纯镁和镁合金的阻尼及微塑变行为研究
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摘要
本文采用热挤压、等通道角挤压(ECAP)和退火处理来改变纯镁的微观组织结构，研究晶粒尺寸和织构等对阻尼和微塑变的影响机制以及它们之间的内在联系，并在纯镁中添加合金元素来研究镁合金的阻尼和微塑变机制。利用光学显微镜(OM)和透射电子显微镜(TEM)等方法观察微观组织变化；通过中子衍射和电子背散射(EBSD)分析织构的演变；采用动态机械分析仪(DMA)研究阻尼性能随应变的变化规律；采用循环拉伸的方法研究微小变形过程中滞弹性应变、摩擦应力和正切弹性模量等的变化规律。揭示了纯镁和镁合金的阻尼和微塑变行为，为新型高阻尼镁合金、耐疲劳镁合金和尺寸稳定镁合金的开发和应用奠定了良好的基础。
挤压后纯镁的平均晶粒尺寸为68μm，随着ECAP变形道次的增加，再结晶程度逐渐提高，晶粒逐渐细化，在250°C下ECAP变形4道次纯镁的平均晶粒尺寸为6μm。挤压态纯镁具有基面平行于挤压方向的织构。随着变形道次的增加，基面极点以TD和ED为轴逐步发生倾转，最终形成基面极点分别偏离ND和TD大约40o和65o的织构，而基面平行于挤压方向的织构逐渐弱化，导致沿挤压方向变形的基面滑移Schmid因子大幅度提高。退火后纯镁的晶粒尺寸逐渐增大，而织构基本保持不变。
纯镁与塑性变形相关的阻尼和微塑变行为可用同样的位错机制来解释，它们有着相同的物理本质。塑性阻尼的两个阶段与循环拉伸微塑变的两个阶段相对应：第一阶段对应塑性应变小于2×10~(-4)(总应变约为8×10~(-4))的区域，Schmid因子较大晶粒内部位错从钉扎点上脱钉并在基面滑移，可动性较大，激活体积较大，加工硬化指数较小；当塑性应变高于2×10~(-4)时，由于位错在同一滑移面上运动而发生缠结和堆积，可动位错密度降低，材料硬化，所以具有较大的加工硬化指数和较小的位错滑移激活体积，这时微塑变进入第二阶段。
在应变小于第一临界应变振幅(1×10~(-4)左右)时，纯镁中位错在弱钉扎点间摆动，加载卸载曲线基本重合为一条直线，此时阻尼性能Q_0~(-1)与应变振幅无关；随应变的增大，位错从弱钉扎点上脱钉，产生应力-应变滞后环，滞弹性应变迅速增加，而正切弹性模量快速降低，此时阻尼性能Q_h~(-1)随应变振幅逐渐增大。这两个滞弹性阶段可以用G-L位错模型来解释；当高于第二临界应变振幅(5×10~(-4)左右)后，G-L曲线偏离直线，基面位错从强钉扎点上脱钉，发生微小塑性变形，阻尼性能Q_p~(-1)快速提高；当高于第三临界应变振幅(9×10~(-4)左右)后，由于位错的缠结和堆积，摩擦应力逐渐增大，而损失模量的增加速度变缓。后两个微小塑性变形阶段需要用微塑变位错模型来解释。随着晶粒尺寸或基面滑移Schmid因子的增大，纯镁的阻尼性能、滞弹性应变和损失弹性模量增大，而正切弹性模量、摩擦应力和背应力降低。
选取在镁中固溶的Al和不固溶的Si两类合金元素来制备Mg-Al和Mg-Si合金，并以纯镁作为对比研究合金元素的添加对镁的阻尼和微塑变影响规律。添加完全固溶的1%Al元素会极大的降低位错线上弱钉扎点间距，从而提高镁合金在较小应变下的正切弹性模量，摩擦应力和背应力，但同时也会导致阻尼性能的急剧降低；添加基本不固溶的Si元素不会降低镁的低应变阻尼性能，但会生成Mg2Si第二相，明显的减小铸态镁合金的晶粒尺寸，严重阻碍脱钉后的位错运动，导致较大应变下阻尼性能的降低，同时提高正切弹性模量，摩擦应力和背应力。
在位错增殖的临界应变振幅时，假设位错在强钉扎点钉扎下循环过程中所扫过的面积为圆形。根据功能原理可计算出此临界应变下纯镁和镁合金中的可动位错密度以及位错线上强弱钉扎点间的平均距离，变形后纯镁和镁合金中的可动位错密度为10~(12)m~(-2)数量级，退火态和铸态材料中可动位错密度为10~(10)-10~(11)m~(-2)数量级。
Hot extrusion, equal channel angular pressing (ECAP) and subsequentannealing treatment were performed on pure Mg to modify the microstructure. Theeffects of grain size and texture of pure Mg on microplastic deformation, dampingcapacity and the relationship between them were studied in detail. Two kinds ofalloy elements were added to Mg to study the microplasticity and damping capacityof Mg alloys. The microstructures characteristics of pure Mg and Mg alloys wereobserved by OM and TEM. The texture evolutions were analyzed by neutrondiffraction and EBSD. The strain dependence of damping capacities were studied byDMA. The developments of anelastic strain, friction stress and secant elasticmodulus, etc. in microstrain region were studied by cyclic loading-unloading tests.This paper revealed the damping and microplastic behaviors of pure Mg and Mgalloys, established a good foundation to develop high damping capacity, high fatigueand high dimensional stability Mg based materials.
The average grain size of pure Mg after extrusion was about68μm. Withincreasing the ECAP passes, the degree of recrystallization was increased, whichresulted in the refinement of grain size. The grain size of pure Mg after ECAPprocessing for4passes at the temperature of250°C was refined to6μm. Theas-extruded Mg exhibits a texture with {0001}1010nearly parallel to theextrusion direction (ED). With increasing the ECAP passes, the basal poles rotatedaround transverse direction (TD) and ED axis and formed a maximum componentlocating at about40ofrom normal direction (ND) and65ofrom TD. The texturecomponent with basal planes parallel to ED becomes much weaker. The textureevolution after ECAP led to a increase of the Schmid factor for basal slip towardsED. The grain size was gradually increased, but the texture almost kept stable afterannealing treatments.
The microplasticity and damping related to plasticity of pure Mg have the samephysical mechanism and can be explain by the same dislocation model. The twostages of plastic damping and microplasticity correspond to each other. The firststage relates to the plastic strain below2×10~(-4)(the total strain of about8-9×10~(-4)),the basal dislocations in favorable oriented grains break away from strong pinningpoints and slide on the basal planes, which results in the larger volume activationsbut smaller hardening exponents; the microplastic deformation process enters intothe second region when the plastic strain is above2×10~(-4), the tangle and piled-upof dislocations lead to the decrease of mobile dislocations, the materials are characterized by the larger hardening exponents but smaller volume activations.
When the strain is smaller than the first critical strain amplitude (about1×10-4),the materials almost deform in an elastic manner. At this time the dislocationsreversibly motion between the weak pinning points, the loading-unloading curvesalmost coincide with the elastic line, and the damping capacityQ1
0is amplitudeindependent; the dislocations unpin from weak pinning points with the increase ofstrain amplitude, as a result of the rapid increase of anelastic strain and decrease ofsecant elastic modulus. At this time the damping capacityQ1
hincreases graduallywith the increase of amplitude. The first two anelastic stages can be explained bydislocation theory developed by Granato and Lücke (G-L); however, a deviationfrom the straight line occurs beyond the second strain amplitude (about5×10-4), thebasal dislocations break away from strong pinning points, which results in theoccurrence of microplastic deformation. In this case the damping capacityQ_p~(-1)rapidly increases; due to the tangle and piled-up of dislocations above the thirdstrain amplitude (about9×10~(-4)), the friction stress gradually increases but theincrease speed of modulus defect slows down. The last two microplasticdeformation processes should be explained in terms of the microplastic dislocationmodel. With increasing the grain size or the Schmid factor of basal slip, the dampingcapacity, anelastic strain and elastic modulus increase defect, but the secant elasticmodulus, friction and back stresses decrease.
Two kinds of Mg alloys which containing elements that possess differentsolubilities in Mg were designed and fabricated. The effects of alloy elements ondamping and microplasyicity of Mg alloys were studied with the comparison of pureMg. The distance of weak pinning points on dislocations was reduced when Alelement was added to Mg, so Mg-1Al alloy exhibited the largest secant elasticmodulus and friction stress but the lowest damping capacity at low strain region.However, the addition of Si did not result in the derease of damping capacitybecause it was unable to dissolve in Mg. But the second phase Mg_2Si led to themarked decrease of grain size of as-cast Mg-1Si alloy, which blocked the motion ofdislocations and so resulted in the derease of damping capacity at high strain region.While the secant elastic modulus, friction stress and back stress were increased atthe same time.
It is supposed that the area swept by a dislocation pinned by strong pinningpoints is a circle at the critical strain amplitude to generate new dislocation. Theaverage mobile dislocation densities and distances between pinning points wereestimated according to the quantitative relationship between the dissipated energy by dislocation motion against the friction stress on the slip plane and dampingcapacity. The average mobile dislocation densities of pure Mg and Mg alloys afterdeformation are in the order of10~(12)m~(-2), and reduced by one or two orders aftershort time annealing treatments.
挤压后纯镁的平均晶粒尺寸为68μm，随着ECAP变形道次的增加，再结晶程度逐渐提高，晶粒逐渐细化，在250°C下ECAP变形4道次纯镁的平均晶粒尺寸为6μm。挤压态纯镁具有基面平行于挤压方向的织构。随着变形道次的增加，基面极点以TD和ED为轴逐步发生倾转，最终形成基面极点分别偏离ND和TD大约40o和65o的织构，而基面平行于挤压方向的织构逐渐弱化，导致沿挤压方向变形的基面滑移Schmid因子大幅度提高。退火后纯镁的晶粒尺寸逐渐增大，而织构基本保持不变。
纯镁与塑性变形相关的阻尼和微塑变行为可用同样的位错机制来解释，它们有着相同的物理本质。塑性阻尼的两个阶段与循环拉伸微塑变的两个阶段相对应：第一阶段对应塑性应变小于2×10~(-4)(总应变约为8×10~(-4))的区域，Schmid因子较大晶粒内部位错从钉扎点上脱钉并在基面滑移，可动性较大，激活体积较大，加工硬化指数较小；当塑性应变高于2×10~(-4)时，由于位错在同一滑移面上运动而发生缠结和堆积，可动位错密度降低，材料硬化，所以具有较大的加工硬化指数和较小的位错滑移激活体积，这时微塑变进入第二阶段。
在应变小于第一临界应变振幅(1×10~(-4)左右)时，纯镁中位错在弱钉扎点间摆动，加载卸载曲线基本重合为一条直线，此时阻尼性能Q_0~(-1)与应变振幅无关；随应变的增大，位错从弱钉扎点上脱钉，产生应力-应变滞后环，滞弹性应变迅速增加，而正切弹性模量快速降低，此时阻尼性能Q_h~(-1)随应变振幅逐渐增大。这两个滞弹性阶段可以用G-L位错模型来解释；当高于第二临界应变振幅(5×10~(-4)左右)后，G-L曲线偏离直线，基面位错从强钉扎点上脱钉，发生微小塑性变形，阻尼性能Q_p~(-1)快速提高；当高于第三临界应变振幅(9×10~(-4)左右)后，由于位错的缠结和堆积，摩擦应力逐渐增大，而损失模量的增加速度变缓。后两个微小塑性变形阶段需要用微塑变位错模型来解释。随着晶粒尺寸或基面滑移Schmid因子的增大，纯镁的阻尼性能、滞弹性应变和损失弹性模量增大，而正切弹性模量、摩擦应力和背应力降低。
选取在镁中固溶的Al和不固溶的Si两类合金元素来制备Mg-Al和Mg-Si合金，并以纯镁作为对比研究合金元素的添加对镁的阻尼和微塑变影响规律。添加完全固溶的1%Al元素会极大的降低位错线上弱钉扎点间距，从而提高镁合金在较小应变下的正切弹性模量，摩擦应力和背应力，但同时也会导致阻尼性能的急剧降低；添加基本不固溶的Si元素不会降低镁的低应变阻尼性能，但会生成Mg2Si第二相，明显的减小铸态镁合金的晶粒尺寸，严重阻碍脱钉后的位错运动，导致较大应变下阻尼性能的降低，同时提高正切弹性模量，摩擦应力和背应力。
在位错增殖的临界应变振幅时，假设位错在强钉扎点钉扎下循环过程中所扫过的面积为圆形。根据功能原理可计算出此临界应变下纯镁和镁合金中的可动位错密度以及位错线上强弱钉扎点间的平均距离，变形后纯镁和镁合金中的可动位错密度为10~(12)m~(-2)数量级，退火态和铸态材料中可动位错密度为10~(10)-10~(11)m~(-2)数量级。
Hot extrusion, equal channel angular pressing (ECAP) and subsequentannealing treatment were performed on pure Mg to modify the microstructure. Theeffects of grain size and texture of pure Mg on microplastic deformation, dampingcapacity and the relationship between them were studied in detail. Two kinds ofalloy elements were added to Mg to study the microplasticity and damping capacityof Mg alloys. The microstructures characteristics of pure Mg and Mg alloys wereobserved by OM and TEM. The texture evolutions were analyzed by neutrondiffraction and EBSD. The strain dependence of damping capacities were studied byDMA. The developments of anelastic strain, friction stress and secant elasticmodulus, etc. in microstrain region were studied by cyclic loading-unloading tests.This paper revealed the damping and microplastic behaviors of pure Mg and Mgalloys, established a good foundation to develop high damping capacity, high fatigueand high dimensional stability Mg based materials.
The average grain size of pure Mg after extrusion was about68μm. Withincreasing the ECAP passes, the degree of recrystallization was increased, whichresulted in the refinement of grain size. The grain size of pure Mg after ECAPprocessing for4passes at the temperature of250°C was refined to6μm. Theas-extruded Mg exhibits a texture with {0001}1010nearly parallel to theextrusion direction (ED). With increasing the ECAP passes, the basal poles rotatedaround transverse direction (TD) and ED axis and formed a maximum componentlocating at about40ofrom normal direction (ND) and65ofrom TD. The texturecomponent with basal planes parallel to ED becomes much weaker. The textureevolution after ECAP led to a increase of the Schmid factor for basal slip towardsED. The grain size was gradually increased, but the texture almost kept stable afterannealing treatments.
The microplasticity and damping related to plasticity of pure Mg have the samephysical mechanism and can be explain by the same dislocation model. The twostages of plastic damping and microplasticity correspond to each other. The firststage relates to the plastic strain below2×10~(-4)(the total strain of about8-9×10~(-4)),the basal dislocations in favorable oriented grains break away from strong pinningpoints and slide on the basal planes, which results in the larger volume activationsbut smaller hardening exponents; the microplastic deformation process enters intothe second region when the plastic strain is above2×10~(-4), the tangle and piled-upof dislocations lead to the decrease of mobile dislocations, the materials are characterized by the larger hardening exponents but smaller volume activations.
When the strain is smaller than the first critical strain amplitude (about1×10-4),the materials almost deform in an elastic manner. At this time the dislocationsreversibly motion between the weak pinning points, the loading-unloading curvesalmost coincide with the elastic line, and the damping capacityQ1
0is amplitudeindependent; the dislocations unpin from weak pinning points with the increase ofstrain amplitude, as a result of the rapid increase of anelastic strain and decrease ofsecant elastic modulus. At this time the damping capacityQ1
hincreases graduallywith the increase of amplitude. The first two anelastic stages can be explained bydislocation theory developed by Granato and Lücke (G-L); however, a deviationfrom the straight line occurs beyond the second strain amplitude (about5×10-4), thebasal dislocations break away from strong pinning points, which results in theoccurrence of microplastic deformation. In this case the damping capacityQ_p~(-1)rapidly increases; due to the tangle and piled-up of dislocations above the thirdstrain amplitude (about9×10~(-4)), the friction stress gradually increases but theincrease speed of modulus defect slows down. The last two microplasticdeformation processes should be explained in terms of the microplastic dislocationmodel. With increasing the grain size or the Schmid factor of basal slip, the dampingcapacity, anelastic strain and elastic modulus increase defect, but the secant elasticmodulus, friction and back stresses decrease.
Two kinds of Mg alloys which containing elements that possess differentsolubilities in Mg were designed and fabricated. The effects of alloy elements ondamping and microplasyicity of Mg alloys were studied with the comparison of pureMg. The distance of weak pinning points on dislocations was reduced when Alelement was added to Mg, so Mg-1Al alloy exhibited the largest secant elasticmodulus and friction stress but the lowest damping capacity at low strain region.However, the addition of Si did not result in the derease of damping capacitybecause it was unable to dissolve in Mg. But the second phase Mg_2Si led to themarked decrease of grain size of as-cast Mg-1Si alloy, which blocked the motion ofdislocations and so resulted in the derease of damping capacity at high strain region.While the secant elastic modulus, friction stress and back stress were increased atthe same time.
It is supposed that the area swept by a dislocation pinned by strong pinningpoints is a circle at the critical strain amplitude to generate new dislocation. Theaverage mobile dislocation densities and distances between pinning points wereestimated according to the quantitative relationship between the dissipated energy by dislocation motion against the friction stress on the slip plane and dampingcapacity. The average mobile dislocation densities of pure Mg and Mg alloys afterdeformation are in the order of10~(12)m~(-2), and reduced by one or two orders aftershort time annealing treatments.
引文
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