Mg-Gd（-Mn-Sc）合金拉伸变形时的塑性失稳行为及变形机制的研究
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摘要
Mg-Gd系合金因其优良的高温力学性能,成为高温耐热镁合金中最具潜力的合金系之一,引起了越来越多国内外的重视。近年来,德国的Mordike等开发了一类新型抗蠕变Mg-Gd-Mn-Sc系列镁合金,该系列合金在300℃的抗蠕变性能要优于现有的最成功的耐热商用WE系列镁合金。国内外对Mg-Gd系合金的时效硬化效应、时效析出序列、高温拉伸强度以及高温抗蠕变性能开展了大量的研究。Mg-Gd-Mn-Sc合金系仅有德国的Mordike等进行过研究,研究内容主要集中在该合金的抗蠕变性能及其组织方面。为了能再进一步提高Mg-Gd(-Mn-Sc)合金的机械性能、塑性变形能力以及发展应用潜力,对Mg-Gd(-Mn-Sc)合金塑性变形过程中的组织及性能的相关性、合金的塑性失稳行为以及塑性变形机制的研究相当重要。然而,目前国内外几乎没有这些方面的报道,很有必要对此进行深入系统的研究。
本文以Mg-5wt.%Gd合金以及Mg-5wt.%Gd-1wt.%Mn-0.7wt.%Sc合金为对象,采用光学显微镜(OM)、扫描电子显微镜(SEM)、电子背散射衍射分析技术(EBSD).透射电镜(TEM)、三维原子探针(3DAP)等分析手段,通过硬度、室温高温拉伸试验,研究了Mn、Sc添加元素对Mg-5wt.%Gd铸造合金组织及时效硬化性能的影响；重点研究了Mg-5wt.%Gd(-1wt.%Mn-0.7wt.Sc)合金在室温、150-300℃中温以及350-400℃高温下拉伸时出现的几种不同类型的塑性变形失稳现象；并系统研究了Mg-5wt.%Gd合金在不同温度和应变速率下拉伸时的显微组织演变和塑性变形机制。研究结果表明：
(1)Mg-5Gd-1 Mn-0.7Sc合金在460-580℃保温的过程中有大量弥散的Mn-Sc金属间化合物析出,该相对合金硬度的提高没有明显作用。经过固溶工艺制度的优化,610℃／24h被选定为该四元合金理想的固溶处理制度,在该制度下升温过程中析出的Mn-Sc金属间化合物几乎完全重新溶入基体,大幅度增加了基体中的Mn、Sc元素固溶度。Mg-5Gd合金和Mg-5Gd-1Mn-0.7Sc合金经过520℃／24h固溶处理后,在后续的200℃以上温度时效时,不具有硬化效果。在经过610℃／24h固溶处理后,在后续的200℃以上温度时效,二元合金仍然不具有硬化效应；而四元合金的硬化效果却得到显著增强,在200、250和300℃下时效的峰值硬度比固溶处理淬火态的硬度值分别提高了63%、44%和33%,这归因于沿基面密集析出的含Mn和Gd的薄片状亚稳相以及圆片状Mn2Sc相的强化作用。
(2)Mg-5Gd(-1Mn-0.7Sc)挤压态样品在室温拉伸时,出现了两种类型的屈服应力下降(Yield drop)现象。第一种是起始屈服点现象(Initial yield point),它的产生是由于挤压态样品在拉伸前的热处理过程中,样品内部的位错密度得到大幅度减小,在塑性变形之初,样品内可动位错密度大量增加所至。第二种是应变时效屈服点现象,它的产生是由于卸载后的热处理过程中,固溶原子气团将可动位错钉扎,重新加载后位错挣脱固溶原子团的钉扎作用所至。虽然引起这两种室温屈服应力下降的直接原因不同,然而它们的产生都基于一个相同的物理本质：即拉伸前样品内部的可动位错密度都很小,当拉伸进入塑性变形阶段时,样品内的可动位错密度突然增加引起样品的塑性变形速率突然增大至超过外界设定的变形速率,从而导致了屈服应力下降现象的发生。
(3)Mg-5Gd(-1Mn-0.7Sc)挤压态样品在150-300℃的温度和1.67×10-4s-1-1.67×10-2s-1的应变速率范围内拉伸时,应力-应变曲线呈现出锯齿状的特征-PLC效应。低温时效处理可析出大量β’相,减小基体中固溶原子的含量,从而显著地减弱了PLC效应,甚至导致了PLC效应的完全消失。本研究合金的PLC效应归因于固溶原子和位错间的交互作用,即动态应变时效效应。
(4)与Mg-5Gd二元合金相比,Mg-5Gd-1Mn-0.7Sc合金的室温应变时效屈服应力下降的幅度更大,高温PLC效应的临界应变更小、应力上下波动的幅度更大,这是由于应变时效效应与固溶原子团和位错之间的相互作用是密切相关的。虽然Gd原子对应变时效效应起了主要的作用,但四元合金中,Mn原子很可能与Gd原子一起,共同形成Cottrell气团。四元合金中Gd与Mn原子共同形成的原子团对位错的钉扎作用,或者说应变时效作用就要强于二元合金中由Gd原子单独形成的固溶原子团对位错的钉扎作用。
(5)Mg-5Gd挤压态合金在350-400℃、不同应变速率下进行拉伸变形,当应变速率低于4.4×10-4s-1时,应力-应变曲线上呈现出锯齿状特征。随着应变速率的下降,锯齿状应力上下波动的幅度更加明显。当应变速率达到或高于8.8x10-4s-1时,锯齿状现象消失,得到光滑的应力-应变曲线。该锯齿状曲线不属于PLC现象。
(6)在400℃、8.8×10-3s-1的较高应变速率下拉伸变形时,动态再结晶现象发生,合金的晶粒得到很大程度的细化；动态再结晶是高应变速率下变形时的主要软化机制,同时动态再结晶带来的晶粒细化作用进一步加速了软化进程。在400℃、8.8×10-5s-1的较低应变速率下拉伸时,没有发生动态再结晶,而是出现了晶粒的正常长大现象。TEM观察显示,高应变速率下变形后的样品中位错密度高、位错堆积和缠结的现象明显；而低应变速率下变形后的样品中位错密度很低,不能给动态再结晶提供足够的驱动力。动态回复是400℃、低应变速率下变形时的主要软化机制,同时晶粒尺寸的增大会导致合金的强度增加,带来强化效果。动态回复产生的软化作用和晶粒粗化带来的强化作用之间相互竞争,从而导致了400℃、低应变速率下锯齿状拉伸曲线的产生。
(7)在350-400℃、低应变速率下拉伸至一定变形量后缓慢卸载继而重新加载时,会产生屈服应力增大的现象；且卸载速度越慢,屈服应力提高的幅度越大。高应变速率下该现象不存在。这归因于低应变速率拉伸时,应力作用下晶粒长大的程度远大于无应力作用时的情况。缓慢卸载后样品的晶粒尺寸要大于快速卸载继而在无载荷条件下保温的样品的晶粒尺寸。正是由于晶粒增长带来的强化作用,从而导致了应力作用下缓慢卸载的样品的屈服应力提高的幅度更大。
(8)Mg-5Gd挤压态合金室温拉伸时主要的变形模式是基面滑移和孪生,其次为非基面a滑移,a+c锥面滑移只在少数晶粒中参与变形。200℃拉伸时主要的变形模式也是基面滑移和孪生,同时非基面a滑移和锥面a+c滑移在大部分晶粒中启动,协调变形。室温和200℃下,晶粒沿着拉伸方向被逐渐拉长。400℃拉伸时,除了基面滑移以外,非基面a滑移、交滑移、锥面a+c滑移、位错攀移也大量参与了变形,但基面滑移仍然是最主要的滑移系,对拉伸过程中织构的演变起了最重要的作用；同时,400℃下晶界滑移和转动也发挥了重要的作用,致使拉伸过程中晶粒的形貌一直保持着等轴状。
Mg-Gd system is one of the most promising candidates for developing precipitation hardenable magnesium alloys and becomes more and more attractive materials for aerospace and automotive applications due to its high specific strength and excellent mechanical properties at elevated temperatures. In recent years, Mordike et al. developed a new group of creep-resistant Mg-Gd materials microalloyed with Mn and Sc. These newly developed Mg-Gd-Mn-Sc alloys have superior creep resistance at 300℃that are better than those of commercial WE54 or WE43 alloys. The previous investigations on Mg-Gd alloy focused on the age-hardening response, precipitation during ageing, tensile properties at elevated temperatures and creep resistance. The reports on Mg-Gd-Mn-Sc alloy were mainly about the creep resistance and the corresponding microstructures. An improved understanding of the correlation of microstructures and deformation behaviors, the plastic deformability and the plastic instabilities of the Mg-Gd(-Mn-Sc) alloys is of critical importance for any further enhancement in mechanical properties at ambient and elevated temperatures and exploiting application prospect.
In this study a Mg-5wt.%Gd alloy and a Mg-5wt.%Gd-1wt.%Mn-0.7wt.%Sc alloy have been selected as model alloys. The effects of Mn and Sc elements on the microstructure and age-hardening response of Mg-Gd alloy, different kinds of plastic instabilities occurred in the Mg-Gd(-Mn-Sc) alloys at temperatures ranging from room temperature (RT) to 400℃and the microstructure evolution and plastic deformation mechanisms during tensile tests under different temperatures and strain rates have been systematically investigated. Optical microscope (OM), scanning electron microscope (SEM), electron backscatter diffraction (EBSD) equipment, transmission electron microscope (TEM) and three-dimensional atom probe (3DAP) were used to examine the microstructures. The results show that:
(1) With the additions of Mn and Sc elements, a large number of Mn-Sc intermetallic dispersoids precipitate in the quaternary alloy after heat treated at 460～580℃. However, when the quaternary alloy is solution treated at 610℃for 24h, these Mn-Sc dispersoids almost fully dissolve intoα-Mg matrix, which leads to the significant increase of the amount of Mn and Sc atoms in the matrix. The binary alloy exhibite almost no age-hardening response above 200℃, no matter the alloy is solution treated at 520℃or 610℃for 24h before ageing. For the quaternary alloy, no obvious age-hardening above 200℃occurs in the sample suffered from 520℃/24h heat treatment before ageing. However, when the quaternary alloy is solution treated at 610℃for 24h then aged at 200～300℃, a substantial increase in hardness is observed. The peak hardnesses obtained during ageing at 200,250 and 300℃are increased by 63%,44% and 33% respectively. The dense precipitation of very thin Mn-Gd plates and tiny Mn2Sc discs is considered to be mainly responsible for the greatly improved age-hardening response.
(2) Initial yield point and strain ageing yield point are observed during the tensile tests at RT. The initial yield point is associated with the significant decrease in the dislocation density during the proper heat treatment before tensile tests. The strain ageing yield point is ascribed to the interaction between solute atoms and the dislocations. When a proper heat treatment is applied to the unloaded sample, the solute atoms can diffuse to and pin the dislocations. After reloading the sample, the strain ageing yield drop happens when the dislocations unpin the solute atoms atmosphere to move. These two yield points are caused by different phenomena, however, they can be ascribed to the same physical principle. In either case, the result of ageing is to reduce the mobile dislocation density below that value needed to sustain the imposed strain rate. When the tensile test is resumed, an increase in stress initially manifests the mobile dislocation deficiency, but when the dislocations multiply quickly and the plastic strain rate exceeds the imposed strain rate, the yield drop is observed.
(3) Serrated flow is observed when the as-extruded samples of Mg-5Gd(-1Mn-0.7Sc) alloys are tensile tested at temperatures ranging from 150 to 300℃and at strain rates of 1.67×10-4s-1 to 1.67×10-2s-1. The serrated flow phenomenon is significantly influenced by the heat treatment conditions.β' phase precipitates densely and the content of solute atoms decreases greatly during the ageing treatment at low temperatures, which weakens the PLC effect or even leads to the disappearance of the PLC effect. The serrated flow is attributed to dynamic interactions between solute atoms and gliding dislocations, i.e. dynamic strain ageing.
(4) The Mg-Gd-Mn-Sc alloy has a smaller critical strain for serrated flow and larger stress drops for both strain ageing yield point at RT and serrated flow at elevated temperatures than those in the Mg-Gd alloy. This phenomenon may be ascribed to the contribution of Mn atoms. The Mn atoms have a smaller atomic radius than Mg atoms, and when combined with larger atoms of Gd, can provide more effective pinning of dislocations than a single type of solute atoms in the Mg-Gd alloy.
(5) When the Mg-Gd extruded samples tensile tested at 350～400℃under different strain rates, serrated flow is observed at the strain rates lower than 4.4×10-4s-1, and absent at strain rates higher than 8.8×10-4s-1.This serrated flow is not ascribed to the PLC effect.
(6) Dynamic recrystallization (DRX) occurs in the specimen tested at 400℃at higher strain rate of 8.8×10-3s-1, and leads to the grain refinement. DRX is the main softening mechanism for the samples deformed at higher strain rates, and the grain refinement caused by DRX further leads to the softening effect. During the tensile at lower strain rate of 8.8×10-5s-1, no sign for the operation of DRX can be detected but normal grain growth is observed. TEM microstructure observation shows that the dislocation density is large in the samples tested at high strain rate, but the dislocation density in the samples tested at low strain rate is too small to stimulate DRX. Dynamic recovery is the main softening mechanism for the samples deformed at low strain rates. Meanwhile, the grain growth brings the hardening effect. Because of the competition between continuous dynamic recovery (softening factor) and continuous grain growth (hardening factor), the serrated flow is observed.
(7) When the specimens are loaded at 350～400℃, at low strain rates to 1-2% total strain, unloaded slowly, and then reloaded again, an increased yield strength is observed. The magnitude of the increase in the yield strength is larger when a slower unloading speed is applied. The phenomenon is absent at the higher strain rates. This is due to the normal grain growth under stress during slow unloading process, which leads to a strengthen effect when tensile temperature is as high as 350～400℃.
(8) At room temperature basal slip and mechanical twinning are predominant in Mg-5Gd extruded samples to accommodate the plastic strain; Non-basal a slip and a+c pyramidal slip can only operate in the minority of the grains. At 200℃basal slip and mechanical twinning are also mainly deformation mechanisms; Non-basal a slip and a+c pyramidal slip are initiated in most of the grains. The grains are stretched gradually along the tensile direction with increasing strain at RT and 200℃. Basal slip, non-basal a slip, cross slip, a+c pyramidal slip, dislocation climbing are all very active at 400℃, but basal slip is still the most active slip system and plays an important role in the texture evolution. Grain boundary slip and grain rotation also operate at 400℃, which leads to the grains remaining equiaxial shape during the whole tensile process.
本文以Mg-5wt.%Gd合金以及Mg-5wt.%Gd-1wt.%Mn-0.7wt.%Sc合金为对象,采用光学显微镜(OM)、扫描电子显微镜(SEM)、电子背散射衍射分析技术(EBSD).透射电镜(TEM)、三维原子探针(3DAP)等分析手段,通过硬度、室温高温拉伸试验,研究了Mn、Sc添加元素对Mg-5wt.%Gd铸造合金组织及时效硬化性能的影响；重点研究了Mg-5wt.%Gd(-1wt.%Mn-0.7wt.Sc)合金在室温、150-300℃中温以及350-400℃高温下拉伸时出现的几种不同类型的塑性变形失稳现象；并系统研究了Mg-5wt.%Gd合金在不同温度和应变速率下拉伸时的显微组织演变和塑性变形机制。研究结果表明：
(1)Mg-5Gd-1 Mn-0.7Sc合金在460-580℃保温的过程中有大量弥散的Mn-Sc金属间化合物析出,该相对合金硬度的提高没有明显作用。经过固溶工艺制度的优化,610℃／24h被选定为该四元合金理想的固溶处理制度,在该制度下升温过程中析出的Mn-Sc金属间化合物几乎完全重新溶入基体,大幅度增加了基体中的Mn、Sc元素固溶度。Mg-5Gd合金和Mg-5Gd-1Mn-0.7Sc合金经过520℃／24h固溶处理后,在后续的200℃以上温度时效时,不具有硬化效果。在经过610℃／24h固溶处理后,在后续的200℃以上温度时效,二元合金仍然不具有硬化效应；而四元合金的硬化效果却得到显著增强,在200、250和300℃下时效的峰值硬度比固溶处理淬火态的硬度值分别提高了63%、44%和33%,这归因于沿基面密集析出的含Mn和Gd的薄片状亚稳相以及圆片状Mn2Sc相的强化作用。
(2)Mg-5Gd(-1Mn-0.7Sc)挤压态样品在室温拉伸时,出现了两种类型的屈服应力下降(Yield drop)现象。第一种是起始屈服点现象(Initial yield point),它的产生是由于挤压态样品在拉伸前的热处理过程中,样品内部的位错密度得到大幅度减小,在塑性变形之初,样品内可动位错密度大量增加所至。第二种是应变时效屈服点现象,它的产生是由于卸载后的热处理过程中,固溶原子气团将可动位错钉扎,重新加载后位错挣脱固溶原子团的钉扎作用所至。虽然引起这两种室温屈服应力下降的直接原因不同,然而它们的产生都基于一个相同的物理本质：即拉伸前样品内部的可动位错密度都很小,当拉伸进入塑性变形阶段时,样品内的可动位错密度突然增加引起样品的塑性变形速率突然增大至超过外界设定的变形速率,从而导致了屈服应力下降现象的发生。
(3)Mg-5Gd(-1Mn-0.7Sc)挤压态样品在150-300℃的温度和1.67×10-4s-1-1.67×10-2s-1的应变速率范围内拉伸时,应力-应变曲线呈现出锯齿状的特征-PLC效应。低温时效处理可析出大量β’相,减小基体中固溶原子的含量,从而显著地减弱了PLC效应,甚至导致了PLC效应的完全消失。本研究合金的PLC效应归因于固溶原子和位错间的交互作用,即动态应变时效效应。
(4)与Mg-5Gd二元合金相比,Mg-5Gd-1Mn-0.7Sc合金的室温应变时效屈服应力下降的幅度更大,高温PLC效应的临界应变更小、应力上下波动的幅度更大,这是由于应变时效效应与固溶原子团和位错之间的相互作用是密切相关的。虽然Gd原子对应变时效效应起了主要的作用,但四元合金中,Mn原子很可能与Gd原子一起,共同形成Cottrell气团。四元合金中Gd与Mn原子共同形成的原子团对位错的钉扎作用,或者说应变时效作用就要强于二元合金中由Gd原子单独形成的固溶原子团对位错的钉扎作用。
(5)Mg-5Gd挤压态合金在350-400℃、不同应变速率下进行拉伸变形,当应变速率低于4.4×10-4s-1时,应力-应变曲线上呈现出锯齿状特征。随着应变速率的下降,锯齿状应力上下波动的幅度更加明显。当应变速率达到或高于8.8x10-4s-1时,锯齿状现象消失,得到光滑的应力-应变曲线。该锯齿状曲线不属于PLC现象。
(6)在400℃、8.8×10-3s-1的较高应变速率下拉伸变形时,动态再结晶现象发生,合金的晶粒得到很大程度的细化；动态再结晶是高应变速率下变形时的主要软化机制,同时动态再结晶带来的晶粒细化作用进一步加速了软化进程。在400℃、8.8×10-5s-1的较低应变速率下拉伸时,没有发生动态再结晶,而是出现了晶粒的正常长大现象。TEM观察显示,高应变速率下变形后的样品中位错密度高、位错堆积和缠结的现象明显；而低应变速率下变形后的样品中位错密度很低,不能给动态再结晶提供足够的驱动力。动态回复是400℃、低应变速率下变形时的主要软化机制,同时晶粒尺寸的增大会导致合金的强度增加,带来强化效果。动态回复产生的软化作用和晶粒粗化带来的强化作用之间相互竞争,从而导致了400℃、低应变速率下锯齿状拉伸曲线的产生。
(7)在350-400℃、低应变速率下拉伸至一定变形量后缓慢卸载继而重新加载时,会产生屈服应力增大的现象；且卸载速度越慢,屈服应力提高的幅度越大。高应变速率下该现象不存在。这归因于低应变速率拉伸时,应力作用下晶粒长大的程度远大于无应力作用时的情况。缓慢卸载后样品的晶粒尺寸要大于快速卸载继而在无载荷条件下保温的样品的晶粒尺寸。正是由于晶粒增长带来的强化作用,从而导致了应力作用下缓慢卸载的样品的屈服应力提高的幅度更大。
(8)Mg-5Gd挤压态合金室温拉伸时主要的变形模式是基面滑移和孪生,其次为非基面a滑移,a+c锥面滑移只在少数晶粒中参与变形。200℃拉伸时主要的变形模式也是基面滑移和孪生,同时非基面a滑移和锥面a+c滑移在大部分晶粒中启动,协调变形。室温和200℃下,晶粒沿着拉伸方向被逐渐拉长。400℃拉伸时,除了基面滑移以外,非基面a滑移、交滑移、锥面a+c滑移、位错攀移也大量参与了变形,但基面滑移仍然是最主要的滑移系,对拉伸过程中织构的演变起了最重要的作用；同时,400℃下晶界滑移和转动也发挥了重要的作用,致使拉伸过程中晶粒的形貌一直保持着等轴状。
Mg-Gd system is one of the most promising candidates for developing precipitation hardenable magnesium alloys and becomes more and more attractive materials for aerospace and automotive applications due to its high specific strength and excellent mechanical properties at elevated temperatures. In recent years, Mordike et al. developed a new group of creep-resistant Mg-Gd materials microalloyed with Mn and Sc. These newly developed Mg-Gd-Mn-Sc alloys have superior creep resistance at 300℃that are better than those of commercial WE54 or WE43 alloys. The previous investigations on Mg-Gd alloy focused on the age-hardening response, precipitation during ageing, tensile properties at elevated temperatures and creep resistance. The reports on Mg-Gd-Mn-Sc alloy were mainly about the creep resistance and the corresponding microstructures. An improved understanding of the correlation of microstructures and deformation behaviors, the plastic deformability and the plastic instabilities of the Mg-Gd(-Mn-Sc) alloys is of critical importance for any further enhancement in mechanical properties at ambient and elevated temperatures and exploiting application prospect.
In this study a Mg-5wt.%Gd alloy and a Mg-5wt.%Gd-1wt.%Mn-0.7wt.%Sc alloy have been selected as model alloys. The effects of Mn and Sc elements on the microstructure and age-hardening response of Mg-Gd alloy, different kinds of plastic instabilities occurred in the Mg-Gd(-Mn-Sc) alloys at temperatures ranging from room temperature (RT) to 400℃and the microstructure evolution and plastic deformation mechanisms during tensile tests under different temperatures and strain rates have been systematically investigated. Optical microscope (OM), scanning electron microscope (SEM), electron backscatter diffraction (EBSD) equipment, transmission electron microscope (TEM) and three-dimensional atom probe (3DAP) were used to examine the microstructures. The results show that:
(1) With the additions of Mn and Sc elements, a large number of Mn-Sc intermetallic dispersoids precipitate in the quaternary alloy after heat treated at 460～580℃. However, when the quaternary alloy is solution treated at 610℃for 24h, these Mn-Sc dispersoids almost fully dissolve intoα-Mg matrix, which leads to the significant increase of the amount of Mn and Sc atoms in the matrix. The binary alloy exhibite almost no age-hardening response above 200℃, no matter the alloy is solution treated at 520℃or 610℃for 24h before ageing. For the quaternary alloy, no obvious age-hardening above 200℃occurs in the sample suffered from 520℃/24h heat treatment before ageing. However, when the quaternary alloy is solution treated at 610℃for 24h then aged at 200～300℃, a substantial increase in hardness is observed. The peak hardnesses obtained during ageing at 200,250 and 300℃are increased by 63%,44% and 33% respectively. The dense precipitation of very thin Mn-Gd plates and tiny Mn2Sc discs is considered to be mainly responsible for the greatly improved age-hardening response.
(2) Initial yield point and strain ageing yield point are observed during the tensile tests at RT. The initial yield point is associated with the significant decrease in the dislocation density during the proper heat treatment before tensile tests. The strain ageing yield point is ascribed to the interaction between solute atoms and the dislocations. When a proper heat treatment is applied to the unloaded sample, the solute atoms can diffuse to and pin the dislocations. After reloading the sample, the strain ageing yield drop happens when the dislocations unpin the solute atoms atmosphere to move. These two yield points are caused by different phenomena, however, they can be ascribed to the same physical principle. In either case, the result of ageing is to reduce the mobile dislocation density below that value needed to sustain the imposed strain rate. When the tensile test is resumed, an increase in stress initially manifests the mobile dislocation deficiency, but when the dislocations multiply quickly and the plastic strain rate exceeds the imposed strain rate, the yield drop is observed.
(3) Serrated flow is observed when the as-extruded samples of Mg-5Gd(-1Mn-0.7Sc) alloys are tensile tested at temperatures ranging from 150 to 300℃and at strain rates of 1.67×10-4s-1 to 1.67×10-2s-1. The serrated flow phenomenon is significantly influenced by the heat treatment conditions.β' phase precipitates densely and the content of solute atoms decreases greatly during the ageing treatment at low temperatures, which weakens the PLC effect or even leads to the disappearance of the PLC effect. The serrated flow is attributed to dynamic interactions between solute atoms and gliding dislocations, i.e. dynamic strain ageing.
(4) The Mg-Gd-Mn-Sc alloy has a smaller critical strain for serrated flow and larger stress drops for both strain ageing yield point at RT and serrated flow at elevated temperatures than those in the Mg-Gd alloy. This phenomenon may be ascribed to the contribution of Mn atoms. The Mn atoms have a smaller atomic radius than Mg atoms, and when combined with larger atoms of Gd, can provide more effective pinning of dislocations than a single type of solute atoms in the Mg-Gd alloy.
(5) When the Mg-Gd extruded samples tensile tested at 350～400℃under different strain rates, serrated flow is observed at the strain rates lower than 4.4×10-4s-1, and absent at strain rates higher than 8.8×10-4s-1.This serrated flow is not ascribed to the PLC effect.
(6) Dynamic recrystallization (DRX) occurs in the specimen tested at 400℃at higher strain rate of 8.8×10-3s-1, and leads to the grain refinement. DRX is the main softening mechanism for the samples deformed at higher strain rates, and the grain refinement caused by DRX further leads to the softening effect. During the tensile at lower strain rate of 8.8×10-5s-1, no sign for the operation of DRX can be detected but normal grain growth is observed. TEM microstructure observation shows that the dislocation density is large in the samples tested at high strain rate, but the dislocation density in the samples tested at low strain rate is too small to stimulate DRX. Dynamic recovery is the main softening mechanism for the samples deformed at low strain rates. Meanwhile, the grain growth brings the hardening effect. Because of the competition between continuous dynamic recovery (softening factor) and continuous grain growth (hardening factor), the serrated flow is observed.
(7) When the specimens are loaded at 350～400℃, at low strain rates to 1-2% total strain, unloaded slowly, and then reloaded again, an increased yield strength is observed. The magnitude of the increase in the yield strength is larger when a slower unloading speed is applied. The phenomenon is absent at the higher strain rates. This is due to the normal grain growth under stress during slow unloading process, which leads to a strengthen effect when tensile temperature is as high as 350～400℃.
(8) At room temperature basal slip and mechanical twinning are predominant in Mg-5Gd extruded samples to accommodate the plastic strain; Non-basal a slip and a+c pyramidal slip can only operate in the minority of the grains. At 200℃basal slip and mechanical twinning are also mainly deformation mechanisms; Non-basal a slip and a+c pyramidal slip are initiated in most of the grains. The grains are stretched gradually along the tensile direction with increasing strain at RT and 200℃. Basal slip, non-basal a slip, cross slip, a+c pyramidal slip, dislocation climbing are all very active at 400℃, but basal slip is still the most active slip system and plays an important role in the texture evolution. Grain boundary slip and grain rotation also operate at 400℃, which leads to the grains remaining equiaxial shape during the whole tensile process.
引文
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