随着特种车辆车速的不断提升和路况复杂程度的增大，作为特种车辆的关键结构件，扭力轴的服役环境由传统的低应力、低强度和低缓冲转变为高应力、高强度和高缓冲。提升扭力轴的综合力学性能是特种车辆整车轻量化、结构件长寿命化的关键。传统扭力轴的热处理工艺选择奥氏体化后进行中温回火(460 ℃)的调质处理，其显微组织为回火屈氏体组织[1]。周铮[2]发现，某扭力轴实验过程中多次发生脆性断裂，断口的微观形貌显示为沿晶断裂，这是扭力轴的中温回火温度处于第二类回火脆性范围导致的。为了提高扭力轴的力学性能，部分学者采用预扭转工艺，即热处理后在其工作方向扭转一个角度。预扭转使扭力轴的表层产生一定厚度的塑性变形，在表面产生一定的残余应力，使扭力轴表面的总应力下降，从而提高材料的疲劳寿命[3]。周国峰等[4]通过弹塑性有限元方法模拟了扭力轴的预扭转过程，发现预扭转角度为52°时，试样表面产生了厚度为8 mm的塑性变形层。预扭转使试样屈服强度提高，并且表面产生了200 MPa的残余应力，从而扩大了扭力轴弹性工作应力区间(约891 MPa)，并提高了承载能力。另外，形变热处理[5,6]、冷变形[7,8]、表面强化[9~11]等方法已经成为强化结构材料的重要手段。表面滚压强化可以使扭力轴的表层组织发生加工硬化并细化晶粒，降低表面粗糙度，形成残余压应力层。这种残余压应力层在提高表面强度和硬度的同时，能够有效阻止裂纹的形成和扩展，从而提高耐磨性能和抗疲劳性能[12,13]。虽然预扭转工艺和表面滚压等表面强化方法可以增加扭力轴的强度和抗疲劳性能，但是这些方法增加了一步预处理工艺，并且增加了使用成本。
近年来，由于热处理工艺简单、成本廉价，采用高温奥氏体化和低温回火(温度200~240 ℃)的2 GPa超高强钢成为扭力轴的候选材料之一[14]。李保成等[15]研究了回火温度对扭力轴力学性能的影响，发现随着回火温度降低，材料的屈服强度和抗拉强度缓慢增加，而冲击韧性则先减小后增加，在回火温度为240 ℃时冲击韧性最高，达到了60 kJ/m2。然而，相比于传统扭力轴，2 GPa超高强扭力轴的断裂韧性、疲劳等性能对表面缺陷、内部夹杂物和未溶解大尺寸碳化物更为敏感。随着冶金质量和机械加工能力的逐步提升，特种装备用超高强钢中未溶解碳化物的调控成为提升扭力轴综合力学性能的主要方式。然而，关于超高强扭力轴中未溶解碳化物对其综合力学性能影响的研究仍不全面。此外，扭力轴中含有微量的V (质量分数为0.11%~0.16%)，传统的扭力轴在460 ℃回火过程中有微量的纳米尺度VC析出，这将有利于提升扭力轴的屈服强度。然而，在低温回火过程中，尽管VC的析出驱动力得到了进一步增加，但是原子迁移所消耗的能量(ΔGm)大幅增加。根据非均匀形核机制[16,17]可知，析出形核速率与ΔGm负相关，所以低温回火(200~240 ℃)不具备VC的析出热力学条件。但目前关于通过热处理工艺协同调控未溶解碳化物和纳米尺度VC，进而提升扭力轴力学性能的研究鲜有报道。此外，特种车辆运行过程中扭力轴承受循环扭转应力，抗扭转疲劳性能成为开发新型特种车辆用扭力轴的重要指标。研究[18~20]表明，轴向拉压和旋转弯曲疲劳与静态拉伸力学性能密切相关。在多轴疲劳(包括拉伸、压缩和剪切)的相关研究中存在多种强度准则，这些准则通常考虑正应力(应变)与切应力(应变)的不同作用，进而获得等临近失效条件[21,22]。因此，扭转疲劳性能与静态扭转性能密切相关，但目前关于扭力轴用2 GPa超高强钢的显微组织与静态扭转性能之间的关系尚缺乏系统的研究。本工作旨在通过系统的表征和分析，研究奥氏体化温度对2 GPa超高强钢中渗碳体的溶解行为和VC的析出规律的影响，明确奥氏体温度对显微组织及拉伸力学性能的影响，通过分析扭转断裂机制揭示扭力轴用2 GPa超高强钢中显微组织与静态扭转性能之间的关系。
1 实验方法
本实验用钢为一种典型的中碳钢，其主要化学成分(质量分数，%)为：C 0.458，Mn 0.58，Si 0.31，Ni 1.43，Cr 0.87，V 0.16，Mo 0.23，Fe余量。将熔炼并连铸后的坯料加热到1200 ℃保温2 h，随后热轧成直径62 mm的棒材，最后空冷至室温。将热轧棒材加工成直径15 mm的棒状试样，并升温至800~1000 ℃保温30 min，随后油淬至室温，最后在220 ℃回火保温2 h，其热处理工艺如图1所示。根据奥氏体化温度(800、850、900、950和1000 ℃)的不同，热处理试样分别标记为Q800T、Q850T、Q900T、Q950T和Q1000T。利用具有TCFE10和MobFe5数据库的Thermo-Calc软件对实验用钢的析出相特性进行理论计算。
图1
图1
热处理工艺示意图
Fig.1
Schematic of the heat treatment process (Ac3—austenite transformation finish temperature, AC—air cooling, OQ—oil quenching)
在热轧和热处理试样上切取金相试样，经过研磨和机械抛光后，采用4%硝酸酒精溶液(体积分数，下同)进行侵蚀，随后利用Merlin场发射扫描电子显微镜(SEM)进行组织观察。通过ImageJ软件统计每个试样中的析出相尺寸，每个试样统计至少5张SEM像。研磨后的金相试样进行电解抛光，使用的电解液为12.5%高氯酸酒精溶液，电压20 V，时间20 s，并利用配置有Auger电子能谱仪(AES)和电子背散射衍射(EBSD)系统的PHI710扫描Auger纳米探针对试样中局部区域进行成分和结构分析。对试样中的典型区域进行C、Cr、V元素线扫描分析，电压和电流分别为20 kV和10 nA，为了减少碳氢化合物污染的沉积[23]，在线扫描之前先进行轻微的表面溅射。EBSD实验时加速电压为20 kV，步长80 nm。切取不同热处理状态的试样进行研磨，并使用10%高氯酸酒精溶液进行双射流电解双喷，电压32 V，温度-20 ℃。采用配备扫描透射(STEM)模式的Tecnai G2 F20场发射透射电子显微镜(TEM)观察试样中的纳米尺度析出相。
使用带有视频引伸计的MTS Exceed E45力学试验机测量所有试样的静态力学性能，拉伸速率为2 mm/min，每种热处理工艺测试2个拉伸试样，并求平均值。拉伸试样为沿轧制方向加工的标距直径为3 mm、标距长度15 mm的圆柱形拉伸试样。扭转试样为沿轧制方向加工的标距直径为10 mm、标距长度70 mm的圆柱形试样。静态扭转实验时，试样在达到屈服强度之前，扭转速率为15°/min，达到屈服强度之后，扭转速率为30°/min。采用Merlin SEM观察扭转失效试样的断口形貌。
2 实验结果
图2
图2
热轧试样显微组织的SEM像及元素分析结果
Fig.2
SEM images of the hot-rolling sample (a, b), AES selected area (c) and the element profiles (d)
图3
图3
不同热处理试样显微组织的SEM像及渗碳体直径与体积分数随奥氏体化温度的变化
Fig.3
SEM images of the sample subjected to different austenitizing temperatures of 800 oC (a), 850 oC (b), 900 oC (c), 950 oC (d), and 1000 oC (e) and tempering at 220 oC (The corresponding samples are named Q800T, Q850T, Q900T, Q950T, and Q1000T), and change curves of cementite diameter and volume fraction with austenitizing temperature (f) (Arrows in Figs.3a-c represent the cementites)
为了明确奥氏体化温度对析出相成分的影响，对析出相进行线扫描分析，结果如图4所示。Q800T试样中球状渗碳体主要包含C和Cr元素，及少量的V元素，这与热轧试样中渗碳体的成分分布相近(图4a)，均为合金渗碳体(Fe, Cr, V)3C，说明在800 ℃奥氏体化热处理工艺条件下得到的渗碳体具有明显的遗传效应。另外，碳化物与基体之间的成分分布差异明显。Q850T试样中球状渗碳体主要包含C和Cr元素，渗碳体中V的含量与基体组织中的V含量相近(图4b)。与Q800T和Q850T试样不同，当奥氏体化温度升高至900 ℃时，Q900T试样中的球状渗碳体是以C元素为主的Fe3C，并且C与基体组织之间存在缓慢过渡的现象，而Cr和V与基体组织中的成分相同(图4c)。因此，随着奥氏体化温度的升高，渗碳体由合金渗碳体转变为Fe3C渗碳体，其合金元素逐渐由C、Cr和V元素富集转变为C元素富集，这一过程伴随着渗碳体的溶解。图4d为Q900T试样中纳米析出相元素含量的分布。可以看出，不同于球状渗碳体，纳米析出相以V元素为主，含有少量Cr元素。
图4
图4
Q800T、Q850T和Q900T试样析出相的元素分布图
Fig.4
Element profiles of precipitate in the Q800T (a), Q850T (b), and Q900T (c, d) samples
(a-c) element distributions of cementite (d) element distributions of VC
图5
图5
Q800T、Q850T、Q900T和Q950T试样的TEM像和STEM-EDS元素面分布图
Fig.5
TEM images and corresponding STEM-EDS mappings of the Q800T (a), Q850T (b), Q900T (c), and Q950T (d) samples
奥氏体化温度不仅影响析出相的类型和尺寸，还显著影响奥氏体的晶粒尺寸，进而影响马氏体板条尺寸。图6a~e为Q800T、Q850T、Q900T、Q950T和Q1000T试样的取向成像图(orientation image mapping)。图6f为各试样马氏体块尺寸随奥氏体化温度变化的统计图。由于Q800T和Q850T试样中含有大量的渗碳体和VC析出相，阻碍了奥氏体长大，进而细化了马氏体板条尺寸(约530 nm)。随着奥氏体化温度升高至900 ℃，渗碳体含量减少，导致马氏体板条尺寸迅速增加到约562 nm；进一步升高奥氏体化温度，马氏体板条尺寸维持在560~580 nm范围内。除了马氏体板条尺寸外，组织中大角度晶界是影响钢力学性能的重要因素。图7a~e为Q800T、Q850T、Q900T、Q950T和Q1000T试样中小角度晶界(蓝色线表示2°~15°晶界)和大角度晶界(红色线表示15°~45°晶界，黑色线表示> 45°晶界)分布图。图7f为所有试样的大角度晶界所占比例随奥氏体化温度的变化。可见，Q800T试样的大角度晶界比例高达81.5%，这是由于奥氏体化温度较低和未溶解碳化物共同作用下抑制了奥氏体晶粒的长大，进而获得了更细小的马氏体块尺寸。当奥氏体化温度升高到850~900 ℃，奥氏体晶粒尺寸增加，组织中晶界、块界、板条界等大角度晶界比例减小。奥氏体化温度进一步升高到950~1000 ℃时，渗碳体和VC完全溶解导致奥氏体晶粒进一步增加，从而减小了组织中大角度晶界的比例。
图6
图6
Q800T、Q850T、Q900T、Q950T和Q1000T试样的EBSD取向成像图和马氏体块尺寸随奥氏体化温度变化的统计图
Fig.6
EBSD orientation imaging mappings of the Q800T (a), Q850T (b), Q900T (c), Q950T (d), and Q1000T (e) samples and statistical diagram of martensitic block sizes (f)
图7
图7
Q800T、Q850T、Q900T、Q950T和Q1000T试样的晶界分布图和高角度晶界比例统计图
Fig.7
Grain boundary distributions in the Q800T (a), Q850T (b), Q900T (c), Q950T (d), and Q1000T (e) samples (blue line: 2°-15°, red line: 15°-45°, yellow line: > 45°) and statistical diagram of high angle grain boundary ratios (f)
图8为不同热处理试样的单轴拉伸力学性能。Q800T试样具有较高的屈服强度(1617 MPa)，随着奥氏体化温度的升高，试样的屈服强度先缓慢降低；当奥氏体化温度达到900 ℃后，屈服强度迅速降低。与此不同，Q800T试样的抗拉强度最低(1966 MPa)，随着奥氏体化温度升高，实验用钢的抗拉强度先迅速升高，后略微降低并趋于稳定，Q850T试样的抗拉强度最高，达到了2062 MPa，屈服强度为1580 MPa。实验用钢的均匀延伸率与抗拉强度的变化趋势基本一致，而总延伸率则随着奥氏体化温度的升高先大幅增加后逐渐降低，其中，Q850T试样的均匀延伸率和总延伸率最高，分别为8.4%和12.7%。
图8
图8
不同热处理试样的拉伸力学性能
Fig.8
Tensile mechanical properties of the samples with different heat treatments
(a) strength (b) elongation
图9为Q800T、Q850T、Q900T、Q950T和Q1000T试样的扭转剪切应变-扭转强度曲线。所有试样在扭转过程中均呈现连续屈服现象(本工作采用剪切应变为0.3%时对应的扭转强度表示扭转屈服强度)。表1列出了试样的扭转性能，包括剪切模量、扭转屈服强度和抗扭强度。Q800T和Q850T试样具有最高的剪切弹性模量(82.6 GPa)，随着奥氏体化温度升高，剪切弹性模量逐渐减小，最终Q1000T试样的剪切弹性模量降至80.6 GPa。钢的扭转屈服强度随奥氏体化温度的升高先增加后减小，Q850T试样的扭转屈服强度最高，达到了1170 MPa。而抗扭强度随着奥氏体化温度的升高先增加后减小，Q950T试样的抗扭强度最高，达到了1710 MPa。因此，Q850T试样具有优异的综合抗拉和抗扭性能。
图9
图9
不同热处理试样的扭转剪切应力-应变曲线
Fig.9
Torsional shear stress-strain curves of the samples with different heat treatments (Fig.9b is the enlarged view of Fig.9a)
表1 不同热处理试样的扭转性能
Table 1
| T / oC | G / GPa | τp0.3 / MPa | τm / MPa |
|---|---|---|---|
| 800 | 82.6 | 1160 | 1610 |
| 850 | 82.6 | 1170 | 1680 |
| 900 | 81.1 | 1150 | 1690 |
| 950 | 78.8 | 1130 | 1710 |
| 1000 | 80.6 | 1130 | 1690 |
Note:T—austenitizing temperature, G—shear modulus, τp0.3—yield strength at 0.3% shear strain, τm—torsion strength
3 分析与讨论
式中，
图10
图10
由Thermo-Calc计算的VC驱动力随奥氏体化温度的变化
Fig.10
Driving force of VC varies with the austenitizing temperature calculated by Thermo-Calc
Q800T和Q850T试样中高体积分数的析出相(图3a和b)和细小的马氏体板条尺寸(图6)导致较高的屈服强度，进而使其具有较高的剪切模量和扭断屈服强度。为了进一步明确显微组织与扭转性能之间的关系，采用SEM观察了Q800T、Q850T和Q1000T试样扭转断裂后的断口形貌，如图11所示。可见，3个扭转断裂试样的基本特征均为平断断口(图11a、d和g)。Q800T试样边部呈剪切流变断裂形态，这是因为试样承受扭转载荷时，其横截面的应力分布是不均匀的，试样边部的切应力最大。在最大的剪切应力作用下，Q800T试样中未溶解的大尺寸、难剪切的纳米析出物极易引发应力集中，从而导致裂纹过早萌生[28]。与Q800T试样(图11b)不同，Q850T和Q1000T试样边部存在剪切塑性变形区(图11e和h)，其厚度分别为270和374 μm。由于变形前产生塑性变形，剪切塑性变形区中存在塑变滑移、塑性擦伤及拉长韧窝等特征。此外，Q800T和Q850T试样的剪切变形区断口特征以更多平坦区域为主，这些区域是因为沿剪切变形而变弱的滑移面发生断裂导致的。而Q1000T试样的剪切变形区是由浅且被拉长的抛物线形的韧窝组织组成的，并且Q1000T试样中大量的二次裂纹(图11g中箭头所示)可以在裂纹扩展过程中偏转裂纹路径和延缓主裂纹扩展。因此，具有较大尺寸马氏体板条的Q900T、Q950T和Q1000T具有较高的抗剪切强度。
图11
图11
Q800T、Q850T和Q1000T试样扭转断口形貌的SEM像
Fig.11
SEM images of torsional fracture surfaces of the Q800T (a-c), Q850T (d-f), and Q1000T (g-i) samples
(a, d, g) overviews of the fracture surfaces (b, c, e, f, h, i) crack propagation regimes
4 结论
(1) 实验用钢热轧后显微组织为珠光体和少量铁素体。在奥氏体化过程中片层状渗碳体发生球化和溶解，其球化机制以非连续辅助机制为主导。随着奥氏体化温度升高，由(Fe, Cr, V)3C合金渗碳体逐渐转变为Fe3C渗碳体，直至950 ℃渗碳体完全溶解。与渗碳体相似，VC中含有Cr元素，随着奥氏体化温度的升高，VC中的Cr含量减少，950 ℃时VC不再析出。
(2) 奥氏体化过程中渗碳体的球化和纳米VC的析出阻碍了奥氏体晶粒长大，进而细化了马氏体板条尺寸。在析出强化和细晶强化的协同作用下，Q850T试样达到了最佳的综合拉伸力学性能，其屈服强度和抗拉强度分别为1580和2062 MPa，均匀延伸率和总延伸率分别为8.4%和12.7%。
(3) 由于具有大量未溶解渗碳体和细小马氏体板条尺寸，Q800T和Q850T试样的扭转屈服强度和剪切模量最佳，但Q800T试样的抗剪切强度最低。这是因为未溶解的大尺寸碳化物易引发应力集中，导致裂纹过早萌生。马氏体板条更粗大的Q950T和Q1000T试样具有更高的剪切应变和更大的扭转强度，这归因于Q950T和Q1000T试样更大的剪切塑性变形区和剪切韧性断裂主导的失效机制。
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高温淬火对扭力轴强韧性的影响
[J].
Numerical simulation of pre-twist process and limited carrying capacity of a tank torsional axis
[J].
坦克扭力轴的预扭处理及极限承载的数值模拟
[J].
Controllable selection of martensitic variant enables concurrent enhancement of strength and ductility in a low-carbon steel
[J].
Ultrastrong low-carbon nanosteel produced by heterostructure and interstitial mediated warm rolling
[J].
Torsion performance of pearlitic steel wires: Effects of morphology and crystallinity of cementite
[J].
Achieving unprecedented yield strength of 2.2 GPa with high ductility in formed parts using strain-aging
[J].
Effect of helical deformation on fatigue life of torsion shaft by rolling
[J].
滚压加工形成螺旋形变对扭力轴疲劳寿命的影响
[J].
A small roller elliptical low-frequency deep rolling machine tool for high-strength steel torque shaft and its usage method
[P].
一种高强钢扭力轴的小滚轮椭圆低频深滚机床及使用方法
[P].
A phase-transformation-strengthened surface layer on Fe-20Mn-3Al-3Si steel fabricated by mechanical grinding
[J].
In the last decade, transformation induced plasticity (TRIP) aided steels and twinning induced plasticity (TWIP) aided steels have attracted a great deal of attention due to their high strength and exceptional ductility at room temperature. In the present work, a TRIP/TWIP austenite steel with the composition as Fe-20Mn-3Al-3Si has been investigated by means of surface mechanical grinding treatment (SMGT) since it is an alloy with low stacking fault energy. The e-M and/or a-M transformations can easily occur in the steel when deformed at room temperature. It is shown that a gradient phase-transformation-strengthened (PTS) surface layer can be formed on bulk Fe-20Mn-3Al-3Si steel by SMGT at room temperature, due to the martensitic transformation of γ→ε→α on its surface. The formation of martensitic phases is dependent on the strain applied and the orientation of grains. The more the {111} or the {110} planes of grains are parallel to the specimen surface plane, the easier the martensitic laths are to be formed in the gains. The total thickness of the PTS surface layer, which formed in the process of turning at room temperature, can be more than 400 μm, but its microstructure and hardness change with depth. The top PTS layer consists of nanometer sized grains while the sublayer contains a great many of martensite laths mixed with deformation twins. With increasing of the depth, the numbers of laths decrease. Correspondingly, the micro-hardness continuously decreases with the depth from 450 HV to about 220 HV (the hardness of the matrix). The formed PTS layer has a good thermodynamic stability, so that its microstructure and hardness almost do not change after annealed at 400 ℃ for 1 h.
机械磨擦制备的Fe-20Mn-3Al-3Si钢表面相变强化层
[J].
Surface rolling FE model for numerical simulation
[J].
滚压有限元模型数值模拟
[J].
The effects of deep rolling and shot peening on fretting fatigue resistance of Aluminum-7075-T6
[J].
The discuss for temper-temperature of steel 45CrNiMoV
[J].
45CrNiMoV钢回火温度讨论
[J].
Correlation between microstructure and properties of new heat-resistant alloy SP2215
[J].
新型耐热合金SP2215组织与性能的关联性
[J].
An optimization criterion for fatigue strength of metallic materials
[J].
High cycle fatigue S-N curve prediction of steels based on transfer learning guided long short term memory network
[J].
Crack branching and deflection in AISI 4340 steel under cyclic torsional loading
[J].
Electron beam-induced sample contamination in the SEM
[J].
Phase field study of the diffusional paths in pearlite-austenite transformation
[J].In the development of advanced steel, accurate and detailed knowledge about the kinetics of phase transformations and microstructure formation is critical. The critical issue in pearlite-austenite transformation is the consideration of diffusional paths of the alloy element. Simulation has been an available method to study the diffusion of alloy elements and the migration rate of the phase boundary in the complex morphological evolution of austenite growth. The isothermal pearlite-austenite transformations at 720 and 740oC in Fe-0.6%C-2%Mn (mass fraction) alloy were studied by phase-field methods based on MICRESS. At different temperatures, the effects of diffusional paths on the austenite transformation were discussed. To achieve a semiquantitative verification of the simulated results, the migration rates of the austenite/pearlite boundary at 720 and 740oC were estimated from the experimental kinetics curves by fitting the JMA equation. By measuring the Mn profile in austenite, the modes of the austenization at 720 and 740oC can be verified as partitioned local equilibrium (PLE) and non-partitioned local equilibrium (NPLE) modes. The heterogeneous distribution of Mn in austenite at 740oC can be observed with STEM-EDS. However, the homogenous distribution of Mn can be found near the pearlite/austenite boundary in austenite at 720oC. The cases considering the γ, α, and interface-diffusional paths were simulated by phase-field methods to compare with the migration rates of the austenite/pearlite boundary. Because carbon is an interstitial element in steel and has an interstitial diffusional mechanism, it can be speculated reasonably that the diffusion of C mainly proceeded in austenite and ferrite through the considerations of the atom-size of carbon and the experimental results. Phase-field methods were used to study Mn diffusion in the lamellar pearlite-austenite transformation. With the analysis of the experimental estimations, the interface-diffusional path was observed as the dominant path for the Mn diffusion. It is because the Mn atoms have greater diffusivity in interfaces than in γ or α-diffusional paths. Furthermore, the diffusional activation energy is closely related to the diffusivity of Mn at the interface. Moreover, compared with the γ-diffusional path, the diffusional flux of Mn in ferrite is much larger than that in austenite. Thus, it can be concluded that the contribution of the α-diffusional path to the migration rate of the pearlite/austenite boundary is larger than that of the γ-diffusional path. As a result, considering the α-diffusional path in the thermodynamics analysis under NPLE mode makes more sense. However, ignoring the interface- and α-diffusional path, which is different from the traditional cognition in PLE mode, will result in a magnitude error for the thermodynamics analysis.
珠光体-奥氏体相变中扩散通道的相场法研究
[J].
Effect of austenitizing temperature on microstructures and mechanical properties of Si-Mn hot-rolled plate after quenching and partitioning treatment
[J].The production of quenching and partitioning (Q&P) steel using hot-rolled steel instead of cold-rolled steel can significantly reduce the manufacturing process time and cost. However, the initial microstructures of hot-rolled and cold-rolled steels are different, which affect the microstructures and mechanical properties of Q&P steel. Because most studies used Q&P steel prepared from cold-rolled steel, the microstructures and mechanical properties of Q&P steel prepared from hot-rolled steel are unclear. This study examines the microstructures and mechanical properties of Q&P Si-Mn steel prepared from hot-rolled steel as a function of the austenitizing temperature. The results showed that the ferrite in the Q&P Si-Mn steel produced from the hot-rolled steel had lath-type and blocky-type morphologies. The observed ferrite morphology could influence the morphology of the adjacent retained austenite. The lath-type and blocky-type ferrite surrounding the retained austenite was mainly observed as the thin lath and blocky types, respectively. The ferrite and retained austenite contents decreased with increasing austenitizing temperature. In addition, the corresponding yield and tensile strengths increased gradually with a concomitant decrease in elongation and the product of strength and elongation. When the austenitizing temperature was 810oC, the product of strength and elongation of the Q&P Si-Mn steels produced from hot-rolled steel reached 28.36 GPa·%, which was approximately 36% higher than that of Q&P980 produced industrially from cold-rolled steel. The higher product of strength and elongation of Q&P Si-Mn steel produced from hot-rolled steel may be related to the different morphologies of ferrite, which might control the morphology and stability of the adjacent retained austenite. These experimental results could provide a theoretical basis for preparing Q&P steel from hot-rolled steel instead of cold-rolled steel.
奥氏体化温度对Si-Mn钢热轧板淬火-配分处理后显微组织和力学性能的影响
[J].通过分析不同奥氏体化温度下热轧板制备的淬火-配分(Q&P)钢的显微组织和力学性能，探讨了奥氏体化温度对其显微组织与力学性能的影响。结果表明，热轧板制备Q&P Si-Mn钢中的铁素体呈带状和块状2种形貌，铁素体形貌会影响相邻区域残余奥氏体形貌，其中带状铁素体相邻区域残余奥氏体以薄带状为主，块状铁素体相邻区域残余奥氏体以块状为主。随着奥氏体化温度的升高，铁素体和残余奥氏体体积分数下降，对应的Q&P Si-Mn钢屈服强度、抗拉强度逐渐升高，延伸率及强塑积均逐渐下降。在810℃奥氏体化后，热轧板制备的Q&P Si-Mn钢强塑积达到28.36 GPa·%，相较于工业生产采用冷轧板制备的980 MPa级Q&P Si-Mn钢，强塑积提高了约36%。热轧板制备Q&P Si-Mn钢具有更高的强塑积与其组织中不同形貌铁素体能调控相邻区域残余奥氏体形貌及稳定性相关。
Quantitative three-dimensional characterization of pearlite spheroidization
[J].
Mechanisms of pearlite spheroidization: insights from 3D phase-field simulations
[J].
Strain hardening mediated by coherent nanoprecipitates in ultrahigh-strength steels
[J].
