1 实验方法
首先在 SRJX-1000型真空感应炉中按照名义成分(Fe-0.35C-9Mn-3.2Al)熔炼获得铸锭。经过1250℃均匀化退火2 h后，锻造成横截面积为100 mm × 30 mm的板坯试样并冷却至室温，开锻和终锻温度分别为1250和850℃。之后再将板坯试样置于1200℃保温2 h，再经过7道次轧制(开轧温度880℃，终轧温度不低于850℃)得到厚度为5 mm的热轧钢板，随后水冷至室温得到热轧淬火态实验用钢。分别对上述钢板在600、650和700℃下保温1 h后水冷至室温。为便于后文叙述，将上述临界退火态样品分别编号为IA600、IA650和IA700。
采用火花线切割沿轧制方向取样加工获得板状拉伸试样，平行段尺寸为40 mm × 12.5 mm × 4 mm (长×宽×厚)。使用CMT-5305型万能试验机测试其单轴拉伸性能，应变速率0.001 s-1。每种状态至少测试3次并取均值，以保证数据的可靠性。从各状态和不同应变试样取样进行显微组织表征。所获样品均截取自靠近标距中心处部分。
使用D/max-2500/PC型X射线衍射仪(XRD)测试不同状态样品的XRD谱。扫描范围为40°~100°，扫描速率为2°/min。根据文献[1]中所报道的方法，选取(200) γ 、(200) α 、(220) γ 、(220) α 和(311) γ 衍射峰的相对积分强度计算奥氏体含量。采用EVO-18场发射扫描电镜(SEM)对轧制面的组织特征进行SEM和电子背散射衍射(EBSD)精细表征。在液氮低温环境中用电解抛光法制备EBSD样品。所用电解液为10%HClO4 + 90%C2H5OH (体积分数)的混合溶液。电解电压、电流和时间分别为15 V、0.06 A和15 s。由于存在明显的组织不均匀，本工作所采用的电解参数无法较好地解析出偏析带，着重分析非偏析区。
2 实验结果
2.1 微观组织特征
如图1a所示，热轧淬火态中锰钢的显微组织由黑色偏析带(箭头所示，平整凸出区域)和灰白的非偏析区(红色方框对应右侧放大图)构成。初步推断，偏析带为粗晶奥氏体且富含Mn元素[15]；非偏析区主要为细小板条马氏体，Mn含量相对较低。在本工作中，具有bcc结构的相可能为铁素体或马氏体，而奥氏体为fcc结构。EBSD相图(图1b)表明，非偏析区主要由bcc相构成。而非偏析区的反极图(IPF)显示出明显的原奥氏体边界(PABs)特征，如图1c所示。综合上述结果可知，热轧淬火态中锰钢由奥氏体(偏析带)和马氏体(非偏析区) 2相构成，其中根据XRD谱(图2)计算结果估算出偏析带含量为50.6% (体积分数)。
图1
图1
热轧淬火态中锰钢的SEM像、EBSD相分布图和反极图(IPF)
Fig.1
SEM image (a), electron backscattered diffraction (EBSD) phase map (b), and inverse pole figure (IPF) (c) of hot-rolled and quenched medium-Mn steel specimen (Inset in Fig.1a shows the high magnified image; M—martensite, HAGBs—high-angle grain boundaries, LAGBs—low-angle grain boundaries, PABs—prior-austenite grain boundaries, RD—rolling direction, TD—transverse direction)
图2
图2
热轧淬火态中锰钢的XRD谱
Fig.2
XRD spectrum of hot-rolled and quenched medium-Mn steel specimen
临界退火态中锰钢样品拉伸前后的XRD谱和奥氏体体积分数如图3所示。相比于IA600试样，IA650和IA700试样的(111) γ 和(200) γ 衍射峰强度相对更强。3种临界退火态中锰钢拉伸前的RA体积分数分别为61.2% (IA600)、65.9% (IA650)和71.8% (IA700)，拉伸后的RA体积分数分别为33.0% (IA600)、28.0% (IA650)和46% (IA700)。由于发生不同程度的SIMT，临界退火态实验用钢经拉伸变形后，RA体积分数显著降低。IA600、IA650和IA700试样中马氏体转变率分别为45.9%、57.7%和36.1%。
图3
图3
临界退火态中锰钢拉伸前后的XRD谱和奥氏体含量
Fig.3
XRD spectra (a) and retained austenite (RA) contents (b) of inter-critical annealed mendium-Mn steels before and after tensile deformation
图4
图4
3种临界退火态中锰钢的SEM像、EBSD相图和相应的IPF
Fig.4
SEM images (a-c), EBSD phase maps (d-f) and corresponding IPFs (g-i) of IA600 (a, d, g), IA650 (b, e, h), and IA700 (c, f, i) samples (Austenite has similar orientation in a single packet, which is roughly distinguished by packet boundaries (white dotted lines) with the adjacent one; insets in Figs.4a-c show enlarged images of corresponding non-segregation zones)
图4d~i为临界退火态中锰钢非偏析区的EBSD相图和对应IPF。值得注意的是，由XRD谱估算出的RA体积分数是在整个较大拍摄区域(偏析区+非偏析区)中的总的相占比。相应的SEM像表明，偏析区中奥氏体为主体相，而其边缘处或存在少量铁素体。明显地，由于2个区域的电解抛光参数不同，EBSD图像采集区域组织为非偏析区，因此在非偏析区进行完全的奥氏体逆转变之前，根据此方法计算出的RA含量会低于根据XRD谱计算出的RA含量。临界退火温度越高，非偏析区奥氏体逆转变程度越大，根据上述2种方法计算的RA含量均呈现出相同趋势。
在600℃下临界退火，逆奥氏体转变程度较低，奥氏体含量较少且晶粒细小，所形成的铁素体主相多呈大块状。在650℃下临界退火，奥氏体大多在铁素体边界形核[17]。可以清楚看到，IA650试样中逆转变奥氏体具有多种形态和较宽的尺寸范围(图4e)，铁素体数量减少，且形态多为条状。在700℃下临界退火，更多奥氏体沿着铁素体边界形核长大，相互连接成网状，铁素体形态为细小条状。由于晶粒形状不规则，本工作采用面积法统计平均晶粒尺寸。奥氏体平均晶粒尺寸分别为0.099 μm2 (IA600)、0.17 μm2 (IA650)和0.23 μm2 (IA700)，铁素体的平均晶粒尺寸分别为0.24 μm2 (IA600)、0.15 μm2 (IA650)和0.081 μm2 (IA700)。即随着临界退火温度提高，奥氏体平均晶粒尺寸增大，铁素体/马氏体平均晶粒尺寸减小。值得注意的是，热处理后非偏析区内的逆转变奥氏体晶粒也呈现出原奥氏体边界特征，即具有相同的晶体学取向(图4g~i)。
2.2 力学性能
图5a为临界退火态中锰钢的工程应力-应变曲线。所有试样均表现出连续屈服及均匀变形段内的塑性失稳(锯齿状应力波动)[18]。IA600试样的抗拉强度(UTS)、总延伸率(TE)和PSE分别为1150 MPa、18.4%和21.2 GPa∙%。IA700试样的UTS、TE和PSE分别为930.2 MPa、25.5%和23.8 GPa·%。IA650试样的UTS和TE分别为1180.5 MPa和61.4%，PSE高达72.5 GPa·%。图5b~d为采用相同方法获得的临界退火态中锰钢加工硬化率(WHR)曲线和真应力-应变曲线。WHR曲线大致分为3个阶段，分别用S1、S2和S3表示。文献[19,20]指出，S1阶段WHR急剧降低，主要与含有少量位错的铁素体快速变形有关，WHR在S2阶段缓慢下降，对应着铁素体变形软化和少量马氏体转变的交叠。在S3阶段，WHR快速波动，与大量马氏体相变以及铁素体和马氏体的协同变形有关[21]；IA650试样的S3阶段应变宽度达到0.2838，远高于IA600的0.0913和IA700试样的0.1119，对应着更为持久的TRIP效应。此外，IA650样品的WHR波动幅度随着真应变增加而逐渐增大。因此可初步认为，IA650样品的WHR波动与马氏体相变诱发加工硬化及后续应力松弛有关。
图5
图5
3种临界退火态中锰钢的工程应力-应变曲线和相应的加工硬化曲线
Fig.5
Engineering stress-strain curves (a) and corresponding work-hardening curves of IA600 (b), IA650 (c), and IA700 (d) samples (WHR—work hardening rate; S1-S3 show the different stages)
3 分析与讨论
3.1 拉伸过程中的组织演变
结合XRD谱(图3)可知，临界退火态中锰钢的力学性能与SIMT效应密切相关，即取决于形变过程中马氏体转变率。图6显示了IA650样品不同工程应变下的相分布和晶粒取向分布图。在10%工程应变下，部分奥氏体晶粒的内部取向发生明显变化(图6b中黑色箭头所指)。图6a中插图为选取某个奥氏体晶粒绘制的局部取向差(KAM)分布。在取向交替处，奥氏体的KAM值明显更高，反映出因变形导致的晶内取向旋转，遵循着特定的滑移面和滑移方向，证实可以改善延展性[22]。同时，部分bcc相晶粒也发生了取向旋转。在20%应变下，奥氏体含量明显降低，且依然可以清晰看到包及包界。经40%工程应变后，SIMT导致奥氏体含量进一步降低，某些bcc粗晶内部取向也发生明显变化(如图6e和f中的黑色圆圈表示)。值得注意的是，该应变下形成了如图6f所示条状细晶区(bundle zone)以及碎块状晶区(fragment zone)，其用白色虚线加以划分。条状细晶区由沿拉伸方向的条状bcc相晶粒及内部奥氏体组成，而碎块状晶区由形状不规则、尺寸各异的bcc相晶粒及内部奥氏体组成。碎块晶区存在大量亚晶界，其演化与位错活动强烈相关。当工程应变达到60%，上述亚结构交替分布的特征更为明显，如图6h所示。更大范围的SEM像(图7)清晰显示出2种亚结构的形貌差异。在更高应变下，条状细晶区的宽度更窄，条状bcc晶粒更加细小，而碎块状晶区包含更多的不规则细小bcc相晶粒，并具有更多连续分布的亚晶界。亚晶界在更大的应变条件下又将进一步演化成大角度晶界，引起晶粒的碎化，这与冷轧中锰钢[22]相似，即包被拉长并转动到择优取向。可以认为，有利取向包内晶粒旋转获得条状细晶区，而不利取向包内晶粒逐渐形成碎块状晶区。因此在大应变下相邻包倾向形成2种亚结构的交替分布形态。此外，经60%工程应变后，碎块状晶区内仍含有少量奥氏体(图6g和h)，表明SIMT并非同步完成。由此推断，IA650样品临近断裂时，2种亚结构交替分布的现象广泛存在。
图6
图6
IA650样品在不同应变下的相分布图和相应的IPF
Fig.6
EBSD phase maps (a, c, e, g) and corresponding IPFs (b, d, f, h) of IA650 sample with different strains (Inset in Fig.6a is the kernel average misorientation (KAM) map of austenite grain in the selected area; the arrows in Fig.6b show that the internal orientation of some austenite grains changed significantly; the black solid circles in Figs.6e and f represent the orientation change inside a large α grain, while the dotted lines in Figs.6f and h represent the boundaries between bundle zone (I) and fragment zone (II)) (a, b) 10% strain (c, d) 20% strain (e, f) 40% strain (g, h) 60% strain
图7
图7
IA650样品非偏析区在拉伸后的SEM像
Fig.7
SEM image of bundle zone (I) and fragment zone (II) in IA650 sample after tensile deformation
取向旋转是晶粒转动从而利于变形的内有方式[22]。在变形初期，载荷配分至包内奥氏体和铁素体晶粒。相对更易整体取向旋转的细晶奥氏体，多系滑移导致某些RA粗晶内部取向陆续改变，从而容易在取向交替处产生应变集中(图6b)[22]。有关拉伸过程中的原位EBSD研究[23]表明，不同位置的奥氏体具有不同的相变行为：(1) 位于马氏体晶间的薄膜状奥氏体在变形初期发生取向旋转，在后期转变为马氏体；(2) 位于三叉界面或更多晶粒连接处的奥氏体因晶界处局部应力集中而在变形前期转变；(3) 而对于完全镶嵌在单一铁素体晶粒内的奥氏体晶粒，马氏体转变行为与薄膜奥氏体相近。在本工作中，细小条状奥氏体晶粒内部并未出现取向交替现象。当应变累计达到相变临界切应变时，粗晶奥氏体势必转变为马氏体。而对细晶奥氏体而言，取向旋转与发生相变的先后顺序更多取决于所处的位置。考虑到临界剪切应变差异，各RA晶粒的应变积累量不尽相同。粗晶奥氏体因力学稳定性较低，变形初期即转变为马氏体。此外，块状奥氏体因多晶变形，在相界面产生较高的剪切应力和应变，为马氏体相变提供了形核驱动力，并由相界向内部逐步转变，导致RA晶粒细化。
3.2 强韧机制
图8
图8
IA600和IA700样品拉断后的相分布图和IPF
Fig.8
EBSD phase maps (a, c) and IPFs (b, d) of IA600 (a, b) and IA700 (c, d) samples after tensile deformation
图9
图9
IA650中样品中偏析带在拉伸前后的SEM像
Fig.9
SEM images of segregation band in IA650 sample before (a) and after (b) tension (The arrows and circle represent the segregation area before and after tension, respectively)
为阐明IA650样品的强化机制，本工作做了如下假设：(1) 经临界退火(600~700℃)处理，偏析带完全为奥氏体组织；(2) 拉伸变形过程中非偏析区奥氏体均完全转变。根据XRD谱和EBSD结果，估算出偏析带奥氏体在拉伸过程的马氏体转化率，如图10所示。IA650样品的偏析带奥氏体具有最高的马氏体转化率(49%)，这可能是使得IA650样品具有最优综合力学性能的重要原因。若非偏析区RA稳定性过低而快速转变，导致α'/γ界面易萌生微裂纹，而后扩展直至断裂失效，此时偏析区的RA来不及充分转变，引起其马氏体转化率较低。因此，合理调控不同区域的奥氏体稳定性，通过更加充分的非连续TRIP效应，可以获得更优的强塑匹配。
图10
图10
计算的3种临界退火态中锰钢中偏析带在拉伸后的马氏体转变率
Fig.10
Calculated martensite transformation ratios of segregation band in IA600, IA650, and IA700 samples
结合XRD谱(图3b)可知，临界退火态中锰钢的力学性能与SIMT密切相关，即取决于形变过程中马氏体转变率，其中IA650试样表现出更加充分的TRIP效应。虽然IA700试样具有最高的RA体积分数(71.8%)，但是拉伸断裂后仍保留了相当数量奥氏体(46%)。
碳化物的析出和溶解对调控奥氏体稳定性具有重要影响。碳化物析出先于奥氏体逆转变，延缓了奥氏体形核[24]。研究[25]表明，逆奥氏体转变主要发生在加热阶段，而非保温阶段。根据平衡相图，由于温度较低，IA600样品中仍然存在未溶解碳化物，“贫碳”的逆转变的细小奥氏体晶粒的力学稳定性相对较低。在形变初期，逆转变奥氏体快速转变为马氏体，偏析带奥氏体未充分转变，对PSE的贡献度较低。此外，在变形过程中IA600样品包结构未发生取向旋转，bcc粗晶内应变累积程度较低，未引起LAGBs发生明显演变，而微裂纹倾向在粗大bcc相(图8a)的晶界处产生[22]。由于粗晶临近晶粒的尺寸较小(图8a)，具有高密度晶界，提供较大的扩展路径和阻力[26]，这可能导致微裂纹沿着粗晶晶界扩展并最终引发断裂。此外，碳化物的存在也可能是IA600试样过早断裂失效的诱因。
而对IA700试样而言，尽管碳化物得到最大程度的溶解，但逆转变奥氏体含量增加且相互连接成网状结构(图4f)，因此平均C含量降低和晶粒粗化带来了逆转变奥氏体力学稳定性的双重叠加降低。奥氏体和铁素体的形态分别为网状和细条(图4f)。晶粒碎化只能发生在转变形成的网状马氏体或条状铁素体中。前已述及，条状细晶区和碎块状晶区的交替分布依赖于相邻包。虽然网状奥氏体具有良好的塑性变形能力，可促进特征区交替分布的形成，但却无法通过连续SIMT缓解应变集中。在变形过程中，少量低稳定性的RA快速转变，而又缺乏大尺寸bcc晶粒来承受应变，从而导致条状bcc晶粒的快速破碎(图8c)。虽然高密度亚晶界阻碍了位错运动和裂纹扩展[27]，但缺乏SIMT来松弛应力，难以阻止断裂失效的产生。
本实验条件下，偏析区RA具有较高的力学稳定性和良好的塑性变形能力，而非偏析区RA力学稳定性较低，转变为马氏体后，晶界和亚晶界易成为裂纹萌生核心。通过合理控制逆奥氏体转变程度，使偏析带和非偏析区奥氏体稳定性具有合理梯度分布，可以避免非偏析区奥氏体过早发生SIMT，避免微裂纹过早地在非偏析区产生，从而促进偏析带内TRIP效应充分产生，由此获得了优异的强塑匹配。
4 结论
(1) Fe-0.35C-9Mn-3.2Al热轧淬火中锰钢经适宜的临界退火处理(650℃、1 h后水淬，IA650)，获得了高达72.5 GPa·%的PSE，主要强韧机制为偏析带奥氏体和逆转变奥氏体的协同TRIP效应。
(2) 碳化物的析出和溶解对非偏析区逆奥氏体转变具有重要影响，经650℃临界退火处理，逆转变奥氏体具有多种形态和广泛的晶粒尺寸分布，形成了奥氏体稳定性的梯度分布。
(3) IA650试样形变过程中，条状细晶区和碎块状晶区亚结构交替分布和协同变形，通过取向旋转，抑制并延缓了逆转变奥氏体的SIMT；促进了偏析区RA的大量转变，最终获得了优异的强度和韧性匹配。
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Twinning-induced plasticity (TWIP) steels
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Smaller is less stable: Size effects on twinning vs. transformation of reverted austenite in TRIP-maraging steels
[J].
Microstructural evolution and mechanical properties of Ni-containing light-weight medium-Mn TRIP steel processed by intercritical annealing
[J].
Role of stress-assisted martensite in the design of strong ultrafine-grained duplex steels
[J].
Spectral TRIP enables ductile 1.1 GPa martensite
[J].
The morphology and crystallography of lath martensite in alloy steels
[J].
The morphology and crystallography of lath martensite in Fe-C alloys
[J].
Influence of refined hierarchical martensitic microstructures on yield strength and impact toughness of ultra-high strength stainless steel
[J].
The hierarchical martensitic features in ultra-high strength stainless steel (UHSSS), including the prior austenite grains, martensite packets, blocks and laths with the descending size, were refined to various extents by employing different thermomechanical processes and then carefully characterized. Their relation to yield strength and impact toughness was analyzed. We conclude that the refinement of martensitic structures could lead to the significant increase of yield strength, which follows the Hall-Petch relation with the effect grain size defined by high angle boundaries (HABs). Impact toughness of UHSSS depends on the frequency and capability for retained austenite (RA) grains at both HABs and martensite lath boundaries to trap the propagating cracks via strain-induced transformation, in which the film-like RA grains at lath boundaries appear to make the greater contribution.
Alloying effects on microstructure formation of dual phase steels
[J].
Novel medium-Mn (austenite + martensite) duplex hot-rolled steel achieving 1.6 GPa strength with 20% ductility by Mn-segregation-induced TRIP mechanism
[J].
Achieving excellent strength-ductility and impact toughness combination by cyclic quenching in medium Mn TRIP-aided steel
[J].
On the origin of dynamic strain aging in twinning-induced plasticity steels
[J].
Microstructural evolution and tensile properties of 70 GPa·% grade strong and ductile hot-rolled 6Mn steel treated by intercritical annealing
[J].
Influence of annealing temperature on microstructure and tensile property of cold-rolled Fe-0.2C-11Mn-6Al steel
[J].
Mechanical properties and austenite stability in hot-rolled 0.2C-1.6/3.2Al-6Mn-Fe TRIP steel
[J].
Multi-scale characterization of austenite reversion and martensite recovery in a cold-rolled medium-Mn steel
[J].A medium-Mn steel (Fe-12Mn-3Al-0.05C wt%) was designed using Thermo-Calc (R) simulations to balance the fraction and stacking fault energy of reverted austenite. lntercritical annealing for 0.5, 8 and 48 h was carried out at 585 degrees C to investigate the microstructural evolution. X-ray diffraction (XRD), electron backscatter diffraction (EBSD), 3-dimensional EBSD, energy-dispersive spectroscopy via scanning transmission electron microscopy (STEM-EDS) and atom probe tomography (APT) enable characterization of phase fraction, grain area, grain morphology and alloy partitioning. An increase in annealing time from 0.5 h to 48 h increases the amount of ultrafine-grained (UFG) reverted austenite from 3 to 40 vol %. EBSD and TEM reveal multiple morphologies of UFG austenite (equiaxed, rod-like and plate-like). In addition, most of the remaining microstructure consists of recovered alpha'-martensite that resembles the cold-rolled state, as well as a relatively small fraction of UFG ferrite (i.e., only a small amount of martensite recrystallization occurs). Multi-scale characterization results show that the location within the cold-rolled microstructure has a strong influence on boundary mobility and grain morphology during austenite reversion. Results from APT reveal Mn-decoration of dislocation networks and low-angle lath boundaries in the recovered alpha'-martensite, but an absence of Mn-decoration of defects in the vicinity of austenite grains, thereby promoting recovery. STEM-EDS and APT reveal Mn depletion zones in the ferrite/recovered alpha'-martensite near austenite boundaries, whereas gradients of C and Mn co-partitioning are visible within some of the austenite grains after annealing for 0.5 h. Relatively flat C enriched austenite boundaries are present even after 8 h of annealing and indicate certain boundaries possess low mobility. At later stages the growth of austenite followed the local equilibrium (LE) model such that the driving force between two equilibrium phases moves the mobile interface, as confirmed by DICTRA simulations (a Thermo-Calc (R) diffusion module). The sequence of austenite reversion is: (i) formation of Mn- and C-enriched face-centered-cubic nuclei from decorated dislocations and/or particles; (ii) co-partitioning of Mn and C and (iii) growth of austenite controlled by the LE mode. (C) 2019 Acta Materialia Inc. Published by Elsevier Ltd.
Strain partitioning and strain localization in medium manganese steels measured by in situ microscopic digital image correlation
[J].
Influence of carbide precipitation and dissolution on the microstructure of ultra-fine-grained intercritically annealed medium manganese steel
[J].
Austenite formation and mechanical behavior of a novel TRIP-assisted steel with ferrite/martensite initial structure
[J].
Void growth by dislocation emission
[J].
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