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WO2024260333A1 - 一种高强韧高淬透性齿轴用钢及其制造方法 - Google Patents

一种高强韧高淬透性齿轴用钢及其制造方法 Download PDF

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Publication number
WO2024260333A1
WO2024260333A1 PCT/CN2024/099794 CN2024099794W WO2024260333A1 WO 2024260333 A1 WO2024260333 A1 WO 2024260333A1 CN 2024099794 W CN2024099794 W CN 2024099794W WO 2024260333 A1 WO2024260333 A1 WO 2024260333A1
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steel
strength
hardenability
toughness
temperature
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French (fr)
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赵四新
高加强
黄宗泽
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Baoshan Iron and Steel Co Ltd
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Baoshan Iron and Steel Co Ltd
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    • 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
    • C22C38/32—Ferrous alloys, e.g. steel alloys containing chromium with boron
    • 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
    • C21D7/00—Modifying the physical properties of iron or steel by deformation
    • C21D7/13—Modifying the physical properties of iron or steel by deformation by hot working
    • 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
    • C21D8/06—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of rods or wires
    • 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
    • C22C33/06—Making ferrous alloys by melting using master 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
    • C22C38/02—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
    • C22C38/12—Ferrous alloys, e.g. steel alloys containing tungsten, tantalum, molybdenum, vanadium, or niobium
    • 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
    • C22C38/20—Ferrous alloys, e.g. steel alloys containing chromium with copper
    • 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
    • C22C38/22—Ferrous alloys, e.g. steel alloys containing chromium with molybdenum or tungsten
    • 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
    • C22C38/26—Ferrous alloys, e.g. steel alloys containing chromium with niobium or tantalum
    • 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
    • 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
    • C22C38/42—Ferrous alloys, e.g. steel alloys containing chromium with nickel with copper
    • 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
    • C22C38/44—Ferrous alloys, e.g. steel alloys containing chromium with nickel with molybdenum or tungsten
    • 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
    • C22C38/46—Ferrous alloys, e.g. steel alloys containing chromium with nickel 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
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/48—Ferrous alloys, e.g. steel alloys containing chromium with nickel with niobium or tantalum
    • 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
    • C22C38/50—Ferrous alloys, e.g. steel alloys containing chromium with nickel with 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
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/54—Ferrous alloys, e.g. steel alloys containing chromium with nickel with boron

Definitions

  • the present invention relates to a high-strength steel material and a manufacturing method thereof, and in particular to a steel material for a gear shaft and a manufacturing method thereof.
  • high-strength and tough gear shaft steel can meet the high technical requirements of automobile lightweight materials.
  • the main technical difficulty of high hardenability MnCr carburized gear steel is how to improve the hardenability while avoiding the gear size dispersion caused by excessive hardenability bandwidth, and at the same time ensuring that the gear does not have mixed crystals and coarse grains after high-temperature carburizing.
  • a Chinese patent document with the publication number CN101096742A, the publication date of January 2, 2008, and the name "High Strength Automotive Gear Steel” discloses a high-strength automotive gear steel.
  • the steel contains alloy elements such as Nb, V, and Al, which refine the original austenite grains.
  • the mass percentage of the components is: C: 0.20-0.40, Si: 0.20-0.50, Mn: 0.50-1.00, Cr: 0.80-1.30, Nb: 0.015-0.080, V: 0.030-0.090, Mo: 0.15-0.55, Al: 0.015-0.050, and the rest is Fe and unavoidable impurities.
  • the grain size, hardenability and bandwidth of the gear steel are optimized.
  • Nb Ti composite microalloyed high temperature carburizing gear steel
  • the steel composition is: C: 0.17-0.22%, Si: 0.20-0.35%, Mn: 0.9-1.10%, P: ⁇ 0.025%, S: 0.020-0.035%, Cr: 1.05-1.30%, Al: 0.015-0.035%, Ti: 0.02-0.06%, Nb: 0.02-0.06%, the balance is iron and unavoidable impurities. It increases the gear carburizing temperature or shortens the carburizing time by controlling the content of microalloying elements such as Nb, Ti and Al.
  • One of the purposes of the present invention is to provide a high-strength, tough and high-hardenability gear shaft steel, which can obtain high-strength gear shaft steel with high hardenability, narrow hardenability bandwidth and good high-temperature grain stability at a low cost by optimizing the component system of the gear shaft steel, especially reasonably controlling the content of microalloy elements and nitrogen in the gear steel.
  • the present invention provides a high-strength, high-toughness and high-hardenability steel for gear shafts, which contains the following chemical elements in the following mass percentages in addition to Fe and inevitable impurities:
  • the mass percentage of each chemical element is:
  • the balance is Fe and unavoidable impurities
  • C In the high-strength, toughness and high hardenability gear shaft steel described in the present invention, C is a necessary component in the steel, and it is also one of the most important elements affecting the hardenability of the steel.
  • High-strength, toughness and high hardenability gear shaft steel requires surface strength as well as sufficient core impact toughness.
  • the C content in the steel is too low, less than 0.16%, the strength of the steel is insufficient and good hardenability requirements cannot be guaranteed; accordingly, the C content in the steel should not be too high.
  • the C content in the steel is too high, the toughness requirements of the gear core cannot be met, and too high a C content is detrimental to the plasticity of the steel, especially for carburized gear steel with a high Mn content.
  • the mass percentage of C is controlled between 0.16 and 0.22%.
  • the Si element can not only better eliminate the adverse effects of iron oxide on steel, but also dissolve into ferrite to strengthen ferrite, thereby improving the strength, hardness, wear resistance, elasticity and elastic limit of steel.
  • the Si element will increase the Ac 3 temperature of steel, and due to poor thermal conductivity, there is a risk of cracking and a tendency to decarburization.
  • the mass percentage of Si is controlled between 0.10 and 0.40%.
  • Mn In the high-strength, toughness and high-hardenability steel for gear shafts described in the present invention, Mn is one of the main elements that affect the hardenability of steel.
  • the Mn element has a very good deoxidation ability, which can reduce the iron oxide in the steel and can effectively increase the output of the steel.
  • Mn can dissolve into ferrite, improve the strength and hardness of the steel, and allow the steel to obtain pearlite with finer layers and higher strength when it is cooled after hot rolling.
  • Mn can also form MnS with S in the steel, which can eliminate the harmful effects of S. It has the ability to form and stabilize austenite structure in steel, which can strongly increase the hardenability of steel and reduce the high-temperature toughness of steel.
  • the mass percentage of Mn is controlled between 0.86 and 1.24%.
  • Cr is one of the main alloying elements added to the steel of the present invention. Cr can significantly improve the hardenability, strength, wear resistance and other properties of the steel. In addition, Cr can also reduce the activity of the C element in the steel, which can prevent decarburization during heating, rolling and heat treatment. However, too high Cr will significantly reduce the toughness of the quenched and tempered steel, and form coarse carbides distributed along the grain boundaries. Therefore, in the high-strength, toughness and hardenability gear shaft steel described in the present invention, the mass percentage of the Cr element is controlled between 0.95 and 1.44%.
  • Al is a grain refining element.
  • the combination of Al and N can further refine the grains and improve the toughness of the steel.
  • Grain refinement plays an important role in improving the mechanical properties of steel, especially strength and toughness.
  • grain refinement also helps to reduce the hydrogen embrittlement sensitivity of steel.
  • the Al content in the steel should not be too high. Too high Al content tends to increase the chance of inclusions in the steel. Therefore, in the high-strength, toughness and hardenability steel for gear shafts described in the present invention, the mass percentage of Al is controlled between 0.02 and 0.05%.
  • Ti Although Ti can form fine precipitates when added to steel, when the Ti content in the steel is too high, coarse and angular TiN particles will be formed during the smelting process, reducing the impact toughness of the steel. Therefore, in the high-strength, high-toughness and high-hardenability gear shaft steel of the present invention, the Ti content is controlled to be 0.015-0.039%.
  • Nb element is added to the steel to form a fine precipitate phase, thereby inhibiting the recrystallization of the steel and effectively refining the grains. It should be noted that the Nb content in the steel should not be too high. When the Nb content in the steel is too high, coarse NbC particles will be formed during the smelting process, which will reduce the impact toughness of the steel. Therefore, in the high-strength, toughness and hardenability gear shaft steel described in the present invention, the mass percentage of the Nb element is controlled to be 0.001-0.034%.
  • N In the high-strength, toughness and hardenability gear shaft steel described in the present invention, N is an interstitial atom, which can combine with the microalloy in the steel to form MN-type precipitates, which can pin the grain boundaries at high temperatures, thereby inhibiting the growth of austenite grains.
  • the mass percentage of the N element is controlled between 0.006 and 0.015%.
  • B Boron can greatly improve the hardenability of steel, and the required content is relatively small. Its role is generally alloy Hundreds or even thousands of times the boron content of the element, with significant economic effects. Moreover, boron steel can be quenched in water, which not only saves quenching oil but also easily obtains martensitic structure, thus making the boron-containing steel have good strength and hardness. As long as the boron content is appropriate, the production process is appropriate, and complete hardening is guaranteed, the plasticity and toughness will not be significantly reduced. However, the B element is easy to segregate, which will cause large fluctuations in the hardenability of the steel. Therefore, in the high-strength, high-toughness and high-hardenability gear shaft steel described in the present invention, the B element content is controlled between 0.0006 and 0.0034%.
  • Al, Nb, Ti and N are all main grain refining elements.
  • the present invention controls the contents of Al, Nb, Ti and N in the gear steel and the microalloying element coefficient r M/N so that the microalloying elements and the excess N element form precipitates, thereby inhibiting the growth of austenite grains in the high temperature stage.
  • the mass percentage of each impurity element satisfies at least one of the following items:
  • P, O, H and Ca are all impurity elements in steel.
  • the content of impurity elements in steel should be reduced as much as possible.
  • P is easily segregated at the grain boundaries in the steel, which will reduce the grain boundary binding energy and deteriorate the impact toughness of the steel. Therefore, in some embodiments of the present invention, the P content can be controlled to P ⁇ 0.030%.
  • O O easily forms oxides and composite oxides with Al element in steel, destroying the continuity of steel, reducing the uniformity of the structure and the low-temperature impact energy and fatigue performance. Based on this, in some embodiments of the present invention, the O element content can be controlled to O ⁇ 0.0020%.
  • H will accumulate at defects in steel, especially in steel with a tensile strength level exceeding 1000 MPa, and hydrogen-induced delayed fracture will occur. Therefore, in some embodiments of the present invention, the H element content can be controlled to H ⁇ 0.0002%.
  • Ca In the high-strength, high-toughness, high-hardenability gear shaft steel of the present invention, Ca element is easy to form inclusions
  • the Ca content can be controlled to be Ca ⁇ 0.0034%, for example, ⁇ 0.003%.
  • the high-strength, toughness and hardenability gear shaft steel of the present invention further contains at least one of the following chemical elements:
  • At least one of the elements S, Ni, Mo, Cu and V may be further added to further improve the performance of the high-strength, toughness and hardenability gear shaft steel of the present invention.
  • S generally exists as an impurity element in steel, which reduces the plasticity and toughness of the steel.
  • a certain content of S element can form non-metallic inclusions with Mn to improve the cutting performance of the steel.
  • the mass percentage of S is controlled to 0 ⁇ S ⁇ 0.04%, for example, 0.001% ⁇ S ⁇ 0.04%.
  • Ni exists in the steel in the form of solid solution, which can effectively improve the low-temperature impact resistance of the steel.
  • the mass percentage of Ni can be preferably controlled to 0 ⁇ Ni ⁇ 0.25%, for example, 0.03% ⁇ Ni ⁇ 0.25%.
  • Mo can be dissolved in the steel, which is beneficial to improving the hardenability of the steel and improving the strength of the steel. Under high temperature tempering, Mo will also form fine carbides to further improve the strength of the steel, and the combined effect of molybdenum and manganese can significantly improve the stability of austenite. Based on this, in the high-strength, toughness and hardenability gear shaft steel described in the present invention, the mass percentage of Mo can be preferably controlled to 0 ⁇ Mo ⁇ 0.10%, for example, 0.01% ⁇ Mo ⁇ 0.10%.
  • Cu In the high-strength, toughness and hardenability steel for gear shafts described in the present invention, Cu can improve the strength of the steel and is beneficial to improving the weather resistance and corrosion resistance of the steel. However, the Cu content in the steel should not be too high. If the Cu content in the steel is too high, it will be enriched in the grain boundaries during the heating process, resulting in weakening of the grain boundaries and cracking. Therefore, in the high-strength, toughness and hardenability steel for gear shafts described in the present invention, the mass percentage of Cu can be preferably controlled to 0 ⁇ Cu ⁇ 0.20%, for example, 0.03% ⁇ Cu ⁇ 0.20%.
  • V In the high hardenability gear shaft steel of the present invention, V can effectively improve the hardenability of the steel.
  • the V element can form precipitates with the C element or the N element in the steel, thereby further improving the strength of the steel. However, if the content of C and V is too high, coarse VC particles will be formed. Based on this, in the high hardenability gear shaft steel of the present invention, the mass percentage of V element can be controlled to 0 ⁇ V ⁇ 0.03%, for example, 0.005% ⁇ V ⁇ 0.03%.
  • the high-strength, high-toughness, high-hardenability gear shaft steel of the present invention maintains an austenite grain size of 5 to 8 after high-temperature carburizing heat treatment.
  • the conditions of the simulated high-temperature carburizing heat treatment test for detecting the austenite grain size may be: first heating to 1200°C, keeping warm for 40 minutes, water cooling, then heating to 1000°C at a rate of 500-800°C for 40-60 minutes, keeping warm for 4 hours, and water quenching.
  • J9mm hardenability of the high-strength, toughness and hardenability steel for gear shafts described in the present invention is 34-42HRC.
  • the high-strength, high-toughness and high-hardenability gear shaft steel of the present invention has a tensile strength R m ⁇ 1100MPa, a yield strength R p0.2 ⁇ 980MPa, an elongation after fracture A ⁇ 12%, a cross-sectional shrinkage ⁇ 50%, and a Charpy impact energy A ku ⁇ 55 J after high-temperature carburizing heat treatment.
  • the conditions for simulating high-temperature carburizing heat treatment for testing mechanical properties may be: heating at 880 ⁇ 10°C for 90min, oil quenching + heating at 870 ⁇ 10°C for 90min, oil quenching + heating at 200 ⁇ 10°C for 150min, tempering and air cooling.
  • the tensile strength R m of the high-strength, high-toughness and high-hardenability gear shaft steel of the present invention after high-temperature carburizing heat treatment can be 1100 to 1400 MPa.
  • the yield strength R p0.2 of the high-strength, high-toughness and high-hardenability gear shaft steel of the present invention after high-temperature carburizing heat treatment can be 980 MPa to 1250 MPa.
  • the high-strength, high-toughness and high-hardenability gear shaft steel of the present invention may have an elongation after high-temperature carburizing heat treatment of 12 to 15%.
  • the cross-sectional shrinkage of the high-toughness and high-hardenability gear shaft steel of the present invention after high-temperature carburizing heat treatment can be 50 to 62%.
  • the Charpy impact energy A ku of the high-strength, high-toughness and high-hardenability gear shaft steel of the present invention after high-temperature carburizing heat treatment can be 55 to 130 J.
  • another object of the present invention is to provide a method for manufacturing high-strength, toughness and high hardenability steel for gear shafts.
  • the manufacturing method is simple to produce, and the obtained high-strength, toughness and high hardenability steel for gear shafts has the characteristics of high hardenability and high strength and toughness.
  • the present invention also proposes a method for manufacturing the high-strength, toughness and hardenability gear shaft steel as described above, which comprises the steps of:
  • the steel billet is first heated to no higher than 700°C in the preheating section, then further heated to no higher than 980°C in the first heating section, and then heated to 950-1200°C in the second heating section after insulation. After insulation, it enters the soaking section, and the temperature in the soaking section is 1050-1250°C;
  • the steel billet is slowly heated at a rate of 50-300°C/h, first heated to no higher than 700°C in the preheating section, then continued to be heated to no higher than 980°C in the first heating section, and continued to be heated to 950-1200°C in the second heating section, and then entered the soaking section, where the temperature is 1050-1250°C;
  • the heating step adopts a unique process, wherein the soaking section temperature is relatively high, because the relatively high soaking section temperature can improve the composition uniformity and organizational uniformity of the continuous casting billet during the diffusion process of billet heating.
  • the precipitation phase has a relatively fast solid solution rate. Therefore, the high heating temperature will cause more dissolution of the original undissolved precipitation phase particles in the steel, increase the concentration of microalloying elements in the matrix, and precipitate more and more dispersed particles during subsequent cooling.
  • only after the heating temperature is increased upwards can the final rolling or final forging temperature be increased, so that the austenite recovery recrystallization after rolling is more sufficient and the precipitation phase distribution is more uniform.
  • the smelting in step (1) can be performed by electric furnace smelting or converter smelting, and then refined and vacuum treated.
  • a vacuum induction furnace can also be used for smelting.
  • the charge for electric furnace smelting can be low P, S scrap steel, offcuts and high-quality pig iron;
  • the alloy can be prepared as ferrochrome, low-phosphorus ferromanganese, ferromolybdenum, etc.;
  • the reducing agent can include: calcium carbide, carbon powder and aluminum powder; during the oxidation period: slag is frequently removed by slag flow;
  • the slag tapping conditions can be controlled as: the slag tapping temperature is 1630-1660°C; P ⁇ 0.015%;
  • the steel tapping conditions can be controlled as: the steel tapping temperature is 1630-1650°C; [P] ⁇ 0.010%, [C] ⁇ 0.03%.
  • the molten steel can be refined in the ladle refining furnace to remove harmful gases and inclusions in the steel. Control the seating of the ladle, measure the temperature and analyze it, and adjust the argon pressure according to the situation.
  • the ladle refining furnace (Ladle Furnace, LF) can be fed with Al for initial deoxidation, and then alloy blocks can be added and stirred for 5 to 10 minutes.
  • T 1650 ⁇ 1670 ° C
  • vacuum degassing can be performed to ensure that [O] ⁇ 0.0020%, [H] ⁇ 0.00015%.
  • the vacuum degree of vacuum degassing can be controlled to ⁇ 66.7 Pa and maintained for not less than 15 minutes,
  • the bag temperature can be controlled to be 1550-1570°C, thereby reducing Lowering the bag temperature accelerates element diffusion, which is beneficial to further alleviate dendrite segregation.
  • step (2) casting can be performed by die casting or continuous casting.
  • the high-temperature molten steel in the ladle passes through the protective sleeve and is poured into the tundish, and the superheat of the tundish can be controlled to be 20 to 40°C.
  • the tundish is completely cleaned before use, and the inner surface is refractory coated and free of cracks; the molten steel in the tundish passes through the continuous casting crystallizer and is fully electromagnetically stirred, and qualified continuous casting billets with a cross-sectional size of 140mm ⁇ 140mm to 320mm ⁇ 425mm can be cast.
  • step (2) the pouring speed can be controlled to be 0.6-2.1 m/min according to different billet sizes. Then, the continuous casting billet is put into a slow cooling pit for slow cooling, and the slow cooling time can be no less than 24 hours.
  • the heating temperature of the preheating section can be 600-700°C, and the temperature of the first heating section can be 900-980°C.
  • step (3) of the manufacturing method described in the present invention the heat-averaging section needs to be kept warm for a period of time, and the heat-averaging section heat-averaging time can be 3 to 12 hours.
  • step (3) of the manufacturing method of the present invention after heating in the first heating section, insulation may or may not be performed, and then the second heating section is performed, and the insulation time of the first heating section can be 0 to 3 hours, such as 0.5 hours, 1 hour, or 2 hours.
  • step (3) of the manufacturing method of the present invention after heating in the second heating section, insulation may or may not be performed, and then the equalization section is entered, and the insulation time of the second heating section can be 0 to 3 hours, such as 0.5 hours, 1 hour, or 2 hours.
  • step (3) of the manufacturing method described in the present invention slow heating is performed at a rate of 50 to 300°C/h, first heated to no higher than 700°C in the preheating section, then continued to heat to no higher than 980°C in the first heating section without insulation, and continued to heat to 950 to 1200°C in the second heating section without insulation, and entered the equalizing section.
  • This process can be carried out in a walking beam heating furnace.
  • step (4) of the manufacturing method of the present invention the start forging or rolling temperature is controlled to be 1050-1250° C., and the final forging or rolling temperature is controlled to be ⁇ 900° C. In some embodiments, in step (4) of the manufacturing method of the present invention, the final forging or rolling temperature is 900-1000° C.
  • the start forging or rolling temperature is controlled between 1050 and 1250°C, and the final forging or rolling temperature is controlled to be ⁇ 900°C, because: this process further facilitates N to dissolve from the ⁇ solid solution and combine with microalloying elements in the steel to form nitrides.
  • the solubility of N in ⁇ -Fe is less than that in ⁇ -Fe, and due to the stimulation of phase transformation, two peaks of precipitation are caused. If the final forging or final rolling temperature is low, the peak of the precipitation phase will precipitate, which will cause uneven distribution of the precipitation phase and insufficient recovery and recrystallization, resulting in anisotropy in the structure. Therefore, the final forging or final rolling temperature is controlled to be ⁇ 900°C. In addition, increasing the final forging or final rolling temperature will result in finer grains. The finer grains increase the difference between the average grain diameter of ferrite after the transformation of supercooled austenite and the spacing between manganese-rich bands, reducing the tendency of manganese-rich bands to form pearlite, thereby reducing the banded structure.
  • step (4) of the manufacturing method of the present invention after the steel billet is taken out of the furnace, high-pressure water can be used to descale and remove oxide skin.
  • step (4) of the manufacturing method of the present invention when forging is performed, the steel can be directly forged to the final product size.
  • the steel billet can be directly rolled to the final product size, or the steel billet can be first rolled to a specified intermediate billet size and then heated and rolled to the final product size.
  • the steel billet is first rolled into an intermediate billet (the size can be 140mm ⁇ 140mm ⁇ 260mm ⁇ 260mm), and the final rolling temperature of the intermediate billet is controlled to be 1000 ⁇ 1050°C; then the intermediate billet is heated according to the following process: the intermediate billet is first heated to 680 ⁇ 700°C in the preheating section, then heated to 1050 ⁇ 1100°C in the first heating section, and then heated to 1200 ⁇ 1220°C in the second heating section, and the heating rate can be 300 ⁇ 500°C/h; then it enters the soaking section, the temperature of the soaking section is 1200 ⁇ 1250°C, and the holding time of the soaking section can be 3 ⁇ 5h; then the intermediate billet after heat treatment is rolled into a finished product, and the final rolling temperature of the finished product is controlled to be ⁇ 900°C (for example, 900 ⁇ 1000°C).
  • the intermediate billet is first rolled into an intermediate billet (the size can be 140mm ⁇ 140mm ⁇ 260mm ⁇ 260mm), and the final rolling temperature of the intermediate billet is controlled
  • the present invention can develop a gear shaft steel with high hardenability through reasonable chemical composition design combined with optimized process.
  • the bars rolled or forged with the high-strength and high-hardenability gear shaft steel can be effectively processed into gears, and can have excellent strength and toughness after subsequent downstream high-temperature carburizing heat treatment.
  • the high-strength, toughness and hardenability gear shaft steel of the present invention controls the content of microalloying elements and nitrogen, and strictly controls the atomic molar ratio, while adding an appropriate amount of Nb element to hinder the abnormal growth of austenite grains, thereby increasing the austenite grain coarsening temperature of the gear steel, so that the grain size of the gear steel can remain stable at level 5 to 8 after high-temperature carburizing at 1000°C for 4 hours, and various performances meet the performance indicators of gear shaft steel.
  • composition and process design of the high-strength, toughness and hardenability gear shaft steel described in the present invention are reasonable. By controlling the content of micro-alloy elements in the steel, large particles of harmful inclusions in the steel are avoided to ensure stable production quality of the steel, reduce the production cost of the steel, and realize mass production on the bar production line.
  • the high-strength, toughness and hardenability gear shaft steel of the present invention can shorten the carburizing time and reduce the gear shaft production cost when it is subsequently used in the production of gear shafts, and has broad industrial application prospects.
  • the high-strength, toughness and hardenability gear shaft steels of Examples 1-8 are all prepared by the following steps:
  • the smelting can be carried out in a 50kg vacuum induction furnace, a 150kg vacuum induction furnace or a 500kg vacuum induction furnace, or in an electric furnace smelting + refining outside the furnace + vacuum degassing, or in a converter smelting + refining outside the furnace + vacuum degassing.
  • the steel billet is first heated to no higher than 700°C in the preheating section, and then further heated to no higher than 980°C in the first heating section. After insulation, it is further heated to 950-1200°C in the second heating section. After insulation, it enters the soaking section, and the temperature of the soaking section is 1050-1250°C. After insulation, subsequent rolling or forging is carried out.
  • Forging or rolling Control the start forging or rolling temperature to 1050-1250°C, and control the final forging or rolling temperature to ⁇ 900°C.
  • Example 1 According to the chemical composition shown in Tables 1-1 and 1-2 below, smelting was carried out in a 50kg vacuum induction furnace. Molten steel was cast into an ingot, heated and forged to form a blank. The ingot was first heated to 700°C in the preheating section, then continued to be heated to 900°C in the first heating section, and then continued to be heated to 950°C in the second heating section after being kept warm for 0 hours. After being kept warm for 1 hour, it entered the soaking section, the temperature of the soaking section was 1050°C, and after being kept warm for 3 hours, subsequent forging was carried out, and the final forging temperature was controlled to be 910°C, and finally forged into a ⁇ 60mm bar.
  • Example 2 According to the chemical composition shown in Tables 1-1 and 1-2 below, smelting was carried out in a 150kg vacuum induction furnace. The molten steel was cast into an ingot, heated and forged into a blank. The ingot was first heated to 650°C in the preheating section, then continued to be heated to 950°C in the first heating section, kept warm for 0.5h, and then continued to be heated in the second heating section. The steel bar is heated to 1100°C in the first stage and kept at this temperature for 0h before entering the soaking stage. The temperature of the soaking stage is 1200°C. After keeping this temperature for 5h, the subsequent forging is carried out. The final forging temperature is controlled at 1000°C and the steel bar is finally forged into a ⁇ 90mm bar.
  • Example 3 According to the chemical composition shown in Tables 1-1 and 1-2 below, smelting was carried out in a 500kg vacuum induction furnace. Molten steel was cast into an ingot, heated and forged to form a blank. The ingot was first heated to 600°C in the preheating section, then continued to be heated to 980°C in the first heating section, and then continued to be heated to 1200°C in the second heating section after being kept warm for 3 hours. After being kept warm for 3 hours, it entered the soaking section, the temperature of the soaking section was 1250°C, and after being kept warm for 12 hours, subsequent forging was carried out, and the final forging temperature was controlled to be 1000°C, and finally forged into a ⁇ 120mm bar.
  • Example 4 According to the chemical composition shown in Tables 1-1 and 1-2, the steel is smelted in an electric furnace, refined and vacuum treated, and then cast into a 280mm ⁇ 280mm continuous casting billet.
  • the continuous casting billet is heated in a walking beam heating furnace at a rate of 300°C/h. It is first heated to 620°C in the preheating section, then continued to be heated to 950°C in the first heating section, and continued to be heated to 1150°C in the second heating section. It enters the soaking section with a temperature of 1200°C. After being kept warm for 4 hours, it is rolled. The billet is removed from the heating furnace and then rolled. The final rolling temperature is controlled to be 970°C, and finally rolled into a ⁇ 80mm bar.
  • Example 5 According to the chemical composition shown in Tables 1-1 and 1-2, the steel is smelted in an electric furnace, refined and vacuum treated, and then cast into a 320mm ⁇ 425mm continuous casting billet.
  • the continuous casting billet is heated in a walking beam furnace and slowly heated at a rate of 150°C/h. First, it is heated to 600°C in the preheating section, then continued to be heated to 950°C in the first heating section, and continued to be heated to 1200°C in the second heating section. It enters the soaking section, and the soaking section temperature is 1230°C. After keeping warm for 4.5 hours, subsequent rolling is carried out.
  • the billet is rolled after being descaled by high-pressure water after leaving the heating furnace, and rolled into an intermediate billet.
  • the first final rolling temperature (i.e., the intermediate billet final rolling temperature) is controlled to be 1050°C, and the size of the intermediate billet is 220mm ⁇ 220mm. Then the intermediate billet is placed in a walking beam furnace for heating, slowly heated at a speed of 400°C/h, heated to 680°C in the preheating section, 1050°C in the first heating section, and 1200°C in the second heating section. After keeping warm, it enters the equalizing section with an equalizing temperature of 1220°C. It is taken out of the furnace after keeping warm for 3.5 hours, and rolling begins after descaling with high-pressure water.
  • the second final rolling temperature (i.e. the final rolling temperature of the finished product) is controlled to 950°C, and the specification of the finished bar is ⁇ 90mm.
  • Example 6 According to the chemical composition shown in Tables 1-1 and 1-2, electric furnace smelting, refining and vacuum treatment are carried out, and then cast into 280mm ⁇ 280mm continuous casting billets, and the continuous casting billets are controlled to be heated in a walking beam heating furnace, slowly heated at a rate of 300°C/h, first heated to 680°C in the preheating section, and then continued to heat to 900°C in the first heating section, and continued to heat to 1180°C in the second heating section, and entered the equalizing section.
  • the temperature of the equalizing section is 1200°C, and the subsequent rolling is carried out after keeping warm for 4.5 hours.
  • the billet is rolled after being descaled by high-pressure water after leaving the heating furnace, and rolled into an intermediate billet.
  • the first final rolling temperature (i.e., the final rolling temperature of the intermediate billet) is controlled to be 1000°C, the size of the intermediate billet is 140mm ⁇ 140mm. Then the intermediate billet is placed in a walking beam heating furnace and heated slowly at a speed of 500°C/h. First, it is preheated to 700°C, the first heating section is heated to 1100°C, the second heating section is heated to 1220°C, and then enters the soaking section with a soaking temperature of 1220°C. After 3.5h of heat preservation, it is taken out of the furnace, and after descaling with high-pressure water, rolling is started. The second final rolling temperature (i.e. the finished product final rolling temperature) is controlled to 920°C, and the finished bar specification is ⁇ 25mm.
  • Example 7 According to the chemical composition shown in Tables 1-1 and 1-2, the steel is smelted in a converter, refined and vacuum treated, and then cast into a molded billet.
  • the billet is heated in a walking beam heating furnace at a rate of 50°C/h. It is first heated to 620°C in the preheating section, then continued to be heated to 950°C in the first heating section, and continued to be heated to 1150°C in the second heating section. It enters the soaking section with a temperature of 1200°C. After being kept warm for 8 hours, it is rolled. The billet is removed from the heating furnace and then rolled. The final rolling temperature is controlled to be 970°C, and finally rolled into a ⁇ 90mm bar.
  • Example 8 According to the chemical composition shown in Tables 1-1 and 1-2, the steel is smelted in a converter, refined and vacuum treated, and then cast into a molded billet.
  • the billet is heated in a walking beam heating furnace at a rate of 100°C/h. It is first heated to 600°C in the preheating section, then continued to be heated to 950°C in the first heating section, and continued to be heated to 1200°C in the second heating section. It enters the soaking section, and the soaking section temperature is 1230°C. After keeping warm for 7 hours, subsequent rolling is carried out.
  • the billet is rolled after being descaled by high-pressure water after leaving the heating furnace, and rolled into an intermediate billet.
  • the first final rolling temperature (i.e., the intermediate billet final rolling temperature) is controlled to be 1050°C, and the size of the intermediate billet is 260mm ⁇ 260mm. Then place the intermediate billet in a walking beam heating furnace and slowly heat it at a speed of 300°C/h.
  • the preheating section is heated to 680°C, the first heating section is heated to 1050°C, the second heating section is heated to 1200°C, and then enters the equalizing section with an equalizing temperature of 1220°C. After keeping warm for 5 hours, it is taken out of the furnace and rolled after descaling with high-pressure water.
  • the second final rolling temperature (i.e. the final rolling temperature of the finished product) is controlled to 950°C, and the specification of the finished bar is ⁇ 60mm.
  • Comparative Examples 1 and 2 are commercially available materials, which have been smelted and refined in an electric furnace to ensure the purity of the commercial materials.
  • Comparative Example 3 According to the chemical composition shown in Tables 1-1 and 1-2, smelting was carried out in a 50kg vacuum induction furnace. The molten steel was cast into an ingot, heated and forged, and heated in a box furnace. The ingot was heated to 1100°C at a rate of 300°C/h, and then forged after 3 hours of heat preservation. The final forging temperature was controlled to be 910°C, and finally forged into a ⁇ 60mm bar.
  • Comparative Example 4 According to the chemical composition shown in Tables 1-1 and 1-2, electric furnace smelting, refining and vacuum treatment were performed, and then cast into 320mm ⁇ 425mm continuous casting billets.
  • the continuous casting billets were controlled to be heated in a walking beam heating furnace at a rate of 150°C/h, first heated to 600°C in the preheating section, and then heated to 600°C in the second heating section.
  • the first heating section continues to heat to 950°C, and then continues to heat to 1200°C in the second heating section, and enters the soaking section, the soaking section temperature is 1230°C, and the subsequent rolling is carried out after keeping warm for 4.5h.
  • the billet is rolled after being descaled by high-pressure water from the heating furnace, and rolled into an intermediate billet.
  • the first final rolling temperature i.e., the final rolling temperature of the intermediate billet
  • the size of the intermediate billet is 220mm ⁇ 220mm.
  • the intermediate billet is placed in a walking beam heating furnace, slowly heated at a speed of 400°C/h, heated to 680°C in the preheating section, heated to 1050°C in the first heating section, and heated to 1200°C in the second heating section, and enters the soaking section, the soaking temperature is 1220°C, and the temperature is kept warm for 6h.
  • the second final rolling temperature i.e., the final rolling temperature of the finished product
  • the finished bar specification is ⁇ 90mm.
  • Table 1-1 and Table 1-2 list the mass percentages of the chemical elements of the high hardenability gear shaft steels of Examples 1-8 and the comparative steels of Comparative Examples 1-4.
  • Table 1-1 (wt.%, the balance is Fe and other unavoidable impurities except P, O, H and Ca)
  • Table 2 lists the specific process parameters of the high-strength, toughness and hardenability gear shaft steels of Examples 1-8 and the comparative steels of Comparative Examples 1-4 in the above process steps.
  • Example 5 As can be seen from Table 2, in Example 5, Example 6, Example 8 and Comparative Example 4, the steel billet is first rolled to a specified intermediate billet size during rolling, and then intermediately heated and rolled again to a final finished product size.
  • the high-strength, toughness and hardenability gear shaft steels of Examples 1-8 and the comparative steels of Comparative Examples 1-4 were sampled, and subjected to simulated high-temperature carburizing heat treatment tests to perform mechanical property tests and austenite grain size tests, and the test results are listed in Table 3. Among them:
  • Simulated high temperature carburizing heat treatment test for detecting austenite grain size the sample is heated to 1200°C, kept at this temperature for 40min, water-cooled, then heated to 1000°C at a rate of 600°C in 50min, kept at this temperature for 4h, and quenched by water.
  • Austenite grain size test Evaluate austenite grain size according to standard ASTM E112.
  • Hardenability test The steels of the examples and comparative examples were sampled and prepared from hot-rolled round steel according to the national standard GB/T 225, and the end hardenability test (Jominy test) was performed with reference to GB/T 5216, with the normalizing temperature at 920 ⁇ 10°C and the quenching temperature at 870 ⁇ 5°C.
  • the Rockwell hardness test was performed according to GB/T 230.2 to obtain the hardness value (HRC) at a specific position, such as the hardness at 9mm from the quenching end, i.e. J9mm.
  • Table 3 lists the test results of the high-strength, toughness and hardenability gear shaft steels of Examples 1-8 and the comparative steels of Comparative Examples 1-4.
  • the hardenability of the high-toughness and high-hardenability gear shaft steels of Examples 1-8 of the present invention at a representative position J9mm is 34-42HRC, with high hardenability and narrow hardenability bandwidth.
  • the tensile strength Rm of each example after simulated high-temperature carburizing heat treatment is greater than 1100MPa
  • the yield strength Rp0.2 is greater than 980MPa
  • the elongation after fracture A is greater than 12%
  • the cross-sectional shrinkage is greater than 50%
  • the Charpy impact energy Aku is greater than 55J.
  • the bars rolled or forged using this high-strength, toughness and hardenability gear shaft steel can be effectively processed into gear shafts. After high-temperature carburizing heat treatment by downstream users, they have high strength and toughness, and can be effectively used in high-end components such as automotive gearboxes or reducers for new energy vehicles and industrial reducers, and have good use prospects and value.
  • the grain size of comparative example 1 after simulated high-temperature carburizing is level 0, that is, the grains grow abnormally and cannot meet the use requirements.
  • Comparative Example 2 Although the comparative steel of Comparative Example 2 does not have the mixed crystal phenomenon, its grains are fine and hardenability is low after the simulated high-temperature carburizing heat treatment, and it does not meet the requirement of high hardenability. In addition, the strength of Comparative Example 2 is low.
  • the mixed crystal phenomenon (level 4) was also observed in the comparative steel of comparative example 4 after the simulated high temperature carburizing heat treatment at a temperature of 1000°C, where 0(4) means that the average grain size is level 0, while some areas are coarsened to level 4.
  • 0(4) means that the average grain size is level 0, while some areas are coarsened to level 4.
  • the impact energy of comparative example 4 is also relatively low.

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Abstract

本发明公开了一种高强韧高淬透性齿轴用钢，其含有Fe和不可避免的杂质，以及质量百分含量如下的下述各化学元素：C：0.16～0.22％，Si：0.10～0.40％，Mn：0.86～1.24％，Cr：0.95～1.44％，Al：0.02～0.05％，Ti：0.015～0.039％，Nb：0.001～0.034％，N：0.006～0.015％，B：0.0006～0.0034％；其微合金元素系数rM/N的范围为1.5～5.0，其中rM/N=(10*[Nb]/93+[Ti]/480+[Al]/27)/([N]/14)。此外，本发明还公开了上述高强韧高淬透性齿轴用钢的制造方法。本发明所述的高强韧高淬透性齿轴用钢具有较高淬透性及较窄淬透性带宽和良好高温晶粒稳定性。

Description

一种高强韧高淬透性齿轴用钢及其制造方法 技术领域
本发明涉及一种高强度钢材及其制造方法，尤其涉及一种齿轴用钢及其制造方法。
背景技术
汽车工业的发展对汽车零部件提出了越来越高的要求。其中，高强韧、高疲劳寿命和高温稳定且经济高效的齿轮是重要的发展方向。
相应地，高强韧齿轴用钢可满足汽车轻量化对材料的高技术要求。此外，为了保证齿轮的淬火强度，通常还需要对齿轴用钢的淬透性提出要求。
高淬透性MnCr系渗碳齿轮钢的主要技术难题是在提高淬透性的同时，如何避免淬透性带宽过大导致的齿轮尺寸散差，且同时满足高温渗碳后齿轮不出现混晶和晶粒粗大现象。
例如：公开号为CN101096742A，公开日为2008年1月2日，名称为“高强度汽车用齿轮钢”的中国专利文献公开了一种高强度汽车用齿轮钢。其钢中复合加入了Nb、V、Al等合金元素，细化了原始奥氏体晶粒，其成分质量百分比为：C：0.20～0.40，Si：0.20～0.50，Mn：0.50～1.00，Cr：0.80～1.30，Nb：0.015～0.080，V：0.030～0.090，Mo：0.15～0.55，Al：0.015～0.050，其余为Fe和不可避免的杂质。通过加入微量的Nb、V后，来优化齿轮钢的晶粒度、淬透性及其带宽。
又例如：公开号为CN103361559A，公开日为2013年10月23日，名称为“一种Nb、Ti复合微合金化高温渗碳齿轮钢”的中国专利文献，公开了一种Nb、Ti复合微合金化高温渗碳齿轮钢，钢的组分为：C：0.17～0.22％，Si：0.20-0.35％，Mn：0.9～1.10％，P：≤0.025％，S：0.020～0.035％，Cr：1.05～1.30％，Al：0.015～0.035％，Ti：0.02～0.06％，Nb：0.02～0.06％，余量为铁与不可避免的杂质。其通过控制Nb、Ti及Al等微合金元素含量，提高齿轮渗碳温度或缩短渗碳时间。
然而，上述专利文献并没有完全解决高强度齿轴用钢的淬透性以及带宽控制的问题。
发明内容
本发明的目的之一在于提供一种高强韧高淬透性齿轴用钢，其通过优化齿轴用钢成分体系，尤其是合理控制齿轮钢中微合金元素与氮元素的含量，可以较低的成本获得具有较高淬透性、较窄淬透性带宽和良好高温晶粒稳定性的高强度齿轴用钢。
为了实现上述目的，本发明提出了一种高强韧高淬透性齿轴用钢，其除了Fe和不可避免的杂质以外还含有质量百分含量如下的下述各化学元素：
C：0.16～0.22％，
Si：0.10～0.40％，
Mn：0.86～1.24％，
Cr：0.95～1.44％，
Al：0.02～0.05％，
Ti：0.015～0.039％，
Nb：0.001～0.034％，
N：0.006～0.015％，
B：0.0006～0.0034％；
其微合金元素系数rM/N的范围为1.5～5.0，其中rM/N=(10*[Nb]/93+[Ti]/480+[Al]/27)/([N]/14)，式中各化学元素均代入该化学元素质量百分含量的百分号前面的数值。
进一步地，在本发明所述的高强韧高淬透性齿轴用钢中，其各化学元素质量百分含量为：
C：0.16～0.22％，
Si：0.10～0.40％，
Mn：0.86～1.24％，
Cr：0.95～1.44％，
Al：0.02～0.05％，
Ti：0.015～0039％，
Nb：0.001～0.034％，
N：0.006～0.015％，
B：0.0006～0.0034％；
余量为Fe和不可避免的杂质；
其微合金元素系数rM/N的范围为1.5～5.0，其中rM/N=(10*[Nb]/93+[Ti]/480+[Al]/27)/([N]/14)，式中各化学元素均代入该化学元素质量百分含量的百分号前面的数值。
在本发明所述高强韧高淬透性齿轴用钢中，各化学元素的设计原理具体如下所述：
C：在本发明所述的高强韧高淬透性齿轴用钢中，C是钢中所必需的成分，同时其也是影响钢的淬透性最主要的元素之一。高强韧高淬透性齿轴用钢需要表面强度的同时也需要足够的心部冲击韧性，当钢中C元素含量太低时，低于0.16％时，则钢材的强度不足，且不能保证良好的淬透性要求；相应地，钢中C元素含量也不宜太高，当钢中C元素含量太高时，无法满足齿轮心部韧性的需求，且C含量过高对钢材的塑性不利，特别是对Mn含量较高的渗碳齿轮钢，C含量大于0.22％时不利于钢的加工性能。因此，在本发明所述的高强韧高淬透性齿轴用钢中，将C的质量百分比控制在0.16～0.22％之间。
Si：在本发明所述的高强韧高淬透性齿轴用钢中，Si元素不仅能够更好地消除氧化铁对钢的不良影响，其也能溶入铁素体，使铁素体强化，提高钢的强度、硬度、耐磨性和弹性及弹性极限。同时，需要注意的是，Si元素会提高钢的Ac3温度，因导热性较差，有开裂风险以及脱碳倾向。基于此，综合考虑Si的有益效果和不利影响，在本发明所述的高强韧高淬透性齿轴用钢中，将Si的质量百分比控制在0.10～0.40％之间。
Mn：在本发明所述的高强韧高淬透性齿轴用钢中，Mn是影响钢淬透性的主要元素之一。Mn元素的脱氧能力很好，其可以还原钢中的氧化铁，能够有效提高钢的产量。Mn能溶入铁素体，提高钢的强度和硬度，并使钢材在热轧后冷却时得到片层较细、强度较高的珠光体。此外，Mn还能与钢中的S形成MnS，可以消除S的有害作用，其具有使钢形成和稳定奥氏体组织的能力，可以强烈增加钢的淬透性，还能减低钢的高温韧性。当钢中Mn元素含量过小时，钢材的淬透性不足；而当钢中Mn元素含量过高时，则会使钢材的热塑性变差， 影响生产，且钢材在水淬时容易发生裂纹。因此，在本发明所述的高强韧高淬透性齿轴用钢中，将Mn的质量百分比控制在0.86～1.24％之间。
Cr：在本发明所述的高强韧高淬透性齿轴用钢中，Cr是本发明钢中添加的主要合金元素之一，Cr可以显著提高钢的淬透性以及强度、耐磨性等性能。另外，Cr还能降低钢中C元素的活度，可以防止加热、轧制和热处理过程中的脱碳，但是过高的Cr会明显降低淬火及回火钢材的韧性，形成粗大的沿晶界分布的碳化物。因此，在本发明所述的高强韧高淬透性齿轴用钢中，将Cr元素的质量百分比控制在0.95～1.44％之间。
Al：在本发明所述的高强韧高淬透性齿轴用钢中，Al属于细化晶粒元素。Al元素与N配合可进一步细化晶粒，并提高钢材的韧性。晶粒细化在提高钢的力学性能尤其是强度和韧性方面有重要的作用，同时晶粒细化还有助于降低钢的氢脆敏感性。但需要注意的是，钢中Al元素含量不宜过高，过高含量的Al易增加钢中夹杂物产生的机会。因此，在本发明所述的高强韧高淬透性齿轴用钢中，将Al元素的质量百分比控制在0.02～0.05％之间。
Ti：Ti加入钢中虽然可以形成细小析出相，但钢中Ti元素含量过高时，则会在冶炼过程中形成粗大的带棱角的TiN颗粒，降低钢的冲击韧性。因此在本发明所述的高强韧高淬透性齿轴用钢中，控制Ti元素含量在0.015～0.039％。
Nb：在本发明所述的高强韧高淬透性齿轴用钢中，Nb元素加入钢中，能够形成细小析出相，从而起到对钢再结晶的抑制作用，可以有效细化晶粒。需要注意的是，钢中Nb元素含量不宜过高，当钢中Nb含量过高时，在冶炼过程中会形成粗大的NbC颗粒，反而降低钢材的冲击韧性。因此，在本发明所述的高强韧高淬透性齿轴用钢中，将Nb元素的质量百分比控制在0.001～0.034％。
N：在本发明所述的高强韧高淬透性齿轴用钢中，N为间隙原子，其可以与钢中的微合金结合形成MN型析出物，在高温下能够钉扎晶界，从而抑制奥氏体晶粒长大。当钢中N元素含量较低时，则形成的MN少，所起到的钉扎作用不明显；而当钢中N元素含量过高时，则容易在炼钢中富集，降低钢的韧性。因此，在本发明所述的高强韧高淬透性齿轴用钢中，将N元素的质量百分比控制在0.006～0.015％之间。
B：硼能极大地提高钢的淬透性，而且所需含量较少，其作用为一般合金 元素的几百倍乃至上千倍，具有显著的经济效果。而且硼钢可以采用水淬，既节省淬火用油又容易获得马氏体组织，因而使含硼钢具有良好的强度和硬度。只要硼含量合适、生产工艺恰当、保证完全淬透，就不会明显降低塑性和韧性。但B元素容易偏聚，会引起钢材淬透性较大的波动。因此在本发明所述的高强韧高淬透性齿轴用钢中，控制B元素含量在0.0006～0.0034％之间。
另外，重要的是，在本发明所述的高强韧高淬透性齿轴用钢中，需要控制微合金元素系数rM/N的范围为1.5～5.0，其中rM/N=(10*[Nb]/93+[Ti]/480+[Al]/27)/([N]/14)，式中的各化学元素均代入该化学元素的质量百分含量的百分号前面的数值。
在本发明中，Al、Nb、Ti和N均是主要的细化晶粒元素，本发明通过控制齿轮钢中Al、Nb、Ti、N的含量以及微合金元素系数rM/N，使得微合金元素与多余的N元素形成析出物，从而在高温阶段抑制奥氏体晶粒长大。
进一步地，在本发明所述的高强韧高淬透性齿轴用钢的不可避免的杂质中，各杂质元素的质量百分含量满足下述各项的至少其中之一：
P≤0.030％，
O≤0.0020％，
H≤0.0002％，
Ca≤0.0034％。
在上述技术方案中，P、O、H和Ca均为钢中的杂质元素，在技术条件允许情况下，为了获得性能更好且质量更优的钢材，应尽可能降低钢中杂质元素的含量。其中：
P：P容易在钢中晶界处偏聚，会降低晶界结合能，恶化钢的冲击韧性，因此在本发明的一些实施方式中，可以控制P含量为P≤0.030％。
O：O容易与钢中的Al元素形成氧化物以及复合氧化物，破坏钢材连续性，降低组织均匀性和低温冲击功及疲劳性能，基于此，在本发明的一些实施方式中，可以控制O元素含量为O≤0.0020％。
H：H会在钢中缺陷处聚集，尤其是在抗拉强度级别超过1000MPa的钢中，会发生氢致延迟断裂。因此，在本发明的一些实施方式中，可以控制H元素含量为H≤0.0002％。
Ca：在本发明所述的高强韧高淬透性齿轴用钢中，Ca元素容易形成夹杂 物，进而影响最终产品的疲劳性能，因此可以控制Ca元素含量为Ca≤0.0034％，例如≤0.003％。
进一步地，在本发明所述的高强韧高淬透性齿轴用钢中，其还含有下述各化学元素的至少其中之一：
0<S≤0.04％，0<Ni≤0.25％，0<Mo≤0.10％，0<Cu≤0.20％，0<V≤0.03％。
可选地，在本发明所述的高强韧高淬透性齿轴用钢中，可以进一步添加S、Ni、Mo、Cu、V元素的至少其中一种，以进一步地提高本发明所述高强韧高淬透性齿轴用钢的性能。其中：
S：S一般作为钢中的杂质元素存在，会降低钢的塑性和韧性。然而，在本发明所述的高淬透性齿轴用钢中，一定含量的S元素可与Mn形成非金属夹杂物，以改善钢材的切削性能。基于此，在本发明所述的高淬透性齿轴用钢中，将S的质量百分比控制在0<S≤0.04％，例如0.001％≤S≤0.04％。
Ni：在本发明所述的高强韧高淬透性齿轴用钢中，Ni在钢中以固溶形式存在，其可以有效提高钢的低温冲击性能。但需要注意的是，过高的Ni含量会导致钢材中的残留奥氏体含量过高，而降低钢的强度。因此，在本发明所述的高强韧高淬透性齿轴用钢中，可以优选地将Ni的质量百分比控制为0<Ni≤0.25％，例如0.03％≤Ni≤0.25％。
Mo：在本发明所述的高强韧高淬透性齿轴用钢中，Mo可在钢中固溶，其有利于提高钢的淬透性，提高钢材强度。在较高的温度回火下，Mo还会形成细小的碳化物进一步提高钢的强度，钼与锰的联合作用又可以显著提高奥氏体的稳定性。基于此，在本发明所述的高强韧高淬透性齿轴用钢中，可以优选地将Mo的质量百分比控制为0<Mo≤0.10％，例如0.01％≤Mo≤0.10％。
Cu：在本发明所述的高强韧高淬透性齿轴用钢中，Cu可以提高钢材的强度，并有利于提高钢材的耐候性及耐腐蚀能力。然而，钢中Cu元素含量不宜过高，如果钢中Cu含量过高，则在加热过程中会富集在晶界，导致晶界弱化以致开裂。因此，在本发明所述的高强韧高淬透性齿轴用钢中，可以优选地将Cu的质量百分比控制为0<Cu≤0.20％，例如0.03％≤Cu≤0.20％。
V：在本发明所述的高淬透性齿轴用钢中，V可以有效提高钢的淬透性。在钢中V元素可以与C元素或N元素形成析出物，从而进一步提高钢的强度。 然而，如果C元素和V元素含量过高，则会形成粗大的VC颗粒。基于此，在本发明所述的高淬透性齿轴用钢中，可以将V元素的质量百分比控制在0<V≤0.03％，例如0.005％≤V≤0.03％。
进一步地，本发明所述的高强韧高淬透性齿轴用钢在高温渗碳热处理后的奥氏体晶粒度保持在5～8级。其中，用于检测奥氏体晶粒度的模拟高温渗碳热处理试验的条件可以是：先加热至1200℃，保温40min，水冷，然后以40-60min加热500-800℃的速度加热至1000℃，保温4h，水冷淬火。
进一步地，本发明所述的高强韧高淬透性齿轴用钢的J9mm淬透性为34～42HRC。
进一步地，本发明所述的高强韧高淬透性齿轴用钢在高温渗碳热处理后的抗拉强度Rm≥1100MPa、屈服强度Rp0.2≥980MPa、断后伸长率A≥12％，断面收缩率≥50％，夏比冲击功Aku≥55J。其中，用于检测力学性能的模拟高温渗碳热处理的条件可以是：880±10℃加热90min油淬+870±10℃加热90min油淬+200±10℃加热150min回火空冷。
进一步地，本发明所述的高强韧高淬透性齿轴用钢在高温渗碳热处理后的抗拉强度Rm可以为1100～1400MPa。
进一步地，本发明所述的高强韧高淬透性齿轴用钢在高温渗碳热处理后的屈服强度Rp0.2可以为980MPa～1250MPa。
进一步地，本发明所述的高强韧高淬透性齿轴用钢在高温渗碳热处理后的断后伸长率A可以为12～15％。
进一步地，本发明所述的高强韧高淬透性齿轴用钢在高温渗碳热处理后的断面收缩率可以为50～62％.
进一步地，本发明所述的高强韧高淬透性齿轴用钢在高温渗碳热处理后的夏比冲击功Aku可以为55～130J。
相应地，本发明的另一目的在于提供一种高强韧高淬透性齿轴用钢的制造方法，该制造方法生产简单，所获得的高强韧高淬透性齿轴用钢具有淬透性高、强韧性高的特点。
为了实现上述目的，本发明还提出了如上文所述的高强韧高淬透性齿轴用钢的制造方法，其包括步骤：
(1)冶炼；
(2)铸造；
(3)加热：钢坯首先在预热段加热至不高于700℃，然后在第一加热段继续加热至不高于980℃，保温后继续在第二加热段加热至950～1200℃，保温后进入均热段，均热段温度为1050～1250℃；
或者钢坯以50～300℃/h的速度缓慢加热，首先在预热段加热不高于700℃，然后在第一加热段继续加热至不高于980℃，继续在第二加热段加热至950～1200℃，进入均热段，均热段温度为1050～1250℃；
(4)锻造或轧制。
在本发明所述的制造方法中，同现有技术相比，其加热步骤采用了独特的工艺，其中均热段温度较高，这是因为较高的均热段温度能够在钢坯加热的扩散过程中提高连铸坯的成分均匀性和组织均匀度。同时在此温度下，析出相有着较快的固溶速度。因此，加热温度高将使钢中原始未溶的析出相粒子有更多的溶解，使基体中微合金元素浓度增加，在以后冷却时析出更多更弥散的粒子。此外只有将加热温度向上提高以后，才能使终轧或终锻温度提高，从而使轧后奥氏体回复再结晶更充分，析出相分布更均匀。
本发明所述的制造方法中，步骤(1)中的冶炼可以采用电炉冶炼或转炉冶炼，并经过精炼及真空处理。当然在一些其他的实施方式中，也可以采用真空感应炉进行冶炼。
此外，在步骤(1)中，电炉冶炼的炉料可以选用低P、S废钢、切头及优质生铁；合金可以准备铬铁、低磷锰铁、钼铁等；还原剂可以包括：电石、碳粉和铝粉；在氧化期：勤流渣去P；可以控制出渣条件为：出渣温度为1630～1660℃；P≤0.015％；可以控制出钢条件为：出钢温度为1630～1650℃；[P]≤0.010％，[C]≥0.03％。
电炉冶炼或转炉冶炼完成后，可以在钢包精炼炉上进行钢液精炼，以去除钢中的有害气体和夹杂物。控制钢包入座、测温并分析，可以根据情况调整氩气压力。钢包精炼炉(Ladle Furnace，LF)初脱氧可以喂Al，然后可以补加合金块搅拌5～10分钟。当钢液测温T＝1650～1670℃时，可以进行真空脱气，以保证[O]≤0.0020％、[H]≤0.00015％。在一个具体的实例中，真空脱气的真空度可以控制为≤66.7Pa，且保持不低于15分钟，
另外，在步骤(1)中，可以控制吊包温度为1550～1570℃，由此由于降 低了吊包温度，加快了元素扩散，有利于进一步减轻枝晶偏析。
此外，在步骤(2)中，铸造可以采用模铸或连铸。连铸浇注过程中钢包内高温钢液通过保护套管，浇进中间包，可以控制中间包过热度为20～40℃。中间包使用前完全清理、内表面为耐火涂层且不得有裂缝；中间包内的钢液经连铸结晶器，电磁搅拌充分，可以浇注出140mm×140mm～320mm×425mm断面尺寸的合格连铸坯。
在步骤(2)中，可以依据不同的方坯尺寸控制浇注速度为0.6～2.1m/min。然后，使连铸坯进缓冷坑缓冷，缓冷时间可以不少于24小时。
进一步地，在本发明所述的制造方法的步骤(3)中，预热段加热温度可以为600～700℃，第一加热段温度可以为900～980℃。
进一步地，在本发明所述的制造方法的步骤(3)中，在均热段需保温一段时间，均热段保温时间可以为3～12h。
在本发明所述的制造方法的步骤(3)中，进行第一加热段加热后，可以进行或不进行保温，然后进行第二段加热，第一加热段保温时间可以为0～3h，例如0.5h、1h、2h。在本发明所述的制造方法的步骤(3)中，在进行第二加热段加热后，可以进行或不进行保温，然后进入均热段，第二加热段保温时间可以为0～3h，例如0.5h、1h、2h。
在另外一些实施方案中，在本发明所述的制造方法的步骤(3)中，以50～300℃/h的速度缓慢加热，首先在预热段加热至不高于700℃，然后在第一加热段继续加热至不高于980℃，不进行保温，继续在第二加热段加热至950～1200℃，不进行保温，进入均热段，此过程可以在步进梁式加热炉中进行。
进一步地，在本发明所述的制造方法的步骤(4)中，控制开锻或开轧温度为1050～1250℃，控制终锻或终轧温度≥900℃。在一些实施方式中，在本发明所述的制造方法的步骤(4)中，终锻或终轧温度为900～1000℃。
在这种实施方式下，控制开锻或开轧温度在1050～1250℃之间，并控制终锻或终轧温度≥900℃，是因为：该工艺进一步有利于N从γ固溶体中脱溶并与钢中的微合金元素结合成氮化物。
需要说明的是，N在α-Fe中的溶解度小于在γ-Fe中的溶解度，且由于受相变的激发而造成析出量的二个峰值，如果终锻或终轧温度低，析出相的峰值析出，会造成析出相分布不均匀以及回复再结晶不充分而产生组织上的各向异 性，所以控制终锻或终轧温度≥900℃。另外提高终锻或终轧温度，会得到较细的晶粒，晶粒细小增大了过冷奥氏体转变后的铁素体平均晶粒直径和富锰带带间距之间的差别，减轻了富锰带形成珠光体的趋势，从而减轻了带状组织。
在上述技术方案中，在本发明所述制造方法的步骤(4)中，在钢坯出炉后，可以采用高压水除鳞去氧化皮。
此外，在本发明所述制造方法的步骤(4)中，当进行锻造时，可以直接锻造至最终成品尺寸。当进行轧制时，既可以采用钢坯直接轧制到最终成品尺寸，也可以采用钢坯先轧制到指定的中间坯尺寸，再进行加热和轧制到最终成品尺寸。
在一些实施方案中，在本发明所述制造方法的步骤(4)中，先将钢坯轧制成中间坯(尺寸可以为140mm×140mm～260mm×260mm)，控制中间坯终轧温度为1000～1050℃；再对中间坯按照以下工艺进行加热：中间坯首先在预热段加热至680～700℃，然后在第一加热段加热至1050～1100℃，然后在第二加热段加热至1200～1220℃，加热速度可以为300～500℃/h；然后进入均热段，均热段温度为1200～1250℃，均热段保温时间可以为3～5h；然后将加热处理后的中间坯轧制成成品，控制成品终轧温度≥900℃(例如900～1000℃)。
本发明所述的高强韧高淬透性齿轴用钢及其制造方法相较于现有技术具有如下所述的优点以及有益效果：
(1)本发明通过合理的化学成分设计并结合优化工艺，可以开发出具有高淬透性的齿轴用钢，采用该高强韧高淬透性齿轴用钢轧制或锻造好的棒材能够有效加工成齿轮，经过后续下游的高温渗碳热处理，可以具有优异的强韧性。
(2)本发明所述的高强韧高淬透性齿轴用钢通过控制微合金元素系数与氮元素的含量，并且严格控制原子摩尔比，同时加入适量Nb元素，以阻碍奥氏体晶粒的异常长大，提高了齿轮钢的奥氏体晶粒粗化温度，使该齿轮钢在高达1000℃高温渗碳4小时后晶粒度仍稳定保持在5～8级，各项性能达到齿轴用钢的使用性能指标。
(3)本发明所述高强韧高淬透性齿轴用钢的成分和工艺设计合理，通过控制钢中微合金元素的含量，从而避免钢材中出现大颗粒有害夹杂，以保证钢材稳定的生产质量，降低了钢材的生产成本，实现在棒材产线上的批量生产。
(4)本发明所述高强韧高淬透性齿轴用钢的淬透性及奥氏体晶粒度和成 本竞争力等方面均优于现有技术，可在保证高淬透性及窄带宽等性能的前提下，控制钢材中合金元素的种类和数量，提高钢材的适用性。
(5)本发明所述高强韧高淬透性齿轴用钢在后续用于生产制造齿轴时，可以缩短渗碳时间，降低齿轴生产成本，具有广阔的工业应用前景。
具体实施方式
下面将结合具体的实施例对本发明所述的高强韧高淬透性齿轴用钢及其制造方法做进一步的解释说明，然而该解释和说明并不对本发明的技术方案构成不当限定。
实施例1-8的高强韧高淬透性齿轴用钢均采用以下步骤制得：
(1)按照下述表1-1和1-2所示的化学成分进行冶炼和浇铸：其中冶炼可以采用50kg真空感应炉、150kg真空感应炉或500kg真空感应炉进行冶炼，也可以采用电炉冶炼+炉外精炼+真空脱气的方式进行冶炼，还可以采用转炉冶炼+炉外精炼+真空脱气的方式进行冶炼。
(2)铸造。
(3)加热：钢坯首先在预热段加热至不高于700℃，然后在第一加热段继续加热至不高于980℃，保温后继续在第二加热段加热至950～1200℃，保温后进入均热段，均热段温度为1050～1250℃，保温后进行后续轧制或锻造。
(4)锻造或轧制：控制开锻或开轧温度为1050～1250℃，控制终锻或终轧温度≥900℃。
进一步地，实施例1-8高淬透性齿轴用钢和对比例1-4钢的具体工艺过程如下所述：
实施例1：按照下述表1-1和1-2所示的化学成分在50kg真空感应炉上进行冶炼。钢水浇铸成钢锭，加热并经锻造开坯，控制钢锭首先在预热段加热至700℃，然后在第一加热段继续加热至900℃，保温0h后继续在第二加热段加热至950℃，保温1h后进入均热段，均热段温度为1050℃，保温3h后进行后续锻造，控制终锻温度为910℃，最终锻造成Φ60mm棒料。
实施例2：按照下述表1-1和1-2所示的化学成分在150kg真空感应炉上进行冶炼。钢水浇铸成钢锭，加热并经锻造开坯，控制钢锭首先在预热段加热至650℃，然后在第一加热段继续加热至950℃，保温0.5h后继续在第二加热 段加热至1100℃，保温0h后进入均热段，均热段温度为1200℃，保温5h后进行后续锻造，控制终锻温度为1000℃，最终锻造成Φ90mm棒料。
实施例3：按照下述表1-1和1-2所示的化学成分在500kg真空感应炉上进行冶炼。钢水浇铸成钢锭，加热并经锻造开坯，控制钢锭首先在预热段加热至600℃，然后在第一加热段继续加热至980℃，保温3h后继续在第二加热段加热至1200℃，保温3h后进入均热段，均热段温度为1250℃，保温12h后进行后续锻造，控制终锻温度为1000℃，最终锻造成Φ120mm棒料。
实施例4：按照表1-1和1-2所示的化学成分电炉冶炼，并进行精炼和真空处理，而后浇铸成280mm×280mm连铸坯，控制连铸坯在步进梁式加热炉中加热，以300℃/h的速度缓慢加热，首先在预热段加热至620℃，然后在第一加热段继续加热至950℃，继续在第二加热段加热至1150℃，进入均热段，均热段温度为1200℃，保温后4h后进行轧制。钢坯出加热炉经高压水除鳞后开始轧制，控制终轧温度为970℃，最终轧制成Φ80mm棒料。
实施例5：按照表1-1和1-2所示的化学成分电炉冶炼，并进行精炼和真空处理，而后浇铸成320mm×425mm连铸坯，控制连铸坯在步进梁式炉中加热，以150℃/h的速度缓慢加热，首先在预热段加热至600℃，然后在第一加热段继续加热至950℃，继续在第二加热段加热至1200℃，进入均热段，均热段温度为1230℃，保温4.5h后进行后续轧制。钢坯出加热炉经高压水除鳞后开始轧制，轧制成中间坯，控制第一终轧温度(即中间坯终轧温度)为1050℃，中间坯尺寸220mm×220mm。而后再将中间坯置于步进梁式炉中加热，以400℃/h的速度缓慢加热，预热段加热至680℃，第一加热段加热至1050℃，第二加热段加热至1200℃，保温后进入均热段，均热温度1220℃，保温3.5h后出炉，经高压水除鳞后开始轧制，控制第二终轧温度(即成品终轧温度)为950℃，成品棒材规格为Φ90mm。
实施例6：按照表1-1和1-2所示的化学成分电炉冶炼，并进行精炼和真空处理，而后浇铸成280mm×280mm连铸坯，控制连铸坯在步进梁式加热炉中加热，以300℃/h的速度缓慢加热，首先在预热段加热至680℃，然后在第一加热段继续加热至900℃，继续在第二加热段加热至1180℃，进入均热段，均热段温度为1200℃，保温4.5h后进行后续轧制。钢坯出加热炉经高压水除鳞后开始轧制，轧制成中间坯，控制第一终轧温度(即中间坯终轧温度)为 1000℃，中间坯尺寸140mm×140mm。而后再将中间坯置于步进梁式加热炉中加热，以500℃/h的速度缓慢加热，首先预热至700℃，第一加热段加热至1100℃，第二加热段加热至1220℃，进入均热段，均热温度1220℃，保温3.5h后出炉，经高压水除鳞后开始轧制，控制第二终轧温度(即成品终轧温度)为920℃，成品棒材规格为Φ25mm。
实施例7：按照表1-1和1-2所示的化学成分转炉冶炼，并进行精炼和真空处理，而后浇铸成模铸坯，控制铸坯在步进梁式加热炉中加热，以50℃/h的速度缓慢加热，首先在预热段加热至620℃，然后在第一加热段继续加热至950℃，继续在第二加热段加热至1150℃，进入均热段，均热段温度为1200℃，保温8h后进行轧制。钢坯出加热炉经高压水除鳞后开始轧制，控制终轧温度为970℃，最终轧制成Φ90mm棒料。
实施例8：按照表1-1和1-2所示的化学成分转炉冶炼，并进行精炼和真空处理，而后浇铸成模铸坯，控制铸坯在步进梁式加热炉中加热，以100℃/h的速度缓慢加热，首先在预热段加热至600℃，然后在第一加热段继续加热至950℃，继续在第二加热段加热至1200℃，进入均热段，均热段温度为1230℃，保温7h后进行后续轧制。钢坯出加热炉经高压水除鳞后开始轧制，轧制成中间坯，控制第一终轧温度(即中间坯终轧温度)为1050℃，中间坯尺寸260mm×260mm。而后再将中间坯置于步进梁式加热炉中，以300℃/h的速度缓慢加热，预热段加热至680℃，第一加热段加热至1050℃，第二加热段加热至1200℃，进入均热段，均热温度1220℃，保温5h后出炉，经高压水除鳞后开始轧制，控制第二终轧温度(即成品终轧温度)为950℃，成品棒材规格为Φ60mm。
对比例1和对比例2来自市售的商品材，其经过了电炉冶炼和精炼处理，以保证商品材的纯净度。
对比例3：按照表1-1和1-2所示的化学成分在50kg真空感应炉上进行冶炼。钢水浇铸成钢锭，加热并经锻造开坯，在箱式炉中加热，控制钢锭按300℃/h的速度加热到1100℃，保温3h后进行后续锻造，控制终锻温度为910℃，最终锻造成Φ60mm棒料。
对比例4：按照表1-1和1-2所示的化学成分电炉冶炼，并进行精炼和真空处理，而后浇铸成320mm×425mm连铸坯，控制连铸坯在步进梁式加热炉中加热，以150℃/h的速度缓慢加热，首先在预热段加热至600℃，然后在第 一加热段继续加热至950℃，继续在第二加热段加热至1200℃，进入均热段，均热段温度为1230℃，保温4.5h后进行后续轧制。钢坯出加热炉经高压水除鳞后开始轧制，轧制成中间坯，控制第一终轧温度(即中间坯终轧温度)为1050℃，中间坯尺寸220mm×220mm。而后将中间坯置于步进梁式加热炉中，以400℃/h的速度缓慢加热，在预热段加热至680℃，第一加热段加热至1050℃，第二加热段加热至1200℃，进入均热段，均热温度1220℃，保温6h，出炉经高压水除鳞后开始轧制，控制第二终轧温度(即成品终轧温度)为950℃，成品棒材规格为Φ90mm。
表1-1和表1-2列出了实施例1-8的高淬透性齿轴用钢和对比例1-4的对比钢的各化学元素的质量百分配比。
表1-1(wt.％，余量为Fe和除P、O、H和Ca以外的其他不可避免的杂质)
表1-2.(wt.％，余量为Fe和除P、O、H和Ca以外的其他不可避免的杂质)


注：rM/N＝(10*[Nb]/93+[Ti]/480+[Al]/27)/([N]/14)，式中的各化学元素均代入该化学元
素的质量百分含量的百分号前面的数值。
表2列出了实施例1-8的高强韧高淬透性齿轴用钢和对比例1-4的对比钢在上述工艺步骤中的具体工艺参数。
表2.

从表2中可以看出，实施例5、实施例6、实施例8和对比例4在轧制时均是将钢坯先轧制到指定的中间坯尺寸，而后再次进行中间加热和二次轧制到最终成品尺寸。
为了验证本发明所述的高强韧高淬透性齿轴用钢的性能，将得到的实施例1-8的高强韧高淬透性齿轴用钢和对比例1-4的对比钢分别取样，并进行模拟高温渗碳热处理试验后，以进行力学性能测试和奥氏体晶粒度检测，将检测结果列于表3中。其中：
用于检测奥氏体晶粒度的模拟高温渗碳热处理试验：试样加热至1200℃，保温40min，水冷，然后以50min加热600℃的速度加热至1000℃，保温4h，水冷淬火。
奥氏体晶粒度检测：按照标准ASTM E112评定奥氏体晶粒度。
力学性能测试：按照GB/T 2975-2018《钢及钢产品力学性能试验取样位置及试样制备》制备样品，参考GB/T 3077-2015，制备Φ15mm毛坯，用于检测力学性能的模拟高温渗碳热处理试验：880℃加热90min油淬+870℃加热90min油淬+200℃加热150min回火空冷，并按照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，测得抗拉强度Rm、屈服强度Rp0.2、断后伸长率A和断面收缩率Z，同时采用GB/T 229-2007《金属材料夏比摆锤冲击试验方法》测试各实施例和对比例在室温夏比冲击功Aku。
此外，还将得到的实施例1-8的高强韧高淬透性齿轴用钢和对比例1-4的对比钢分别取样，进行淬透性测试和硬度测试，将测试试验结果同样列于表3中。其中：
淬透性测试：各实施例钢和对比例钢按照国家标准GB/T 225从热轧圆钢上取样、制样，参考GB/T 5216进行末端淬透性测试(Jominy试验)，控制正火温度920±10℃，淬火温度870±5℃。根据GB/T 230.2进行洛氏硬度测试，得到特定位置的硬度值(HRC)，比如距离淬火端9mm处的硬度，即J9mm。
表3列出了实施例1-8的高强韧高淬透性齿轴用钢和对比例1-4的对比钢的测试试验结果。
表3.


注：表3中每一个实施例和对比例的冲击功具有斜杠前后的两个数值表示两次测量结果。
从表3中可以看出，本发明所述实施例1-8的高强韧高淬透性齿轴用钢经过1000℃模拟高温渗碳热处理后，其奥氏体晶粒度都维持在5～8级范围内，未见混晶、晶粒异常粗大等现象，因此其具有良好的高温晶粒稳定性。
此外，从表3中还可以看出，本发明各实施例1-8的高强韧高淬透性齿轴用钢其代表性位置J9mm淬透性均为34～42HRC，具有较高的淬透性及较窄的淬透性带宽。另外，各实施例在模拟高温渗碳热处理后的抗拉强度Rm均大于1100MPa、屈服强度Rp0.2均大于980MPa、断后伸长率A均≥12％，断面收缩率均大于50％，夏比冲击功Aku均大于55J。
由此可见，采用该高强韧高淬透性齿轴用钢轧制或锻造好的棒材能够有效加工成齿轴，经过下游用户的高温渗碳热处理，其具有高强韧性，从而可以有效应用于汽车用变速箱或新能源车用减速器及工业减速器等高端零部件中，具有良好的使用前景和价值。
与本发明各实施例不同的是，对比例1模拟高温渗碳后的晶粒度为0级，即晶粒异常长大，不能满足使用要求。
对比例2的对比钢虽然未出现混晶现象，但其在模拟高温渗碳热处理后晶粒细小，淬透性较低，达不到高淬透性的要求。此外，对比例2的强度较低。
对比例3的对比钢在1000℃的温度下模拟高温渗碳热处理后观察到了混晶现象，其中5(0)表示平均晶粒度为5级，部分区域粗化为0级。
而对比例4的对比钢在1000℃的温度下模拟高温渗碳热处理后也观察到了混晶现象(4级)，其中0(4)表示平均晶粒度为0级，而部分区域粗化为4级。此外，对比例4的冲击功也较低。
此外，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 (11)

  1. 一种高强韧高淬透性齿轴用钢，其含有Fe和不可避免的杂质，其特征在于，还含有质量百分含量如下的下述各化学元素：
    C：0.16～0.22％，Si：0.10～0.40％，Mn：0.86～1.24％，Cr：0.95～1.44％，
    Al：0.02～0.05％，Ti：0.015～0.039％，Nb：0.001～0.034％，N：0.006～0.015％，B：0.0006～0.0034％；
    其微合金元素系数rM/N的范围为1.5～5.0，其中rM/N=(10*[Nb]/93+[Ti]/480+[Al]/27)/([N]/14)，式中各化学元素均代入该化学元素质量百分含量的百分号前面的数值。
  2. 如权利要求1所述的高强韧高淬透性齿轴用钢，其特征在于，其各化学元素质量百分含量为：
    C：0.16～0.22％，Si：0.10～0.40％，Mn：0.86～1.24％，Cr：0.95～1.44％，
    Al：0.02～0.05％，Ti：0.015～0.039％，Nb：0.001～0.034％，N：0.006～0.015％，B：0.0006～0.0034％；余量为Fe和其他不可避免的杂质；
    其微合金元素系数rM/N的范围为1.5～5.0，其中rM/N＝(10*[Nb]/93+[Ti]/480+[Al]/27)/([N]/14)，式中各化学元素均代入该化学元素质量百分含量的百分号前面的数值。
  3. 如权利要求1或2所述的高强韧高淬透性齿轴用钢，其特征在于，在不可避免的杂质中，P≤0.030％、O≤0.002％、H≤0.0002％、Ca≤0.0034％。
  4. 如权利要求1或2所述的高强韧高淬透性齿轴用钢，其特征在于，其还含有下述各化学元素的至少其中之一：0<S≤0.04％，0<Ni≤0.25％，0<Mo≤0.10％，0<Cu≤0.20％，0<V≤0.03％。
  5. 如权利要求1或2所述的高强韧高淬透性齿轴用钢，其特征在于，其在高温渗碳热处理后的奥氏体晶粒度保持在5～8级。
  6. 如权利要求1或2所述的高强韧高淬透性齿轴用钢，其特征在于，其J9mm淬透性为34～42HRC。
  7. 如权利要求1或2所述的高强韧高淬透性齿轴用钢，其特征在于，其在高温渗碳热处理后的抗拉强度Rm≥1100MPa、屈服强度Rp0.2≥980MPa、断后伸长率A≥12％，断面收缩率≥50％，夏比冲击功Aku≥55J。
  8. 如权利要求1-7中任意一项所述的高强韧高淬透性齿轴用钢的制造方法，其特征在于，其包括步骤：
    (1)冶炼；
    (2)铸造；
    (3)加热：钢坯首先在预热段加热至不高于700℃，然后在第一加热段继续加热至不高于980℃，保温后继续在第二加热段加热至950～1200℃，保温后进入均热段，均热段温度为1050～1250℃；
    (4)锻造或轧制。
  9. 如权利要求8所述的制造方法，其特征在于，在步骤(4)中，控制开锻或开轧温度为1050～1250℃，控制终轧温度或终锻温度≥900℃。
  10. 如权利要求8所述的制造方法，其特征在于，在步骤(4)中，直接轧制或锻造至成品尺寸。
  11. 如权利要求8所述的制造方法，其特征在于，在步骤(4)中，先轧制到中间坯尺寸，再进行中间加热，然后轧制到最终成品尺寸。
PCT/CN2024/099794 2023-06-19 2024-06-18 一种高强韧高淬透性齿轴用钢及其制造方法 Pending WO2024260333A1 (zh)

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Cited By (1)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
CN120138484A (zh) * 2025-03-27 2025-06-13 重庆望江工业有限公司江苏分公司 一种优化风电齿轮钢成分的微合金化方法

Citations (7)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
EP1167561A2 (en) * 2000-06-28 2002-01-02 Mitsubishi Steel Muroran Inc. Carburizing and carbonitriding steel
CN101096742A (zh) 2006-06-28 2008-01-02 宝山钢铁股份有限公司 高强度汽车用齿轮钢
CN103361559A (zh) 2013-07-18 2013-10-23 首钢总公司 一种Nb、Ti复合微合金化高温渗碳齿轮钢
CN113755752A (zh) * 2021-08-24 2021-12-07 江苏利淮钢铁有限公司 一种高强韧性工程机械轮体用30Mn2CrTiB钢及其生产方法
CN114000055A (zh) * 2021-10-21 2022-02-01 山东钢铁股份有限公司 一种硼微合金化齿轮钢及其制备方法
CN114635086A (zh) * 2022-03-17 2022-06-17 襄阳金耐特机械股份有限公司 一种高强韧性铸钢
CN116254470A (zh) * 2023-02-17 2023-06-13 山东钢铁股份有限公司 一种Nb-Ti微合金化Cr-Mn系齿轮钢及其制备方法

Family Cites Families (3)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
CN104789892B (zh) * 2015-03-20 2017-03-08 宝山钢铁股份有限公司 具有优异低温冲击韧性的低屈强比高强韧厚钢板及其制造方法
CN115261715A (zh) * 2021-04-29 2022-11-01 宝山钢铁股份有限公司 一种高温渗碳齿轴用钢及其制造方法
CN115369315A (zh) * 2021-05-21 2022-11-22 宝山钢铁股份有限公司 一种高温渗碳高淬透性齿轮用钢及其制造方法

Patent Citations (7)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
EP1167561A2 (en) * 2000-06-28 2002-01-02 Mitsubishi Steel Muroran Inc. Carburizing and carbonitriding steel
CN101096742A (zh) 2006-06-28 2008-01-02 宝山钢铁股份有限公司 高强度汽车用齿轮钢
CN103361559A (zh) 2013-07-18 2013-10-23 首钢总公司 一种Nb、Ti复合微合金化高温渗碳齿轮钢
CN113755752A (zh) * 2021-08-24 2021-12-07 江苏利淮钢铁有限公司 一种高强韧性工程机械轮体用30Mn2CrTiB钢及其生产方法
CN114000055A (zh) * 2021-10-21 2022-02-01 山东钢铁股份有限公司 一种硼微合金化齿轮钢及其制备方法
CN114635086A (zh) * 2022-03-17 2022-06-17 襄阳金耐特机械股份有限公司 一种高强韧性铸钢
CN116254470A (zh) * 2023-02-17 2023-06-13 山东钢铁股份有限公司 一种Nb-Ti微合金化Cr-Mn系齿轮钢及其制备方法

Cited By (2)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
CN120138484A (zh) * 2025-03-27 2025-06-13 重庆望江工业有限公司江苏分公司 一种优化风电齿轮钢成分的微合金化方法
CN120138484B (zh) * 2025-03-27 2026-01-30 重庆望江工业有限公司 一种优化风电齿轮钢成分的微合金化方法

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