WO2018010494A1 - 一种过共析钢轨及其制备方法 - Google Patents

一种过共析钢轨及其制备方法 Download PDF

Info

Publication number
WO2018010494A1
WO2018010494A1 PCT/CN2017/085651 CN2017085651W WO2018010494A1 WO 2018010494 A1 WO2018010494 A1 WO 2018010494A1 CN 2017085651 W CN2017085651 W CN 2017085651W WO 2018010494 A1 WO2018010494 A1 WO 2018010494A1
Authority
WO
WIPO (PCT)
Prior art keywords
weight
rail
steel
temperature
rolling
Prior art date
Application number
PCT/CN2017/085651
Other languages
English (en)
French (fr)
Inventor
韩振宇
邹明
郭华
陶功明
汪渊
王春建
贾济海
Original Assignee
攀钢集团攀枝花钢铁研究院有限公司
Priority date (The priority date is an assumption and is not a legal conclusion. Google has not performed a legal analysis and makes no representation as to the accuracy of the date listed.)
Filing date
Publication date
Application filed by 攀钢集团攀枝花钢铁研究院有限公司 filed Critical 攀钢集团攀枝花钢铁研究院有限公司
Priority to AU2017295527A priority Critical patent/AU2017295527B2/en
Priority to US16/317,416 priority patent/US20190226040A1/en
Priority to BR112019000331-2A priority patent/BR112019000331A2/pt
Publication of WO2018010494A1 publication Critical patent/WO2018010494A1/zh

Links

Images

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
    • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
    • C21D9/04—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for rails
    • 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
    • C22C38/12—Ferrous alloys, e.g. steel alloys containing tungsten, tantalum, molybdenum, vanadium, or niobium
    • 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
    • C21D6/002—Heat treatment of ferrous alloys containing Cr
    • 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
    • C21D6/005—Heat treatment of ferrous alloys containing Mn
    • 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
    • C21D6/008—Heat treatment of ferrous alloys containing Si
    • 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
    • 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
    • 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
    • C22C38/14—Ferrous alloys, e.g. steel alloys containing titanium or zirconium
    • 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
    • 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
    • C22C38/24—Ferrous alloys, e.g. steel alloys containing chromium with vanadium
    • 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
    • C22C38/28—Ferrous alloys, e.g. steel alloys containing chromium with titanium or zirconium
    • 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
    • C21D2211/003—Cementite
    • 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
    • C21D2211/009—Pearlite

Definitions

  • the invention relates to the field of railways, and in particular to a hypereutectoid steel rail and a preparation method thereof.
  • Pearlitic rails are the most widely used and widely used products in the railway sector.
  • the existing carbon steel rails and microalloyed rails with over-eutectoid carbon content have a maximum strength of only 1300-1400 MPa even after heat treatment.
  • the rail head part is easy to form too fast wear, nuclear injury, peeling off the block and other damage, the service life of the rail is greatly reduced and the driving safety is endangered. Therefore, in the case of irreversible transportation conditions, it is very important to improve the toughness of the rail and thus extend the service life of the rail. For heavy-duty railway small radius curve sections, wear has gradually become the primary factor affecting life.
  • the most effective way to improve wear performance is to increase the hardness of the rail.
  • the most effective and cheapest way to increase the hardness is to increase the carbon content in the steel.
  • an increase in carbon content will result in a decrease in the toughness and plasticity index of the rail, and the existing ductile plasticity is close to the lower limit of the safe service of the rail. If this index is further reduced, the risk of brittle fracture of the rail under the same service conditions will increase.
  • the use of heat treatment technology relying on the strong fine grain strengthening effect, can improve the toughness and plasticity of the rail while improving the hardness of the rail. It has been proved that the heat-treated rail usually has higher toughness and comprehensive performance than the hot-rolled rail of the same composition.
  • the object of the present invention is to overcome the problem of poor toughness and comprehensive performance of the hypereutectoid steel rail in the prior art, and to provide a hypereutectoid steel rail and a preparation method of the hypereutectoid steel rail.
  • the inventors of the present invention found in the course of research that for high carbon content rails, refining the austenite grains and finally refining the pearlite layer spacing not only contributes to the strength, but also has a significant benefit to the improvement of the ductile plasticity index; The final effect of refining austenite grains on toughness and plasticity is far less than that of fine grain strengthening. However, for hypereutectoid steels with higher carbon content, if the ductile plasticity has reached the limit, it can be further refined. Austenitic grains are important for maintaining high strength of the rail while achieving higher toughness and plasticity.
  • CN102803536A High-carbon steel rail with excellent ductility and manufacturing method thereof
  • 1 low-temperature reheating after rail rolling but there is a problem that coarse carbides remain in the austenite grain inside, On the contrary, the toughness and plasticity of the pearlite structure after the accelerated cooling is lowered.
  • reheating also has economic problems such as high manufacturing cost and low production efficiency; 2 the precipitates pin the austenite grain boundaries and refine the austenite grains, thereby improving the toughness of the final rail;
  • the effect of the austenitic state on the properties of the final product is limited. However, when the carbon content exceeds 0.85%, the influence of the austenite state on the properties of the product cannot be ignored.
  • the inventors of the present invention also found in the research process that if it is desired to refine the austenite grains and finally obtain a rail product with excellent performance, it is necessary to uniformly control the rolling temperature and the finishing temperature to be finally prepared. Rails with excellent toughness and comprehensive performance under eutectoid conditions, if combined with the control of the chemical composition and cooling process of the billet, will produce better performance rails.
  • the present invention provides a method of preparing a hypereutectoid steel rail, wherein the method comprises rolling a steel slab containing V and Ti, wherein the steel slab contains 0.85-0.94 by weight based on the total weight of the steel slab % carbon, and start rolling temperature T start and a final finish rolling temperature T and the relationship between the content of vanadium [V] and a titanium content [Ti] satisfy the following formula:
  • T final 750 + b ([V] + 5 [Ti]),
  • [V] is 0.03-0.08 wt%
  • [Ti] is 0.011-0.02 wt%
  • [V]+5 [Ti] is 0.12-0.14 wt% based on the total weight of the billet.
  • the invention also provides a hypereutectoid steel rail prepared by the above method.
  • the hypereutectoid steel rail prepared by the method of the invention has excellent toughness comprehensive properties (such as toughness and contact fatigue resistance), in particular, the microstructure can be obtained as pearlite + trace secondary cementite. Eutectoid rails.
  • FIG. 1 is a view showing the microstructure of a hypereutectoid steel rail obtained in Example 1 of the present invention.
  • the present invention provides a method of preparing a hypereutectoid steel rail, wherein the method comprises rolling a steel slab containing V and Ti, wherein the steel slab contains 0.85-0.94% by weight based on the total weight of the steel slab carbon, and start rolling temperature T start and a final finish rolling temperature T and the relationship between the content of vanadium [V] and a titanium content [Ti] satisfy the following formula:
  • T final 750 + b ([V] + 5 [Ti]),
  • [V] is 0.03-0.08 wt%
  • [Ti] is 0.011-0.02 wt%
  • [V]+5 [Ti] is 0.12-0.14 wt% based on the total weight of the billet.
  • C is the most important and cheapest element for improving the hardness and pearlite transformation of pearlite rail.
  • the C content is less than 0.85%, even if the accelerated cooling after rolling is used, the high strength and corresponding force required for heavy-duty railway cannot be obtained. Wear resistance; when the C content is more than 0.94%, secondary cementite distributed along the grain boundary will still be precipitated at the prior austenite grain boundary, which deteriorates the impact toughness of the rail, and the contact fatigue resistance of the rail is greatly reduced. Therefore, the C content is limited to 0.85% to 0.94%.
  • V When V is at room temperature, the solubility in steel is very low, and in the hot rolling process, such as in the austenite grain boundary or other areas, it is precipitated in the form of fine grained vanadium carbonitride, or in steel. Ti composite precipitation, inhibiting the growth of austenite grains, thereby achieving the purpose of refining the grain to improve performance; when [V] is less than 0.03 wt%, the precipitation of V-containing carbonitride is limited, and it is difficult to exert a strengthening effect; When V] is more than 0.08% by weight, coarse carbonitrides are easily formed, and the ductility of the rail is deteriorated. Therefore, [V] is 0.03-0.08% by weight.
  • the main role of Ti in steel is to refine austenite grains during heating, rolling and cooling, and ultimately increase the elongation and stiffness of the rail, which is one of the important additive elements of the present invention.
  • [Ti] is less than 0.011% by weight, the amount of carbides formed in the rail is extremely limited; when [Ti] is more than 0.02% by weight, on the one hand, since Ti is a strong carbonitride forming element, the excessive TiC generated will cause The hardness of the rail is too high.
  • the TiC is more concentrated and enriched to form coarse carbides, which not only reduces the ductile plasticity, but also makes the contact surface of the rail easy to crack and cause fracture under the impact load. Therefore, [Ti] is from 0.011 to 0.02% by weight.
  • V, Ti in steel and C, N and other affinity and the formation of carbides in the number and form are significantly different, but formed
  • the role of carbonitrides in refining austenite grains is similar.
  • V in a steel with a lower N content, the V content of the solid solution in the ferrite matrix exceeds 50%, and in the steel with a higher N content, the V dissolved in the steel is more than 20%, and the remaining 70 % is precipitated as vanadium carbonitride.
  • the improvement of performance by adding V or Ti alone is not significant, such as adding 0.09% of V while the strength of the rail without adding Ti can still reach 1350 MPa, but the elongation is generally less than 10%, and Ti is added separately. Microalloying, the strength of the rail cannot meet the requirement of 1350 MPa, and therefore, the slab of the present invention contains both V and Ti, and [V] + 5 [Ti] is 0.12 - 0.14 wt%.
  • [V] is from 0.045 to 0.055% by weight
  • [Ti] is from 0.014 to 0.02% by weight
  • [V] + 5 [Ti] is from 0.12 to 0.13% by weight.
  • the steel slab of the above composition can be obtained by a conventional method in the art, for example, smelting molten steel containing the above components by a converter or an electric furnace, and being subjected to aluminum-free deoxidation, refining outside the furnace, vacuum degassing treatment, and continuous casting to 250 mm ⁇ 250 mm -
  • the 450mm ⁇ 450mm section billet is cooled and heated into the heating furnace.
  • the billet heating temperature is greater than 1200°C, and the steel billet is heated to no more than 3h to ensure that the billet cross-section billet temperature is evenly discharged, and the scale is removed to obtain the billet of the present invention.
  • the specific process I will not repeat them here.
  • the size, distribution and morphology of the precipitated phase are significantly different depending on the difference between the rolling temperature and the finishing temperature, and the different [V] and [Ti].
  • the rail performance obtained varied significantly, the inventors of the present invention found that, when the start rolling temperature T and the open end and the finishing temperature T [V] +5 [Ti] satisfy the following expression, the steel can be ensured
  • the fine and dispersed vanadium carbonitride and titanium carbonitride are fully dissolved and precipitated, so that the prepared hypereutectoid steel rail has excellent toughness and plasticity and contact fatigue resistance.
  • 780 ⁇ a ⁇ 800, 470 ⁇ b ⁇ 500 is preferable.
  • the steel slab may further contain Si: 0.4 to 0.9% by weight, Mn: 0.7 to 1.3% by weight, Cr: 0.2 to 0.6% by weight, P ⁇ 0.02% by weight, S ⁇ based on the total weight of the slab. 0.00% by weight, N: 0.06-0.09 by weight.
  • Si is present as a solid solution strengthening element in steel in the ferrite and austenite to increase the strength of the structure.
  • the invention preferably controls the content of Si within the above range, can enhance the solid solution effect, enhance the ductility of the rail, optimize the lateral performance of the rail, and is beneficial to improving the safety of use of the rail.
  • Mn can form a solid solution with Fe in the steel slab, and the strength of ferrite and austenite can be improved.
  • Mn is a carbide forming element. After entering the cementite, it can partially replace the Fe atom, increase the hardness of the carbide, and finally increase the hardness of the steel.
  • it is preferable to control the content of Mn within the above range, and it is possible to prevent the toughness and plasticity from being affected by the excessive hardness of the carbide in the steel while ensuring a high strengthening effect.
  • Cr and Fe in the slab can form a continuous solid solution and form a plurality of carbides with C, which is one of the main strengthening elements in the steel.
  • Cr can uniform the distribution of carbides in the steel and improve the wear properties of the steel.
  • the present invention preferably controls the content of Cr to the above range, which is advantageous for improving the ductility of the rail.
  • the N content it is preferred to control the N content to the above range, which is advantageous for improving the toughness and comprehensive properties of the rail under room temperature conditions.
  • the slab contains Si: 0.55-0.65 wt%, Mn: 1.25-1.3 wt%, Cr: 0.4-0.55 wt%, P ⁇ 0.014 wt%, S ⁇ based on the total weight of the slab. 0.005 wt%, N: 0.06-0.07 wt.
  • the method may further comprise: after the rolling, when the surface layer of the rail head surface temperature T decreases below the final finish rolling temperature T is lower than 20-50 deg.] C, rapid cooling of the rail head When the rapid cooling process causes the rail surface temperature T surface layer to be 450-550 ° C, the rapid cooling is stopped and air cooling is continued to room temperature.
  • the cooling rate of the rapid cooling is not particularly limited and may be a conventional choice in the art.
  • the cooling rate of the rapid cooling is preferably 2 to 5 ° C / s, preferably 4.5 to 4.9 ° C / s.
  • the embodiment of the rapid cooling is not particularly limited and may be a conventional choice in the art.
  • the rapid cooling embodiment may be to apply a cooling medium to the top surface and both sides of the rail head of the rail. .
  • the selection of the cooling medium is not particularly limited and may be a conventional choice in the art as long as the cooling purpose of the present invention can be attained.
  • the cooling medium is a compressed air and/or water mist mixture.
  • the present invention also provides a hypereutectoid steel rail prepared by the above method. It should be understood that the hypereutectoid steel rail provided by the present invention has the same composition as the steel slab described above.
  • the hypereutectoid steel rail of the invention has excellent toughness and plasticity and contact fatigue resistance.
  • room temperature means “25 ° C”.
  • the slabs used in the following examples and comparative examples contained chemical compositions as shown in Table 1, in which, with the exception of the elements in Table 1, the balance was Fe and unavoidable impurities:
  • the 1# to 6# billets in Table 1 were rolled into 60 kg/m rails according to the rolling and rapid cooling processes shown in the corresponding numbers in Table 2, and rapidly cooled by compressed air, and then the above-mentioned finished rails were air-cooled to At room temperature, hypereutectoid rails A1 to A6 were obtained.
  • the 1# to 6# billets in Table 1 were rolled into 60 kg/m steel rails according to the rolling and rapid cooling processes shown in the corresponding numbers in Table 3, and rapidly cooled by compressed air, and then the above-mentioned finished rails were air-cooled to At room temperature, hypereutectoid rails D1 to D6 were obtained.
  • microstructure Inspection Method the microstructure of the hypereutectoid rail was measured by MeF3 optical microscope. The microstructure was measured as shown in Table 4. The microstructure of A1 is shown in Figure 1. Show.
  • P+Fe 3 C II (micro) refers to pearlite + trace secondary cementite
  • P+Fe 3 C II (less) refers to pearlite + small amount of secondary cementite
  • trace and “small amount” “It is a relative relationship, and the "trace” indicates less amount relative to "small amount”, mainly to show that the secondary cementite in the microstructure of the hypereutectoid rail in the embodiment is less than that in the comparative example.

Landscapes

  • Chemical & Material Sciences (AREA)
  • Engineering & Computer Science (AREA)
  • Materials Engineering (AREA)
  • Mechanical Engineering (AREA)
  • Metallurgy (AREA)
  • Organic Chemistry (AREA)
  • Physics & Mathematics (AREA)
  • Thermal Sciences (AREA)
  • Crystallography & Structural Chemistry (AREA)
  • Heat Treatment Of Steel (AREA)
  • Heat Treatment Of Articles (AREA)
  • Metal Rolling (AREA)

Abstract

一种过共析钢轨及其制备方法，包括将含有V和Ti的钢坯进行轧制，其中，以钢坯的总重量为基准，钢坯含有0.85-0.94重量％的碳，开轧温度T 开和终轧温度T 终与钒含量[V]和钛含量[Ti]的关系满足以下公式：T 开＝1100+a([V]+5[Ti])，T 终＝750+b([V]+5[Ti])，其中，500≤a≤800，300≤b≤500；以钢坯的总重量为基准，[V]为0.03-0.08重量％，[Ti]为0.011-0.02重量％，且[V]+5[Ti]为0.12-0.14重量％。该方法制备的过共析钢轨具有优良的强韧综合性能。

Description

一种过共析钢轨及其制备方法 技术领域
本发明涉及铁路领域，具体地，涉及一种过共析钢轨及其制备方法。
背景技术
珠光体钢轨是铁路领域用量最大、使用范围最广的产品。随着铁路轴重及行车密度的不断增加，现有过共析碳含量的普通碳素钢轨及微合金化钢轨即使经过热处理后，最高强度仅为1300-1400MPa，在轮轨间多向复杂重载荷作用下，轨头部位易形成过快磨耗、核伤、剥离掉块等伤损，钢轨的服役寿命大幅降低并危及行车安全。因此，在苛刻的运输条件不可逆转的情况下，提高钢轨的强韧性能，进而延长钢轨的服役寿命就显得十分重要。对于重载铁路小半径曲线路段而言，磨耗已逐步成为影响寿命的首要因素。研究证实，提高磨耗性能最有效的方式是提高钢轨的硬度。现有珠光体钢轨条件下，提高硬度最有效、最廉价的方式是提高钢中的碳含量。然而，碳含量的提高将导致钢轨的韧塑性指标降低，而现有的韧塑性已接近钢轨安全服役的下限，如果该指标进一步降低将增加钢轨在同等服役条件下发生脆断的风险。采用热处理技术，依托强烈的细晶强化效应，可在提高钢轨强硬度的同时提高钢轨的韧塑性，实践证明，热处理钢轨通常拥有较相同成分热轧钢轨更高的强韧综合性能。即使如此，如果碳含量进一步提高，例如达到0.85％以上，即使热处理技术对韧塑性的提升幅度也是有限的，需要通过新技术的应用来满足过共析碳含量条件下钢轨强韧性能的需要。
发明内容
本发明的目的是克服现有技术中过共析钢轨的强韧综合性能不佳的问题，提供一种过共析钢轨，以及该过共析钢轨的制备方法。
本发明的发明人在研究过程中发现，对于高碳含量钢轨，细化奥氏体晶粒并最终细化珠光体片层间距不仅对强度有贡献，对韧塑性指标的提高也有显著益处；尽管细化奥氏体晶粒最终对韧塑性的提升效果远不及细晶强化的效果，但对于碳含量更高的过共析钢轨，当韧塑性已趋于极限的条件下，如能进一步细化奥氏体晶粒，则对钢轨保持高强度同时获得更高的韧塑性是有重要意义的。对于奥氏体晶粒的细化， CN102803536A《延展性优良的珠光体系高碳钢轨及其制造方法》中描述了如下方法：①钢轨轧制后进行低温再加热，但存在的问题是奥氏体晶粒内部融化残留有粗大碳化物，反而使加速冷却后的珠光体组织的韧塑性降低。同时，再加热还存在制造成本高、生产效率低等经济性问题；②析出物钉扎奥氏体晶界，细化奥氏体晶粒，从而改善最终钢轨的强韧性；研究表明，对于过共析碳含量的钢轨，奥氏体状态对最终产品性能的影响是有限的。然而，当碳含量超过0.85％时，奥氏体状态对产品的性能的影响不容忽视。
本发明的发明人在研究过程中还发现，若要达到细化奥氏体晶粒并最终获得性能优良的钢轨产品，需对开轧温度、终轧温度进行统一协调控制，才能最终制备出过共析条件下强韧综合性能优良的钢轨，如果同时结合对钢坯的化学成分和冷却工艺的控制，则会制备出性能更优的钢轨。
由此，本发明提供一种制备过共析钢轨的方法，其中，该方法包括将含有V和Ti的钢坯进行轧制，其中，以钢坯的总重量为基准，所述钢坯含有0.85-0.94重量％的碳，开轧温度T开和终轧温度T终与钒含量[V]和钛含量[Ti]的关系满足以下公式：
T开＝1100+a([V]+5[Ti])，
T终＝750+b([V]+5[Ti])，
其中，500≤a≤800，300≤b≤500；
以钢坯的总重量为基准，[V]为0.03-0.08重量％，[Ti]为0.011-0.02重量％，且[V]+5[Ti]为0.12-0.14重量％。
本发明还提供了上述方法制备得到的过共析钢轨。
采用本发明所述方法制备的过共析钢轨具有优良的强韧综合性能(例如韧塑性和耐接触疲劳性能)，特别是可以得到显微组织结构为珠光体+微量二次渗碳体的过共析钢轨。
本发明的其它特征和优点将在随后的具体实施方式部分予以详细说明。
附图说明
附图是用来提供对本发明的进一步理解，并且构成说明书的一部分，与下面的具体实施方式一起用于解释本发明，但并不构成对本发明的限制。在附图中：
图1是本发明实施例1得到的过共析钢轨的显微组织图。
具体实施方式
以下对本发明的具体实施方式进行详细说明。应当理解的是，此处所描述的具体 实施方式仅用于说明和解释本发明，并不用于限制本发明。
在本文中所披露的范围的端点和任何值都不限于该精确的范围或值，这些范围或值应当理解为包含接近这些范围或值的值。对于数值范围来说，各个范围的端点值之间、各个范围的端点值和单独的点值之间，以及单独的点值之间可以彼此组合而得到一个或多个新的数值范围，这些数值范围应被视为在本文中具体公开。
本发明提供了一种制备过共析钢轨的方法，其中，该方法包括将含有V和Ti的钢坯进行轧制，其中，以钢坯的总重量为基准，所述钢坯含有0.85-0.94重量％的碳，开轧温度T开和终轧温度T终与钒含量[V]和钛含量[Ti]的关系满足以下公式：
T开＝1100+a([V]+5[Ti])，
T终＝750+b([V]+5[Ti])，
其中，500≤a≤800，300≤b≤500；
以钢坯的总重量为基准，[V]为0.03-0.08重量％，[Ti]为0.011-0.02重量％，且[V]+5[Ti]为0.12-0.14重量％。
C是珠光体钢轨提高强硬度、促进珠光体转变的最重要、最廉价的元素，当C含量小于0.85％时，即使采用轧后加速冷却，也无法获得重载铁路所需的高强度及相应的耐磨损性能；当C含量大于0.94％时，在原奥氏体晶界处仍将析出沿晶界分布的二次渗碳体，恶化钢轨的冲击韧性，钢轨耐接触疲劳性能大幅降低。因此，C含量限定在0.85％-0.94％。
V处于室温条件下时，在钢中的溶解度很低，而在热轧过程中如存在于奥氏体晶界或其它区域，以细化颗粒状的钒碳氮化物形式析出，或与钢中的Ti复合析出，抑制奥氏体晶粒的生长，从而达到细化晶粒提高性能的目的；当[V]小于0.03重量％时，含V碳氮化物析出有限，难以发挥强化效果；当[V]大于0.08重量％时，易形成粗大的碳氮化物，恶化钢轨的韧塑性。因此，[V]为0.03-0.08重量％。
Ti在钢中的主要作用是细化加热、轧制及冷却时的奥氏体晶粒，最终增加钢轨的延伸率和刚度，是本发明重要的添加元素之一。当[Ti]小于0.011重量％时，在钢轨中形成的碳化物数量极为有限；当[Ti]大于0.02重量％时，一方面由于Ti是强碳氮化物形成元素，产生的TiC偏多将使钢轨硬度过高，另一方面TiC偏多将偏聚富集形成粗大碳化物，不仅降低韧塑性，还使得钢轨在冲击载荷作用下接触面易于开裂并导致断裂。因此，[Ti]为0.011-0.02重量％。
V、Ti在钢中与C、N等亲和力以及形成碳化物的数量和形态有明显区别，但形成 碳氮化物细化奥氏体晶粒的作用是相似的。以V为例，N含量较低的钢中，固溶于铁素体基体的V含量超过50％，而N含量较高的钢中，固溶于钢中的V为20％多，剩余70％均以钒碳氮化物的形式析出。在本发明中，单独添加V或Ti对性能的改善均不显著，如添加0.09％的V的同时不添加Ti的钢轨强度仍能够达1350MPa，但延伸率一般低于10％，而单独添加Ti微合金化，钢轨的强度无法达到1350MPa要求，因此，本发明所述钢坯中同时含有V和Ti，并且，[V]+5[Ti]为0.12-0.14重量％。
优选的情况下，[V]为0.045-0.055重量％，[Ti]为0.014-0.02重量％，且[V]+5[Ti]为0.12-0.13重量％。
根据本发明，上述组成的钢坯可以通过本领域的常规方法获得，例如采用转炉或电炉冶炼含上述成分的钢水，经无铝脱氧、炉外精炼、真空脱气处理，连铸为250mm×250mm-450mm×450mm断面钢坯后冷却进入加热炉中加热，钢坯加热温度大于1200℃，保温至不超过3h确保钢坯断面坯料温度均匀后出炉，去除氧化铁皮，便可获得本发明所述钢坯，具体的过程在此不再赘述。
根据本发明，在钢坯轧制为钢轨的过程中，随着开轧温度和终轧温度的不同，结合不同的[V]和[Ti]，析出相的尺寸、分布及形貌差异明显，进而在获得的钢轨性能也存在明显差异，经本发明的发明人研究发现，当开轧温度T开和终轧温度T终与[V]+5[Ti]的关系满足上式时，可以确保钢中细小、弥散分布的钒碳氮化物和钛碳氮化物充分溶解和析出，从而使制备出的过共析钢轨具有优良的韧塑性及耐接触疲劳性能。为了确保析出相对韧塑性及耐疲劳性能的贡献最大化，优选情况下，在上式中，780≤a≤800，470≤b≤500。
根据本发明，以钢坯的总重量为基准，所述钢坯还可以含有Si：0.4-0.9重量％，Mn：0.7-1.3重量％，Cr：0.2-0.6重量％，P≤0.02重量％，S≤0.008重量％，N：0.06-0.09重量‰。
根据本发明，Si作为钢中的固溶强化元素存在于铁素体和奥氏体中提高组织的强度。本发明优选将Si的含量控制在上述范围，可以强化固溶效果，增强钢轨的韧塑性，优化钢轨的横向性能，有利于提高钢轨的使用安全性。
根据本发明，Mn可以和钢坯中的Fe形成固溶体，提高铁素体和奥氏体的强度。同时，Mn又是碳化物形成元素，进入渗碳体后可部分替代Fe原子，增加碳化物的硬度，最终增加钢的硬度。本发明优选将Mn的含量控制在上述范围，可以在确保较高强化效果的同时，避免因为钢中碳化物硬度过高而影响韧塑性。
根据本发明，Cr与钢坯中的Fe能形成连续固溶体，并与C形成多种碳化物，是钢中的主要强化元素之一。此外，Cr能均匀钢中碳化物的分布，提高钢材的磨损性能。本发明优选将Cr的含量控制在上述范围，有利于提高钢轨的韧塑性。
根据本发明，优选将N含量控制在上述范围，有利于提高室温条件下钢轨的强韧综合性能。
进一步优选的情况下，以钢坯的总重量为基准，所述钢坯含有Si：0.55-0.65重量％，Mn：1.25-1.3重量％，Cr：0.4-0.55重量％，P≤0.014重量％，S≤0.005重量％，N：0.06-0.07重量‰。
根据本发明的一种优选实施方式，所述方法还可以包括：在轧制之后，当轨头表层温度T表层降低至比终轧温度T终低20-50℃时，对轨头进行快速冷却；当所述快速冷却的过程使得轨头表层温度T表层为450-550℃时，停止快速冷却并继续空冷至室温。
根据本发明，控制快速冷却开始时的轨头表层温度T表层和快速冷却停止时轨头表层温度在上述范围内，可以防止钢轨中可能会形成沿晶界分布的二次渗碳体，有利于提高钢轨的韧塑性。
根据本发明，对所述快速冷却的冷却速度没有特别的限定，可以为本领域常规的选择。但是，为了减少二次渗碳体析出，提高钢轨的强度指标和耐磨损性能，优选情况下，所述快速冷却的冷却速度为2-5℃/s，优选为4.5-4.9℃/s。
根据本发明，对所述快速冷却的实施方式没有特别的限定，可以为本领域的常规选择，例如所述快速冷却的实施方式可以是在所述钢轨的轨头顶面和两个侧面施加冷却介质。
根据本发明，对所述冷却介质的选择没有特别的限定，可以为本领域的常规选择，只要能达到本发明冷却目的即可。优选情况下，所述冷却介质为压缩空气和/或水雾混合液。
本发明还提供了由上述方法制备得到的过共析钢轨。应当理解的是，本发明提供的过共析钢轨具有与上述钢坯组成一样的组成。本发明所述过共析钢轨具有优良的韧塑性和耐接触疲劳性能。
以下将通过实施例对本发明进行详细描述。以下实施例中，“室温”是指“25℃”。
以下实施例和对比例所采用的钢坯含有的化学成分如表1中所示，其中，除了表1中的元素以外，余量为Fe和不可避免的杂质：
表1
Figure PCTCN2017085651-appb-000001
实施例1-6
按照表2对应编号所示的轧制和快速冷却工艺将表1中的1#至6#钢坯轧制成60kg/m的钢轨，并用压缩空气进行快速冷却，然后将上述完成处理的钢轨空冷至室温，得到过共析钢轨A1至A6。
表2
Figure PCTCN2017085651-appb-000002
对比例1-6
按照表3对应编号所示的轧制和快速冷却工艺将表1中的1#至6#钢坯轧制成60kg/m的钢轨，并用压缩空气进行快速冷却，然后将上述完成处理的钢轨空冷至室温，得到过共析钢轨D1至D6。
表3
Figure PCTCN2017085651-appb-000003
测试例1
根据以下方法对实施例1-11和对比例1-6制备的过共析钢轨A1-A11和D1-D6进行性能检测，具体地：
按GB/T228.1-2010《金属材料室温拉伸试验方法》测定过共析钢轨的拉伸性能，测得Rm(抗拉强度)、A％(伸长率)，结果如表4所示；
按本领域的常规方法测定过共析钢轨的轨头常温U型冲击性能，结果如表4所示；
按GB/T 13298-1991《金属显微组织检验方法》采用MeF3光学显微镜测定过共析钢轨的显微组织，测得显微组织结果如表4所示，A1的显微组织如图1所示。
表4
Figure PCTCN2017085651-appb-000004
注：P+Fe3CⅡ(微)是指珠光体+微量二次渗碳体；P+Fe3CⅡ(少)是指珠光体+少量二次渗碳体；“微量”和“少量”是相对的关系，“微量”相对于“少量”表示的量更少，主要为了表明实施例中过共析钢轨显微组织中的二次渗碳体比对比例中的更少。
从实施例1-6和对比例1-6的结果比较可看出，显微组织检验结果表明，所有钢轨组织均为珠光体+二次渗碳体(Fe3CⅡ)，不同的是，采用本发明成分及相应工艺条件下沿晶界析出的二次渗碳体数量更少，更有利于提高钢轨的韧塑性及服役安全性。从性能指标来看，在实施例与对比例钢轨抗拉强度相近的情况下，本发明所述过共析钢轨的延长率与轨头常温U型冲击性能大幅提高，更有利于提高钢轨的耐接触疲劳性能。
以上详细描述了本发明的优选实施方式，但是，本发明并不限于上述实施方式中 的具体细节，在本发明的技术构思范围内，可以对本发明的技术方案进行多种简单变型，这些简单变型均属于本发明的保护范围。
另外需要说明的是，在上述具体实施方式中所描述的各个具体技术特征，在不矛盾的情况下，可以通过任何合适的方式进行组合，为了避免不必要的重复，本发明对各种可能的组合方式不再另行说明。
此外，本发明的各种不同的实施方式之间也可以进行任意组合，只要其不违背本发明的思想，其同样应当视为本发明所公开的内容。

Claims (10)

  1. 一种制备过共析钢轨的方法，其特征在于，该方法包括将含有V和Ti的钢坯进行轧制，其中，以钢坯的总重量为基准，所述钢坯含有0.85-0.94重量％的碳，开轧温度T开和终轧温度T终与钒含量[V]和钛含量[Ti]的关系满足以下公式：
    T开＝1100+a([V]+5[Ti])，
    T终＝750+b([V]+5[Ti])，
    其中，500≤a≤800，300≤b≤500；
    以钢坯的总重量为基准，[V]为0.03-0.08重量％，[Ti]为0.011-0.02重量％，且[V]+5[Ti]为0.12-0.14重量％。
  2. 根据权利要求1所述的方法，其中，[V]为0.045-0.055重量％，[Ti]为0.014-0.02重量％，且[V]+5[Ti]为0.12-0.13重量％。
  3. 根据权利要求1或2所述的方法，其中，以钢坯的总重量为基准，所述钢坯还含有Si：0.4-0.9重量％，Mn：0.7-1.3重量％，Cr：0.2-0.6重量％，P≤0.02重量％，S≤0.008重量％，N：0.06-0.09重量‰。
  4. 根据权利要求3所述的方法，其中，以钢坯的总重量为基准，所述钢坯含有Si：0.55-0.65重量％，Mn：1.25-1.3重量％，Cr：0.4-0.55重量％，P≤0.014重量％，S≤0.005重量％，N：0.06-0.07重量‰。
  5. 根据权利要求1或2所述的方法，其中，所述方法还包括：在轧制之后，当轨头表层温度T表层降低至比终轧温度T终低20-50℃时，对轨头进行快速冷却；当所述快速冷却的过程使得轨头表层温度T表层为450-550℃时，停止快速冷却并继续空冷至室温。
  6. 根据权利要求5所述的方法，其中，所述快速冷却的冷却速度为2-5℃/s。
  7. 根据权利要求6所述的方法，其中，所述快速冷却的冷却速度为4.5-4.9℃/s。
  8. 根据权利要求5所述的方法，其中，所述快速冷却的实施方式是在所述钢轨的轨头顶面和两个侧面施加冷却介质。
  9. 根据权利要求8所述的方法，其中，所述冷却介质为压缩空气和/或水 雾混合液。
  10. 由权利要求1-9中任意一项所述的方法制备得到的过共析钢轨。
PCT/CN2017/085651 2016-07-14 2017-05-24 一种过共析钢轨及其制备方法 WO2018010494A1 (zh)

Priority Applications (3)

Application Number Priority Date Filing Date Title
AU2017295527A AU2017295527B2 (en) 2016-07-14 2017-05-24 A hypereutectoid rail and manufacturing method thereof
US16/317,416 US20190226040A1 (en) 2016-07-14 2017-05-24 Hypereutectoid rail and manufacturing method thereof
BR112019000331-2A BR112019000331A2 (pt) 2016-07-14 2017-05-24 trilho hipereutectoide e método de produção do mesmo

Applications Claiming Priority (2)

Application Number Priority Date Filing Date Title
CN201610552631.3 2016-07-14
CN201610552631.3A CN106086663B (zh) 2016-07-14 2016-07-14 一种过共析钢轨及其制备方法

Publications (1)

Publication Number Publication Date
WO2018010494A1 true WO2018010494A1 (zh) 2018-01-18

Family

ID=57220113

Family Applications (1)

Application Number Title Priority Date Filing Date
PCT/CN2017/085651 WO2018010494A1 (zh) 2016-07-14 2017-05-24 一种过共析钢轨及其制备方法

Country Status (5)

Country Link
US (1) US20190226040A1 (zh)
CN (1) CN106086663B (zh)
AU (1) AU2017295527B2 (zh)
BR (1) BR112019000331A2 (zh)
WO (1) WO2018010494A1 (zh)

Families Citing this family (8)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
CN106086663B (zh) * 2016-07-14 2018-03-06 攀钢集团攀枝花钢铁研究院有限公司 一种过共析钢轨及其制备方法
CN107739806B (zh) * 2017-10-10 2019-10-11 攀钢集团研究院有限公司 高韧塑性过共析钢轨及其制造方法
CN107675084B (zh) * 2017-10-10 2019-05-10 攀钢集团研究院有限公司 高碳高强韧性珠光体钢轨及其制造方法
CN107739983A (zh) * 2017-10-30 2018-02-27 攀钢集团攀枝花钢铁研究院有限公司 一种过共析钢轨及其生产方法
CN107779768A (zh) * 2017-10-31 2018-03-09 攀钢集团攀枝花钢铁研究院有限公司 用于生产耐腐蚀高速铁路用钢轨的方法
CN107747040A (zh) * 2017-10-31 2018-03-02 攀钢集团攀枝花钢铁研究院有限公司 耐腐蚀高速铁路钢轨制备方法
CN112718855A (zh) * 2020-11-30 2021-04-30 攀钢集团攀枝花钢铁研究院有限公司 一种过共析钢轨及其制备方法
CN115094338B (zh) * 2022-07-27 2023-09-22 内蒙古科技大学 一种过共析钢轨用钢及其制备方法

Citations (5)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
JP2005171327A (ja) * 2003-12-11 2005-06-30 Nippon Steel Corp 耐表面損傷性および耐内部疲労損傷性に優れたパーライト系レールの製造方法およびレール
CN102220545A (zh) * 2010-04-16 2011-10-19 攀钢集团有限公司 耐磨性和塑性优良的高碳高强热处理钢轨及其制造方法
CN104195433A (zh) * 2014-09-02 2014-12-10 攀钢集团攀枝花钢铁研究院有限公司 一种高强韧性珠光体钢轨及其生产方法
CN105018705A (zh) * 2015-08-11 2015-11-04 攀钢集团攀枝花钢铁研究院有限公司 一种过共析钢轨及其制备方法
CN106086663A (zh) * 2016-07-14 2016-11-09 攀钢集团攀枝花钢铁研究院有限公司 一种过共析钢轨及其制备方法

Family Cites Families (2)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
JP4469248B2 (ja) * 2004-03-09 2010-05-26 新日本製鐵株式会社 耐摩耗性および延性に優れた高炭素鋼レールの製造方法
CN104060075B (zh) * 2014-07-14 2016-05-04 攀钢集团攀枝花钢铁研究院有限公司 提高钢轨硬化层深度的热处理方法

Patent Citations (5)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
JP2005171327A (ja) * 2003-12-11 2005-06-30 Nippon Steel Corp 耐表面損傷性および耐内部疲労損傷性に優れたパーライト系レールの製造方法およびレール
CN102220545A (zh) * 2010-04-16 2011-10-19 攀钢集团有限公司 耐磨性和塑性优良的高碳高强热处理钢轨及其制造方法
CN104195433A (zh) * 2014-09-02 2014-12-10 攀钢集团攀枝花钢铁研究院有限公司 一种高强韧性珠光体钢轨及其生产方法
CN105018705A (zh) * 2015-08-11 2015-11-04 攀钢集团攀枝花钢铁研究院有限公司 一种过共析钢轨及其制备方法
CN106086663A (zh) * 2016-07-14 2016-11-09 攀钢集团攀枝花钢铁研究院有限公司 一种过共析钢轨及其制备方法

Also Published As

Publication number Publication date
US20190226040A1 (en) 2019-07-25
AU2017295527A1 (en) 2019-01-31
CN106086663B (zh) 2018-03-06
CN106086663A (zh) 2016-11-09
AU2017295527B2 (en) 2020-05-21
BR112019000331A2 (pt) 2019-04-16

Similar Documents

Publication Publication Date Title
WO2018010494A1 (zh) 一种过共析钢轨及其制备方法
AU2015353251B2 (en) Low-alloy high-strength high-tenacity steel panel and method for manufacturing same
CN106011648B (zh) 一种齿轮钢及其生产方法
JP2016534230A (ja) 高硬度熱間圧延鋼材製品及びその製造方法
CN109722611A (zh) 一种低屈强比超高强度连续油管用钢及其制造方法
WO2016045264A1 (zh) 一种高成形性的冷轧超高强度钢板、钢带及其制造方法
CN110527910B (zh) 含稀土Ce及Nb元素的耐腐蚀纳米珠光体钢轨制备方法
CN104046907B (zh) 一种屈服强度≥960MPa精轧螺纹钢筋及生产方法
CN107130172B (zh) 布氏硬度400hbw级整体硬化型高韧性易焊接特厚耐磨钢板及其制造方法
JP6152375B2 (ja) 低温靭性及び硫化水素応力腐食割れ抵抗性に優れた圧力容器用鋼材、その製造方法及び深絞り製品の製造方法
JP2021507091A (ja) 少なくとも１００ｍｍの厚さを有する鋼セクション及びその製造方法
AU2015205937A1 (en) Pearlitic steel rail with high strength and toughness and producing method thereof
CN111850399B (zh) 具有良好耐磨性耐蚀塑料模具钢及其制备方法
CN114807772B (zh) 一种时效强化的高强韧轻质钢及其制造方法
CN114517254B (zh) 一种船舶用耐低温球扁钢及其制备方法
JP2011504549A (ja) 深絞り用高張力鋼及びその製造方法
JP2024060017A (ja) 超高強度ばね用線材、鋼線及びその製造方法
JP6600412B2 (ja) 強度及び冷間加工性に優れた非調質線材及びその製造方法
US10113219B2 (en) Nano-pearlite rail and process for manufacturing same
CN107988553B (zh) 一种高韧性高速车轴用钢及其热处理方法
JP4220830B2 (ja) 靭性および延性に優れたパーライト系レールおよびその製造方法
JP7197707B2 (ja) 耐磨耗性及び高温強度に優れた車両のブレーキディスク用鋼材及びその製造方法
JP2001003140A (ja) 靭性および延性に優れた高強度パーライト系レールおよびその製造方法
CN110578041B (zh) 添加稀土Ce及Nb元素的耐腐蚀超细珠光体材料
JP2002363696A (ja) 靭性および延性に優れたパーライト系レールおよびその製造方法

Legal Events

Date Code Title Description
121 Ep: the epo has been informed by wipo that ep was designated in this application

Ref document number: 17826833

Country of ref document: EP

Kind code of ref document: A1

NENP Non-entry into the national phase

Ref country code: DE

REG Reference to national code

Ref country code: BR

Ref legal event code: B01A

Ref document number: 112019000331

Country of ref document: BR

ENP Entry into the national phase

Ref document number: 2017295527

Country of ref document: AU

Date of ref document: 20170524

Kind code of ref document: A

ENP Entry into the national phase

Ref document number: 112019000331

Country of ref document: BR

Kind code of ref document: A2

Effective date: 20190108

122 Ep: pct application non-entry in european phase

Ref document number: 17826833

Country of ref document: EP

Kind code of ref document: A1