Al-Zn-Mg-Cu合金热处理工艺及组织性能研究
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摘要
Al-Zn-Mg-Cu铝合金具有比强度高、价格低廉及加工性能良好等优点,被广泛用作航空航天领域的结构材料。近年来,随着航空航天事业的不断发展和国防现代化建设步伐的加快,我国对大型飞机的需求也日益紧迫,这就急需高性能的Al-Zn-Mg-Cu铝合金板材。本论文工作依托国防重点攻关项目《××××铝合金板材研制》,以航空用Al-Zn-Mg-Cu合金为研究对象,采用热加工模拟、室温拉伸、硬度、电导率、疲劳裂纹扩展、剥落腐蚀和应力腐蚀等实验方法及金相显微镜(OM)、电子显微镜(SEM、 TEM)、差热分析(DSC)、背散射电子衍射(EBSD)和X射线衍射(XRD、 Micro-XRD)等检测手段,系统研究了合金在不同热加工工艺和热处理条件下的显微组织、性能及其变化规律,并从理论上进行了分析与讨论,旨在为制定和优化合金热加工工艺、热处理工艺以及改善合金性能提供理论依据。
试验结果表明,铸态Al-Zn-Mg-Cu合金中Zn、Mg和Cu等元素偏聚在晶界处,形成了含有η相(MgZn2)、S相(Al2CuMg)和T相(AlZnMgCu)的低熔点共晶组织,该共晶组织的熔点为475.2℃。铸态合金经过465℃/24h均匀化处理后,低熔点共晶相基本被消除,而Al7Cu2Fe相则很难溶入到基体中。Cr元素在Al-Zn-Mg-Cu合金中形成了Al18Mg3Cr2相(E相),主要分布在基体中。经过均匀化处理后,E相没有发生明显变化。Er元素在Al-Cu-Mg-Er合金凝固过程中主要偏聚在晶界处,形成了Al8Cu4Er三元相。Al8Cu4Er相在均匀化过程中不能被消除,而是在轧制过程中被轧碎,沿晶界呈条带状分布。本文发现在Al-Cu-Mg-Er合金基体中有少量的二次析出的Al3Er相存在。
研究了Al-Zn-Mg-Cu合金的热变形行为,结果表明变形温度低于390℃时,合金组织仅发生动态回复；变形温度为420℃时,合金组织已发生了动态再结晶。合金适宜的热加工温度范围为330℃～400℃,变形速率为O.1s-1。根据热变形结果,计算出合金的热变形激活能为192.6KJ/mol,并据此构建了合金的热变形本构方程。确定了合金的最佳T761工艺为125℃/3h+170℃/10h,在此条件下,合金的抗拉强度、屈服强度、延伸率和电导率分别为498MPa、429MPa、12.1%和40.2%IACS。确定了合金的最佳回归再时效工艺为120℃/25h+190℃/10min+120℃/25h,在此条件下,合金的抗拉强度、屈服强度、延伸率和电导率分别为554MPa、507MPa、16.5%和35.4%IACS,与T6态合金性能相当。采用上述T761工艺处理的Al-Zn-Mg-Cu合金板材满足AMS-4085B对合金性能的要求,并已经成功应用于某大型飞机上。
研究发现,在应力比为0.1、加载频率为10Hz条件下,RRA态Al-Zn-Mg-Cu合金的疲劳裂纹扩展速率要小于T761态合金。RRA态合金晶内析出了位错可切割的GP区和η’相以及在晶界附近形成了较窄的无析出带,导致在裂纹扩展的Paris区形成了大量的二次裂纹使裂纹扩展路径复杂,从而提高了疲劳裂纹的扩展抗力。而T761态合金晶内则析出了位错不可切割的粗大η’相以及在晶界附近形成了较宽的无析出带,因而在裂纹扩展的Paris区有沿晶扩展的二次裂纹存在,使疲劳裂纹扩展抗力降低。
通过对T761态Al-Zn-Mg-Cu合金疲劳短裂纹扩展的EBSD分析发现,短裂纹扩展路径的偏折是由于相邻晶粒的取向差较大而造成的。对T761态Al-Zn-Mg-Cu合金疲劳长裂纹尖端的塑性区进行了微区XRD分析,没有发现疲劳裂纹尖端处的第二相回溶现象。EBSD观察发现,在长裂纹扩展时,裂纹的实际扩展路径倾向于沿晶界扩展。这表明,在疲劳裂纹尖端处,位于晶界的η平衡相不因位错的往复滑移而发生回溶。
阐明了T761态Al-Zn-Mg-Cu合金疲劳裂纹的扩展方式,并提出了裂纹扩展方式转变的晶体学模型。随着疲劳裂纹长度的增加,T761态Al-Zn-Mg-Cu合金的裂纹扩展方式由短裂纹的穿晶扩展转变为长裂纹的沿晶扩展为主。短裂纹扩展受剪切机制的控制,因此在该阶段无析出带对裂纹扩展的影响可以忽略。而长裂纹扩展是受双滑移机制控制的,由于晶界处存在较宽的无析出带,是合金微观结构中的薄弱部位,诱导长裂纹沿晶界扩展。
研究表明,RRA态、T761态和T6态Al-Zn-Mg-Cu合金的剥落腐蚀性能依次降低,剥落腐蚀等级分别为EA、EA和EB。T6态合金中连续分布的晶界η相促进了阳极溶解通道的形成,对剥落腐蚀性能不利。而T761态和RRA态合金中断续分布的晶界η相则提高了合金的剥落腐蚀性能。与T761态合金相比,RRA态合金晶界η相中较高的铜含量减缓了晶界η相的阳极溶解,导致合金的剥落腐蚀性能提高。还研究了不同RRA工艺处理合金在482MPa应力腐蚀环境下的断裂时间,结果表明经过120℃/25h+190℃/10min+120℃/25h时效处理的合金由于具有较高的屈服强度,为507MPa,而使得断裂时间较长,达到260h。
Al-Zn-Mg-Cu aluminum alloys have been widely used as the structural materials in the aerospace industry due to their high specific strength, low price and excellent processability. In recent years, with the rapid development of the aerospace industry and the national defense modernization, there is a great demand for large aircraft, which make it desirable to use Al-Zn-Mg-Cu plates with high properties. This work is based on "National Defense Key Research Project" titled as "Research of X X X X aluminum alloy plate". In order to provide conferences for the optimization of hot working and heat treatment processes and the improvement of properties, the heat treatment process, microstructure and properties of an Al-Zn-Mg-Cu alloy were systematically studied by means of hot working simulation, tension, microhardness, conductance measurement, fatigue crack propagation test, exfoliation corrosion test, stress corrosion test, optical microscopy (OM), electronic microscopy (SEM, TEM), differential scanning calorimetry (DSC), electron back-scattered diffraction (EBSD) and X-ray diffraction (XRD, Micro-XRD). In addition, the results were analyzed and discussed theoretically.
The experimental results indicated that the eutectic structures consisting of η phase (MgZn2), S phase (Al2CuMg) and T phase (AlZnMgCu) were found to form in the grain boundary of as cast Al-Zn-Mg-Cu alloy due to the segregation of Zn, Mg and Cu, and the melting point of the eutectic structure was475.2℃. The primary eutectic structure of the alloy homogenized at465℃for24h dissolved into the matrix and the residual particles were Al7Cu2Fe. The triangular Al18Mg3Cr2phase (E phase) which distributed mainly in matrix formed during solidification. Moreover, there were no obvious changes on the E precipitates during the homogenization. It was found that the major existing form of erbium in Al-Cu-Mg-Er alloy was Al8Cu4Er phase, which formed in grain boundary during solidification. Al8Cu4Er particles could not be eliminated during homogenization and were crushed up along with the grain boundaries after hot rolling. A few secondary Al3Er precipitates were observed in the matrix in this investigation.
The hot deformation behavior of Al-Zn-Mg-Cu alloy was studied and the results showed that when the deformation temperature was below390℃, the dominative restoring mechanism was dynamic recovery. However, when the temperature was420℃, dynamic recrystallization has occurred. The suitable hot working temperature range of the alloy was330℃-400℃with a strain rate of0.1s-1. Based on the results, the hot deformation activation energy was calculated which was192.6KJ/mol, and the constitutive equation was also given. The optimal T761ageing treatment was125℃/3h+170℃/10h. Under this condition, the tensile strength, yield strength, elongation and conductivity were498MPa,429MPa,12.1%and40.2%IACS, respectively. The optimal RRA treatment of the alloy was120℃/25h+190℃/10min+120℃/25h, and under this condition the tensile strength, yield strength, elongation and conductivity were554MPa,507MPa,16.5%and35.4%IACS, which were closely comparable to those of T6-treated alloy. The sheet of T761-treated Al-Zn-Mg-Cu alloy could meet the needs of AMS-4085B, and was satisfactorily applied to a large aircraft.
It was found that the fatigue crack propagation (FCP) rate of RRA-treated Al-Zn-Mg-Cu alloy was slower than that of T761-treated alloy at a stress ratio of0.1with a sine-wave loading frequency of lOHz. The formation of more secondary cracks led to a more complex crack propagation path, thereby reducing the fatigue crack driving force and enhancing the crack propagation resistance in the RRA-treated alloy with shearable precipitates. For the T761-treated alloy with coarse η'precipitates and wide PFZs, secondary cracks were found to propagate along the grain boundaries, which was associated with the higher FCP rate.
Short crack propagation of T761-treated Al-Zn-Mg-Cu alloy was investigated by EBSD analysis and it was found that the formation of zigzag crack was ascribed to the high misorientation of adjacent grains. The plastic deformation zone surrounding the long crack tip of T761-treated Al-Zn-Mg-Cu alloy was analyzed by Micro-XRD, and the dissolution of precipitates was not observed. The results of EBSD analysis showed that long crack tended to propagate along grain boundary, therefore, the dissolution of n precipitates on grain boundary was not found on the condition of reciprocating motion of dislocations.
The fatigue crack propagation mode of T761-treated Al-Zn-Mg-Cu alloy was revealed and a crystallographic model of the transition of crack propagation mode was given. It was showed that short crack tended to transgranularly propagate in the alloy, whereas intergranular long crack propagation was observed. Short crack propagation was predominated by single shear, and the effect of PFZs on the crack propagation could be ignored in this stage. However, long crack propagation was predominated by duplex slip mechanism. Moreover, the presence of wide PFZs accelerated the growth of long crack which propagated along the grain boundary.
The research showed that the exfoliation corrosion resistances of RRA-treated, T761-treated and T6-treated Al-Zn-Mg-Cu alloys were decreased in turn, and the degrees of exfoliation corrosion were EA-, EA and EB, respectively. The enhanced corrosion resistance of T761-treated alloy and RRA-treated alloy was ascribed to the discontinued GBPs, whereas the continued GBPs presented in T6-treated alloy resulted in the severe corrosion. Compared with the T761-treated alloy, the anode dissolution of GBPs in RRA-treated alloy was abated due to the higher copper content of GBPs, which led to the enhanced corrosion resistance of RRA-treated alloy. The stress corrosion resistance of the alloy under various RRA treatments on the corrosion condition of482MPa was also studied and the results showed that the alloy treated at120℃/25h+190℃/10min+120℃/25h had the longest fracture time of260h due to the higher yield strength of507MPa.
试验结果表明,铸态Al-Zn-Mg-Cu合金中Zn、Mg和Cu等元素偏聚在晶界处,形成了含有η相(MgZn2)、S相(Al2CuMg)和T相(AlZnMgCu)的低熔点共晶组织,该共晶组织的熔点为475.2℃。铸态合金经过465℃/24h均匀化处理后,低熔点共晶相基本被消除,而Al7Cu2Fe相则很难溶入到基体中。Cr元素在Al-Zn-Mg-Cu合金中形成了Al18Mg3Cr2相(E相),主要分布在基体中。经过均匀化处理后,E相没有发生明显变化。Er元素在Al-Cu-Mg-Er合金凝固过程中主要偏聚在晶界处,形成了Al8Cu4Er三元相。Al8Cu4Er相在均匀化过程中不能被消除,而是在轧制过程中被轧碎,沿晶界呈条带状分布。本文发现在Al-Cu-Mg-Er合金基体中有少量的二次析出的Al3Er相存在。
研究了Al-Zn-Mg-Cu合金的热变形行为,结果表明变形温度低于390℃时,合金组织仅发生动态回复；变形温度为420℃时,合金组织已发生了动态再结晶。合金适宜的热加工温度范围为330℃～400℃,变形速率为O.1s-1。根据热变形结果,计算出合金的热变形激活能为192.6KJ/mol,并据此构建了合金的热变形本构方程。确定了合金的最佳T761工艺为125℃/3h+170℃/10h,在此条件下,合金的抗拉强度、屈服强度、延伸率和电导率分别为498MPa、429MPa、12.1%和40.2%IACS。确定了合金的最佳回归再时效工艺为120℃/25h+190℃/10min+120℃/25h,在此条件下,合金的抗拉强度、屈服强度、延伸率和电导率分别为554MPa、507MPa、16.5%和35.4%IACS,与T6态合金性能相当。采用上述T761工艺处理的Al-Zn-Mg-Cu合金板材满足AMS-4085B对合金性能的要求,并已经成功应用于某大型飞机上。
研究发现,在应力比为0.1、加载频率为10Hz条件下,RRA态Al-Zn-Mg-Cu合金的疲劳裂纹扩展速率要小于T761态合金。RRA态合金晶内析出了位错可切割的GP区和η’相以及在晶界附近形成了较窄的无析出带,导致在裂纹扩展的Paris区形成了大量的二次裂纹使裂纹扩展路径复杂,从而提高了疲劳裂纹的扩展抗力。而T761态合金晶内则析出了位错不可切割的粗大η’相以及在晶界附近形成了较宽的无析出带,因而在裂纹扩展的Paris区有沿晶扩展的二次裂纹存在,使疲劳裂纹扩展抗力降低。
通过对T761态Al-Zn-Mg-Cu合金疲劳短裂纹扩展的EBSD分析发现,短裂纹扩展路径的偏折是由于相邻晶粒的取向差较大而造成的。对T761态Al-Zn-Mg-Cu合金疲劳长裂纹尖端的塑性区进行了微区XRD分析,没有发现疲劳裂纹尖端处的第二相回溶现象。EBSD观察发现,在长裂纹扩展时,裂纹的实际扩展路径倾向于沿晶界扩展。这表明,在疲劳裂纹尖端处,位于晶界的η平衡相不因位错的往复滑移而发生回溶。
阐明了T761态Al-Zn-Mg-Cu合金疲劳裂纹的扩展方式,并提出了裂纹扩展方式转变的晶体学模型。随着疲劳裂纹长度的增加,T761态Al-Zn-Mg-Cu合金的裂纹扩展方式由短裂纹的穿晶扩展转变为长裂纹的沿晶扩展为主。短裂纹扩展受剪切机制的控制,因此在该阶段无析出带对裂纹扩展的影响可以忽略。而长裂纹扩展是受双滑移机制控制的,由于晶界处存在较宽的无析出带,是合金微观结构中的薄弱部位,诱导长裂纹沿晶界扩展。
研究表明,RRA态、T761态和T6态Al-Zn-Mg-Cu合金的剥落腐蚀性能依次降低,剥落腐蚀等级分别为EA、EA和EB。T6态合金中连续分布的晶界η相促进了阳极溶解通道的形成,对剥落腐蚀性能不利。而T761态和RRA态合金中断续分布的晶界η相则提高了合金的剥落腐蚀性能。与T761态合金相比,RRA态合金晶界η相中较高的铜含量减缓了晶界η相的阳极溶解,导致合金的剥落腐蚀性能提高。还研究了不同RRA工艺处理合金在482MPa应力腐蚀环境下的断裂时间,结果表明经过120℃/25h+190℃/10min+120℃/25h时效处理的合金由于具有较高的屈服强度,为507MPa,而使得断裂时间较长,达到260h。
Al-Zn-Mg-Cu aluminum alloys have been widely used as the structural materials in the aerospace industry due to their high specific strength, low price and excellent processability. In recent years, with the rapid development of the aerospace industry and the national defense modernization, there is a great demand for large aircraft, which make it desirable to use Al-Zn-Mg-Cu plates with high properties. This work is based on "National Defense Key Research Project" titled as "Research of X X X X aluminum alloy plate". In order to provide conferences for the optimization of hot working and heat treatment processes and the improvement of properties, the heat treatment process, microstructure and properties of an Al-Zn-Mg-Cu alloy were systematically studied by means of hot working simulation, tension, microhardness, conductance measurement, fatigue crack propagation test, exfoliation corrosion test, stress corrosion test, optical microscopy (OM), electronic microscopy (SEM, TEM), differential scanning calorimetry (DSC), electron back-scattered diffraction (EBSD) and X-ray diffraction (XRD, Micro-XRD). In addition, the results were analyzed and discussed theoretically.
The experimental results indicated that the eutectic structures consisting of η phase (MgZn2), S phase (Al2CuMg) and T phase (AlZnMgCu) were found to form in the grain boundary of as cast Al-Zn-Mg-Cu alloy due to the segregation of Zn, Mg and Cu, and the melting point of the eutectic structure was475.2℃. The primary eutectic structure of the alloy homogenized at465℃for24h dissolved into the matrix and the residual particles were Al7Cu2Fe. The triangular Al18Mg3Cr2phase (E phase) which distributed mainly in matrix formed during solidification. Moreover, there were no obvious changes on the E precipitates during the homogenization. It was found that the major existing form of erbium in Al-Cu-Mg-Er alloy was Al8Cu4Er phase, which formed in grain boundary during solidification. Al8Cu4Er particles could not be eliminated during homogenization and were crushed up along with the grain boundaries after hot rolling. A few secondary Al3Er precipitates were observed in the matrix in this investigation.
The hot deformation behavior of Al-Zn-Mg-Cu alloy was studied and the results showed that when the deformation temperature was below390℃, the dominative restoring mechanism was dynamic recovery. However, when the temperature was420℃, dynamic recrystallization has occurred. The suitable hot working temperature range of the alloy was330℃-400℃with a strain rate of0.1s-1. Based on the results, the hot deformation activation energy was calculated which was192.6KJ/mol, and the constitutive equation was also given. The optimal T761ageing treatment was125℃/3h+170℃/10h. Under this condition, the tensile strength, yield strength, elongation and conductivity were498MPa,429MPa,12.1%and40.2%IACS, respectively. The optimal RRA treatment of the alloy was120℃/25h+190℃/10min+120℃/25h, and under this condition the tensile strength, yield strength, elongation and conductivity were554MPa,507MPa,16.5%and35.4%IACS, which were closely comparable to those of T6-treated alloy. The sheet of T761-treated Al-Zn-Mg-Cu alloy could meet the needs of AMS-4085B, and was satisfactorily applied to a large aircraft.
It was found that the fatigue crack propagation (FCP) rate of RRA-treated Al-Zn-Mg-Cu alloy was slower than that of T761-treated alloy at a stress ratio of0.1with a sine-wave loading frequency of lOHz. The formation of more secondary cracks led to a more complex crack propagation path, thereby reducing the fatigue crack driving force and enhancing the crack propagation resistance in the RRA-treated alloy with shearable precipitates. For the T761-treated alloy with coarse η'precipitates and wide PFZs, secondary cracks were found to propagate along the grain boundaries, which was associated with the higher FCP rate.
Short crack propagation of T761-treated Al-Zn-Mg-Cu alloy was investigated by EBSD analysis and it was found that the formation of zigzag crack was ascribed to the high misorientation of adjacent grains. The plastic deformation zone surrounding the long crack tip of T761-treated Al-Zn-Mg-Cu alloy was analyzed by Micro-XRD, and the dissolution of precipitates was not observed. The results of EBSD analysis showed that long crack tended to propagate along grain boundary, therefore, the dissolution of n precipitates on grain boundary was not found on the condition of reciprocating motion of dislocations.
The fatigue crack propagation mode of T761-treated Al-Zn-Mg-Cu alloy was revealed and a crystallographic model of the transition of crack propagation mode was given. It was showed that short crack tended to transgranularly propagate in the alloy, whereas intergranular long crack propagation was observed. Short crack propagation was predominated by single shear, and the effect of PFZs on the crack propagation could be ignored in this stage. However, long crack propagation was predominated by duplex slip mechanism. Moreover, the presence of wide PFZs accelerated the growth of long crack which propagated along the grain boundary.
The research showed that the exfoliation corrosion resistances of RRA-treated, T761-treated and T6-treated Al-Zn-Mg-Cu alloys were decreased in turn, and the degrees of exfoliation corrosion were EA-, EA and EB, respectively. The enhanced corrosion resistance of T761-treated alloy and RRA-treated alloy was ascribed to the discontinued GBPs, whereas the continued GBPs presented in T6-treated alloy resulted in the severe corrosion. Compared with the T761-treated alloy, the anode dissolution of GBPs in RRA-treated alloy was abated due to the higher copper content of GBPs, which led to the enhanced corrosion resistance of RRA-treated alloy. The stress corrosion resistance of the alloy under various RRA treatments on the corrosion condition of482MPa was also studied and the results showed that the alloy treated at120℃/25h+190℃/10min+120℃/25h had the longest fracture time of260h due to the higher yield strength of507MPa.
引文
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