Microstructure evolution of Al-4Cu-Mg alloy during semi-solid treatment

多尺度铝合金变形组织演变建模研究进展1王冠1,2，卞东伟1，寇琳媛1，易杰2，刘志文2，李落星2（1.宁夏大学机械工程学院，银川750021；2.湖南大学，汽车车身先进设计制造国家重点实验室，长沙410082；）摘要：铝合金在热成型过程中，微观组织会发生晶粒长大、晶粒不均匀变形、动态再结晶等一系列复杂的演化，而这些材料内部微观结构的改变，会直接影响到铝合金的综合性能。

通过掌握变形过程中微观组织演变的物理本质，来达到控制微观组织及产品性能的目的，已经越来越受到材料研究者的重视。

本文综述了铝合金变形组织演变建模的研究现状，重点介绍了多尺度模拟方法，同时指出了研究中存在的问题，展望了铝合金变形组织演变建模的发展趋势。

关键词：铝合金；微观组织演变；多尺度建模；热压缩变形；Research Progress in Multi-scale modelling of microstructure evolution during hot deformation ofaluminum alloyWANG Guan1,2, BIAN Dong-wei1, KOU Lin-yuan1, YI Jie2, LIU Zhi-wen2, LI Luo-xing2(1.College of Mechanical Engineering, Ningxia University, Yinchuan 750021, China;2.State Key Laboratory of Advanced Design and Manufacture for Vehicle body, Hunan University,Changsha 410082;)Abstract:During the hot forming process of aluminum alloy, microstructure will occur in a series of complex evolution such as grain growth, inhomogeneous deformation, dynamic recrystallization and which will directly affect the comprehensive properties of aluminum alloy. By mastering the physical essence of the microstructure evolution during heat deformation, to achieve the purpose of controlling the microstructure and the properties of the products has been paid more and more attention by the researchers of materials. This paper summarizes the research status quo of modelling of microstructure evolution during hot deformation of aluminum alloy, especially for the multi-scale simulation method, and points out the problems existing in current research and forecast the development trend of modelling of microstructure evolution during hot deformation of aluminum alloy.Key words: Aluminum alloy; Microstructure evolution; Multi-scale modelling; Hot compression deformation；铝合金具有密度低、比强度高、耐腐蚀性好、可循环利用等优点，被公认为汽车轻量化的理想材料。

刍议温度对铝合金蠕变行为的影响摘要：在不同温度下实施了蠕变温度实验，并从微观角度出发，重点探讨了蠕变温度对铝合金的蠕变应力变化及变形情况。

结果表明在125-200摄氏度的蠕变过程中，铝合金蠕变寿命都接近100小时的情况下，蠕变应力会随着蠕变温度的逐渐升高而表现出明显下降的趋势。

在125-175摄氏度蠕变过程中，合金的变形机制突出反映在大量位错在晶内的滑移。

关键词：温度铝合金蠕变应力变形近年来，随着我国航空航天事业的迅速发展及相关技术的不断进步，材料耐高温性能也得到了大幅度提升。

在巡航速度情况下，涂漆蒙皮的温度一般可以达到99—157摄氏度之间，而在紧急条件下，短时俯冲状态可能达到204摄氏度的高温。

长期处于热环境下的蠕变作用就会在一定程度上对飞机的材料结构带来影响。

所以加强温度对铝合金蠕变行为的影响不仅具有一定的学术价值，也具有重要的现实意义。

一、实验准备及方法二、实验结果及讨论（一）蠕变温度与铝合金蠕变应力之间的关系。

实验结果如图1所示，铝合金在四种温度下的蠕变寿命分别达到了106、96、100以及115小时。

从图中我们不难看出，当蠕变寿命都临近于100 小时时，蠕变应力随着蠕变温度的逐步提高出现了明显下降。

和125 摄氏度时相比，在蠕变温度为150 摄氏度的环境下，蠕变应力下降幅度不高，为为9.3%左右；但是当蠕变温度为175 摄氏度时，蠕变应力开始出现了较为显著的下降，为30.3%左右；当蠕变温度达到200 摄氏度时，下降幅度为45.8%。

（二）蠕变温度和合金变形机制之间的关系。

在四种蠕变温度情况下，合金断口周围平行于轧制面的金相组织情况如图2所示。

从图中我们可以看出，在125摄氏度合金的断口处，晶内开始存在有波浪状的滑移带，从（a）（b）中我们可以看到滑移带除了存在着互相平行的单向滑移带，还存在部分相交的交叉滑移带，但是整体来看以单向滑移带为主。

在175摄氏度环境下，如（c）所示，晶内交叉滑移带开始出现显著增加，滑移带也明显增粗，同时间距也更宽。

6061铝合金中富铁相在均匀化过程中的相变机理杜鹏;闫晓东;李彦利;沈健【摘要】采用金相(OM)、扫描电镜(SEM)、能谱(EDS)和透射电镜(TEM),研究6061铝合金中富铁相在均匀化过程中的转变和析出行为.结果表明:Mn元素直接参与6061铝合金中富铁相的相变过程,使富铁相由板条状的β-AlFeSi相转变成颗粒状的α-Al(FeMn)Si相,在560℃未发现明显的β-Al5FeSi→α-Al8Fe2Si的相变过程;在均匀化过程中,析出块状Al8Fe2Si相和颗粒状Al167.8Fe44.9Si23.9相,其中,Al167.8Fe44.9Si23.9相的析出速度受β-Al5FeSi→α-Al8Fe2Si的相变过程影响.%The phase transformation and precipitation behavior of iron-rich phase of 6061 aluminum alloy during the homogenization were investigated by optical microscopy(OM), scanning electronmicroscopy(SEM), energy dispersive spectrum(EDS) and transmission electron microscopy (TEM). The results show that the element of Mn is directly involved in the phase transformation of iron-rich phase during the homogenization, which makes the needle shaped β-AlFeSi phase transform into particle shaped α-Al(MnFe)Si phase. There is not evident phase transformation of β-Al5FeSi phase → α-Al8Fe2Si phase at 560 ℃. Block shaped phase Al8Fe2Si and granular shaped phaseAl167.8Fe44.9Si23.9 precipitate during the homogenization, and the precipitate rate of Al167.8Fe44.9Si23.9 phase is affected by the transition process ofβ-Al5FeSi phase → α-Al8Fe2Si phase.【期刊名称】《中国有色金属学报》【年(卷),期】2011(021)005【总页数】7页(P981-987)【关键词】6061铝合金;富铁相;相变;均匀化【作者】杜鹏;闫晓东;李彦利;沈健【作者单位】北京有色金属研究总院,北京,100088;北京有色金属研究总院,北京,100088;北京有色金属研究总院,北京,100088;北京有色金属研究总院,北京,100088【正文语种】中文【中图分类】TG166.36061铝合金作为一种中等强度铝合金，因其具有良好的塑性、耐蚀性和着色性能而广泛应用于建筑装饰、交通运输和航空航天等领域。

高纯无氧铜的晶界迁移行为及其晶粒长大机制高纯无氧铜的晶界迁移行为及其晶粒生长机制1. 引言高纯无氧铜是一种重要的工程材料，具有良好的导电性和热导性。

在制造电子设备、电力传输系统和化学工艺装备等领域具有广泛的应用。

高纯无氧铜的性能主要由其晶界迁移行为和晶粒生长机制决定。

本文旨在探讨高纯无氧铜的晶界迁移行为及其晶粒生长机制。

2. 高纯无氧铜的晶界迁移行为晶界迁移是指晶界位置在固态材料中发生改变的过程。

高纯无氧铜中，晶界迁移由两个主要因素驱动：体动力学效应和力学应力。

体动力学效应是指晶界迁移是由于原子在固态材料中的扩散运动，主要受温度和时间的影响。

力学应力是指晶界迁移是由于外部应力的作用，如热循环等。

晶界迁移过程中，晶界位置的变化使得晶粒的形状和尺寸发生改变。

3. 高纯无氧铜的晶粒生长机制晶粒生长是指晶体中的晶粒逐渐增长并形成较大晶粒的过程。

在高纯无氧铜中，晶粒生长的主要机制有两种：晶界扩散和气液固相变。

晶界扩散是指晶界附近的原子扩散，使得晶界迁移速率增加并促进晶粒生长。

气液固相变是指在高纯无氧铜中气体的溶解和析出，从而引发晶界迁移和晶粒生长。

4. 高纯无氧铜晶界迁移行为的研究方法为了研究高纯无氧铜的晶界迁移行为，研究者使用了多种实验方法和理论模型。

实验方法包括金相显微镜观察、原子力显微镜观察、电子背散射衍射等。

这些实验方法可以直接观察晶界的迁移过程和晶粒的生长过程。

理论模型主要是基于晶界迁移的动力学模型，如弥散选择模型和非饱和模型。

5. 高纯无氧铜晶粒生长机制的研究方法高纯无氧铜晶粒生长机制的研究主要使用了相场模型和分子动力学模拟。

相场模型是通过数学模拟晶粒长大的过程，可以研究晶粒的形状和尺寸变化。

分子动力学模拟是通过计算原子之间的相互作用力和位移，模拟晶粒生长的过程。

这些模拟方法可以预测晶粒长大的趋势和速率。

6. 结论通过对高纯无氧铜晶界迁移行为及其晶粒生长机制的研究，我们可以更好地理解并控制高纯无氧铜的性能。

第47卷第5期燕山大学学报Vol.47No.52023年9月Journal of Yanshan UniversitySept.2023㊀㊀文章编号:1007-791X (2023)05-0398-13金属基自润滑复合材料固体润滑剂研究进展邹㊀芹1,2,王㊀鹏1,徐江波1,李艳国2,∗(1.燕山大学机械工程学院,河北秦皇岛066004;2.燕山大学亚稳材料制备技术与科学国家重点实验室,河北秦皇岛066004)㊀㊀收稿日期:2022-05-25㊀㊀㊀责任编辑:唐学庆基金项目:丹凤朝阳人才支持计划(丹人才办[2019]3号);河北省高等学校科学研究重点项目(ZD2021099)㊀㊀作者简介:邹芹(1978-),女,安徽淮北人,博士,教授,博士生导师,主要研究方向为超硬及特种陶瓷材料㊁摩擦磨损;∗通信作者:李艳国(1978-),男,河北唐山人,博士,副研究员,主要研究方向为金属基复合材料,Email:lyg@㊂摘㊀要:固体润滑剂在金属基自润滑复合材料中的应用正在迅速增加,特别是在极端环境(高温㊁高负载等)条件下工作的耐磨材料㊂目前,金属基自润滑复合材料中常使用的固体润滑剂主要有无机层状固体润滑剂㊁金属及其化合物㊁MAX 金属陶瓷㊁有机物固体润滑剂㊁碳纳米材料固体润滑剂㊁多元复合固体润滑剂等,其种类很多,且各自有其适用的环境和基体㊂根据基体材料以及工况环境选择相匹配的固体润滑剂,可以保证金属基自润滑复合材料具有良好的减摩耐磨效果㊂针对上述内容,本文综述了金属基自润滑复合材料采用的固体润滑剂种类㊁基本性质㊁优缺点㊁润滑机理,总结了固体润滑剂的适用温度及其在金属基自润滑复合材料中的应用情况,并对金属基自润滑复合材料固体润滑剂的发展趋势进行了展望㊂关键词:金属基自润滑复合材料;固体润滑剂;润滑机理;研究进展;展望中图分类号:TB331㊀㊀文献标识码:A㊀㊀DOI :10.3969/j.issn.1007-791X.2023.05.0030㊀引言固体润滑剂[1]是金属基自润滑复合材料的重要组成部分,在金属基自润滑复合材料中的应用具有很长的历史㊂早在19世纪初期[2-3],石墨和Pb 已经作为润滑剂用于低速运转的机器上㊂20世纪30年代,添加固体润滑剂的铁基自润滑轴承在德国出现㊂20世纪60年代,添加MoS 2的金属基自润滑复合材料逐渐产生,并对超音速飞机的问世起到了重要的推动作用[4]㊂到目前为止,由于固体润滑剂可在一些特殊工况下(见表1)起润滑作用,这对高新技术的发展起到了重要的推动作用[5]㊂金属基自润滑复合材料固体润滑剂种类很多,包括无机层状固体润滑剂㊁金属及其化合物㊁MAX 金属陶瓷㊁有机物固体润滑剂㊁多元复合固体润滑剂等,其各有优缺点,且仍处于不断发展阶段㊂表1㊀固体润滑剂的适用场景Tab.1㊀Applicable scenaries of solid lubricants适用场景具体应用高负载滑动场景重型机械中的摩擦部件高温环境下磨损场景航空航天发动机㊁导弹燃油泵等摩擦部件强辐射环境下摩擦场景核电站㊁卫星等设备上的裸露活动部件强腐蚀性介质中摩擦场景化学反应器轴承,压缩机螺丝等部件摩擦接触表面导电场景电刷㊁受电弓滑板等灰尘或碎片环境中工作场景矿山机械和织机机械中的摩擦部件需要保证清洁的摩擦场景食品机械㊁纺织机械等摩擦部件微颤环境下的摩擦场景汽车和飞机上的摩擦部件1㊀无机层状固体润滑剂1.1㊀石墨石墨价格低廉,在潮湿环境中由于水的氢离第5期邹㊀芹等㊀金属基自润滑复合材料固体润滑剂研究进展399㊀子和氢氧根离子的饱和导致层间范德华键减弱,从而促进了层间分裂,在金属表面形成一层具有减摩作用的润滑膜[6],使得其可在潮湿环境提供有效润滑㊂目前,石墨作为金属基自润滑复合材料固体润滑剂的研究主要集中在改善不同钢种在不同工业应用中的摩擦磨损性能上,而制备时石墨与部分金属基体(Cu㊁Al等)润湿性较差,导致两者界面结合变差,影响复合材料的力学性能以及摩擦学性能,另外使用过程中产生的高温会导致石墨氧化和烧蚀,严重影响润滑效果[6-8]㊂对石墨进行金属化改性,如采用金属(Ni㊁Cu等)包覆石墨的办法,能有效改善石墨与基体的界面结合,同时防止石墨氧化和腐蚀,改善石墨高温润滑效果,从而提高复合材料摩擦学性能,扩大使用范围㊂张鑫等[9]采用Cu包覆石墨制备了Cu基粉末冶金摩擦材料,其材料表面形成的摩擦膜主要为氧化膜,而采用普通石墨时,由于材料表面较多的石墨会抑制氧化反应,会形成石墨膜,其对材料表面的保护效果不及氧化膜㊂但相对于原基体,两种材料摩擦性能均有明显提高㊂Zhao等[10]证明了石墨与青铜无法充分润湿,而加入Ni或Cu包覆石墨的复合材料可以明显提高石墨与基体的结合性,Ni包覆石墨青铜基材料具有更稳定的摩擦系数㊁更低的磨损率㊁更高的维氏硬度,包覆石墨的Ni也可以提高复合材料的耐蚀性㊂牛志鹏等[11]发现加入镀Ni石墨可以降低石墨与Al的润湿角,提高基体的力学性能,降低复合材料的摩擦系数和磨损率,使金相组织变得更加致密㊂但石墨表面光滑且亲水性差,难以实现完全包覆㊂罗虞霞等[12]发现,采用机械化整形处理石墨表面,可以获得更为完整的Ni包覆层㊂冀国娟等[13]发现,在石墨表面进行微氧化以及在化学包覆反应溶液中加入醇类表面活性剂,均可提高包覆率㊂综上,采用金属包覆石墨作为固体润滑剂可显著提高其高温润滑特性㊂然而,石墨表面包覆金属层的完整性是决定其润滑性能的关键因素㊂故进一步提高石墨表面包覆金属层的完整性以及连续性将继续成为研究的重点㊂1.2㊀BNBN导电性能强㊁热稳定性高,在大气环境中适用温度为500~800ħ,是高温自润滑材料的优良润滑剂㊂其润滑机理为[14-15]:高于500ħ时,BN 会在摩擦过程中剥落而转移到摩擦表面形成润滑膜,起减摩作用㊂蒋冰玉等[16]以Ni-Cr合金为基体材料,BN为固体润滑剂,制备出燃气轮机中减摩耐磨用的高温自润滑复合材料㊂目前,尽管BN 是一种人们熟知的高温固体润滑剂,但由于其存在有效性差㊁不可润湿等问题,使得人们对于BN 单独应用在金属基自润滑复合材料上的报道较少,其常与其他固体润滑剂协同润滑[17]㊂2㊀金属及其化合物2.1㊀金属常见的金属固体润滑剂有Pb㊁Al㊁Ag㊁Au㊁Sn㊁Bi㊁In等,其具有纯度高㊁原料易得㊁低温环境不会丧失润滑性能等优点㊂金属固体润滑剂在强辐射㊁真空㊁低温等极端工作条件非常适合作为金属基自润滑复合材料的固体润滑剂使用,常与Cu㊁Al㊁TiAl等金属基体组成复合材料㊂其润滑机理为:在摩擦热的作用下,由于热膨胀系数不同,金属逐渐从基体内扩散到摩擦表面形成润滑膜,起减摩作用,但其适用环境受温度限制严重㊂Yao等[18]发现,在200ħ时,Ag在剪切应力作用下扩散到摩擦表面,起减摩耐磨作用㊂但在600ħ时Ag完全失去润滑作用(图1)㊂Dong 等[19]发现,Cu-24Pb-x Sn合金的自润滑性能和力学性能随Sn含量的增加而增加,Pb含量的增加有效地削弱了以摩擦系数变化为特征的粘滑现象㊂李聪敏等[20]以Al-Cu-Mg合金为基体,添加低熔点组元Bi后合金抗咬合能力明显提升,发现带状富Bi 相涂覆在磨损表面,起到减摩自润滑作用㊂金属在强辐射㊁真空㊁低温等极端环境仍具有润滑特性,但是也存在着一些缺点,如:Pb本身有毒,对人体和环境都有危害,Ag㊁Au㊁In等金属作为固体润滑剂时成本太高;金属在空气中暴露的时间过长时,易发生氧化反应,影响润滑效果㊂2.2㊀金属氧化物常见的金属氧化物固体润滑剂有PbO㊁CuO㊁MoO3㊁SnO㊁ZnO等㊂金属氧化物是最早应用的高温固体润滑剂,常与Fe㊁Ni㊁NiAl等金属基体组成复合材料㊂由于金属氧化物具有较低的剪切强度,可有效避免摩400㊀燕山大学学报2023擦过程中的咬合现象㊂Peterson 等[21]考察了大量氧化物的高温摩擦学特性,发现PbO 等少数氧化物可实现较宽温度范围内的有效润滑㊂但是,由于PbO 危害环境,国外已限制其应用㊂Zhu 等[22]通过PM 制备了添加氧化物(ZnO /CuO)的NiAl-C-Mo 自润滑材料,发现氧化物在低温时几乎不起减摩作用㊂但当温度达到600ħ时,磨损表面形成了ZnO㊁CuO 和MoO 3层,表现出了良好的减摩耐磨效果㊂结果表明,金属氧化物在高温时润滑效果显著㊂但是,目前关于二组元氧化物的润滑机理还未得到统一㊂图1㊀TiAl 基自润滑复合材料磨损表面的微观结构演变示意图Fig.1㊀Schematic diagram of microstructure evolution of wear surface of TiAl based self-lubricating composite2.3㊀金属氟化物常见的金属氟化物固体润滑剂有CaF 2㊁BaF 2㊁LaF 3等㊂金属氟化物热稳定性良好,从500ħ到1000ħ的温度范围都能起到良好的减摩耐磨作用,其原因主要为金属氟化物在500ħ时经历了由脆性到塑性的转变㊂Longson [23]发现,CaF 2和BaF 2具有良好润滑性的原因是其在摩擦过程中由脆性向塑性转变以及氟元素与金属表面发生化学反应的共同作用㊂尽管对CaF 2和BaF 2润滑机理进行了大量研究,但是对于其转移润滑机理的全面认识还有赖于进一步研究㊂综上,由于金属氟化物特殊的润滑机制导致其在低温时不提供润滑,故单独采用金属氟化物作为金属基自润滑复合材料固体润滑剂的报道很少,其多与石墨㊁Ag 等固体润滑剂复合使用,达到宽温度范围有效润滑的目的㊂2.4㊀金属硫化物常见的金属硫化物固体润滑剂有MoS 2㊁WS 2㊁FeS㊁CrS 等㊂MoS 2属于六方晶系,具有层状结构,常与Fe㊁Al㊁Ag 等金属基体组成复合材料㊂MoS 2在大气环境中适用温度可达350ħ,润滑机理与石墨相似,由于具有低摩擦㊁低接触电阻等优点,广泛用作航空㊁航天机构中的滑动电接触材料[24]㊂WS 2因其良好的热稳定性和抗氧化性而广泛应用于高温环境㊂研究表明[25-27],在大气环境中通过在金属基体中掺入MoS 2或WS 2颗粒可显著提高Ni [25]㊁Al [26]㊁Fe [27]等金属基复合材料的摩擦学性能,使其满足使用要求㊂但是,MoS 2和WS 2会因大气湿度高㊁氧气的存在以及高温而导致润滑性能降低㊂通过掺杂金属或无定形碳可以保护MoS 2边缘位置免受氧化,从而提高MoS 2和WS 2在潮湿或较高温度条件下的摩擦学性能㊂Rigato 等[28]发现在MoS 2层状结构中掺杂Ti 增加了MoS 2层间距离,从而改善了其摩擦学性能㊂此外,研究发现,在MoS 2层状结构中掺杂Ni [29]㊁Cu [30]等金属可提高复合材料在潮湿环境和真空条件下的摩擦磨损性能㊂FeS 与MoS 2相比,具有优异的耐高温特性,因其较疏松的鳞片状结构能储存润滑油,可进一步提升润滑性能㊂尹延国等[31]发现FeS /Cu 基复合材料在在干摩擦过程中,FeS 颗粒聚集在摩擦表面形成一层硫化物固体润滑膜,具有较好的减摩㊁抗粘着作用,在油润滑条件下,润滑油膜和FeS 固体润滑膜可以起协同润滑作用㊂Lu 等[32]采用NiCr /Cr 3C 2和WS 2粉末在Ti 6Al 4V 基体上激光熔覆制备了Ti 2SC /CrS 自润滑耐磨复合涂层,由于原位合第5期邹㊀芹等㊀金属基自润滑复合材料固体润滑剂研究进展401㊀成的自润滑Ti2SC和CrS的存在,自润滑抗磨复合涂层显示出比不添加WS2粉末的抗磨复合涂层更好的摩擦学性能㊂综上,MoS2和WS2在高温真空条件下具有优良的润滑特性,被认为高温真空条件下的首选固体润滑剂㊂在大气环境中,温度低于350ħ时,金属基-MoS2自润滑材料表现出优异的摩擦学性能㊂但是,MoS2在大气环境中高温时容易发生氧化[29-30],限制了其应用环境㊂故如何进一步提高MoS2在潮湿和较高温度条件下的摩擦学性能将继续成为研究的重点㊂2.5㊀金属硒化物常见的金属硒化物固体润滑剂有NbSe2㊂NbSe2导电性能优异,相对摩擦系数低,常与Ag㊁Cu[33-34]等金属基体组成复合材料,广泛应用于电接触领域㊂早在20世纪80年代,美国NASA便采用Ag-NbSe2自润滑材料来制作卫星上的电刷,并取得良好效果㊂Ag-NbSe2自润滑材料具有良好润滑性能的原因[33]为在摩擦热和变形挤压的共同作用下,部分NbSe2转移到摩擦表面,形成了NbSe2润滑膜,起减摩作用㊂孙建荣等[34]发现,高负载㊁真空条件下,添加纤维状NbSe2的Cu-石墨复合材料摩擦系数远低于原复合材料㊂因此, NbSe2常作为真空条件下的固体润滑剂使用㊂3㊀MAX金属陶瓷MAX金属陶瓷因为其原子结构和独特的化学键特性,使MAX金属陶瓷兼具金属和陶瓷的优点,如高硬度㊁高弹性模量,具有良好的抗氧化性㊁耐腐蚀性㊁导电导热性㊁辐照性能㊁高温机械和摩擦学性能等[35]㊂理论计算约有600余种能稳定存在的三元MAX金属陶瓷,如今可以通过实验合成80多种[36],如Ti3SiC2㊁Ti3AlC2㊁Ti2AlC㊁Ti2AlN㊁Ta2AlC等㊂目前,除Ti3SiC2和Ti3AlC2外,对于其他MAX金属陶瓷应用于金属基自润滑复合材料的研究鲜有报道㊂在材料基体中添加一定量的Ti3SiC2/Ti3AlC2颗粒润滑相能够显著提升金属基体的摩擦学性能㊂研究表明[37-39]不同温度下的微观结构以及反应产物对Ti3SiC2㊁Ti3AlC2的润滑性能有重要的影响㊂Zou等[38]用放电等离子烧结制备Ti3SiC2增强TiAl基复合材料,Ti3SiC2均匀分布在TiAl基质中,部分分解形成Ti5Si3和TiC,室温摩擦时复合材料表面形成Ti3SiC2润滑膜,550ħ摩擦时形成Fe-Ti-Al-Si-氧化物润滑膜,起润滑作用㊂朱咸勇等[39]发现,当试验温度低于400ħ在轻载条件下难以形成稳定氧化物润滑膜,其润滑特性主要依赖于特殊的层状形貌,而试验温度超过500ħ会促使材料表面形成氧化物润滑膜,起到减摩耐磨的作用㊂同时,MAX金属陶瓷添加量对复合材料摩擦学性能影响较为显著㊂陈海吉[40]使用放电等离子烧结制备Ti3AlC2/Cu复合材料,研究表明,随着Ti3AlC2添加量增加,复合材料摩擦磨损性能得到提高㊂研究发现当含量过高时会导致其致密度降低,影响摩擦学性能㊂烧结温度对MAX金属陶瓷自润滑复合材料性能也有重要影响㊂Zhou等人[41]发现烧结温度在900ħ以上时,在Cu和Ti3SiC2界面会形成Cu㊁TiC x㊁Ti3SiC2和Cu x Si y混合区从而提高系统的润湿性和耐磨性㊂综上,MAX金属陶瓷应用在摩擦材料的大多数情况下,由于摩擦过程中形成的氧化物润滑膜具有特殊的层状结构,使复合材料润滑效果更好㊂另外,表面改性以及较高的烧结温度可进一步提高其润滑效果㊂4㊀有机固体润滑剂除上述固体润滑剂外,还有一类性能优越㊁可用于极端环境(真空㊁强辐射)条件下的单一固体润滑剂-有机固体润滑剂㊂有机固体润滑剂种类很多,如聚四氟乙烯(PTFE)㊁三聚氰胺氰尿酸盐(MCA)等,但较低的适用温度(-270~275ħ)限制了其在金属基复合材料中的应用㊂PTFE是所有聚合物中摩擦系数最低的[42]㊂其抗剪切强度较低,受剪切力时聚合物链脱开,可提供润滑作用㊂同时,由于含F外壳的存在,其抗咬合性优异,常采用电沉积法与Ni[43]㊁Fe[44]等金属基体组成复合材料㊂MCA润滑特性与MoS2相似,滑动面间极易受力断裂,提供润滑作用㊂Tang 等[43]发现,由于润滑转移层的存在,Ni-Co-PTFE 复合材料显示出良好的摩擦学性能(摩擦系数0.08)㊂Xiang等[44]则指出PTFE的低摩擦系数以及40Cr钢的高强度是40Cr钢-PTFE复合材料具有良好摩擦学性能的重要原因㊂但是PTFE的力402㊀燕山大学学报2023学性能较差,线膨胀系数大,故将PTFE用作固体润滑材料时通常要添加填充物对其进行改性或对金属基体进行阳极氧化处理[45]㊂魏羟等[46]用Pb 粉㊁石墨㊁玻璃纤维填充PTFE制成Cu基镶嵌型关节轴承材料,显示出较好的摩擦磨损性能㊂但李同生等[47]发现,与含铅PTFE镶嵌轴承相比,无铅PTFE镶嵌轴承在工作时所形成的润滑膜最为完整㊁均匀,耐磨性更好㊂同时,对金属基体进行阳极氧化处理改性可进一步提高PTFE与基体金属基体的附着性[45]㊂综上,添加填充物对PTFE进行改性或对金属基体进行阳极氧化处理可大大提高复合材料的机械和摩擦学性能㊂5㊀碳纳米材料固体润滑剂近年来,纳米技术的快速发展推动了金属基自润滑复合材料的开发,出现了新型碳纳米材料固体润滑剂,例如碳纳米管(CNTs)㊁石墨烯(GPLs)等㊂由于其尺寸小,容易进入摩擦接触区域,形成保护摩擦膜,产生自润滑效应㊂同时,界面以下的新型碳纳米材料还可以防止应力集中而引发的严重磨损㊂5.1㊀碳纳米管CNTs具有良好的润滑特性,被认为是金属基自润滑复合材料中石墨的替代品㊂在这方面,有相关报道称已经成功开发了用于汽车工业的CNTs-金属基自润滑复合材料[48]㊂Orowan环化机制以及CNTs与金属基体之间热膨胀失配所产生的位错在增强Al/Cu-CNTs复合材料中起着重要作用[49]㊂为达到预想的润滑效果,CNTs在基体中的均匀分布以及界面调控就显得尤为重要㊂对此,研究者们做了大量的工作㊂2004年,Noguchi等[50]开发了一种新方法制备复合材料,首先让CNTs均匀分布在弹性体基体内,然后用Al来置换弹性体基体,从而保证CNTs均匀分布在Al基体内㊂2019年,周川等[51]采用混酸处理㊁分子水平法结合行星球磨两步混合工艺成功制备出Cu-CNTs复合粉末㊂混酸处理将含O官能团成功引入CNTs表面,提高了CNTs与基体的界面结合㊂以上研究均表明,均匀分布的CNTs可显著提高材料的机械和摩擦学性能㊂5.2㊀石墨烯片GPLs是目前已知最薄㊁最硬㊁导电性能最好的材料,具有良好的润滑特性,同时,可以通过晶粒细化㊁位错强化和应力转移来提高复合材料强度[52]㊂在过去的十多年里,绝大多数报道均表明在基体中均匀分布且结合良好的GPLs能够明显改善金属基复合材料的摩擦学性能㊂但是,聚集状态的GPLs增强效果较差,与石墨薄片几乎无差别㊂研究表明[53-55],不同的因素(例如GPLs的类型㊁含量㊁基体材料㊁混料方法和球磨时间等)会显著影响GPLs在金属基体中的分散性㊂为了保证GPLs均匀地分散在基体中,部分研究者在粉体混合工艺中采用氧化石墨烯代替石墨烯,先得到均匀混合的氧化石墨烯/合金粉体,再通过氧化石墨烯的热还原性质得到高度均匀的还原石墨烯/合金粉体[56]㊂Bastwros等[53]则研究了球磨时间对GPLs增强Al基复合材料的影响㊂发现经过10 min球磨后的材料综合性能反而降低,而60min 球磨后GPLs均匀分散在到Al基体内,在摩擦学性能上,GPLs显示出了良好的增强效果㊂另一方面,化学镀和电化学沉积法制备金属包覆型碳纳米材料,也可以确保GPLs均匀地分散在基体中㊂李远军[55]通过化学镀将纳米铜颗粒负载于还原氧化石墨烯表面的方法来确保其在Cu基体上均匀分布㊂但研究表明,化学镀和电化学沉积法一般仅适用于Cu㊁Ni㊁Ag等电负性较低的金属基体㊂综上,碳纳米材料可显著提高材料摩擦学和机械性能㊂但是,CNTs严重团聚以及与基体结合不牢固会减弱增强效果,甚至导致材料失效㊁降低使用寿命,从而进一步增加制造成本,限制其在金属基自润滑复合材料上的广泛应用㊂这就对制造方法㊁材料尺寸大小以及空间分布提出来更为苛刻的要求,但是,由于弱的层间相互作用,碳纳米管㊁石墨烯在实现超滑方面有很大的潜力[57]㊂因此,目前研究者们对于碳纳米材料固体润滑增强金属基自润滑复合材料的研究也主要集中在这四方面:1)提高碳纳米材料在金属基复合材料中分散的均匀性;2)对碳纳米材料与金属形成的界面组织进行调控;3)掺杂其他固体润滑剂,进一步提高金属的减摩耐磨性能;4)微观尺度上,研第5期邹㊀芹等㊀金属基自润滑复合材料固体润滑剂研究进展403㊀究石墨烯对材料性能的作用机理㊂综上,单一固体润滑剂对使用环境具有选择性,无法实现宽温度范围(25~800ħ)以及多种环境下的有效润滑㊂常见单一固体润滑剂的性能及优缺点见表2[1-57]㊂表2㊀单一固体润滑剂性能及优缺点Tab.2㊀Performance and relative merits of single solid lubricant固体润滑剂适用温度/ħ摩擦系数μ优点存在的问题最新解决方法石墨-270~5500.05~0.3(大气中)廉价㊁减震性良好㊁可在潮湿环境提供有效润滑强度较低,仅在大气环境提供有效润滑对石墨粉末进行表面改性,如镍包覆石墨MoS2-270~3500.006~0.25(大气中)0.001~0.2(真空中)高温真空条件下稳定性优异大气环境易氧化失效掺杂金属或无定形碳BN500~8000.15~0.25(大气中)良好的高温固体润滑剂成本较高,低温润滑性差与低温固体润滑剂协同润滑Ag㊁Au-270~4000.08~0.2(大气中).0.08~0.15(真空中)导电性能优异在酸碱条件下无效,成本高与其他固体润滑剂协同润滑PbO200~6500.1~0.3(大气中)可实现宽温度范围有效润滑有毒物质,摩擦系数较高㊁且形成润滑膜易脱落已被其他固体润滑剂替代CaF2㊁BaF2㊁LaF3500~9000.2~0.4(大气中)可实现高温有效润滑低温润滑性差与低温固体润滑剂协同润滑MAX金属陶瓷400~8000.005(大气中)高温机械和摩擦学性能优异,导电性能良好与Fe等基体复合时,界面结合差,易脱落1)添加增强相;2)对Ti3SiC2㊁Ti3AlC2进行表面改性,如镀铜PTFE-270~2750.04~0.2(大气中)0.04~0.15(真空中)真空润滑性能优异,抗咬合性好300ħ以上失效,不耐高温㊁力学性能较差,线膨胀系数大1)添加填充物对PTFE进行改性;2)对金属基体进行阳极氧化处理碳纳米材料-270~5000.05~0.2(大气中)轻质,可显著提高复合材料机械学㊁摩擦学性能团聚以及界面结合严重影响润滑效果,生产成本高昂1)氧化石墨烯代替石墨烯;2)混酸处理;3)金属包覆碳纳米材料;4)掺杂其他固体润滑剂6㊀多元复合固体润滑剂早在20世纪60年代初,人们就已经发现,两种或者多种固体润滑剂混合使用时,由于不同固体润滑剂之间的协同作用,使得其润滑效果好于其中任何一种固体润滑剂单独作用㊂6.1㊀Ni基自润滑材料的多元复合固体润滑剂在过去的20年中,已经成功开发了一系列Ni 基的高温自润滑复合材料[58-62]㊂该类由Ni基体与固体润滑剂(Ag-BaF2/CaF2/LaF3-金属氧化物/无机盐)组成的自润滑复合材料,在很宽的温度范围(25~800ħ)和高强度(800ħ,500MPa的抗压强度)并存的情况下表现出优异的润滑性能(图2[59])㊂其良好的润滑特性(摩擦系数(0.23~ 0.34)和低磨损率(10-6~10-5mm3N-1m-1)解释为Ag㊁氟化物㊁无机盐的协同作用㊂当高于500ħ时,氟化物中的低共熔物从基体中逸出,发生由脆性到塑性的转变,可进一步提升润滑效果[60]㊂Zhen等[61]指出由于Ag膜的存在,真空环境中该类复合材料摩擦系数和磨损率均低于大气环境中的摩擦系数和磨损率,是一种很有潜力的航空㊁航天材料㊂此外Zhen等[62]的另一份研究表明,在Ag-BaF2-CaF2固体润滑剂的基础上再添质量分数为0.5%~1%的石墨可以使Ni基复合材料获得稳定的摩擦性能(摩擦系数(0.19~0.29)和磨损率(5.3ˑ10-6~2.3ˑ10-5mm3N-1m-1)㊂404㊀燕山大学学报2023图2㊀Ni 基自润滑复合材料的摩擦学性能Fig.2㊀Tribological properties of Ni basedself-lubricating composites6.2㊀Ni 3Al 基自润滑材料的多元复合固体润滑剂进一步研究表明[63-65],该类由Ni 3Al 基体与固体润滑剂(Ag-CaF 2-BaF 2)和增强材料(Cr,Mo 等金属元素)组成的自润滑复合材料,在从室温到1000ħ的宽温度范围内表现出低摩擦系数(μ<0.4)和低磨损率(10-6~10-4mm 3N -1m -1),且具有令人满意的机械性能(硬度>300HV,抗压强度>1000MP)㊂Zhu 等[65]采用热压烧结法制备的Ni 3Al-6.2BaF 2-3.8CaF 2-12.5Ag-20Cr 复合材料实现了室温到1000ħ的有效润滑(摩擦系数(0.24~0.37)和低磨损率(5.2ˑ10-5~2.3ˑ10-4mm 3N -1m -1))㊂Ni 3Al 基体良好的高温机械性能,Ag㊁氟化物㊁无机盐的协同润滑以及Cr 元素对基体的增强作用使得其可以实现更宽温度范围的有效润滑㊂与Ni 基自润滑复合材料相比,Ni 3Al 基自润滑复合材料则可实现更宽温度范围内的有效润滑,其润滑机理见图3[66]㊂6.3㊀TiAl 基自润滑材料的多元复合固体润滑剂近年来,由于航空㊁航天工业的需要,科研人员制备了一系列基于TiAl 基的高温自润滑复合材料[67-69]㊂该类由TiAl 基体与固体润滑剂(Ag-Ti 3SiC 2-BaF 2/CaF 2)组成的自润滑复合材料,具有硬度高(>500HV)㊁轻质(ρ<3.9g /cm 3)等优点㊂结果表明[66-68],Ag-Ti 3SiC 2-BaF 2-CaF 2润滑体系在宽温度范围内下具有良好的协同效应:低温时,银扩散到金属基体的摩擦表面形成了一层富Ag 的摩擦膜,高温时,由于BaF 2㊁CaF 2的挤压和Ti 的氧化,在摩擦表面形成了一层含氟化物和氧化物的摩擦膜㊂但是,从室温到800ħ的宽温度范围内其摩擦系数(μ>0.3)和磨损率(10-4mm 3N -1m -1)较高,摩擦学性能有待进一步提高㊂图3㊀宽温度范围内Ni 3Al 基自润滑复合材料的润滑机理Fig.3㊀Lubrication mechanism of Ni 3Al based self-lubricating composites in a wide temperature range㊀㊀综上,可得出:1)多元复合固体润滑剂的协同作用在宽温度范围内对改善复合材料的摩擦学性能起重要作用;2)选择高温机械性能优异的金属基体以及适当添加Cr㊁Mo 等金属元素可实现更宽温度范围的有效润滑;3)Ag 与氟化物/无机盐/MAX 金属陶瓷材料等高温固体润滑剂的组合具有极佳的协同润滑作用㊂6.4㊀Fe /Cu /Ag 等金属基自润滑材料的多元复合固体润滑剂㊀㊀人们对多元复合固体润滑剂对Fe [70-71]㊁Cu [72]㊁Ag [73]等金属基体性能影响也进行了大量研究㊂Li 等[71]发现以LaF 3和MoS 2作为润滑组元的Fe 基复合材料可显示出超低的摩擦系数(0.09),。

短语1 数值模拟：numerical simulation2 力学性能：mechanical property3 铝合金：aluminum alloy4 应力分析：stress analysis5 钛合金：titanium alloy6 表面处理：surface treatment7 电磁场：electromagnetic field8 抗拉强度：tensile strength9 晶粒细化：grain refinement10 工艺参数：process parameter11 有机合成：organic synthesis12 表面质量：surface quality13 定向凝固：directional solidification14 生产管理：production management15 制备工艺：preparation technology16 拉伸强度：tensile strength17 冷轧：cold rolling18 速度场：Velocity Field19 电子束：Electron beam20 ANSYS软件：ANSYS software21 电磁搅拌：electromagnetic stirring22 铸铁：cast iron23 隔振：vibration isolation24 动力学仿真：Dynamic Simulation25 铜合金：copper alloy26 离心铸造：centrifugal casting27 色差：color difference28 金属基复合材料：metal matrix composites29 应变速率：Strain Rate30 气力输送：pneumatic conveying31 压铸：Die Casting32 金属氧化物：metal oxide33 正电子湮没：Positron annihilation34 热效率：heat efficiency35 凝固组织：solidification structure36 界面反应：interfacial reaction37 模具设计：mold design38 置换通风：displacement ventilation39 镁合金：Mg alloy40 熔模铸造：Investment Casting41 高铬铸铁：high chromium cast iron42 电磁力：electromagnetic force 43 生产实践：production practice44 AZ91D镁合金：AZ91D magnesium alloy45 机械振动：mechanical vibration46 机械系统：mechanical system47 温差：temperature Difference48 传热模型：heat transfer model49 耐磨性能：wear resistance50 硅溶胶：silica sol51 生产系统：production system52 色散关系：dispersion relation53 超声振动：ultrasonic vibration54 知识表达：knowledge representation55 真空系统：Vacuum system56 工艺控制：process control57 TiAl合金：TiAl alloy58 离心力：Centrifugal force59 连续铸造：Continuous Casting60 液压控制：Hydraulic control61 球墨铸铁：nodular cast iron62 流变模型：rheological model63 时效处理：aging treatment64 小波网络：wavelet network65 软件包：software package66 弹簧钢：spring steel67 冷却速率：cooling rate68 铸钢：Cast steel69 水平连铸：horizontal continuous casting70 技术改造：technological transformation71 脉冲电流：pulse current72 凝固过程：Solidification Process73 气缸盖：cylinder head74 制备技术：preparation technology75 复合形法：Complex method76 工艺分析：process analysis77 动力学建模：dynamic modeling78 消失模铸造：Lost Foam Casting79 真空干燥：vacuum drying80 余热：waste heat81 系统控制：system control82 铝硅合金：Al-Si Alloy83 响应面分析法：Response surface methodology84 铸造工艺：casting process85 气缸套：cylinder liner86 SIMPLE算法：SIMPLE algorithm87 工艺优化：technology optimization88 流场：fluid field89 工艺过程：Technological process90 氮化硼：boron nitride91 精密铸造：investment casting92 热循环：thermal cycling93 表面缺陷：Surface defects94 节能技术：energy-saving technology95 低压铸造：Low Pressure Casting96 界面结构：interface structure97 铁水：hot metal98 Al-Cu合金：Al-Cu alloy99 AZ91镁合金：AZ91 magnesium alloy 100 凝固模拟：Solidification simulation101 碳酸钾：potassium carbonate102 等离子弧：plasma arc103 抗裂性：crack resistance104 模锻：die forging105 冲蚀磨损：erosion wear106 注射成形：injection molding107 热压缩变形：hot compression deformation108 激光淬火：laser quenching109 超声检测：ultrasonic inspection110 磨球：Grinding ball111 冷变形：cold deformation112 强韧化：strengthening and toughening 113 气泡：air bubble114 保温时间：holding time115 白口铸铁：white cast iron116 电磁铸造：electromagnetic casting117 断口形貌：fracture morphology118 氢含量：hydrogen content119 浇注温度：pouring temperature120 锥齿轮：bevel gear121 灰铸铁：gray iron122 喷丸：shot peening123 排气系统：exhaust system124 水玻璃：Sodium silicate125 挤压铸造：Squeezing Casting126 密度分布：density distribution127 渣浆泵：slurry pump128 分型面：parting surface 129 A356合金：A356 alloy130 静磁场：static magnetic field131 网格剖分：mesh generation132 电磁连铸：electromagnetic continuous casting133 快速制造：rapid manufacturing134 压铸模：die-casting die135 韧性断裂：ductile fracture136 ADAMS软件：ADAMS software137 弯曲变形：bending deformation138 缸体：cylinder block139 变频控制：frequency conversion control 140 热应力场：thermal stress field141 压铸机：Die Casting Machine142 TiNi合金：TiNi alloy143 碳当量：carbon equivalent144 析出相：precipitated phase145 保温材料：thermal insulation material 146 对甲苯磺酸：p-toluene sulphonic acid 147 组织性能：microstructure and property 148 半固态成形：Semi-solid Forming149 TC4合金：TC4 alloy150 疲劳破坏：fatigue failure151 熔池：molten pool152 超声处理：ultrasonic treatment153 阀体：Valve Body154 压缩变形：Compression Deformation 155 扩散层：Diffusion layer156 缸套：cylinder liner157 铸钢件：steel casting158 性能计算：Performance calculation 159 缸盖：cylinder head160 微波炉：microwave oven161 浇注系统：pouring system162 Al-Zn-Mg-Cu合金：Al-Zn-Mg-Cu alloy 163 炉衬：furnace lining164 规则推理：rule-based reasoning165 在线控制：on-line control166 共晶碳化物：eutectic carbide167 振动频率：vibrational frequency168 TA15钛合金：TA15 titanium alloy169 Cr12MoV钢：Cr12MoV steel170 变形镁合金：wrought magnesium alloy 171 功率超声：power ultrasound172 TiAl基合金：TiAl-based alloy173 Box-Behnken设计：Box-behnken design 174 专业课：specialized course175 金相组织：metallurgical structure176 模具寿命：die life177 研究应用：research and application 178 Al-Mg合金：Al-Mg alloy179 成本优化：cost optimization180 变形激活能：deformation activation energy181 干燥工艺：drying technology182 合金铸铁：alloy cast iron183 模具材料：die material184 铸态组织：as-cast microstructure185 电磁制动：electromagnetic brake186 球铁：ductile iron187 侧架：side frame188 气缸体：cylinder block189 洛伦兹力：Lorentz Force190 微观组织演变：microstructure evolution 191 显微组织：microscopic structure192 共晶组织：Eutectic structure193 冶金质量：metallurgical quality194 热震稳定性：thermal shock resistance 195 强迫对流：forced convection196 切削加工：cutting process197 过共晶Al-Si合金：Hypereutectic Al-Si Alloy198 定量金相：quantitative metallography 199 磁感应强度：Magnetic Flux Density 200 半固态浆料：Semi-solid Slurry201 电磁泵：electromagnetic pump202 超声衰减：Ultrasonic attenuation203 加热时间：heating time204 半连续铸造：Semi-continuous Casting 205 液压站：Hydraulic station206 三元硼化物：ternary boride207 内应力：inner stress208 热裂纹：hot crack209 黄麻纤维：jute fiber210 泡沫陶瓷：foam ceramics211 砂型铸造：Sand casting212 油润滑：oil lubrication213 预热温度：preheating temperature 214 维氏硬度：Vickers Hardness215 高温合金：high-temperature alloy216 拉速：casting speed217 铝熔体：aluminum melt218 异型坯：beam blank219 高钒高速钢：high vanadium high speed steel220 静液挤压：hydrostatic extrusion221 等轴晶：equiaxed grain222 摩擦角：friction angle223 初生相：Primary Phase224 转向节：steering knuckle225 快速成型技术：rapid prototyping technology226 冷坩埚：Cold Crucible227 A357合金：A357 Alloy228 焊接结构：welding structure229 耦合场：coupled field230 AZ80镁合金：AZ80 magnesium alloy 231 止推轴承：thrust bearing232 铝镁合金：Al-Mg alloy233 真空熔炼：vacuum melting234 铝锂合金：aluminum-lithium alloy235 充型过程：filling process236 AZ61镁合金：AZ61 magnesium alloy 237 声流：Acoustic streaming238 金属凝固：metal solidification239 高速钢轧辊：high speed steel roll240 石墨形态：graphite morphology241 磁粉检测：Magnetic particle testing 242 颗粒级配：particle size distribution243 型砂：molding sand244 收缩率：shrinkage rate245 Mg-Li合金：Mg-Li alloy246 自动生产线：automatic production line 247 高频磁场：High Frequency Magnetic Field248 组织与性能：microstructure and property249 连续定向凝固：continuous unidirectional solidification250 充型：mold filling251 失效机制：failure mechanism252 梯度分布：gradient distribution253 制动鼓：Brake drum254 摄动分析：perturbation analysis255 铸造企业：foundry enterprise256 超声波振动：Ultrasonic vibration257 测量系统分析：measurement system analysis258 固溶处理：solution heat treatment259 冷却速度：cooling velocity260 固液混合铸造：solid-liquid mixed casting 261 温度场分布：temperature distribution 262 部分重熔：Partial Remelting263 工艺措施：technological measures264 变形量：deformation amount265 模糊优化设计：Fuzzy optimal design 266 零缺陷：zero defect267 重力分离：gravitational separation268 晶粒：crystal grain269 离心力场：centrifugal force field270 凝固行为：Solidification Behavior271 铝铜合金：Al-Cu alloy272 组织和性能：microstructure and property 273 复合板：composite plate274 Al-Fe合金：Al-Fe alloy275 马氏体不锈钢：martensite stainless steel 276 冷却装置：cooling device277 铝合金车轮：aluminum alloy wheel 278 热应力分析：thermal stress analysis 279 Al含量：Al content280 挤压比：extrusion ratio281 相似准则：similarity criterion282 热疲劳裂纹：thermal fatigue crack283 原子团簇：atomic cluster284 湿型砂：green sand285 AZ91D合金：AZ91D alloy286 6061铝合金：6061 aluminum alloy287 锻造工艺：forging technology288 铸铁件：Iron casting289 表面复合材料：Surface composites 290 盲孔法：blind-hole method291 加热功率：heating power292 铸造合金：Cast Alloy293 低铬白口铸铁：Low chromium white cast iron294 初生硅：primary silicon 295 热节：Hot Spot296 锡青铜：tin bronze297 ZL101合金：ZL101 alloy298 真空感应熔炼：vacuum induction melting299 薄带连铸：strip casting300 真空压铸：vacuum die casting301 缩孔：shrinkage hole302 等温处理：Isothermal Treatment303 平均晶粒尺寸：average grain size304 抽芯：core pulling305 离心浇铸：Centrifugal casting306 铸铁管：cast iron pipe307 感应线圈：induction coil308 冷却介质：Cooling medium309 气体压力：gas pressure310 船用柴油机：marine diesel311 高温强度：high-temperature strength 312 3Cr2W8V钢：3Cr2W8V steel313 缺陷预测：defect prediction314 工艺方案：process scheme315 温度均匀性：temperature uniformity 316 电磁离心铸造：electromagnetic centrifugal casting317 横向应力：transverse stress318 超声声速：ultrasonic velocity319 残留应力：residual stress320 固化工艺：curing process321 精铸：Investment Casting322 铝锭：aluminum ingot323 短路过渡：short circuit transfer324 反重力铸造：counter-gravity casting 325 感应电炉：induction furnace326 稀土Y：rare earth Y327 工艺因素：Technological factor328 双辊铸轧：twin roll casting329 凝固速率：solidification rate330 含氢量：Hydrogen Content331 钢锭：steel ingot332 浆料制备：slurry preparation333 η相：η phase334 衬板：lining board335 压铸件：die casting336 水口堵塞：nozzle clogging337 陶瓷型芯：ceramic core338 车间布局：workshop layout339 安全操作：safe operation340 铸造不锈钢：cast stainless steel341 压铸模具：die casting die342 热裂：Hot Crack343 失效形式：failure form344 成形机理：forming mechanism345 AlSi7Mg合金：AlSi7Mg Alloy346 铸件缺陷：casting defect347 银合金：silver alloys348 反应层：reaction layer349 镍基高温合金：Ni base superalloy350 薄带：thin strip351 覆膜砂：coated sand352 CAE技术：CAE Technique353 性能预测：property prediction354 液态金属：liquid metals355 熔模精密铸造：investment casting356 空气压力：air pressure357 ZA合金：ZA alloy358 凝固传热：Solidification and heat transfer 359 侧向分型：Side Parting360 高温塑性：Hot Ductility361 黑斑：black spot362 点火温度：ignition temperature363 旋压机：spinning machine364 Al-Ti-B中间合金：Al-Ti-B master alloy 365 减排：discharge reduction366 射线检测：radiographic inspection367 耐热：heat resistant368 2024铝合金：2024 aluminum alloy369 技术现状：technology status370 复合变质：complex modification371 蠕墨铸铁：vermicular iron372 机械搅拌：mechanical agitation373 保温炉：holding furnace374 成形技术：forming technology375 碳化硅颗粒：SiC particle376 可锻铸铁：malleable iron377 模型控制：model control378 改性水玻璃：modified sodium silicate 379 熔炼工艺：melting process380 焊补：repair welding 381 异常组织：abnormal structure382 组织细化：structure refinement383 防止措施：preventing measures384 铸渗：Casting infiltration385 BT20钛合金：BT20 titanium alloy386 直流电场：direct current field387 铸造应力：casting stress388 初晶Si：primary Si389 夹紧装置：clamping device390 均衡凝固：Proportional solidification 391 熔模精铸：investment casting392 空心叶片：hollow blade393 ZL201合金：ZL201 alloy394 温轧：warm rolling395 不均匀变形：inhomogeneous deformation396 呋喃树脂砂：furan resin sand397 纸浆：paper pulp398 半连铸：semi-continuous casting399 锻锤：forging hammer400 延伸率：elongation rate401 焊接修复：welding repair402 冶金结合：metallurgical bond403 技术对策：technical measures404 结晶器振动：Mold Oscillation405 厚壁：thick wall406 WC颗粒：WC particles407 预处理技术：pretreatment technology 408 金属零件：metal part409 特种铸造：special casting410 低熔点合金：low melting point alloy 411 水模实验：water model experiment 412 复合管：clad pipe413 插装阀：plug-in valve414 金相试样：Metallographic specimen 415 抗吸湿性：humidity resistance416 近液相线铸造：near-liquidus casting 417 新设计：new design418 电机转子：motor rotor419 CAE：computer aided engineering420 交流变频：AC variable frequency421 下横梁：lower beam422 ZL102合金：ZL102 alloy423 模型参考控制：model reference control424 虚拟对象：virtual object425 加工图：processing maps426 立式离心铸造：vertical centrifugal casting427 抽芯机构：core pulling mechanism428 连铸连轧：casting and rolling429 残留强度：residual strength430 复合铸造：composite casting431 树脂砂：resin bonded sand432 AM60B镁合金：AM60B magnesium alloy 433 铸造CAE：casting CAE434 砂型：sand mould435 熔化：melting process436 高硼铸钢：high boron cast steel437 稳恒磁场：stable magnetic field438 Al-Ti-C晶粒细化剂：Al-Ti-C grain refiner 439 再生技术：regeneration technology 440 压铸工艺：die casting process441 管坯：tube billet442 厚大断面：Heavy section443 保护气体：protective gas444 性能特征：performance characteristics 445 Al-5%Fe合金：Al-5%Fe alloy446 半固态挤压：Semi-solid extrusion447 金属型铸造：Permanent mold casting 448 晶粒组织：grain structure449 综合经济效益：Comprehensive economic benefit450 半固态压铸：semi-solid die casting451 气膜：gas film452 硅酸乙酯：Ethyl Silicate453 自动化生产线：automatic production line454 Mg-Gd-Y-Zr合金：Mg-Gd-Y-Zr alloy455 渗透检测：Penetrant testing456 W-Cu复合材料：W-Cu composites457 存放时间：storage time458 ProCAST软件：ProCAST software459 滑板：sliding plate460 铸造铝合金：casting aluminum alloy 461 水玻璃砂：Water-glass Sand462 电脉冲：Electrical pulse463 蜡模：Wax Pattern464 悬浮铸造：suspension casting 465 D型石墨：D-type graphite466 工艺性能：technological performance 467 Al-1%Si合金：Al-1%Si alloy468 悬浮性：suspension property469 差压铸造：counter-pressure casting 470 工艺原理：process principle471 铸轧：continuous roll casting472 行波磁场：traveling magnetic field473 型壳：Shell Mold474 金属型：permanent mould475 脱模机构：demolding mechanism476 调压铸造：adjusted pressure casting 477 喷砂：sand blasting478 界面换热系数：interfacial heat transfer coefficient479 Al-Mg-Si-Cu合金：Al-Mg-Si-Cu alloy 480 电熔镁砂：fused magnesia481 充型速度：Filling Velocity482 泵体：pump body483 钢锭模：ingot mould484 Cu-Fe合金：Cu-Fe alloy485 辐射力：radiation force486 空化泡：Cavitation bubble487 渣池：slag pool488 原位生成：In-situ Synthesis489 热型连铸：heated-mold continuous casting490 缩松：dispersed shrinkage491 CO2气体保护焊：CO_2 arc welding 492 伺服控制系统：servo system493 端盖：End cover494 铸造技术：casting technology495 水力学模拟：Hydraulics simulation496 再生铝：secondary aluminum497 轴套：axle sleeve498 成形模具：forming die499 抗磨性能：Wear Resistance500 水模拟：water model501 快速铸造：rapid casting502 电磁软接触：electromagnetic soft-contact503 石膏型：plaster mold504 大型铸钢件：heavy steel casting505 移动磁场：traveling magnetic field506 轴承座：bearing seat507 混合稀土：rare earth508 铸态球铁：as-cast nodular iron509 砂芯：sand core510 铸造性能：casting properties511 真空差压铸造：vacuum counter-pressure casting512 玻璃模具：glass mold513 双联熔炼：duplex melting514 设备改进：improvement of equipment 515 铸坯质量：billet quality516 局部加压：Local Pressurization517 旧砂再生：used sand reclamation518 结晶速度：Crystallization rate519 壳体：shell body520 干强度：dry strength521 浇注系统设计：gating system design 522 慢压射：slow shot523 图像分析仪：image analysis system 524 温度曲线：Temperature profile525 水力效率：hydraulic efficiency526 单晶铜：single-crystal copper527 电渣重熔：electroslag refining528 铸造起重机：casting crane529 Cu-Cr合金：Cu-Cr alloys530 堆垛机：stacking machine531 巴氏合金：Babbitt alloy532 自抗扰控制器：auto-disturbance rejection controller(ADRC)533 陶瓷型：ceramic mold534 直流磁场：direct current magnetic field 535 漏气：air leakage536 泡沫陶瓷过滤器：foam ceramic filter 537 过共晶高铬铸铁：Hypereutectic High Cr Cast Iron538 壁厚差：wall thickness difference539 HPb59-1黄铜：HPb59-1 Brass540 旋转喷吹：Spinning Rotor541 水玻璃旧砂：used sodium silicate sand 542 冷却强度：cooling strength543 耐磨铸铁：wear resistant cast iron544 ZA35合金：ZA35 alloy545 钠基膨润土：sodium bentonite546 熔体净化：melt purification 547 油雾润滑：oil-mist lubrication548 初生α相：primary α phase549 铸造生产：foundry production550 高电位：High Potential551 钴基高温合金：cobalt base superalloy 552 Al-Zn-Mg-Cu-Zr合金：Al-Zn-Mg-Cu-Zr alloy553 水平连续铸造：Horizontal continuous casting554 自硬砂：no-bake sand555 微区分析：micro-area analysis556 顺序凝固：sequential solidification557 非枝晶组织：Non-dendritic microstructure558 反变形：reverse deformation559 铬青铜：Chromium bronze560 湿型铸造：green sand casting561 配料计算：burden calculation562 热-力耦合：Thermo-mechanical Coupling 563 浇注时间：Pouring time564 铸造速度：Casting velocity565 亚共晶铝硅合金：Hypoeutectic Al-Si Alloy566 搅拌功率：power consumption567 热电场：thermoelectricity field568 铸铝合金：cast aluminum alloy569 陶瓷型铸造：Ceramic mold casting570 热凝固：Thermal coagulation571 界面压力：interface pressure572 多尺度模拟：multiscale simulation573 输送链：Conveyor Chain574 关键措施：key measures575 冒口系统：Riser system576 开炉：blowing in577 铜锡合金：Cu-Sn alloy578 无铅黄铜：unleaded brass579 球墨铸铁管：ductile cast iron pipe580 二次枝晶间距：secondary dendrite arm spacing581 GA-BP网络：GA-BP network582 铝合金熔体：aluminum alloy melt583 生产条件：production conditions584 铬铁矿砂：chromite sand585 再生效果：regeneration effect586 导向叶片：Guide Vane587 金属管：Metal tube588 空心管坯：hollow billet589 超高强铝合金：ultra-high strength aluminum alloy590 流变曲线：flow curve591 蠕化剂：vermicularizing alloy592 波浪型倾斜板：wavelike sloping plate 593 凝固特性：solidification characteristics 594 磨头：grinding head595 反白口：reverse chill596 黑线：black line597 净化技术：purifying technology598 中间合金：master alloys599 捏合块：Kneading Block600 硅相：silicon phase601 低过热度浇注：low superheat pouring 602 3004铝合金：3004 aluminum alloy603 液态压铸：liquid die casting604 中频感应电炉：intermediate frequency induction electric furnace605 球墨铸铁件：Ductile iron casting606 凝固路径：solidification path607 喷枪：spraying gun608 ZL201铝合金：ZL201 aluminum alloy 609 质量改善：quality improvement610 气路：gas circuit611 补缩设计：Feeding design612 油底壳：Oil sump613 汽缸体：cylinder block614 CREM法：CREM process615 铸造机：Casting machine616 提高措施：improving measure617 SIMA法：SIMA method618 铬系白口铸铁：Chromium white cast iron 619 高合金钢：High alloy steels620 增压系统：pressurization system621 收缩缺陷：shrinkage defect622 卧式离心铸造：Horizontal Centrifugal Casting623 测控仪：measuring and controlling instrument624 精铸件：Investment Castings625 制动阀：Brake valve 626 金属成型：metal forming627 有机纤维：organic fiber628 大气采样器：air sampler629 钢支座：steel bearing630 低频磁场：low frequency magnetic field 631 破坏面：failure surface632 偏轨箱形梁：bias-rail box girder633 数值处理：data processing634 双辊薄带：twin-roll thin strip635 合成铸铁：Synthetic cast iron636 堆冷：stack cooling637 行星轧制：planetary rolling638 铸造缺陷：foundry defect639 二次冷却：second cooling640 炉衬材料：lining material641 弥散强化：dispersion hardening642 2D70铝合金：2D70 aluminum alloy 643 A356铝合金：A356 Al alloy644 元胞自动机方法：Cellular Automaton method645 铸造温度：casting temperature646 铸造涂料：Foundry coating647 耦合模拟：coupled simulation648 充型能力：Filling ability649 复合尼龙粉：nylon composite powder 650 改性纳米SiC粉体：modified SiC nano-powders651 炉外脱硫：external desulfurization652 绿色铸造：green casting653 净化方法：purification method654 制芯：Core making655 铸态球墨铸铁：as-cast ductile iron656 复合轧辊：compound roller657 冷隔：cold shut658 薄壁件：thin-wall part659 铸钢车轮：cast steel wheel660 铁水质量：quality of molten iron661 热物理性能：Thermo-physical properties 662 7050铝合金：7050 Al alloy663 半固态金属加工：semi-solid metal forming664 半固态铸造：semisolid casting665 表面反应：Surface reactions666 KBE：knowledge-based engineering(KBE)667 倾斜板：inclined plate668 弯销：dog-leg cam669 多边形效应：polygonal effect670 脱模剂：releasing agent671 铜包铝线：copper clad aluminum wire 672 球化衰退：nodularization degeneration 673 低过热度：low superheat674 升降机构：lifting mechanism675 SLS：selective laser sintering(SLS)676 溢流槽：spillway trough677 制浆技术：pulping technology678 浇注工艺：casting process679 变形行为：deformation behaviors680 转移涂料：transfer coating681 牵引速度：haulage speed682 WC/钢复合材料：WC/steel composites 683 泡沫模样：foam pattern684 皮下气孔：surface blowhole685 超高强度铝合金：ultrahigh strength aluminum alloy686 薄带铸轧：strip casting687 造型线：moulding line688 工具杆：tool rod689 铸锭组织：ingot microstructure690 复合变质剂：composite modifier691 发热剂：Heating Agent692 液相线半连续铸造：liquidus semi continuous casting693 Mg-Al-Zn合金：Mg-Al-Zn alloy694 洛仑兹力：Lorenz force695 散射比：scattering ratio696 翻转机构：turnover mechanism697 超声铸造：Ultrasonic Casting698 A356：A356 alloy699 Mg-Li-Al合金：Mg-Li-Al alloy700 复合磁场：electromagnetic field701 单缸机：single cylinder engine702 快速产品设计：Rapid Product Design 703 真空阀：Vacuum valve704 界面传热系数：Interfacial heat transfer coefficient705 液态金属冷却：liquid metal cooling 706 散射衰减：scattering attenuation707 电磁场频率：Electromagnetic Frequency 708 半连续铸锭：semicontinuous casting ingot709 凝固补缩：Solidification Feeding710 Mg-Zn合金：Mg-Zn alloy711 连铸-热轧区段：CC-HR region712 TC11钛合金：titanium alloy713 损坏机理：failure mechanism714 元素分布：Distribution of element715 原位TiC颗粒：in-situ TiC particles716 均匀化处理：uniform heat treatment 717 使用要求：application requirement718 初生相形貌：morphology of primary phase719 枝晶形貌：dendritic morphology720 铸造废弃物：foundry waste721 AZ91D：AZ91D Magnesium Alloy722 高压铸造：high pressure die casting 723 细化变质：Refinement and Modification 724 结疤：scale formation725 连续铸轧：continuous casting726 热变形行为：Thermal Deformation Behavior727 壳型铸造：shell mould casting728 消失模：evaporative pattern729 手机外壳：mobile phone shell730 热管技术：heat pipe731 水韧处理：water toughening process 732 阻燃镁合金：Ignition proof magnesium alloys733 除尘装置：dust collector734 悬浮率：suspending rate735 非线性估算法：nonlinear estimation method736 电解铝液：electrolytic aluminum melt 737 双金属复合：bimetal compound738 离心浇注：centrifugal pouring739 抗磨损：abrasion resistance740 薄壁铸件：thin-walled casting741 盖包法球化处理：tundish-cover nodulizing process742 无定形二氧化硅：amorphous silicon dioxide743 排气槽：air vent744 高铬白口铸铁：high chromium cast iron745 熔炼炉：smelting furnace746 过滤机理：Filtration mechanism747 汽车覆盖件模具：auto panel die748 低合金高强度钢：Low-alloy high-strength steel749 精铸模具：investment casting mould 750 铝板带：aluminum plate751 球状石墨：nodular graphite752 铸轧区：casting-rolling zone753 接线盒：junction box754 铁水净化剂：purifying agent for molten iron755 石墨块：graphite block756 优质铸件：high quality casting757 处理温度：treatment temperature758 高尔夫球头：golf head759 固相体积分数：solid volume fraction 760 纳米SiC颗粒：SiC nanoparticle761 检测仪器：testing instrument762 Mg17Al12相：Mg_(17)Al_(12) phase 763 攻关：tackling key problems764 硬化机理：Hardening mechanism765 真空吸铸：vacuum suction766 热分析技术：thermal analysis technology 767 高频调幅磁场：High Frequency Amplitude-modulated Magnetic Field768 坯料制备：blank production769 补缩通道：feeding channel770 水基涂料：water-based coating771 球铁件：Ductile Iron Castings772 稀土Er：rare earth Er773 陶瓷型壳：Ceramic shell774 精密电铸：precision electroforming 775 发气性：Gas evolution776 充型凝固：Mold Filling and solidification 777 铝带：aluminum strip778 新SIMA法：new SIMA method779 AZ91HP镁合金：AZ91HP magnesium alloy780 电子束冷床熔炼：electron beam cold hearth melting781 粘砂：metal penetration782 物理冶金学：physical metallurgy783 砂处理：Sand preparation 784 铸造裂纹：casting crack785 气冲造型：air impact molding786 金属模：metal mould787 磷共晶：phosphor eutectic788 近液相线半连续铸造：nearby liquidus semi-continuous casting789 液固反应：liquid-solid reaction790 呋喃树脂：furane resin791 汽缸盖：Cylinder Cap792 充型模拟：Simulation of mold filling 793 铸造工艺CAD：casting technology CAD 794 粘土砂：Clay sand795 冲天炉熔炼：cupola smelting796 射料充填过程：filling process797 半固态金属：semisolid metals798 大型铸件：heavy casting799 电机端盖：motor cover800 熔铸工艺：casting process801 加入方法：Joined technique802 区域熔化：zone melting803 真空除气：Vacuum Degassing804 相平衡热力学：phase equilibrium thermodynamics805 溢流系统：overflow system806 Al-Ti-C中间合金：Al-Ti-C master alloys 807 晶界碳化物：grain boundary carbide 808 净化装置：purification equipment809 液穴形状：sump shape810 铝合金铸造：Aluminum Alloy Casting 811 修模：Tool modification812 SKD61钢：SKD61 steel813 软化退火：Softening Annealing814 大齿轮：Large Gear815 合金渗碳体：Alloy cementite816 工艺性能试验：technological property tests817 硅碳比：Si/C ratio818 冷却曲线：Cooling Curves819 壁厚不均：non-uniform wall thickness 820 V法铸造：V process821 铸造系统：casting system822 电渣加热：electroslag heating823 残余内应力：residual stress824 表面清理：surface cleaning825 黄斑：macular region826 电磁振荡：Electromagnetic Oscillation 827 初始组织：initial structure828 气密性能：air permeability performance 829 电极调节：electrode adjustment830 气体速度：gas velocity831 抑制方法：suppressing method832 孔洞率：void ratio833 废品率：reject rate834 气动装置：pneumatic actuator835 应急发电机：emergency generator836 缺陷修复：Error repair837 有机高聚物：organic polymer838 理论成果：theoretical achievements 839 凝固曲线：Solidification curve840 元胞自动机法：cellular automaton841 ZL101铝合金：ZL101 Al alloy842 高韧性球墨铸铁：High toughness ductile iron843 搅拌方式：stirring method844 沉积坯尺寸：deposit dimension845 高锌镁合金：high zinc magnesium alloy 846 雕铣机：carves-milling machine847 铸造模拟：Casting simulation848 精益设计：lean design849 无余量精密铸造：Investment Casting 850 热顶铸造：hot-top casting851 羊油：mutton tallow852 压射速度：injection speed853 DOE试验：DOE experiment854 超声波振荡：ultrasonic oscillation855 酯固化：ester cured856 缸盖罩：cylinder head cover857 尺寸变化率：dimension variance rate 858 大型铸铁件：heavy iron castings859 单晶铜线材：copper single crystal wire 860 厚大断面球墨铸铁：heavy section ductile iron861 钛镍合金：Ti-Ni alloy862 实型铸造：Full Mold863 6082合金：6082 Alloy864 奥贝球铁：austenite-bainite nodular-iron 865 白口组织：white microstructure866 铸轧工艺参数：casting process parameters867 铸铁轧辊：cast iron milling roll868 强化处理：strengthen treatment869 半固态成型：semi-solid processing870 深腔：deep cavity871 耐热镁合金：Heat resistant magnesium alloys872 斜滑块：inclined sliding block873 回炉料：recycled scrap874 半固态坯：semi-solid billet875 感应熔炼：inductive melting876 链板：chain board877 含泥量：sediment percentage878 模料：mould material879 复合界面：compounded interface880 铸造方法：casting methods881 模温：mold temperature882 轻合金：light alloys883 增碳工艺：recarburation process884 定位装置：location equipment885 加压速率：pressurization rate886 半固态流变成形：Semi-solid Rheoforming887 复杂铸件：Complicated casting888 高强度灰铸铁：High strength grey cast iron889 针孔度：pinhole degree890 中频感应加热：intermediate frequency induction heating891 石墨转子：graphite rotor892 修磨机：Grinding machine893 动态顺序凝固：dynamic directional solidification894 针状组织：acicular structure895 粒度配比：particle size distribution896 铝合金壳体：aluminum alloy shell897 内冷铁：Internal chill898 铸件质量：quality of casting899 精炼效果：refining effect900 发动机缸体：cylinder body901 增碳剂：carburizing agent902 7005铝合金：7005Al alloys903 复合孕育：Multiple inoculations904 复合孕育剂：compound inoculation905 气孔缺陷：blowhole defect906 铁液质量：quality of molten iron907 钛铝合金：TiAl alloys908 7A09铝合金：7A09 aluminium alloy 909 SiC颗粒增强：SiC particle reinforcement 910 沉淀相：precipitated phases911 铝母线：aluminum bus912 凝固分数：solid fraction913 球化组织：spheroidized microstructure 914 蠕铁：vermicular iron915 组织均匀性：microstructure uniformity 916 压铸型：die-casting die917 镁合金压铸机：magnesium alloy die casting machine918 凝固微观组织：solidification microstructure919 灰铸铁件：Gray iron casting920 最大剪应力：ultimate shear stress921 热挤压成形：hot extrusion922 铝合金铸件：aluminium alloy cast923 抗湿性：humidity resistance924 耳子：rolling edge925 结合面：joint face926 推管：ejector sleeve927 黑点：black spot928 铝铸件：aluminum casting929 固相分数：Solid fraction930 快干硅溶胶：Quick-dry silica sol931 激冷铸铁：Chilled iron932 负压消失模铸造：Negative pressure EPC 933 LC9铝合金：LC9 aluminium alloy934 接触层：Contact layer935 工频炉：main frequency furnace936 消失模涂料：lost foam casting coating 937 高温均匀化：high temperature homogenization938 均热炉：pit furnace939 镁合金轮毂：magnesium wheel940 平砧：flat anvil941 铝合金扁锭：aluminum alloy slab942 凝固界面：solidifying interface943 低温冲击功：Low Temperature Impact Energy944 复合发泡剂：Composite Foaming Agent 945 交叉型芯：Crossed Core946 SCR连铸连轧：SCR continuous casting-rolling947 FS粉：FS powder948 AZ81镁合金：AZ81 alloy949 ZL109活塞：ZL109 piston950 掉砂：dropping sand951 型腔壁厚：cavity wall thickness952 铝件：aluminum part953 导向装置：guide mechanism954 彩色云图：color contour image955 柴油机缸体：Diesel engine cylinder block 956 圆盘铸锭机：casting wheel957 热风冲天炉：Hot-blast cupola958 充氧压铸：pore-free die casting959 铝钛硼细化剂：Al-Ti-B refiner960 保温冒口：Insulating riser961 共晶相：Eutectic phase962 夹砂：sand inclusion963 无冒口铸造：Riserless casting964 充芯连铸：continuous core-filling casting 965 熔体混合：melt mixing966 保护渣道：mold flux channel967 碱性酚醛树脂：alkaline phenolic resins 968 细深孔：Long-deep hole969 行星减速机：planetary reducer970 直接铸型制造：direct casting mold manufacturing971 引锭头：dummy bar head972 静置炉：holding furnace973 工艺出品率：process yield974 真空法：vacuum process975 石灰石砂：limestone sand976 整体浇注：monolithic casting977 混料工艺：mixing procedure978 螺旋套：screwy sheath979 胶凝机理：gelling mechanism980 覆砂铁型：permanent mould with sand facing981 球铁铸件：ductile iron casting982 成型率：molding rate983 球状组织：spherical structure984 电弧冷焊：arc cold welding985 钢液流场：flow field of molten steel。

7050铝合金铸锭均匀化热处理滕广标【摘要】试验研究了7050铝合金铸态及不同温度-时间均匀化处理后的组织演变.研究结果表明:铸态组织中存在严重枝晶偏析,400℃均匀化处理过程中,非平衡凝固共晶相向合金基体持续溶解,在465℃均匀化时,平衡η(MgZn2)相、T (AlZnMgCu)相等大部分共晶相回溶到基体中,晶界明显细化,均匀化效果显著.确定了7050铝合金铸锭最佳均匀化工艺制度为465℃保温24h.【期刊名称】《轻合金加工技术》【年(卷),期】2019(047)008【总页数】5页(P21-25)【关键词】7050铝合金;铸锭;均匀化处理;显微组织【作者】滕广标【作者单位】广东坚美铝型材厂(集团)有限公司,广东佛山528222【正文语种】中文【中图分类】TG146.21;TG166.37050铝合金具有低密度、高强度、良好的韧性及抗应力腐蚀性能好等优点,是航空航天领域重要的结构材料之一[1-2]。

η(MgZn2)相作为其中的主要强化相，在时效过程中析出[3]。

然而，由于该合金的合金化程度高，合金元素往往处于过饱和状态，在铸造过程中，由于冷却速度较快，合金元素快速在晶界处富集，产生枝晶偏析及区域偏析，并且在铸锭内部形成很强的内应力，降低合金塑性，恶化材料的使用性能[4-7]。

因此铸锭加工之前必须经过均匀化热处理，消除不均匀铸锭组织成分及非平衡凝固状态下的残余应力，减少铸锭开裂现象，改善材料的性能。

对于不同牌号的铝合金,前人做了相关的研究工作[8-17]。

由于合金元素含量比例不同，均匀化过程中的显微组织演变也有所不同。

本课题研究了均匀化温度、时间等因素对7050铝合金铸锭均匀化效果的影响，分析了均匀化前后相的转变过程，以便于精确制定7050铝合金铸锭的均匀化热处理工艺制度。

1 试验材料及方法试验用7050铝合金铸锭成分见表1。

使用线切割方法切割成尺寸为20 mm×10 mm×5 mm试样。

锆-4合金高温高压水蒸气氧化行为王志武;宋涛【摘要】Oxidation behavior of the zircaloy-4 at 330, 400 and 500 ℃ and high-pressure water vapor was investigated by discontinuous weighing method. The composition and phases of oxidation films were also analyzed. The results indicate that oxidation dynamic curve obeys parabolic law at all testing temperatures. The oxidation film is nucleated at ravine position, and then covers the surface gradually; the second oxidation film appears as slice-sharp. t-ZrO2 and m-ZrO2 are two main phases in the oxidation films. Oxides rich in metal atoms and lack of oxygen atoms appear at lower temperature due to short supply of oxygen.%利用不连续称重法测定锆-4合金于330,400和500℃高温高压水蒸气氧化动力学曲线,并对氧化膜成分和物相进行分析.研究结果表明:实验温度下氧化动力学曲线符合yn=Kt+C的抛物线规律;氧化膜在沟壑位置形核逐渐覆盖表面,第2层氧化膜以片状开始出现;氧化物主要是t-ZrO2和m-ZrO2 2种,低温时由于氧原子供应不足导致金属原子富余氧原子缺失的氧化物出现.【期刊名称】《中南大学学报（自然科学版）》【年(卷),期】2013(044)002【总页数】5页(P515-519)【关键词】锆-4合金;高温高压水蒸气氧化;动力学曲线;t-ZrO2;m-ZrO2【作者】王志武;宋涛【作者单位】武汉大学动力与机械学院,湖北武汉,730072;武汉大学动力与机械学院,湖北武汉,730072【正文语种】中文【中图分类】TG174.1锆-4合金是核电站中压水堆燃料棒包壳管重要用材，最高使用温度340 ℃，工作压力不超过16 MPa，起着防止放射性物质泄漏到一回路水中的作用[1-3]。

润滑方式对7075铝合金车削表面耐腐蚀的影响裴宏杰;付坤鹏;邹晔;刘成石;王贵成【摘要】7075高强度铝合金经常在大气环境中应用,极易发生腐蚀.7075铝合金零件配合表面需要经过切削加工而成形,其机械加工一般都是在干切削和微量润滑(Minimum Quantity Lubricant,MQL)条件下进行的.为了揭示出润滑条件对7075铝合金加工面耐腐蚀性的影响规律,进行了切削和腐蚀单因素实验.在不同的切削速度和进给量下,对直径30衄的7075铝合金棒料,进行干车削和MQL车削加工,然后对样件进行3个周期的盐雾试验,每个周期72 h.观察其腐蚀表面形貌,统计出腐蚀面积、蚀坑数量、腐蚀损伤平均深度和腐蚀损伤度DOP等特征参数,揭示出7075铝合金在不同润滑条件下车削表面的腐蚀规律.%7075 high strength aluminum alloy is easily to be corroded in atmosphere.In the mechanical manufacturing process,7075 aluminum alloy component surface was got through the dry and MQL machining method.In this paper,with differental cutting speed,and feed rate,both dry and MQL turning experiments were carried out with Φ30 mm 7075 aluminum alloy rod.After that,salt spray test was done for three cycles,each period was 72 hours.Then,the corrosion state of surface morphology was observed,and meanwhile the corrosion area,corrosion pits quantities,the degree of corrosion damage,the average depth and average diameter of corrosion pits were counted.Finally,the corrosion law of machined surface was revealed in the aluminum alloy turning application.【期刊名称】《制造技术与机床》【年(卷),期】2018(000)003【总页数】5页(P122-126)【关键词】7075铝合金;MQL;干切削;腐蚀【作者】裴宏杰;付坤鹏;邹晔;刘成石;王贵成【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;无锡职业技术学院机械技术学院,江苏无锡214121;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】TG7077075高强度铝合金作为航空航天以及船舶制造领域中的主要材料，经常要暴露在大气环境下，在这种特殊环境中金属材料容易发生腐蚀[1]。

Trans.Nonferrous Met.Soc.China29(2019)1816−1823Microstructure evolution of Al−Cu−Mg alloy duringrapid cold punching and recrystallization annealingZe-yi HU1,Cai-he FAN1,Dong-sheng ZHENG1,Wen-liang LIU1,Xi-hong CHEN21.College of Metallurgy and Material Engineering,Hunan University of Technology,Zhuzhou412007,China;2.CRRC Zhuzhou Electric Locomotive Co.,Ltd.,Zhuzhou412007,ChinaReceived31October2018;accepted30June2019Abstract:The microstructure evolution of spray formed and rapidly solidified Al−Cu−Mg alloy with fine grains during rapid cold punching and recrystallization annealing was investigated by transmission electron microscopy(TEM).The results show that the precipitates of fine-grained Al−Cu−Mg alloy during rapid cold punching and recrystallization annealing mainly consist of S phase and a small amount of coarse Al6Mn phase.With the increase of deformation passes,the density of precipitates increases,the size of precipitates decreases significantly,and the deformation and transition bands disappear gradually.In addition,the grains are refined and tend to be uniform.Defects introduced by rapid cold punching contribute to the precipitation and recrystallization,and promote nucleation and growth of S phase and recrystallization.Deformation and transition bands in the coarse grains transform into deformation-induced grain boundary during the deformation and recrystallization,which refine grains,obtain uniform nanocrystalline structure and promote homogeneous distribution of S phase.Key words:Al−Cu−Mg alloy;microstructure evolution;precipitate;recrystallization;deformation band;rapid cold punching1IntroductionAl−Cu−Mg alloy has been widely applied to aerospace and military industries because of its advantage such as high strength,good formability and heat resistance[1,2].Precipitation strengthening and grain refinement are the main strengthening and toughening methods.Under the conventional T6-like heat treatment,the main strengthening phase in Al−Cu−Mg alloy is S phase with low Cu/Mg ratios (Al2CuMg)andθ'phase(Al2Cu)with high Cu/Mg ratios[3−6].Deformation not only makes the material obtain good work hardening properties,but also introduces a large number of dislocations during the deformation.Aging process after deformation can release deformation stress,promote the dispersive nucleation and growth of precipitates,and even change the characteristics and precipitation sequence of precipitates[7−10].STYLES et al[11]investigated the relationship between the decomposition sequence of supersaturated solid solution and the phase in Al−Cu−Mg alloy,and pointed out that the formation time of S phase at higher temperature is much shorter than that at lower temperature.LI et al[12]investigated the effect of pre-deformation on the microstructure of high-purity Al−Cu−Mg alloy and found that the density of S'phase(Al2CuMg)increases while its size decreases with the increase of pre-deformation degree.YIN et al[13]studied the growth behavior of S precipitation phase particles within the grains of high strength Al−Cu−Mg alloy.It was reported that the structural units in the GPB region formed around S phase at higher aging temperature(above180°C)and hindered the growth of S phase along the width direction,resulting in the growth of S phase into columnar crystal.YANG et al[14] studied the effect of applied stress on the precipitation process ofθ'and S phases in Al−Cu−Mg alloy and found that applied stress prevented the precipitation ofθ'phase by changing theθ'/S ratio during the competitive precipitation,thereby the precipitation of S phase was promoted.It can be seen that the microstructure evolution and phase structure characteristics of Al−Cu−Mg alloy have been studied under the conventional deformation and aging temperatures.In the present work,the aluminum alloy cartridgeFoundation item:Project(2019JJ60050)supported by the Natural Science Foundation of Hunan Province,China Corresponding author:Cai-he FAN;Tel:+86-731-22183432;E-mail:369581813@DOI:10.1016/S1003-6326(19)65089-2Ze-yi HU,et al/Trans.Nonferrous Met.Soc.China29(2019)1816−18231817 was prepared by rapid cold punching andrecrystallization annealing,based on the spray formedAl−Cu−Mg alloy with fine grains.The precipitationphase,grain morphology and the evolution ofdeformation band during rapid cold punching andrecrystallization annealing were studied.The interactionbetween precipitation phase and recrystallization,theformation mechanism of deformation band and its effecton grain refinement were discussed.2ExperimentalFine-grained Al−Cu−Mg alloy cylindrical billet wasprepared by spray forming on a self-developed sprayforming device SD380.The chemical composition of thealloy is shown in Table1.The cylindrical billet wasextruded into round bar with a diameter of30mm by a1250T extruding machine at723K and the extrusionratio was15:1.Cylinder samples with20mm in lengthwere cut from the bar by wire-cutting machine and thenplaced in a self-designed stamping die.After four passesof rapid cold punching and recrystallization annealing,the aluminum alloy cartridges were prepared.Theschematic diagram of rapid cold punching is shownin Fig.1.The process of rapid cold punching andrecrystallization annealing is shown in Fig. 2.Inthe case of recrystallization annealing,the correspondingTable1Chemical composition of Al−Cu−Mg alloy(wt.%)Cu Mg Mn Si Fe Al4.51 1.460.56<0.05<0.05Bal.Fig.1Schematic diagrams of rapid cold punching:a—Sample; b—Drawing die;c—Punch Fig.2Process diagram of recrystallization annealing at763K for30min and rapid cold punching at298Kheating rate was623K/min.After holding for30min, the samples were cooled to room temperature and the next cold punching was performed.The process parameters of rapid cold punching are shown in Table2.The specimens were selected on the wall of cartridges for transmission electron microscopy(TEM) analysis.The preparation process of TEM specimens was as follows.The specimens were mechanically ground to 100μm before punching,and then finely ground to 70μm.After that,the specimens were twin jet electropolished in a mixed solution of25%nitric acid and75%methanol at253K using the voltage of20V. All specimens were washed in plasma cleaner Fishione, and their fine structures were observed in a Titan G2 60−300transmission electron microscope.The electron microscopic parameters observed by high-angle annular dark-field scanning transmission electron microscopy (HAADF-STEM)were as follows.Acceleration voltage was200keV,half-convergence angle of electron beam was10mrad,inner half-angle of high-angle annular probe was36mrad,and beam spot diameter was 0.20nm.Table2Process parameters of rapid cold punchingPunch Diameter/mm Velocity/(mm·s−1)One-pass1030Two-pass1425Three-pass2020Four-pass27153Results3.1Precipitation phase characteristicsThe TEM images of the precipitation phases after recrystallization annealing of Al−Cu−Mg alloy specimens with different passes are shown in Fig.3.EDS spectra of precipitation phase in Fig.3(a)are shown in Fig.4.The main precipitation phase of the alloy is S1818Ze-yi HU,et al/Trans.Nonferrous Met.Soc.China29(2019)1816−1823Fig.3TEM images of precipitation phases in Al−Cu−Mg alloy specimens under different conditions:(a)As-extruded;(b)One-pass;(c)Two-pass;(d)Three-pass;(e,f)Four-passFig.4EDS spectra of precipitation phases in Al−Cu−Mg alloy specimens:(a)Position1in Fig.3(a);(b)Position2in Fig.3(a)Ze-yi HU,et al/Trans.Nonferrous Met.Soc.China29(2019)1816−18231819phase(Al2CuMg).Due to the addition of Mn,a small amount of coarse Al6Mn phase can be observed.With the increase of deformation passes,the degree of deformation increases,the density of precipitation phase increases,the size decreases,and the precipitation phase tends to disperse.As a result of high deformation temperature of hot extruded alloy,the coarse Al6Mn phase and S phase can be observed in the alloy (Fig.3(a)).After rapid cold punching of one pass and recrystallization annealing,the size of precipitation phase significantly decreases,especially the Al6Mn phase is refined obviously and its shape is elongated(Fig.3(b)). After rapid cold punching of two or three passes and recrystallization annealing,the sizes of Al6Mn and S phases decrease further,the amount of both phases increases,and the distribution of them in matrix tends to be more uniform.The elongated Al6Mn phase decreases continuously,and the spherical Al6Mn phase increases (Figs.3(c)and(d)).Compared with the cold-punched specimens in the first three passes,the precipitation phases in the specimens through rapid cold punching of four passes and recrystallization annealing need to be observed clearly at a larger multiple.A larger multiple was used to observe the region near the grain boundary (the square area in Fig.3(e)).It is found that the nano-sized precipitation phases are uniformly distributed in the matrix,and the size and morphology of the precipitation phases are basically the same(Fig.3(f)).3.2Grain morphologyFigure5shows TEM micrographs of Al−Cu−Mg alloy specimens after rapid cold punching of different passes and recrystallization annealing.With the increase of deformation passes,both the deformation and recrystallization degrees of the specimens increase,the grain size becomes finer and finer,and the grain structure becomes more uniform.Incomplete recrystallization occurs in as-extruded alloy specimens annealed at763K for30min.Substructure with high dislocation density still exists near the recrystallization zone.The recrystallized grains are mainly in micron size,and a small number of recrystallized grains are found in the vicinity of coarse recrystallized grains(Fig.5(a)).After rapid cold punching of one pass or two passes,the grain size of the specimens is obviously refined,the recrystallization degree increases and the dislocation density decreases,but it is still incomplete recrystallization.The recrystallized grains are mainly nanocrystalline,and some recrystallized grains grow to coarse grains(Figs.5(b)and(c)).After rapid cold punching of three passes and recrystallization annealing, the specimens are fully recrystallized,and the microstructure becomes uniform.There are no coarse recrystallized grains,the recrystallized grains are all nanocrystalline,the average grain size is less than 100nm,and the dislocation density is further reduced (Fig.5(d)).When the specimens are subjected torapid Fig.5TEM images of Al−Cu−Mg alloy specimens under different conditions:(a)As-extruded;(b)One-pass;(c)Two-pass;(d)Three-pass;(e)Four-passZe-yi HU,et al/Trans.Nonferrous Met.Soc.China29(2019)1816−1823 1820cold punching of four passes and recrystallization annealing,the fully recrystallized structure is more uniform and the recrystallized grains are mainly pared with the specimens through rapid cold punching of three passes,the grain size is further refined and the average grain size is less than50nm. Due to the increase of deformation degree and strain rate, the high dislocation density produced by cold punching still exists in the local region(Fig.5(e)).3.3Deformation band and transition bandTEM images of deformation and transition bands in Al−Cu−Mg alloy specimens cold-punched with different passes are presented in Fig.6.Deformation band with about100nm in width can be observed in coarse grains after single-pass rapid cold punching(Fig.6(a)).When the rectangular region in Fig.6(a)is further enlarged,the boundary of the deformation band with about10nm in thickness can be clearly observed,which is the transition band(Fig.6(b)).The transition band is obviously narrowed after rapid cold punching deformation of two passes(Fig.6(c)),and can hardly be observed in the case of rapid cold punching deformation of three passes (Fig.6(d)).4Analysis and discussion4.1Interaction between precipitation phases andrecrystallizationThe dissolving sequence of Al−Cu−Mg alloy from high to low is generally GP region,S'(Al2CuMg), S(Al2CuMg).In this the work,GP zones,S'and S phases are found in Al−Cu−Mg alloy specimens after rapid cold punching of different passes and recrystallization annealing(Fig.7).The GP region consists of Cu and Mg atom pairs enriched on{110}crystal plane[1](Fig.7(a)). These atom pairs strengthen the alloy by pinning dislocations.The S'phase,which is semi-coherent with the matrix,is mainly formed(Fig.7(b)).The non-coherent equilibrium phase of granular S phase is formed undergoing four passes(Fig.7(c)).Previous studies have shown that the precipitation sequence of precipitates in Al−Cu−Mg alloy is related to heating temperature,deformation amount and grain size after deformation[15,16].If the heating temperature can eliminate the high stress caused by strong deformation, the sequence is first transition phase,and then stable phase.If the heating temperature cannot eliminatethe Fig.6TEM images of deformation and transition bands in Al−Cu−Mg alloy specimens under different conditions:(a,b)One-pass;(c)Two-pass;(d)Three-passZe-yi HU,et al/Trans.Nonferrous Met.Soc.China29(2019)1816−18231821 Fig.7Phase morphologies of Al−Cu−Mg alloy specimens:(a)GP zones;(b)S'phase;(c)S phasehigh stress,the grain size is ultra-fine.The transition phase is inhibited,and the stable phase is generated directly when reprecipitated.Based on the spray forming and rapid solidification technique,fine-grained Al−Cu−Mg alloy billet is prepared in this work.The Al−Cu−Mg alloy cartridge is prepared by multi-pass rapid cold punching,high temperature recrystallization annealing,rapid heating and slow cooling.With the increase of deformation passes,the density of precipitation phases increases,and their size decreases significantly(Fig.3).A large number of S'phases can be observed in the specimen,which indicates that the main dissolving process during the forming of cartridge is first transition phase,and then the stable phase(Fig.7). Meanwhile,the grain structure of the specimen tends to be homogeneous and the grains are refined to nanocrystalline(Fig.5).Further analysis shows that the main reason for the above phenomena is the interaction between precipitates and recrystallization under the condition of rapid cold punching and high temperature recrystallization annealing[5,6].Rapid cold punching significantly increases the dislocation density and becomes the most effective absorption source of vacancy, thus increasing the number of vacancies diffused to dislocation.S'phase nucleates preferentially at dislocation.The high density dislocation introduced by rapid cold punching provides the effective nucleation sites for S'phase,thereby the nucleation number of S' phase increases with the increase of rapid cold punching passes.It was reported that dissolving and recrystallization compete and interact with each other during the recrystallization annealing of supersaturated solid solution after deformation,and that the recrystallization process depends on the instantaneous equilibrium of the dissolving and recrystallization[17].In this work,the defect introduced by rapid cold punching promotes dissolving and recrystallization nucleation,and the dissolving phase particles in turn pin the grain boundaries,thereby affecting the recrystallization nucleation and growth,so as to delay the recrystallization.Further study shows that the rapid cold punching and recrystallization annealing process can obviously promote the recrystallization of spray-formed Al−Cu−Mg alloy and refine the microstructure of the specimens.The main reasons are as follows.Firstly, rapid heating in this experiment makes the dissolving particles too late to produce,reduces effectively the recrystallization temperature,and promotes the occurrence of recrystallization,leading to the precipitates formed during the subsequent recrystallization annealing to affect the recrystallization to proceed in the recrystallized grains.Obviously,the recrystallization annealing temperature is the critical factor affecting precipitation and recrystallized grains[18].Secondly, there is a simple geometric relationship between the recrystallized grain size D N and precipitation phase volume fraction f and particle radius r[18,19]:D N≈2rf−1/3(1)As can be seen in Eq.(1),the larger the volume fraction of precipitation phase is,the smaller the radius of particles is,the finer the recrystallized grains are, which is consistent with the present result.Thirdly,the larger the deformation amount is,the larger the dislocation density is,thus increasing the deformation storage energy and the driving force of recrystallization. Meanwhile,the larger the deformation amount is,the greater the maximum orientation difference of the precipitation phase edge is.The larger orientation gradient is conducive to recrystallization nucleation,thus promoting the microstructure homogeneity and grain refinement of the alloy.Finally,certain amount of Mn is added to the alloy,resulting in intragranular segregation due to the grain boundary adsorption phenomenon of Mn. Elongated Al6Mn phase is easily formed during the rapid cold punching and recrystallization annealing(Fig.3(b)). The coarse Al6Mn phase can further improve the deformation storage energy and increase the orientation difference at the edge of precipitates,thereby promoting the occurrence recrystallization.Ze-yi HU,et al/Trans.Nonferrous Met.Soc.China29(2019)1816−1823 18224.2Formation mechanism of deformation band andits effect on grain refinementDuring large deformation of high stacking fault energy aluminum alloy polycrystal,due to different crystallographic orientations of adjacent grains at grain boundaries,the deformation of each grain must be coordinated with the adjacent grains to maintain the continuity of polycrystal deformation.In this experiment, the specimens are divided into different orientation zones in coarse grains,i.e.deformation bands,because of the inhomogeneous stress of grains propagating to adjacent grains or the instability of grains during plastic deformation.Figure8(a)shows the HRTEM image of the deformation and transition bands after rapid cold punching of one pass,and Fig.8(b)shows schematic diagram of deformation and transition bands.As can be seen,region B is the original orientation region of the grain,and region C is the deformation band with different orientations from the original grain.Region A is the transition band with drastic change of orientation from region B to region C,and the orientation change across the transition band A has great gradient orientation. Studies indicate[1,20]that transition band A can be either a wide orientation region or a narrow orientation region,and can be transformed into the large angle grain boundary,i.e.deformation-induced grain boundary, under high strain rate deformation condition.The mechanism of transformation from transition band to deformation-induced grain boundary can effectively refine coarse grains and play an important role in obtaining uniform fine grain structure.Studies[20,21]demonstrate that the occurrence of deformation band in high stacking fault energy polycrystal depends on the microstructure and deformation condition,and that the grain orientation determines the rotation of grains during deformation. The rotation of each part of coarse deformed grains varies greatly under the action of adjacent fine grains, resulting in deformation band or transition band.Lower deformation temperature will increase the inhomogeneity of deformation,so it helps to cause deformation band.In this experiment,the grain morphologies of Al−Cu−Mg alloy specimens under different states are observed (Fig.5).It is found that the coarse grains exist in extruded and one pass or two passes cold-punched specimens to a certain extent.Therefore,the deformation and transition bands can be observed in the coarse grains (Figs.6(a)and(b)).However,in the specimens after cold punching of three or four passes,the grain morphology tends to be consistent,the grain size is remarkably uniform and the deformation band is difficult to be observed in the specimens(Fig.6(d)).It can be seen that the combination of rapid cold punching deformation and recrystallization annealing process,and the formation of deformation and transition bands in the coarse grains under the process plays a decisive role in the grain refinement of the alloy,and there is a high correlation amongthem.Fig.8HRTEM image(a)and schematic diagram(b)of deformation and transition bands in Al−Cu−Mg alloy specimens undergoing one-pass deformation5Conclusions(1)The precipitates of Al−Cu−Mg alloy during rapid cold punching and recrystallization annealing mainly consist of S phase and a small amount of coarse Al6Mn phase.With the increase of deformation passes, the density of precipitates increases and the size decreases significantly.(2)Rapid cold punching promotes the dissolving and recrystallization nucleation of the specimens by introducing defect,and contributes to the nucleation and growth of S'phase and recrystallization,and thus obtaining the nanocrystalline structure and dispersed S phase.(3)Rapid cold punching favors the formation of deformation and transition bands in the coarse grains, and these bands transform into deformation-induced grain boundary during the deformation and recrystallization,thus effectively refining grains and obtaining uniform fine grains.Ze-yi HU,et al/Trans.Nonferrous Met.Soc.China29(2019)1816−18231823References[1]WANG Zhu-tang,TIAN er manual for Al alloys andprocessing version[M].3rd ed.Changsha:Central South University Press,2007.(in Chinese)[2]WILLIAMS J C,STARKE J E.Progress in structural materials foraerospace systems[J].Acta Materialia,2003,51(19):5775−5799. [3]WU Cui-lan,ZHOU Bin,NIU Feng-jiao,DUAN Shi-yun,GONGXiang-peng,CHEN Jiang-hua.Deformation-inducedΩphase precipitation strengthening of AlCuMg alloy with high Cu/Mg atomic ratio[J].Journal of Hunan University,2018,45(6):1−10.(in Chinese)[4]GAZIZOV M,KAIBYSHEV R.Effect of pre-straining on the agingbehavior and mechanical properties of an Al−Cu−Mg−Ag alloy[J].Materials Science and Engineering A,2015,625(14):119−130. [5]ZHAO Y L,YANG Z Q,ZHANG Z.Double-peak age strengtheningof cold-worked2024aluminum alloy[J].Acta Materialia,2013, 61(5):1624−1638.[6]ZUIKO I S,GAZIZOV M R,KAIBYSHEV R O.Effect ofthermomechanical treatment on the microstructure,phase composition,and mechanical properties of Al−Cu−Mn−Mg−Zr alloy [J].Physics of Metals and Metallography,2016,117(9):906−919. [7]XU Fu-shun,ZHANG Jin,DENG Yun-lai,ZHANG Xin-ming.Effectof snake rolling on strength,toughness and microstructure of Al−Cu−Mg alloy plate[J].The Chinese Journal of Nonferrous Metals,2017,27(10):2005−2011.(in Chinese)[8]MIRZAEI M,ROSHAN M R,JENABALI JAHROMI S A.Microstructure and mechanical properties relation in cold rolled Al2024alloy determined by X-ray line profile analysis[J].Materials Science and Engineering A,2015,620:44−49.[9]AN Li-hui,CAI Yang,LIU Wei,YUAN Shi-jian,ZHU Shi-qiang,MENG Fan-cheng.Effect of pre-deformation on microstructure and mechanical properties of2219aluminum alloy sheet by thermomechanical treatment[J].Transactions of Nonferrous Metals Society of China,2012,22(2):370−375.[10]LI Zhou-bing,SHEN Jian,LEI Wen-ping,YAN Liang-ming,LIJun-peng,MAO Bai-ping.Effects of prestretching on precipitated phase and mechanical properties of Al−Cu−Mg−Ag alloy[J].The Chinese Journal of Nonferrous Metals,2010,20(8):1508−1512.(in Chinese)[11]STYLES M J,MARCEAU R K W,BASTOW T J,BRAND H E A,GIBSON M A,HUTCHINSON C R.The competition between metastable and equilibrium S(Al2CuMg)phase during the decomposition of Al−Cu−Mg alloys[J].Acta Materialia,2015,98(1): 64−80.[12]LI Hui-zhong,LIU Ruo-mei,LIANG Xiao-peng,DENG Min,LIAOHui-juan,HUANG Lan.Effect of pre-deformation on micro-structures and mechanical properties of high purity Al−Cu−Mg alloy [J].Transactions of Nonferrous Metals Society of China,2016,26(6): 1482−1490.[13]YIN Mei-jie,CHEN Jiang-hua,LIU Chun-hui.Effect of interruptedageing on mechanical property and microstructure of AA2024alloy [J].The Chinese Journal of Nonferrous Metals,2015,25(12): 3271−3281.(in Chinese)[14]YANG Pei-yong,ZHENG Zi-qia,XU Fu-shun,LI Shi-chen,LI Jian,ZHOU Ming.Effect of external stress on kinetics of precipitation and morphologies of precipitates in Al−Cu−Mg alloy with high Cu/Mg ratio[J].Rare Metals,2006,30(3):324−328.[15]ZHAO Yun-long,YANG Zhi-qing,ZHANG Zhen,SU Guo-yue,MAXiu-liang.Double-peak age strengthening of cold-worked2024 aluminum alloy[J].Acta Materialia,2013,61(6):1624−1638. [16]MOGHANAKI S K,KAZEMINEZHAD M.Effects ofnon-isothermal annealing on microstructure and mechanical properties of severely deformed2024aluminum alloy[J].Transactions of Nonferrous Metals Society of China,2017,27(1): 1−9.[17]ZHOU Ze-peng,ZHANG Jing,DENG Yun-lai,ZHANG Xin-ming.Creep forming heat treatment technology of Al−Cu−Mg alloy[J].The Chinese Journal of Nonferrous Metals,2017,27(8):1607−1614.[18]PONGE D,GOTTSTEIN G.Necklace formation during dynamicrecrystallization:Mechanisms and impact on flow behavior[J].Acta Materialia,1998,46(1):69−80.[19]HUMPHREYS F J.A unified theory of recovery,recrystallizationand grain growth,based on the stability and growth of cellular microstructures—I[J].Acta Materialia,1997,45(10):4231−4240. [20]ZIEGENBEIN A,HAHNER P,NEUHAUSER H.Correlation oftemporal instabilities and spatial localization during Portiven-Le Chatelier deformation of Cu−10at.%Al and Cu−15at.%Al[J].Computational Materials Science,2000,19(3):27−34.[21]JIANG Hui-feng,ZHANG Qing-chuan,JIANG Zhen-yu,ZHAOSi-min,CHEN Zhong-jia,WU Xiao-ping.Investigation on the Portevin-Le Chatelier deformation bands in Al−Cu alloys[J].Journal of Experimental Mechanics,2004,19(4):430−436.Al−Cu−Mg合金在快速冷冲及再结晶退火过程中的显微组织演变胡泽艺1，范才河1，郑东升1，刘文良1，陈喜红21.湖南工业大学冶金与材料工程学院，株洲412007；2.中国中车株洲电力机车有限公司，株洲412007摘要：采用透射电镜技术(TEM)系统研究喷射成形快速凝固细晶Al−Cu−Mg合金在快速冷冲及再结晶退火工艺过程中的显微组织演变。

Al-Cu-Mg-Ag新型耐热铝合金的抗腐蚀性能齐浩;刘晓艳;梁顺星;张喜亮;崔好选;高飞;陈啟怀【摘要】The effects of the aging treatment on the microstructure and the corrosion resistance of Al-Cu-Mg-Ag heat-resistant aluminium alloy were studied by intergranular corrosion, exfoliation corrosion, electrochemical analysis and transmission electron microcopy. The results show that with increasing aging time, the precipitations both in the grains and on the grain boundaries grow, the distribution of the precipitations on the grain boundaries transforms from continuous to discontinuous, and the precipitation free zone (PFZ) widens, the intergranular corrosion resistance and the exfoliation corrosion resistance decrease. The self-corrosion potential of the matrixφmatrix, precipitations on the grain boundariesφθ and the PFZ testedφPFZ both in the intergranular corrosion solution and exfoliation corrosion solution satisfyφmatrix>φθ>φPFZ. The corrosion resistance of Al-Cu-Mg-Ag alloy is determined by PFZ with the lowest self-corrosion potential. With increasing aging time, the PFZ widens with less Cu solution atoms, leading to lower potential and wider corrosion passageway. This results in the larger potential difference between the matrix and the PFZ, and the corrosion resistance correspondingly decreases. The corrosion resistance of Al-Cu-Mg-Ag heat-resistant aluminium alloy changing from high to low are as follows: under-aged, peak-aged, over-aged.%通过晶间腐蚀、剥落腐蚀、电化学和透射电镜等方法研究热处理制度对Al-Cu-Mg-Ag耐热铝合金组织与抗腐蚀性能的影响。

冷拔连续铸造单晶和多晶铜线材的退火行为陈建;严文;刘建康;李炳;范新会【摘要】采用光学金相以及电子背散射衍射技术,对单晶连续铸造以及传统连续铸造2种技术制备的单晶和多晶铜线材冷拔变形组织的再结晶温度、完全再结晶后的晶粒尺寸以及再结晶组织的孪晶界等进行分析,并比较两者之间的差异.研究结果表明,与多晶铜线材相比,单晶铜线材的再结晶温度较高,完全再结晶的晶粒尺寸较大；随变形量的增加,两者之间的再结晶温度以及再结晶后晶粒尺寸的差异减小.随退火温度以及塑性变形量的增加,冷拔铜线材单个晶粒的孪晶数量减少.与冷拔多晶铜线材相比,单晶铜线材再结晶后孪晶数量明显增多,随塑性变形量以及退火温度的增加,冷拔单晶与多晶铜线材单个晶粒内孪晶数量的差异明显减小.【期刊名称】《粉末冶金材料科学与工程》【年(卷),期】2011(016)005【总页数】6页(P641-646)【关键词】单晶铜;再结晶;冷拔;电子背散射衍射【作者】陈建;严文;刘建康;李炳;范新会【作者单位】西安工业大学材料与化工学院,西安710032;西北工业大学应用物理系,西安710072;西安工业大学材料与化工学院,西安710032;西安工业大学材料与化工学院,西安710032;西安工业大学材料与化工学院,西安710032【正文语种】中文【中图分类】TG114.3随着电子工业和通讯技术的迅猛发展，要求超细丝的直径越来越细。

21世纪初，集成电路的键合丝直径已小至20～25 μm，目前已达12～15 μm。

虽然我国超细丝的需求居全球第一，但制备直径在15 μm以下的金属丝还比较困难，且断线率较高，每轴丝的长度小于1 000 m，而国外已达3 000 m。

超细丝采用冷拔变形获得，材料的品质是超细丝制备的关键因素之一。

在变形过程中，晶界不仅阻碍位错运动，还是杂质聚集之处，这会显著降低线材的拉丝性能。

因此，单晶金属线材是制备超细丝的理想材料[1-4]。

制备单晶金属线材所用的单晶连铸技术(Ohno continuous casting，简称OCC)由日本学者大野笃美发明并投入应用[5-6]。

al-4cu-mg合金半固态压缩过程中的微观组
织演变
al-4cu-mg合金半固态压缩过程中的微观组织演变是一个复杂而
重要的过程。

合金在半固态状态下具有高变形性能和良好的成形性能，使得半固态成形技术成为目前先进制造技术的重要方向之一。

在半固态压缩过程中，合金的微观组织演变是由固态颗粒、液态
相和半固态相之间的相互作用所驱动的。

在初期阶段，液态相首先进
入半固态区域，然后由于自重和表面张力的作用，液态相向空隙处移动，形成了空隙，半固态相在空隙中成长。

当固相颗粒密度增加并占
据了大量的空隙时，半固态相会在空隙中脱离固相颗粒并成为孤立的
团簇。

这些团簇会合并，成为连续的半固态相，从而形成了一种网络
结构，并随着流变应力的增加而变得更加致密。

在半固态压缩过程中，随着固相颗粒的消耗和半固态相的形成，
合金的相变性也会发生变化。

在初期阶段，合金中的液态相和过饱和
固溶质会随着半固态相的形成逐渐减少，最终转化为晶粒边缘的固溶体。

同时，随着半固态相的团簇合并，半固态相的成分也会发生变化，最终形成了高密度的紧密堆积结构，达到了半固态压缩的效果。

总之，al-4cu-mg合金半固态压缩过程中的微观组织演变是由多
种因素共同作用的结果，包括固态颗粒、液态相和半固态相的相互作用、流变应力的作用以及相变的影响等。

通过形变时效工艺同时提高Al-Mg-Si-Cu合金强度和电导率陈敬;陈江华;刘春辉;赖玉香;顾媛【摘要】铝是一种优良的导电材料,但由于强度低,其应用受到很大限制.随着铝在电力工业中应用逐渐增加,近年来,越来越多的工作致力于提高铝的导电率与强度的综合性能.通过改变传统T6时效工艺顺序发明一种同时显著提高Al-Mg-Si-Cu合金导电率和强度的形变时效工艺.本文采用显微硬度测量,导电率测试以及透射电镜(TEM)微观结构表征研究了形变时效工艺与传统T6时效工艺制备的材料在综合性能和微观组织上的差异.轧制变形引入的位错在后续时效过程调控析出,析出相形貌的改变是导电率相对T6工艺提高的原因,而残留位错可提高材料强度.【期刊名称】《功能材料》【年(卷),期】2016(047)002【总页数】5页(P2139-2142,2147)【关键词】铝合金;导电率;强度;位错;析出相【作者】陈敬;陈江华;刘春辉;赖玉香;顾媛【作者单位】湖南大学材料科学与工程学院,长沙410082;湖南大学材料科学与工程学院,长沙410082;湖南大学材料科学与工程学院,长沙410082;湖南大学材料科学与工程学院,长沙410082;湖南大学材料科学与工程学院,长沙410082【正文语种】中文【中图分类】TG113;TM241铝在电器制造工业、电线电缆工业和无线电工业中有广泛的用途。

在铝材料中商业纯铝拥有最高的导电率，大约为62%IACS(国际退火铜标准)，然而它的抗拉强度仅仅只有160 MPa左右，这使得它的应用受到很大限制[1-2]。

为了提高铝材料在电力相关行业中的应用，我们必须在保证它优良的导电性能的情况下，尽可能地提高强度。

早期的工作者通过合金化来提高强度，比如添加Mg和Si元素形成的Al-Mg-Si系合金[2]。

虽然合金强度获得了一定的提升，但是金属的导电性能对于其微观结构很敏感。

溶质原子、晶格自振动和缺陷都会成为电子运动的散射源，阻碍电子的运动，从而使铝合金的导电性能下降。

Al-Cu-Mg-Ag耐热铝合金均匀化处理刘晓艳;潘清林;陆智伦;刘畅;何运斌;李文斌【摘要】研究了均匀化温度和均匀化时间对Al-Cu-Mg-Ag耐热铝合金微观组织的影响,优化了合金的均匀化处理制度,并对其均匀化过程进行了动力学分析.结果表明,铸态合金组织中存在严重的枝晶偏析,晶界上有大量残留相,各元素在晶内和晶界分布不均匀.随着均匀化温度的升高或均匀化时间的延长,合金组织中的残留相逐渐溶解,晶界变得稀薄,元素分布趋于均匀.该合金过烧温度为520℃.适宜的的均匀化制度为510℃×12 h,这与均匀化动力学分析得到的结果基本相符.%The effect of the homogenization temperature and time on the microstructure of Al-Cu-MgAg heat-resisted alloy was studied,the homogenization process of the alloy was also optimized,and the kinetic analysis of homogenization was carried out.The results showed that the serious dendrite microstructure existed in Al-Cu-Mg-Ag alloy ingot,and there were many residual phases in grain boundaries and the elements were unevenly distributed from grain boundary to inside.With the increasing of homogenization temperature or time,the residual phases dissolved,the grain boundaries became sparse and all elements became more homogenized.The overburn temperature of the alloy was 520℃,and the suitable homogenizing process of Al-Cu-Mg-Ag alloy was 510℃×12 h,which agreed well with the kinetic analysis of homogenization.【期刊名称】《材料科学与工艺》【年(卷),期】2011(019)004【总页数】5页(P28-32)【关键词】Al-Cu-Mg-Ag合金;均匀化;显微组织;动力学方程【作者】刘晓艳;潘清林;陆智伦;刘畅;何运斌;李文斌【作者单位】中南大学材料科学与工程学院,长沙410083;河北工程大学装备制造学院,河北邯郸056038;中南大学材料科学与工程学院,长沙410083;中南大学材料科学与工程学院,长沙410083;中南大学材料科学与工程学院,长沙410083;中南大学材料科学与工程学院,长沙410083;中南大学材料科学与工程学院,长沙410083【正文语种】中文【中图分类】TG146.212×××系耐热铝合金如2219和2618由于具有较高的强度和良好的耐热性能，被广泛用于航空航天领域.随着航空航天技术的发展，对铝合金材料工作温度的要求也越来越高.有研究表明，向高铜镁比Al－Cu－Mg合金中添加Ag后可改变合金的时效序列，析出1种均匀细小弥散的耐热强化相——Ω相，此相可在200℃以下长期存在而不发生聚集长大［1－2］. 因此，Al－Cu－Mg－Ag新型合金有望满足超音速飞机的经济性要求及耐热性能要求，是超音速飞机备选材料的一个极有希望的发展方向.铝合金在凝固时都存在枝晶偏析［3－4］，元素在晶内和晶界分布不均匀.这种组织和成分的不均匀性会降低合金塑性、恶化其热加工性能，降低成品强度和塑性.因此，必须对合金铸锭进行均匀化处理，以消除或降低化学成分和组织的不均匀性［5－6］，从而改善合金的性能.游文等［7］采用双级均匀化制度对Al－Cu－Mg－Ag合金进行处理，发现在420℃ ×6 h+ 515℃×6 h的工艺下，合金均匀化效果最好，晶内偏析基本消除，晶间组织分布均匀.经热轧、固溶和时效处理，合金的抗拉强度达到470 MPa以上，对应的伸长率也达到8%～10%.因此，采用适当的均匀化处理，对合金成品的性能起着至关重要的作用.目前，对Al－Cu－Mg－Ag合金均匀化工艺研究较多［8－9］，而有关均匀化过程中组织演化的研究很少.本文采用单级均匀化工艺，研究了均匀化温度和均匀化时间对Al－Cu－Mg－Ag合金组织和元素分布的影响，探讨了均匀化过程中的组织演化，得出了适宜的均匀化工艺，并基于扩散理论建立了合金的均匀化动力学方程.实验所用原材料为工业纯铝、纯镁和纯银以及A1－Cu、A1－Mn和A1－Zr中间合金.采用铸锭冶金方法制备了Al－5.3Cu－0.8Mg－0.5Ag－0.3Mn－0.15Zr(质量分数，%)合金铸锭.将铸锭加工成12 mm×12 mm×15 mm的试样，分别在470、480、490、500、510和520℃下均匀化处理24 h，然后在选定的最佳均匀化温度下处理8～48 h.采用飞利浦 Sirion200场发射扫描电镜和POLYVER－MET金相显微镜观察合金的微观组织，用扫描电镜上配套的EDS设备对合金相进行能谱分析.差示扫描量热分析(DSC)在SDTQ600热分析仪上进行.图1为合金的铸态微观组织.由图1可见，铸态组织呈树枝状(图1(a))，在晶界处存在粗大的残留相(图1(b)).这些相大部分呈灰色，经EDS能谱分析可知其成分接近Al2Cu.其中也有小部分颜色较亮的相，能谱分析结果如表1所示.图2为铸态合金的SEM组织与合金中主要元素Cu、Mg和Ag在晶内和晶界的分布情况.由图2可见，合金中的主要元素在晶界处存在不同程度的富集，其元素偏析程度为Cu＞Mg＞Ag.由以上实验结果可知，铸态合金在晶界处存在大量的残留相，各元素在晶内及晶界分布不均匀，必须对其进行均匀化处理.均匀化过程中，扩散系数与温度的关系为式中:D0为与温度无关的系数;T为绝对温度;Q为扩散激活能;R为气体常数.由式(1)可知，均匀化温度越高，扩散系数越大，原子的扩散速度越快，偏析就越容易消除.然而，为了防止过烧，必须先确定合金均匀化的最高温度.图3为合金的DSC曲线.从图3中可以看出，合金铸锭在495.9、526.1和643.4℃处有3个吸热峰.低温495.9℃处有一个微弱的吸热峰，经均匀化处理后完全消失，所以该峰对应的可能是某一个相的溶解温度，此相在均匀化过程中溶入基体.526.1℃处的吸热峰强度随着均匀化温度的升高先增强后逐渐减弱，均匀化温度为510℃时，仍有微弱的吸热峰，当均匀化温度提高到520℃时，吸热峰基本消失.因此，526.1℃是低熔点共晶相的熔化温度，而643.4℃为合金的熔化温度，由此确定合金的均匀化温度应低于526.1℃.图4给出了合金在不同温度均匀化处理24 h后的金相组织.由图4可见，随均匀化温度的升高，枝晶网络逐渐消失，晶界上残留相的分布逐渐变得不连续.经510℃均匀化处理后合金中的枝晶网络减少，大部分残留相溶入基体.当均匀化温度进一步升高到520℃时，合金严重过烧，组织中出现晶界复熔物和复熔三角形.由以上分析可见，在保证合金不过烧的情况下，Al－Cu－Mg－Ag合金适宜的均匀化温度为510℃.图5为合金在510℃下经不同时间均匀化处理的金相组织.由图5可见，随均匀化时间的延长，晶界上粗大的残留相逐渐溶解.在510℃均匀化处理8 h时，合金晶界上粗大的残留相和枝晶网络均有所减少.均匀化处理12 h时，合金组织中的残留相基本溶解，晶界变得稀薄(图5(b))，此后再延长均匀化时间，效果也不明显(图5(c)和(d)).综合考虑合金均匀化处理的有效性与经济性，Al－Cu－Mg－Ag合金适宜的均匀化处理时间为12 h.图6为合金铸锭经510℃×12 h均匀化处理后的线扫描分析结果.从图6中可以看出，经过均匀化处理后，合金中的主要元素Cu、Mg和Ag在晶界的偏析基本消除，从晶界至晶内的分布趋于平稳，但Cu元素在晶界处仍有少量偏聚.图7为铸态合金主要元素沿枝晶间分布的线扫描分析结果.从图7可以看出，Al－Cu－Mg－Ag合金中的主要元素的浓度沿枝晶间呈周期性变化.Hillert等［10］的研究结果表明，在存在偏析的铸态组织中，固溶体内部合金元素的含量比枝晶处的含量低很多，各合金元素的浓度沿枝晶间的分布呈周期性变化.这与本实验结果一致.所以，研究合金中各元素在均匀化过程中的变化规律，只需要研究相邻枝晶间合金元素的扩散规律即可.Shewman［11］认为均匀化过程中，合金中各元素的分布状态可以用余弦函数的傅氏级数分量逼近，表示为其中，¯C为元素平均浓度， L为枝晶间距，A0为初始偏聚振幅，可表示为式(2)和(3)中各变量如图8所示.振幅随着均匀化时间的增加而衰减，可表示为A(t).根据Fick定律和边界条件，A(t)可表示为式中，t为均匀化时间，其余各物理量含义同上.把式(1)代入式(4)，得从式(5)可以看出，随着均匀化温度T的升高或均匀化时间t的延长，枝晶间的偏聚减少，这与本实验中观察到的结果相吻合.通常，当各元素偏聚振幅减小为1%时，均匀化过程结束，则即两边取自然对数，得假设，可得式(9)即为Al－Cu－Mg－Ag合金的均匀化动力学方程.只要给出合金铸锭组织参量便可作出其均匀化动力学曲线.在相同温度下，Mg、Ag和Mn元素的扩散比Cu元素快［12］，因此，主要考虑Cu元素的扩散.将参量D0(Cu)=0.084 cm2/s，Q(Cu)= 136.8 kJ/mol，R=8.31 J/(mol·K)代入式(9)即可作出不同组织参量下Cu元素的均匀化动力学曲线(图9).由图9可见，同一组织参量下，随均匀化温度的升高，合金均匀化时间大大缩短.由定量金相测出Al－Cu－Mg－Ag合金铸态枝晶平均间距L为42 μm.由2.2节实验结果可知，合金铸锭理想的均匀化温度为510℃，代入式(9)计算可得合金的均匀化时间为10 h，与实验结果510℃×12 h基本相符.合金铸锭经均匀化退火处理后，晶界上的残留相基本溶解，但晶界处仍存在少量的Cu偏聚，这可能是由于合金中Cu含量较高引起的.1)Al－Cu－Mg－Ag耐热铝合金铸态组织中存在大量枝晶组织，各元素在晶内和晶界分布不均匀.随均匀化温度的升高或均匀化时间的延长，合金中的非平衡相逐渐溶解，晶界变得稀薄，元素分布趋于均匀.2)该合金的过烧温度为520℃.实验得到的适宜的均匀化制度为510℃×12 h.3)合金的均匀化方程为，由此得到的均匀化制度为510℃ ×10 h，与实验结果基本相符.【相关文献】［1］ RINGER S P，HONO K，POLMEAR I J，et al.Nucleation of precipitates in aged Al－Cu－Mg－(Ag) alloys with high Cu:Mg ratios［J］.Acta Materialia，1996，44(5):1883－1898.［2］ CHESTER R J，POLMEAR I J.Precipitation in Al－Cu－Mg－Ag alloys［J］.Acta Metallurgica，1989，37 (3):777－789.［3］ LI Nian-kui，CUI Jian-zhong.Microstructural evolution of high strength 7B04 ingot during homogenization treatment［J］.Transactions of Nonferrous Metals Society of China，2008，18(4):769－773.［4］ ZHANG Jing，ZUO Rul-in，CHEN Youxing，et al.Microstructure evolution during homogenization of a τ－type Mg－Zn－Al alloy［J］.Journal of Alloys and Compounds，2008，448(1－2):316－320.［5］ TOTIK Y，SADELER R，KAYMAZ I，et al.The effect of homogenisation treatment on cold deformations of AA 2014 and AA 6063 alloys［J］.Journal of Materials ProcessingTechnology，2004，147(1):60－64.［6］ FAN Xi-gang，JIANG Da-ming，MENG Qing-chang，et al.The microstructural evolution of an Al－Zn－Mg－Cu alloy during homogenization［J］.Materials Letters，2006，60(12):1475－1479.［7］游文，余日成，刘志义.Al－Cu－Mg－Ag系高强耐热合金的热加工工艺研究［J］.铝加工，2006，166:12－17.［8］李云涛，刘志义，夏卿坤，等.含铈Al－Cu－Mg－Ag－Mn－Zr铝合金均匀化工艺的研究［J］.金属热处理，2006，31(10):55－57.［9］李云涛，刘志义，夏卿坤，等.Er在Al－Cu－Mg－Ag合金中的存在形式及其均匀化工艺［J］.中南大学学报(自然科学版)，2006，37(6):1043－1047.［10］ HILLERT A.合金扩散和热力学［M］.赖和怡，刘国勋译.北京:冶金工业出版社，1999. ［11］ SHEWMAN P G.Diffusion in Solids［M］.New York:McGraw-Hill，1963.［12］ MONDOLFO L F.Aluminium alloys structure and properties［M］.London-Boston:Butter Worths，1976.。

铜锌镁铝四元水滑石的微观结构及其姜-泰勒畸变王力耕施炜姚萍倪哲明*李远刘娇(浙江工业大学化学工程与材料学院,先进催化材料实验室,杭州310032)摘要:采用密度泛函理论(DFT),选取CASTEP 程序模块,对铜锌镁铝四元水滑石[(M)-IV-LDHs (M=Cu,Zn,Mg,Al)]周期性模型进行几何全优化,从各体系的结构参数、电子排布、Mulliken 电荷布居、结合能、氢键等方面,研究了体系中的姜-泰勒效应、氯离子位置对层板畸变及体系稳定性的影响.优化结果表明,姜-泰勒效应不仅存在于d 轨道未排满的Cu 2+中,也存在于理论上d 轨道排满的Zn 2+与p 轨道未排满的Mg 2+中.氯离子排在金属上方的体系,其金属畸变程度大于氯离子排在非金属上方的体系.同时,对于本文选取的8个CuAl-IV-LDHs 体系,结合能绝对值按照1-8号逐渐降低,体系的稳定性下降,最终转变为不稳定的压扁的八面体构型.这有助于从理论上对含铜四元水滑石的姜-泰勒效应进一步认识.关键词:四元含铜水滑石;密度泛函理论;姜-泰勒效应中图分类号:O641Microstructure and Jahn-Teller Effect of Cu-Zn-Mg-Al Layered DoubleHydroxidesWANG Li-GengSHI WeiYAO PingNI Zhe-Ming *LI YuanLIU Jiao(Laboratory of Advanced Catalytic Materials,College of Chemical Engineering and Materials Science,Zhejiang University of Technology,Hangzhou 310032,P .R.China )Abstract:We propose a periodic interaction model for layered double hydroxides,CuZnMgAl quaternary hydrotalcites [(M)-IV-LDHs (M=Cu,Zn,Mg,Al)].Based on density functional theory the geometries of CuZnMgAl quaternary hydrotalcites were optimized using the CASTEP program.The impacts of the Jahn-Teller effect and the location of chlorine over the layer distortion and stability were investigated by analyzing the geometric parameters,the electronic arrangement,charge populations,binding energies,and hydrogen-bonding.The optimization results showed that the Jahn-Teller effect does not only exist in Cu 2+when its d orbital is partially filled but it also exists in Zn 2+when its d orbital is full as well as in Mg 2+when its p orbital is partially filled.Systems where the chloride is located above the metal show greater metal distortion than systems with anions located above non-metals.Eight systems (Nos.1-8)were chosen for our work and their absolute binding energy values were found to decrease gradually while the stability of the systems became worse.Finally,the systems became unstable and were found to be flattened octahedral forms.These results help us to better understand the Jahn-Teller effect in copper-containing IV-LDHs from theory.Key Words:Copper-containing quaternary hydrotalcite;Density functional theory;Jahn-Teller effect[Article]doi:10.3866/PKU.WHXB20122858物理化学学报(Wuli Huaxue Xuebao )Acta Phys.-Chim.Sin .2012,28(1),58-64JanuaryReceived:September 5,2011;Revised:October 14,2011;Published on Web:October 24,2011.∗Corresponding author.Email:jchx@;Tel:+86-138********ⒸEditorial office of Acta Physico-Chimica Sinica1引言水滑石(LDHs)是一种层状双羟基复合金属氢氧化物,其通式为[M 2+1-x M 3+x (OH)2]x +(A n -)x /n·y H 2O,其中M 2+表示二价金属阳离子,M 3+表示三价金属阳离子,A n -表示层间可交换的阴离子,y 表示层间水分子的个数.1,2位于层板上的金属阳离子可以在一定的比58王力耕等:铜锌镁铝四元水滑石的微观结构及其姜-泰勒畸变No.1例范围内被离子半径相近的金属阳离子同晶取代,从而形成二元、三元、四元水滑石.目前,对于水滑石材料的理论模拟多为二元、三元结构,3-6而对四元水滑石的理论研究涉及较少.随着水滑石制备技术的进步及其应用领域的拓宽,多元水滑石材料也受到了广泛关注,冯拥军7和吴健松8等分别选用不同方法合成了CuNiMgAl 、CuZnMgAl 四元水滑石,并通过X 射线衍射(XRD)等表征方法证明了其晶型的完整性;Wang 等9发现CuZnMgAl 四元水滑石对加氢反应具有催化活性;张军等10采用共沉淀法合成了NiMgAlLa 四元水滑石,并将其应用于甲烷的部分氧化,性能优越,但从实验手段很难精确获得这些四元水滑石的微观结构、电子性质等特性.因此,计算机模拟技术的引入就显得尤为重要.密度泛函理论(DFT)11,12是一种研究多电子体系电子结构的量子力学方法,目前已被广泛应用于计算水滑石体系的结构参数、成键状况、作用能、电子密度以及了解分子间交互作用的电子性质等.Wei 等3以簇模型[M(OH 2)6]n +(M 为Ca 2+、Mg 2+、Cu 2+、Ga 3+等金属阳离子)为基础,采用混合密度泛函B3LYP 方法讨论了体系中的姜-泰勒效应对结构参数的影响,并在此基础上进一步比较了不同阳离子组成的水滑石的结构和性质.本课题组13曾采用密度泛函理论分析了铜锌铝三元水滑石(Cu x Zn 3-x Al-LDHs,x =0-3)周期性模型的结构参数、姜-泰勒效应、Mulliken 电荷布居、结合能和氢键,探究了体系的畸变情况和稳定性.考虑到铜是常见的催化活性组元,具有价格低廉、毒性低等优点,被广泛应用于各种催化反应中,如Glaser 偶联反应、14Heck 反应、15不对称Henry 合成、16糠醛加氢、17碳杂偶联反应18等.所以,研究含铜水滑石的微观结构具有重要的意义.基于以上原因,本文采用密度泛函理论,构建了一系列氯离子插层的铜锌镁铝四元水滑石模型,从结构参数、Mulliken 电荷布居、结合能、氢键等角度研究了铜锌镁铝四元水滑石的姜-泰勒效应及氯离子位置对层板畸变和体系稳定性的影响,为进一步设计和制备具有姜-泰勒效应的水滑石材料提供理论参考.2计算模型与方法本文以2H 堆积模式1构建了铜锌镁铝四元水滑石的主体层板[Cu 2Zn 2Mg 2Al 2(OH)16]2+,层间氯离子可以排布在层板的Top 位、Hcp 位、Fcc 位以及Bridge 位(见图1(A)),氯离子总共在层板表面存在16种结合位置(见图1(B)).本课题组先前工作19表明氯离子位于Bridge 位和Top 位的构型都不能够稳定存在.因此,本文中只考虑氯离子位于Hcp 位和Fcc 位时的8种情况,共搭建了48种金属离子或氯离子不同排布的铜锌镁铝四元水滑石.采用先前工作中证实对水滑石体系适用的计算方法,19,20细节为:选用CASTEP 程序模块,21在LDA-CA-PZ 基组22水平对模型进行几何全优化,原子电子采用超软赝势,23截止能量为330.0eV ,自洽场计算的误差为2×10-6eV ·atom -1,能带结构在布里渊区k 矢量的选取为4×4×1,基态能量选用Pulay 密度混合算法,24整体电荷数为0,同时优化晶胞,其它参数设置为程序的默认值.优化结果表明,在48个铜锌镁铝四元水滑石模型中,Cu 在晶胞顶点,Al 在晶胞中心的8种模型能量相对较低,记作CuAl-IV-LDHs.本文即选取如图2所示的Cu 在顶点,Al 在中心,氯离子不同排布的8种图1Cl -在(M)-IV-LDHs (M=Cu,Zn,Mg,Al)中的几种不同位置Fig.1Models of (M)-IV-LDHs (M=Cu,Zn,Mg,Al)with different positions of Cl-(A)(B)59Acta Phys.-Chim.Sin.2012Vol.28情况进行分析.3结果与讨论3.1结合能分析为估计体系结构的稳定性,现定义CuAl-IV-LDHs体系的结合能ΔE CuAl-IV-LDHs为:ΔE CuAl-IV-LDHs=E CuAl-IV-LDHs-(16E H+16E O+2E Cl+2E Cu+2E Zn+2E Mg+2E Al)其中E CuAl-IV-LDHs为优化后CuAl-IV-LDHs体系的能量, E H=-1169.2kJ·mol-1、E O=-41364.3kJ·mol-1、E Cl=-39180.6kJ·mol-1、E Cu=-129759.8kJ·mol-1、E Zn=-165066.0kJ·mol-1、E Mg=-93770.7kJ·mol-1、E Al=-5114.7kJ·mol-1,为各原子的能量,结果列在表1中.从表1中可以看出,对于本文选取的CuAl-IV-LDHs体系,层间氯离子排布不同会影响其结合能, 1号到8号体系,结合能绝对值逐渐降低,体系的稳定性下降.氯离子在非Mg原子正上方,即CuAl-FccCuAlZn最稳定;氯离子在Mg原子正上方,即CuAl-HcpMg最不稳定.3.2结构参数分析将几何优化得到的CuAl-IV-LDHs的结构参数列入表2,晶胞参数a值和b值的大小主要归因于层板金属离子半径的影响.2即在CuAl-IV-LDHs体系中,主要是与晶胞棱上的Cu－O/Zn－O/Mg－O/ Al－O键长的大小相关.从表2可知,对于单独的CuAl-IV-LDHs层板(层间无氯离子时,即表2中Layer,下列各表中Layer均指单独的CuAl-IV-LDHs 层板),a=0.59003nm,b=0.59442nm.当氯离子插入层板以后,a、b值略微上升,这可能是由于当氯离表1CuAl-IV-LDHs的结合能Table1Binding energy of CuAl-IV-LDHsNo.1 2 3 4 5 6 7 8SystemCuAl-FccCuAlZnCuAl-HcpAlCuAl-FccMgAlCuCuAl-FccMgAlZnCuAl-HcpCuCuAl-HcpZnCuAl-FccCuZnMgCuAl-HcpMgE CuAl-IV-LDHs(kJ·mol-1)-1569771.8-1569757.0-1569753.4-1569752.2-1569742.3-1569724.0-1569697.7-1569696.1ΔE CuAl-IV-LDHs(kJ·mol-1)-23452.2-23437.4-23433.8-23432.6-23422.7-23404.4-23378.1-23376.5表2CuAl-IV-LDHs的晶胞参数Table2Lattice parameters of CuAl-IV-LDHsNo.12345678SystemLayerCuAl-FccCuAlZnCuAl-HcpAlCuAl-FccMgAlCuCuAl-FccMgAlZnCuAl-HcpCuCuAl-HcpZnCuAl-FccCuZnMgCuAl-HcpMga/nm0.590030.611820.611270.600210.599970.612070.619850.599930.61978b/nm0.594420.598360.600700.611790.609740.601000.601790.607680.60194c/nm1.475501.441251.478271.482861.486461.450741.456791.507461.53055图2CuAl-IV-LDHs模型Fig.2Models ofCuAl-IV-LDHs60王力耕等:铜锌镁铝四元水滑石的微观结构及其姜-泰勒畸变No.1子插入层板以后,Cu－O、Al－O、Mg－O、Zn－O的成键布居总体均降低,分别由单独层板的0.236e、0.382e、-1.145e、0.231e降至0.228e、0.374e、-0.973e、0.223e,使得金属－氧键键能下降,从而导致金属－氧键键长增大,a、b值略微上升.晶胞参数c值(层间距d c=0.5c)不仅受到层板金属离子半径的影响,还受到主客体间作用力大小的影响,2即晶胞参数c值为层板厚度和层间通道大小的加和的2倍.从表2中可知,1-8号体系的晶胞参数c值有增大的趋势,这主要是由于氯离子插入层板后,1-8号体系的结合能逐渐降低(数据见表1),说明体系的主客体作用力逐渐下降,从而使得层间通道增大,c值呈增大趋势.3.3Jahn-Teller效应分析1937年,姜(Jahn,H.A.)和泰勒(Teller,E.)指出:在对称的非线性分子中,如果有一个体系的基态有几个简并能级,则是不稳定的,体系会发生畸变,使能级发生改变,以消除简并性,这就是姜-泰勒效应.25水滑石层板类似于水镁石Mg(OH)2结构,它是由MO6八面体共用棱边所形成的,当其层板上的镁被铜等一些金属替换时,水滑石就会由于姜-泰勒效应发生畸变.对于CuAl-IV-LDHs体系中的Cu2+,d轨道电子数理论上为9,则有可能失去dz2或dx2-y2上的部分电子.若失去的是dz2上的电子,则会变成压扁的八面体构型;若失去的是dx2-y2上的电子,则会变成拉长的八面体构型.25结合表3中Cu－O键键长数据可知,1号体系中的6个Cu－O键出现了四个共面的短键和两个与该面垂直的长键,为典型的拉长的八面体构型,随着氯离子放置位置的变化,从1号到8号体系,短键基本上逐渐拉长,长键逐渐缩短,最终出现了四个共面的长键和两个与该面垂直的短键,为典型的压扁的八面体构型,说明1号到8号体系随着层间氯离子位置的改变,层板上的铜由易失去dx2-y2上的电子,逐渐转变为易失去dz2上的电子,体系的稳定性逐渐下降,这一现象与结合能分析结果一致.实验证明,Cu的六配位配合物,以拉长的八面体构型稳定存在,这是因为在无其它能量因素影响时,形成两条长键四条短键比形成两条短键四条长键的总键能要大的缘故.25表4中列出了CuAl-IV-LDHs体系中的原子轨道布居.从表4可以看出,体系中Zn2+的d轨道电子数为9.990e,相比Cu2+的d轨道电子数(9.425e-9.440e)更接近于饱和的d10状态,即Zn-O八面体配合物理论上不应该出现畸变现象.但从表5数据可以看出,在CuAl-IV-LDHs体系中,金属平均畸变角(指12个θOMO与理想六配位角度90°的绝对差值的平均No.1 2 3 4 5 6 7 8SystemLayerCuAl-FccCuAlZnCuAl-HcpAlCuAl-FccMgAlCuCuAl-FccMgAlZnCuAl-HcpCuCuAl-HcpZnCuAl-FccCuZnMgCuAl-HcpMgl Cu－O1/nm0.190030.190360.191290.191510.191250.190650.191870.190020.18904l Cu－O2/nm0.190280.190630.191510.191750.191320.191170.192060.190170.18916l Cu－O3/nm0.196440.198600.197880.198180.198770.200180.200320.203660.21571l Cu－O4/nm0.196460.198950.197970.198340.198950.200310.200410.203770.21572l Cu－O5/nm0.226240.234620.233580.233430.233300.230250.230640.228530.21704l Cu－O6/nm0.226300.234690.234870.233970.234030.232060.230800.228970.21819表3CuAl-IV-LDHs中的Cu－O键长Table3Cu－O bond lengths in CuAl-IV-LDHsNo.1 2 3 4 5 6 7 8SystemLayerCuAl-FccCuAlZnCuAl-HcpAlCuAl-FccMgAlCuCuAl-FccMgAlZnCuAl-HcpCuCuAl-HcpZnCuAl-FccCuZnMgCuAl-HcpMgQ(Al-p)/e0.8900.8900.8900.9000.8900.9000.9000.9300.900Q(Mg-p)/e5.5705.7705.7105.6705.7005.7205.7255.6705.610Q(Cu-p)/e0.4100.4200.4200.4000.4200.3900.4400.4100.450Q(Cu-d)/e9.4109.4309.4309.4309.4259.4409.4309.4309.430Q(Zn-p)/e0.6500.6500.6600.6800.6600.6700.6300.6500.660Q(Zn-d)/e9.9909.9909.9909.9909.9909.9909.9909.9909.990表4CuAl-IV-LDHs中的原子轨道布居Table4Atomic orbital populatins in CuAl-IV-LDHs61Acta Phys.-Chim.Sin.2012Vol.28值)ΔθOZnO值却大于ΔθOCuO值,在Zn-O八面体中也观察到了姜-泰勒畸变.同时,体系中没有d轨道电子的镁、铝配合物也存在一定程度的金属畸变,且ΔθOMgO值仅略低于ΔθOCuO值.据此可推测,不仅只有d轨道电子的不均匀排布会发生姜-泰勒畸变,镁的p轨道电子的不均匀排布也存在发生姜-泰勒畸变的可能性,因为只要p x、p y、p z中任一轨道上的电子缺失或者多余,导致三个轨道上的电子排布不相同,相应轴上的M－O键就会缩短或者拉长,从而导致畸变的产生.图3为CuAl-IV-LDHs体系平均金属畸变角的分布情况,总体上来说,氯离子排在金属上方的体系(2、5、6、8号体系)其金属八面体畸变程度大于氯离子排在非金属上方的体系,故氯离子排在非金属上方更有利于体系的稳定性.且从表5数据可知,1号到8号体系的平均金属畸变基本呈逐渐增大趋势,说明层板上的金属八面体畸变越来越严重.3.4Mulliken电荷布居分析Mulliken布居是Mulliken26提出的表示电子在各原子轨道上分布情况的方法,它可以间接地讨论分子内相互作用力的强弱,尤其对同一系列的分子十分奏效.因此,为了进一步研究CuAl-IV-LDHs体系的作用力,对其进行了电子分析,得到的Mulliken 电荷布居(对上下的层板、客体氯离子、各原子及键的电荷布居作了平均处理)列入表6和表7中.从表6、7可以看出,当CuAl-IV-LDHs体系中插入氯离子后,Cu、Zn、Mg、Al所带的电荷数均减少,说明氯离子的电荷向层板发生了转移,转移的数量大小顺序为Cu>Zn>Al>Mg,层板与层间氯离子的静电作用增强.且Cu－O、Al－O、Mg－O、Zn－O键布居的绝对值也基本上减少,故层板上M－O键的离子键作用力亦减弱,说明氯离子的引入,虽增加表5CuAl-IV-LDHs中的金属畸变角Table5Metal distortion angle of CuAl-IV-LDHsNo.1 2 3 4 5 6 7 8SystemLayerCuAl-FccCuAlZnCuAl-HcpAlCuAl-FccMgAlCuCuAl-FccMgAlZnCuAl-HcpCuCuAl-HcpZnCuAl-FccCuZnMgCuAl-HcpMgAverage distortion of metal/(°)ΔθOMgO/(°)6.60810.0997.6567.4247.7789.8868.2497.7639.2228.298ΔθOZnO/(°)8.8907.65310.08210.50010.2377.9819.7989.8518.7859.309ΔθOCuO/(°)7.6708.6249.0768.8108.9859.0989.0928.8529.1738.820ΔθOAlO/(°)3.2623.9033.6333.8123.5644.3334.1874.8224.2383.973Average metal distortion/(°)6.6087.5707.6127.6377.6417.8257.8327.8227.855ΔθOMO:systemʹs average metal distortion angle图3CuAl-IV-LDHs平均金属畸变情况Fig.3Average metal distortion ofCuAl-IV-LDHsNo.12345678SystemLayerCuAl-FccCuAlZnCuAl-HcpAlCuAl-FccMgAlCuCuAl-FccMgAlZnCuAl-HcpCuCuAl-HcpZnCuAl-FccCuZnMgCuAl-HcpMgQ(Cu－O)/e0.2360.2370.2300.2280.2230.2170.2270.2270.228Q(Al－O)/e0.3820.3740.3720.3680.3660.3700.3720.3700.374Q(Mg－O)/e-1.145-1.013-1.023-0.992-0.968-0.956-0.966-0.963-0.973Q(Zn－O)/e0.2310.2380.2270.2250.2290.2290.2140.2120.223Q(H－O)/e0.6100.5860.5880.5950.5900.6000.5900.6040.606表6CuAl-IV-LDHs的Mulliken成键布居Table6Mulliken bond population of CuAl-IV-LDHs62王力耕等:铜锌镁铝四元水滑石的微观结构及其姜-泰勒畸变No.1了主客体的作用力,但会相对地削弱层板自身的作用力.同时,纵向比较,从1号到8号体系,Cl -所带的负电荷呈增大趋势,说明从客体向主体层板转移的电子呈下降趋势,整体层板电荷增大,说明随着体系稳定性的减弱,主客体的静电作用力呈增强趋势,而层板自身的成键作用减弱.3.5CuAl-IV-LDHs 的氢键结构分析氢键会对材料的微观结构造成较大的影响,它是一种广泛存在的分子间弱作用力,是特殊的分子间或分子内作用.它是由极性很强的X －H 键上的氢原子与另一个键(可存在于同一种分子或另一种分子中)上电负性很强、原子半径较小的Y 原子(如F 、N 、O 等)的孤对电子之间相互吸引而形成的一种键(用X －H …Y 表示).27-29一般情况下,氢键具有方向性和饱和性,氢键的键长越短、键角越接近于180°,氢键的强度越强.铜锌镁铝四元水滑石体系中存在多重的氢键,它与一般的氢键性质不同,层间氯离子并不只与层板上的一个羟基中的氢原子形成氢键,而是与多个羟基上的氢原子形成多重氢键.本文选取的CuAl-IV-LDHs 体系的氢键数据列于表8和图4中.由表8和图4可见,氯离子排在非金属上方的体系氢键强度强于氯离子排在金属上方的体系.故氯离子排在非金属上方更有利于体系的稳定性.此外,1号体系CuAl-FccCuAlZn 的单个氢键的强度较大,且共形成了12个氢键,故其强度较大;而8号体系CuAl-HcpMg 不管是形成的单个氢键的强度还是氢键数目,都是不利于体系的稳定性的.总体上来说,从1号到8号体系,氢键强度逐渐减弱.4结论本文采用赝势平面波法CASTEP 计算了(M)-IV-LDHs (M=Cu,Zn,Mg,Al)的结构和能量,选取了表7CuAl-IV-LDHs 的Mulliken 电荷布居Table 7Mulliken charge population of CuAl-IV-LDHsNo.12345678System LayerCuAl-FccCuAlZn CuAl-HcpAlCuAl-FccMgAlCu CuAl-FccMgAlZn CuAl-HcpCu CuAl-HcpZnCuAl-FccCuZnMg CuAl-HcpMg Q (H)/e 0.4010.3880.3860.3860.3880.3880.3910.3910.390Q (O)/e -0.952-0.949-0.952-0.953-0.954-0.953-0.952-0.955-0.958Q (Mg)/e 2.1201.9101.9701.9981.9551.9501.9752.0102.055Q (Al)/e 1.5801.5501.5601.5601.5601.5701.5601.5701.565Q (Zn)/e 1.0400.9700.9600.9700.9700.9900.9900.9900.980Q (Cu)/e 0.6700.6150.6200.6400.6500.6300.6500.6400.650Q (Layer)/e-0.5570.5820.6320.6070.620.6870.6980.671Q (Cl)/e --0.590-0.600-0.635-0.620-0.630-0.685-0.690-0.685表8CuAl-IV-LDHs 的氢键结构参数Table 8Hydrogen bond parameters of CuAl-IV-LDHsNo.12345678System CuAl-FccCuAlZn CuAl-HcpAlCuAl-FccMgAlCu CuAl-FccMgAlZn CuAl-HcpCu CuAl-HcpZnCuAl-FccCuZnMg CuAl-HcpMgBond length/nm0.227700.228910.223600.233060.230510.226310.232900.24106Charge/e 0.0650.0600.0700.0580.0570.0650.0610.048Number of hydrogen bonds1212812128811Bond angle/(°)154.951150.791153.026151.452152.698151.330146.558144.939图4CuAl-IV-LDHs 氢键键长和电荷关系图Fig.4Relation of charges and hydrogen bond lengths ofCuAl-IV-LDHsThe number x (y )in the figure means that there are y hydrogen bonds for the system x .refers to anion on the nonmetal atoms andrefers to anion on the metalatoms.63Acta Phys.-Chim.Sin.2012Vol.28能量相对较低的CuAl-IV-LDHs体系,探讨了体系的姜-泰勒效应以及氯离子位置对层板畸变的影响,结论如下:(1)对于CuAl-IV-LDHs体系,姜-泰勒效应不仅存在于d轨道未排满的Cu2+中,在p轨道未排满的Mg2+、Al3+中也可能存在.对于1号到8号体系的平均金属畸变呈逐渐增大趋势,说明层板上的金属八面体畸变越来越严重,不利于层板中金属八面体的稳定性.(2)层板中的铜原子与六个氧原子配位,从1号到8号体系,六个Cu－O键先出现四长键二短键的分化,为拉长的稳定的八面体,再逐渐发生构型转变,最后转变为四短键二长键的压扁的不稳定的八面体形式.(3)通过对CuAl-IV-LDHs体系的氢键、结合能的分析,氯离子排在非金属上方的体系氢键强度强于氯离子排在金属上方的体系,故氯离子排在非金属上方更有利于体系的稳定性.对于1号到8号体系,总结合能绝对值逐渐降低,体系稳定性下降,氯离子在非Mg原子正上方,即CuAl-FccCuAlZn最稳定;氯离子在Mg原子正上方,即CuAl-HcpMg最不稳定.References(1)Cavani,F.;Trifiro,F.;Vaccari,A.Catal.Today1991,11,173.(2)Duan,X.;Zhang,F.Z.Intercalation and Assembly Chemistry ofInorganic Supramolecular Materials;Science Press:Beijing,2009.[段雪,张法智.无机超分子材料的插层组装化学.北京:科学出版社,2009.](3)Yan,H.;Lu,J.;Wei,M.;Ma,J.;Li,H.;He,J.;Evans,D.G.;Duan,X.J.Mol.Struct.-Theochem2008,866,34.(4)Becke,A.D.Chem.Phys.1993,98,5648.(5)Lee,C.;Yang,W.;Parr,R.G.Phys.Rev.B1988,37,785.(6)Yan,H.;Wei,M.;Ma,J.;Li,F.;Evans,D.G.;Duan,X.J.Phys.Chem.A2009,113,6133.(7)Feng,Y.J.;Li,D.Q.;Li,C.X.;Wang,Z.G.Acta ChimicaSinica2003,61,78.[冯拥军,李殿卿,李春喜,王子镐.化学学报,2003,61,78.](8)Wu,J.S.;Xiao,Y.K.;Lin,Y.P.;Liang,H.Q.;Li,C.Y.;He,H.Y.Journal of Synthetic Crystals2010,39,817.[吴健松,肖应凯,林意萍,梁海群,李春银,何海英.人工晶体学报,2010,39,817.](9)Wang,X.P.;Yu,J.J.;Cheng,J.;Hao,P.Z.;Xu,P.Z.Environ.Sci.Technol.2008,42,614.(10)Zhang,J.;Zhao,N.;Wei,W.;Sun,Y.H.The Research on UsingQuaternary Hydrotalcites for the Partial Qxidation of Methane.In The Proceedings of The Seventh National Conference onCatalysis Science and Technology,The Seventh NationalConference on Catalysis Science and Technology,Taiyuan,2009;pp319-321.[张军,赵宁,魏伟,孙予罕.新型四元水滑石在甲烷部分氧化中的应用研究.第七届全国催化剂制备科学与技术研讨会会议论文集,第七届全国催化剂制备科学与技术研讨会,太原,2009:319-321.](11)Heermann,puter Simulation Methods in TheoreticalPhysics;Springer-Verlag Press:Heidelberg,1990;pp387-439.(12)Leach,A.R.Molecular Modelling:Principles and Applications;Addison Wesley Longman Limitted Press:Essex,2001;pp26-454.(13)Ni,Z.M.;Yao,P.;Liu,X.M.;Wang,Q.Q.;Xu,Q.Chem.J.Chin.Univ.2010,31,2438.[倪哲明,姚萍,刘晓明,王巧巧,胥倩.高等学校化学学报,2010,31,2438.](14)Tang,J.Y.;Jiang,H.F.;Deng,G.H.;Zhou,.Chem.2005,25,1503.[唐金玉,江焕峰,邓国华,周磊.有机化学,2005,25,1503.](15)Suresh,I.;Vinay,V.T.J.Mol.Catal.A:Chem.2000,157,275.(16)Matthew,D.J.;Christine,J.C.;Mary,F.M.;Paul,R.R.;David,A.;Joann,W.;David,C.J.Mol.Catal.A:Chem.2010,325,8.(17)Surapas,E.R.Catal.Lett.2011,141,784.(18)Xu,J.H.;Man,Q.S.;Lin,Y.C.;Li,Y.Y.;Feng,Y.S.Chin.J.Org.Chem.2010,30,9.[许建华,蔄秋实,林义成,李源源,冯乙巳.有机化学,2010,30,9.](19)Xu,Q.;Ni,Z.M.;Mao,J.H.J.Mol.Struct:Theochem.2009,915,122.(20)Xu,Q.;Ni,Z.M.;Pan,G.X.;Chen,L.T.;Liu,T.ActaPhys.-Chim.Sin.2008,24,601.[胥倩,倪哲明,潘国祥,陈丽涛,刘婷.物理化学学报,2008,24,601.](21)Segall,M.D.;Linda,P.;Probert,M.;Pickard,C.;Hasnip,P.;Clark,S.;Payne,M.J.Phys.-Condes.Matter2002,14,2717.(22)Ceperley,D.M.;Aider,B.J.Phys.Rev.Lett.1980,45,566.(23)Vanderbilt,D.Phys.Rev.B1990,41,7892.(24)Kresse,G.;Furthmiiller,J.Phys.Rev.B1996,54,11169.(25)Pan,D.K.;Zhao,C.D.;Zheng,Z.X.The Structure of Matter;Higher Education Press:Beijing,1989;pp329-330.[潘道皑,赵成大,郑载兴.物质结构.北京:高等教育出版社,1989:329-330.](26)Mulliken,R.S.J.Chem.Phys.1955,23,1833.(27)Scheiner,S.Hydrogen Bonding;Oxford University Press:NewYork,1997.(28)Jeffrey,G.A.An Introduction to Hydrogen Bond;OxfordUniversity Press:New York,1997.(29)Desiraju,G.;Steiner,T.The Weak Hydrogen Bond;OxfordUniversity Press:New York,1999.64。

Microstructure evolution and mechanical properties of1 000 MPa cold rolled dual-phase steelZHAO Zheng-zhi(赵征志), JIN Guang-can(金光灿), NIU Feng(牛枫), TANG Di(唐荻), ZHAO Ai-min(赵爱民) Engineering Research Institute, University of Science and Technology Beijing, Beijing 100083, ChinaReceived 10 August 2009; accepted 15 September 2009Abstract: The microstructure evolution of 1 000 MPa cold rolled dual-phase (DP) steel at the initial heating stages of the continuous annealing process was analyzed. The effects of different overaging temperatures on the microstructures and mechanical properties of 1 000 MPa cold rolled DP steel were investigated using a Gleeble−3500 thermal/mechanical simulator. The experimental results show that ferrite recovery and recrystallization, pearlite dissolution and austenite nucleation and growth take place in the annealing process of ultra-high strength cold rolled DP steel. When being annealed at 800 ℃ for 80 s, the tensile strength and total elongation of DP steel can reach 1 150 MPa and 13%, respectively. The microstructure of DP steel mainly consists of a mixture of ferrite and martensite. The steel exhibits low yield strength and continuous yielding which is commonly attributed to mobile dislocations introduced during cooling process from the intercritical annealing temperature.Key words: cold rolled dual-phase steel; microstructure evolution; recrystallization; mechanical property; overaging temperature1 IntroductionAdvanced high-strength steels (AHSS) have been used in the automotive industry as a solution for the weight reduction, safety performance improvement and cost saving. Among them, the dual-phase (DP) steels, whose microstructure mainly consists of ferrite and martensite, are an excellent choice for applications where low yield strength, high tensile strength, continuous yielding, and good uniform elongation are required [1−4].The continuous annealing process to produce cold rolled DP steels typically has the following stages: heating to the intercritical temperature region, soaking in order to allow the nucleation and growth of austenite, slow cooling to the quench temperature, rapid cooling to transform the austenite into martensite, overaging, and air cooling. The amount and morphology of the constituents formed depend on such annealing parameters. The effects of the retained austenite, ferrite, and martensite morphologies on the mechanical behavior of DP steels have been intensively investigated[5−9]. As we all known, overaging treatment is an important process during the production of dual-phase steel. It can reduce the hardness of martensite and improve the comprehensive mechanical properties of DP steel [10−14].The purpose of the present research was to study the microstructure evolution of cold rolled DP steel at the initial heating stages of the continuous annealing process using a Gleeble simulator. At the same time, the effects of overaging temperature on the mechanical properties of DP steel were also studied. The microstructures of specimens simulated on a Gleeble simulator, were analyzed using scanning electron microscopy (SEM) and transmission electron microscopy (TEM).2 ExperimentalThe chemical compositions of the experimental steel (mass fraction, %) were: 0.14−0.17C, 0.40−0.60Si, 1.70−1.90Mn, 0.02−0.04Nb, 0.40−0.60Cr, ≤0.010P, ≤0.010S, 0.02−0.06Al and balance Fe. Firstly, experimental steels were smelted in a 50 kg vacuum induction furnace. After smelting, experimental steels were forged into 35 mm×100 mm×100 mm cubic samples. The forged slabs were reheated to 1 200 ℃and soaked for 1 h. The hot rolled thickness was 3.5 mm after 6 passes rolling. The finish rolling temperature was about 880 ℃. The coiling temperature was 620 ℃. After being pickled in hydrochloric acid, the hot rolledFoundation item: Project(2006BAE03A06) supported by the National Key Technology R&D Program during the 11th Five-Year Plan Period Corresponding author: ZHAO Zheng-zhi; Tel: +86-10-62332617; E-mail: zhaozhzhi@ZHAO Zheng-zhi, et al/Trans. Nonferrous Met. Soc. China 19(2009) s563−s568 s564bands were cold rolled to the final thickness of 1.0 mm, and the reduction was about 70%. Finally, the cold rolled sheets were cut into the samples for the simulation of continuous annealing experiment.The microstructure evolution at the initial steps of the continuous annealing process was studied using a Gleeble 1500 simulator. The steel was heated at 10 ℃/sto the different heating temperatures (550, 630, 670, 710, 730, 750 and 780 ℃) and held for 20 s followed by water-quenching. The effects of different overaging temperatures on the microstructures and mechanical properties of DP steel were investigated using a Gleeble 3500 simulator. The processing schedules and parameters used are shown in Fig.1. The soaking temperature of intercritical region was set at 800 ℃, soaking time is 80 s; after a slow cooling, the samples were rapidly cooled to 240, 280, 320 and 360 ℃, respectively and soaked for 300 s; at last, the samples were air cooled to the room temperature.Fig.1 Continuous annealing process of DP steelAfter heat treatment, the steel sheet would be cut into standard tensile specimens (length 200 mm, gauge length 50 mm). The tensile test was performed with CMT4105-type tensile test machine to test mechanical properties. The longitudinal cold rolling plane sections of samples after annealing were prepared and etched with 4% natal. The microstructure was analyzed by scanning electron microscopy (SEM). Some samples were analyzed using transmission electron microscopy (TEM).3 Results and discussion3.1 Mechanical properties and microstructures ofsamples after hot-rolling and continuousannealingTable 1 shows the tensile test data for the two samples after hot-rolling and continuous annealing in terms of yield strength, ultimate tensile strength and total elongation. When the annealing temperature is 800 ℃and soaking time is 60 s, the tensile strength reaches 1 110 MPa and the total elongation reaches 12%. Compared with the hot-rolled samples, the yield strength and total elongation of sample after annealing are similar, but the tensile strength increases by about 450 MPa. The yield ratio decreases obviously. The engineering uniaxial tensile stress—strain curve of the sample after continuous annealing is characterized by very uniform plastic flow until necking. There is no physical yield point and yield point extension, that is, the steel exhibits continuous yielding which is commonly attributed to mobile dislocations introduced during cooling from the intercritical annealing temperature. Many dislocation sources come into action at low strain and plastic flow begins simultaneously through the specimen, thereby suppressing discontinuous yielding[15].Table 1 Mechanical properties of samples after hot rolling and annealingConditionYieldstrength/MPaTensilestrength/MPaYieldratio*Totalelongation/% Hot rolling555 665 0.83 16 Annealing540 1110 0.49 12* Yield ratio is defined as the ratio of yield strength to tensile strength.The microstructures of the hot-rolled and cold-rolled samples are shown in Fig.2. It can be observed that hot rolled steel features a band microstructure, i.e. pearlite band in a ferrite grain matrix. The ferrite grain size is measured to be 5.0−9.0 µm. After cold rolling, the microstructure consists of elongated grains of ferrite and deformed colonies of pearlite (Fig.2(b)). After cold-rolling, there is an increase in the stored energy of the steel due to the high dislocation density and this provides the driving pressure for the ferrite recrystallization during annealing process. The total ferrite grain boundary area increases and the cementite laminar structure in pearlite is broken down. The latter has been shown to promote spheroidization of cementite during subsequent annealing process.The SEM micrograph of the sample after annealing is given in Fig.3(a). The microstructure of DP steel consists of a mixture of ferrite, martensite, martensite/austenite constituent. There is also some bainite in the microstructure. The martensite islands are homogeneously distributed in ferrite matrix. The DP steel has finer grain size and the size of ferrite grain and martensite island are about 1.0−2.0 µm. Some martensite islands have a bright white circle around the edge, and the center of martensite is of irregular black structure.ZHAO Zheng-zhi, et al/Trans. Nonferrous Met. Soc. China 19(2009) s563−s568 s565Fig.2 Microstructures of steel after hot rolling (a) and cold rolling (b)Fig.3 SEM images (a) and TEM micrograph (b) of steel after continuous annealingThe main reason is the manganese partitioning will occur during the continuous annealing process. During the heating process, a high-Mn side lap forms around austenite, which makes the hardenability of austenite island edge higher than that of the center. So, it makes high-Mn side lap form around martensite in the cooling process. The volume fraction of martensite is about 40%, which is the main reason for DP steel with a higher strength. After the continuous annealing process, band structure is significantly improved, which plays an important role in improving the performance of DP steel.The fine structures of martensite and ferrite are shown in Fig.3(b) by the TEM observation. The lath martensite is fine, and is relatively clean; at the same time, a very high density of dislocations can be observed in the ferrite grain adjacent to martensite. These dislocations are generated in order to accommodate transformation induced strain built between martensite transformed by quenching and retained ferrite. In addition, they are known to be mobile and play an important role on rapid, extensive strain hardening of DP steel from the onset of its plastic deformation.3.2 Microstructure evolution at initial steps ofcontinuous annealing processThe microstructure evolution at the initial stages of the continuous annealing process is very important for producing the ultra-high strength DP steel. During the annealing process of high strength DP steel, ferrite recovery and recrystallization, pearlite dissolution and austenite nucleation and growth will occur. When the sample is heated to 550 , the℃microstructure has no visible change as compared with the cold rolled sample. The ferrite grain is stretched along the rolling direction significantly; lamellar pearlite is stretched along the rolling direction too. At the same time, there are some carbide particles in the ferrite matrix, as shown in Fig.4(a). At this temperature, the recrystallization nucleus was not found in the structure. So, at this stage the sample is still at the recovery stage. When the heating temperature is 630 , the℃recrystallization nucleus begins to appear in the microstructure. The nucleus of crystal appears mainly nearby the large deformation ferrite (Fig.4(b)). The recrystallization nucleus is fine and equiaxed. Large deformation storage power is present in the large deformation region. So, recrystallization nucleus forms in this region firstly. With the heating temperature increasing, the recrystallization nucleus begins to grow. Therefore, the size of recrystallization is uneven at this stage, as shown in Fig.4(c). When the heating temperature is 670 ℃, the deformation structure still exists in the microstructure. With the temperature increasing, the deformed ferrite grains are replaced by recrystallization ferrite grains. When the heating temperature is 710 , the d℃eformation structure has already vanished, which is replaced by theZHAO Zheng-zhi, et al/Trans. Nonferrous Met. Soc. China 19(2009) s563−s568 s566equiaxed recrystallization grain. So, the process of recrystallization completes basically. In the ferrite recrystallization process, the pearlite transforms to granular from lamellar gradually.When the heating temperature is 730 ,℃it begins to enter the two-phase region; and the ferrite and spheroidised carbides begin to transform to austenite. A small amount of austenite nucleates in the original pearlite region, as shown in Fig.4(e). Austenite nucleates mainly in the ferrite and pearlite grain boundary; and a part of austenite also nucleates in the carbide particles of ferrite. After austenite nucleation, it begins to grow rapidly. At this stage, the pearlite dissolves rapidly. When the temperature reaches 750 , the austenite℃transformation occurs obviously. The bright white particle which distributes in the ferrite matrix is the martensite island. The martensite transforms from austenite during the rapid cooling process. At the same time, a small amount of martensite particles can also be observed in ferrite; and there are still some non-dissolved carbide particles in the ferrite matrix. The initial austenite growing-up is mainly controlled by the carbon Fig.4Microstructure evolutions duringcontinuous heating process: (a) 550 ℃; (b)630 ℃; (c) 670 ℃; (d) 710 ℃; (e) 730 ℃; (f)750 ℃; (g) 780 ℃ZHAO Zheng-zhi, et al/Trans. Nonferrous Met. Soc. China 19(2009) s563−s568 s567diffusion in the austenite, and the diffusion path is along the pearlite/austenite interface. When the annealing temperature is 780 , the austenite volume increase℃s, and the number of carbide particles is reduced gradually. There is only a very small amount of carbide particles distributing in ferrite matrix.3.3 Effect of overaging temperature onmicrostructure and mechanical properties ofDP steelThe overaging is a temper treatment to harden martensite in the dual-phase steel, reduce the hardness of martensite and improve the comprehensive mechanical properties[16]. Fig.5 shows the effect of overaging temperature on the mechanical properties of dual-phase steel. All the samples are intercritically annealed at 800℃ with different overaging temperatures. As can be seen from Fig.5, the highest tensile strength is achieved in the sample overaged at 280 ℃. The yield strength is 560 MPa, the tensile strength is 1 150 MPa, and the total elongation reaches 13%. The good combination of high strength and toughness properties is obtained. And then, with the increase of overaging temperature, the yield strength and tensile strength of samples decrease, while the total elongation increases. When the overaging temperature reaches 360 ℃, the tensile strength decreases, the yield strength does not change significantly. The mechanical properties of sample cannot meet the necessary requirements of CR980DP. At the same time, the stress—strain curve of the steel shows discontinuous yielding behaviour and develops yield plateaus.Fig.6 shows the SEM microstructures with different overaging temperatures. It can be seen that the microstructure mainly consists of dark grey ferrite grains and white martensite. When the overaging temperature is 360 ℃, the martensite boundary is fuzzier than that of sample overaged at 320 ℃, and there are more carbides, which is due to the effects of tempering on the martensite, such as the volume contraction of martensite during the tempering, the changes of the martensite strength and additional carbon clustering or precipitation near the ferrite and martensite interfaces.Fig.5 Effects of different overaging temperatures on mechanical propertiesFig.6 SEM images of microstructures of DP steel overaged at different temperatures: (a) 240 ℃; (b) 280 ℃; (c) 320 ℃; (d) 360 ℃ZHAO Zheng-zhi, et al/Trans. Nonferrous Met. Soc. China 19(2009) s563−s568 s5684 Conclusions1) When the DP steel is annealed at 800 ℃ for 80 s and overaged at 280 ℃, the tensile strength and total elongation of ultra-high strength dual-phase steel can reach 1 150 MPa and 13%, respectively.2) The microstructure of DP steel consists of a mixture of ferrite, martensite, martensite/austenite constituent. There are also some bainites in the microstructure. The martensite islands are homogeneously distributed in ferrite matrix. The ferrite and martensite island grain size are about 1.0−2.0 µm. When the overaging temperature reaches 360 ℃, the tensile strength decreases, the yield strength does not change significantly. The mechanical properties of sample cannot meet the necessary requirements of CR980DP. At the same time, the steel shows discontinuous yielding behaviour and develops yield plateaus.References[1]KANG Yong-lin. Quality control and formability of the mordernMotor plate [M]. Beijing: Metallurgical Industry Press, 1999.[2]LIU Peng, JIN Xian-zhe. The development and research ofautomobile steel plate [J]. Shanxi Metallurgy, 1997(2): 32−33.[3]MA Ming-tu, WU Bao-rong. Dual-phase steel-the physical andmechanical metallurgy [M]. Beijing: Metallurgical Industry Press,1988.[4]LLEWELLYN D T，HILLS D J. Dual phase steels [J]. Ironmakingand Steelmaking, 1996(6): 471−478.[5]SARKAR P P. Microstructural influence on the electrochemicalcorrosion behaviour of dual phase steels in 3.5% NaCl solution [J].Materials Letters, 2005(59): 2488−2491. [6]ROCHA R O, MELO T M F, PERELOMA E V, SANTOS D B.Microstructural evolution at the initial stages of continuous annealingof cold rolled dual-phase steel [J]. Materials Science and EngineeringA, 2005, 391: 296−304.[7]MA C, CHEN D L, BHOLE S D, BOUDREAU G, LEE A, BIRO E.Microstructure and fracture characteristics of spot-welded DP600 steel [J]. Materials Science and Engineering A, 2008, 485: 334−346.[8]SUN Shou-jin, Martin P. Manganese partitioning in dual-phase steelduring annealing [J]. Materials Science and Engineering A, 2000, 276: 167−174.[9]ZHU Xiao-dong, WANG Li. Effect of the continuous annealingparameters on the mechanical properties of cold rolled Si-Mn dualphase steel [C]//CSM 2003 Annual Meeting Proceedings, 2003: 684−688.[10]MOHAMMAD R A, EKRAMI A. Effect of ferrite volume fractionon work hardening behavior of high bainite dual phase (DP) steels [J].Materials Science and Engineering A, 2008, 477: 306−310.[11]HA VV A K Z, CEYLAN K, HUSEYIN A. Investigation of dual phasetransformation of commercial low alloy steels: Effect of holding timeat low inter-critical annealing temperatures [J]. Materials Letters, 2008, 62: 2651−2653.[12]DOU Ting-ting, KANG Yong-lin, YU Hao, KUANG Shuang, LIURen-dong, YAN Ling. Microstructural evolution of cold rolled dualphase steel during initial stages of continuous annealing [J]. Heat Treatment of Metal, 2008, 33(3): 31−35.[13]CHEN Hui-feng, ZHANG Qing-fen, AN Jia-shen. Recrystallizationcharacteristic of IF steel during rapid heating [J]. Journal of East China University of Metallurgy, 1999, 16(1): 21−23.[14]YANG D Z, BROWNEL E L, MATLOCK D K, et al. Ferriterecrystallization and austenite formation in cold rolled intercriticallyannealed steel [J]. Metallurgical Transactions A, 1985, 16A: 1385−391.[15]SULEYMAN G. Static strain ageing behaviour of dual phase steels[J]. Materials Science and Engineering A, 2008, 486: 63−71.[16]KUANG Shuang, KANG Yong-lin, YU Hao, LIU Ren-dong, YANLing. Experimental study on microstructure evolution in continuousannealing of cold-rolled dual phase steels [J]. Iron and Steel, 2007,42(11): 65−73.(Edited by CHEN Ai-hua)。

0前言航空航天装备向着高可靠性、更长寿命的发展趋势对材料提出了更高的性能要求。

铝合金具有高比强度、优良的耐蚀性及热塑性等优点，已经广泛地应用于航空航天装备结构件中。

2014铝合金是一种典型的航空航天装备用材料，具有良好的热成型性能、焊接性能，常常以锻件、挤压件等多种形式应用于承力结构件[1]。

目前，2014铝合金的大多研究集中于热处理制度和焊接工艺的研究[2-5]，而合金性能的优劣主要源于成分和组织的优化。

2014合金成分范围较宽，成分差异和不同的加工工艺及热处理方式使其表现出不同的综合性能。

本文通过设计三种不同成分的2014铝合金，研究了Cu、Mg、Si 三种合金元素对2014合金组织和性能的影响，以期为该合金性能多样化及综合性能的提高提供有益指导。

1实验材料与方法在2014合金AMS4133E 标准成分范围内，按照Cu、Mg、Si 在合金中的不同作用，设计了三种成分的2014铝合金（标记为A、B 和C 合金），其化学成分见表1。

A 合金成分为标准成分范围的中间值；B 合金同时增加Cu、Mg 含量；C 合金同时减少Cu、Mg、Si 含量。

通过对比三种合金，研究主合金元素对2014合金组织与性能的影响。

经半连续铸造获得直径为ϕ200mm 的合金铸锭，铸锭经均匀化热处理后热挤压成截面为125mm×25mm 的型材，挤压比为10，空冷至室温。

采用502℃×5h 固溶处理和177℃×8h 时效处理，获得T6态型材。

Cu 、Mg 、Si 元素含量对2014铝合金组织与性能的影响林茂1，曹海龙2，田宇兴2，吴浩2（1.西北铝业有限责任公司，定西748111；2.中铝材料应用研究院有限公司，北京102209）摘要：通过设计三种合金成分，研究了Cu、Mg、Si 元素含量对2014铝合金组织与性能的影响。

研究结果表明，随着Cu、Mg、Si 含量的增大，铸态组织的共晶相增多；当Cu 含量低于4.2%时，502℃×30h 均匀化退火可以使Al 2Cu 完全回溶。

捅要摘要采用微合金化的方法来提高铝合金的综合性能是提高铝材质量的基础研究中的重要课题。

作者所在的课题组对稀±元素Ｅｒ在铝合金中的作用进行了系统的预研工作，发现Ｅｒ在铝合金中具有积极作用。

为进一步研究稀土元素Ｅｒ在铝合金中的作用及其作用机理，本论文采用铸锭冶金法制备了六种成分不同的ＡＪ．Ｚｎ．Ｍｇ（．Ｅｒ）合金，通过硬度测试、拉伸力学性能测试、金相显微组织观察、ｘ射线衍射分析、扫描电镜观察、透射电镜观察及能谱分析等手段研究了稀土元素Ｅｒ对Ａｌ—Ｚｎ—Ｍｇ合金在不同状态下的力学性能、显微组织，热稳定性能以及时效特性的影响；初步探讨了稀土元素Ｅｒ在Ａ１．Ｚｎ．Ｍｇ合金中的存在形式，对合金显微组织的细化机理，对合金力学性能的强化机理．以及含Ｅｒ的Ａ１．Ｚｎ—Ｍｇ合金的再结晶形核机制。

结果表明：稀土元素Ｅｒ可以明显地提高ＡＩ．Ｚｎ．Ｍｇ合金的力学性能，其中添加０．１％Ｅｒ的合金，其力学性能上升的幅度最大，约为１５％。

随着Ｅｒ含量的增加，合金力学性能继续提高，但当Ｅｒ含量超过Ｏ．４％后，硬度和强度的增加不再明显。

Ｅｒ还可以抑制合金的再结晶过程，使合金的再结晶温度提高约５０。

Ｃ。

添加Ｅｒ后，Ａ１．Ｚｎ．Ｍｇ合金达到硬度峰值所用的时间缩短了，时效硬度峰值也明显提高。

关键词稀土元素Ｅｒ：ＡＩ－Ｚｎ—Ｍｇ合金：Ａ１３Ｅｒ；时效北京工业大学工学顾士学位论文ＡＢＳＴＲＡＣＴＨｏｗｔＯｉｍｐｒｏｖｅｔｈｅｓｙｎｔｈｅｓｉｓｐｒｏｐｅｒｔｉｅｓｏｆｔｈｅａｌｕｍｉｎｕｍａｌｌｏｙｓｂｙｍｅａｎｏｆｍｉｃｒｏａｌｌｏｙｉｎｇｉｓｏｎｅｏｆｔｈｅｉｍｐｏｒｔａｎｔｂａｓｉｃｒｅｓｅａｒｃｈｉｔｅｍｓ．Ｉｎｏｕｒｅａｒｌｉｅｒｓｔｕｄｉｅｓｃｏｎｓｉｄｅｒａｂｌｅｅｆｆｏｒｔｓｈａｖｅｂｅｅｎｓｐｅｎｔｏｎｔｈｅｅｆｆｅｃｔｓｏｆｔｈｅｒａｒｅ—ｅａｃｈｅｌｅｍｅｎｔＥｒｂｉｕｍｏｎｔｈｅａｌｕｍｉｎｕｍａｎｄｉｔｓａｌｌｏｙｓ．ＴｈｅｒｅｓｕｌｔｓｓｈｏｗｔｈａｔＥｒｂｉｕｍｉｓｇｏｏｄｆｏｒｔｈｅｐｒｏｐｅｒｔｙｉｍｐｒｏｖｅｍｅｎｔｏｆｔｈｅａｌｕｍｉｎｕｍａｎｄｉｔｓａｌｌｏｙｓ．ＦｏｒｆｕｒｔｈｅｒｓｔｕｄｙｏｎｔｈｅｅｆｆｅｃｔｓｏｆＥｒｏｎｔｈｅａｌｕｍｉｎｕｍａｎｄｉｔｓａｌｌｏｙｓ，ｓｉｘｄｉｆｆｅｒｅｎｔｃｏｍｐｏｓｉｔｉｏｎｓｏｆＡＩ—Ｚｎ—ＭｇａｌｌｏｙｓｃｏｎｔａｉｎｉｎｇＥｒｗｅｒｅｐｒｅｐａｒｅｄｂｙｉｎｇｏｔｍｅｔａｌｌｕｒｇｙ．ＴｈｅｉｎｆｌｕｅｎｃｅｏｆｄｉｆｆｅｒｅｎｔｃｏｎｔｅｎｔｓｏｆＥｒａｄｄｉｔｉｏｎｏｎｔｈｅｍｅｃｈａｎｉｃａｌｐｒｏｐｅｒｔｉｅｓｕｎｄｅｒｄｉｆｆｅｒｅｎｔｓｔａｔｅｓｍｉｃｒｏｓｔｒｕｃｔｕｒｅ，ｈｅａｔｒｅｓｉｓｔａｎｃｅａｎｄｔｈｅａｇｉｎｇｐｒｏｐｅｒｔｉｅｓｏｆＡＩ—Ｚｎ—Ｍｇａｌｌｏｙｓｈａｓｂｅｅｎｓｔｕｄｉｅｄｂｙｈａｒｄｎｅｓｓｔｅｓｔ，ｔｅｎｓｉｌｅｐｒｏｐｅｒｔｉｅｓｍｅａｓｕｒｅｍｅｎｔ，Ｘ—ｒａｙｄｉｆｆｒａｃｔｉｏｎ（ＸＲＤ）ａｎａｌｙｓｉｓ，ｏｐｔｉｃａｌｍｉｃｒｏｓｃｏｐｙ，ｓｃａｎｎｉｎｇｅｌｅｃｔｒｏｎｍｉｃｒｏｓｃｏｐｙ（ＳＥＭ），ｔｒａｎｓｍｉｓｓｉｏｎｅｌｅｃｔｒｏｎｍｉｃｒｏｓｃｏｐｙ（ＴＥＭ）ａｎｄｅｎｅｒｇｙｄｉｓｐｅｒｓｉｖｅｘ—ｒａｙｓｐｅｃｔｒｏｓｃｏｐｙ（ＥＤｘｓ）ａｎａｌｙｓｉｓＴｈｅｅｘｐｅｒｉｍｅｎｔａｌｒｅｓｕｌｔｓｓｈｏｗｔｈａｔ：ｔｈｅｍｅｃｈａｎｉｃａｌｐｒｏｐｅｒｔｉｅｓｏｆｔｈｅＡＩ－Ｚｎ－ＭｇａｌｌｏｙｃａｎｂｅｏｂｖｉｏｕｓｌｙｉｍｐｒｏｖｅｄｂｙＥｒａｄｄｉｔｉｏｎ．ＷｈｅｎｔｈｅＥｒｃｏｎｔｅｎｔｉｓ０．１ｗｔ％，ｔｈｅｓｔｒｅｎｇｔｈｉｎｃｒｅｍｅｎｔｉｓｔｈｅｌａｒｇｅｓｔ，ｔｈａｔｉｓａｂｏｕｔ１５％．ＴｈｅｎｔｈｅｉｎｃｒｅｍｅｎｔｒｅｄｕｃｅｓｗｈｅｎｔｈｅＥｒｃｏｎｔｅｎｔｉｓｏｖｅｒＯ．４ｗｔ．％．ｔｈｏｕｇｈｔｈｅｓｔｒｅｎｇｔｈｃｏｎｔｉｎｕｅｓｔｏｉｎｃｒｅａｓｅｃｏｍｐａｒｅｄｔｏｔｈａｔｏｆｔｈｅＡＩ—Ｚｎ—ＭｇａｌｌｏｙｆｒｅｅｏｆＥｒ．ＷｉｔｈＥｒａｄｄｉｔｉｏｎ，ｔｈｅｒｅｃｒｙｓｔａｌｌｉｚａｔｉｏｎｏｆｔｈｅＡＩ—Ｚｎ—Ｍｇａｌｌｏｙｉｓｓｕｐｐｒｅｓｓｅｄ，ｔｈａｔｉｓ，ｔｈｅｒｅｃｒｙｓｔａｌｌｉｚａｔｉｏｎｔｅｍｐｅｒａｔｕｒｅｏｆＥｒ－ｄｏｐｅｄＡ１－Ｚｎ—Ｍｇａｌｌｏｙｉｎｃｒｅａｓｅｓａｂｏｕｔ５０。