长寿命服役条件下DZ125合金的蠕变行为及影响因素
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摘要
通过对DZ125镍基合金进行不同工艺的热处理、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在不同条件下的相形筏化时间，通过晶格错配度的计算，研究了不同状态合金中′/两相的晶格应变程度。通过微观形貌观察及衍衬分析，研究了合金在蠕变期间的微观变形特征与断裂机制，得出如下主要结论：
铸态DZ125镍基合金的组织结构主要由基体、相、共晶组织以及块状碳化物组成，在枝晶干/间区域存在明显的成分偏析及/两相的尺寸差别。合金经完全热处理后，元素偏析程度及晶格错配度有所减小，但仍在枝晶干/间区域存在不同尺寸的′相，尺寸约为0.4m的细小立方相均匀分布在枝晶干区域，尺寸约为1—1.2m的粗大立方相存在于枝晶间区域，并有块状碳化物存在于枝晶间区域，其放射状或筛网状的共晶组织存在于枝晶间区域。
在中温/高应力蠕变期间，合金中的′相不形成筏状组织；而在高温/低应力蠕变期间，合金中的立方′相转变成与施加应力轴垂直的筏状结构。合金在1040℃/137MPa蠕变3h，相转变成与应力轴垂直的N型筏状结构。采用热力学方法计算出元素在不同条件蠕变期间的扩散迁移速率，并预测出合金在840℃和760℃蠕变期间相的筏形化时间各自需近400h和3000h，即：随蠕变温度下降，相的筏形化时间延长。
在中温/高应力蠕变期间，该合金的变形机制是位错在基体中滑移和剪切相，其中，剪切进入相的位错可以分解，形成两肖克莱不全位错加层错的位错组态，切入相的位错也可以从{111}面交滑移至{100}晶面，形成具有非平面芯的K-W锁，可以有效抑制位错在{111}面滑移，提高合金的蠕变抗力。在高温/低应力蠕变条件下，合金在稳态蠕变期间的变形机制是位错在基体中滑移和攀移越过相，其中，在位错攀移期间，位错的割阶易于形成，空位的形成和扩散是位错攀移的控制环节。蠕变后期，合金的变形机制是位错在基体中滑移和剪切进入筏状相，且在{111}面滑移。蠕变期间，分布在/两相界面的六边形或四边形位错网络，可释放晶格错配应力，减缓应力集中，提高合金的蠕变抗力。
高温蠕变的后期，合金中的裂纹首先在晶界处萌生与扩展，且不同形态晶界具有不同的损伤特征，其中，沿应力轴呈45角的晶界承受较大剪切应力，是易于使其产生蠕变损伤的主要原因；而加入的元素Hf，可促进细小粒状相沿晶界析出，可抑制晶界滑移，提高晶界强度，是使合金蠕变断裂后，断口呈现非光滑特征的主要原因。
与传统工艺热处理相比，随固溶温度提高至1260℃，合金中难熔元素的偏析程度及晶格错配度明显减小，在枝晶间区域的粗大相可完全溶解。经时效处理后，高体积分数的细小立方相均匀分布在枝晶干和枝晶间区域，可完全消除合金中的共晶组织，使合金中原大尺寸块状碳化物发生分解，并沿晶界弥散析出细小碳化物，可抑制晶界滑移。因此，与传统工艺热处理相比，高温固溶处理可改善合金的组织均匀性，提高合金蠕变抗力和蠕变寿命。
By means of heat treatments at different regimes, creep property measurement andmicrostructure observation, the influence of the heat treatment on the microstructure and creepperformance of DZ125nickel-based superalloy has been investigated. By thermodynamiccalculations, the rafting time of phase has been measured and predicted. And the strainextent of lattices for the′/double phase alloy at different states has been studied by means ofthe calculation of the lattice mismatch. The deformation and fracture mechanisms of the alloyduring creep are studied by microstyructue observation and contrast analysis od dislocationconfiguration. Some main conclusions have be obtained given as follows:
The microstructure of the as-cast DZ125nickel-base superalloy is mainly composed of matrix,′phase, eutectic and carbides. And the obvious elements segregation and nonuniformof′and phases in size exist in the dendrite arm and inter-dendrite regions. After the alloly isfull heat treated, the segregation extent of refractory elements between the dendritic arm/interdendritic regions and the lattice mismatch of/phases decrease. But the obviousdifference of′phase in size still exists in the dendritic arm and inter-dendritic regions, thefine cuboidal′precipitates is uniformly distributed in the dendritic regions, while the coarseones are distributed in the interdendritic regions. And the block-like carbides and radial ormesh-like eutectic microstructure distribute in the inter-dendritic regions.
During the creep at intermediate temperatures, the′phase in the alloy can not transforminto the rafted structure. While during creep at high temperatures, the cuboidal′phase in thealloy are transformed into the rafted structure along the direction perpendicular to the stressaxis. After crept for3h at1040°C/137MPa, the′phase in the alloy is transformed into theN-type rafted structure. The diffusion migration rate of the elements in the alloy at varioustemperatures can be calculated, by means of thermodynamic calculations, to forecast therafting time of the′phase in the alloy at different conditions. It is indicated according to thecalculation that the rafting time of′-phase prolongs as the creep temperatures decrease.Furthermore, the needed times of the′phase in the alloy during creep at840and760℃arecalculated to be400hand3000h, respectively.
The deformation mechanism of the alloy during creep at intermediate temperature is thedislocations slipping in the matrix and shearing into phase. Thereinto, the dislocationsshearing into the phase may be decomposed to form the configuration of two Shockleypartial dislocations plus stacking faults. Moreover, the super-dislocations shearing into phase may cross-slip from {111} plane to (100) plane to form the configuration of K-Wdislocation locking, which can effectively hinder dislocation slipping on {111} plane toimprove the creep resistance of alloy. Undef the conditions of high temperature and lowerstress, the dislocations slipping in the matrix and climbing over the rafted phase is thoughtto be the deformation mechanism of the alloy during the steady state creep. Thereinto, duringthe dislocations climbing, the dislocation jogs are easily formed, and the formation anddiffusion of vacancies are the control links for dislocation climbing. At the latter stage of creep,the deformation mechanism of the alloy is the dislocations slipping in matrix and shearinginto the phase. During creep, the hexagonal and quadrilateral dislocations networks locatedat the/interfaces can release mismatch stress of the lattice and delay the stressconcentration to improve the creep resistance of the alloy.
In the latter stage of creep at high termperature, the cracks in the alloy are firstly initiatedand propagated along the grain boundaries, and the various damage features display in thegrain boundary regions with different morphologies. Thereinto, the grain boundaries beingabout45angles relative to the stress axis support the bigger shear stress, which is the mainreason for promoting the occurrence of creep damage. The addition of Hf element maypromote the precipitation of the fine carbides along grain boundaries to restrain the boundariessliding, which may enhance the bonding strength of the grain boundaries. This is the mainreason for grain boundaries displaying the non-smooth surfaces after creep rupture of thealloy.
Compared to the conventional heat treatment regime, when the solution temperatureenhances to1260°C, the segregation extent of refractory elements between the dendritic/inter-dendritic regions and misfits of/phases decreases obviously, and the coarse phasein the inter-dendritic regions may be completely dissolved. After aging treatment, the fine precipitates with high volume fraction are dispersedly distributed in the dendrite andinter-dendrite regions, and the eutectic structure can be completely eliminated. Moreover, theoriginal blocky-like carbides in the alloy can be decomposed, and the fine particle-like carbides can precipitate along the boundaries to inhibit boundary slipping. Consequently,compared to the conventional heat treatment regime, the high-temperature solution treatmentcan improve the homogeneity of the microstructure in the alloy, which may enhance the creepresistance to prolong the creep life of the alloy.
铸态DZ125镍基合金的组织结构主要由基体、相、共晶组织以及块状碳化物组成，在枝晶干/间区域存在明显的成分偏析及/两相的尺寸差别。合金经完全热处理后，元素偏析程度及晶格错配度有所减小，但仍在枝晶干/间区域存在不同尺寸的′相，尺寸约为0.4m的细小立方相均匀分布在枝晶干区域，尺寸约为1—1.2m的粗大立方相存在于枝晶间区域，并有块状碳化物存在于枝晶间区域，其放射状或筛网状的共晶组织存在于枝晶间区域。
在中温/高应力蠕变期间，合金中的′相不形成筏状组织；而在高温/低应力蠕变期间，合金中的立方′相转变成与施加应力轴垂直的筏状结构。合金在1040℃/137MPa蠕变3h，相转变成与应力轴垂直的N型筏状结构。采用热力学方法计算出元素在不同条件蠕变期间的扩散迁移速率，并预测出合金在840℃和760℃蠕变期间相的筏形化时间各自需近400h和3000h，即：随蠕变温度下降，相的筏形化时间延长。
在中温/高应力蠕变期间，该合金的变形机制是位错在基体中滑移和剪切相，其中，剪切进入相的位错可以分解，形成两肖克莱不全位错加层错的位错组态，切入相的位错也可以从{111}面交滑移至{100}晶面，形成具有非平面芯的K-W锁，可以有效抑制位错在{111}面滑移，提高合金的蠕变抗力。在高温/低应力蠕变条件下，合金在稳态蠕变期间的变形机制是位错在基体中滑移和攀移越过相，其中，在位错攀移期间，位错的割阶易于形成，空位的形成和扩散是位错攀移的控制环节。蠕变后期，合金的变形机制是位错在基体中滑移和剪切进入筏状相，且在{111}面滑移。蠕变期间，分布在/两相界面的六边形或四边形位错网络，可释放晶格错配应力，减缓应力集中，提高合金的蠕变抗力。
高温蠕变的后期，合金中的裂纹首先在晶界处萌生与扩展，且不同形态晶界具有不同的损伤特征，其中，沿应力轴呈45角的晶界承受较大剪切应力，是易于使其产生蠕变损伤的主要原因；而加入的元素Hf，可促进细小粒状相沿晶界析出，可抑制晶界滑移，提高晶界强度，是使合金蠕变断裂后，断口呈现非光滑特征的主要原因。
与传统工艺热处理相比，随固溶温度提高至1260℃，合金中难熔元素的偏析程度及晶格错配度明显减小，在枝晶间区域的粗大相可完全溶解。经时效处理后，高体积分数的细小立方相均匀分布在枝晶干和枝晶间区域，可完全消除合金中的共晶组织，使合金中原大尺寸块状碳化物发生分解，并沿晶界弥散析出细小碳化物，可抑制晶界滑移。因此，与传统工艺热处理相比，高温固溶处理可改善合金的组织均匀性，提高合金蠕变抗力和蠕变寿命。
By means of heat treatments at different regimes, creep property measurement andmicrostructure observation, the influence of the heat treatment on the microstructure and creepperformance of DZ125nickel-based superalloy has been investigated. By thermodynamiccalculations, the rafting time of phase has been measured and predicted. And the strainextent of lattices for the′/double phase alloy at different states has been studied by means ofthe calculation of the lattice mismatch. The deformation and fracture mechanisms of the alloyduring creep are studied by microstyructue observation and contrast analysis od dislocationconfiguration. Some main conclusions have be obtained given as follows:
The microstructure of the as-cast DZ125nickel-base superalloy is mainly composed of matrix,′phase, eutectic and carbides. And the obvious elements segregation and nonuniformof′and phases in size exist in the dendrite arm and inter-dendrite regions. After the alloly isfull heat treated, the segregation extent of refractory elements between the dendritic arm/interdendritic regions and the lattice mismatch of/phases decrease. But the obviousdifference of′phase in size still exists in the dendritic arm and inter-dendritic regions, thefine cuboidal′precipitates is uniformly distributed in the dendritic regions, while the coarseones are distributed in the interdendritic regions. And the block-like carbides and radial ormesh-like eutectic microstructure distribute in the inter-dendritic regions.
During the creep at intermediate temperatures, the′phase in the alloy can not transforminto the rafted structure. While during creep at high temperatures, the cuboidal′phase in thealloy are transformed into the rafted structure along the direction perpendicular to the stressaxis. After crept for3h at1040°C/137MPa, the′phase in the alloy is transformed into theN-type rafted structure. The diffusion migration rate of the elements in the alloy at varioustemperatures can be calculated, by means of thermodynamic calculations, to forecast therafting time of the′phase in the alloy at different conditions. It is indicated according to thecalculation that the rafting time of′-phase prolongs as the creep temperatures decrease.Furthermore, the needed times of the′phase in the alloy during creep at840and760℃arecalculated to be400hand3000h, respectively.
The deformation mechanism of the alloy during creep at intermediate temperature is thedislocations slipping in the matrix and shearing into phase. Thereinto, the dislocationsshearing into the phase may be decomposed to form the configuration of two Shockleypartial dislocations plus stacking faults. Moreover, the super-dislocations shearing into phase may cross-slip from {111} plane to (100) plane to form the configuration of K-Wdislocation locking, which can effectively hinder dislocation slipping on {111} plane toimprove the creep resistance of alloy. Undef the conditions of high temperature and lowerstress, the dislocations slipping in the matrix and climbing over the rafted phase is thoughtto be the deformation mechanism of the alloy during the steady state creep. Thereinto, duringthe dislocations climbing, the dislocation jogs are easily formed, and the formation anddiffusion of vacancies are the control links for dislocation climbing. At the latter stage of creep,the deformation mechanism of the alloy is the dislocations slipping in matrix and shearinginto the phase. During creep, the hexagonal and quadrilateral dislocations networks locatedat the/interfaces can release mismatch stress of the lattice and delay the stressconcentration to improve the creep resistance of the alloy.
In the latter stage of creep at high termperature, the cracks in the alloy are firstly initiatedand propagated along the grain boundaries, and the various damage features display in thegrain boundary regions with different morphologies. Thereinto, the grain boundaries beingabout45angles relative to the stress axis support the bigger shear stress, which is the mainreason for promoting the occurrence of creep damage. The addition of Hf element maypromote the precipitation of the fine carbides along grain boundaries to restrain the boundariessliding, which may enhance the bonding strength of the grain boundaries. This is the mainreason for grain boundaries displaying the non-smooth surfaces after creep rupture of thealloy.
Compared to the conventional heat treatment regime, when the solution temperatureenhances to1260°C, the segregation extent of refractory elements between the dendritic/inter-dendritic regions and misfits of/phases decreases obviously, and the coarse phasein the inter-dendritic regions may be completely dissolved. After aging treatment, the fine precipitates with high volume fraction are dispersedly distributed in the dendrite andinter-dendrite regions, and the eutectic structure can be completely eliminated. Moreover, theoriginal blocky-like carbides in the alloy can be decomposed, and the fine particle-like carbides can precipitate along the boundaries to inhibit boundary slipping. Consequently,compared to the conventional heat treatment regime, the high-temperature solution treatmentcan improve the homogeneity of the microstructure in the alloy, which may enhance the creepresistance to prolong the creep life of the alloy.
引文
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