| 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 at1150and 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 of12<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. |
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