用室温变形方式为冲压而非轧制，其他内容均与实施例1相同。
实施例12：
所用室温变形方式为锻造而非轧制，其他内容均与实施例1相同。
实施例13：
所用室温变形方式为拉拔而非轧制，其他内容均与实施例1相同。
实施例14：
所用室温轧制道次变形量为0.01，其他内容均与实施例1相同。
实施例15：
所用室温轧制道次变形量为0.1，累计变形量为0.3，其他内容均与实施例1相同。
实施例16：
所用室温轧制累计变形量为0.15，其他内容均与实施例1相同。
实施例17：
所用室温轧制道次变形量为0.01，累计变形量为0.1，其他内容均与实施例1相同。
实施例18：
室温轧制后所用退火工艺为550℃退火10min，其他内容均与实施例1相同。
实施例19：
室温轧制后所用退火工艺为50℃退火100h，其他内容均与实施例1相同。
实施例20：
室温轧制后所用退火工艺为150℃退火50h，其他内容均与实施例1相同。
对比例1：
所用材料成分为0.2％的C，0.25％的N，0.03％的P，0.003％的S，0.5％的Si，1.0％的Mn，15％的Cr，5％的Ni，其余为Fe其他内容均与实施例1相同。
对比例2：
所用材料成分为0.05％的C，0.1％的N，0.03％的P，0.003％的S，0.5％的Si，1.0％的Mn，20％的Cr，5％的Ni，其余为Fe其他内容均与实施例1相同。
对比例3：
所用材料成分为0.05％的C，0.1％的N，0.03％的P，0.003％的S，0.5％的Si，1.0％的Mn，17％的Cr，9％的Ni，其余为Fe其他内容均与实施例1相同。
对比例4：
所用材料成分为0.05％的C，0.1％的N，0.03％的P，0.003％的S，0.5％的Si，1.0％的Mn，13％的Cr，5％的Ni，其余为Fe其他内容均与实施例1相同。
对比例5：
所用材料成分为0.05％的C，0.1％的N，0.03％的P，0.003％的S，0.5％的Si，1.0％的 Mn，17％的Cr，3％的Ni，其余为Fe其他内容均与实施例1相同。
对比例6：
所用室温轧制每道次变形量为0.2，其他内容均与实施例1相同。
对比例7：
所用室温轧制累计变形量为0.5，其他内容均与实施例1相同。
对比例8：
室温轧制后所用退火工艺为650℃退火24h，其他内容均与实施例1相同。
对比例9：
室温轧制后所用退火工艺为250℃退火5min，其他内容均与实施例1相同。
上述实施例的结果如表1和表2所示：
表1为上述实施例长时间低温退火所得试样的马氏体含量和尺寸
Figure PCTCN2017087156-appb-000001
Figure PCTCN2017087156-appb-000002
表2为上述实施例室温变形所得试样和长时间低温退火所得试样的屈服强度与延伸率
Figure PCTCN2017087156-appb-000003
Figure PCTCN2017087156-appb-000004
表3为上述实施例长时间低温退火所得试样的三维原子探针测试所得元素含量
Figure PCTCN2017087156-appb-000005
结果分析如下：
实施例1～8是考察成分对马氏体形态、含量和尺寸的影响作用的例子。实施例1-8均得到了纺锤体马氏体，且含量在1％～20％之间，长轴在100～1000nm，短轴在20～500nm。而对比例1由于C和N含量过高，产生了大量Cr化合物；对比例2由于Cr含量过高，扩大了铁素体区，导致奥氏体区太小，马氏体含量较高且相互连接在一起形成块状马氏体；对比例3由于Ni含量过高，显著扩大奥氏体区且奥氏体过于稳定无法在室温加工变形阶段发生应变诱发马氏体效应，组织中无马氏体；对比例4由于Cr含量太低，相当于Ni含量过高，使得奥氏体过于稳定无法在室温加工变形阶段发生应变诱发马氏体效应，组织中无马氏体；对比例5由于Ni含量太低，奥氏体过于不稳定，在固溶冷却过程完全转变为马氏体组织；以上结果说明只有材料成分符合本发明所公开的范围，才能获得合理的马氏体含量和尺寸。
实施例1,9,10是考察固溶处理方式对材料微结构和性能的影响作用的例子。无论是淬水还是淬油，只要在本发明规定的保温温度和时间范围内都能获得理想的马氏体形态、含量和尺寸的微观结构，并在低温热处理后表现出强度显著提高而且不降低延伸率的优异的力学性能。
实施例1，11～13是考察室温变形方式对材料微结构和性能影响作用的例子。无论是轧制、挤压、锻压还是拉拔均能获得理想的马氏体形态、含量和尺寸的微观结构，并在低热温处理后表现出强度显著提高而且不降低延伸率的优异的力学性能。
实施例1,14～17是考察室温道次变形量和累计变形量对材料微结构和性能影响作用的例 子。只要是在本发明所限定的道次变形量范围和累计变形量范围，均可获得理想的马氏体形态、含量和尺寸的微观结构，并在低温热处理后表现出强度显著提高而且不降低延伸率的优异的力学性能。由实施例1和实施例14可知，道次变形量越小，所得马氏体尺寸越小，强化效果相对越明显。当道次变形量超过限定范围，如对比例6，则无法实现低温处理后表现出强度显著提高而且不降低延伸率的效果。由实施例1和实施例15可知，在所限定范围内累计变形量越大，马氏体含量越高，强化效果也越明显，且能保持低温退火后延伸率不降低的特性。当累计变形量超过限定，如对比例7，虽然马氏体含量高强化明显，但是延伸率很低，无法达到高强高韧目标。
实施例1，18～20是考察室温变形后退火温度和时间对材料微结构和性能影响作用的例子。只要在本发明所限定的退火温度范围和时间范围，均可得理想的马氏体形态、含量和尺寸的微观结构，并在低温热处理后表现出强度显著提高而且不降低延伸率的优异的力学性能。若退火温度超过限定范围，如对比例8，则将导致马氏体逆相变回奥氏体，并产生Cr化合物，严重劣化材料性能。若退火时间过短，如对比例9，则合金元素尚未来得及扩散富集。如表3所示，尚未产生明显的元素偏聚层，无法实现强度显著提高而且不降低延伸率的效果。

Claims (7)

  1. 一种不锈钢，该不锈钢包含如下特征：
    (1)所述不锈钢含有重量百分比为0.01％～0.1％的C，0.05％～0.2％的N，不高于0.03％的P，不高于0.003％的S，0.5％～1％的Si，1.0％～2.0％的Mn，15％～17％的Cr，5％～7％的Ni，其余为Fe；
    (2)所述不锈钢包含奥氏体和应变诱发马氏体组织，其中马氏体为不规则的近似纺锤体形状，其长轴平均尺寸在50～1000nm之间，短轴平均尺寸在20～500nm之间，马氏体在不锈钢中的体积百分比为0.1％～20％；马氏体与奥氏体的界面存在一个元素偏聚层，该偏聚层厚度为1～20nm，层内Ni、Mn、N、Si元素的含量分别是各元素在不锈钢中平均含量的1.2～3倍。
  2. 一种如权利要求1所述的不锈钢的加工方法，包括如下步骤：
    (a)将化学元素组成符合要求的原料进行固溶处理，冷却得到试样；原料的化学元素组成为：重量百分含量为0.01％～0.1％的C，0.05％～0.2％的N，不高于0.03％的P，不高于0.003％的S，0.5％～1％的Si，1.0％～2.0％的Mn，15％～17％的Cr，5％～7％的Ni，其余为Fe；
    (b)将试样在室温下进行一定程度的变形，变形过程采用多道次小变形逐步增加变形量的方式进行，每道次的截面收缩率增量为0.01～0.1，累计总截面收缩率应符合公式(1)
    1-exp{-β[1-exp(-αε)]n}<0.3  (1)
    ε是截面收缩率，
    α是与层错能相关的参数，
    β是与该试样材料马氏体相变化学驱动能相关参数，
    n是指前因子；
    (c)对经步骤(b)处理的试样进行退火处理，退火温度为50～550℃，退火时间为10min～100h，冷却得到不锈钢。
  3. 如权利要求2所述的不锈钢的加工方法，其特征在于：所述不锈钢的加工方法由步骤(a)～(c)组成。
  4. 如权利要求2或3所述的不锈钢的加工方法，其特征在于：步骤(a)中，固溶处理的温度为1050℃～1150℃，保温时间为1min～2h。
  5. 如权利要求2或3所述的不锈钢的加工方法，其特征在于：步骤(a)中，冷却方式为淬水或者淬油。
  6. 如权利要求2或3所述的不锈钢的加工方法，其特征在于：步骤(b)中，变形方式为轧制、冲压、锻造或拉拔。
  7. 如权利要求2或3所述的不锈钢的加工方法，其特征在于：步骤(c)中，加热方式为随炉升温，冷却方式为随炉冷却或空冷。
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