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Electron Channeling Contrast Imaging (ECCI) - 牛津仪器

EBSD.cn

电子背散射衍射

EBSD原理详解

EBSD是什么?EBSD初学者EBSD系统的基本组成EBSD花样的形成解释衍射花样EBSD系统标定自动标定的基础晶体学的基本概念

EBSD技术

EBSD探测器标定技术高精度EBSD集成EDS3D EBSD透射菊池衍射（TKD)原位EBSD电子通道衬度成像显示EBSD数据有用的EBSD参考文献

EBSD应用

EBSD提供什么信息?增材制造汽车和航空航天生物成像和生命科学能源生产和存储物证鉴定与环境地质学、岩石学和采矿金属、合金、复合材料和陶瓷半导体、微电子和存储器

提示和技巧

数据采集

EBSD几何设置优化SEM设置EBSD花样采集

样品制备

引言样品切割样品镶嵌样品研磨抛光电解抛光及腐蚀离子束抛光技术预防荷电效应/导电镀膜样品保存各种材料的制备方法

技术

牛津仪器EBSD产品

CMOS探测器系列AZtecHKL采集软件AZtecCrystal处理软件

牛津仪器技术详解

光纤板光学系统和灵敏度直接电子探测与间接电子探测EBSD探测器Tru-I标定技术高精度标定模式高级相区分修正伪对称母相晶粒重构位错分析花样匹配

培训

入门教程视频"操作方法"视频AZtecCrystal 培训

电子背散射衍射 (EBSD)

EBSD技术

电子通道衬度成像

ECCI的原理

图册

电子通道衬度成像

电子通道衬度成像（ECCI），是在扫描电子显微镜（SEM）下，通过电子通道效应进行成像的技术。ECCI图像的衬度由晶粒的晶体学取向决定。ECCI图像可用于EBSD采集前，快速预览样品的显微组织。当晶格与入射电子相对取向合适时，ECCI图像可以用来成像并表征样品中的单根位错，提供有关滑移系和变形机理的信息。虽然“ECCI”是一个较新的术语，但通过背散射电子（BSE）进行取向衬度成像的技术，已有几十年的发展历史。在上世纪七八十年代，电子通道花样（ECP）已经用于在SEM中分析材料和地质样品的晶体学，其方法和采集“通道显微图像”一样，例如参考文献Joy et al. (1982), Journal of Applied Physics 53, R81。随着EBSD技术在90年代逐步成熟，研究人员开始利用安装在受EBSD分析要求，高度倾转样品下方的BSE探测器。这些向前散射的探测器（FSD，前散射探测器），基于晶面通道效应的电子，同样生成了取向衬度成像。ECCI技术的应用广泛，如今多利用固定于极靴下的BSE探测器，对相对低倾转的样品成像。以下标签页介绍了ECCI关键原理及一些应用案例。

ECCI的原理图册

ECCI的主要原理在于电子穿过晶态材料晶面的通道效应。随着入射电子束和晶粒之间的取向关系改变（例如，不同取向的晶粒），BSE信号强度会改变，因此在灰度图像上，不同的晶粒产生明暗对比度，如以下展示的。

ECCI技术已经普遍使用多年，并形成了常规取向衬度成像的基础。一般使用标准的极靴安装的BSE探测器，或安装在EBSD探测器荧光屏下方的前散射探测器。

然而，如果样品的取向对特定某个晶粒来讲，位于所谓的“双束”条件下，那么晶格取向任何微小的改变（例如受到单根位错存在的影响）将会引起BSE强度显著变化。

所以，位错将会相对所在的晶粒，呈现明显的衬度变化。

变形并部分再结晶的Ni基高温合金的ECCI图像，在高样品倾转角下，通过FSD探测器采集，视场宽度约300 mm。

UO2中位错网的ECCI图像，结果和解释发表在Mansour et al. (2018), Ceramics International, 45 (15), pp.18666-18671。

这种方法有时候被称作“受控ECCI”（简称c-ECCI）。首先利用EBSD或ECP获取特定晶粒的取向，然后根据理论计算理想的样品取向，使该晶粒满足双束条件。

此外，已知晶格取向，也可以用来确定成像的位错类型。c-ECCI的成像条件苛刻，需要高精度5轴样品台，通常在SEM中实际操作时，还需要专门的电动子样品台。

ECCI是一项强大的技术。和同等的透射电子显微镜（TEM）技术相比，它优势明显，可以直接观察块状样品的抛光面，并且样品表面的观察区域更大。然而，利用ECCI表征精细的位错比较耗时，而且和EBSD分析一样，它对样品表面的质量要求极高。即便如此，ECCI技术也有不适合的材料，特别是易被氧化的材料，如镁合金、铝合金等。

含金属间化合物的轧制双相不锈钢的低倾斜ECCI

使用FSD探测器采集的双相钢的彩色ECCI

GaN薄膜中的线位错和原子台阶。

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ECCI 电子通道衬度成像简介-CSDN博客

ECCI 电子通道衬度成像简介

最新推荐文章于 2022-08-16 22:03:01 发布

余京泰

最新推荐文章于 2022-08-16 22:03:01 发布

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扫描电镜

本文链接：https://blog.csdn.net/tju_yjt/article/details/112731588

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扫描电镜

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1 篇文章

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文章目录

ECCI简介ECCI的应用优势ECCI的应用现状ECCI文献分析国内研究单位

ECCI简介

ECCI，全写为electron channeling contrast imaging，即电子通道衬度成像，目前有部分研究人员利用该技术对块状样品表面的位错进行表征。这里是对检索文献过程中的收获做一些记录，以便督促自己学习，也方便后期整理。

ECCI的应用优势

ECCI技术是在扫描电镜下的表征技术，而常规位错表征常使用的为透射电镜。在金属塑性变形和断裂的研究领域，与常规透射电镜相比，ECCI显然具有一下对比优势：

相比于透射样品制样过程，ECCI不需要将样品减薄至电子透明，其制样过程为非破坏的方式，即利用ECCI表征位错后，样品还可以根据需求继续试验；由于其不需要电子透明，ECCI技术可表征区域的面积显著高于TEM表征区域，即可以提供相对整体和具有一定统计意义的位错表征结果；

由于其具有以上的优势，可以预见的，ECCI技术将在以下方面有较大的施展空间：

原位试验，包括单轴变形，多轴变形，循环变形等试验，ECCI提供了原位表征能力，即可以关注同一区域在不同变形阶段的位错表现；位错统计，ECCI提供了较大的表征面积，可以支持研究者对一定范围内的位错进行定量或者半定量的统计结果；特定位置位错分析，由于其不需要减薄的特点，可以很方便的实现对裂尖、特定取向晶粒、应力集中位置等特殊位置的位错进行表征，而避免了FIB提片的昂贵费用；环境试验，ECCI技术可以提供较好的腐蚀环境下的变形和失效样品位错的表征。

ECCI的应用现状

在wed of science中检索，可以发现ECCI作为关键词的论文中有四篇高被引论文，按年份排序如下：

Gutierrez-Urrutia I, Zaefferer S, Raabe D. The effect of grain size and grain orientation on deformation twinning in a Fe–22 wt.% Mn–0.6 wt.% C TWIP steel[J]. Materials Science and Engineering: A, 2010, 527(15): 3552-3560.被引468次(谷歌学术)Gutierrez-Urrutia I, Raabe D. Dislocation and twin substructure evolution during strain hardening of an Fe–22 wt.% Mn–0.6 wt.% C TWIP steel observed by electron channeling contrast imaging[J]. Acta Materialia, 2011, 59(16): 6449-6462.被引516次(谷歌学术)Yao M J, Pradeep K G, Tasan C C, et al. A novel, single phase, non-equiatomic FeMnNiCoCr high-entropy alloy with exceptional phase stability and tensile ductility[J]. Scripta Materialia, 2014, 72: 5-8.被引334次(谷歌学术)Deng Y, Tasan C C, Pradeep K G, et al. Design of a twinning-induced plasticity high entropy alloy[J]. Acta Materialia, 2015, 94: 124-133.被引335次(谷歌学术)

以上四篇文献均为马普所的工作，可以看出，该技术的主要应用单位为马普所。马普所还是强呀，以下我们欣赏一下这四篇文献中ECCI的表征图片：

文献1：文献2：文献3：文献4：从中我们可以发现，马普所利用ECCI技术对变形机制进行了非常细致的分析，值得我们学习。此外，这些ECCI表征位错的图像非常清晰，质量极高，甚至使用了双束条件对位错进行分析，可以说是非常恐怖了，我辈要奋起直追呀。

ECCI文献分析

通过web of science的文献分析功能，整理了以ECCI为主题，同时对研究方向进行精炼：

TS= ECCI 精炼依据: 研究方向: ( MATERIALS SCIENCE OR PHYSICS OR CRYSTALLOGRAPHY OR METALLURGY METALLURGICAL ENGINEERING OR MICROSCOPY )

可以检索到251篇文章，对其进行分析如下：出版年：从出版年的分析中，我们可以发现，该技术发展历史并不长，但是发展速度非常迅速，2016年之后，每年都保持在20篇以上的文献发表，2018年达到最高值29篇，近两年分别为2019年22篇，2020年26篇。

出版刊物：从出版刊物中可以发现，最高发文量的三个期刊分别为MATERIALS SCIENCE AND ENGINEERING A ；ACTA MATERIALIA；SCRIPTA MATERIALIA；符合该技术主要应用于金属材料变形的特点，同时研究水平也较高；此外值得注意的是，该技术在MATERIALS CHARACTERIZATION；ULTRAMICROSCOPY；两个期刊中也有较多发表，说明该技术具有一定难度，引起了大家在技术层面的讨论。

研究机构：从研究机构上看，最主要的研究机构还是马普所：MAX PLANCK SOCIETY 和 MAX PLANCK INST EISENFORSCH GMBH ，其次包括：美国密西根大学 MICHIGAN STATE UNIVERSITY 和 MICHIGAN STATE UNIV 、法国国家科学研究中心 CENTRE NATIONAL DE LA RECHERCHE SCIENTIFIQUE CNRS、德国弗莱贝格工业大学 TECHNICAL UNIVERSITY FREIBERG 等。主要分布在德国、美国和法国等高水平研究单位，国内研究机构较少。

国内研究单位

以ECCI作为主题，可以检索到国内机构的12篇文章，按发表时间顺序包括中科院金属所、重庆大学、北京科技大学、大连理工大学、中国科学院上海应用物理研究所钍熔盐反应堆中心、西北工业大学、中南大学等7家单位。

中科院金属所：

Li Y, Li S X, Li G Y. Deformation bands and dislocation structures of [1 5 5] coplanar double-slip-oriented copper single crystal under cyclic deformation[J]. Materials Science and Engineering: A, 2004, 372(1-2): 75-80.

Sun S J, Tian Y Z, Lin H R, et al. Transition of twinning behavior in CoCrFeMnNi high entropy alloy with grain refinement[J]. Materials Science and Engineering: A, 2018, 712: 603-607.

Yang H K, Tian Y Z, Zhang Z F. Revealing the mechanical properties and microstructure evolutions of Fe–22Mn–0.6 C–(x) Al TWIP steels via Al alloying control[J]. Materials Science and Engineering: A, 2018, 731: 61-70.

重庆大学：

Zhang Z, Chen D, Zhao H, et al. A comparative study of clock rolling and unidirectional rolling on deformation/recrystallization microstructure and texture of high purity tantalum plates[J]. International Journal of Refractory Metals and Hard Materials, 2013, 41: 453-460.

北京科技大学：

Chen J, Dong J, Zhang M, et al. Deformation mechanisms in a fine-grained Udimet 720LI nickel-base superalloy with high volume fractions of γ′ phases[J]. Materials Science and Engineering: A, 2016, 673: 122-134.

Jiang H, Dong J, Zhang M, et al. A study on the effect of strain rate on the dynamic recrystallization mechanism of alloy 617B[J]. Metallurgical and Materials Transactions A, 2016, 47(10): 5071-5087.

大连理工大学：

Gutierrez-Urrutia I, Li C L, Emura S, et al. Study of {332}< 113> twinning in a multilayered Ti-10Mo-xFe (x= 1–3) alloy by ECCI and EBSD[J]. Science and Technology of Advanced Materials, 2016, 17(1): 220-228.

中国科学院上海应用物理研究所钍熔盐反应堆中心：

Han F, He S, Liu M, et al. Hydrogen embrittlement susceptibility of a Ni-16Mo-7Cr base superalloy[J]. Materials Science and Engineering: A, 2018, 733: 291-298.

西北工业大学：

Zhu B, Xue X, Kou H, et al. The nucleation of microcracks under tensile stress in multi-phase high Nb-containing TiAl alloys[J]. Intermetallics, 2019, 106: 13-19.

中南大学：

Guo L, Wu W, Ni S, et al. Strengthening the FeCoCrNiMo0. 15 high entropy alloy by a gradient structure[J]. Journal of Alloys and Compounds, 2020: 155688.

Wang Z, Gu J, An D, et al. Characterization of the microstructure and deformation substructure evolution in a hierarchal high-entropy alloy by correlative EBSD and ECCI[J]. Intermetallics, 2020, 121: 106788.

Wang Z, Guo L, Xia W, et al. An SEM-based approach to characterize the microstructural evolution in a gradient CoCrFeNiMo0. 15 high-entropy alloy[J]. Materials Characterization, 2020, 161: 110169.

以上展示中可以看出，真正把这个技术用在位错领域做的很好的主要是：中科院金属所，中南大学。金属所做这项技术较早，04年就有第一篇工作，但是之后18年才有新的工作。中南大学主要是马普所合作，2020年发表了三篇，主要做高熵合金梯度结构材料，考虑到三篇工作都是最新发表的，不排除后续还有相关工作的报道。

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ECCI 电子通道衬度成像简介

文章目录ECCI简介ECCI的应用优势ECCI的应用现状ECCI文献分析国内研究单位ECCI简介ECCI，全写为electron channeling contrast imaging，即电子通道衬度成像，目前有部分研究人员利用该技术对块状样品表面的位错进行表征。这里是对检索文献过程中的收获做一些记录，以便督促自己学习，也方便后期整理。ECCI的应用优势ECCI技术是在扫描电镜下的表征技术，而常规位错表征常使用的为透射电镜。在金属塑性变形和断裂的研究领域，与常规透射电镜相比，ECCI显然具有一下对比

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复型技术只能对样品表面性貌进行复制，不能揭示晶体内部组织结构信息，受复型材料本身尺寸的限制，电镜的高分辨率本领不能得到充分发挥，萃取复型虽然能对萃取物相作结构分析，但对基体组织仍是表面性貌的复制。在这种情况下，样品减薄技术具有许多特点，特别是金属薄膜样品：可以最有..

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电子通道效应及电子通道衬度成像（ECCI）原理？ - 知乎

电子通道效应及电子通道衬度成像（ECCI）原理？ - 知乎首页知乎知学堂发现等你来答切换模式登录/注册材料科学晶体扫描电子显微镜电子显微镜透射电镜电子通道效应及电子通道衬度成像（ECCI）原理？有大神可以讲解一下吗？包括原理及应用，谢谢！显示全部关注者2被浏览865关注问题写回答邀请回答好问题添加评论分享1 个回答默认排序嗑学家仪器仪表制造业从业人员关注透射电镜对于样品制备要求较高，厚度一般要求100nm以下且观察区域较小。因此，随着高性能扫描电子显微镜SEM技术的进步，越来越多的研究者选择通过SEM配备的背散射电子探测器，基于电子通道衬度原理（ECCI）表征材料。再进一步结合EBSD和或摇摆电子束功能技术可以对衍射条件（g矢量）、位错类型及伯氏矢量等进行深入的分析，由于其高通量、高效率、高分辨率的特点，在金属、陶瓷及半导体材料的晶体缺陷研究中应用越来越广泛。电子通道效应是指当入射电子束与晶格满足布拉格衍射条件时，晶格点阵对电子的反射大大减弱，大量电子得以穿透晶格，呈现出“通道”效应（如图1所示）。对于完整晶体，当入射电子束相对于晶格的入射角度连续变化时，会出现通道效应和背散射效应的交替变化，最终在BSD探测器上叠加后显现出类似于菊池花样的衍射图像，称为电子通道花样（ECP），该现象最早是在1967年D.G.Coates用SEM观察锗、硅等单晶体时发现的。图1、a. 背散射衬度（亮） b. 通道衬度（暗）c. 晶粒内存在缺陷电子通道效应和入射电子束与晶体晶面的夹角相关，当电子束入射方向小于θB（布拉格衍射角）背散射电子数量较大，进入晶体的几率较小，属于禁道，在BSE图像上显示较亮的区域；当电子束入射方向大于θB （布拉格衍射角）背散射电子数量较小，进入晶体的几率较大，属于通道，在BSE图像上显示较暗的区域。多晶材料包含不同取向的晶粒，呈现出不同的图像衬度，因此不需要进行样品腐蚀便可通过ECCI得到不同晶粒的取向衬度；对于单晶或者多晶材料，当存在晶体缺陷（位错、层错、晶界）或者样品经过形变处理，通过摇摆电子束法获取指定晶粒的通道效应，晶粒与缺陷相对于入射电子束的夹角不同，从而呈现出不同的衬度。图2、通道效应的衬度差示意图电子通道效应主要来源于500埃的表面层，可以反映样品表面的缺陷，同时样品表面状态也会影响电子通道图像的质量，为了获得优异的ECC图像，样品制备可以采用振动抛光、电解抛光或者氩离子抛光等方法获取表面无氧化、无脏污、无应力的镜面。因为电子通道衬度效应非常弱，且要满足一定的电子光学条件，尤其衬度分辨率及角分辨率是电子通道成像质量的重要影响因素，对于电镜提出了更高要求，尤其是镜筒光路设计、探测器的灵敏度、样品台精度等。基于以上要求，如何获得质量优异的ECC图像？01为了获得观察晶粒的ECP，采用适当的入射电子扫描方式（光路设计）使入射电子束相对于晶面连续改变其入射角。目前有2种方法（电子束扫描法和电子束摇摆法），其中电子束扫描法适合于单晶样品，电子束摇摆法适用于单晶或者多晶样品，如图3所示。图3、不同的入射电子扫描方式a. 电子束扫描法 b. 电子束摇摆法02因为电子通道效应衬度较低，图像的衬度分辨率是重要的影响因素。对于加速电压、探针束流等参数设定，提供如下参考（不同材料对应的参数也会有区别，需要根据实际情况调整）：加速电压：20-30kv探针束流：1-4nA图像分辨率：最好大于 1536X1092Dwell Time ：10-20us案例对于多晶材料，比如金属及陶瓷可以观察其取向衬度及变形后缺陷，如下图4a为SLM制备316L不锈钢的微观组织，将图中红框位置放大（图4b）可以看到产生通道效应晶粒内部的大量位错胞，图像放大之后（图4c）可见析出相（黑色颗粒）与位错相互关系。图4、SLM制备316不锈钢的ECCI压电陶瓷材料在外加电场或机械应力作用下，其晶胞结构发生微小形变，发生正负电荷分离和极化，在材料内部形成电畴，常用于压电器和传感器等，电畴的形成基于材料的晶体结构变化。通过ECCI除了可以观察压电陶瓷晶粒取向衬度外，还可以观察其内部电畴分布情况，图5为Na0.5K0.5NbO3压电陶瓷外加电场后采集的电畴ECCI结果。图5、Na0.5K0.5NbO3压电陶瓷电畴 ECCI最后关注嗑学家，不定期发出你所不知道的科学世界！编辑于 2023-12-02 21:43赞同添加评论分享收藏喜欢收起

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Electron channeling contrast imaging - ECCI | Max-Planck-Institut für Eisenforschung GmbH

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Electron channeling contrast imaging - ECCI

ECCI is an imaging technique in scanning electron microscopy based on electron channelling applying a backscatter electron detector. It is used for direct observation of lattice defects, for example dislocations or stacking faults, close to the surface of bulk samples.

Individual dislocations and dense dislocation walls in a Fe 3% Si alloy, fatigued by 180 cycles of cyclic loading of 0.5 % of strain. Collaboration with J. Bouquerel, ESCIL, Lille.

Understanding ECCI is done best by looking into the physics of interaction of an electron beam with a crystal: When high energy primary beam electrons of a scanning electron microscope enter into a crystalline sample they form a standing electron-density wave inside the lattice which is coherent with the crystal lattice, the so-called primary wave field. Depending on the direction of the primary electron beam with respect to the lattice the maxima of the electron density waves may change from a position at the atomic nuclei to one in between them. In the first case strong backscattering of electrons out of the primary wave field occurs, while in the latter case only few electrons are backscattered. As a consequence, the backscatter signal carries information about the crystal lattice and its orientation relative to the primary beam. Minimum backscattering occurs when the primary beam almost exactly fullfills the Bragg angle with one of the lattice planes. The electrons then travel deep into the crystal without intense interaction with it. This case is called electron channelling.

If a defect, e.g. a dislocation or a stacking fault is present in the crystal then the coherency of the channelling primary electron wave field with the lattice is disturbed and strong backscattering occurs at the position of the defect. As a result, the defect is visible as a bright feature on an otherwise dark background when the sample is observed with a backscatter electron detector.

ECCI shows very comparable contrast features to dark-field TEM with the advantage that the images are obtained on bulk samples rather than on thin foils. In contrast, resolution and contrast are less good than in TEM.

When conventionally applying the ECCI technique for defect observation in a polycrystalline sample, it is not known which lattice plane leads to the observed defect contrast and whether the exact channeling conditions are excited. This can be corrected for by measuring the crystal orientation with EBSD or with ECP (electron channeling pattern). When ECP is available (which is rarely the case on modern microscopes) then the channeling contrast can be directly selected by tilting the sample until a channeling (or Kikuchi) line runs through the centre of the pattern. Subsequently, the microscope is run in scanning mode and the desired channelling contrast is observed. When the orientation is measured by EBSD, a software is required to emulate the theoretically obtained EC pattern based on the measured orientation. This software also emulates the sample tilt required to obtain good channelling conditions. The software-determined tilt conditions are subsequently set to the microscope stage and channelling contrast is observed.

We call the control of the channelling conditions via ECP or EBSD technique as "Electron channeling contrast imaging under controlled diffraction conditions", short cECCI. The figure below shows a schematics of this technique: (1) determination of a crystal orientation via EBSD. (2) Simulation of an ECP for the position of interest using the computer program TOCA. This program is also used to determine the tilt angles required for two-beam conditions. (3) Positioning of the sample as determined in step (2). (4) and (5) Observation of the channelling contrast. (6) Crystallographic interpretation of the channelling contrast using, for example, TOCA.

Schematics of cECCI by combination of EBSD and ECCI (see text for details)

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In-Situ Electron Channeling Contrast Imaging under Tensile Loading: Residual Stress, Dislocation Motion, and Slip Line Formation | Scientific Reports

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In-Situ Electron Channeling Contrast Imaging under Tensile Loading: Residual Stress, Dislocation Motion, and Slip Line Formation

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Published: 14 February 2020

In-Situ Electron Channeling Contrast Imaging under Tensile Loading: Residual Stress, Dislocation Motion, and Slip Line Formation

Keiichiro Nakafuji1, Motomichi Koyama2 & Kaneaki Tsuzaki1,3

Scientific Reports

volume 10, Article number: 2622 (2020)

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Mechanical propertiesMetals and alloys

AbstractElastoplastic phenomena, such as plastic deformation and failure, are multi-scale, deformation-path-dependent, and mechanical-field-sensitive problems associated with metals. Accordingly, visualization of the microstructural deformation path under a specific mechanical field is challenging for the elucidation of elastoplastic phenomena mechanisms. To overcome this problem, a dislocation-resolved in-situ technique for deformation under mechanically controllable conditions is required. Thus, we attempted to apply electron channeling contrast imaging (ECCI) under tensile loading, which enabled the detection of lattice defect motions and the evolution of elastic strain fields in bulk specimens. Here, we presented the suitability of ECCI as an in-situ technique with dislocation-detectable spatial resolution. In particular, the following ECCI-visualized plasticity-related phenomena were observed: (1) pre-deformation-induced residual stress and its disappearance via subsequent reloading, (2) heterogeneous dislocation motion during plastic relaxation, and (3) planar surface relief formation via loading to a higher stress.

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IntroductionMetals have been extensively used in structural applications due to their superior ductility and toughness. These characteristics are attributed to plasticity-driven strain evolution and associated stress re-distribution. The plasticity is originated from dislocation glides, which have a variable behavior sensitive to local shear stress, depending on external load and stress concentration. Therefore, in order to understand ductility and toughness, the plasticity mechanisms under certain mechanical conditions need to be shown. In this context, observations using a dislocation-resolved in-situ technique and done under mechanically controllable conditions are the most suitable approach.From a spatial resolution perspective, transmission electron microscopy (TEM) is an effective technique for resolving dislocations1,2,3,4. However, regarding in-situ deformation experiments, the special boundary conditions in a thin foil-type specimen may be problematic, giving unwanted signals such as effects of image force5,6,7,8,9. Consequently, in addition to the TEM approach, there is a high demand for a dislocation-resolved in-situ technique suitable for bulk specimens.Electron channeling contrast imaging (ECCI) enables visualization of dislocations10,11,12,13, stacking faults6,11,12,14, twins11,12,15,16, and elastic strain fields using a field-emission scanning electron microscope. Thus, ECCI can be used to characterize bulk specimens, allowing for the use of the conventional geometry of the mechanical test specimens13,17,18,19. Therefore, dislocation-resolved ECCI was chosen as a mechanically specified in-situ technique for the analysis of deformation. Moreover, since lattice defects can be detected under optimal specimen surface orientation, nanometer-scale deformation heterogeneity can be visualized in an observational area greater than a 10 µm square.Therefore, in this study, we investigated the suitability of ECCI as an in-situ characterization technique, while presenting its applicability for in-situ observation of a bulk metallic specimen under tensile loading.Results and DiscussionVisualization of residual stress variationFigure 1a shows an electron channeling contrast (ECC) image without external stress. This image was taken after the sample was submitted to pre-deformation of 2% tensile strain (corresponding to 246 MPa of external stress), followed by mechanical polishing. The alloy used in this experiment was an fcc ferrous alloy with the chemical composition of Fe15Mn10Cr8Ni in mass%. The loading and reloading process is shown in Fig. S1. As shown in Fig. 1a, a wide white area appears around the circular hole, corresponding to the black area in the upper right side. Additionally, Fig. 1b shows the microstructure with external stress after reloading to 199 MPa (elastic regime) and subsequent displacement holding for 8 min. As seen, the white area disappears after the reloading in the elastic regime. Note that significant dislocation movement was not observed during the reloading, as recognized by comparing Fig. 1a,b. Thus, the change in the electron channeling contrast from Fig. 1a,b indicates that the local contrast difference around the circular hole in Fig. 1a arises from the presence of residual stress.Figure 1Change in electron channeling contrast associated with relaxation of residual compressive stress. Electron channeling contrast images of the 2% pre-deformed and mechanically polished specimen (a) under an unloading condition and (b) after reloading to 199 MPa and displacement holding for 8 min. The incident beam direction was [0.51 –0.35 0.79] near the [1 –1 2] direction. The initial 1-μm-radius circular hole corresponds to the black areas in the upper right of the images. (c) von Mises equivalent stress (σeq) map of the pre-deformed specimen without external stress. The general-purpose finite element program ANSYS 17.0 (http://www.ansys.com/) was used to depict (c).Full size imageTo clarify the origin of the white region in Fig. 1a, we utilized a three-dimensional finite element method and found that residual stress exists around the circular hole after 2% pre-straining and unloading (Fig. 1c). The 2% pre-straining causes significant plastic strain and higher strain near the circular hole due to the stress concentration. Thus, the region surrounding the stress concentration source consequently compresses the region with high plastic deformation during the unloading process. Therefore, the elastic deformation attributed to the residual compressive stress produces elastic lattice strain and lattice rotation.Furthermore, the electron channeling contrast has a strong dependence on the backscatter electron signal because of the electron channeling mechanism. This signal is originated from the angular difference of the incident beam, starting from a Bragg position, with respect to low-index lattice planes, which are nearly parallel to the beam. Hence, the lattice strain and lattice rotation change the angular difference, and consequently, the electron channeling contrast.Therefore, the white region in Fig. 1a originated from the residual compressive stress. In turn, tensile reloading relaxes this residual compressive stress, resulting in the disappearance of the white region observed in Fig. 1b. To our knowledge, this is the first successful report of residual stress variation visualization using ECCI.

In-situ observation of dislocation motionNext, we determined the microstructure-dependent heterogeneous motion of dislocations. As previously discussed, the region near the hole was preferentially deformed due to stress concentration. However, dislocations around the hole were not observed during the displacement holding test conducted after reloading to 199 MPa (shown in the area marked by white lines in Fig. 2a,b). Moreover, Fig. 1 showed a compressive residual stress around the hole after the 2% straining and unloading. Therefore, we hypothesize that the compressive residual stress decreases the local resolved shear stress produced by the reloading, which, in turn, reduces the driving force for dislocation motion near the hole. This was supported by the lack of evidence of dislocation motion near the circular hole area (seen in Fig. 2c1–c4).Figure 2Dislocation distribution around the circular hole during the displacement holding test performed after reloading to 199 MPa (elastic regime). (a) Overview of the area observed using ECCI, after displacement holding for 8 min. (b) ECC image after displacement holding for 24 min. (c1)–(c4) Sequential ECC images of the area marked by white lines in (a,b) after displacement holding for (c1) 8, (c2) 15, (c3) 18, and (c4) 24 min. A movie of (c) is available in the supplemental materials.Full size imageContrastingly, we found dislocation motions during the displacement holding in regions considerably away from the hole (shown as Figs 3a1 in 2a). Figure 3a1–a6 show sequential ECC images of the area marked by the white dashed lines in Fig. 2a after reloading to 199 MPa (elastic regime) and subsequent displacement holding for 8, 11, 15, 18, 21, and 24 min, respectively. These figures picture two different dislocations on the {1 1 1} planes of the fcc ferrous alloy. It is possible to see that these dislocations ((i), yellow arrows and (ii), blue arrows) moved during the holding. Notably, dislocation motion was not observed when the stress level was low as shown in Fig. 2c1–c4. That is, the electron beam during the imaging did not induce any dislocation motion under the present observational condition. In other words, the dislocation motion shown in this study was attributed to external stress effects. Also, since the displacement holding was conducted at an ambient temperature, we concluded that these movements were originated from a glide process and not from a climb process.Figure 3Dislocation motions in the area marked in Fig. 2a, after reloading to 199 MPa and displacement holding for (a1) 8, (a2) 11, (a3) 15, (a4) 18, (a5) 21, and (a6) 24 min. (b) Simulated electron channeling pattern with the incident beam direction (i.e., [0.51 –0.35 0.79]). (c) Stereographic projection representing <1 1 1>, <1 1 0>, and <1 1 2> directions, where the yellow and green dashed lines indicate the traces of dislocations (i) and (ii) in (a1). The orange curve indicates the (1 1 1) plane as a great circle. (d) Schematic dislocation loop indicating the possible two cases of dislocation segment (ii) in (a1). A movie of (a) is available in the supplemental materials.Full size imageThus, we tried to identify the activated slip system of the dislocations. Unfortunately, the resulting Burgers vectors could not be determined through g·b = 0 analysis, due to the conditions used for the incident beam direction - [0.51 –0.35 0.79] near the [1 –1 2] pole (Fig. 3b). With this beam direction, we could not obtain {1 1 1} or {0 0 2} diffraction vectors that had sufficiently strong backscattered electron intensity for ECCI except for the (1 –1 –1) vector. Therefore, we used the Schmid factors in order to estimate the slip system (summarized in Table 1). According to Schmid factor criteria, the slip plane where the dislocations glided was the (1 1 1) plane, with a slip direction of [−1 0 1] or [1 –1 0]. With this information, we then examined the dislocation characteristics based on the dislocation morphology. The dislocation line directions shown in Fig. 3a1 were estimated from both the slip plane and the dislocation line traces in the stereographic projection (shown in Fig. 3c). Any dislocation line trace was obtained through the projection trace of partial and perfect dislocations on the specimen surface. Note that the line traces of dislocations (Fig. 3a(i,ii)) are parallel, and that the angle between the line trace and the tensile direction [−0.42 0.70 0.58] is 43°. Thus, the dislocations are on the plane of which the normal direction is n, as shown in the following formula:$$n=[\,-0.42\,\,0.70\,\,0.58\,]\,\sin \,{43}^{\circ }+[\,0.75\,\,0.63\,\,-0.21\,]\,\cos \,{43}^{\circ }=[\,0.27\,\,0.94\,\,0.24\,]$$

(1)

Table 1 Schmid factors of the slip system for the observed grain. The incident beam direction was [0.51 –0.35 0.79] and the tensile direction was [−0.42 0.70 0.58].Full size table([0 42 0 70 0 58]: right direction in Fig 3c, [0 75 0 63 0 21] : downward direction in Fig 3c).The dislocations are also on the (1 1 1) plane from the Schmid factor discussion and hence the dislocation line vector t can be determined as follows:$$t=[\,1\,\,1\,\,1\,]\times n=[\,-0.69\,\,0.02\,\,0.67\,]$$

(2)

When the dislocations have a ± a/2[−1 0 1] Burgers vector, the angle between the Burgers vector and the dislocation line vector t is 2° and they have a screw character. When the dislocations have a ± a/2[1 –1 0] Burgers vector, the angle between the Burgers vector and the dislocation line vector t is 122° and they have a mixed character.Once the dislocation line vector was determined, the visible depth for ECCI could be evaluated. When the dislocation line vector t is given as [−0.69 0.02 0.67] in Eq. (2), the visible depth in the sample becomes a value between 20 and 30 nm using the line lengths of the projected dislocations (130 nm) and the inclination angles of the dislocation lines with respect to the surface (10°). The typical value of the maximum depth of visibility has been reported to vary between 50 and 100 nm for scanning electron microscope conditions11, which is greater than the range obtained in our study. However, the range difference could come from the sample used in the reported study containing sufficient chromium (10 mass%) to form a thin passive film of chromium oxide, leading to a maximum depth of visibility smaller than normally expected.Next, we examined the validity of the dislocation motion observed during the displacement holding test performed after reloading to 199 MPa (elastic regime) (Fig. 3a). The dislocations in Fig. 3a1 are either segment (I) or (II) of the dislocation loops, schematically depicted by the red lines in Fig. 3d. Note that the specimen was pre-deformed to a 2% tensile strain and then mechanically polished with a layer of 30 µm in thickness. Hence, the dislocation source was considered to either be above or below the dislocation segment. Case I shows that the source still exists in the specimen (below the dislocation segment), while case II indicates that the source was in the polished region (above the segment). Fig. 3a3–a6 depict dislocation (ii) as a dot, indicating that the dislocation line changed to be approximately perpendicular to the specimen’s surface. With this position change, the dislocation loop shrinks in case II, whereas in case I the loop expands for the present dislocation motion. Note that the dislocation loops must expand for plastic relaxation during the displacement holding. Thus, it is provable that segment (ii) observed in Fig. 3a1 corresponds to case I in the schematic of Fig. 3d. It is also supported by the length of the dislocation line. When the dislocation line is approximately perpendicular to the surface, namely the line length of the projected dislocation is smallest on the (1 1 1) plane, the value of the inclination angle of the dislocation line with respect to the surface is 33° and the value of the line length of the projected dislocation is between 10 and 20 nm. This is reasonable for the expanding motion of the dislocation loop, with the dislocation source below the segment.However, the image force from the specimen surface was not considered in the aforementioned discussion. To evaluate the image force in a simple manner, it was assumed that dislocation (ii) had a screw character with a [−1 0 1] Burgers vector and the distance between dislocation (ii) and the image dislocation was 40 nm, corresponding to twice the visible depth for ECCI (20–30 nm). With this assumption and a shear modulus of G = 78.9 GPa20, we obtained shear stress of 68 MPa in the [−1 0 1] direction on the (1 1 1) plane, which arises from the image dislocation. The magnitude was approximately the same as the resolved shear stress arising from the external load: 70 MPa (199 MPa × 0.35). Therefore, when considering the image force, case II in the schematic of Fig. 3d is also possible for the case of dislocation segment (ii) shown in Fig. 3a1.

In-situ observation of slip line formationAfter displacement holding for 24 min (Fig. 3a6), the specimen was further loaded to 223 MPa, with the loading stage still macroscopically elastic, and held at constant displacement for 281 min (Fig. S1). After the second displacement holding and unloading, a new type of defect, line patterns, was determined using secondary electron imaging (shown in Fig. 4b, corresponding to the area marked by white lines in Fig. 4a). Secondary electron imaging is a surface-relief-sensitive technique. Hence, the observed line patterns were attributed to surface relief, (i.e. slip lines). The observed lines are along the (1 1 1) plane, which is the slip plane with the maximum Schmid factor (Table 1). Note that the angle between the specimen surface and the (1 1 1) plane is 57°. Similar line patterns were not observable using ECCI during the stage after reloading to 199 MPa and displacement holding for 24 min (Fig. 4c1). However, these patterns were observed after further loading to 223 MPa and displacement holding for 5 min (Fig. 4c2), indicating that these line patterns were induced by a stress increase.Figure 4Formation of line patterns and dislocation motions during the displacement holding test performed after further loading to 223 MPa (elastic regime). (a) ECC image depicting nearly the same area shown in Fig. 3a. (b) Secondary electron image of the area marked in (a) taken after the second displacement holding and unloading. This image was obtained at an acceleration voltage of 10 kV, a probe current of 10 nA, and a working distance of 10.3 mm. (c1)–(c3) ECC images of the area in (b). (c1) was obtained after loading to 199 MPa and displacement holding for 24 min. (c2) and (c3) were obtained after further loading to 223 MPa and displacement holding for 5 and 14 min, respectively. The red arrows show the shape change of the dislocation line in a high dislocation density wall. (d1)–(d4) ECC images of the area marked in (c3) after loading to 233 MPa and displacement holding for (d1) 14, (d2) 24, (d3) 33, and (d4) 36 min. Movies of Fig. 4c,d are available in the supplemental materials (Figs. S5 and S6). Wider region figures of Fig. 4 are available in Fig. S7.Full size imageNotably, Fig. 4 shows that both motion and morphology are different between the single dislocations and line patterns. For single dislocations, there are two motion patterns depending on the dislocation density. In the region of a high dislocation density wall, significant movement was hardly observed, but a shape change in the dislocation line was determined (red arrows in Fig. 4c1–c2). However, some single dislocations show significant movement in cell core regions with a low dislocation density (yellow and blue arrows in Fig. 4d1–d4). The single dislocation indicated by the blue arrow in Fig. 4d1 were emitted from the specimen surface, as shown in Fig. 4d2. The single dislocation indicated by the yellow arrow moved to the right (Fig. 4d1–d3) and stopped at the high dislocation density wall (Fig. 4d4). Therefore, we concluded that the motions of the single dislocations are restricted by the high dislocation density wall.In contrast to the single dislocation motions, the line patterns in Fig. 4 are seen to pass through the high dislocation density wall. Because deformation-induced martensite and twin plates are able to grow through dislocation cells and walls21,22, the observed line patterns possibly correspond to the group motion of partial dislocations. Since the fcc alloy used in this work has a low stacking fault energy of 20 mJ/m2 23,24,25, twinning or fcc-to-hcp martensitic transformation with a group motion of Shockley partial dislocations can occur21,26,27,28,29. Therefore, we determined the Schmid factors for partial dislocation motion (summarized in Table 2). This table indicates that the group motion of partial dislocations most easily occurs on the (1 1 1) plane. The following are possible reactions derived from the Table 2:$$\frac{1}{2}{[-101]}_{(111)}\,\to \,\frac{1}{6}{[-211]}_{(111)}+\frac{1}{6}{[-1-12]}_{(111)}$$

(3)

$$\frac{1}{2}{[-110]}_{(111)}\,\to \,\frac{1}{6}{[-211]}_{(111)}+\frac{1}{6}{[-12-1]}_{(111)}$$

(4)

Table 2 Schmid factors of the twinning system for the observational area of the ECCI.Full size tableFrom these features, it is very possible to conclude that the ECCI detected the group motion of partial dislocations for twinning or fcc-to-hcp martensitic transformation.In summary, we successfully characterized stress and lattice defect motion/formation behaviors in a bulk specimen using a scanning electron microscope-based technique. In particular, the dislocation-resolved ECCI technique was applied under loading conditions, which had never been reported before and demonstrated some patently clear advantages. Firstly, scanning electron microscope-based ECCI was performed in a bulk specimen and is, thereby, applicable for a specimen geometry used for normal tensile tests. Additionally, lattice strain and rotation were shown to alter ECCI, which facilitates the visualization of the variation in residual stress distribution, depending on remote stress conditions. Finally, ECCI enabled visualization of both dislocation glides and surface relief formation under loading. Therefore, since the observational area is at a grain-size-scale, the heterogeneous lattice defect motion and slip line formation in the crystal grain can be kinetically characterized using the approach reported herein.However, the in-situ observations using ECCI were limited to the elastic regime to avoid distinct surface relief formation that could significantly disturb the quality of lattice images. Nevertheless, we believe that ECCI can be advantageously applied for analysis of more severe plastic deformations, provided that the orientation or the mechanical conditions are confined to induce in-plane plastic deformation.MethodsMaterial preparationWe prepared an Fe15Mn10Cr8Ni austenitic steel (mass%). An ingot was prepared by vacuum induction melting. The ingot was forged and caliber rolled at 1373 K. Subsequently, the steel was solution-treated at 1273 K for 1 h followed by water quenching to suppress uncontrolled precipitation and segregation. The microstructure was fully austenitic prior to deformation, and showed deformation-induced ε martensitic transformations17,20,23. The solution-treated bar was cut into the desired shape by electrical discharge machining. The gauge shape of the tensile specimens was 2 mmw × 1 mmt × 10 mml.Displacement holding experiment under SEMLoading history schematics is shown in Fig. S1. Briefly, a specimen was firstly pre-deformed to 2% tensile strain (corresponding to 246 MPa: σ*) at room temperature with an initial strain rate of 5 × 10−4 s−1 (estimated from the cross-head speed); this was achieved using an in-situ tensile stage (TSL Solutions CO., Ltd) outside the scanning electron microscope (Merlin, Carl Zeiss). The strain was measured using a strain gauge. Then, the pre-deformed specimen was mechanically ground and polished, which reduced the specimen’s thickness by 30 µm. Afterwards, the pre-deformed specimen with the tensile stage was set into the field emission scanning electron microscope for performing the in-situ ECCI observations. ECCI was conducted in the region on the polished specimen surface, wherein a circular hole with a radius of 1 µm existed. The hole is associated with the presence of manganese oxide (MnO). The sample contains many spherical oxides with a diameter ranging from submicron to several microns. ECCI was operated at an acceleration voltage of 30 kV and a probe current of 2 nA at a working distance of 2.6 mm. When the surface orientation was optimized for a Bragg’s condition, the contrast in the resulting ECC image appears dark. Therefore, local residual stress can be visualized as a contrast gradation from dark to bright in the ECC image.The pre-deformed specimen was held at a constant displacement for 24 min after reloading to 199 MPa (approximately σ* × 0.8), and subsequently for 281 min after reloading to 223 MPa (approximately σ* × 0.9) using the field-emission scanning electron microscope. During the displacement holding tests, we observed dislocation motion using ECCI. In this work, all ECCI was performed at 30 kV and 2 nA.Thereafter, lattice defects and associated crystallographic features such as {1 1 1} traces were characterized using electron backscatter diffraction and secondary electron imaging. The electron backscatter diffraction measurement was conducted at 20 kV and 10 nA and a beam step size of 0.3 µm. The secondary electron imaging was conducted at an acceleration voltage of 10 kV, a probe current of 10 nA, and a working distance of 10.3 mm.Finite element methodThe visualized residual stress distribution in the ECC image was simulated by calculating von Mises equivalent stress using the general-purpose finite element program ANSYS 17.0 (http://www.ansys.com/). The advantage of equivalent stress is expressing stress as a scalar, because the electron channeling contrast change via the residual stress is also a scalar criterion. The analysis conditions of the finite element method are shown in Fig. S2. Assuming that the observed hole is a hemisphere 1 µm in radius in the center of the specimen surface, one-quarter of the specimen was used in the analysis to take advantage of the symmetry and large deformations were applied. The material properties of the specimen are as follows: Young’s modulus E = 200 GPa, Poisson’s ratio = 0.27, and yield strength σY = 162 MPa20. Tangent modulus Y after the yielding point was set to 1.90 GPa from the result of a tensile test as shown in Fig. S2c.

Data availability

The datasets generated during and/or analyzed during the current study are available from the corresponding authors on reasonable request.

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Download referencesAcknowledgementsThis work was financially supported by JSPS KAKENHI (JP16H06365 and JP17H04956) and the support is greatly appreciated. MK is grateful to “Kazato Research Foundation” for the financial support.Author informationAuthors and AffiliationsDepartment of Mechanical Engineering, Kyushu University, Motooka 744, Nishi-ku, Fukuoka, 819-0395, JapanKeiichiro Nakafuji & Kaneaki TsuzakiInstitute for Materials Research, Tohoku University, 2-1-1 Katahira, Aoba-ku, Sendai, Miyagi, 980-8577, JapanMotomichi KoyamaElements Strategy Initiative for Structural Materials (ESISM), Kyoto University, Yoshida-honmachi, Sakyo-ku, Kyoto, 606-8501, JapanKaneaki TsuzakiAuthorsKeiichiro NakafujiView author publicationsYou can also search for this author in

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PubMed Google ScholarContributionsM.K. and K.T. designed the research. K.N. performed most experiments. K.N. wrote the main manuscript text and M.K. and K.T. revised it. All authors discussed and commented on the manuscript.Corresponding authorCorrespondence to

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Reprints and permissionsAbout this articleCite this articleNakafuji, K., Koyama, M. & Tsuzaki, K. In-Situ Electron Channeling Contrast Imaging under Tensile Loading: Residual Stress, Dislocation Motion, and Slip Line Formation.

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Electron channeling contrast imaging, ECCI, SEM, dislcoation, TWIP, twin, grain boundary

WelcomeSearch this websiteUseless Information about materialsWhy we need strong materialsThe Metal MenTeachingSustainable MetallurgyGoals of Sustainable MetallurgyCarbon Footprint of Green SteelThe Green Steel DealGreen Steel: the role of scrapPellets for Hydrogen-Based Green Steel MakingGreen Steel: Direct reduction of iron ore with hydrogenGreen Steel: Reduction of iron ore by hydrogen plasmaSimulation of Hydrogen-based direct reductionCombining direct and plasma reductionSustainable steel production: Strip casting of steelGreen steel making with ammoniaWhat is red mud and why is it dangerous?Sustainable metals for electromobilityGreen AluminiumMaking Aluminium from ScrapEBSD and 3D EBSDECCIDigital Image Correlation DICDigital Image Correlation for multiphase alloysDigital Image Correlation and crystal plasticity simulationDissipative deformation heatingAtom probe tomographyAtom probe tomography cryogenic UHV prepAtom probe and atomistic simulationsCorrelative atom probe tomographyAtom probe tomography on steelsAtom probe tomography on metallic glassField Ion Microscopy - Mapping single atomsCrystallographic TexturesGrain boundary segregation engineeringMultiscale modelingAb initio modelingDDD: Discrete dislocation dynamicsICMEShear bandingLinear ComplexionsNanoindentation and pillar compressionIndentation pile-upRecrystallization and grain growth simulationConstitutive modelingChemistry dependence of constitutive modelsCrystal Plasticity Finite Element MethodDislocation-based Crystal Plasticity Finite Element MethodPolycrystal mechanicsRVE and homogenization simulationSheet forming simulationYield Surface SimulationLarge scale crystal plasticity amd ansiotropy simulationPhase field and Ginzburg-Landau modelingJoint crystal plasticity and phase field modelsLattice Boltzmann model (LBM)Steels: Brief introductionSteels: facts, figures, environmentBainite steelsMartensite alloys and transformationsPearlitic steelsQP - Quench - partitioning steelsTRIPLEX steelsElectrical steels (Fe-3%Si)Soft magnetic alloysSoft magnetic steels for electromobilityODS Eurofer steelsNanolaminatesNanolaminate steelsNanosteelsSteels for additive manufacturingDual phase steelsMedium Mn steelsHydrogen and hydrogen embrittlementHydrogen embrittlement of medium Mn steelStainless steelsCombinatorial steel designUltrafine grained steelsTextures of SteelsTextures of dual phase steelsMetallurgical Alloy DesignMetastability Alloy DesignTWIP steelsTRIP steelsShape Memory SteelsCombining TWIP and TRIPLow density and weight reduced steelsHigh stiffness steelsIron-AluminidesHigh strength compositesThermoelectric half-Heusler alloysMagnesium alloysMagnesium applicationsCopper alloys and Copper-Nano-CompositesAluminium alloysAluminium alloys for aerospace applicationsHigh Entropy Alloys - OverviewMetastable High Entropy AlloysNon-Equiatomic High Entropy AlloysTWIP High Entropy AlloysDual Phase High Entropy AlloysHigh Entropy Alloys and HydrogenInterstitial High Entropy AlloysHigh Entropy Alloys by Text MiningHigh Entropy SteelsBi-directional TRIP High Entropy AlloysMedium Entropy AlloysCombinatorial discovery of high-entropy alloysTitanium alloysNickel Alloy 718SuperalloysCobalt AlloysNiobiumTantalumMultilayer Nitride CoatingsImportant Alloy ClassesBauschinger EffectLaser Additive ManufacturingAdditive Manufacturing with SteelSegregation Engineering in Additive ManufacturingWelding MicrostructuresSeminconductors Solar CellsBiological (natural) MaterialsChitin structureHierarchical Architecture of ChitinPolymer MaterialsDramatic material failureMovies and AnimationsRSS feeds MSE JournalsLecturesReprintsSteel properties and overview ImagesMaterials Science Image GalleryLarge projectsHistory of materialsUseful linksInfos for Materials ScientistsBasic structure information on metalsGlossary of materials scienceCantus MirabilisMetals in poetry3D modelsScience cartoonsAnimated gifs for scienceContactApplicantsMax-Planck-InstitutScientific CVJournal publicationsReview publicationsBooksResearch projects

ECCI (Electron Channeling Contrast Imaging) in the SEM and its application to high Mn steels Coupling of electron channeling with EBSD: towards the quantitative characterization of deformation structures in the SEM Abstract

The coupling of electron channeling contrast imaging (ECCI) with EBSD provides an efficient and fast approach

to perform ECCI of crystal defects, such as dislocations, cells, and stacking faults, under controlled diffraction conditions with

enhanced contrast. From a technical point of view, the ECCI technique complements two of the main electron microscopy techniques, namely, EBSD and TEM. In this review, we provide several application

examples of the EBSD-based ECCI approach on microstructure characterization, namely, characterization of single dislocations, measurement of dislocation densities and

characterization of dislocation substructures in deformed bulk materials. We make use of a two-beam Bloch wave approach to interpret the channeling contrast associated to crystal

defects. The approach captures the main features observed in the experimental contrast associated to stacking faults and dislocations.

Introduction to electron channeling contrast imaging

The characterization of deformation substructures is commonly performed by transmission electron microscopy techniques (TEM). The high spatial and angular resolution of TEM allow the

characterization of individual crystal defects such as dislocations, stacking faults and point defects, as well as more complex configurations formed by dislocations and deformation twins [1-3].

However, TEM presents some drawbacks such as sample preparation, where extremely thin samples are needed, and reduced observation area. The common methods for TEM sample preparation, namely,

electropolishing, ion milling and focused ion beam (FIB), are time consuming and may even modify the microstructure by introducing point defects (for instance, foil thinning by ion milling [4]) or

relaxing internal stresses (bulk thinning prior to electropolishing [5]). Furthermore, the resulting TEM foil offers only a limited observation area (~ 100 × 100 mm2) making it difficult to perform

quantitative microstructural characterization required in many topics in materials science. While TEM has proven to be a very versatile tool for the characterization of crystal defects, the

limitations outlined above have led to the development of diffraction-based methods such as X-ray diffraction (XRD) [6, 7], synchrotron radiation-based techniques [8-10], and 2D/3D electron backscatter diffraction (EBSD) techniques [11-15]. An alternative electron microscopy technique for characterizing bulk

deformed materials is electron channeling contrast imaging (ECCI) [16-18]. ECCI is a scanning electron microscope (SEM) technique that makes use of the strong dependence of the backscatter electron

(BSE) signal on the orientation of the crystal lattice planes with respect to the incident electron beam due to the electron channeling mechanism. Interestingly, slight local distortions in the

crystal lattice produced by the strain fields associated to dislocations yield a modulation in the BE signal that can be detected by available BSE detectors. Accordingly, crystal defects such as

dislocations can be imaged by ECCI in the SEM. According to the dynamical theory of electron diffraction, the optimum diffraction contrast for imaging crystal defects in ECCI is obtained by orienting

the crystal exactly into Bragg condition for a selected set of intense diffracting lattice planes [19]. However, since the early observations in 1967 by Coates of Kikuchi-like reflection patterns in

the SEM, the so-called electron channeling patterns (ECPs) [20], the impact of ECCI in materials science has been rather limited. This is due to the specific tilting conditions required to image

crystal defects in ECCI.

To date, two main ECCI set-ups have been developed, namely, forward-scatter electron (FSE) and backscatter electron (BSE)-based methods. FSE-based approaches make use of high-tilt EBSD-type

geometries to enhance the intensity of thermally diffuse (or phonon) scattered electrons. The FS detector records the contrast created by the Kikuchi bands intersecting the detector. This

configuration, originally used by Czernuszka et al. [21] to image dislocations in Si and Ni3Ga, has been adapted by Wilkinson’s group [16, 22, 23] and Picard et al. [24-26] to examine dislocations in

SiGe/Si epitaxial layers [22], and GaN and SiC thin films [23-26]. BSE-based methods make use of electron channeling patterns to optimize the diffraction contrast of the crystal defect image. These

patterns can be obtained after running the microscope at low magnification (so-called electron channeling patterns, ECPs) [17, 18, 20] or after rocking the incoming beam on the material surface by a

deflection focusing technique (so-called selected-area channeling patterns, SACPs) [27-29]. The main limitation of these approaches is the large observation area, typically >10 mm, and the low

spatial resolution, 2-5 mm. This shortcoming limits the application of ECPs/SACPs-based ECCI set-ups to lightly deformed coarse polycrystals and single crystals [16, 29-32]. Recently, we have

proposed the use of EBSD to obtain the rocking parameters of the sample surface under the incoming beam (tilt and rotation angles) where the crystal is oriented into optimum diffraction contrast

[33]. Due to the high spatial resolution of EBSD (35±5 nm [34]), crystal defects can be imaged in sub-micrometer areas of deformed bulk materials. In the last few years, we have applied the

EBSD-based ECCI approach to several cases, namely, the characterization of complex deformation structures such as dislocation and nano-twin substructures in bulk deformed high-Mn TWIP steels (TWIP:

twinning induced plasticity) [35-39], the investigation of dislocation patterns [36, 37, 40], the measurement of dislocation densities in deformed alloys [41, 42] and the characterization of

individual dislocations and stacking faults [43, 44]. These studies have demonstrated the feasibility of the EBSD-based ECCI approach on the quantitative characterization of complex deformation

structures in bulk materials. From a technical point of view, the ECCI technique complements two of the main electron microscopy techniques, namely, EBSD and conventional diffraction-based TEM. The

present paper summarizes the experimental configuration of the EBSD-based ECCI set-up and illustrates its use with several experimental studies. We also provide a physical-based model of ECC contrast

associated with crystals defects.

References:

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A study on the geometry of dislocation patterns in the surrounding of nanoindents in a TWIP steel using electron channeling contrast imaging and discrete dislocation dynamics simulationsJ.-l. Zhang, S. Zaefferer, D. Raabe

Materials Science & Engineering A 636 (2015) 231-242

J.-l. Zhang, S. Zaefferer, D. Raabe Materials Science & Engineering A 636 (2015) 231-242 A study on the geometry of dislocation patterns in the surrounding of nanoindents in a TWIP steel using electron channeling contrast imaging and dislocation modeling.

Instrumented nanoindentation is an excellent tool for characterising the mechanical properties and the deformation behaviour of materials at nano/micro-scales [1,2]. The mechanical data,

such as hardness and elastic modulus, obtained by nanoindentation are, however, not easily interpretable in terms of macroscopic mechanical properties because of the complex stress and

strain field developed during the test. Although numerous studies were carried out to understand the strain field formed beneath and besides the indenter, a number of questions about

the associated dislocation activities are still unresolved. It is known, for example from 3D EBSD investigations, that fields of different crystallographic rotations are formed below

an indent. Transmission electron microscopy (TEM) observations of nanoindentation provided detailed information of the mechanisms associated with localized deformation. All these

experimental investigations have certain limitations though: The 3D EBSD observations reveal the existence of geometrically necessary dislocations in terms of the detected rotation

patterns but fail to show the actual complete deformation pattern in terms of the underlying dislocations. TEM observations show the true dislocation arrangements but they suffer from

the fact that thin foils have to be used for that and, hence, either lateral or depth directional information cannot be obtained. To a certain extent the missing information can

be complemented by crystal plasticity simulations, as have been carried out in some studies. This approach, however, suffers from adjustable parameters such as latent hardening and

cross hardening parameters, and uncertainties associated with boundary condition treatment such as friction. A significant number of investigations have been reported on

the crystallographic orientation induced patterns around nanoindents

in materials with fcc structure. Most of the works were conducted by comparison of finite element simulations with secondary electron (SE) images or atomic force microscopy

(AFM) profiles. Irrespective of the shape of the indenter used in these studies, four- and two-fold symmetries for {100}- and {110}-oriented crystals were reported. For {111}-oriented

crystals both, six-fold [9] and three-fold [11] symmetries were suggested. In these works the surface topographical pile-up or sink-in patterns were used for comparison between simulation and experimental results. However, detailed information on how these patterns are

actually formed in terms of the underlying dislocation activities was not provided. In addition, the variation of hardness and/or elastic modulus with indentation depth, the so

called indentation size effect, raises a lot of difficulties with obtaining real values of the mechanical properties. In order to understand the size effect numerous studies have been

carried out in investigating the size of the plastic deformation zone as a function of the indentation depth and/or

indenter size. Most of the works were done by numerical calculations or by topographical pile-up or sink-in pattern observations using atomic force microscopy (AFM) and scanning

electron microscopy (SEM). Also, the real lateral extension of the plastic deformation zone was not sufficiently studied in detail yet. In order to overcome some of the mentioned

experimental and

simulation difficulties we used a newly designed technique, referred to as electron channeling contrast imaging under controlled diffraction conditions (cECCI). It allows the direct

observations of crystal defects like dislocations or stacking faults close to the surface of bulk samples. This technique, which has

similarities to dark field TEM, is applied in an SEM and allows probing approximately the first 50–100 nm of material below the surface.

In the present work we applied the cECCI technique to study dislocation structures in the surrounding of nanoindents in a steel with twinning induced plasticity (TWIP) and fcc crystal

structure. Nanoindentation was carried out either in a load-controlled mode with a maximum load or in displacement-controlled mode up to different indentation depths on {100}- {110}-

and {111}-oriented grains in a polycrystalline specimen. The aim of the study is to

understand in more detail the formation of the defects below and next to the indent and to contribute in this way to a better understanding of the pattern formation process. In order to be

able to interpret our experimental results we also applied discrete dislocation dynamics (DDD) simulations using the Parallel Dislocation Simulator (ParaDiS) code, which shows, to a

certain

extent, which slip systems are active and what the shape and the type of the resulting dislocations in the network is. These data were compared with our experimental results.

Here, electron channeling contrast imaging under controlled diffraction conditions (cECCI) is used since it enables observation of crystal defects, especially dislocations, stacking faults

and nano-twins, close to the surface of bulk samples. In this work cECCI has been employed to observe defects around nanoindents into the surface of {100}-, {110}-, {111}-oriented grains in

a Fe–22Mn–0.65C (wt%) TWIP steel sample (fcc crystal structure, stacking fault energy 20 mJ/m²) using a cone-spherical indenter. The dislocation patterns

show four- and two-fold symmetries for the {100}- and {110}-orientation, and a three-fold symmetry for the {111}-orientation which is, however, difficult to observe. Discrete dislocation

dynamics (DDD) simulations of the indentation were carried out to complement the static experimental investigations.

The simulations were carried out with both, cross-slip disabled and enabled conditions, where the former were found to match to the experimental results better, as may be expected for an fcc

material with low stacking fault energy. The 3-dimensional geometry of the dislocation patterns of the different

indents was analysed and discussed with respect to pattern formation mechanisms. The force–displacement curves obtained during indentation showed a stronger strain hardening for the

{111} oriented crystal than that for the other orientations. This is in contrast to the behaviour of, for example,

copper and is interpreted to be due to planar slip. Irrespective of orientation and indentation depth the radius of the plastically deformed area was found to be approximately 4 times larger

than that of the indenter contact area.

On the origin of creep dislocations in a Ni-base, single-crystal superalloy: an ECCI, EBSD, and dislocation dynamics-based studyActa Materialia 109 (2016) 151-161

On the origin of creep dislocations in a Ni-base, single-crystal superalloy: an ECCI, EBSD, and dislocation dynamics-based studyRam et al Ni base superalloy creep dislc[...] PDF-Dokument [2.2 MB]

Nickel-base superalloys are employed in high-temperature applications,

particularly for jet propulsion and power conversion. To increase creep resistance, these alloys are produced as single crystals. Single-crystal superalloys are often fabricated by

directional solidification techniques using competitive dendrite growth.

This work investigates the origin of creep dislocations in a Ni-base, single crystal superalloy subject to creep at an intermediate stress and temperature. Employing high angular resolution

electron backscatter diffraction (HR-EBSD), electron channeling contrast imaging under controlled diffraction conditions

(cECCI) and discrete dislocation dynamics (DDD) modelling, it is shown that low-angle boundaries - which correspond to dendrite boundaries or their residues after annealing - are not the major

sources of creep dislocations. At the onset of creep deformation, they are the only active sources. Creep dislocations are emitted from them and percolate into the dislocation-depleted

crystal. However, the percolation is very slow. As creep deformation proceeds, before the boundary-originated dislocations move further than a few micrometers away from their source

boundary, individual dislocations that are spread throughout the undeformed microstructure become active and emit avalanches of creep dislocations in boundary-free regions, i.e. regions

farther than a few micrometer away from boundaries. Upon their activation, the density of creep dislocations in boundary-free regions soars by two orders of magnitude; and the entire

microstructure becomes deluged with creep dislocations. The total area of boundary-free regions is several times the total area of regions covered by boundary-originated creep

dislocations. Therefore, the main sources of creep dislocations are not low-angle boundaries but individual, isolated dislocations in boundary-free regions.

Effects of strain amplitude, cycle number and orientation on low cycle fatigue microstructures in austenitic stainless steel studied by electron channelling contrast imagingActa Materialia 87 (2015) 86--99Acta Materialia vol 87 (2015) 86 ECCI fa[...] PDF-Dokument [3.0 MB]

Substructure analysis on cyclically deformed metals is typically performed by time-consuming transmission electron microscopy probing,

thus limiting such studies often to a single parameter. Here, we present a novel approach which consists in combining electron backscatter diffraction

(EBSD), digital image correlation and electron channelling contrast imaging (ECCI), enabling us to systematically probe a large matrix of different

parameters with the aim of correlating and comparing their interdependence. The main focus here is to identify the influence of cycle number, initial

grain orientation and local strain amplitude on the evolving dislocation patterns. Therefore, experiments up to 100 cycles were performed on a polycrystalline

austenitic stainless steel with local strain amplitudes between 0.35% and 0.95%. EBSD and ECCI maps reveal the individual influence of

each parameter while the others remained constant. We find that the dislocation structures strongly depend on grain orientation. Dislocation structures in grains with double-slip (<112> //

LD, <122> // LD and <012> // LD) and multiple-slip (<111> // LD, M <011> // LD and <001> // LD)

orientations with respect to the loading direction (LD) are characterized under the variation of strain amplitude and cycle number.

Revealing the strain-hardening behavior of twinning-induced plasticity steels: Theory, simulations, experimentsActa Materialia 61 (2013) 494-510

David R. Steinmetz, Tom Jäpel, Burkhard Wietbrock, Philip Eisenlohr, Ivan Gutierrez-Urrutia, Alireza Saeed–Akbari, Tilmann Hickel, Franz Roters, Dierk Raabe

Revealing the strain-hardening behavior of twinning-induced plasticity steels: Theory, simulations, experimentsActa Materialia 61 (2013) 494 ab initio [...] PDF-Dokument [1.8 MB]

Here we present a multiscale dislocation density-based constitutive model for the strain-hardening behavior in twinning-induced plasticity (TWIP) steels. The approach is a physics-based

strain rate- and temperature-sensitive model which reflects microstructural investigations of twins and dislocation structures in TWIP steels. One distinct advantage of the approach is that the

model parameters, some of which are derived by ab initio predictions, are physics-based and known within an order of magnitude. This allows more complex microstructural information to be

included in the model without losing the ability to identify reasonable initial values and bounds for all parameters.

Dislocation cells, grain size and twin volume fraction evolution are included. Particular attention is placed on the mechanism by which new deformation twins are nucleated, and a new formulation

for the critical twinning stress is presented. Various temperatures were included in the parameter optimization process. Dissipative heating is also considered. The use of physically justified

parameters enables the identification of a universal parameter set for the example of an Fe–22Mn–0.6C TWIP steel.

Coupling of Electron Channeling with EBSD: Toward the Quantitative Characterization of Deformation Structures in the SEMJOM, Vol. 65, No. 9, 2013

Coupling of Electron Channeling with EBSD: Toward the Quantitative Characterization of Deformation Structures in the SEM

I. GUTIERREZ-URRUTIA,S. ZAEFFERER, D. RAABEJOM Vol. 65 No. 9 2013 Overview Electron[...] PDF-Dokument [812.0 KB]

Electron Channeling Contrast Imaging: Application examples of the ECCI technique: (a) characterization of single dislocations, (b) stacking faults, (c–d) dislocation substructures, and (e) nanotwins; (f) measurement of dislocation densities

The coupling of electron channeling contrast imaging (ECCI) with electron

backscatter diffraction (EBSD) provides an efficient and fast approach to

perform ECCI of crystal defects, such as dislocations, cells, and stacking

faults, under controlled diffraction conditions with enhanced contrast. From a

technical point of view, the ECCI technique complements two of the main

electron microscopy techniques, namely, EBSD and conventional diffraction based transmission electron microscopy. In this review, we provide several

application examples of the EBSD-based ECCI approach on microstructure

characterization, namely, characterization of single dislocations, measurement

of dislocation densities, and characterization of dislocation substructures

in deformed bulk materials. We make use of a two-beam Bloch wave approach

to interpret the channeling contrast associated with crystal defects. The approach captures the main features observed in the experimental contrast

associated with stacking faults and dislocations. More specific, we use a new approach to electron channelling contrast imaging (ECCI), referred to as “controlled ECCI”, or cECCI. This method improves the existing ECCI method in a way that it

uses EBSD to determine the crystal orientation. Based on this orientation measurement the optimum sample alignment for obtaining good channelling contrast is calculated. A dedicated eucentric

goniometer stage is then used to move the sample into the calculated position for imaging of dislocations and other crystal defects in the SEM. The approach offers excellent opportunities for the

efficient quantification of substructure features at a large field of view that were not accessible so far to SEM characterization. It can be combined with EBSD maps so that we are now capable of

conducting detailed microstructure quantification mappings of orientations together with its dislocation substructure in the same experiment. Time consuming TEM investigations of dislocation

structures that provide a small field of view and that are affected by the thin slice preparation typical of TEM thus may become obsolete in a number of cases.

I. Gutierrez-Urrutia, D. Raabe, Dislocation and twin substructure evolution during strain hardening of an Fe–22 wt.% Mn–0.6 wt.% C TWIP steel observed by

electron channeling contrast imaging, Acta Materialia 59 (2011) 6449–6462

I. Gutierrez-Urrutia, S. Zaefferer, D. Raabe: The effect of grain size and grain orientation on deformation twinning in a Fe–22 wt.% Mn–0.6 wt.% C TWIP steel,

Mater. Sc. Engin. A 527 (2010) 3552-3560

I. Gutierrez-Urrutia, S. Zaefferer, D. Raabe: Electron channeling contrast imaging of twins and dislocations in twinning-induced plasticity steels under

controlled diffraction conditions in a scanning electron microscope, Scripta Mater. 61 (2009) 737-740

I. Gutierrez-Urrutia, D. Raabe: Viewpoint Set no. 50: Twinning Induced Plasticity Steels: Grain size effect on strain hardening in twinning-induced plasticity

steels,

Scripta Materialia, Volume 66, Issue 12, June 2012, Pages 992–996

ECCI Analysis on TWIP steel

Dislocations in the SEM by using ECCI and EBSD (Gutierrez-Urrutia, Zaefferer, Raabe: Scripta Mater. 61 (2009) 737)

For further details

please see also Stefan Zaefferer's group at the Max-Planck Institut.

Dislocation and twin substructure evolution during strain hardening of an Fe–22 wt.% Mn–0.6 wt.% C TWIP steel observed by electron channeling contrast imagingActa Materialia 59 (2011) 6449-6462

Dislocation and twin substructure evolution during strain hardening of an Fe–22 wt.% Mn–0.6 wt.% C TWIP steel observed by electron channeling contrast imaging

I. Gutierrez-Urrutia, D. RaabeActa Materialia 59 (2011) 6449 ECCI TWIP[...] PDF-Dokument [2.5 MB]

Acta Materialia 59 (2011) 6449 Dislocation and twin substructure evolution during strain hardening of an Fe–22 wt.% Mn–0.6 wt.% C TWIP steel observed by electron channeling contrast imaging I. Gutierrez-Urrutia, D. Raabe

High-manganese steels have received much interest in recent years due to their outstanding mechanical properties combining high strength and ductility. This property profile is

attributed to their high strain hardening capacity. High-manganese steels are typically austenitic steels, i.e. face-centered cubic (fcc) alloys, with a high Mn content (above 20%

wt.%) and additions of elements such as carbon (<1 wt.%), silicon (<3 wt.%) and aluminum (<10 wt.%). This steel grade exhibit different hardening mechanisms, such as

transformation-induced plasticity (TRIP), twinning-induced plasticity (TWIP) or microband-induced plasticity (MBIP). The activation of these mechanisms is strongly dependent on the stacking fault energy. TRIP is observed in very low stacking fault steels (below 20 mJ / m^2) and is associated with the transformation of austenite (fcc phase) into epsilon-martensite (hexagonal close-packed phase), which in turn further acts as nucleus of a'-martensite (body-centered cubic or tetragonal phase). TWIP is observed in medium stacking fault energy steels (20–40 mJ/ m^2) and is characterized by the formation of deformation twins with nanometer

thickness. MBIP has been recently reported in steel grades with high stacking fault energy (90 mJ / m^2) and is attributed to the formation of microbnds, which are in-grain shear

zones that are confined by geometrically necessary boundaries or conventional grain boundaries.

These microstructure features (e-martensite plates, deformation twins and microbands) lead to a remarkable variety of strain hardening phenomena as they all act as effective obstacles for dislocation glide. High-manganese TWIP steels are characterized by a hierarchical microstructure refinement that includes complex dislocation and twin substructures, and their interactions. Although there are some previous studies on the strain hardening behavior in TWIP steels, the details of the underlying kinetics of the substructure evolution and its correspondence to the stress–strain

and strain hardening evolution is not yet fully understood. Most of these works analyze strain hardening in terms of a dislocation mean free path (MFP) approach,

focusing essentially on a single microstructure parameter, namely,

the twin spacing. These works attribute the high strain hardening rate at intermediate strains (0.1–0.2 true strain) to twin spacing refinement. The increasing density of deformation

twin boundaries and the strong effect they

have on dislocation glide leads to the so-called “dynamic Hall–Petch effect”. However, our analysis reveals that the deformed microstructure of these alloys is too complicated to be

reduced to a single microstructure parameter and,

therefore, a detailed analysis of the contribution of dislocation and twin substructures, as well as their interactions, to strain hardening is required.

One important limitation in the characterization of TWIP steels is the complexity of the microstructure, which involves features of different length scales: deformation twins with

thicknesses of some tens of nanometers and dislocation substructures extending over several micrometers. As a consequence of this scale discrepancy, quantitative microstructure

characterization by

conventional electron microscopy techniques such as electron backscatter diffraction (EBSD) or transmission electron microscopy (TEM) is limited due to the angular resolution (EBSD)

and the small field of view (TEM), respectively. In Here, therefore, we make use of electron channeling contrast imaging (ECCI), which is conducted in a scanning electron microscope

(SEM), to perform a quantitative characterization of the deformation microstructure of TWIP steel. The ECCI technique has been established as an excellent tool for examining

complex deformation microstructures of metallic materials, revealing

microstructure features such as deformation twins, stacking faults and complex dislocation arrangements from a wide field of view directly in the SEM. The reason for the recent

improvement in the ECCI technique lies in its combination with EBSD. This allows us to efficiently identify optimum contrast conditions and, therefore, produce ECCI images of crystal

defects under controlled diffraction

conditions. Here, we study the kinetics of the substructure evolution and its correspondence to the strain hardening

evolution of an Fe–22 wt.% Mn–0.6 wt.% C TWIP steel during tensile deformation by means of electron channeling

contrast imaging (ECCI) combined with electron backscatter diffraction (EBSD). The contribution of twin and

dislocation substructures to strain hardening is evaluated in terms of a dislocation mean free path approach

involving several microstructure parameters, such as the characteristic average twin spacing and the dislocation

substructure size. The analysis reveals that at the early stages of deformation (strain below 0.1 true strain) the dislocation substructure provides a high strain hardening rate with hardening coefficients of about G/40 (G is the shear modulus). At

intermediate strains (below 0.3 true strain), the dislocation mean free path refinement due to deformation

twinning results in a high strain rate with a hardening coefficient of about G/30. Finally, at high strains

(above 0.4 true strain), the limited further refinement of the dislocation and twin substructures reduces the

capability for trapping more dislocations inside the microstructure and, hence, the strain hardening decreases.

Grains forming dislocation cells develop a self-organized and dynamically refined dislocation cell structure which follows the similitude principle but with a smaller similitude constant than that found in medium to high stacking fault energy alloys. We

attribute this difference to the influence of the stacking fault energy on the mechanism of cell

formation.

Hydrogen-assisted failure in a twinning-induced plasticity steel studied under in situ hydrogen charging by electron channeling contrast imagingWe study hydrogen embrittlement in Fe–18Mn–1.2%C (wt.%) twinning-induced plasticity steel, focusing on the influence of deformation twins on hydrogen-assisted cracking.Acta-Mater-hydrogen-induced-cracking-201[...] PDF-Dokument [4.0 MB]

Multistage strain hardening through dislocation substructure and twinning in a high strength and ductile weight-reduced Fe–Mn–Al–C steelHere we study the kinetics of the deformation structure evolution and its contribution to the strain hardening of a Fe–30.5Mn–2.1Al–1.2C (wt.%) steel during tensile deformation by means of trans2012-Acta FeMnAlC-multi-strain-hardening[...] PDF-Dokument [1.8 MB]

High Mn steel characterized by EBSD, ECCI and TEM (Materials Science and Engineering A 527 (2010) 3552).

Reprints on applications of the ECCI method (Electron Channeling contrast imaging) to the microstructure of high Mn TWIP and TRIPLEX

steels.

Application of the ECCI method to weight reduced TRIPLEX steels.

Understanding strain-hardening in twinning-induced plasticity (TWIP) steels: Theory, simulations, experimentsHere we present a multiscale dislocation density-based constitutive model for the strain-hardening behavior in twinning-induced plasticity (TWIP) steels. The approach is a physics-based strain rate- aActa-Materialia-Vol-61-2013-modeling-TWI[...] PDF-Dokument [1.8 MB]

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Coupling of Electron Channeling with EBSD: Toward the Quantitative Characterization of Deformation Structures in the SEM

Published: 16 July 2013

Volume 65, pages 1229–1236, (2013)

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Coupling of Electron Channeling with EBSD: Toward the Quantitative Characterization of Deformation Structures in the SEM

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I. Gutierrez-Urrutia1, S. Zaefferer1 & D. Raabe1

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AbstractThe coupling of electron channeling contrast imaging (ECCI) with electron backscatter diffraction (EBSD) provides an efficient and fast approach to perform ECCI of crystal defects, such as dislocations, cells, and stacking faults, under controlled diffraction conditions with enhanced contrast. From a technical point of view, the ECCI technique complements two of the main electron microscopy techniques, namely, EBSD and conventional diffraction-based transmission electron microscopy. In this review, we provide several application examples of the EBSD-based ECCI approach on microstructure characterization, namely, characterization of single dislocations, measurement of dislocation densities, and characterization of dislocation substructures in deformed bulk materials. We make use of a two-beam Bloch wave approach to interpret the channeling contrast associated with crystal defects. The approach captures the main features observed in the experimental contrast associated with stacking faults and dislocations.

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IntroductionThe characterization of deformation substructures is commonly performed by transmission electron microscopy (TEM) techniques. The high spatial and angular resolution of TEM allows for the characterization of individual crystal defects such as dislocations, stacking faults, and point defects, as well as for more complex configurations formed by dislocations and deformation twins.1–3 However, TEM presents some drawbacks such as sample preparation, where extremely thin samples are needed, and reduced observation area. The common methods for TEM sample preparation, namely, electropolishing, ion milling, and focused ion beam (FIB), are time consuming and may even modify the microstructure by introducing point defects (for instance, foil thinning by ion milling4) or relaxing internal stresses (bulk thinning prior to electropolishing).5 Furthermore, the resulting TEM foil offers only a limited observation area (~100 × 100 μm2), making it difficult to perform the quantitative microstructural characterization required by many topics in materials science. While TEM has proven to be a very versatile tool for the characterization of crystal defects, the limitations outlined earlier have led to the development of diffraction-based methods such as x-ray diffraction (XRD),6,7 synchrotron radiation-based techniques,8–10 and 2D/3D electron backscatter diffraction (EBSD) techniques.11–15 An alternative electron microscopy technique for characterizing bulk deformed materials is electron channeling contrast imaging (ECCI).16–18 ECCI is a scanning electron microscope (SEM) technique that makes use of the strong dependence of the backscatter electron (BSE) signal on the orientation of the crystal lattice planes with respect to the incident electron beam due to the electron channeling mechanism. Interestingly, slight local distortions in the crystal lattice produced by the strain fields associated with dislocations yield a modulation in the BSE signal that can be detected by available BSE detectors. Accordingly, crystal defects such as dislocations can be imaged by ECCI in the SEM. According to the dynamical theory of electron diffraction, the optimum diffraction contrast for imaging crystal defects in ECCI is obtained by orienting the crystal exactly into Bragg condition for a selected set of diffracting lattice planes.19 However, since the early observations in 1967 by Coates of Kikuchi-like reflection patterns in the SEM, the so-called electron channeling patterns (ECPs),20 the impact of ECCI in materials science has been rather limited. This is due to the specific tilting conditions required to image crystal defects in ECCI.To date, two main ECCI setups have been developed, namely, forward-scatter electron (FSE) and BSE methods. FSE-based approaches make use of high-tilt, EBSD-type geometries to enhance the intensity of thermally diffuse (or phonon) scattered electrons. The FSE detector records the contrast created by the Kikuchi bands intersecting the detector. This configuration, originally used by Czernuszka et al.21 to image dislocations in Si and Ni3Ga, has been adapted by Wilkinson’s group16,22,23 and by Picard et al.24–26 to examine dislocations in SiGe/Si epitaxial layers22 and GaN and SiC thin films.23–26 BSE-based methods make use of electron channeling patterns to optimize the diffraction contrast of the crystal defect image. These patterns can be obtained after running the microscope at low magnification (so-called electron channeling patterns, ECPs)17,18,20 or after rocking the incoming beam on the material surface by a deflection focusing technique (so-called selected-area channeling patterns, SACPs).27–29 The main limitation of these approaches is the large observation area, typically >10 μm, and the low spatial resolution, 2 μm–5 μm. This shortcoming limits the application of ECPs/SACPs-based ECCI setups to lightly deformed coarse polycrystals and single crystals.16,29–32 Recently, we have proposed the use of EBSD to obtain the rocking parameters of the sample surface under the incoming beam (tilt and rotation angles) where the crystal is oriented into optimum diffraction contrast.33 Due to the high spatial resolution of EBSD (35 ± 5 nm34), crystal defects can be imaged in submicrometer areas of deformed bulk materials. In the last few years, we have applied the EBSD-based ECCI approach to several cases, namely, the characterization of complex deformation structures such as dislocation and nano-twin substructures in bulk deformed high-Mn TWIP steels (TWIP: twinning induced plasticity),35–39 the investigation of dislocation patterns,36,37,40 the measurement of dislocation densities in deformed alloys,41,42 and the characterization of individual dislocations and stacking faults.43,44 These studies have demonstrated the feasibility of the EBSD-based ECCI approach on the quantitative characterization of complex deformation structures in bulk materials. From a technical point of view, the ECCI technique complements two of the main electron microscopy techniques, namely, EBSD and conventional diffraction-based TEM. The present paper summarizes the experimental configuration of the EBSD-based ECCI setup and illustrates its use with several experimental studies. We also provide a physical-based model of electron channeling contrast associated with crystals defects.Experimental ConfigurationThe EBSD-based ECCI observations were carried out with a Cross Beam XB 1540 FIB instrument (Carl Zeiss SMT AG, Germany) consisting of a Gemini-type field emission gun (FEG) electron column combined with a FIB device (Orsay Physics). This instrument is equipped with a TSL OIM EBSD system. EBSD maps were measured at 15 kV acceleration voltage and a working distance of 15 mm. ECCI was conducted at acceleration voltages between 10 kV and 30 kV, using a solid-state, four-quadrant BSE detector. Working distances of 5 mm–7 mm were used with sample tilts ranging from –10° to 30°. The microscope was run in the “high current” mode, and an objective lens aperture of 120 μm was used. The experimental electron beam parameters were the following: electron beam, 7 nA–10 nA; half-beam convergence angle, 5 mrad–6 mrad; and spatial resolution, 3 nm–3.5 nm.Figure 1 shows a sketch of the EBSD-based ECCI setup. This setup is made up of two sample configurations, namely, ECCI position (a) and EBSD position (b). In the ECCI configuration, the sample surface is positioned perpendicular to the incident electron beam. The initial settings are 0° tilt/rotation angles. In the EBSD position, the sample is tilted 20° from the incident electron beam using a conventional EBSD geometry.16 We have selected a low-tilt ECCI configuration, originally introduced by Crimp’s group,31,45 instead of a high-tilt, EBSD-type configuration because of the convenient position for performing crystal orientation experiments (i.e., tilt and rotation experiments), the higher spatial resolution (~10 nm against ~50 nm43), the enhanced electron channeling contrast (~10 mrad against ~160 mrad43), and the isotropic interaction volume. In the high-tilt, EBSD-type configuration, the interaction volume is anisotropic leading to a significantly lower lateral resolution in the direction along the surface perpendicular to the tilt axis.Fig. 1(a–b): Sketch of the EBSD-based ECCI configuration. The setup is made up of two sample configurations, namely, low-tilt ECCI position (a) and EBSD position (b). (c): Illustration of the operation mode of the ECCI setup to perform ECCI under controlled diffraction conditionsFull size image

The operation mode of the EBSD-based ECCI setup is as follows: First, EBSD scans are performed on sample areas previously selected by ECCI at the low-tilt configuration. Electron channeling patterns are thereafter calculated from the crystal orientation determined by EBSD by means of computer software43,46 (Fig. 1c). The simulations were performed using a kinematical approach except those involving the width and intensity of Kikuchi bands, where a two-beam dynamical approach was instead applied.43 This code also calculates the tilt and rotation angles required to obtain the ECC image of a crystal defect under two-beam diffraction conditions for one well-defined set of lattice planes. This diffraction condition is obtained after superimposing the primary beam direction on one of the simulated Kikuchi lines far out of any principal zone axis. Subsequently, the microscope stage is positioned according to these predetermined tilt and rotation angles, and the ECC image of the selected area at optimum contrast is attained.Theoretical BackgroundElectron Scattering Mechanisms Related to Electron ChannelingThe most relevant scattering mechanisms involved in the formation of electron channeling contrast are the following: Bragg scattering, core-loss scattering, and multiple scattering.43,47 Bragg scattering is an elastic and coherent scattering event that has a large cross section for small scattering angles in the order of few degrees. Bragg scattering is the mechanism that leads to the formation of Bloch wave fields and, consequently, to conventional Bragg diffraction. Core-loss scattering processes correspond to inelastic scattering events involving large energy losses (between 10 eV and 10 keV) and very small scattering angles. These scattering processes lead to the formation of electromagnetic radiation (x-rays, light), plasmons, or secondary electrons.48 Finally, and very important for backscattering, inelastic and phonon scattering events combine to a chain of scattering events comprised as multiple scattering. This is characterized by large scattering angles and a wide energy distribution of electrons. Multiple scattering is commonly considered the main scattering event in the calculation of the spatial and spectral energy distribution of backscattered electrons in Monte-Carlo electron trajectory simulation programs.49

The formation of the channeling contrast can be understood as follows:43 The primary beam electrons that interact with the surface crystal are coherently (Bragg) scattered by the atoms of the crystal lattice and form a lattice-coherent electron wave. Depending on the orientation of the crystal lattice with respect to the primary electron beam direction, the coherent electron field contains density maxima at atom positions or further away, resulting in a different number of inelastic scattering events by core-loss scattering. The resulting low-energy electrons may, by multiple scattering, eventually leave the sample and be recorded as backscattered electrons with a broad energy distribution but with an orientation-dependent intensity.Electron Channeling Contrast Associated with Lattice DefectsThe electron channeling contrast associated with crystalline defects has been simulated by different approaches, namely, the numerical integration of two-beam Howie–Whelan equations50 and Bloch wave-based models.19,51,52 Recently, we have made use of a two-beam Bloch wave approach that captures the main features observed in the experimental contrast associated with stacking faults and dislocations.43 In this model, the orientation contribution to the total backscatter intensity can be written as:43,48

$$ \Updelta \eta = \frac{{N\sigma_{\text{B}} }}{2\pi }\xi_{0}^{'} \left( { - \frac{{\left( {w + \xi_{0}^{'} /\xi_{g}^{'} } \right)}}{{1 + w^{2} - \left( {\xi_{0}^{'} /\xi_{g}^{'} } \right)^{2} }} + \frac{w}{{1 + w^{2} + \left[ {(1 + w)\left( {\xi_{0}^{'} /\xi_{g}^{'} } \right)^{2} } \right]}}} \right) $$

(1)

where N is the number of atoms per unit volume, σ

B is the cross section for backscattering, and w = s·ξ

g with ξ

g being the extinction distance and s the excitation error that indicates the angular deviation of the primary electron beam wave vector, k

0, from the exact Bragg position of the lattice planes with reciprocal lattice vector g. As example, we have calculated the backscattering intensity associated with a stacking fault (SF) laying in a depth z

SF below the surface (Fig. 2a).43 In this calculation, we have assumed that the atoms below the SF are shifted by exactly half an atomic plane distance. This figure reveals that the intensity is highest close to the surface, which is associated with the intersection of the SF with the sample surface. The intensity profile is subsequently attenuated with characteristic but weak oscillations. The frequency of these oscillations, measured with respect to the distance, z, to the entry surface, corresponds to the effective extinction distances obtained under the current electron beam conditions, $ \xi_{g}^{\text{eff}} $. Figure 2b shows an ECC image of a stacking fault in a high-Mn steel. The oscillations are well visible. From the number of oscillations, it is therefore possible to estimate the observation depth of crystal defects: For g = (111), E = 20 keV, and an fcc-iron crystal structure, it yields ξ

= 13 nm. In Fig. 2b, three bright lines are visible. Accordingly, the total visibility of the stacking fault is about five times this value, ~60 nm–70 nm. Interestingly, similar probe depths have been estimated from theoretical approaches in Fe alloys.42

Fig. 2Comparison of calculated (a) and experimental electron channeling contrast of a stacking fault. The calculation was performed by means of a two-beam Bloch wave approach. The experimental image was taken from a stacking fault in a high-Mn steel, observed under g = (111) (g: diffraction vector)43

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Application Examples of ECCI to the Characterization of Deformation StructuresIn recent years, we have successfully applied the EBSD-based ECCI approach to the characterization of several microstructural features, namely, single dislocations and stacking faults,42,43 nanotwins,35–37,43 and complex dislocation configurations, such as cells, cell blocks, Taylor lattices, and microbands.35–38,40 It has been also proved as a valuable technique to measure dislocation densities.41–43 Figure 3 illustrates several application examples of the ECCI technique (details can be found in Refs. 40,41,43). In this section, we provide further details of two particular application examples, namely, the characterization of single dislocations and dislocation substructures.Fig. 3Experimental application examples of the ECCI technique: (a) characterization of single dislocations, (b) stacking faults, (c–d) dislocation substructures, and (e) nanotwins; (f) measurement of dislocation densities (details can be found in Refs. 40, 41, 43)Full size image

Characterization of Single DislocationsSingle dislocations can be clearly imaged by ECCI at suitable diffraction conditions. ECC images of dislocations correspond to projected dislocation lines within the interaction probe depth, typically <100 nm. The optimum dislocation contrast is attained by orienting the crystal into the channeling condition, which is close to the exact Bragg condition but with a small negative value of w (Eq. 1).19,43 In this diffraction condition, dislocations appear as bright lines on a dark background. Dislocation contrast becomes weaker when deviating from the exact channeling conditions. Figure 4a illustrates, as an example, an ECC image of dislocations taken after orienting the crystal into Bragg condition in a lightly deformed FeSi alloy with bcc crystal structure.43 Three different dislocation sets are visible labeled as (i)–(iii). Two of them, namely, (i) and (ii), appear as straight long lines. Assuming that long dislocation segments in bcc structures usually contain a screw character with a 〈111〉-type Burgers vector, we can index the corresponding Burgers vectors and line directions, provided that their inclination angles with respect to the surface are known. Under the present diffraction conditions and acceleration voltage, the probe depth is ~80 nm. Assuming that the visible dislocations intersect the sample surface and considering the corresponding projected dislocation lines, the inclination angles for the dislocation lines (i) and (ii) are ~10° and 15°–20°, respectively. These angles correspond to the inclination angles of the [1$ \bar{1} $1] and [$ \bar{1} $11] directions on the line traces (i) and (ii), respectively (Fig. 4b). The dislocations labeled as (iii) appear as small bright dots. These dislocations lie almost exactly perpendicular to the sample surface and clearly correspond to screw dislocations with [111] dislocation line directions and Burgers vectors. The dislocation lines (iv), although almost invisible, can be clearly attributed to [11$ \bar{1} $] line directions and Burgers vectors. As in ECCI, the g·b = 0 invisibility criterion holds32,53 (g: diffraction vector, b: Burgers vector), and we can associate these dislocations with those fulfilling the invisibility criterion.Fig. 4(a): ECC image of dislocations taken after orienting the crystal into Bragg condition in a lightly deformed FeSi alloy.43 (b): Stereographic projection of the orientation indicating the three 〈111〉 directions that correspond to line direction and Burgers vectors of the clearly visible dislocationsFull size image

Measurement of Dislocation DensitiesAverage dislocation densities can be determined by counting the number of dislocations intersecting the observation area or using the relationship ρ = 2 N/Lt, where N is the number of dislocation lines intersecting a grid of total line length L on the corresponding ECC image and t is the probe depth. The latter is a standard relationship used in TEM when individual dislocations can be clearly distinguished.3,54 The probe depth for imaging dislocations in ECCI has been estimated to be about 5ξ

(calculations were performed using a dynamical diffraction approach by Wilkinson et al.22). Figure 5 shows the average dislocation densities measured by ECCI in two to three grains of a Fe-3 wt.% Si alloy deformed to a macroscopic stress level of 500 MPa under two different diffraction conditions, namely, two-beam conditions with one set of hkl planes at Bragg orientation and three-beam conditions with two sets of hkl planes out-of-Bragg.42 The higher average density obtained by the second diffraction condition is ascribed to the larger number of excited reflectors, which provides, due to a channeling mechanism, a higher amount of visible dislocations. This figure also includes the average dislocation density determined by bright-field TEM in a sample deformed to the same macroscopic stress level. The average dislocation density was measured from TEM micrographs using the same method as in the present work.55 These findings reveal that the average dislocation density estimated from ECCI is in the same range as that determined from BF-TEM (BF: bright-field). In particular, the average dislocation density measured by ECCI in two-beam conditions is about two times smaller than that determined by BF-TEM. This result is well known from corresponding TEM-based estimates.3 Figure 5 also reveals that three-beam diffraction conditions with two sets of hkl planes out-of-Bragg provide a better estimation of the average dislocation density by ECCI, which is close to that obtained by bright-field TEM.Fig. 5Average dislocation densities determined in a Fe-3 wt.% Si alloy tensile deformed to a macroscopic stress of 500 MPa by two different electron microscopy techniques, namely, ECCI and TEM. Average dislocation densities were measured by ECCI under two different diffraction conditions: two-beam conditions with one set of hkl planes in the Bragg orientation and three-beam diffraction conditions with two sets of hkl planes out-of-Bragg (BF: bright field)Full size image

Characterization of Dislocation Substructures in Deformed Bulk MaterialsThe investigation of dislocation substructures has played an essential role in many topics in materials science in the last 50 years, such as strain hardening, polycrystal and crystal plasticity, internal stresses, and recrystallization. Due to the low misorientation of the dislocation boundaries involved (typically <5°), its investigation has been mainly carried out by TEM. This has significantly limited the number of studies devoted to the quantitative characterization of dislocation patterns. In particular, this limitation has a pronounced influence on the understanding of polycrystal plasticity and strain hardening. According to dislocation mean free path-based theories of strain hardening, not only the dislocation pattern spacing must be evaluated but also its crystallographic orientation dependence with respect to the macroscopic stress as well.56 We have recently explored the characteristics of the ECCI technique into the investigation of the evolution of dislocation substructures upon deformation in high-Mn steels.35–38,40 Dislocation configurations such as highly dense dislocation walls, cells, cell blocks, Taylor lattices, and microbands can be readily imaged and identified. Figure 6 shows an example of the crystallographic orientation dependence of the dislocation substructure in a Fe-30.5Mn-2.1Al-1.2C (wt.%) high-Mn steel tensile strained to 0.2 true strain/950 MPa at an initial strain rate of 5 × 10−4 s−1.37 Due to the wide field of view of the SEM, the evaluation of a large number of grains by ECCI is not as time consuming as in TEM. In this example, the crystallographic orientation dependence of the dislocation patterns, namely, cell blocks (CBs) and cells (DCs), was investigated in 30 individual grains with a representative orientation distribution that matches the overall crystallographic texture. Green and red dots correspond to CBs (green) and DCs (red), respectively. The data indicate that cell patterning has strong crystal orientation dependence. Equiaxed cells are only developed in grains oriented close to 〈001〉//TA directions (Fig. 6c). The remaining analyzed grains contain a well-developed cell block structure (Fig. 6b). The shape of such structures depends on the number of active sip systems. The characterization of these dislocations provides new insights into the influence of alloying elements on the underlying dislocation mechanisms and, accordingly, into the strain hardening mechanisms.37

Fig. 6(a): Inverse pole figure along the tensile axis direction showing experimental grain orientations of a Fe-30.5Mn-2.1Al-1.2C (wt.%) high-Mn steel tensile strained to 0.2 true strain/950 MPa (red dots: cells (DCs); green dots: cell blocks (CBs)), (b and c) ECCI images of (b) CBs and (c) DCs, respectively37

Full size image

ConclusionsThe present article reviews the most relevant experimental and theoretical aspects associated with ECCI. The following conclusions can be drawn:

The coupling of ECCI with EBSD provides an efficient and fast approach to perform ECCI of crystal defects, such as dislocations and stacking faults, under controlled diffraction conditions with enhanced contrast. From a technical point of view, the ECCI technique complements two of the main electron microscopy techniques, namely, EBSD and conventional diffraction-based TEM.

We make use of a two-beam Bloch wave approach to interpret the channeling contrast associated with crystal defects. The approach captures the main features observed in the experimental contrast associated with stacking faults and dislocations.

We provide several application examples of the EBSD-based ECCI approach on microstructure characterization, namely, characterization of single dislocations, measurement of dislocation densities, and characterization of dislocation substructures in deformed bulk materials.

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D. Kuhlmann-Wilsdorf, Dislocations in Solids. vol. 11, eds. F.R.N. Nabarro and M.S. Duesbery (2002).Download referencesAcknowledgementsThe authors acknowledge the financial support by the German Research Foundation in the framework of SFB 761 “steel ab initio.” The authors also acknowledge intense discussions on the two-beam dynamical theory with Dr. A. Winkelmann (MPI Halle) and on the general application of ECCI with Prof. M. Crimp (Michigan State University).Author informationAuthors and AffiliationsMax-Planck-Institut für Eisenforschung, 40237, Dusseldorf, GermanyI. Gutierrez-Urrutia, S. Zaefferer & D. RaabeAuthorsI. Gutierrez-UrrutiaView author publicationsYou can also search for this author in

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I. Gutierrez-Urrutia.Rights and permissionsReprints and permissionsAbout this articleCite this articleGutierrez-Urrutia, I., Zaefferer, S. & Raabe, D. Coupling of Electron Channeling with EBSD: Toward the Quantitative Characterization of Deformation Structures in the SEM.

JOM 65, 1229–1236 (2013). https://doi.org/10.1007/s11837-013-0678-0Download citationReceived: 15 May 2013Accepted: 17 June 2013Published: 16 July 2013Issue Date: September 2013DOI: https://doi.org/10.1007/s11837-013-0678-0Share this articleAnyone you share the following link with will be able to read this content:Get shareable linkSorry, a shareable link is not currently available for this article.Copy to clipboard

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KeywordsCrystal DefectDislocation SubstructureTWIP SteelDiffraction ConditionIncident Electron Beam

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