热处理对FGH95镍基合金组织结构与蠕变行为的影响
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摘要
通过对不同温度热等静压(HIP)制备的FGH95粉末镍基合金进行不同工艺热处理、蠕变性能测试及组织形貌观察,研究热处理工艺对合金组织结构和蠕变性能的影响；通过热力学计算预测合金中碳化物的析出行为,通过XRD曲线分析,研究了热处理工艺与合金中γ'、Y两相晶格应变、蠕变性能之间的依赖关系,通过微观形貌观察及衍衬分析,研究合金在蠕变期间的微观变形特征与断裂机制。得出如下结论：
FGH95镍基合金的组织结构主要由Y基体、γ'相及碳化物组成,采用不同工艺热处理可使合金中γ'相、碳化物具有不同的尺寸、形态及分布,其中,γ'相按析出特征分为一次、二次和三次γ'相,采用电解萃取方法测定出合金中γ'相的体积分数约为47～48%,且合金中可析出(Nb,Ti)C和(Cr,W)23C6碳化物。随着热等静压(HIP)温度升高,合金中原始颗粒边界(PPB)区域的一次粗大γ'相尺寸和数量减小,不同温度热等静压(HIP)合金经1155℃固溶、550℃盐浴及时效处理,晶界处仅存在少量粗大γ'相,并有较多粒状碳化物在晶内及晶界弥散析出；1120℃CHIP合金经1150℃固溶和时效处理后,颗粒边界区域仍存在较多粗大γ'相和γ'相贫化区,随固溶温度提高,粗大γ'相及γ'相贫化区数量减少。当固溶温度提高到1160℃,合金中粗大γ'相完全溶解,γ'相贫化区消失,其获得的组织是高体积分数细小γ'相在晶内弥散分布,并有粒状碳化物在晶内及沿晶界不连续析出；经1165℃固溶,合金的晶粒尺寸明显长大,并有片状碳化物膜沿晶界连续析出。热力学计算确定出NbC、TiC、Cr23C6和W23C6碳化物自合金中析出的温度分别为1353K、1090K、1009K和1536K；在583℃计算出合金中析出TiC、NbC、WC碳化物的相变驱动力为：-9.77J·mol-1、-11.23J·mol-1和13.48J·mol-1,形核驱动力为：-80.69J·mol-1、-59.95J·mol-1、-36.01Jmol-1;与元素Nb、Ti相比,Cr是弱碳化物形成元素,其中,析出TiC、NbC有较大的形核驱动力,是使合金中析出TiC、NbC或(Nb、Ti)C相的主要原因。随热等静压(HIP)温度提高,合金中γ、Y'两相的晶格常数值略有增加,而错配度逐渐减小,经完全热处理后,合金错配度由0.3322%减少到0.2838%；随固溶温度提高,合金中Y、γ'两相的晶格常数和错配度逐渐减小；与油浴处理合金相比,盐浴处理合金中Y、γ'两相具有较高的晶格错配度,是使合金具有较高蠕变抗力的原因之一。长期时效期间,合金中细小γ'相尺寸及晶格常数随着时效温度和时间的增加略有增大,而合金中γ、γ'两相间的错配度略有减小；与时效时间相比,时效温度对合金中γ、γ'两相的尺寸及晶格常数、错配度具有较大影响。
不同温度热等静压(HIP)合金经1155℃固溶和550℃盐浴处理后,1180℃HIP合金具有较好的蠕变性能；随着固溶温度的升高,合金中原始颗粒边界区域的尺寸逐渐减小,且当碳化物沿晶界呈粒状分布时,合金具有较长的蠕变寿命,而当碳化物沿晶界呈片状连续析出时,合金的蠕变寿命大幅度降低；随时效温度的提高和时间延长,合金的蠕变寿命逐渐降低。在实验条件下,分别测出经工艺16、工艺5和工艺15制备合金的蠕变激活能分别为：Q1=597-2kJ·mol-、Q2=578.6kJ·mol-1和Q3=480.3kJ·mol-1。
“油浴”处理合金的变形特征是位错在晶内发生单取向或双取向滑移；“盐浴”处理合金的变形机制是微孪晶变形、位错剪切或绕过Y'相,其中,当<110>位错切入Y或γ'相时,可分解形成1/6<112>肖克莱不全位错(或1/3<112>超肖克莱不全位错)+层错的位错组态；且合金中的微孪晶是由1/2<110>位错在γ基体中分解,形成“两1/6<112>肖克莱不全位错+层错”所组成。同时,蠕变期间,合金中可形成六角形或四边形的位错网络,且在γ'/γ两相界面形成的位错网可释放晶格错配应力,减缓应力集中,提高合金的蠕变抗力。合金在蠕变后期,“油浴”处理合金中晶内可发生双取向滑移,且在滑移迹线区域有细小碳化物析出,并在γ'相贫化区发生晶界组织的碎化；而在盐浴处理合金中晶内发生单取向滑移,随着蠕变进行,位错在晶界处塞积,并引起应力集中,致使裂纹在晶界处萌生并沿晶界扩展是合金的蠕变断裂机制。
In the paper, by means of heat treatment at different regimes, creep propertiesmeasurement and microstructure observation, the effects of heat treatment on microstructureand creep behaviors of FGH95superalloy prepared by hot isostatic pressing (HIP) at differenttemperatures have been investigated. The precipitating behaviors of the carbides are forecastedby thermodynamic method. By means of XRD curves measurement, the dependence of the heattreatment regimes on the J c laJttice strain and creep properties of alloy is investigated. In thefurther, the deformation features and fracture mechanisms of the alloy during creep are analyzedby microstructure observation and diffraction contrast analysis of dislocations configuration.
The results show that the microstructure of FGH95alloy consists of Jmatrix, J cphase andcarbides, and the various sizes, morphologies and distribution of them may be obtained bydifferent heat treatment regimes. Thereinto, the J cparticles may be divided into primary,secondary and tertiary J cphase according to the precipitating features, and the volume fractionof J cphase is measured to be47~48%by electrolytic extraction method, and the precipitatedcarbides are identified as (Nb, Ti)C and (Cr, W)23C6phases. As the HIP temperature increases,the amount and size of coarser J cphase in the previous particle boundaries (PPB) decrease.When the alloy is solution treated at1155and cooled in molten salt at550, a few ofcoarse J cphase distributes in the boundary regions, and some granular carbide precipitatedispersedly in the grains and boundaries. After the1120HIP alloy is solution treated at1150
and aged, some coarse J cprecipitates and J-cfree zone appear in the previous particleboundaries (PPB), and the amount of them decreases with the solution temperature increasing.When the solution temperature enhances to1160, the coarse J cphase is completely dissolvedto disappear the J-cfree zone. As the solution temperature increases to1165, the grain size ofthe alloy increase obviously and the carbide in the form of the films precipitates along theboundaries. By means of thermodynamic method, the precipitating temperatures of NbC, TiC, Cr23C6and W23C6carbides are defined to be1353K,1090K,1009K and1536K, respectively.Moreover, the driving force values of the phase transformation for TiC, NbC and WC carbidesin the alloy at583are calculated to be-9.77J mol-1,-11.23J m ol-1and-13.48J m ol-1,respectively, and the nucleating driving force of TiC, NbC and WC carbides are calculated to be-80.69J mol-1,-59.95J mol-1and-36.01J mol-1, respectively. Thereinto, the fact that TiC, NbCcarbides possess the bigger nucleating driving force is thought to be the main reasons for theTiC, NbC and (Nb, Ti)C carbides precipitating in the alloy.
As the HIP temperature enhances, the parameters of J and J cphases in alloy increaseslightly, and the misfits increase gradually. After full heat treatment, the misfits between Jand J cphases decrease from0.3322%to0.2838%. As the solution temperature increases, theparameters and misfits of J and J cphase decrease gradually. Compared with the “oil cooling”alloy, the “molten salt cooling” alloy possesses a bigger misfit, which is one of the reasons forthe alloy having a better creep resistance. During long-term aging, the size and parameters of J cphase increase with the aging temperature and time, while the misfit in alloy decreases slightly.And compared to the aging time, the aging temperature has a bigger effect on the size of J cphase and misfits in the alloy.
After the different temperature HIP alloys are solution treated at1155and cooled inmolten salt at550,1180HIP alloy possesses better creep property. As the solutiontemperature enhances, the size of PPB regions decreases gradually, and the granular carbidesare precipitated along boundaries, which results in the better creep resistance of alloy. But thecreep lifetime of the alloy decreases obviously once the carbide films appears along boundaries.In the range of the applied stresses and temperatures, the creep activation energy of the alloyprepared by technique16,5and15are calculated to be Q1=597.2kJ mol-1, Q2=578.6kJ m ol-1and Q3=480.3kJ m ol-1, respectively.
The deformation features of the “oil bath” treated alloy during creep are that the single ordouble orientations slipping of dislocations activated in the J matrix. And the deformationmechanism of the “salt bath” treated alloy during creep is that dislocations slipping in the Jmatrix or shearing into the J cphase and the micro-twinning deformation, Thereinto, the <110>dislocation shearing into Jor J cphases may decompose to form the configuration of16<112>Shockley partials or13<112> super-Shockley partials plus stacking fault. The activating micro-twinning in the alloy during creep consists of16<112> Shockley partials and stacking fault, which is attributed to the decomposition of1
2<110> dislocation shearing into J or J cphases.Moreover, the dislocation networks with hexagon and quadrangle may be formed in the alloyduring creep. And the dislocation networks formed in the J/c J interface may release the latticestrain energy to decrease the stress concentration for enhancing the creep resistance of the alloy.In the later stage of creep, the slipping traces with double oriented feature appear within thegrains, and the fine carbides precipitate along the slipping traces in the “oil bath” treated alloy,and the fine grains appear in the free-J cphase zone of the boundary. However, the slippingtraces with single oriented feature appear in the “salt bath” treated alloy in the later stage ofcreep, and significant amount of dislocations are piled up in the boundaries regions to bring thestress concentration, which may results in the initiation and propagation of the cracks alongboundaries, this is thought to be the fracture mechanism of the alloy during creep.
FGH95镍基合金的组织结构主要由Y基体、γ'相及碳化物组成,采用不同工艺热处理可使合金中γ'相、碳化物具有不同的尺寸、形态及分布,其中,γ'相按析出特征分为一次、二次和三次γ'相,采用电解萃取方法测定出合金中γ'相的体积分数约为47～48%,且合金中可析出(Nb,Ti)C和(Cr,W)23C6碳化物。随着热等静压(HIP)温度升高,合金中原始颗粒边界(PPB)区域的一次粗大γ'相尺寸和数量减小,不同温度热等静压(HIP)合金经1155℃固溶、550℃盐浴及时效处理,晶界处仅存在少量粗大γ'相,并有较多粒状碳化物在晶内及晶界弥散析出；1120℃CHIP合金经1150℃固溶和时效处理后,颗粒边界区域仍存在较多粗大γ'相和γ'相贫化区,随固溶温度提高,粗大γ'相及γ'相贫化区数量减少。当固溶温度提高到1160℃,合金中粗大γ'相完全溶解,γ'相贫化区消失,其获得的组织是高体积分数细小γ'相在晶内弥散分布,并有粒状碳化物在晶内及沿晶界不连续析出；经1165℃固溶,合金的晶粒尺寸明显长大,并有片状碳化物膜沿晶界连续析出。热力学计算确定出NbC、TiC、Cr23C6和W23C6碳化物自合金中析出的温度分别为1353K、1090K、1009K和1536K；在583℃计算出合金中析出TiC、NbC、WC碳化物的相变驱动力为：-9.77J·mol-1、-11.23J·mol-1和13.48J·mol-1,形核驱动力为：-80.69J·mol-1、-59.95J·mol-1、-36.01Jmol-1;与元素Nb、Ti相比,Cr是弱碳化物形成元素,其中,析出TiC、NbC有较大的形核驱动力,是使合金中析出TiC、NbC或(Nb、Ti)C相的主要原因。随热等静压(HIP)温度提高,合金中γ、Y'两相的晶格常数值略有增加,而错配度逐渐减小,经完全热处理后,合金错配度由0.3322%减少到0.2838%；随固溶温度提高,合金中Y、γ'两相的晶格常数和错配度逐渐减小；与油浴处理合金相比,盐浴处理合金中Y、γ'两相具有较高的晶格错配度,是使合金具有较高蠕变抗力的原因之一。长期时效期间,合金中细小γ'相尺寸及晶格常数随着时效温度和时间的增加略有增大,而合金中γ、γ'两相间的错配度略有减小；与时效时间相比,时效温度对合金中γ、γ'两相的尺寸及晶格常数、错配度具有较大影响。
不同温度热等静压(HIP)合金经1155℃固溶和550℃盐浴处理后,1180℃HIP合金具有较好的蠕变性能；随着固溶温度的升高,合金中原始颗粒边界区域的尺寸逐渐减小,且当碳化物沿晶界呈粒状分布时,合金具有较长的蠕变寿命,而当碳化物沿晶界呈片状连续析出时,合金的蠕变寿命大幅度降低；随时效温度的提高和时间延长,合金的蠕变寿命逐渐降低。在实验条件下,分别测出经工艺16、工艺5和工艺15制备合金的蠕变激活能分别为：Q1=597-2kJ·mol-、Q2=578.6kJ·mol-1和Q3=480.3kJ·mol-1。
“油浴”处理合金的变形特征是位错在晶内发生单取向或双取向滑移；“盐浴”处理合金的变形机制是微孪晶变形、位错剪切或绕过Y'相,其中,当<110>位错切入Y或γ'相时,可分解形成1/6<112>肖克莱不全位错(或1/3<112>超肖克莱不全位错)+层错的位错组态；且合金中的微孪晶是由1/2<110>位错在γ基体中分解,形成“两1/6<112>肖克莱不全位错+层错”所组成。同时,蠕变期间,合金中可形成六角形或四边形的位错网络,且在γ'/γ两相界面形成的位错网可释放晶格错配应力,减缓应力集中,提高合金的蠕变抗力。合金在蠕变后期,“油浴”处理合金中晶内可发生双取向滑移,且在滑移迹线区域有细小碳化物析出,并在γ'相贫化区发生晶界组织的碎化；而在盐浴处理合金中晶内发生单取向滑移,随着蠕变进行,位错在晶界处塞积,并引起应力集中,致使裂纹在晶界处萌生并沿晶界扩展是合金的蠕变断裂机制。
In the paper, by means of heat treatment at different regimes, creep propertiesmeasurement and microstructure observation, the effects of heat treatment on microstructureand creep behaviors of FGH95superalloy prepared by hot isostatic pressing (HIP) at differenttemperatures have been investigated. The precipitating behaviors of the carbides are forecastedby thermodynamic method. By means of XRD curves measurement, the dependence of the heattreatment regimes on the J c laJttice strain and creep properties of alloy is investigated. In thefurther, the deformation features and fracture mechanisms of the alloy during creep are analyzedby microstructure observation and diffraction contrast analysis of dislocations configuration.
The results show that the microstructure of FGH95alloy consists of Jmatrix, J cphase andcarbides, and the various sizes, morphologies and distribution of them may be obtained bydifferent heat treatment regimes. Thereinto, the J cparticles may be divided into primary,secondary and tertiary J cphase according to the precipitating features, and the volume fractionof J cphase is measured to be47~48%by electrolytic extraction method, and the precipitatedcarbides are identified as (Nb, Ti)C and (Cr, W)23C6phases. As the HIP temperature increases,the amount and size of coarser J cphase in the previous particle boundaries (PPB) decrease.When the alloy is solution treated at1155and cooled in molten salt at550, a few ofcoarse J cphase distributes in the boundary regions, and some granular carbide precipitatedispersedly in the grains and boundaries. After the1120HIP alloy is solution treated at1150
and aged, some coarse J cprecipitates and J-cfree zone appear in the previous particleboundaries (PPB), and the amount of them decreases with the solution temperature increasing.When the solution temperature enhances to1160, the coarse J cphase is completely dissolvedto disappear the J-cfree zone. As the solution temperature increases to1165, the grain size ofthe alloy increase obviously and the carbide in the form of the films precipitates along theboundaries. By means of thermodynamic method, the precipitating temperatures of NbC, TiC, Cr23C6and W23C6carbides are defined to be1353K,1090K,1009K and1536K, respectively.Moreover, the driving force values of the phase transformation for TiC, NbC and WC carbidesin the alloy at583are calculated to be-9.77J mol-1,-11.23J m ol-1and-13.48J m ol-1,respectively, and the nucleating driving force of TiC, NbC and WC carbides are calculated to be-80.69J mol-1,-59.95J mol-1and-36.01J mol-1, respectively. Thereinto, the fact that TiC, NbCcarbides possess the bigger nucleating driving force is thought to be the main reasons for theTiC, NbC and (Nb, Ti)C carbides precipitating in the alloy.
As the HIP temperature enhances, the parameters of J and J cphases in alloy increaseslightly, and the misfits increase gradually. After full heat treatment, the misfits between Jand J cphases decrease from0.3322%to0.2838%. As the solution temperature increases, theparameters and misfits of J and J cphase decrease gradually. Compared with the “oil cooling”alloy, the “molten salt cooling” alloy possesses a bigger misfit, which is one of the reasons forthe alloy having a better creep resistance. During long-term aging, the size and parameters of J cphase increase with the aging temperature and time, while the misfit in alloy decreases slightly.And compared to the aging time, the aging temperature has a bigger effect on the size of J cphase and misfits in the alloy.
After the different temperature HIP alloys are solution treated at1155and cooled inmolten salt at550,1180HIP alloy possesses better creep property. As the solutiontemperature enhances, the size of PPB regions decreases gradually, and the granular carbidesare precipitated along boundaries, which results in the better creep resistance of alloy. But thecreep lifetime of the alloy decreases obviously once the carbide films appears along boundaries.In the range of the applied stresses and temperatures, the creep activation energy of the alloyprepared by technique16,5and15are calculated to be Q1=597.2kJ mol-1, Q2=578.6kJ m ol-1and Q3=480.3kJ m ol-1, respectively.
The deformation features of the “oil bath” treated alloy during creep are that the single ordouble orientations slipping of dislocations activated in the J matrix. And the deformationmechanism of the “salt bath” treated alloy during creep is that dislocations slipping in the Jmatrix or shearing into the J cphase and the micro-twinning deformation, Thereinto, the <110>dislocation shearing into Jor J cphases may decompose to form the configuration of16<112>Shockley partials or13<112> super-Shockley partials plus stacking fault. The activating micro-twinning in the alloy during creep consists of16<112> Shockley partials and stacking fault, which is attributed to the decomposition of1
2<110> dislocation shearing into J or J cphases.Moreover, the dislocation networks with hexagon and quadrangle may be formed in the alloyduring creep. And the dislocation networks formed in the J/c J interface may release the latticestrain energy to decrease the stress concentration for enhancing the creep resistance of the alloy.In the later stage of creep, the slipping traces with double oriented feature appear within thegrains, and the fine carbides precipitate along the slipping traces in the “oil bath” treated alloy,and the fine grains appear in the free-J cphase zone of the boundary. However, the slippingtraces with single oriented feature appear in the “salt bath” treated alloy in the later stage ofcreep, and significant amount of dislocations are piled up in the boundaries regions to bring thestress concentration, which may results in the initiation and propagation of the cracks alongboundaries, this is thought to be the fracture mechanism of the alloy during creep.
引文
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