梯度结构铜铝合金的室温加工硬化行为

根据Hall-Petch公式可推测具有均质结构金属材料的力学性能[1, 2] 纳米材料[3, 4]的强度和硬度都非常高，但是很难生成极其细化的晶胞以存储更多的位错，因此往往很早就失稳进入了颈缩阶段 粗晶材料的均匀延伸性好，加工硬化能力强，但是其强度较低 近年来，朱运田、卢柯等[5~10]提出了异质结构材料的概念，即在金属材料中引入一个异质区域以同时得到纳米材料的高强度和粗晶材料的高延展性，从而实现金属材料强度-塑性的良好匹配

异质结构材料，分为异质层状结构材料、梯度结构材料、核壳结构材料、双相结构材料以及多模态结构材料[9] 在异质结构材料的塑性变形过程中，异质区的变形是不均匀的 在软质区产生背应力[11]，在硬质区产生前应力[12]，其共同作用产生异质变形诱导强化，增强应变硬化，有助于提高异质结构材料的全局屈服强度并保持延展性 异质变形诱导应力和应变硬化的微观机理研究尚处于前沿阶段，梯度结构层中界面影响区的形成和演变尚不十分清楚 科研人员以数字图像相关法开展塑性变形过程的力学研究[13, 14]，为梯度结构材料的力学行为的解释和变形机理的阐述提供了直观的表征支持 本文研究Cu-4.5%(质量分数)Al合金梯度结构材料的加工硬化行为，探究室温下准静态拉伸优异的屈服强度-均匀延伸性组合的微观机理，通过表征表面机械研磨处理(Surface mechanical attrition treatment，SMAT)后合金板的梯度层由表层到芯部形成的微观组织的差异并结合异质变形诱导强化等理论，系统说明高强度、高塑性潜在的微观塑性变形的力学行为，以较低的加工成本换取极佳的强塑性匹配，探索梯度结构材料的工业化生产和应用方式

1 实验方法

实验用铜铝合金为纯铜中固溶了4.5%的Al元素 将热锻和冷锻成厚度为4 mm的铜铝合金板放置在温度为923 K的管式真空炉中退火2 h，制备出均质结构的粗晶铜铝合金板 再将其在液氮下表面机械研磨[7] 2 min，制备出具有双面梯度结构的铜铝合金板

使用型号为SHIMADZU Universal Tester的力学试验机进行单轴拉伸测试，应变速率为5×10-4 s-1 拉伸试样标距部分的尺寸为15 mm×5 mm×4 mm(图1a)，用240#到2000#砂纸将打磨电火花线切割残留痕迹的截面 数字图像相关法(Digital image correlation，DIC)与拉伸试验同步进行，图1b给出了斑点与分析概览



图1拉伸样品尺寸、数字图像相关法斑点及观测区域以及微观表征区域

Fig.1Size of tensile text samples (a), DIC speckle and observation area (b) and the area of microscopic characterization (c)

用仪器型号Optic Microscope (OM) Leica DM 5000的光学显微镜对样品进行微观表征 制备所用样品时，依次用400#、800#、2000#、5000#砂纸将其打磨，然后将研磨抛光的表面腐蚀余额1 min，腐蚀液是用氯化铁5 g+50 mL稀盐酸+100 mL水配制；将样品研磨抛光后，再用离子抛光仪抛光以去除机械损伤层用于扫描电镜观察，所用仪器是型号为FE-SEM, NOVA Nano SEM 450并带有电子背散射衍射(Electron backscattered diffraction，EBSD)信号探头的扫描电镜，其操作电压为250 KV；用于透射电镜(Transmission electron microscope，TEM)观察的样品，需要沿深度取宽约1000 μm的梯度截面，用2000#砂纸将其机械减薄至50 um，再镶嵌在双联铜环上离子减薄使其有良好的可观察薄区，离子减薄参数为5 keV 8°处理2 h、4 keV 5°处理0.5 h、2.5 keV 3°处理2 h，透射电镜的型号为JEM-2100 Plus TEM，操作电压为200 kV 微观表征区域的选区示意图，在图1c中给出

2 结果和讨论

Cu-4.5%Al的合金板在液氮温度下SMAT处理2 min，为了维持材料整体的连续性，在其表面层产生了沿深度呈梯度递减分布的几何必要位错(Geometrically necessary dislocations，GNDs)(图2a) 这符合金属材料的塑性变形理论，统计同一深度的平均局部取向差(Kernel average misorientation，KAM)值并基于应变梯度理论[15, 16]，其GND的值为

ρGND=2KAMaveμb(1)

式中ρGND是GNDs的值，KAMave为局部取向差的平均值，μ为EBSD的扫描步长，b为位错的柏氏矢量



图2梯度层的微观结构

Fig.2Microstructure of gradient layer (a) KAM map of the depth ~300 μm depth, characterized by EBSD, (b) The GND distribution map with depth, calculated by the average KAM statistics in (a), (c~e) TEM bright field image of ~30 μm, ~153 μm, ~246 μm depth, respectively

图2b表明，统计出来的GND密度沿深度呈梯度降低，在~50 μm处趋势放缓最终在~250 μm深度处趋于平稳 这表明，处理后的样品加工影响深度可能比~250 μm更大，其最表面的GND约为6.8×1013 m-2，芯部区域的GND约为2.4×1013 m-2，表层和芯部GND的差距约2倍 该成分的层错能(Stacking fault energy，SFE)很低只有12 mJ/m2(纯铜的层错能有78 mJ/m2)，加工温度对微观结构的影响各不相同[17]，生成位错的临界分切应力随温度降低而增大，生成层错(Stacking faults，SFs)的临界分切应力随温度的降低而降低[18] 在此加工温度下占主导地位的变形机制由层错取代了位错[19, 20]，从而促进了{111}晶面族层错的形成 用TEM表征了~30 μm、~153 μm、~246 μm 3个深度，观察到极其细小的层错密度沿深度梯度递减，而梯度层中的位错也从靠近表层的大量林位错(图2c)演变成靠近芯部缠结的位错(图2d)，在芯部甚至能观察到单个分布的位错(图2e)，证实了微观结构随深度变化呈现缺陷密度梯度特征

对SMAT处理前后的拉伸试样进行准静态单轴拉伸测试(图3a)，短时SMAT处理将Cu-4.5%Al合金的屈服强度从~90 MPa提高到~170 MPa，均匀延伸率从~54%降低至~45%，在真应力应变曲线(图3a)背景的金相上可观察到加工后腐蚀液对位错滑移带的点蚀痕迹[21, 22] 加工后屈服强度的提高可以归因于层错，对梯度层~30 μm深度处捕捉了一张放大数倍的TEM明场像(图3b)，可以发现a2<110>全位错沿{111}面滑移会分解成2个a6<211>肖克利不全位错 由于Cu-4.5%Al合金的层错能较低，两个不全位错之间的层错宽度增大，因此异号不全位错很难碰到而湮灭，因此在梯度层留下了大量的层错(图3c) 在相同晶面上产生的层错堆叠在一起，形成了有一定原子层宽度的纳米孪晶[23, 24](Nano twins，NTs)，从面缺陷演变成了体缺陷；在不同晶面上产生的层错和层错之间发生反应，两个拓展位错在各自滑移面相向移动 当每个拓展位错中的一个肖克利不全位错运动到滑移面的交截线时，位错反应产生了一个新的纯刃型位错，将这两个可动的拓展位错在此处固定住，此时的混合位错组态称为面角位错(Lomer-cottrell dislocation，L-C dislocation)(图3b)，由三个不全位错和两片层错所构成 对于面心立方晶体(Face centered cubic，FCC)结构的金属加工硬化起重大作用，能有效钉扎和塞积位错 由此可知，屈服强度大大提高的直接原因，可归于低层错能合金在液氮下塑性应变产生的由层错组成的微观结构以及层错对位错的塞积[25, 26]



图3SMAT处理前后的真应力-应变曲线对比(背景为梯度层的金相显微图)、在梯度层~30 μm深度处的TEM明场像以及拓展位错的形成-全位错分解为不全位错

Fig.3Comparison of true stress-strain curves before and after SMAT treatment (a, metallographic image with gradient layer in the background); TEM bright field image at gradient layer ~30 μm depth (b); Extended dislocation formation-decomposition of full dislocations into partial dislocations (c)

为此，图4给出了准静态单轴拉伸测试试样扫描的断口形貌，图4a和4b分别给出了梯度结构试样和粗晶试样的断口 断口芯部的韧窝很大且均匀(图4a2、4b2)，表明芯部的塑性变形充分且稳定 从断口边缘可观察到，粗晶试样的表层和芯部特征相同，而梯度试样的特征就有所不同 断口边缘从表及芯存在过渡的形貌(图4a3)，表明材料变形的非均质的特性 梯度层虽然有韧窝，但不均匀且较浅(图4a4)，说明单一梯度层在颈缩阶段前SMAT处理的影响 晶胞的加工硬化的潜力降低了[27]，不能继续均匀变形，失去了和芯部同步塑性变形的能力 梯度层失稳进入颈缩阶段，使试样提前进入颈缩阶段[28]



图4梯度结构试样和粗晶试样的断口特征

Fig.4Observation of fracture features of gradient-structure sample (a) and coarse-grain sample (b)

SMAT处理前后拉伸试样断口形貌的差异表明，剪切带过早的形核并集中使梯度层的加工硬化不能保持，为此分析了拉伸过程中剪切带的演变 结果表明，材料的加工硬化并不是整体同步进行，而是加工硬化更强的区域将应力分散到加工硬化较弱的区域[13, 29] 观察了粗晶试样DIC从屈服后到~30%应变阶段，如图5所示，剪切带的形核在边缘产生(图5a的0.78%应变点)，并不是在边缘立即深化导致微裂纹产生，加工硬化将应力分散到芯部(图5a的4.54%应变点) 从图5a可见，早期沿Y轴产生的剪切带与同一Y轴高度上的应变大致相近，较突出的应变集中区通过应力分散将附近应变较小的区域加工硬化(见图5a的29.95%应变点)，最终在颈缩阶段前试样整体达到加工硬化能力的最大值，不能继续加工硬化保持均匀的塑性变形 此时剪切带从试样边缘区域深化产生微裂纹，试样进入不稳定变形的颈缩阶段，直至断裂 依据均匀塑性变形时体积恒定原则，沿Y轴正应变较大的区域在X轴收缩的负应变更多 在全局上，均匀分布的剪切带，是试样保持较好均匀塑性变形必不可少的



图5粗晶试样的剪切带演变

Fig.5Evolution of shear bands of CG sample (a) plastic strain distribution along the Y-axis at each global strain; (b) plastic strain distribution along the X-axis at each global strain

具有梯度结构试样(图6)，其剪切带形核也从边缘开始，而梯度层中的剪切带形核比粗晶试样更明显 随着塑性变形的进行，表层内很集中的应变区(图6a)从表层传递到芯部 从图6a中9.61%应变图像可见，芯部的变形量比表层多，因为试样刚从屈服后进入塑性变形阶段，梯度层内的缺陷密度梯度较高，表层与芯部之间的弹性-塑性阶段相互作用引起的长程内应力导致了过渡界面的高应力集中[13, 14] 芯部内较高的加工硬化使强度提高从而降低了流应力差异，粗晶芯部稳定了剪切带并阻止了向芯部传递 这导致在梯度层形核的剪切带先向标距段未加工硬化的部分传递(图6a的0.34%应变点)，标距段全局都加工硬化后梯度层的剪切带核才开始向芯部扩散，造成应变局域化(图6a的9.61%应变点) 其微观机理是，表层与芯部存在应变差异 为了维持材料的整体连续性，在芯部开动的弗兰克-里德位错源产生了大量的几何必要位错并向梯度层发射(图7a) 但是，梯度层靠近表层的区域有高密度的几何必要位错，因此从芯部发射过来的几何必要位错在梯度层的过渡界面塞积[12]，有序排列塞积的几何必要位错使靠近表层的梯度层产生前应力 前应力使梯度层在承受的流应力之外附加了来自芯部的应力，梯度层也反馈背应力施加在芯部 芯部加工硬化的强度加上背应力，可分担芯部承受的部分流应力 因此，芯部变相地被背应力强化了 梯度层与芯部之间的异质结构诱导强化行为使剪切带在整个标距部分均匀分布，芯部发生的几何必要位错弥补了一部分表层晶胞产生位错能力的不足，从而使试样在塑性变形过程中在保持了较高的均匀延伸性的同时使屈服强度大大提高



图6梯度结构试样的剪切带演变

Fig.6Evolution of shear bands of GS sample (a) plastic strain distribution along the Y-axis at each global strain; (b) plastic strain distribution along the X-axis at each global strain



图7GND塞积的示意图,在软区中诱导背应力，在硬区中诱导前应力[12]以及有限元模拟平均应变为0.2%应力和应变沿厚度方向的变化[30]

Fig.7Schematics of a GND pile-up, inducing back stress in the soft domain, which in turn induces forward stress in the hard domain[12] (a) andstress and strain variations along thickness direction on 0.2% average strain by finite element simulation (FEM)[30] (b)

异质结构诱导的强化行为，不只是在强-弱的过渡界面上通过前-背应力实现的 双面梯度结构的强-弱-强的双面结构分布推迟了应变局域化[30](如图7b)，有助于应力均匀分布和避免应变集中 应力应变的均匀分布，使试样的延展性提高

3 结论

将Cu-4.5%Al合金板材进行机械研磨后，在其表面形成了厚度约为250 μm的梯度结构层，从表层到芯部的位错结构分别为林位错、缠结的位错、单个的位错，层错结构在靠近表层区域(~30 μm)生成了纳米孪晶 Cu-4.5%Al合金的拉伸试样经准静态单轴拉伸后，其屈服强度提高了~80 MPa，均匀延伸率降低9% 屈服强度提高的直接原因，是梯度层中的纳米孪晶和面角位错对可动位错的存储、塞积的作用；间接原因，是异质结构诱导强化产生的额外的强化作用 均匀延伸性降低程度很低，可以归因于双面梯度结构对剪切带均匀分布的贡献，梯度层与芯部之间的异质结构诱导强化行为使剪切带在整个标距段均匀分布，在芯部发生的几何必要位错弥补了一部分表层晶胞产生位错能力的不足，从而使试样在塑性变形过程中在保持了一定均匀延伸性的同时大大提高了屈服强度、推迟了应变局域化，从而使标距段继续加工硬化，避免了在较早阶段失稳进入颈缩阶段

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The microstructure and mechanical properties of extruded Mg-alloy of Mg-1Al-0.4Ca-0.5Mn-0.2Zn (mass fraction, %) were systematically investigated. As indicated by the results, the incomplete dynamic recrystallization occurred for the alloys extruded at 260℃ (denoted as AXMZ1000-260) and 290°C (AXMZ1000-290) with recrystallized grain sizes of 0.75 μm and 1.2 μm, respectively. The two alloys have high-density G.P. regions and spherical nano-phases, which can effectively inhibit the dislocation motion and provide abundant nucleation sites for dynamic recrystallization. Moreover, the nano-phases precipitated along grain boundaries can restrain the migration of grain boundary and restrict the growth of DRXed grains, which results in the ultrafine grains with a size of 0.75 μm in AXMZ1000-260 alloy. The strength of the alloy decreases with the increase of extrusion temperature, and the change of elongation is not obvious. The yield strength and elongation of alloys extruded at 260℃ and 290℃ are approximately 322 MPa and 343 MPa, as well as 13.4% and 13%, respectively. The dynamic precipitation and recovery process are promoted by the increasing extrusion temperature, and a high-density G.P. zones and spherical nano-phases are accumulated in the alloy. At the same time, many dislocations are transformed into LAGBs by dynamic recovery, and the unDRXed areas are subdivided into dense lamellar subgrains. The nano-phases and LAGBs can effectively hinder the newly generated dislocation motion, which is the major reason that the alloy extruded at 290℃ still have a high yield strength and the change of ductility is not obvious. Furthermore, TEM observations show that the pinning effect of G.P. zones can impede the dynamic recovery to certain extent, resulting in a high number of residual dislocations in the alloy, which is conducive to the improvement of the yield strength.

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The semi-solid ZCuSn10 alloy billets were prepared with strain induced melt activated (SIMA) method involved with hot rolling and reheating process. The microstructure evolution process and spheroidizing mechanism of α(Cu) phase were studied by means of optical microscope, scanning electron microscope and image analysis software. The results show that when a hot rolled ZCuSn10 copper alloy billet with a deformation rate 16% was reheated at 930℃, of which the semi-solid primary phase spheroidized gradually with the increasing holding time; while the average grain size of the copper alloy decreases firstly with time from 68.24 μm for 8 min to 62.31 μm for 10 min and then increases to 71.09 μm for 25 min; the liquid fraction increases from 18.14% for 8 min to 25.32% for 25 min; the shape factor decreases firstly with time from 2.91 for 8 min to 1.67 for 15 min and then increases to 2.43 for 25 min. The alloy exhibits the best semi-solid microstructure for 15 min holding with an average grain size 65.64 μm, a liquid fraction 23.66% and a shape factor 1.67. The microstructure evolution mechanism involves with merge of grains and growth as well as atom diffusion leading to grain growth and spheroidization.

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采用热轧与重熔的SIMA法制备了ZCuSn10铜合金半固态坯料 先通过多向热轧对ZCuSn10铜合金坯料进行预变形, 然后对其进行半固态温度区间保温不同时间的二次加热处理 使用光学显微镜、扫描电镜、能谱及图像分析软件等手段研究了坯料二次加热过程中微观组织的演化, 并分析了α(Cu)球化组织的形成机制 结果表明：在热轧变形量为16%、加热温度为930℃时, 随着保温时间的延长半固态ZCuSn10铜合金坯料初生α(Cu)逐渐发生球化, 平均晶粒直径先减小后增大, 由二次加热保温8 min的68.24 μm先减小至10 min的62.31 μm然后增大至25 min的71.09 μm; 保温10 min时平均晶粒直径最小, 液相率由保温8 min时的18.14%逐渐增加到保温25 min时的25.32%; 形状因子随着保温时间的延长先减小后增加, 由保温 8 min时的2.91先减小至15 min时的1.67, 然后增大到保温25 min时的2.43 保温15 min时的半固态组织最优, 其平均晶粒直径为65.64 μm、液相率为23.66%、形状因子为1.67 在半固态铜合金二次加热过程中, 组织演变的主要机制是加热前期的晶粒合并长大和液相增加后的原子扩散导致的晶粒长大并球化

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