WO2021223758A1 - 一种可形成复合耐蚀层的变形高温合金及其制备工艺 - Google Patents

一种可形成复合耐蚀层的变形高温合金及其制备工艺 Download PDF

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WO2021223758A1
WO2021223758A1 PCT/CN2021/092490 CN2021092490W WO2021223758A1 WO 2021223758 A1 WO2021223758 A1 WO 2021223758A1 CN 2021092490 W CN2021092490 W CN 2021092490W WO 2021223758 A1 WO2021223758 A1 WO 2021223758A1
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temperature
alloy
hours
resistant layer
composite corrosion
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French (fr)
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严靖博
谷月峰
袁勇
杨征
张醒兴
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西安热工研究院有限公司
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    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C19/00—Alloys based on nickel or cobalt
    • C22C19/03—Alloys based on nickel or cobalt based on nickel
    • C22C19/05—Alloys based on nickel or cobalt based on nickel with chromium
    • C22C19/051—Alloys based on nickel or cobalt based on nickel with chromium and Mo or W
    • C22C19/056—Alloys based on nickel or cobalt based on nickel with chromium and Mo or W with the maximum Cr content being at least 10% but less than 20%
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C1/00—Making non-ferrous alloys
    • C22C1/02—Making non-ferrous alloys by melting
    • C22C1/023—Alloys based on nickel
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22F—CHANGING THE PHYSICAL STRUCTURE OF NON-FERROUS METALS AND NON-FERROUS ALLOYS
    • C22F1/00—Changing the physical structure of non-ferrous metals or alloys by heat treatment or by hot or cold working
    • C22F1/10—Changing the physical structure of non-ferrous metals or alloys by heat treatment or by hot or cold working of nickel or cobalt or alloys based thereon

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  • the invention belongs to the field of high-temperature alloys, and specifically relates to a deformed high-temperature alloy capable of forming a composite corrosion-resistant layer and a preparation process thereof, and is particularly suitable for over/reheaters of advanced ultra-supercritical units of thermal power plants, as well as hydrogen-producing converters and high-temperature gas Key components of cold reactor, etc.
  • the processing and preparation are relatively complicated, which poses new challenges to the selection of superalloys.
  • the common solid solution strengthening element W in superalloys has an obvious tendency to macrosegregate, while the precipitation strengthening element Nb is easy to form a CrNbN phase (Z phase) at the grain boundary, which in turn harms the processing and preparation properties of gold pipes. Therefore, when selecting high-temperature alloys as advanced ultra-supercritical pipe components, it is necessary to ensure the high-temperature performance of the alloy while taking into account its processing performance.
  • the content of W element in the alloy composition is strictly controlled to reduce the red hardness and macro segregation of the material, and the Mo element is used to achieve the effect of solid solution strengthening.
  • the addition of Mo element is not conducive to the coal ash corrosion resistance of the alloy, so it is necessary to ensure that the alloy has sufficient content of corrosion resistance elements such as Al and Cr at the same time.
  • the pure Al 2 O 3 layer is unstable in the coal ash corrosive environment, so it is necessary to ensure that the alloy has a sufficient Cr content, but this will also cause problems such as the nucleation of the Z phase at the grain boundary. Therefore, it is necessary to reasonably adjust the amount of Nb element added to improve the stability of Ni 3 Al ( ⁇ ') while avoiding the precipitation of harmful phases at the grain boundary due to excessive addition of it and jeopardizing the processing performance of the alloy.
  • the purpose of the present invention is to develop a deformed high-temperature alloy capable of forming a composite corrosion-resistant layer and its preparation process.
  • a deformed high-temperature alloy capable of forming a composite corrosion-resistant layer includes: Cr: 16-19%, Co: 10-15%, Ti: 0.5-1.5%, Al: 3.5-4.5%, W: ⁇ 0.5%, Mo: ⁇ 5.0%, Si: ⁇ 0.5%, Mn: ⁇ 0.5%, Nb: 0.5-1.0%, C: 0.04-0.07%, and the balance is Ni.
  • a preparation process of a deformed high-temperature alloy capable of forming a composite corrosion-resistant layer includes the following steps:
  • Alloy smelting According to mass percentage, Cr: 16-19%, Co: 10-15%, Ti: 0.5-1.5%, Al: 3.5-4.5%, W: ⁇ 0.5%, Mo: ⁇ 5.0% , Si: ⁇ 0.5%, Mn: ⁇ 0.5%, Nb: 0.5-1.0%, C: 0.04-0.07%, the balance is Ni; added to a vacuum induction furnace, alloy smelting under vacuum and argon protection, Refining, and finally obtaining alloy ingots;
  • High-temperature rolling homogenize the alloy ingot at 1160-1200°C for 24-72 hours, and then roll it at a high temperature of 50-100°C above the Ni 3 Al ⁇ ' dissolution temperature, and the amount of deformation per pass Not less than 25%, and the total deformation is not less than 70%;
  • a further improvement of the present invention is that in step 1), the mass percentage of N element in the alloy ingot is not more than 0.03%, and the mass percentage of P and S is not more than 0.03%.
  • a further improvement of the present invention is that before step 2), the alloy ingot is kept at 950-1020°C for 0.5-1.0 hours before step 2) is performed.
  • a further improvement of the present invention is that the temperature rise rate is not higher than 10°C/min to 950-1020°C.
  • a further improvement of the present invention is that in step 2), when high temperature rolling is performed, the outside of the alloy ingot is sheathed with 304 stainless steel with a thickness of 0.5-1.0 mm.
  • a further improvement of the present invention is that in step 2), after each pass of deformation is completed, the furnace is returned to the furnace for 15-20 minutes and then the next rolling is performed.
  • step 3 is: heating the rolled alloy with the furnace to a temperature of 30-70°C above the ⁇ 'dissolution temperature for 3-5 hours, and then air cooling to room temperature after completion; The alloy is heated to 300-350°C below the ⁇ 'melting temperature for 3-9 hours and then air-cooled, and finally heated to 200-250°C below the ⁇ 'melting temperature for 1-3 hours and then air-cooled.
  • the present invention has the following beneficial effects:
  • the present invention can form a composite oxide layer, exert the excellent protective effect of alumina on the matrix under high temperature conditions, and avoid direct contact with the outer coal ash layer at the same time.
  • the W element content in the alloy composition is controlled to reduce the red hardness and macro segregation of the material, and the Mo element is used to achieve the solid solution strengthening effect, which guarantees the high temperature performance of the alloy while taking into account its processing performance.
  • Figure 1 is a photo of the organization of Example 1;
  • Figure 2 shows the morphology of the ⁇ 'enhanced phase in Example 1;
  • Figure 3 is the test result of the thermal expansion coefficient of Example 1;
  • Figure 4 shows the impact toughness fracture surface of Example 1
  • Figure 5 is a photo of the tissue of Example 2.
  • Figure 6 is a tissue photograph of Example 2 after 1000 hours of heat exposure
  • Figure 7 is a cross-sectional photo of Example 2 coal ash corrosion after 500 hours
  • Figure 8 is a photo of the organization of the comparative example
  • Figure 9 is a photo of the tissue of the comparative example exposed to heat for 1000 hours.
  • Figure 10 is a cross-sectional photograph of the coal ash of the comparative example after 500 hours of corrosion.
  • the invention is developed in response to the requirements of advanced ultra-supercritical thermal power units, and can meet the performance requirements of high-temperature components such as superheaters/reheaters.
  • the alloy composition meets the following requirements in terms of mass percentage: Cr: 16-19%, Co: 10-15%, Ti: 0.5-1.5%, Al: 3.5-4.5%, W: ⁇ 0.5%, Mo: ⁇ 5.0%, Si: ⁇ 0.5%, Mn: ⁇ 0.5%, Nb: 0.5-1.0%, C: 0.04-0.07%, the balance is Ni; homogenization treatment after melting, hot rolling, and final heat treatment.
  • the tensile yield strength of the alloy at room temperature and 850°C of the present invention is higher than 820MPa and 520MPa, respectively.
  • the present invention includes the following steps:
  • Alloy smelting Use vacuum induction furnace for alloy smelting, and make sure that the vacuum degree is lower than 5*10 -3 before introducing high-purity argon gas. The electroslag remelting process is then used for refining, and finally an alloy ingot is obtained for processing.
  • the alloy meets the following requirements: in terms of mass percentage, the alloy in terms of mass percentage, including: Cr: 16-19%, Co: 10-15 %, Ti: 0.5 to 1.5%, Al: 3.5 to 4.5%, W: ⁇ 0.5%, Mo: ⁇ 5.0%, Si: ⁇ 0.5%, Mn: ⁇ 0.5%, Nb: 0.5 to 1.0%, C: 0.04 ⁇ 0.07%, the balance is Ni; the N element content of the alloy after electroslag remelting is not more than 300ppm, and the P and S content is not more than 0.03%;
  • the average grain size of the alloy after heat treatment is 30-50 ⁇ m, and a large number of dispersed ⁇ 'phases are precipitated in the crystal, and the average diameter does not exceed 50nm.
  • the discontinuous carbides are distributed on the grain boundary, and the size of the discontinuous carbide does not exceed 5 ⁇ m.
  • the tensile yield strength of the alloy at room temperature and 850°C is higher than 820MPa and 520MPa, respectively. At the same time, it has excellent processing properties and structural stability. No harmful phases are precipitated during heat exposure at 850°C, and after heat exposure at this temperature for 1350 hours The tensile yield strength at room temperature and 850°C is higher than 630MPa and 370MPa respectively.
  • a composite layer of Cr 2 O 3 and Al 2 O 3 is formed on the surface of the alloy, and its weight changes Less than 0.2mg/cm 2 .
  • the alloy is smelted by a vacuum induction furnace, and the obtained alloy includes Cr: 16%, Co: 15%, Ti: 1.1%, Al: 4.1%, Si: 0.3%, Mn: 0.3%, Nb: 1.0 by mass percentage. %, C: 0.04%, the balance is Ni.
  • Use vacuum induction furnace for alloy smelting and ensure that the vacuum degree is lower than 5*10 -3 before introducing high-purity argon gas.
  • the electroslag remelting process is then used for refining, and finally an alloy ingot is obtained for processing. Ensure that the N element content of the alloy after electroslag remelting is not higher than 300ppm, and the P and S content is not higher than 0.03%.
  • the alloy is homogenized at 1160°C for 24 hours, and then rolled at a temperature of 100°C above the ⁇ 'dissolution temperature.
  • the total deformation is 95%, and the deformation per pass is not less than 30%.
  • the heating rate is controlled to 10°C/min, and after the temperature is increased to 950 for 0.5 hours, the heating is continued to the specified temperature for homogenization treatment.
  • the rolled alloy is heated in the furnace to 70°C above the ⁇ 'dissolution temperature for 4 hours, and then air-cooled to room temperature after completion. Subsequently, the alloy was heated to 300°C below the ⁇ 'dissolution temperature for 8 hours and then air-cooled, and finally heated to 200°C below the ⁇ 'dissolution temperature for 2 hours and then air-cooled.
  • Figure 1 and Figure 2 are photos of the structure and ⁇ 'phase morphology after heat treatment in Example 1. It can be seen that after the heat treatment, the average grain size does not exceed 50 ⁇ m, and a large number of dispersed ⁇ 'phases are precipitated in the crystals, and the average diameter is not More than 50nm.
  • the discontinuous carbides are distributed on the grain boundary, and the size of the discontinuous carbide does not exceed 5 ⁇ m.
  • Figures 3 and 4 show the thermal expansion coefficient test results and impact toughness fracture surface of Example 1.
  • the thermal expansion coefficients at 800 and 850°C do not exceed 15.49 and 16.48*10 -6 K -1 , respectively, and the room temperature impact toughness is 74J/cm 2 .
  • a large number of dimple features can be observed on the surface of the impact fracture, indicating that it has good fracture toughness.
  • the alloy is smelted by a vacuum induction furnace, and the obtained alloy includes Cr: 17%, Co: 10%, Ti: 1.5%, Al: 4.0%, Si: 0.2%, Mo: 5.0%, Mn: 0.2 by mass percentage. %, Nb: 1.0%, C: 0.04%, and the balance is Ni.
  • Use vacuum induction furnace for alloy smelting and ensure that the vacuum degree is lower than 5*10 -3 before introducing high-purity argon gas.
  • the electroslag remelting process is then used for refining, and finally an alloy ingot is obtained for processing. Ensure that the N element content of the alloy after electroslag remelting is not higher than 300ppm, and the P and S content is not higher than 0.03%.
  • the alloy is homogenized at 1200°C for 24 hours, and then rolled at a temperature of 100°C above the ⁇ 'dissolution temperature.
  • the total deformation is 70%, and the deformation per pass is not less than 25%.
  • the heating rate is controlled to 10°C/min, and the temperature is increased to 1020 for 0.5 hours and then the temperature is increased to the specified temperature for homogenization.
  • the rolled alloy is heated in the furnace to 70°C above the ⁇ 'dissolution temperature for 4 hours, and then air-cooled to room temperature after completion. Subsequently, the alloy was heated to 300°C below the ⁇ 'dissolution temperature for 8 hours and then air-cooled, and finally heated to 200°C below the ⁇ 'dissolution temperature for 2 hours and then air-cooled.
  • the yield strength of the alloy at room temperature and 850°C is 879MPa and 570MPa, respectively.
  • the tensile yield strength at room temperature and 850°C is higher than 688MPa and 418MPa, respectively.
  • Figures 5 and 6 are the structure photos of the heat-treated state and the long-term thermal exposure state of Example 2. It can be seen that no obvious harmful phases are precipitated in the crystal grains during heat exposure at 850°C, which proves that it has excellent structural stability.
  • Figure 7 is a cross-sectional photo of the alloy of Example 2 after being corroded by coal ash at 850°C high temperature flue gas environment (N 2 -15% CO 2 -3.5% O 2 -0.1% SO 2 ) for 500 hours. It can be seen that the corrosion layer is from the outer layer. Composed of chromium oxide and inner alumina, its weight gain is only 0.16mg/cm 2 .
  • Alloy smelting According to mass percentage, Cr: 16%, Co: 15%, Ti: 1.5%, Al: 3.5%, W: 0.5%, Mo: 3.0%, Si: 0.3%, Mn: 0.1% , Nb: 0.5%, C: 0.04%, the balance is Ni; added to a vacuum induction furnace, alloy smelting and refining under the protection of vacuum and argon gas, and finally an alloy ingot is obtained; the quality of N in the alloy ingot is 100%
  • the content of P and S is not more than 0.03%, and the mass percentage of P and S is not more than 0.03%.
  • High temperature rolling After raising the alloy ingot to 950°C for 1.0 hour at a heating rate of 10°C/min, the alloy ingot is homogenized at 1160°C for 72 hours, and then the outside thickness is 0.5 -1.0mm 304 stainless steel is sheathed, and high temperature rolling is performed at 50°C above the dissolution temperature of Ni 3 Al ⁇ '. The deformation of each pass is not less than 25%, and the total deformation is not less than 70%; After completion, return to the furnace for 20 minutes and then proceed to the next rolling.
  • Alloy smelting In terms of mass percentage, Cr: 17%, Co: 12%, Ti: 1%, Al: 4%, W: 0.1%, Mo: 5.0%, Si: 0.2%, Mn: 0.4% , Nb: 1%, C: 0.04%, the balance is Ni; added to a vacuum induction furnace, alloy smelting and refining under the protection of vacuum and argon gas, and finally an alloy ingot is obtained; the quality of N in the alloy ingot is 100%
  • the content of P and S is not more than 0.03%, and the mass percentage of P and S is not more than 0.03%.
  • High temperature rolling After raising the alloy ingot to 1020°C at a heating rate of 3°C/min for 0.5 hours, the alloy ingot is homogenized at 1120°C for 24 hours, and then the outside thickness is 0.5 -1.0mm 304 stainless steel is sheathed, and high temperature rolling is performed at 70°C above the dissolution temperature of Ni 3 Al ⁇ ', and the deformation of each pass is not less than 25%, and the total deformation is not less than 70%; After completion, return to the furnace for 15 minutes and then proceed to the next rolling.
  • Alloy smelting In terms of mass percentage, Cr: 19%, Co: 10%, Ti: 0.5%, Al: 4.5%, Mo: 1.0%, Si: 0.5%, Mn: 0.5%, Nb: 0.7% , C: 0.05%, the balance is Ni; added to a vacuum induction furnace, alloy smelting and refining under the protection of vacuum and argon gas, and finally an alloy ingot is obtained; the mass percentage of N element in the alloy ingot is not higher than 0.03%, the mass percentage content of P and S is not higher than 0.03%.
  • High temperature rolling After raising the alloy ingot to 1000°C at a heating rate of 1°C/min for 0.7 hours, then homogenizing the alloy ingot at 1180°C for 50 hours, and then adopting a thickness of 0.5 on the outside. -1.0mm 304 stainless steel is sheathed, and high temperature rolling is carried out at 100°C above the dissolution temperature of Ni 3 Al ⁇ ', and the deformation of each pass is not less than 25%, and the total deformation is not less than 70%; After completion, return to the furnace for 17 minutes and perform the next rolling.
  • the alloy is smelted by a vacuum induction furnace, and the obtained alloy includes Cr: 21%, Co: 15%, Ti: 1.8%, Al: 4.5%, Si: 0.2%, Mo: 7.0%, Mn: 0.2 by mass percentage. %, Nb: 0.5%, C: 0.04%, the balance is Ni.
  • Use vacuum induction furnace for alloy smelting and ensure that the vacuum degree is lower than 5*10 -3 before introducing high-purity argon gas.
  • the electroslag remelting process is then used for refining, and finally an alloy ingot is obtained for processing. Ensure that the N element content of the alloy after electroslag remelting is not higher than 300ppm, and the P and S content is not higher than 0.03%.
  • the alloy is homogenized at 1200°C for 24 hours, and then rolled at a temperature of 100°C above the ⁇ 'dissolution temperature.
  • the total deformation is 70%, and the deformation per pass is not less than 25%.
  • the heating rate is controlled to 10°C/min, and the temperature is increased to 1020 for 0.5 hours and then the temperature is increased to the specified temperature for homogenization.
  • the rolled alloy is heated in the furnace to 70°C above the ⁇ 'dissolution temperature for 4 hours, and then air-cooled to room temperature after completion. Subsequently, the alloy was heated to 300°C below the ⁇ 'dissolution temperature for 8 hours and then air-cooled, and finally heated to 200°C below the ⁇ 'dissolution temperature for 2 hours and then air-cooled.
  • the yield strength of the alloy at room temperature and 850°C after heat treatment is 930MPa and 586MPa, respectively, and the tensile yield strength at room temperature and 850°C after 1000 hours of heat exposure at 850°C is higher than 955MPa and 438MPa, respectively.
  • Figures 8 and 9 are the microstructure photos of the heat treatment state and the long-term thermal exposure state of the comparative example. It can be seen that a large number of harmful phases appear in the crystal grains during the heat exposure at 850°C, indicating that the microstructure stability is poor.
  • Figure 10 is a cross-sectional photo of the comparative alloy after being corroded by coal ash in a high temperature flue gas environment (N 2 -15% CO 2 -3.5% O 2 -0.1% SO 2 ) at 850°C for 500 hours. A complete alumina layer is not formed on the inside. , So the weight loss is as high as 0.96mg/cm 2 .

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Abstract

一种可形成复合耐蚀层的变形高温合金及其制备工艺，按质量百分比计包括：Cr：16～19％，Co：10～15％，Ti：0.5～1.5％，Al：3.5～4.5％，W：≤0.5％，Mo：≤5.0％，Si：≤0.5％，Mn：≤0.5％，Nb：0.5～1.0％，C：0.04～0.07％，余量为Ni；熔炼后均匀化处理，热轧，最后热处理。本发明的合金室温及850℃拉伸屈服强度分别高于820MPa与520MPa，同时其具有优异的加工性能及组织稳定性，在850℃热暴露期间无有害相析出，且在该温度下经热暴露1350小时后室温及850℃拉伸屈服强度分别高于630MPa与370MPa。

Description

一种可形成复合耐蚀层的变形高温合金及其制备工艺 技术领域
本发明属高温合金领域，具体涉及一种可形成复合耐蚀层的变形高温合金及其制备工艺，特别适用于火电先进超超临界机组过/再热器等，以及制氢转化炉、高温气冷堆关键部件等。
背景技术
随着我国用电需求不断增加，能源紧缺及环境污染问题日益凸显，发展高效、节能、环保发电方式的需求越发紧迫。火力发电作为我国长期以来最主要的发电技术，提高机组蒸汽参数被认为是解决上述问题最有效的途径。以往大量实践表明，材料的服役性能是制约锅炉机组蒸汽参数提高的最主要原因，而作为火电机组锅炉中服役工况最严苛的关键部件，过/再热器、主蒸汽管、集箱等对材料的服役性能提出了极高的要求。上述部件在服役期间将承受高温蠕变、热疲劳、氧化及高温烟气腐蚀等多重因素的影响。随着火电机组主蒸汽参数的大幅提高，开发出可以满足高参数机组关键部件使用性能需求的高温合金材料已成为火力发电行业亟待解决的课题。
针对高参数火电机组锅炉关键部件对材料使用性能的需求，目前国外已进行了大量研究，并已开发出一系列新型合金材料，如美国特殊金属公司开发的Inconel 740H、美国哈氏公司开发的Haynes 282、德国蒂森克虏伯公司开发的CCA 617、英国Rolls-Royce公司开发的Nimonic 263、日本日立公司开发的FENIX700、日本东芝公司开发的TOS1X、日本三菱公司开发的LTESR700等镍基变形高温合金。而研究表明，材料强度的不断提高往往以牺牲合金加工性能为代价。由于过/再热器、主蒸汽管、集箱等部件多以管材为主，加工制备相对 较为复杂，对高温合金的选材提出了新的挑战。例如，高温合金中常见的固溶强化元素W具有明显的宏观偏析倾向，同时析出强化元素Nb易于在晶界形成CrNbN相(Z相)，进而危害和金管材的加工制备性能。因此，选用高温合金作为先进超超临界管材部件时，需要在保障合金高温性能的同时兼顾其加工性能。
为确保材料的制备加工性能，严格控制合金成分中W元素含量以降低材料红硬性及宏观偏析等问题，并采用Mo元素从而达到固溶强化的效果。然而，Mo元素的添加不利于合金的抗煤灰腐蚀性能，因此需要同时保障合金中具备足够的Al、Cr等耐蚀元素含量。然而，单纯的Al 2O 3层在煤灰腐蚀环境下不稳定，因此合金中需保证具有足够的Cr含量，但这也同时将造成Z相在晶界形核等问题。因此，还需合理调整Nb元素添加量，在改善Ni 3Al(γ’)稳定性的同时，避免因其过量添加造成晶界有害相析出并危害合金加工性能。
发明内容
本发明的目的在于开发一种可形成复合耐蚀层的变形高温合金及其制备工艺。
为了实现以上发明目的，本发明所采用的技术方案为：
一种可形成复合耐蚀层的变形高温合金，按质量百分数计，包括：Cr：16～19％，Co：10～15％，Ti：0.5～1.5％，Al：3.5～4.5％，W：≤0.5％，Mo：≤5.0％，Si：≤0.5％，Mn：≤0.5％，Nb：0.5～1.0％，C：0.04～0.07％，余量为Ni。
一种可形成复合耐蚀层的变形高温合金的制备工艺，包括以下步骤：
1)合金冶炼：按质量百分数计，将Cr：16～19％，Co：10～15％，Ti：0.5～1.5％，Al：3.5～4.5％，W：≤0.5％，Mo：≤5.0％，Si：≤0.5％，Mn：≤0.5％，Nb：0.5～1.0％，C：0.04～0.07％，余量为Ni；加入到真空感应炉中，在真空以及氩气保护下进行 合金冶炼，精炼，最终获得合金铸锭；
2)高温轧制：对合金铸锭在1160-1200℃进行24-72小时的均匀化处理，然后将其在Ni 3Alγ’溶解温度以上50-100℃进行高温轧制，每道次变形量不低于25％，总变形量不低于70％；
3)热处理。
本发明进一步的改进在于，步骤1)中，合金铸锭中N元素质量百分含量不高于0.03％，P与S质量百分含量均不高于0.03％。
本发明进一步的改进在于，进行步骤2)前，将合金铸锭在950-1020℃保温0.5-1.0小时后再进行步骤2)。
本发明进一步的改进在于，以不高于10℃/min的升温速率升至950-1020℃。
本发明进一步的改进在于，步骤2)中，进行高温轧制时，合金铸锭外部采用厚度0.5-1.0mm的304不锈钢包套。
本发明进一步的改进在于，步骤2)中，在每道次变形完成后回炉保温15-20min后进行下一道次轧制。
本发明进一步的改进在于，步骤3)的具体过程为：将轧制后的合金随炉升温至γ’溶解温度以上30-70℃范围内保温3-5小时，完成后空冷至室温；随后将合金加热至γ’溶解温度以下300-350℃范围内保温3-9小时后空冷，最后加热至γ’溶解温度以下200-250℃范围内保温1-3小时后空冷。
与现有技术相比，本发明具有的有益效果：
由于氧化铝在煤灰腐蚀环境中不稳定，容易失去其保护基体的效果。所以本发明通过控制合金中的Cr、Al含量及其相互比例，可以形成复合氧化层，发挥出氧化铝在高温条件下对基体优异的保护作用，同时避免其与外层煤灰层直接接触。此外，本发明中通过控制合金成分中W元素含量以降低材料红硬性及 宏观偏析等问题，并采用Mo元素从而达到固溶强化的效果，在保障合金高温性能的同时兼顾其加工性能。同时，合理调整Nb元素添加量，在改善Ni 3Al(γ’)稳定性的同时，避免过量添加造成晶界有害相析出并危害对合金加工性能。通过调整控制合金中Al、Cr等耐蚀元素含量及其比例，促进煤灰腐蚀环境下复合耐蚀层的形成，最终获得一种具有优异高温强度、抗腐蚀/氧化性能且兼具良好组织稳定性及加工性能的新型高温合金。
附图说明
图1为实施例1组织照片；
图2为实施例1中γ’强化相形貌；
图3为实施例1热膨胀系数测试结果；
图4为实施例1冲击韧性断口表面；
图5为实施例2组织照片
图6为实施例2热暴露1000小时组织照片
图7为实施例2煤灰腐蚀500小时后截面照片
图8为对比例组织照片；
图9为对比例热暴露1000小时组织照片。
图10为对比例煤灰腐蚀500小时后截面照片。
具体实施方式
下面结合实施例对本发明作进一步详细说明。
本发明是针对先进超超临界火电机组要求而开发的，可满足过热器/再热器等高温部件的使用性能需求。合金成分按质量百分比满足如下范围要求：Cr：16～19％，Co：10～15％，Ti：0.5～1.5％，Al：3.5～4.5％，W：≤0.5％，Mo：≤5.0％，Si：≤0.5％，Mn：≤0.5％，Nb：0.5～1.0％，C：0.04～0.07％，余量为Ni；熔炼后均匀化处理，热轧，最后热处理。本发明的合金室温及850℃拉伸屈服强度分别 高于820MPa与520MPa，同时其具有优异的加工性能及组织稳定性，在850℃热暴露期间无有害相析出，且在该温度下经热暴露1350小时后室温及850℃拉伸屈服强度分别高于630MPa与370MPa。此外，合金经850℃高温烟气环境(N 2-15％CO 2-3.5％O 2-0.1％SO 2)腐蚀500小时后表面形成Cr 2O 3与Al 2O 3复合层，其重量变化小于0.2mg/cm 2。
本发明包括以下步骤：
1)合金冶炼：采用真空感应炉进行合金冶炼，通入高纯氩气前确保真空度低于5*10 -3。随后采用电渣重熔工艺精炼，最终获得合金铸锭以备加工；其中，合金满足以下要求：按质量百分数计，合金按质量百分数计，包括：Cr：16～19％，Co：10～15％，Ti：0.5～1.5％，Al：3.5～4.5％，W：≤0.5％，Mo：≤5.0％，Si：≤0.5％，Mn：≤0.5％，Nb：0.5～1.0％，C：0.04～0.07％，余量为Ni；合金经电渣重熔后N元素含量不高于300ppm，P、S含量不高于0.03％；
2)高温轧制：对合金在1160-1200℃进行24-72小时的均匀化处理，然后将其在Ni 3Al(γ’)溶解温度以上50-100℃进行高温轧制，总变形量不低于70％，每道次变形量不低于25％；其中，合金均匀化升温过程中控制其升温速率不高于10℃/min，并在升至950-1020℃保温0.5-1.0小时后继续升温至指定温度进行均匀化处理；轧制过程中合金锭外部采用厚度0.5-1.0mm的304不锈钢包套以减缓出炉后温度下降速率，并在变形完成后回炉保温15-20min后进行下一道次轧制；
3)热处理：将轧制后的合金随炉升温至γ’溶解温度以上30-70℃范围内保温3-5小时，完成后空冷至室温；随后将合金加热至γ’溶解温度以下300-350℃范围内保温3-9小时后空冷，最后加热至γ’溶解温度以下200-250℃范围内保温 1-3小时后空冷。
合金经过热处理后平均晶粒尺寸30-50μm，晶内大量析出弥散分布的γ’相，其平均直径不超过50nm。晶界分布不连续碳化物，其尺寸最大不超过5μm。
合金室温及850℃拉伸屈服强度分别高于820MPa与520MPa，同时其具有优异的加工性能及组织稳定性，在850℃热暴露期间无有害相析出，且在该温度下经热暴露1350小时后室温及850℃拉伸屈服强度分别高于630MPa与370MPa。此外，合金经850℃高温烟气环境(N 2-15％CO 2-3.5％O 2-0.1％SO 2)腐蚀500小时后表面形成Cr 2O 3与Al 2O 3复合层，其重量变化小于0.2mg/cm 2。
实施例1
合金在确保满足强度性能的同时，限制W、Nb等元素的添加以改善其加工性能。利用真空感应炉对合金进行熔炼，获得的合金按质量百分比包括：Cr：16％，Co：15％，Ti：1.1％，Al：4.1％，Si：0.3％，Mn：0.3％，Nb：1.0％，C：0.04％，余量为Ni。采用真空感应炉进行合金冶炼，通入高纯氩气前确保真空度低于5*10 -3。随后采用电渣重熔工艺精炼，最终获得合金铸锭以备加工。确保合金经电渣重熔后N元素含量不高于300ppm，P、S含量不高于0.03％。
对合金在1160℃进行24小时的均匀化处理，然后将其在γ’溶解温度以上100℃进行高温轧制，总变形量95％，每道次变形量不低于30％。合金均匀化升温过程中控制其升温速率10℃/min，并在升至950保温0.5小时后继续升温至指定温度进行均匀化处理。
将轧制后的合金随炉升温至γ’溶解温度以上70℃保温4小时，完成后空冷至室温。随后将合金加热至γ’溶解温度以下300℃保温8小时后空冷，最后加热至γ’溶解温度以下200℃保温2小时后空冷。
图1与图2为实施例1热处理后组织及γ’相形貌照片，可以看出经过热处理后其平均晶粒尺寸不超过50μm，晶内大量析出弥散分布的γ’相，其平均直径不超过50nm。晶界分布不连续碳化物，其尺寸最大不超过5μm。
图3与图4为实施例1热膨胀系数测试结果及冲击韧性断口表面，其在800、850℃热膨胀系数分别不超过15.49与16.48*10 -6K -1，室温冲击韧性74J/cm 2。对冲击断口表面观察可见大量韧窝特征，表明其具有良好的断裂韧性。
实施例2
合金在确保满足强度性能的同时，限制W、Nb等元素的添加以改善其加工性能。利用真空感应炉对合金进行熔炼，获得的合金按质量百分比包括：Cr：17％，Co：10％，Ti：1.5％，Al：4.0％，Si：0.2％，Mo：5.0％，Mn：0.2％，Nb：1.0％，C：0.04％，余量为Ni。采用真空感应炉进行合金冶炼，通入高纯氩气前确保真空度低于5*10 -3。随后采用电渣重熔工艺精炼，最终获得合金铸锭以备加工。确保合金经电渣重熔后N元素含量不高于300ppm，P、S含量不高于0.03％。
对合金在1200℃进行24小时的均匀化处理，然后将其在γ’溶解温度以上100℃进行高温轧制，总变形量70％，每道次变形量不低于25％。合金均匀化升温过程中控制其升温速率10℃/min，并在升至1020保温0.5小时后继续升温至指定温度进行均匀化处理。
将轧制后的合金随炉升温至γ’溶解温度以上70℃保温4小时，完成后空冷至室温。随后将合金加热至γ’溶解温度以下300℃保温8小时后空冷，最后加热至γ’溶解温度以下200℃保温2小时后空冷。合金经热处理后室温及850℃屈服强度分别为879MPa与570MPa，经850℃热暴露1350小时后室温及850℃拉伸 屈服强度分别高于688MPa与418MPa。
图5与图6为实施例2热处理态及长期热暴露态的组织照片，可见其在850℃热暴露期间晶粒内部无明显有害相析出，证实了其具备优异的组织稳定性。
图7为实施例2合金经850℃高温烟气环境(N 2-15％CO 2-3.5％O 2-0.1％SO 2)腐蚀500小时煤灰腐蚀后截面照片，可见其腐蚀层由外层氧化铬及内侧氧化铝组成，其增重仅0.16mg/cm 2。
实施例3
1)合金冶炼：按质量百分数计，将Cr：16％，Co：15％，Ti：1.5％，Al：3.5％，W：0.5％，Mo：3.0％，Si：0.3％，Mn：0.1％，Nb：0.5％，C：0.04％，余量为Ni；加入到真空感应炉中，在真空以及氩气保护下进行合金冶炼，精炼，最终获得合金铸锭；合金铸锭中N元素质量百分含量不高于0.03％，P与S质量百分含量均不高于0.03％。
2)高温轧制：将合金铸锭以10℃/min的升温速率升至950℃保温1.0小时后，再对合金铸锭在1160℃进行72小时的均匀化处理，然后将其外部采用厚度0.5-1.0mm的304不锈钢包套后，在Ni 3Alγ’溶解温度以上50℃进行高温轧制，每道次变形量不低于25％，总变形量不低于70％；在每道次变形完成后回炉保温20min后进行下一道次轧制。
3)热处理：将轧制后的合金随炉升温至γ’溶解温度以上30℃范围内保温3小时，完成后空冷至室温；随后将合金加热至γ’溶解温度以下300℃范围内保温9小时后空冷，最后加热至γ’溶解温度以下200℃范围内保温3小时后空冷。
实施例4
1)合金冶炼：按质量百分数计，将Cr：17％，Co：12％，Ti：1％，Al：4％， W：0.1％，Mo：5.0％，Si：0.2％，Mn：0.4％，Nb：1％，C：0.04％，余量为Ni；加入到真空感应炉中，在真空以及氩气保护下进行合金冶炼，精炼，最终获得合金铸锭；合金铸锭中N元素质量百分含量不高于0.03％，P与S质量百分含量均不高于0.03％。
2)高温轧制：将合金铸锭以3℃/min的升温速率升至1020℃保温0.5小时后，再对合金铸锭在1120℃进行24小时的均匀化处理，然后将其外部采用厚度0.5-1.0mm的304不锈钢包套后，在Ni 3Alγ’溶解温度以上70℃进行高温轧制，每道次变形量不低于25％，总变形量不低于70％；在每道次变形完成后回炉保温15min后进行下一道次轧制。
3)热处理：将轧制后的合金随炉升温至γ’溶解温度以上40℃范围内保温5小时，完成后空冷至室温；随后将合金加热至γ’溶解温度以下350℃范围内保温3小时后空冷，最后加热至γ’溶解温度以下250℃范围内保温1小时后空冷。
实施例5
1)合金冶炼：按质量百分数计，将Cr：19％，Co：10％，Ti：0.5％，Al：4.5％，Mo：1.0％，Si：0.5％，Mn：0.5％，Nb：0.7％，C：0.05％，余量为Ni；加入到真空感应炉中，在真空以及氩气保护下进行合金冶炼，精炼，最终获得合金铸锭；合金铸锭中N元素质量百分含量不高于0.03％，P与S质量百分含量均不高于0.03％。
2)高温轧制：将合金铸锭以1℃/min的升温速率升至1000℃保温0.7小时后，再对合金铸锭在1180℃进行50小时的均匀化处理，然后将其外部采用厚度0.5-1.0mm的304不锈钢包套后，在Ni 3Alγ’溶解温度以上100℃进行高温轧制，每道次变形量不低于25％，总变形量不低于70％；在每道次变形完成后回炉保 温17min后进行下一道次轧制。
3)热处理：将轧制后的合金随炉升温至γ’溶解温度以上70℃范围内保温4小时，完成后空冷至室温；随后将合金加热至γ’溶解温度以下320℃范围内保温5小时后空冷，最后加热至γ’溶解温度以下220℃范围内保温2小时后空冷。
对比例
利用真空感应炉对合金进行熔炼，获得的合金按质量百分比包括：Cr：21％，Co：15％，Ti：1.8％，Al：4.5％，Si：0.2％，Mo：7.0％，Mn：0.2％，Nb：0.5％，C：0.04％，余量为Ni。采用真空感应炉进行合金冶炼，通入高纯氩气前确保真空度低于5*10 -3。随后采用电渣重熔工艺精炼，最终获得合金铸锭以备加工。确保合金经电渣重熔后N元素含量不高于300ppm，P、S含量不高于0.03％。
对合金在1200℃进行24小时的均匀化处理，然后将其在γ’溶解温度以上100℃进行高温轧制，总变形量70％，每道次变形量不低于25％。合金均匀化升温过程中控制其升温速率10℃/min，并在升至1020保温0.5小时后继续升温至指定温度进行均匀化处理。
将轧制后的合金随炉升温至γ’溶解温度以上70℃保温4小时，完成后空冷至室温。随后将合金加热至γ’溶解温度以下300℃保温8小时后空冷，最后加热至γ’溶解温度以下200℃保温2小时后空冷。合金经热处理后室温及850℃屈服强度分别为930MPa与586MPa，经850℃热暴露1000小时后室温及850℃拉伸屈服强度分别高于955MPa与438MPa。
图8与9为对比例热处理态及长期热暴露态的组织照片，可见其在850℃热暴露期间晶粒内部出现大量有害相，表明其组织稳定性较差。
图10为对比例合金经850℃高温烟气环境(N 2-15％CO 2-3.5％O 2-0.1％SO 2)腐 蚀500小时煤灰腐蚀后截面照片，其内侧未形成完整氧化铝层，因此失重高达0.96mg/cm 2。

Claims (8)

  1. 一种可形成复合耐蚀层的变形高温合金，其特征在于：按质量百分数计，包括：Cr：16～19％，Co：10～15％，Ti：0.5～1.5％，Al：3.5～4.5％，W：≤0.5％，Mo：≤5.0％，Si：≤0.5％，Mn：≤0.5％，Nb：0.5～1.0％，C：0.04～0.07％，余量为Ni。
  2. 一种可形成复合耐蚀层的变形高温合金的制备工艺，其特征在于，包括以下步骤：
    1)合金冶炼：按质量百分数计，将Cr：16～19％，Co：10～15％，Ti：0.5～1.5％，Al：3.5～4.5％，W：≤0.5％，Mo：≤5.0％，Si：≤0.5％，Mn：≤0.5％，Nb：0.5～1.0％，C：0.04～0.07％，余量为Ni；加入到真空感应炉中，在真空以及氩气保护下进行合金冶炼，精炼，最终获得合金铸锭；
    2)高温轧制：对合金铸锭在1160-1200℃进行24-72小时的均匀化处理，然后将其在Ni 3Alγ’溶解温度以上50-100℃进行高温轧制，每道次变形量不低于25％，总变形量不低于70％；
    3)热处理。
  3. 根据权利要求2所述的一种可形成复合耐蚀层的变形高温合金的制备工艺，其特征在于，步骤1)中，合金铸锭中N元素质量百分含量不高于0.03％，P与S质量百分含量均不高于0.03％。
  4. 根据权利要求2所述的一种可形成复合耐蚀层的变形高温合金的制备工艺，其特征在于，进行步骤2)前，将合金铸锭在950-1020℃保温0.5-1.0小时后再进行步骤2)。
  5. 根据权利要求4所述的一种可形成复合耐蚀层的变形高温合金的制备工艺，其特征在于，以不高于10℃/min的升温速率升至950-1020℃。
  6. 根据权利要求2所述的一种可形成复合耐蚀层的变形高温合金的制备工艺，其特征在于，步骤2)中，进行高温轧制时，合金铸锭外部采用厚度0.5-1.0mm的304不锈钢包套。
  7. 根据权利要求2所述的一种可形成复合耐蚀层的变形高温合金的制备工艺，其特征在于，步骤2)中，在每道次变形完成后回炉保温15-20min后进行下一道次轧制。
  8. 根据权利要求2所述的一种可形成复合耐蚀层的变形高温合金的制备工艺，其特征在于，步骤3)的具体过程为：将轧制后的合金随炉升温至γ’溶解温度以上30-70℃范围内保温3-5小时，完成后空冷至室温；随后将合金加热至γ’溶解温度以下300-350℃范围内保温3-9小时后空冷，最后加热至γ’溶解温度以下200-250℃范围内保温1-3小时后空冷。
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