Mg-Gd-Y-Zr（-Ca）合金的微观组织演变、性能和断裂行为研究
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摘要
Mg-Gd系、Mg-Dy系和Mg-Tb系等高性能重稀土镁合金,由于它们优异的比强度和良好的耐热性能,对航天航空、军事工业和赛车等领域是极具吸引力的。其中,Mg-Gd系合金因其具有出众的时效硬化特性和高达250oC的耐热温度,成为最具潜力的合金之一。近年来,为了进一步改善力学性能和降低昂贵的Gd金属的用量,Y、Nd、Sm、Sc和Zn等元素被添加到Mg-Gd合金当中,展现了多元发展的趋势。但是,简单的Mg-Gd二元合金及较为复杂的Mg-Gd-Y三元合金成分-结构-性能之间关系的研究还远不够系统全面,尤其是它们的时效析出过程及其与力学性能之间的关系未完全被理解,这阻碍了复杂多元Mg-Gd系合金的发展。此外,国内关于Mg-Gd系合金的研究与开发尚处于刚起步阶段,因此深入开展Mg-Gd系合金的微观组织、热处理工艺与性能研究十分必要。
本文以配制的数种Mg-(6-12)Gd-(1-3)Y-Zr、Mg-15Gd和Mg-8Y合金(wt%)为研究对象,采用微机数据采集系统、电感耦合等离子直读光谱仪(ICP)、光学显微镜(OM)、金相图像分析仪、差示扫描量热仪(DSC)、差热分析(DTA)、X射线衍射仪(XRD)、带能谱分析(EDAX)的扫描电子显微镜(SEM)和透射电子显微镜(TEM)等分析手段和TEM微衍射(Microdiffraction)技术,通过硬度、室温高温拉伸、蠕变性能和腐蚀性能试验,主要系统地研究了不同Gd含量、热处理工艺及热挤压形变热处理工艺对Mg-xGd-3Y-Zr合金(6≤x≤12)的显微组织、力学性能、蠕变性能、耐蚀性能和断裂行为的影响,分析和探讨了合金的强化机制,重点研究了时效析出相结构、形态、尺寸和分布的演变过程。另外,首次考察了Ca对该合金的组织和性能,尤其是蠕变性能和耐蚀性能的影响。研究目的是为高性能重稀土镁合金的进一步开发和应用提供理论和实践依据。
铸造Mg-Gd-Y-Zr(-Ca)合金的原始态(F)、固溶态(T4)和峰值时效态(T6)的组织依次由α-Mg + Mg24(Gd, Y)5共晶化合物,α-Mg过饱和固溶体+方块相,和α-Mg +方块相+β′析出相组成,而且所有状态都不变地包含了Zr核。方块相是富含Gd和Y的固溶体,具有面心立方结构(fcc),晶格常数a = 5.25?。共晶化合物的组成是Mg24(Gd, Y)5,具有体心立方结构(bcc),晶格常数a = 11.2?。
通过采用优化的热处理工艺,铸造Mg-xGd-3Y-Zr合金(6≤x≤12)的室温抗拉强度在Gd含量为10%时达到峰值(370MPa),并具有241MPa的屈服强度和4.0%的延伸率的最佳组合。所有Gd含量≥9%的Mg-xGd-3Y-Zr(-Ca)合金的高温力学性能均优于WE54合金。通过采用合适的挤压工艺显著细化晶粒,并联合采用6%冷加工硬化和时效强化手段,合金含量最高的GW123K获得了屈服强度和抗拉强度分别为436MPa和491MPa的最高强度指标,但与此同时延伸率也下降到了3.6%。铸造T6和挤压T5态Mg-Gd-Y-Zr合金在从室温到200oC的拉伸试验温度段,强度只发生了平缓而轻微地下降;过了250oC则都急剧下跌,且T5态下降速度比T6态更快,与此同时延伸率都大幅升高。GW123K和GW102K合金挤压T5态的瞬间高温拉伸强度都高于2618耐热铝合金和WE54商业耐热镁合金。
Mg-Gd-(Y)-Zr合金的过饱和固溶体S.S.S.S.(cph)随着时效时间的延长依如下顺序进行四阶段分解和转化:β″(D019)→β′(cbco)→β1(fcc)→β(fcc),与前人三阶段的研究结果β″(D019)→β′(cbco)→β(fcc)有所不同。具有三角分布的棱柱面片状析出相β′是合金硬度峰值状态的主要强化相。过时效阶段,β1在处于分解状态中的β′相的颈缩处形核并沿另一方向进行生长。随后,β1原位转变成平衡相β,并且保持原来与基体的位向关系。这一过程最终导致了析出相的位向由β′的棱镜面{2110}α向另一方位棱镜面β的{1010}α转变。
原子结构模型表明,β″和β′相的原子堆垛方式相对于α-Mg并没有改变,而只是Gd和Y等溶质原子重新进行了长周期有序排列,所以β″和β′相在整个空间上都是与基体高度共格的,而β1相在惯析面上几乎与基体完全共格,但在与惯析面相垂直的顶面和底面至多形成半共格界面,直到最终的平衡相β仍很有可能在惯析面上与基体形成半共格界面甚至共格界面。
在200oC/200MPa条件下,各铸造T6态合金的蠕变抗力在300h内依下列次序递增:GW63K < GW83K < WE54 < GW103K < GW113K,其中高Gd含量的GW103K和GW113K合金的蠕变抗力优于WE54,GW113K的最小蠕变速率最低,比WE54合金降低了40%以上。在200MPa的高应力条件下,Mg-Gd-Y-Zr(-Ca)合金可以应用于175oC以下温度的工程应用场合。
首次对铸造和挤压Mg-Gd-Y合金的进行了抗蠕变性能对比研究。发现细晶挤压T5态Mg-Gd-Y合金的最小蠕变速率相对于铸造T6态急剧提高,幅度可达2个数量级以上。在200oC/160MPa的次高温、高应力的蠕变条件下挤压T5态Mg-Gd-Y合金就表现出了明显的扩散蠕变特征,只经过不到75h就形成了很宽的晶界无析出区,而且合金的蠕变伸长主要由析出相贫化区引起。
在175-200oC/160-200MPa的温度和应力范围内,Mg-Gd-Y-Zr(-Ca)合金铸造T6和挤压T5态的蠕变激活能Qc均在190-256 kJ/mol之间,均大于镁的自扩散激活能;在应力指数方面,铸造T6态合金均高于相应的挤压T5态合金,挤压T5态合金的n在4-6之间,铸造T6态合金的n在6-10之间。TEM对位错的观测结果表明,200oC/160MPa蠕变条件下,一级角锥面{1011}和二级角锥面{1012}等非基面滑移以及它们与基面之间的交滑移得以激活。
对铸造T6态Mg-xGd-3Y-Zr合金而言,盐雾腐蚀失重速率随着Gd含量的增加经历了先升后降的过程。耐腐蚀性能以最低合金含量的GW63K合金为最佳,GW83K次之,GW103K最差,之后再提高Gd含量,到了GW123K合金,耐腐蚀性能反而又略微增强。铸造T6态的GW63K和GW83K合金的耐腐蚀性能优于铸态AZ91D,而T6态的GW103K和GW123K合金的耐腐蚀性能与铸态AZ91D合金相当。相对于铸造T6态,挤压T5态合金晶粒显著细化,晶界面积大幅增加,因而造成自腐蚀电位降低和盐雾腐蚀失重速率急剧升高。
铸造T4、T6和挤压T5态室温断口都可归入准解理断裂之列,而沿晶断裂只发生在高含量合金的铸造F态室温断口当中。随着温度的升高,或加之热加工使晶粒进一步细化,合金的断裂方式逐步由脆性向韧性转化。无论是铸造T6态还是挤压T5态,从室温到200oC,Mg-Gd-Y-Zr合金的断裂方式以准解理断裂为主;在200oC以上,随着温度的升高,断口形貌中微孔聚集型断裂的比重逐渐增加,在250oC合金的断裂方式为微孔聚集+准解理断裂混合型,当温度高达300oC时变成以微孔聚集型断裂为主。
位错在晶界、第二相或夹杂物等障碍物前塞积产生应力集中的机制合理地解释了所研究合金的裂纹起源。经过修正后的Griffith公式则很好地解释了合金随成分和状态变化的裂纹扩展阻力。
铸造T4态的屈服强度的来源在工业纯镁的固有屈服强度的基础上,还包括了固溶强化和细晶强化,其中固溶强化发挥了主要作用,其强化贡献量? s与溶质原子浓度c的关系式为? s ? 905c2/3。铸造T6态在T4态的基础上,以损失大部分固溶强化为代价,产生了显著的析出强化效应,大幅提高了屈服强度,而且各合金的析出强化贡献量均大于总屈服强度的50%。挤压T5态相对于铸造T6态提高了细晶强化的贡献量,并且出现了织构强化,使屈服强度继续获得了明显改善;虽然析出强化的相对比重有所下降,但仍保持最大的贡献量。
Mg-Gd-Y系合金在200oC以下的温度之所以表现出显著的析出强化效应是因为其主要强化相β′具有高效阻碍基面位错滑移的形状和位向,较大的体积分数,良好的析出相/基体共格界面和较高的热稳定性等。时效Mg-Gd-Y合金的强度依靠共格强化和Orowan机制的联合作用达到了峰值状态。在过时效状态下,合金主要受Orowan绕过机制控制。
添加0.4-0.6wt%的Ca能增强Mg-Gd-Y-Zr合金的蠕变抗力并显著提高其耐蚀性能,但严重降低了延伸率,这些变化主要与Ca偏聚于晶界有关。
High-performance magnesium-heavy rare earth alloys, such as those based on Mg-Gd, Mg-Dy, Mg-Tb systems and so on, are very attractive for aerospace, armament and racing automotive industries because of their high specific strength and good thermal stability. Among them, Mg-Gd system is one of the most promising candidates due to the remarkable age-hardening response and very good thermal stability of the main strengthening phase at up to 250oC. Recently, Mg-Gd alloys with the addition of Y, Nd, Sm, Sc, and Zn, were investigated to further improve the mechanical properties or reduce the high content of expensive Gd, which exhibited the development trend of multi-elements alloying. However, there is lack of general and systematic investigation for Mg-Gd and Mg-Gd-Y alloys and the relationship among the compositions, microstructure and properties, especially the precipitation process and its relationship with mechanical properties, which will obstruct the development of the complex Mg-Gd series alloys.
Several Mg-(6-12)Gd-(1-3)Y-Zr, Mg-15Gd and Mg-8Y (wt%) alloys were prepared. Effects of variant content of Gd, heat treatment and thermal-mechanical process on the microstructure, mechanical properties, creep resistance, corrosion resistance and fracture behavior of Mg-xGd-3Y-Zr (6≤x≤12) alloys were mainly investigated, by computer data collection system, optical microscopy (OM), image analysis apparatus with a image analysis software, X–ray diffractometer (XRD), inductively coupled plasma analyzer (ICP), Differential Thermal Analysis (DTA), Differential Scanning Calorimeter (DSC), scanning electron microscopy (SEM) and transmission electron microscopy (TEM) with energy dispersive X-ray analyses (EDAX) and microdiffraction etc.. The strengthening mechanism of the alloys was analyzed and discussed, and microstructural evolution during aging, including the morphology, structure, size and distribution of the precipitates, was studied in detail. Furthermore, the effect of a small Ca addition on microstructure and properties, especially creep resistance and corrosion resistance, was investigated for the first time. The purpose of the present work is to provide theoretical and practical results for the development of high-performance magnesium-heavy rare earth alloys.
The microstructure evolution of cast Mg-Gd-Y-Zr(-Ca) alloys from as-cast to T4 to T6 conditions involves solid solution + eutectic compound→supersaturated solid solution + cuboidal phase→solid solution +β′precipitates + cuboidal phase. In addition, Zirconium cores exist in all these conditions. The cuboidal phase is Gd and Y rich solid solution with a f.c.c. structure (Fm3m cubic, a = 5.25?), which are observed mostly in grain interior but partially at grain boundary. The eutectic compound has a b.c.c. structure with a = 11.2? and a composition of Mg24(Gd, Y)5.
The peak tensile properties of cast-T6 Mg-xGd-3Y-Zr (6≤x≤12) alloys is such that the ultimate tensile strength (UTS), tensile yield strength (TYS) and elongation is 370MPa, 241MPa and 4.0% respectively, which is attained by the optimized heat treatment. UTS and TYS of Mg-xGd-3Y-Zr(-Ca) alloys containing a high Gd content with x≥9% are remarkablely superior to those of WE54, especially in the temperature range from room temperature to 200oC. A very high strength of extruded-T5 GW123K alloy with UTS=491MPa, TYS=436MPa and elongation=3.6%, is achieved by appropriate hot extrusion process, the following 6% cold working hardening and age strengthening at room temperature. The strengths of Mg-xGd-3Y-Zr(-Ca) both in cast-T6 and extruded-T5 conditions decline very slowly from room temperature to 200oC. However, at the temperature of 250oC or more, the strengths steeply decrease, and the difference in strength between cast-T6 and extueded-T5 becomes small at 300oC. The instant tensile strengths of extruded-T5 GW123K and GW102K alloys are higher than those of forged-T6 2618 aluminum alloy and forged-T5 WE54 magnesium alloys at the temperatures from 20oC to 250oC.
The decomposition ofα-Mg supersaturated solid solution (S.S.S.S., cph) in Mg-Gd-(Y)-Zr alloys with increasing aging time is as follows:β″(D019)→β′(cbco)→β1(fcc)→β(fcc), which is similar to that of Mg-Gd-Nd, Mg-Dy-Nd and Mg-Y-Nd alloys, but different from previously reported three stage sequence: S.S.S.S.→β″(D019)→β′(cbco)→β(fcc). It is found that the metastableβ″andβ′phases coexist in the matrix at the very early stage of ageing. Peak age-hardening is attributed to the precipitation of prismaticβ′plates in a triangular arrangement. At the over-aged stage,β1 phase appears to take place via an in situ transformation from a decomposedβ′phase but grows in a direction different from the previous one ofβ′phase. Continued ageing makes theβ1 phase transform in situ to the equilibriumβphase and the orientation relationship between the precipitate and matrix phases is retained through the in situ transformation of theβ1 phase.
The atomic models based on both the microdiffraction and crystallography analysis indicate that the stacking type of atoms inβ″andβ′phases are the same as that of the hcp matrix and the precipitation makes progress simply by long-period ordering of RE atoms such as Gd and Y. Both of them keep a perfect coherency with the matrix. The interface between theβ1 phase and matrix is near to perfectly coherent in the ( 1 100)αhabit plane of theβ1 plates, but only semi-coherent in the (1120 )αwhich is vertical to the habit plane. Probably so does the interface between theβphase and matrix.
In 200oC/200MPa test condition within 300h, the creep resistance of cast-T6 alloys increases in the following order: GW63K < GW83K < WE54 < GW103K < GW113K, where the creep resistance of GW103K or GW113K alloy is higher than that of WE54 alloy, and the minimum creep rate of GW113K is lower than that of WE54 alloy by over 40%. Under a constant stress of 200MPa, Mg-Gd-Y-Zr(-Ca) alloys can endure the applying temperature lower than 175oC.
Comparative investigation on the creep resistance of cast and extruded Mg-Gd-Y alloys has been carried out. The minimum creep rates of extruded-T5 alloys with fine grains can be two orders of magnitude more than those of cast-T6 alloys. Under the test condition of 200oC/160MPa, the extruded-T5 alloys have exhibited the characteristic of diffusion creep, and formed a very wide precipitate free zone (PFZ), which results in the main large creep strain.
Within the temperature and applied stress range of 175-200oC/160-200MPa, the activation energy for creep of the cast-T6 and extruded-T5 alloys similarly varies from 190-256kJ/mol, which is higher than the activation energy for self diffusion of magnesium. The stress exponent n of the cast-T6 alloys (n=6-10) is higher than that of the extruded-T5 alloys (n=4-6). It can be showed from the result of TEM observation that the non-basal slip planes of dislocations, including the first order pyramidal plane {1011} and the second order plane {1012 }, and the cross-slip between them and basal plane, can be activated under the creep condition of 200oC/160MPa.
The weight loss corrosion rate of cast-T6 Mg-xGd-3Y-Zr alloys in the salt spray test increases first and decreases later, and the corrosion resistance of the alloys descends in the following order: GW63K < GW83K < GW123K < GW103K. The corrosion resistance of cast-T6 GW63K or cast-T6 GW83K alloy is superior to that of as cast AZ91D, and the corrosion resistance of cast-T6 GW103K or cast-T6 GW123K alloy is comparable with that of as cast AZ91D. Compared with cast-T6 alloys, the self-corrosion potential and weight loss corrosion rate of cast-T5 alloys increase dramatically due to the obvious grain refinement.
The rupture of cast-T4, cast-T6 and ectruded-T5 alloys mainly belongs to quasi-cleavage fracture, and the inter-granular fracture cleavage fracture only occurs in the as-cast alloys with high Gd content. With the elevated temperature, or the grain refinement by hot extrusion, the fracture style of the alloys transforms gradually from brittleness to toughness. From room temperature to 200oC, the rupture of the cast-T6 and extruded-T5 alloys is mainly quasi-cleavage fracture; over 200oC, the proportion of the fracture due to microcavities coalescing increases with the elevated temperature; the rupture of the alloys consists of mixed microcavities coalescing and quasi-cleavage fracture at 250oC, but microcavities coalescing becomes the main fracture mechanism at 300oC.
The dislocations piling-up model can be successfully applied to explain the initiation mechanism of cleavage microcracks, and the modified Griffith equation to explain well the resistance to crack propagation which vary with the composition and status of the alloys.
Two contributions to the TYS above that of pure Mg, can be identified for the cast-T4 alloys, namely solid solution strengthening and grain refinement strengthening, where the solid solution strengthening plays the main strengthening role in the cast-T4 alloys. The relationship between TYS (σs) and the atom concentration (c) is such that ? s ? 905c2/3. In spite of sacrificing most of solid solution strengthening, the TYS of cast-T6 alloys is greatly elevated from cast-T4 alloys due to the remarkable precipitation strengthening, which is main strengthening contribution over 50% of total TYS to all the cast-T6 alloys. The contribution from grain refinement strengthening increases and the one from texture strengthening occurs in the extruded-T5 alloys, and both of these two contributors lead to the increase of TYS in the extruded-T5 alloys, compared to cast-T6 alloys. Although the proportion of strengthening contribution due to precipitation strengthening decreases, it keeps as the largest contributor.
At the temperature lower than 200oC, the outstanding precipitation strengthening in aged Mg-Gd-Y alloys is mainly attributed to the favorable shape and orientation ofβ′precipitates which make them obstruct effectively the basal dislocation slip, the bigger volume fraction of the precipitates, perfectly coherent interface between the precipitates and the matrix, and good thermal stability of the precipitates. Coherent strengthening and Orowan mechanism both contribute to the peak hardness of aged Mg-Gd-Y alloys. However, Orowan mechanism becomes the main contributor to TYS at the over-aged stage.
A small addition of 0.4-0.6wt% Ca improves the creep and corrosion resistances of Mg-Gd-Y-Zr alloys but severely deteriorates the elongation, which is related to the presence of the Ca segregation at grain boundaries.
本文以配制的数种Mg-(6-12)Gd-(1-3)Y-Zr、Mg-15Gd和Mg-8Y合金(wt%)为研究对象,采用微机数据采集系统、电感耦合等离子直读光谱仪(ICP)、光学显微镜(OM)、金相图像分析仪、差示扫描量热仪(DSC)、差热分析(DTA)、X射线衍射仪(XRD)、带能谱分析(EDAX)的扫描电子显微镜(SEM)和透射电子显微镜(TEM)等分析手段和TEM微衍射(Microdiffraction)技术,通过硬度、室温高温拉伸、蠕变性能和腐蚀性能试验,主要系统地研究了不同Gd含量、热处理工艺及热挤压形变热处理工艺对Mg-xGd-3Y-Zr合金(6≤x≤12)的显微组织、力学性能、蠕变性能、耐蚀性能和断裂行为的影响,分析和探讨了合金的强化机制,重点研究了时效析出相结构、形态、尺寸和分布的演变过程。另外,首次考察了Ca对该合金的组织和性能,尤其是蠕变性能和耐蚀性能的影响。研究目的是为高性能重稀土镁合金的进一步开发和应用提供理论和实践依据。
铸造Mg-Gd-Y-Zr(-Ca)合金的原始态(F)、固溶态(T4)和峰值时效态(T6)的组织依次由α-Mg + Mg24(Gd, Y)5共晶化合物,α-Mg过饱和固溶体+方块相,和α-Mg +方块相+β′析出相组成,而且所有状态都不变地包含了Zr核。方块相是富含Gd和Y的固溶体,具有面心立方结构(fcc),晶格常数a = 5.25?。共晶化合物的组成是Mg24(Gd, Y)5,具有体心立方结构(bcc),晶格常数a = 11.2?。
通过采用优化的热处理工艺,铸造Mg-xGd-3Y-Zr合金(6≤x≤12)的室温抗拉强度在Gd含量为10%时达到峰值(370MPa),并具有241MPa的屈服强度和4.0%的延伸率的最佳组合。所有Gd含量≥9%的Mg-xGd-3Y-Zr(-Ca)合金的高温力学性能均优于WE54合金。通过采用合适的挤压工艺显著细化晶粒,并联合采用6%冷加工硬化和时效强化手段,合金含量最高的GW123K获得了屈服强度和抗拉强度分别为436MPa和491MPa的最高强度指标,但与此同时延伸率也下降到了3.6%。铸造T6和挤压T5态Mg-Gd-Y-Zr合金在从室温到200oC的拉伸试验温度段,强度只发生了平缓而轻微地下降;过了250oC则都急剧下跌,且T5态下降速度比T6态更快,与此同时延伸率都大幅升高。GW123K和GW102K合金挤压T5态的瞬间高温拉伸强度都高于2618耐热铝合金和WE54商业耐热镁合金。
Mg-Gd-(Y)-Zr合金的过饱和固溶体S.S.S.S.(cph)随着时效时间的延长依如下顺序进行四阶段分解和转化:β″(D019)→β′(cbco)→β1(fcc)→β(fcc),与前人三阶段的研究结果β″(D019)→β′(cbco)→β(fcc)有所不同。具有三角分布的棱柱面片状析出相β′是合金硬度峰值状态的主要强化相。过时效阶段,β1在处于分解状态中的β′相的颈缩处形核并沿另一方向进行生长。随后,β1原位转变成平衡相β,并且保持原来与基体的位向关系。这一过程最终导致了析出相的位向由β′的棱镜面{2110}α向另一方位棱镜面β的{1010}α转变。
原子结构模型表明,β″和β′相的原子堆垛方式相对于α-Mg并没有改变,而只是Gd和Y等溶质原子重新进行了长周期有序排列,所以β″和β′相在整个空间上都是与基体高度共格的,而β1相在惯析面上几乎与基体完全共格,但在与惯析面相垂直的顶面和底面至多形成半共格界面,直到最终的平衡相β仍很有可能在惯析面上与基体形成半共格界面甚至共格界面。
在200oC/200MPa条件下,各铸造T6态合金的蠕变抗力在300h内依下列次序递增:GW63K < GW83K < WE54 < GW103K < GW113K,其中高Gd含量的GW103K和GW113K合金的蠕变抗力优于WE54,GW113K的最小蠕变速率最低,比WE54合金降低了40%以上。在200MPa的高应力条件下,Mg-Gd-Y-Zr(-Ca)合金可以应用于175oC以下温度的工程应用场合。
首次对铸造和挤压Mg-Gd-Y合金的进行了抗蠕变性能对比研究。发现细晶挤压T5态Mg-Gd-Y合金的最小蠕变速率相对于铸造T6态急剧提高,幅度可达2个数量级以上。在200oC/160MPa的次高温、高应力的蠕变条件下挤压T5态Mg-Gd-Y合金就表现出了明显的扩散蠕变特征,只经过不到75h就形成了很宽的晶界无析出区,而且合金的蠕变伸长主要由析出相贫化区引起。
在175-200oC/160-200MPa的温度和应力范围内,Mg-Gd-Y-Zr(-Ca)合金铸造T6和挤压T5态的蠕变激活能Qc均在190-256 kJ/mol之间,均大于镁的自扩散激活能;在应力指数方面,铸造T6态合金均高于相应的挤压T5态合金,挤压T5态合金的n在4-6之间,铸造T6态合金的n在6-10之间。TEM对位错的观测结果表明,200oC/160MPa蠕变条件下,一级角锥面{1011}和二级角锥面{1012}等非基面滑移以及它们与基面之间的交滑移得以激活。
对铸造T6态Mg-xGd-3Y-Zr合金而言,盐雾腐蚀失重速率随着Gd含量的增加经历了先升后降的过程。耐腐蚀性能以最低合金含量的GW63K合金为最佳,GW83K次之,GW103K最差,之后再提高Gd含量,到了GW123K合金,耐腐蚀性能反而又略微增强。铸造T6态的GW63K和GW83K合金的耐腐蚀性能优于铸态AZ91D,而T6态的GW103K和GW123K合金的耐腐蚀性能与铸态AZ91D合金相当。相对于铸造T6态,挤压T5态合金晶粒显著细化,晶界面积大幅增加,因而造成自腐蚀电位降低和盐雾腐蚀失重速率急剧升高。
铸造T4、T6和挤压T5态室温断口都可归入准解理断裂之列,而沿晶断裂只发生在高含量合金的铸造F态室温断口当中。随着温度的升高,或加之热加工使晶粒进一步细化,合金的断裂方式逐步由脆性向韧性转化。无论是铸造T6态还是挤压T5态,从室温到200oC,Mg-Gd-Y-Zr合金的断裂方式以准解理断裂为主;在200oC以上,随着温度的升高,断口形貌中微孔聚集型断裂的比重逐渐增加,在250oC合金的断裂方式为微孔聚集+准解理断裂混合型,当温度高达300oC时变成以微孔聚集型断裂为主。
位错在晶界、第二相或夹杂物等障碍物前塞积产生应力集中的机制合理地解释了所研究合金的裂纹起源。经过修正后的Griffith公式则很好地解释了合金随成分和状态变化的裂纹扩展阻力。
铸造T4态的屈服强度的来源在工业纯镁的固有屈服强度的基础上,还包括了固溶强化和细晶强化,其中固溶强化发挥了主要作用,其强化贡献量? s与溶质原子浓度c的关系式为? s ? 905c2/3。铸造T6态在T4态的基础上,以损失大部分固溶强化为代价,产生了显著的析出强化效应,大幅提高了屈服强度,而且各合金的析出强化贡献量均大于总屈服强度的50%。挤压T5态相对于铸造T6态提高了细晶强化的贡献量,并且出现了织构强化,使屈服强度继续获得了明显改善;虽然析出强化的相对比重有所下降,但仍保持最大的贡献量。
Mg-Gd-Y系合金在200oC以下的温度之所以表现出显著的析出强化效应是因为其主要强化相β′具有高效阻碍基面位错滑移的形状和位向,较大的体积分数,良好的析出相/基体共格界面和较高的热稳定性等。时效Mg-Gd-Y合金的强度依靠共格强化和Orowan机制的联合作用达到了峰值状态。在过时效状态下,合金主要受Orowan绕过机制控制。
添加0.4-0.6wt%的Ca能增强Mg-Gd-Y-Zr合金的蠕变抗力并显著提高其耐蚀性能,但严重降低了延伸率,这些变化主要与Ca偏聚于晶界有关。
High-performance magnesium-heavy rare earth alloys, such as those based on Mg-Gd, Mg-Dy, Mg-Tb systems and so on, are very attractive for aerospace, armament and racing automotive industries because of their high specific strength and good thermal stability. Among them, Mg-Gd system is one of the most promising candidates due to the remarkable age-hardening response and very good thermal stability of the main strengthening phase at up to 250oC. Recently, Mg-Gd alloys with the addition of Y, Nd, Sm, Sc, and Zn, were investigated to further improve the mechanical properties or reduce the high content of expensive Gd, which exhibited the development trend of multi-elements alloying. However, there is lack of general and systematic investigation for Mg-Gd and Mg-Gd-Y alloys and the relationship among the compositions, microstructure and properties, especially the precipitation process and its relationship with mechanical properties, which will obstruct the development of the complex Mg-Gd series alloys.
Several Mg-(6-12)Gd-(1-3)Y-Zr, Mg-15Gd and Mg-8Y (wt%) alloys were prepared. Effects of variant content of Gd, heat treatment and thermal-mechanical process on the microstructure, mechanical properties, creep resistance, corrosion resistance and fracture behavior of Mg-xGd-3Y-Zr (6≤x≤12) alloys were mainly investigated, by computer data collection system, optical microscopy (OM), image analysis apparatus with a image analysis software, X–ray diffractometer (XRD), inductively coupled plasma analyzer (ICP), Differential Thermal Analysis (DTA), Differential Scanning Calorimeter (DSC), scanning electron microscopy (SEM) and transmission electron microscopy (TEM) with energy dispersive X-ray analyses (EDAX) and microdiffraction etc.. The strengthening mechanism of the alloys was analyzed and discussed, and microstructural evolution during aging, including the morphology, structure, size and distribution of the precipitates, was studied in detail. Furthermore, the effect of a small Ca addition on microstructure and properties, especially creep resistance and corrosion resistance, was investigated for the first time. The purpose of the present work is to provide theoretical and practical results for the development of high-performance magnesium-heavy rare earth alloys.
The microstructure evolution of cast Mg-Gd-Y-Zr(-Ca) alloys from as-cast to T4 to T6 conditions involves solid solution + eutectic compound→supersaturated solid solution + cuboidal phase→solid solution +β′precipitates + cuboidal phase. In addition, Zirconium cores exist in all these conditions. The cuboidal phase is Gd and Y rich solid solution with a f.c.c. structure (Fm3m cubic, a = 5.25?), which are observed mostly in grain interior but partially at grain boundary. The eutectic compound has a b.c.c. structure with a = 11.2? and a composition of Mg24(Gd, Y)5.
The peak tensile properties of cast-T6 Mg-xGd-3Y-Zr (6≤x≤12) alloys is such that the ultimate tensile strength (UTS), tensile yield strength (TYS) and elongation is 370MPa, 241MPa and 4.0% respectively, which is attained by the optimized heat treatment. UTS and TYS of Mg-xGd-3Y-Zr(-Ca) alloys containing a high Gd content with x≥9% are remarkablely superior to those of WE54, especially in the temperature range from room temperature to 200oC. A very high strength of extruded-T5 GW123K alloy with UTS=491MPa, TYS=436MPa and elongation=3.6%, is achieved by appropriate hot extrusion process, the following 6% cold working hardening and age strengthening at room temperature. The strengths of Mg-xGd-3Y-Zr(-Ca) both in cast-T6 and extruded-T5 conditions decline very slowly from room temperature to 200oC. However, at the temperature of 250oC or more, the strengths steeply decrease, and the difference in strength between cast-T6 and extueded-T5 becomes small at 300oC. The instant tensile strengths of extruded-T5 GW123K and GW102K alloys are higher than those of forged-T6 2618 aluminum alloy and forged-T5 WE54 magnesium alloys at the temperatures from 20oC to 250oC.
The decomposition ofα-Mg supersaturated solid solution (S.S.S.S., cph) in Mg-Gd-(Y)-Zr alloys with increasing aging time is as follows:β″(D019)→β′(cbco)→β1(fcc)→β(fcc), which is similar to that of Mg-Gd-Nd, Mg-Dy-Nd and Mg-Y-Nd alloys, but different from previously reported three stage sequence: S.S.S.S.→β″(D019)→β′(cbco)→β(fcc). It is found that the metastableβ″andβ′phases coexist in the matrix at the very early stage of ageing. Peak age-hardening is attributed to the precipitation of prismaticβ′plates in a triangular arrangement. At the over-aged stage,β1 phase appears to take place via an in situ transformation from a decomposedβ′phase but grows in a direction different from the previous one ofβ′phase. Continued ageing makes theβ1 phase transform in situ to the equilibriumβphase and the orientation relationship between the precipitate and matrix phases is retained through the in situ transformation of theβ1 phase.
The atomic models based on both the microdiffraction and crystallography analysis indicate that the stacking type of atoms inβ″andβ′phases are the same as that of the hcp matrix and the precipitation makes progress simply by long-period ordering of RE atoms such as Gd and Y. Both of them keep a perfect coherency with the matrix. The interface between theβ1 phase and matrix is near to perfectly coherent in the ( 1 100)αhabit plane of theβ1 plates, but only semi-coherent in the (1120 )αwhich is vertical to the habit plane. Probably so does the interface between theβphase and matrix.
In 200oC/200MPa test condition within 300h, the creep resistance of cast-T6 alloys increases in the following order: GW63K < GW83K < WE54 < GW103K < GW113K, where the creep resistance of GW103K or GW113K alloy is higher than that of WE54 alloy, and the minimum creep rate of GW113K is lower than that of WE54 alloy by over 40%. Under a constant stress of 200MPa, Mg-Gd-Y-Zr(-Ca) alloys can endure the applying temperature lower than 175oC.
Comparative investigation on the creep resistance of cast and extruded Mg-Gd-Y alloys has been carried out. The minimum creep rates of extruded-T5 alloys with fine grains can be two orders of magnitude more than those of cast-T6 alloys. Under the test condition of 200oC/160MPa, the extruded-T5 alloys have exhibited the characteristic of diffusion creep, and formed a very wide precipitate free zone (PFZ), which results in the main large creep strain.
Within the temperature and applied stress range of 175-200oC/160-200MPa, the activation energy for creep of the cast-T6 and extruded-T5 alloys similarly varies from 190-256kJ/mol, which is higher than the activation energy for self diffusion of magnesium. The stress exponent n of the cast-T6 alloys (n=6-10) is higher than that of the extruded-T5 alloys (n=4-6). It can be showed from the result of TEM observation that the non-basal slip planes of dislocations, including the first order pyramidal plane {1011} and the second order plane {1012 }, and the cross-slip between them and basal plane, can be activated under the creep condition of 200oC/160MPa.
The weight loss corrosion rate of cast-T6 Mg-xGd-3Y-Zr alloys in the salt spray test increases first and decreases later, and the corrosion resistance of the alloys descends in the following order: GW63K < GW83K < GW123K < GW103K. The corrosion resistance of cast-T6 GW63K or cast-T6 GW83K alloy is superior to that of as cast AZ91D, and the corrosion resistance of cast-T6 GW103K or cast-T6 GW123K alloy is comparable with that of as cast AZ91D. Compared with cast-T6 alloys, the self-corrosion potential and weight loss corrosion rate of cast-T5 alloys increase dramatically due to the obvious grain refinement.
The rupture of cast-T4, cast-T6 and ectruded-T5 alloys mainly belongs to quasi-cleavage fracture, and the inter-granular fracture cleavage fracture only occurs in the as-cast alloys with high Gd content. With the elevated temperature, or the grain refinement by hot extrusion, the fracture style of the alloys transforms gradually from brittleness to toughness. From room temperature to 200oC, the rupture of the cast-T6 and extruded-T5 alloys is mainly quasi-cleavage fracture; over 200oC, the proportion of the fracture due to microcavities coalescing increases with the elevated temperature; the rupture of the alloys consists of mixed microcavities coalescing and quasi-cleavage fracture at 250oC, but microcavities coalescing becomes the main fracture mechanism at 300oC.
The dislocations piling-up model can be successfully applied to explain the initiation mechanism of cleavage microcracks, and the modified Griffith equation to explain well the resistance to crack propagation which vary with the composition and status of the alloys.
Two contributions to the TYS above that of pure Mg, can be identified for the cast-T4 alloys, namely solid solution strengthening and grain refinement strengthening, where the solid solution strengthening plays the main strengthening role in the cast-T4 alloys. The relationship between TYS (σs) and the atom concentration (c) is such that ? s ? 905c2/3. In spite of sacrificing most of solid solution strengthening, the TYS of cast-T6 alloys is greatly elevated from cast-T4 alloys due to the remarkable precipitation strengthening, which is main strengthening contribution over 50% of total TYS to all the cast-T6 alloys. The contribution from grain refinement strengthening increases and the one from texture strengthening occurs in the extruded-T5 alloys, and both of these two contributors lead to the increase of TYS in the extruded-T5 alloys, compared to cast-T6 alloys. Although the proportion of strengthening contribution due to precipitation strengthening decreases, it keeps as the largest contributor.
At the temperature lower than 200oC, the outstanding precipitation strengthening in aged Mg-Gd-Y alloys is mainly attributed to the favorable shape and orientation ofβ′precipitates which make them obstruct effectively the basal dislocation slip, the bigger volume fraction of the precipitates, perfectly coherent interface between the precipitates and the matrix, and good thermal stability of the precipitates. Coherent strengthening and Orowan mechanism both contribute to the peak hardness of aged Mg-Gd-Y alloys. However, Orowan mechanism becomes the main contributor to TYS at the over-aged stage.
A small addition of 0.4-0.6wt% Ca improves the creep and corrosion resistances of Mg-Gd-Y-Zr alloys but severely deteriorates the elongation, which is related to the presence of the Ca segregation at grain boundaries.
引文
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