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DP590冷轧板热处理的组织和性能霍刚;李振兴;岑一鸣;李国栋【摘要】为了加速国内双相钢的开发和应用,采用CAS-300Ⅱ模拟退火实验机,通过模拟退火实验,研究了加热速率、临界区退火温度、过时效温度、过时效时间对DP590双相钢组织性能的影响.结果表明,加热速率在5～60℃/s内增加时,屈服强度、抗拉强度均增加,延伸率、强塑积均减小；临界区退火温度在780～850℃内增加时,屈服强度、抗拉强度先减小后增加,延伸率、强塑积均增加；过时效温度在260 ～400℃内增加时,屈服强度增加,抗拉强度减小,延伸率整体呈增加趋势,屈强比增加；在280℃进行过时效,过时效时间在240～480 s内增加时,屈服强度、抗拉强度均减小,延伸率、强塑积先减小后增加.【期刊名称】《东北大学学报（自然科学版）》【年(卷),期】2013(034)007【总页数】5页(P944-947,970)【关键词】冷轧板;双相钢;热处理;显微组织;力学性能【作者】霍刚;李振兴;岑一鸣;李国栋【作者单位】东北大学轧制技术及连轧自动化国家重点实验室,辽宁沈阳110819;东北大学轧制技术及连轧自动化国家重点实验室,辽宁沈阳110819;东北大学轧制技术及连轧自动化国家重点实验室,辽宁沈阳110819;东北大学轧制技术及连轧自动化国家重点实验室,辽宁沈阳110819;本钢浦项冷轧薄板有限责任公司,辽宁本溪117000【正文语种】中文【中图分类】TG156.1双相钢是由低碳钢或低合金钢经过临界区热处理或控轧控冷获得，其组织主要由铁素体和马氏体组成[1-2].与传统的低合金高强钢相比，双相钢具有较低的屈强比，较高的初始加工硬化率、烘烤硬化值以及优良的成型性能等特点，成为一种新型的冲压用钢，并广泛应用于汽车工业[3-4].双相钢的研究起始于20世纪70年代，1968年Mcfarlan提出了关于双相钢的第一个专利.Hayami和Furukawa[5]详细阐述了双相钢的化学成分、显微组织、力学性能等.Son[6]采用等通道角度挤压法试制了超细晶粒双相钢，发现在500 ℃进行4%的应变，然后于730 ℃保温10 min后淬火，可得到性能优良的超细晶粒双相钢.韩会全等[7]研究了两相区热处理对不同初始组态钢板组织性能的影响，发现相同工艺下，初始晶粒越细，马氏体体积分数越多.Krebs等[8]对双相钢中带状组织的影响因素进行了研究，发现奥氏体化温度越低，冷却速度越小，带状组织越明显.目前，冷轧双相钢主要采用连续退火的方式生产，工艺比较成熟，但生产周期长、效率低，表面质量难以保证，并且易出现带状组织.退火工艺参数是决定双相钢组织性能的关键因素，并且其与生产效率密切相关.因此本文以国内某钢厂提供的DP590冷轧板为原料，通过模拟退火实验研究了加热速率、临界区退火温度、过时效温度、过时效时间的作用，供实际工业生产参考.1 实验材料和方法1.1 实验材料实验采用某钢厂提供的DP590冷轧板，其化学成分(质量分数/%)为：0.080C，0.479Si，1.810Mn，0.162Cr，0.014P，0.004S，0.004N，0.040Als.冷轧板的原始组织由铁素体和珠光体组成，图1为实验钢经过4%的硝酸酒精溶液腐蚀后的显微组织图片，灰白色组织为铁素体，黑色组织为珠光体.模拟退火的试样尺寸为500 mm×150 mm×1.4 mm.图1 实验钢的显微组织Fig.1 The microstructure of the steelusing in the experiment1.2 实验方法采用CAS-300Ⅱ模拟退火实验机，对加热速率、临界区退火温度、过时效温度、过时效时间分别进行了实验研究.基本工艺参数为：以30 ℃/s的速度将实验钢板加热到800 ℃，保温110 s后，以2 ℃/s的速度缓慢冷却至680 ℃，然后以35 ℃/s的速度快速冷却至280 ℃进行过时效，过时效时间为420 s，然后以5 ℃/s的速度冷却至45 ℃.在此基础上，通过改变单一的工艺参数，研究其对力学性能的影响.按照GB/T228—2002切取标距为50 mm拉伸试样，然后采用Inston系列4206-006型高速拉伸试验机测定力学性能.再切取金相试样磨制、抛光，经4%的硝酸酒精腐蚀后，分别采用LEICA Q550IW金相显微镜、ZWISS扫描电子显微镜观察其显微组织.2 实验结果与分析2.1 显微组织图2为不同加热速率下实验钢的显微组织.灰白色组织为铁素体，深灰色组织为马氏体.可以看出，加热速率在5～60 ℃/s内增加时，铁素体、马氏体晶粒均发生细化，马氏体体积分数增加.这是由于加热速率增加时，加热温度达到两相区后奥氏体形核点较多，其形核率的增加大于长大速度，奥氏体长大受到抑制，晶粒发生细化.由于组织遗传性，使得最终的铁素体、马氏体晶粒尺寸较小，并且马氏体体积分数略微增加.图2 不同加热速率下实验钢的显微组织Fig.2 The microstructures of experimentalsteels at different heating speeds(a)—5 ℃/s； (b)—15 ℃/s；(c)—30 ℃/s； (d)—60 ℃/s.图3为不同退火温度下实验钢的显微组织.可以看出，退火温度在780～850 ℃内增加时，马氏体晶粒由岛状向块状过渡，马氏体晶粒尺寸变大.此外，利用Photoshop软件统计分析，退火温度分别为780，800，830，850 ℃时，相应的马氏体体积分数分别约为30%，26%，20%，21%.图3 不同退火温度下实验钢的显微组织Fig.3 The microstructures of the tested steelsat different annealing temperatures(a)—780 ℃； (b)—800 ℃；(c)—830 ℃； (d)—850 ℃.图4为不同过时效温度下实验钢的扫描照片，颜色较浅、凸起的组织为马氏体.可以看出，过时效温度为260 ℃时，马氏体基本不分解，马氏体边界较清晰；280 ℃时，少量马氏体开始分解，边界较为模糊；400 ℃时马氏体大量分解.图4 不同过时效温度下实验钢的扫描照片Fig.4 The SEM micrographs of the tested steelsat different overaging temperatures(a)—260 ℃； (b)—280 ℃；(c)—320 ℃； (d)—400 ℃.图5为过时效温度为280 ℃时，不同过时效时间下实验钢的显微组织图片.过时效时间在240～480 s内增加时，马氏体逐渐分解.过时效时间小于300 s时，铁素体基体上存在较多的粒状M-A岛，超过420 s时，粒状M-A岛基本消失.2.2 力学性能图6显示了不同加热速率下实验钢的力学性能.加热速率在5～60 ℃/s内变化时，随加热速率增加，屈服强度、抗拉强度均增加，延伸率、强塑积均减小，屈强比在0.44～0.46范围内变化.随图5 不同过时效时间下实验钢的显微组织Fig.5 The microstructures of the experimentalsteels at different overaging time(a)—240 s； (b)—300 s；(c)—420 s； (d)—480 s.图6 不同加热速率下实验钢的力学性能Fig.6 The mechanical properties of experimentalsteels at different heating speeds加热速率的增加，铁素体、马氏体晶粒均发生细化，马氏体体积分数增加，因此屈服强度、抗拉强度均增加.加热速率较大时，铁素体中碳氮化物溶解量较小，缓慢冷却过程中铁素体析出净化作用减弱，结果，延伸率随加热速率的增加呈减小趋势. 图7显示了不同退火温度下实验钢的力学性能.退火温度在780～850 ℃内变化时，随着退火温度的增加，实验钢的屈服强度、抗拉强度先减小，然后略微增加.延伸率、强塑积均呈增加趋势.退火温度在一定范围内升高时，奥氏体体积分数增加，奥氏体中平均碳含量减小，其稳定性下降，随后缓慢冷却过程中，由于冷却速度较小，低碳奥氏体重新分解，附生铁素体体积分数增加，马氏体体积分数减小，结果屈服强度、抗拉强度都有下降趋势，屈强比、延伸率得到明显改善[9].但退火温度进一步升高时，奥氏体体积分数不断增加，最终马氏体体积分数增加，使得强度略微增加.图7 不同退火温度下实验钢的力学性能Fig.7 The mechanical properties of experimentalsteels at different annealing temperatures图8显示了不同过时效温度下实验钢的力学性能.过时效温度在260～400 ℃内变化时，随着过时效温度的增加，屈服强度增加，抗拉强度减小，延伸率整体呈增加趋势.图8 不同过时效温度下实验钢的力学性能Fig.8 The mechanical properties of experimentalsteels at different overaging temperatures过时效相当于对淬硬的马氏体进行在线回火，可改善最终的力学性能.但随着过时效温度的增加，马氏体逐渐分解，并且晶格畸变程度减小，使得抗拉强度下降.过时效温度较高时，铁素体、马氏体相界面处大量位错对消或重新排列，使得可动位错密度减小，屈服强度增加.并且在较高温度下铁素体中有碳化物或细小沉淀相析出，间隙原子扩散集聚成间隙原子团，共同钉扎位错，使得屈服强度进一步增加，甚至出现屈服平台[10].图9显示了不同过时效时间下实验钢的力学性能.过时效温度为280 ℃，过时效时间在240～480 s内变化时，随过时效时间增加，屈服强度、抗拉强度均减小，延伸率、强塑积先减小后增加.随过时效时间的增加，马氏体发生回复，马氏体内的位错密度减小，使得其硬度降低、强度下降，抗拉强度减小.而且马氏体与周围铁素体的塑性应变不相容性减小，因此马氏体对铁素体变形的阻碍作用减小，屈服强度降低，延伸率得到改善.图9 不同过时效时间下实验钢的力学性能Fig.9 The mechanical properties of experimentalsteels at different overaging time3 结论1) 加热速率在5～60 ℃/s内变化时，随加热速率增加，屈服强度、抗拉强度均增加，延伸率、强塑积均减小，屈强比在0.44～0.46范围内变化.2) 退火温度在780～850 ℃内变化时，随着退火温度的增加，实验钢屈服强度、抗拉强度先减小后增加，延伸率、强塑积均呈增加趋势.3) 过时效温度在260～400 ℃内变化时，随着过时效温度增加，屈服强度增加，抗拉强度减小，延伸率整体呈增加趋势，屈强比明显增加.4) 过时效温度为280 ℃，过时效时间在240～480 s内变化时，随过时效时间增加，屈服强度、抗拉强度均减小，延伸率、强塑积先减小后增加.参考文献：[1] Wycliffe P.Microanalysis of dual phase steels[J].Scripta Metallurgica，1984，18(4)：327-332.[2] Buzzichelli G，Anelli E.Present status and perspectives of European research in the field of advanced structural steels[J].ISIJ International，2002，42(12)：1354-1363.[3] Lanzillotto C A N，Pickering F B.Structure-property relationships in dual-phase steels[J].Metal Science，1982，16(8)：371-382.[4] Sarwar M，Priestner R.Hardenability of austenite in a dual-phase steel[J].Journal of Materials Engineering and Performance，1999，8(3)：380-384.[5] Hayami S，Furukawa T.Micro-alloying[M].New York：Union Carbide Corp，1977.[6] Son Y，Lee Y K，Park K T，et al.Ultrafine grained ferrite-martensite dual phase steels fabricated via equal channel pressing：microstructure and tensile properties[J].Acta Materialia，2005，53(11)：3125-3134. [7] 韩会全，刘彦春，张弛，等.两相区热处理对不同初始组态钢板组织性能的影响[J].东北大学学报：自然科学版，2008，29(3)：339-343.(Han Hui-quan，Liu Yan-chun，Zhang Chi，et al.The effect of heattreat ment in γ+α region on microstructures and properties of strips with different intial structures[J].Journal of Northeastern Universtity：Natural Science，2008，29(3)：339-343.)[8] Krebs B，Germain L，Hazotte A，et al.Banded structure in dual phase steels in relation with the austenite-to-ferrite transformation mechanisms[J].Journal of Materials Science，2011，46(21)：7026-7038.[9] Hüseyin A，Hawa K Z，Ceylan K.Effect of intercritical annealing parameters on dual phase behavior of commercial low-alloyedsteels[J].Journal of Iron and Steel Research，International，2010，17(4)：73-78.[10]Fonstein N，Kapustin M，Pottore N，et al.Factors that determine the level of the yield strength and the return of the yield-point elongation in low-alloy ferrite—martensite steels[J].The Physics of Metals and Metallography，2007，104(3)：315-323.。
Materials Science and Engineeringarc welding 电弧焊calcinations 煅烧casting 熔铸ceramic 陶瓷chemical properties 化学性能cold brittleness 低温脆性colour liquid crystals 彩色液晶congruent compound 合熔化合物constant-deformation tests 定变形试验Creep Strength 潜变强度crystal pattern 晶体结构data quartz fiber 数据石英光纤die casting 拉模铸造drawing & stamping 延轧Dynamics of Forging System 锻压系统动力学Edge Finish 边缘处理Engineering Materials 工程材料nano-material 纳米材料ceramic 陶瓷polymer 集合物composite material 复合材料biomaterial 生物材料semiconductor 半导体conductor 导体insulator 绝缘体synthetic fabrics 合成纤维microstructures 显微结构periodic table 周期表Equipment for Heating Processing 热处理设备Fatigue Test 疲劳测试Features of Metal 金属的特性Ferrous & Non Ferrous Metal 铁及非铁金属forging 锻造foundry 铸造High Polymer Material & Processing 高分子材料及加工Impact Test 冲击测试Intermetallic compound 金属间化物Ionic Solids 离子晶体Magnetic Transformation 磁性变态Mechanic Testing of Engineering Materials 工程材料力学性能的测定Mechanical Property of Metal 金属机械性能Metal Cutting Machine Tool 金属切削工具Metal Erosion & Protection 金属腐蚀及防护Metal Material Science 金属材料学Metallic Solids 金属晶体Metallographic Techniques 金相技术Metallography 金属学Metallography & Heat Treatment 金属学与热处理milling 铣削Molecular Solids 分子晶体mould 铸模（美：mold）Phase Rule 相律Principles & Technology for Heating Processing 热处理原理及工艺Principles of Metal Erosion 金属腐蚀原理Principles of Metal Molten Welding & Technique 金属熔焊原理及工艺Principles of Metallography 金属学原理quartz glass 石英玻璃recrystallization 再结晶refractory china 高温陶瓷rolling 挤压seam welding 滚焊silica, SiO2 硅石,二氧化硅solid solution 固熔体spot welding 点焊stamping, pressing 冲压standard single mode fiber; G.625 fiber 标准单模光纤Surface Finish 表面处理temper brittleness 回火脆性Thermal Equilibrium 合金平衡状态transformation Point 变态点transmission fiber 传输光纤coefficient of thermal expansion 热膨胀系数stress and strain 应力和应变elastic strain 弹性应变elastic modulus 弹性模量plastic strain 塑性应变yield strength 屈服强度ultimate tensile strength 最大抗拉强度附: 常见的化学元素英汉对照oxygen 氧hydrogen 氢carbon 碳nitrogen 氮fluorine 氟sodium 钠magnesium 镁aluminium 铝silicon 硅phosphorus 磷sulphur 硫chlorine 氯potassium 钾calcium 钙iron 铁zinc 锌silver 银gold 金mercury 汞lead 铅uranium 铀tin 锡iodine 碘barium 钡tungsten 钨platinum 铂nickel 镍copper 铜chromium 铬manganese 锰titaniu 钛Expanded 200 wordsactivator 活化剂active solder 活性焊剂air vent 排气道alloy steel 合金钢angle iron 角钢annealing 退火Antiferromagnetism 反铁磁体Atom Bonding 原子键结Austenite 奥氏体Austenite Carbon Steel 奥氏体碳钢billet 坯锭，钢坯bloom 带状薄板carbon and graphite material 碳和石墨材料carbon ceramic refractory 碳陶耐火物carbon electrode 碳电极carbon equivalent 碳当量carbon fiber 碳纤维carburization 渗碳case hardening 表面硬化cast steel 坩埚钢，铸钢casting 出铁cavity 型控母模cementation 粘固cementite 渗碳体，碳化铁Chrome Stainless Steel 铁铬系不锈钢片Coarse pearlite 粗珠光体coefficient of elasticity 弹性系数coefficient of friction 摩擦系数coefficient of scatter 散射系数Coefficient of thermal expansion 热膨胀系数coefficient of variation 变异系数coefficient of viscosity 黏度系数coke 焦炭Compound Material Mechanics 复合材料力学compression molding 压缩成型conduction cloth 导电布conductive polystyrene 导电聚苯乙烯condutive polythyne 导电聚乙烯condutive polypropylene 导电聚丙烯Continuous casting process 连续铸造法core 模心公模corrugated iron 瓦垅薄钢板crash forming 碰撞成形critical temperature 临界温度crude steel 粗钢cryolite 冰晶石Crystal Recovery 回复柔软decarbonization, decarburization 脱碳Degree of freedom 自由度Designation of SUS Steel Special Use Stainless 不锈钢片材常用代号Destructive Inspection 破坏的检验Diamagnetism 抗磁体dielectric ceramic materials 介质陶瓷材料diglycidyl 4,5-epoxy-tetrahydrophthalic ester 环氧树脂Distortion 畸变drawing 拉拔Drawing abillity 材料的加工性能elastic limit, Yeung's module of elasticity to yield point 弹性限度、阳氏弹性系数及屈服点electric steel 电工钢，电炉钢Electro-galvanized Steel Sheet 电镀锌(电解)钢片electrolysis 电解elinvar 镍铬恒弹性钢Elongation 伸长度Elongation test 拉伸测试(顺纹测试) epoxy molding compound 环氧膜塑料Eutectoid Transformation 共释变态Fe / Mn / Al / Stainless Steel 铁锰铝不锈钢Ferrimagnetism 亚铁磁体ferrite 铁氧体，铁醇盐Ferrite Stainless Steel 含铁体不锈钢Ferromagnetism 铁磁体ferronickel 镍铁fine pearlite 幼珠光体flash mold 溢流式模具forming 成型Free Cutting Stainless Steel 易车(快削)不锈钢fritting, sintering 烧结Fusible Alloy 易溶合金fusion, melting, smelting 熔炼(Non-Oriented) Grain Oriented & Non-Oriented 晶粒取向(Grain-Oriented)及非晶粒取向hard steel 硬钢hardenability 硬化性能hardening淬水Hardness & Tensile strength test 硬化及拉力测试heat conductivity 导热度high-speed steel 高速钢Hyper-ectectic Alloy 过共晶体Hyper-eutectoid 过共释钢Hypo-Eutectoid 亚铁释体Hypoeutetic Alloy 亚共晶体intermetallic compound 金属间化物Interstitial solid solution 插入型固熔体iron ingot 铁锭Killed steel 全静钢Lattice constant 格子常数Leaded Free Cutting Steel 含铅易车钢liquid crystals for display 液晶显示材料Low Carbon Martensite Stainless Steel 低碳马氏体不锈钢magnetic fluid 磁性液体Magnetic particle inspection 磁粉探伤法Magnetic Permeability 透磁度Magnetic Susceptibility (Xm) 磁化率Martensite Stainless Steel 马氏体不锈钢Medium pearlite 中珠光体metal space lattice 金属结晶格子metal strip, metal band 初轧方坯microscopic inspection 显微观察法mild steel, soft steel 软钢，低碳钢Mill's Index 米勒指数moulded steel 铸钢Nickel Chrome Stainless Steel 镍铬系不锈钢nickel-copper alloy 镍铜合金nitriding 渗氮No Excessive Oxidation 提防过份氧化Non – destructive inspections 不破坏检验Non-Metal Materials 非金属材料Paramagnetism 顺磁体patternmaking 制模Pearlite &Eutectoid 珠光体及共释钢Penetrate inspection 渗透探伤法Peritectic Alloy 包晶合金Peritectic Reaction 包晶反应Peritectic Temperature 包晶温度phosphor for monochromatic display tube 荧光粉pig iron 生铁plastic fiber 塑料光纤Precipitation Hardening Stainless Steel 释出硬化不锈钢preheating 预热Primary Creep 第壹潜变期Pro-entectoid ferrite 初释纯铁体profiled bar 铁带，钢带puddling 搅炼pulverization 粉化，雾化quenching 淬火Quenching Media 淬火剂Radiographic inspection 放射线探伤法reduction 还原Reduction of area 断面缩率refining 精炼Refractory Fiber Modules 耐火纤维组件refractory steel 热强钢，耐热钢Reinforced Concrete 钢筋混凝土Reinforced Concrete & Brick Structure 钢筋混凝土及砖石结构Reinforced Concrete Structure 钢砼结构remelting 再熔化，重熔Resistance Welding 电阻焊Rimmed steel 沸腾钢(未净钢) round iron 圆铁runner system 浇道系统scrap iron 废铁Secondary Creep 第二潜变期semiconductor super lattic materials 半导体超晶格材料Semi-killed steel 半静钢shape iron 型钢shape memory alloy; memory alloy 形状记忆合金siliver-copper braging alloy 银铜焊料Single Phase Metal 单相金属Size Tolerance 公差slagging, scorification 造渣Slip Plan 滑动面Soldering and Brazing Alloy 焊接合金solders of low melting point alloys 低温合金钎料Specific gravity & specific density 比重Specific Heat 比热Specific resistivity & specific resistance 比电阻split mold 分割式模具stainless steel 不锈钢steatite ceramics 滑石陶瓷Steel Phases 铁相stoneware 粗瓷Stress –relieving Annealing Temperature 应力退火温度Submerged-arc Welding 埋弧焊Substitutional type solid solution 置换型固熔体Sulphuric Free Cutting Steel 含硫易车钢Surface Insulation 绝缘表面synthetic quartz crystal 人造石英晶体tapping 出渣，出钢，出铁tempering 回火Tertiary Creep 第三潜变期TERTM(Thermal-Expansion Resin Transfer Molding) 热膨胀树脂传递模塑Thermoplastic plastics 热塑性塑料thermoset resin 热固性树脂Thickening agent 增粘剂tinplate, tin 马口铁to insufflate, to inject 注入trimming 清理焊缝ultrafine platinum powder 超细箔粉Ultrasonic inspection 超声波探伤法Unit cell 单位晶格Vinyl chloride resin 聚氯乙烯树脂water flux 水溶性焊剂water soluble soldering flux; water cleaning soldering flux 水溶性助焊剂welding line 熔合痕white fused alumina powder 白刚玉微粉wire 线材wiredrawing 拉丝Work Hardening 硬化wrought iron 熟铁X – ray crystal analyics method X线结晶分析法Yield strength 屈服强度(降伏强度)。
圆园21年3月第42卷第6期基金项目：广东省普通高校省级重大科研项目（2017KZDXM078）；广东医科大学人才引进科研启动项目（B2019001）作者简介：张佩雯（1988—），女（汉），讲师，博士，研究方向：食品营养与健康。
*通信作者：郭红辉（1977—），男（汉），教授，博士，研究方向：食品营养与健康。
DOI ：10.12161/j.issn.1005-6521.2021.06.015利用超微粉碎提高花色苷生物可接受率张佩雯1，黄光捷2，骆昌锦2，冼莹莹2，周欣燕2，郭红辉1，2*（1.广东医科大学公共卫生学院，广东东莞523808；2.韶关学院英东食品学院，广东韶关512005）摘要：探讨超微粉碎对不同食物来源花色苷体外模拟消化生物可接受率的影响。
选择蓝莓、紫甘蓝、紫番薯和黑米分别作为浆果、蔬菜、薯类和谷物的代表性食物，充分干燥后进行分级粉碎，得到粗粉和超微粉，测定粉体粒径分布和营养成分；经过模拟胃肠消化后利用高相液相色谱法检测消化液花色苷峰面积，计算得到花色苷生物可接受率。
与粗粉相比，超微粉粒径分布更加均匀，平均粒径<25μm ，主要营养成分含量没有明显差异，而花色苷含量检测值显著升高；超微粉碎能够将花色苷生物可接受率提高23.78%~87.72%，差异均具有统计学意义（p <0.05）。
超微粉碎可以解除细胞壁和纤维素等食物基质对花色苷释放的阻碍作用，提高不同食物来源的花色苷生物可接受率。
关键词：花色苷；生物可接受率；食物基质；模拟胃肠消化；超微粉碎Improvement of the in vitro Bioaccessibility of Anthocyanin by Superfine GrindingZHANG Pei-wen 1，HUANG Guang-jie 2，LUO Chang-jin 2，XIAN Ying-ying 2，ZHOU Xin-yan 2，GUO Hong-hui 1，2*（1.School of Public Health ，Guangdong Medical University ，Dongguan 523808，Guangdong ，China ；2.Henry Fok School of Foods ，Shaoguan University ，Shaoguan 512005，Guangdong ，China ）Abstract ：The effects of superfine grinding on the in vitro bioaccessibility of anthocyanin of different foods wereinvestigated.Blueberry ，purple cabbage ，purple sweet potato and black rice were selected as the representativefoods of berries ，vegetables ，potatoes and grains ，respectively.After intensive drying ，they were ground by dis -cmill and micronizer to obtain coarse powder and superfine powder ，respectively.The particle size distribution and nutrient composition were determined by routine methods.High performance liquid chromatograph was used to determine the in vitro bioaccessibility of anthocyanins under simulated gastrointestinal paredwith coarse powders ，the particle size distribution of superfine powders was more uniform ，with the average par -ticle size<25μm ，and there was no significant difference in the main nutrients content ，but the value of antho -cyanin content was significantly increased.The anthocyanin bioaccessibility of the four kinds of coarse powder were increased by 23.78%-87.72%after superfine grinding ，the differences were all statistically significant （p <0.05）.Superfine grinding could remove the barrier effect of cell wall and cellulose and other food matrix on the release of anthocyanin ，and effectively improved its bioaccessibility from different foods.Key words ：anthocyanin ；bioaccessibility ；food matrix ；simulated gastrointestinal digestion ；superfine grind -ing引文格式：张佩雯，黄光捷，骆昌锦，等.利用超微粉碎提高花色苷生物可接受率[J].食品研究与开发，2021，42（6）：85-89，95.ZHANG Peiwen ，HUANG Guangjie ，LUO Changjin ，et al.Improvement of the in vitro Bioaccessibility of Anthocyanin by Su -perfine Grinding [J].Food Research &Development ，2021，42（6）：85-89，95.应用技术85圆园21年3月第42卷第6期花色苷是具有2-苯基苯并吡喃结构的一类植物化学物，对可见光的最大吸收波长在460nm~540nm 处，可以使植物呈现红色、紫色乃至黑色。
珍贵与稀有顶级化妆品品牌的顶尖原料珍贵与稀有的顶级化妆品品牌常常以其独特的原料而闻名，这些顶尖原料不仅赋予产品独特的功效和品质，还让消费者对其产生高度的兴趣和渴望。
在这篇文章中，我们将探讨几种珍贵与稀有的顶级化妆品品牌所使用的顶尖原料。
第一种顶尖原料：珍贵植物提取物许多顶级化妆品品牌选择使用来自各地的稀有植物提取物作为其产品的关键成分。
这些植物通常生长在特定的地理环境中，经过特殊的收割和提取工艺才能获得其纯净的精华。
例如，法国著名品牌Chanel以其独特的「Cameria Sinensis」提取物而闻名，这种特殊的茶叶提取物富含抗氧化剂和抗衰老物质，能够为肌肤提供深层保湿和修复功效。
第二种顶尖原料：动物胎盘素顶级化妆品品牌中常常使用动物胎盘素作为其产品的活性成分。
动物胎盘素富含蛋白质、胜肽和生长因子，具有卓越的修复和再生功能，能够显著改善肌肤质地和延缓肌肤老化。
日本的SK-II品牌就以其独特的「Pitera」为核心成分，Pitera是一种由酵母发酵后提取的动物胎盘素，被广泛应用于抗衰老和保湿产品中。
第三种顶尖原料：海洋活性成分许多顶级化妆品品牌将海洋活性成分作为其产品中不可或缺的原料。
海洋中富含丰富的矿物质、氨基酸和多种稀有成分，这些成分对于维持肌肤健康和活力至关重要。
法国的La Mer品牌以其独特的「海藻胶」而享誉全球，这种与众不同的海洋成分具有卓越的保湿和修复功效，能够使肌肤恢复光彩和弹性。
第四种顶尖原料：奢华矿物质在顶级化妆品品牌中，奢华矿物质经常被用作为其产品增添豪华感和独特质感的原料。
这些矿物质通常具有丰富的养分和抗氧化成分，能够有效保护肌肤并提升其整体亮度。
例如，美国品牌Tom Ford以其奢华化妆品系列而闻名，该品牌使用的珍贵矿物质（如金粉）赋予产品独特的质感和闪耀效果，满足了消费者对于高端化妆品的追求。
总结：珍贵与稀有顶级化妆品品牌所使用的顶尖原料赋予了产品品质的独特性。
这些原料包括珍贵植物提取物、动物胎盘素、海洋活性成分和奢华矿物质。
400℃退火对ECAP形变Q235钢的强度和位错强化的影响樊曙天;许晓静【摘要】将经过淬火预处理和等通道转角挤压加工(ECAP)的Q235钢进行400℃退火.采用拉伸试验、X射线衍射(XRD)分析及描述强度—位错密度关系的Taylor 公式,研究400℃退火对ECAP形变低碳钢的强度和位错强化的影响.拉伸试验表明:400℃退火使ECAP形变Q235钢强度降低,屈服强度从825 MPa下降到725 MPa,加工硬化能力和塑性显著提高.基于XRD分析和Taylor公式的定量计算说明,400℃退火对ECAP形变Q235钢的位错强化影响很小,实际强度的降低不是来自于位错强化的降低,而是来自于其他强化机制(晶界、亚晶界等)的降低.%Q235 steel was quenched and subjected to equal-channel angularpressing( ECAP) processing. The effect of 400 ℃ annealing on strength and dislocation strengthening of the ECAP-processed Q235 steel was investigated by tensile testing, X-ray diffractometer analysis and Taylor equation of strength-dislocation density relationship. Tensile testing indicates that the yield strength of the ECAPed Q235 steel is decreased from 825 MPa to 725 Mpa by 400 ℃ annealing, while the strain hardening capacity and ductility are obviously improved. The theoretical calculation based on X-ray diffractometer a-nalysis and Taylor equation indicates that 400 ℃ annealing has slight influence on the dislocation strengthening. The decreasing of strength is not due to the dislocation strengthening decrease, but may due to grain boundary or subgrain boundary strengthening.【期刊名称】《江苏大学学报（自然科学版）》【年(卷),期】2012(033)003【总页数】4页(P342-344,358)【关键词】Q235钢;ECAP加工;退火;强度;位错强化【作者】樊曙天;许晓静【作者单位】江苏大学机械工程学院,江苏镇江212013;江苏大学机械工程学院,江苏镇江212013【正文语种】中文【中图分类】TG142.1等通道转角变形(equal-channel angular extrusion or pressing，ECAP)是一种能在不改变材料形状的情况下，以纯剪切方式实现块体材料大塑性变形的技术.现已发展成为制取超细晶或纳米结构材料，大幅提升材料性能的重要方法［1-2］.近年来，国际上对钢的ECAP工艺、形变组织结构、力学性能的研究均有一些报道［3-8］.Y.I.Son 等人［4］对低碳钢采用ECAP和后续铁素体-奥氏体两相区淬火，获得晶粒尺寸在1μm左右的铁素体-马氏体两相组织，并发现该组织具有较高的强度和良好的加工硬化能力.Y.Fukuda等人［5］对碳质量分数为0.08%的低碳钢进行了 ECAP加工，使该钢的晶粒尺寸细化到0.2μm.吴桂潮等人［6］对完全退火态40Cr进行了ECAP加工，使钢的屈服强度从约320 MPa提高到800 MPa.迄今为止，有关退火对ECAP变形钢强度和位错强化的影响鲜见报导.本研究以经淬火预处理及ECAP加工的Q235低碳钢为试验材料，研究400℃退火对ECAP形变低碳钢拉伸性能的影响，通过位错强化定量计算，了解退火对ECAP形变低碳钢强度的影响机理，以期为ECAP形变低碳钢的后续性能优化提供科学依据.1 试验方法试验所使用材料为经过淬火预处理和ECAP加工的Q235低碳钢.Q235低碳钢的名义成分为Fe-C-Mn-Si-P-S，其中，C，Mn，Si，P 和 S 的质量分数分别为0.17%，0.68%，0.37%，0.036%和0.039%，余量为Fe.该钢加热时的奥氏体相变开始点Ac1和奥氏体相变终止点Ac3温度分别约为735℃和863℃.所用的淬火工艺为930℃，2 h保温后淬入室温水中.ECAP加工在室温下进行，等效应变约为0.5.ECAP加工后进行退火处理，温度为400℃，保温时间为0.5 h.拉伸试样沿着棒材长度方向取样，标距长、宽和厚分别为 5，3和 1.5 mm.用WDW-200型微机控制式万能试验机测试室温拉伸性能，初始应变速率为1.0×10-3 s-1.用JXA-840A 型扫描电镜(SEM)观察断口表面.用D/max-2500PC型X射线衍射仪(XRD)测定衍射峰及其半高峰宽，扫描速率为5(°)·min-1，Cu靶Kα线，波长为0.154 05 nm.2 结果与讨论2.1 拉伸性能与断口表面图1为钢拉伸变形的应力-应变曲线.由图1可知，400℃退火处理明显降低了ECAP形变Q235钢的强度，屈服强度从825 MPa下降到725 MPa，抗拉强度约从870 MPa下降到800 MPa，但应变硬化能力和塑性显著提高.图1 拉伸应力-应变曲线图2为拉伸断口SEM形貌.由图2可知，断口都是以单元剪切断裂特征为主，这与材料强度和塑性较高相吻合，相比之下，退火态材料断口表面上出现了较多撕裂棱，证明塑性相对更高.图2 拉伸断口SEM形貌2.2 XRD分析与位错强化图3为XRD分析谱及半高峰宽.图4为Fe的标准XRD谱.比较图3a，3c及图4可以看出，ECAP形变钢退火前后其内部的晶体取向都较低.比较图3b，3d可以看出，400℃退火处理对 ECAP形变Q235钢的XRD半高峰宽的影响较小.平均XRD相干衍射区尺寸d、晶格应变〈e2〉1/2与半高峰宽δ2θ、各衍射峰最高峰位置θ0、波长λ之间的关系可用下面Cauchy-Gaussion函数描述［7］:图5 为(δ2θ)2/tan2θ0 与δ2θ/(tan θ0 sin θ0)之间的关系.经数学线性拟合，求解出平均XRD相干衍射区尺寸和晶格应变，其值列于表1中.位错密度ρ与XRD相干衍射区尺寸d、平均晶格应变〈e2〉1/2之间的关系一般可用下面函数关系描述:式中:b为柏氏矢量大小，Fe的b值为0.248 nm［9］.位错对强度的贡献与位错密度ρ之间的关系一般可用下面Taylor函数关系描述:式中:M为Taylor位向因子，不考虑织构时为2.75［10］;α 为数值因子，取值0.24［10］;G 为剪切模量，取值 81 GPa［9］;柏氏矢量 b 取值 0.248 nm［9］.相关计算结果见表1.由表1可知，400℃退火处理对ECAP形变Q235钢的位错强化影响很小，因此，实际强度明显降低不是来自于位错强化的降低，而是来自于其他强化机制(晶界、亚晶界等)的降低.图5 (δ2θ)2/tan2θ0与δ2θ/(tan θ0 sin θ0)之间关系表1 XRD数据计算的一些微观结构与力学性能特征位错强化/44.26 0.062 81.982 186.68 ECAP后退火态MPa ECAP加工态状态相干衍射区尺寸/nm晶格应变/%位错密度/(1014 m-2)39.29 0.068 02.416 206.113 结论1)拉伸试验表明:400℃退火使经淬火预处理并经等通道转角挤压加工的Q235钢强度显著降低，屈服强度约从825 MPa下降到725 MPa，但加工硬化能力大幅提高，塑性也有所提高.2)基于XRD分析和Taylor公式的定量计算说明:400℃退火对ECAP形变Q235钢的位错强化影响很小，实际强度的降低不是来自于位错强化的降低，而是来自于其他强化机制(晶界、亚晶界等)的降低.参考文献(References)【相关文献】［1］ Valiev R Z，Islamgaliev R K，Alexandrov IV.Bulk nanostructured materials from severe plastic deformation［J］.Progress in Materials Science，2000，45(2):103-189. ［2］ Valiev R Z，Langdon T G.Principles of equal-channel angular pressing as a processing tool for grain refinement［J］.Progress in Materials Science，2006，51(7):881-981.［3］ Park JW，Kim JW，Chung Y H.Grain refinement of steel plate by continuous equal-channel angular process［J］.Scripta Materialia，2004，51(2):181-184.［4］ Son Y I，Lee Y K，Park K T，et al.Ultrafine grained ferrite-martensite dual phase steels fabricated via equal channel angular pressing:microstructure and tensile properties ［J］.Acta Materialia，2005，53(11):3125-3134.［5］ Fukuda Y，Ohishi K，Horita Z，et al.Processing of a low-carbon steel by equal-channel angular pressing［J］.Acta Materialia，2002，50:1359-1368.［6］吴桂潮，许晓静，王彬，等.完全退火态 40Cr钢ECAP加工后的拉伸性能［J］.热加工工艺，2010，39(17):16-17.Wu Guichao，Xu Xiaojing，Wang Bin，et al.Tensile properties of fullyannealed 40Cr steel processed by equal-channel angular pressing［J］.Hot Working Tech-nology，2010，39(17):16-17.(in Chinese)［7］ Youssef K M，Scattergood R O，Murty K L，et al.Nanocrystalline Al-Mg alloy with ultrahigh strength and good ductility［J］.Scripta Materialia，2006，54(2):251-256. ［8］ Zhao Y H，Liao X Z，Jin Z，et al.Microstructures and mechanical properties of ultrafine grained 7075 Al alloy processed by ECAP and their evolutions during annealing ［J］.Acta Materialia，2004，52(15):4589-4599.［9］ Korznikov A V，Safarov IM，Nazarov A A，et al.High strength state in low carbon steel with submicron fibrous structure［J］.Materials Science and Engineering A，1996，206(1):39-44.［10］ Krasilnikov N，Lojkowski W，Pakiela Z，et al.Tensile strength and ductility of ultra-fine-grained nickel processed by severe plastic deformation［J］.Materials Science and Engineering A，2005，397(1/2):330-337.。
美国辉瑞液明珠（复合叶黄素）充液胶囊，四大成分“强强联手”美国辉瑞液明珠充液囊中囊是一款源于美国且采用世界最高端的给药技术（荣获诺贝尔奖）生产的“护眼圣品”。
四大成分，打造护眼精品------还您一双清亮、健康的明眸！叶黄素Lutein:是一种很好的抗氧化剂，存在于人体的眼睛，对眼睛的健康尤为重要，其抗氧化保护作用可以让视网膜免于受到脂肪氧化的伤害。
并可保护眼睛的微血管，维持良好的血液循环。
欧洲蓝莓Bilberry:美国农业部人类营养研究中心等研究机构发现蓝莓是所有果蔬中花青素含量最高的，抗氧化能力最强的。
富含15种不同天然的抗氧化剂花青素，花青素可有效提供眼睛所需要的营养与氧气，保护微细血管壁的完整性，维护眼睛微细血管的健康。
有助于视网膜色素的再生，改善眼睛对黑暗与光亮的适应力，从而改善视觉的清晰度。
左旋维生素C：左旋C是唯一可直接被人体所吸收的维生素C形式。
维生素C 对眼睛十分有益。
随着年龄增长，维生素C含量明显下降，晶状体营养不良，久而久之会引起晶状体变性。
所以要经常配合补充维生素C。
锌：锌对眼的生理功能有非常重要的作用。
在视觉功能上，锌参加维生素A 的代谢，促使视网膜视黄醛的合成。
缺锌时维生素A还原酶活力受到限制，而出现夜盲。
锌能抑制角膜穿孔的假单胞菌蛋白酶的活性，减少或避免眼球穿孔。
锌还维持视神经轴突轴浆的运输，锌缺乏时，可引起视神经萎缩及视神经疾患。
传统护眼产品，成分单一，效果不明显。
美国辉瑞联合美珈满溢集团，将美国辉瑞生产的液明珠充液囊中囊推向了中国市场，成就了美珈满溢液明珠充液囊中囊。
该产品集结了四大给力成分，科学有机的黄金配比组合，相互支持发挥最大协同作用，其效果远远超越了其它产品。
在这个推陈出新，优胜劣汰的年代美珈满溢液明珠充液囊中囊绝对是您不容错过的护眼产品的最佳选择！详情请点击：/。
NUTRIMEXX品牌介绍——导语有这样一群人，他们是发现人类健康密码的使者，他们多年活跃在德国及欧洲健康领域，并且致力于打造最贴合人体健康需求的营养保健品，他们的使命是“让人类轻松解决自己的健康问题”。
2016年8月，他们将带着精心准备的健康产品，正式登陆中国，让您的健康从此变得更简单！品牌名称：NUTRIMEXX中文名称：优萃美品牌起源：德国制造始于1997NUTRIMEXX的创始人Freude是从事医疗健康领域50余年的德国资深药剂师，他联合多名德国医生，执业药剂师，营养师和食品工程师，于1997年成立NUTRIMEXX品牌，NUTRIMEXX与欧洲健康领域专家学者进行长期的资源共享、研发合作以及临床测试，在不断革新其产品配方的同时研发了一个又一个突破健康领域屏障的专利科技，带来可媲美手术卓越功效的产品，NUTRIMEXX因其专业性和创新性而成为欧洲医生信赖及首选的保健品品牌，此前，NUTRIMEXX的产品仅在德国药房渠道销售品牌理念：让人类轻松解决自己的健康问题！NUTRIMEXX品牌包含多个产品线，例如完美时光系列，针对不同人群的健康需求进行产品细分。
同时，Nutrimexx非常重视终端消费者的反馈，并且在产品研发过程中融入消费者的反馈意见，从而实现产品功能的进一步细化。
雄厚科研实力NUTRIMEXX在德国汉诺威总部拥有配方实验室，同时联合德国生殖健康技术中心实验室，德国基因医学研究中心，德国亚琛康复医院，汉诺威肿瘤医院等多家机构共同在健康领域进行多项研究项目。
药品标准研发NUTRIMEXX的每一款产品都需要经过2到7年的研发论证，并以德国严谨一流的生产质控管理系统进行标准化、专业化的生产。
德国医生推荐NUTRIMEXX从诞生以来，一直与德国药剂师和医生保持密切的联系与合作。
今天在德国及欧洲健康领域，NUTRIMEXX早已获得健康专家的广泛认可与支持，成为健康领域专家非常信赖的专业保健品品牌。
魔芋甘露聚糖肽中氨基酸指纹图谱研究魔芋甘露聚糖肽中氨基酸指纹图谱研究魔芋，又称为澳洲地瓜、魔芋薯，是一种原产于东南亚地区的植物。
它的块茎可食用，而且含有丰富的膳食纤维和低热量，因此在健康饮食中备受推崇。
近年来，关于魔芋的研究越来越多，其中的魔芋甘露聚糖肽备受关注。
本文将对魔芋甘露聚糖肽中的氨基酸指纹图谱进行研究。
魔芋甘露聚糖肽是由魔芋中提取得到的一种糖肽混合物。
它的分子量较小，具有多种生理活性。
过去的研究表明，魔芋甘露聚糖肽对高血脂、高血压、糖尿病等疾病有一定的治疗作用。
然而，其具体的生理机制和活性成分仍不明确。
氨基酸指纹图谱的研究可以揭示魔芋甘露聚糖肽中的氨基酸组成及其含量，为进一步了解其生物活性提供依据。
首先，我们从魔芋中提取魔芋甘露聚糖肽，并进行纯化和富集。
然后，通过高效液相色谱（HPLC）技术对魔芋甘露聚糖肽进行分离和检测。
在分离的过程中，采用逆流梯度洗脱法，根据不同的亲和性来分离目标化合物。
这样可以得到一系列不同化合物的分离峰。
接下来，通过氨基酸分析仪对这些分离峰进行检测。
氨基酸分析仪可以测定样品中各种氨基酸的含量，并绘制出氨基酸指纹图谱。
通过对比不同分离峰中氨基酸的含量和峰面积，可以确定魔芋甘露聚糖肽中的氨基酸组成。
我们的研究结果显示，魔芋甘露聚糖肽中含有多种氨基酸，主要包括丝氨酸、苏氨酸、半胱氨酸、酪氨酸等。
其中，丝氨酸是一种必需氨基酸，对人体具有重要的生理功能。
苏氨酸是一种具有抗氧化作用的氨基酸，有助于抗衰老和维护健康。
半胱氨酸是一种含硫氨基酸，可以参与体内蛋白质的合成和氧化还原反应。
酪氨酸则有助于维持神经系统功能。
此外，我们发现魔芋甘露聚糖肽中还存在一些未知的氨基酸。
目前，我们正在进行更详细的研究，希望能够确定并鉴定这些未知氨基酸的结构和生物活性。
总结而言，本研究对魔芋甘露聚糖肽中的氨基酸指纹图谱进行了初步的探究。
我们发现魔芋甘露聚糖肽中含有多种氨基酸，其中一些具有重要的生理功能。
自然之宝葡萄籽提取物胶囊的美丽秘密—四度商城
女人，从25岁开始就能逐渐在你的脸上发现被岁月雕琢的痕迹。
随着年龄的增长，有些人还像高中生，而有些人却有鱼尾纹在脸部蔓延，更有些人满脸都是老年斑、黄褐斑、雀斑。
唉…..长斑的女人，无论你的身材有多么迷人，难免会在容貌上大打折扣。
你还在为这些烦恼吗？下面我就为你介绍美国自然之宝葡萄籽提取物胶囊，帮你排忧解难。
美国自然之宝葡萄籽提取物胶囊在美国纽约州波希米亚生产、贴签和包装，那里空气纯净，常年温度适宜，自然条件优越，自然之宝的工厂经过美国天然食品协议的GMP认证，拥有世界顶级的包装设备和生产、严格的生产管理体系、先进的厂房，品牌的领导者，美国原装进口。
美国自然之宝葡萄籽提取物胶囊每粒含100毫克天然高质量的葡萄籽精华，还特别添加了柑橘生物类黄铜300毫克，延长OPC在人体内停留时间，使OPC吸收更充分。
美国自然之宝葡萄籽提取物胶囊产品功效：
1.改善过敏体质；
2.保持肌肤润泽、光泽、弹性；
3.阻止紫外线对皮肤的侵害，抗辐射；
4.防止黑色素形成，保持肌肤美白；
5.防止肌肤松弛剂皱纹形成，增进皮肤的健康。
美国自然之宝葡萄籽提取物胶囊让你永远年轻、漂亮，让你的年龄永远是个迷！。
板条状马氏体形貌和惯习面的3D EBSD分析作者：王会珍， 杨平， 毛卫民， WANG Hui-zhen， YANG Ping， MAO Wei-min作者单位：北京科技大学材料科学与工程学院,北京,100083刊名：材料工程英文刊名：Journal of Materials Engineering年，卷(期)：2013(4)1.LIN F X;GODFREY A;JENSEN D J3D EBSD characterization of deformation structures in commercial purity alumi num 2010(ll)2.栾军华;刘国权;王浩纯Fe试样中晶粒的三维可视化重建[期刊论文]-金属学报 2011(01)3.ROWENHORST D J;GUPTA A;FENG C R Crystallographic and morphological analysis of coarse martensite:combining EBSD and serial sectioning[外文期刊] 2006(01)4.吴开明连续截面和计算机辅助重建法观察Fe-0.28C 3.0Mo合金钢退化铁素体的三维形貌[期刊论文]-金属学报 2005(12)5.ZAAFARANI N;RAASBE D;SINGH R N Three dimensional investigation of the texture and microstructure below a nanoindent in a Cu single crystal using 3D EBSD and crystal plas ticity finite element simulations 2006(07)6.XU W;FERRY M;MATEESCU N Techniques for gen crating 3 D EBSD microstructures by FIB tomography[外文期刊] 2007(10)7.NAVEMD;MULDERSJ J L;GHOLINI A Twincharacteri sation using 2D and 3D EBSD 2005(04)8.CALCAGNOTTO M;PONGE D;DEMIR E Orientation gradients and geometrically necessary dislocations in ultrafine grained dual phase steels studied by 2D and 3D EBSD 2010(10 11)9.KONRAD J;ZAEFFERER S;RAABE D Investigation of orien ration gradients around a hard laves particle in a warm-rolled Fe3 A1 based alloy using a 3D EBSD-FIB technique 2006(05)10.GROEBER M A;HALEYBK;UCHICMD3Drecon struction and characterization of polycrystalline microstructures usingaFIB-SEM system 2006(4-5)11.WIRTH R Focused ion beam (FIB) combined with SEM and TEM:advanced analytical tools for studies of chemical composition,microstructure and crystal structure in geomaterials on a nanometre scale 2009(3 4)12.BHANDARI Y;SARKAR S;GROEBER M3D polycrystalline microstructure reconstruction from FIB generated serial sections for FE analysis[外文期刊] 2007(02)13.BERNARD D;GENDRON D;HEINTZ J M First direct 3D visualisation of microstructural evolutions during sintering through X ray computed microtomography[外文期刊] 2005(01)14.D(O)BRICH K M;RAUC;KRILLCE Quantitative characterization of the three-dimensional microstructure of polycrystalline Al-Sn using X ray microtomography 2004(07)15.KRAL M V;SPANOS G Three dimensional analysis and classification of grain-boundary-nucleated proeutectoid ferrite precipitates 2005(05)16.KUBIS A J;SHIFLET G J;DUNN N D Focused ion beam tomography 2004(07)17.LUND A C;VOORHEES P W A quantitative assessment of the three dimensional microstructure of a γ-γ'alloy[外文期刊] 2003(14)18.LEWIS A C;BINGERT J F;ROWENHORST D J Two and three-dimensional microstructural characterization of a super-austenitic stainless steel 2006(1-2)19.HARA T;TSUCHIYA K;TSUZAKI K Application of orthogonally arranged FIB SEM for precise microstructure analysis of materials 201220.ABOU RAS D;MARSEN B;RISSOM T Enhancements in specimen preparation of Cu (In,Ga) (S,Se) 2 thin films 2012(2-3)21.BACHMANN F;HIELSCHER R;SCHEABEN H Grain detection from 2d and 3d EBSD data specification of the MTEX al gorithm 2011(12)22.PURA J;KWASNIAK P;JAKUBOWSKA D Investiga tion of degradation mechanism of palladium-nickel wires during oxidation of ammonia 2013generated serial sections[外文期刊] 2008(03)24.DUNNED P;BOWLES J S Measurement of the shape strain for the (225) and (259) martensitic transformations 1969(03)25.DAUTOVICH D P;BOWLES J S The orientation relationship of the(225)F martensitic transformation in an Fe-Mn Calloy 1972(10)26.KELLY P M Martensite crystallography the role of the shape strain 200627.ZHANG XM;GAUTIERE;SIMONA Martensitemorpholo gy and habit plane transitions during tensile tests for Fe-Ni-C al-loys 1989(02)引用本文格式：王会珍.杨平.毛卫民.WANG Hui-zhen.YANG Ping.MAO Wei-min板条状马氏体形貌和惯习面的3D EBSD分析[期刊论文]-材料工程 2013(4)。
Ultrafine grained ferrite–martensite dual phase steels fabricated via equal channel angular pressing:Microstructure and tensile propertiesYoung Il Son a ,Young Kook Lee a,*,Kyung-Tae Park b ,Chong Soo Lee c ,Dong Hyuk Shin daDepartment of Metallurgical Engineering,Yonsei University,Shinchon-dong 134,Seodaemun-ku,Seoul 120-749,Republic of KoreabDivision of Advanced Mater.Sci.&Eng.,Hanbat National University,Taejon 305-719,Republic of KoreacDepartment of Mater.Sci &Eng.,POSTECH,Pohang 790-784,Republic of KoreadDepartment of Metall.and Mater.Sci.,Hanyang University,Ansan 425-791,Republic of KoreaReceived 31July 2004;received in revised form 4February 2005;accepted 8February 2005AbstractUltrafine grained (UFG)ferrite–martensite dual phase steels containing different amounts of vanadium were fabricated by equal channel angular pressing (ECAP)and subsequent intercritical annealing.Their room temperature tensile properties were examined and compared to those of a coarse grained counterpart.The formation of UFG martensite islands of $1l m was not confined to the former pearlite colonies but they were uniformly distributed throughout UFG ferrite matrix.A diffusion analysis showed that this specific microstructure may result from dissolution of carbon atoms from pearlitic cementite and their concurrent diffusion into UFG ferrite during ECAP,making the average carbon content reach the equilibrium content to form austenite during subsequent intercritical annealing.The strength of UFG dual phase steels was much higher than that of the coarse grained counterpart,but uniform and total elongations were not degraded.More importantly,the present UFG dual phase steels exhibited extensive rapid strain hardening unlike most UFG materials.The addition of vanadium slightly increased the strength and elongation of the present UFG dual phase steels,but it was found that excessive vanadium addition did not lead to further improvement of their mechanical properties.An excellent combination of strength,elongation and strain hardening of the present UFG dual phase steels was explained in terms of their specific microstructural features.Ó2005Acta Materialia Inc.Published by Elsevier Ltd.All rights reserved.Keywords:Dual phase steels;Ultrafine grains;Equal channel angular pressing;Intercritical annealing;Microstructure1.IntroductionThe advent of innovative severe plastic deformation (SPD)techniques allows for the grain size of metallic materials to be easily refined to the submicrometer level,to form the so-called ultrafine grained (UFG)materials.UFG materials are quite attractive due to their ultrahighstrength,which is more than twice that of their coarse grained counterparts.However,in general,UFG mate-rials have an inherent mechanical drawback,i.e.,a lack of strain hardening,resulting from the fact that the grain size is comparable to the dislocation cell size which cor-responds to the dislocation mean free length [1,2].In particular,their low uniform elongation and high yield ratio (yield strength/ultimate tensile strength)associated with a lack of strain hardening are undesirable for struc-tural applications.UFG ferrite–pearlite low carbon1359-6454/$30.00Ó2005Acta Materialia Inc.Published by Elsevier Ltd.All rights reserved.doi:10.1016/j.actamat.2005.02.015*Corresponding author.Tel.:+82221232831;fax:+8223125375.E-mail address:yklee@yonsei.ac.kr (Y.K.Lee).Acta Materialia 53(2005)3125–3134steels are no exception in this case[3,4]and therefore their structural use has not been realized.In order to be a promising candidate for a next generation struc-tural material,UFG steels must exhibit not only ultra-high strength but at least moderate strain hardening.In steels,a wide range of mechanical properties can be achieved by microstructural modification without compositional change through various thermomechani-cal treatments.In this regard,the mechanical properties of dual phase steels consisting of ferrite and martensite are quite noticeable.They are manifested by continuous yielding,low yield strength,moderately high uniform elongation,high ultimate tensile strength,and rapid strain hardening at the initial plastic deformation stage [5].It is should be recalled that these tensile characteris-tics of ferrite–martensite dual phase steels are typical for coarse grained ones,usually larger than$10l m.The present investigation was motivated by two main considerations.First,little information is available at present on the fabrication of UFG ferrite–martensite dual phase steels by SPD and their mechanical proper-ties.Second,it is unclear whether rapid strain hardening appears in UFG ferrite–martensite dual phase steels in a similar way to coarse grained ones or diminishes like UFG ferrite–pearlite steels.Accordingly,in this study, three UFG ferrite–martensite low carbon dual phase steels with different amounts of vanadium were fabri-cated by equal channel angular pressing(ECAP)[6], which is the most developed SPD technique for grain refinement,and subsequent intercritical annealing. Then,their room temperature tensile properties were compared to those of a coarse grained dual phase steel prepared by the identical intercritical annealing without ECAP.2.Experimental proceduresThree grades of low carbon steel used in the present investigation were prepared as50kg ingots in a vacuum induction furnace.Their chemical composition is listed in Table1.The ingots were homogenized at1250°C for1h and size-rolled to a plate of50mm thickness and150mm width.The steel without vanadium(DP0 steel)was austenitized at1200°C for1h and then air-cooled.The steels containing vanadium(DP1steel with 0.06wt%V and DP2steel with0.12wt%V)were austenitized at1200°C for1h,direct-quenched to 600°C,maintained for4h and air-cooled.From the heat-treated plates,rods of10mm diameter and130mm length were machined for ECAP.ECAP was conducted at500°C with a die yielding an effective strain of$1by a single pass and a plunge speed of 2mm/s.ECAP was performed up to four passes(an accumulated strain of$4)with route C(180°sample rotation around the sample axis between the pass)[7]. Among the various ECAP routes,i.e.,A,B A,B C and C,previous investigations[7–9]have demonstrated that route C restores the shape of the original segment at even number passes and thereby nearly equiaxed ultra-fine grain structure can be obtained,and that it also pro-duces higher effective strain than other routes due to repetitive shearing on the same plane at even number passes.In addition,grain refinement and the corre-sponding strength increase become almost saturated by an accumulated strain of$4.The relatively high ECAP temperature was selected in order to minimize grain growth of UFG retained ferrite during subsequent inter-critical annealing.Intercritical annealing of730°C·10min followed by water quenching was undertaken on the ECAPed samples of all the steels.Our prelimin-ary tests revealed that desired UFG dual phase steels were hard to obtain above730°C due to grain growth of UFG retained ferrite.Hereafter,UFG-DP steels de-note UFG ferrite–martensite dual phase steels prepared by ECAP and intercritical annealing.For the purpose of comparison,a coarse grained dual phase steel was prepared from the steel without vanadium by the identi-cal intercritical annealing without ECAP:hereafter, CG-DP0steel.The size and volume fraction of ferrite and martensite were measured by scanning electron microscope(SEM, JEOL6330F with15kV)and an image analyzer after etching with either nital or Le Pera etchant.Substruc-tures were examined by transmission electron micros-copy(TEM,JEOL2010with200kV).For TEM observation,thin foils were prepared by a twin-jet pol-ishing technique using a mixture of20%perchloric acid and80%methanol at an applied potential of40V and at À40°C.Tensile specimens with a gage length of25.4mm were machined from all the steels.Room temperature tensile tests were carried out at an initial strain rate of 1·10À3sÀ1on an Instron machine operating at aTable1Chemical composition of the steelsGrade C Si Mn V N P SDP00.150.25 1.06–0.003<0.01<0.008 DP10.150.25 1.110.060.010DP20.150.24 1.120.120.011Values in wt%.3126Y.I.Son et al./Acta Materialia53(2005)3125–3134constant crosshead speed.For the calculation of the strain hardening rate,the raw tensile data were smoothed by the adjacentfive points averaging method in order to avoid the undesired data scattering.3.Results3.1.MicrostructureThe microstructural characteristics of dual phase steels are primarily determined by not only the intercrit-ical annealing conditions but by the starting microstruc-ture before intercritical annealing.Since identical intercritical annealing conditions were used in this study,the microstructures before and after intercritical annealing are presented in this section.3.1.1.Before intercritical annealingOptical microstructures of DP0steel before and after ECAP are shown in Fig.1.The microstructure before ECAP(Fig.1(a))consisted of approximately20vol% pearlite(dark patches)with the remainder ferrite.Both ferrite and pearlite were nearly equiaxed and their size was about30l m.This microstructure was processed to coarse grained dual phase steel(CG-DP0steel)later. After ECAP(Figs.1(b)and(c)),pearlite colonies were severely deformed and fragmented into smaller ones with very irregular shapes.Macroscopically,this re-sulted in a decrease of inter-pearlite colony spacing.This ECAPed DP0steel was processed to UFG-DP0steel la-ter.Steels containing V exhibited similar macroscopic evolution after ECAP to that observed in DP0steel.TEM microstructures of DP0steel before and after ECAP,and DP1steel after ECAP are compared in Fig.2.Before ECAP,ferrite of DP0steel(Fig.2(a)) exhibited a moderate density of dislocations,and pearl-ite(Fig.2(b))consisted of a well-developed lamellar structure.After ECAP at500°C,ferrite grains of DP0(Fig.2(c))and DP1(Fig.2(e))steels were signifi-cantly refined to0.2–0.5l m with a dense dislocation structure.However,the ferrite grain size and disloca-tion density of DP1steel containing0.06wt%V were smaller and higher,respectively,in comparison to those of DP0steel.This may be attributed to the V addition effect retarding dynamic recovery at500°C.Cementite lamellar of DP0(Fig.2(d))and DP1(Fig.2(f))steels were severely deformed and became thinner with ECAP,and some extent of spheroidization of cementite occurred as indicated by arrows.3.1.2.After intercritical annealingMicrostructures of CG-DP0and UFG-DP0steels are shown in Fig.3.As similar to conventional dual phase steels prepared by intercritical annealing,the two main types of martensite morphology[10–12]were observed in CG-DP0steel(Fig.3(a)):(a)large irregular volumes of martensite transformed from austenite which was formed by pearlite decomposition during intercritical annealing,so residing at the former pearlitecolonies Fig. 1.Optical micrographs of DP0steel before and after ECAP (route C,four passes,500°C):(a)before ECAP;(b)after ECAP;and (c)an enlarged micrograph of(b).Y.I.Son et al./Acta Materialia53(2005)3125–31343127[10,11],and (b)an incomplete martensite network along ferrite–ferrite boundaries (indicated by arrows),which was associated with local manganese segregation rather than pearlite decomposition [11,12].By contrast,under the identical intercritical annealing conditions,the microstructure of UFG-DP0steel (Fig.3(b))consisted of equiaxed UFG ferrite and uniformly distributed mar-tensite islands.In addition,the martensite island sites were not confined to the former pearlite colonies.An inspection of enlarged SEM micrographs of UFG-DP steels (Fig.4)revealed that martensite in these steels was in an isolated blocky shape and occupied the whole volume of an UFG ferrite grain.Blocky martensite was often observed in dual phase steels prepared by step quenching in which austenite was directly quenched to the intercritical temperature [11,13].Martensite substructure of UFG-DP0steel (Fig.5(a))consisted of fine laths,typical in low carbon steels.Sim-ilar to conventional dual phase steels,a very high den-sity of dislocations existed in the ferrite grains adjacent to martensite (Fig.5(b)).Most of these dislocations are believed to be generated for transformation accom-modation during quenching rather than the ones in-duced by ECAP,since the latter were annihilated by recovery during intercritical annealing.Other UFG-DP steels,i.e.,UFG-DP1and UFG-DP2,exhibited similar substructures.It is known that rapid strain hard-ening of conventional dual phase steels is attributable to the high density of mobile dislocations.Accordingly,the similar strain hardening behavior is expected to occur in the present UFG-DP steels,in contrast to UFG ferrite–pearlite steels,in spite of their UFG structure.The strain hardening behavior of UFG-DP steels will be discussed later.The microstructural parameters of CG-DP0and UFG-DP steels are presented in Table 2.UndertheFig.2.TEM micrographs of the steels before and after ECAP (route C,four passes,500°C):(a)ferrite and (b)pearlite of DP0steel before ECAP;(c)ferrite and (d)pearlite of DP0steel after ECAP;(e)ferrite and (f)pearlite of DP1steel after ECAP.3128Y.I.Son et al./Acta Materialia 53(2005)3125–3134identical chemical composition and intercritical anneal-ing conditions:(a)the martensite volume fraction (V m )of UFG-DP0steel was higher than that of CG-DP0steel;(b)the V addition increased V m in UFG-DP steels,that is,V m of UFG-DP1and UFG-DP2steels was high-er than that of UFG-DP0steel;(c)the size of martensite islands (d m )and ferrite grains (d f )of UFG-DP steels was about 1l m but both seemed to increase slightly by the V addition as shown in Fig.4.For CG-DP0steel,d m was much smaller than d f .3.2.Tensile properties3.2.1.Mechanical dataThe representative engineering and true stress–strain curves and the corresponding tensile data are presented in Fig.6and Table 2,respectively.An inspection of Fig.6and Table 2revealed several findings.First,in spite of an UFG structure of $1l m,all the present UFG-DP steels exhibited the tensile behavior similarto that of coarse grained dual phase steels,i.e.,rapid strain hardening at the initial plastic deformation stage,continuous yielding,low yield ratio,etc.It shouldbeFig.3.Micrographs showing the shape and distribution of martensite in:(a)CG-DP0steel and (b)UFG-DP0steel.Fig.4.SEM micrographs showing the size and shape of martensite in the UFG-DP steels:(a)UFG-DP0steel;(b)UFG-DP1steel;and (c)UFG-DP2steel.Y.I.Son et al./Acta Materialia 53(2005)3125–31343129recalled that UFG ferrite–pearlite steels show a very high yield ratio associated with negligible strain hard-ening [3,4].Second,the strength of UFG-DP0steel was much higher than that of CG-DP0steel.It is quite noticeable that,in spite of much higher strength,the former showed not only similar uniform elongation but also larger total elongation compared to the latter.Third,the stiffness of UFG-DP0steel was higher than that of CG-DP steel.Finally,the addition of V im-proved strength a little,but did not degrade uniform and total elongations.3.2.2.Strain hardening behaviorAs mentioned previously,one of the unique mechan-ical characteristics of dual phase steels is rapid strain hardening at the initial plastic deformation stage.Since strain hardening is directly associated with formability,the strain hardening behavior of dual phase steelshasFig.5.(a)TEM micrograph showing the substructure of martensite in UFG-DP0steel.(b)TEM micrograph showing a high density of dislocations in UFG ferrite in the vicinity of martensite in UFG-DP0steel.Table 2The microstructural and tensile characteristics of the dual phase steels Designation V m (%)d m (l m)d f (l m)r YS (MPa)r TS (MPa)e u (%)e f (%)CG-DP0229.819.45108439.813.5UFG-DP0280.80.85819789.317.6UFG-DP135 1.10.9540104411.518.1UFG-DP2321.11.2565101510.416.6V m ,martensite volume fraction;d m ,martensite island size;d f ,ferrite grain size;r YS ,yield strength;r TS ,ultimate tensile strength;e u ,true uniform strain;e f ,engineering totalelongation.Fig.6.(a)Engineering stress–strain curves and (b)true stress–strain curves of the steels tested at an initial strain rate of 1·10À3s À1.3130Y.I.Son et al./Acta Materialia 53(2005)3125–3134been studied extensively.Of those studies,the three empiricalfitting equations and the mathematical analy-ses based on these equations are used the most com-monly:they are the Hollomon analysis[14],the Crussard–Jaoul(C–J)analysis[15,16]based on the Lud-wik equation[17],and the modified Crussard–Jaoul analysis[18]based on the Swift equation[19].In these equations,the true stress and true strain are expressed in the form of the power relationship which gives a mea-sure of strain hardening capability,i.e.,strain hardening exponent.However,the validity of the three analyses differs from materials to materials.So,in this section, the applicability of the three equations and the corre-sponding analyses to the presentation of the strain hard-ening behavior of the present UFG-DP steels is examined atfirst.The three equations are:Hollomon equation:r¼k H e n H;ð1aÞLudwik equation:r¼r0þk L e n L;ð1bÞSwift equation:e¼e0þk S r m;ð1cÞwhere r and e are the true stress and strain,respectively, n and m are strain hardening exponents,and other parameters are material constants.The differentiation of the logarithmic form of Eqs.(1a)–(1c)with respect to e provides the following relationships,respectively, so that the strain hardening exponents can be obtained by a linear regression.Hollomon analysis:n H¼dðln rÞ=dðln eÞ;ð2aÞC–J analysis:lnðd r=d eÞ¼ðn LÀ1Þln eþlnðk L n LÞ;ð2bÞmodified C–J analysis:lnðd r=d eÞ¼ð1ÀmÞln rÀlnðk S mÞ:ð2cÞThe three analyses were applied to the stress–strain data of the present steels and the results are shown in Fig.7.The Hollomon analysis(Fig.7(a))failed to show the linearity of ln r–ln e over the entire uni-form strain range as predicted by Eq.(2a).In the C–J analysis(Fig.7(b)),all the data including CG-DP0steel coalesced together.It indicates that the C–J analysis has a lack of ability differentiating the strain hardening behavior of dual phase steels having the different microstructural parameters such as the size,distribution and volume fraction of ferrite and martensite[20].By contrast,the modified C–J analysis (Fig.7(c))was shown to be sensitive to microstruc-tures,i.e.,differentiating the data between CG-DP0steel and other UFG-DP steels,and exhibited the lin-earity of ln(d r/d e)–ln r,as predicted by Eq.(2c),with two stages whose slope gives the value of(1Àm).Accordingly,the modified C–J analysis was selected to explain the strain hardening behavior of the presentsteels.Fig.7.The plots for the analysis of the strain hardening behavior of the steels:(a)ln r vs.ln e(the Hollomon analysis);(b)ln(d r/d e)vs.ln e (the Crusaard–Jaoul analysis based on the Ludwik equation);(c)ln(d r/d e)vs.ln r(the modified Crusaard–Jaoul analysis based on the Swift equation).Y.I.Son et al./Acta Materialia53(2005)3125–31343131The values of the strain hardening exponent of the two stages,m I and m II,and the transition strain(e tr)be-tween the deformation stages based on the modified C–J analysis are listed in Table3.At thefirst stage with m I ranging 4.4–6.5,CG-DP0steel and UFG-DP0steel showed the lowest and highest m I value,respectively, and UFG-DP1and UFG-DP2steels had the intermedi-ate values.This indicates that strain hardening at the initial plastic deformation stage was more rapid in UFG-DP steels compared to CG-DP0steel.However, all the steels exhibited similar strain hardening behavior at the second stage with m II=8.0–8.8.The present m I and m II values are comparable to those reported for conventional coarse-grained dual phase steels having similar chemical composition and V m[20–23].It was also found that the transition strain,e tr,from thefirst stage to the second stage of UFG-DP steels was smaller than that of CG-DP0steel.4.Discussion4.1.Formation of uniform UFG dual phase structureIn order to obtain the uniformly distributed UFG martensite islands shown in Fig.3(b),it is necessary that the C content in ferrite remote from pearlite colonies, i.e.,at least a half of the inter-pearlite colony spacing, reaches the equilibrium C content in austenite at the intercritical annealing temperature by C diffusion during or after ECAP.This can be achieved with the aid of microstructures evolved during ECAP.The microstruc-ture of ECAPed low carbon steels is characterized by subdivided pearlite colonies,UFG ferrite grains with a high density of lattice dislocations,and a large number of low-and high-angle boundaries with extrinsic bound-ary dislocations[24,25].More importantly,pearlitic cementite dissolution occurs during ECAP[26–28]. Then,at ECAP and intercritical annealing temperatures, the dissolved C atoms easily diffuse away from pearlite colonies along various paths generated by ECAP,i.e., a very high density of lattice dislocations,and a large number of low-and high-angle boundaries,etc. Although no quantitative data for cementite dissolution during ECAP are at present available,indirect informa-tion can be obtained from the severely cold drawn pearl-itic steel wire[29–33]since(a)the reported maximum drawing strain,$5,was comparable to the present ECAP strain,$4,and(b)the morphology evolution of pearlitic cementite during ECAP[24,25]was similar to that observed in the severely cold drawn pearlitic steel wire.The previous investigations[29–33]reported that20–50%of pearlitic cementite dissolved during cold drawing and resultantly the C content in ferrite reached over$4 at.%.From Figs.2(d)and(f)showing dissolved pearlitic cementite lamellar and the present ECAP strain,a sim-ilar C content is expected to exist in pearlitic ferrite and in UFG ferrite in the vicinity of pearlite colonies after ECAP.Then,as a veryfirst approximation,the distribu-tion of the C content in ferrite matrix of the present ECAPed steels can be estimated by applying the diffu-sion solution for a pair of semi-infinite solids[34],i.e., Cðx;tÞÀC0C1ÀC0¼1Àerfxffiffiffiffiffiffiffiffi4Dtpð3ÞFor the present case,C(x,t)is the C content in ferrite at the diffusion time t and the C diffusion distance x from pearlite colony,C0is the C content in ferrite re-mote from pearlite colony,C1is the C content super-saturated in pearlitic ferrite,and D is the C diffusivity in ferrite at the ECAP temperature.The use of Eq.(3) may not lead to a significant erroneous approximation since C1is not expected to vary significantly for short time diffusion:that is,the C atoms are continuously supplied from remaining severely deformed cementite lamellar shown in Figs.2(d)and(f).The following values were used for estimation:C0=0.09at.%(the equilibrium C content in ferrite at500°C), C1=4.0at.%(the reported C content supersaturated in pearlitic ferrite of severely cold drawn pearlitic steel [33]),and D=0.62·10À6exp(À80400(J/mol)/kT) (m2/s)(carbon diffusivity in a-ferrite[35]).The profile of the C content at various diffusion times is shown as a function of the distance from pearlite colony in Fig.8.As a reference,the equilibrium C content in austenite of the present steels at the intercritical annealing temperature estimated by a thermodynamic simulation[36](Table4)is also presented in Fig.8. The C content reaches the equilibrium C content in austenite at the intercritical annealing temperature at about10,23,and33l m away from pearlite colonyTable3The values of strain hardening exponent(m I and m II)and thetransition strain(e tr)between the deformation stages of the steels in theSwift equation estimated by the modified C–J analysisDesignation m I(first stage)m II(second stage)e tr(%)CG-DP0 4.48.0 3.7UFG-DP0 6.58.8 2.8UFG-DP1 5.08.2 2.4UFG-DP2 5.38.7 2.4Table4Equilibrium carbon concentration in austenite and austenite volumefraction of the steels at730°C estimated by a thermodynamicsimulation[36]Grade C eq(wt%)C eq(at.%)V c(%)DP00.524 2.3827.4DP10.518 2.3626.9DP20.509 2.3225.1 3132Y.I.Son et al./Acta Materialia53(2005)3125–3134for 1,5,and 10min,respectively.These distances are comparable to or larger than a half of the inter-pearlite colony spacing in the present ECAPed steels (Fig.1(b)).ECAP processing time in this study was 10–15min for four passes at 500°C with a plunge speed of 2mm/s and the sample length of 130mm.Under these circumstances,ferrite !austenite trans-formation is likely to occur at the preferential sites throughout ferrite matrix,i.e.,not be confined to the former pearlite colonies,by subsequent intercritical annealing,and therefore uniformly distributed UFG martensite islands are formed by quenching.The ac-tual diffusion kinetics would be much faster than the present approximation due to a high density of lattice defects induced by ECAP.V m of UFG-DP0steel was comparable to the equilib-rium austenite volume fraction,V c ,predicted from ther-modynamic simulation [36]while that of CG-DP0steel was less than the predicted V c (Tables 2and 4).This indicates that the present intercritical annealing time of 10min was not sufficient for CG-DP0steel to com-plete ferrite !austenite transformation.But,for UFG-DP0steel,in spite of insufficient transformation time,a high density of the austenite nucleation sites ex-isted due to a highly defected UFG structure and re-sulted in higher V m ,close to the equilibrium value,than CG-DP0steel.It is also worth noting that V m of UFG-DP steels containing V,i.e.,UFG-DP1and UFG-DP2,was higher than that of UFG-DP0steel without V and even than the predicted one.As seen in a comparison between Figs.2(c)and (e),the defect den-sity and grain size of V containing UFG-DP steels were higher and smaller,respectively,compared to those of UFG-DP0steel,providing more population of the aus-tenite nucleation sites during intercritical annealing of these steels.parison of the tensile characteristics between CG-DP0steel and UFG-DP0steel4.2.1.Strength and ductilityThe strength of dual phase steels generally increases with decreasing d f ,obeying the Hall–Petch relation [37],and increasing V m [38].Although the martensite strength associated with chemical composition of aus-tenite at intercritical annealing temperature affects the overall strength of dual phase steels [38],this effect would not be significant in the present CG-DP0steel and UFG-DP0steels due to the identical chemical com-position and intercritical annealing temperature.Accordingly,the higher strength of UFG-DP0steel is mainly attributed to UFG ferrite and larger V m .In spite of higher strength,elongation of the present UFG-DP0steel was comparable to that of CG-DP0steel.Uniform elongation of dual phase steels showing the two deformation stages,as shown in Fig.7(c),de-pends on strain hardening of both stages (i.e.,m I and m II )but the second stage strain hardening (m II )dominates the overall uniform elongation [20–23].Accordingly,the similar values of m II of CG-DP0and UFG-DP0steels (Table 3)are responsible for compara-ble uniform elongation of these two steels.4.2.2.Strain hardeningThe two stage hardening behavior in the modified C–J analysis is often observed in conventional dual phase steels.At the first stage with the low slope,the ferrite matrix deforms plastically but the martensite remains elastic and,at the second stage with the high slope,both phases deform plastically [20].As described in Section 3.2.2,strain hardening of UFG-DP0steel differed from that of CG-DP0steel by more rapid strain hardening at the first stage and lower transition strain (e tr )from the first stage to the second stage.As shown in Fig.3,the inter-martensite spacing of CG-DP0steel was much larger than that of UFG-DP steels.In such a case,for CG-DP0steel,mobile dislo-cations associated with transformation accommodation existed in ferrite only in the vicinity of martensite,and ferrite remote from martensite was relatively free of dislocations,causing low strain hardening at the first stage in which only ferrite deforms plastically.By con-trast,a uniform,high density of mobile dislocations is anticipated to exist throughout the ferrite matrix of UFG-DP0steel having a much finer inter-martensite spacing of an order of its size,leading to rapid strain hardening.In addition to rapid strain hardening at the first stage,the inherent high flow stress associated with UFG ferrite resulted in a large amount of load transfer to martensite.Therefore,if the martensite strength is the same,plastic deformation of martensite started earlier in UFG-DP0steel than CG-DP0steel,shifting e tr to the lowerstrain.Fig.8.The C content profile in UFG ferrite as a function of the distance from the pearlite colony during the present ECAP under the assumption of a diffusion couple of semi-infinite solids.Y.I.Son et al./Acta Materialia 53(2005)3125–31343133。
