不同取向镍基单晶合金的蠕变行为与微观变形机制
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摘要
本文通过测定[001]、[011]以及[111]取向单晶合金在不同条件下的蠕变曲线及组织形貌观察,研究了不同取向单晶合金的蠕变行为及微观变形机制,得出以下结论:
[001]取向单晶合金的横断面枝晶呈现整齐的“+”字花样;[011]取向合金在横截面沿垂直方向的枝晶尺寸较长,沿水平方向的枝晶尺寸较短;而[111]取向的二次枝晶呈60°、或120°夹角分布。[001]、[011]和[111]取向单晶合金经完全热处理后,其组织结构均是立方γ′相沿<100>取向规则排列,其中,[011]取向合金中立方γ′相在(100)晶面呈45o角排列,经高温蠕变后,γ′相在三维空间沿<100>特定取向形成纤维状筏形组织,其γ基体相连续填充在筏状γ′相之间。[111]取向合金经高温蠕变后,γ′相在(?101)晶面沿与[111]应力轴方向近似呈约35°角形成筏状组织,形成的筏状γ′相在三维空间交错生长,γ基体相连续填充在筏形γ′相之间。
在1040℃/137MPa条件下,[001]、[111]和[011]取向单晶合金的蠕变寿命依次降低;根据不同条件蠕变性能的比较,确定出在高、中温蠕变条件下,[001]取向合金具有较好的蠕变抗力和较长的蠕变寿命,而[011]取向合金的蠕变性能较差。在实验的温度和应力范围内,测定出不同取向单晶合金的表观蠕变激活能分别为Q[001] = 469.56kJ/mol、Q[011] = 396.54kJ/mol和Q[111] = 411.12kJ/mol;应力指数分别为n[001] = 4.77、n[011] = 4.10和n[111] = 5.56。
[001]取向单晶合金在稳态蠕变期间的变形机制是位错攀移越过筏状γ′相,[001]和[011]取向单晶合金在蠕变后期的变形机制是<110>超位错剪切筏状γ′相;而[111]取向单晶合金在蠕变期间,应变量较大,有较多位错切入筏状γ′相,使其相邻γ′相之间产生取向差,并形成亚晶结构。在蠕变后期,[011]、[111]取向单晶合金首先在筏形γ′/γ两相界面之间出现裂纹,并沿垂直于应力轴方向逐渐扩展,是其合金在高温蠕变期间的主要断裂机制。
In the paper, by means of measurating of creep curves and microstructure observation, an investigation has been made into the microstructure evolution regularity of the single crystal superalloy with different orientations during creep, and the influence of the crystal orientations on the creep behaviors of the single crystal nickel base superalloy.
Results show that the regular“+”dendritic morphology is displayed in the cross-section of single crystal nickel base superalloy with [001] orientation, and the longer and shorter dendritic are displayed in the cross-section of [011]orientation alloy along the vertical and horizontal directions, and the asymmetry distribution of the dendritic directions along the 60o and 120o angles arrangment is displayed in the cross-section of [111] orientation alloy. After fully heat treated, the microstructure of the superalloys with [001], [011] and [111] orientations consists of that the cubicγ′phase is regularly arranged along the <100> directions, thereinto, the cubicγ′phase in the [011] orientation superalloy is arranged along the direction of 45o angle on (100) crystal plane, and the strip-like raftedγ′phase is formed along [001] orientation during creep. and theγmatrix phase is continuously filled between the raftedγ′phase. The cubicγ′phase in the superalloy with [111] orientation is arranged along the direction of 53o angle on (1?21) plane. During tensile creep, the meah-like raftedγ′phase is formed [111] orientation alloy along the direction of 35o angle on the (?101) plane.
The creep lifetimes of the superalloy with [001], [111] and [011] orientations reduce in turn, under the condition of 1040℃/137MPa. Compared to the [011] and [111] orientations alloys, [001] orientation alloy displays a better creep resistance and longer creep lifetime under the conditions of high and intermediate temperatures. In the range of the experimental temperatures and stresses, the strain rates of the superalloys with [001], [011] and [111] orientations during steaty state creep obey to Dorn rate equation. In the further, the activation energies of the superalloys with [001]、[011] and [111] orientations are calculated to be Q[001] = 469.56kJ/mol, Q[011] = 396.54 kJ/mol and Q[111] = 411.12kJ/mol respectively, and the stress exponents are calculated to be n[001] =4.77、n[011] =4.10 and n[111] =5.56 respectively. The dislocation climbing over theγ′raft phase is thought to be the deformed mechanism of the alloy with [001] orientation during the steady state creep. In the later stage of creep, the deformed mechanism of the alloys with [001] and [011] orientations is <110> super-dislocation shearing into the raftedγ′phase. Significant amount of dislocations appear in the [111] orientation alloy during creep due to the bigger strain, which result in the form of the subgrain as creep goes on.
In the later stage of creep, the microcracks are initiated in the interfaces of the raftedγ′/γphase in the [011] and [111] orientations supperalloy, and propagated on the interfaces of the raftedγ′/γphase along the direction vertical to the applied stress axis as creep goes on, which is thought to be the fracture mechanism of the alloys during creep.
[001]取向单晶合金的横断面枝晶呈现整齐的“+”字花样;[011]取向合金在横截面沿垂直方向的枝晶尺寸较长,沿水平方向的枝晶尺寸较短;而[111]取向的二次枝晶呈60°、或120°夹角分布。[001]、[011]和[111]取向单晶合金经完全热处理后,其组织结构均是立方γ′相沿<100>取向规则排列,其中,[011]取向合金中立方γ′相在(100)晶面呈45o角排列,经高温蠕变后,γ′相在三维空间沿<100>特定取向形成纤维状筏形组织,其γ基体相连续填充在筏状γ′相之间。[111]取向合金经高温蠕变后,γ′相在(?101)晶面沿与[111]应力轴方向近似呈约35°角形成筏状组织,形成的筏状γ′相在三维空间交错生长,γ基体相连续填充在筏形γ′相之间。
在1040℃/137MPa条件下,[001]、[111]和[011]取向单晶合金的蠕变寿命依次降低;根据不同条件蠕变性能的比较,确定出在高、中温蠕变条件下,[001]取向合金具有较好的蠕变抗力和较长的蠕变寿命,而[011]取向合金的蠕变性能较差。在实验的温度和应力范围内,测定出不同取向单晶合金的表观蠕变激活能分别为Q[001] = 469.56kJ/mol、Q[011] = 396.54kJ/mol和Q[111] = 411.12kJ/mol;应力指数分别为n[001] = 4.77、n[011] = 4.10和n[111] = 5.56。
[001]取向单晶合金在稳态蠕变期间的变形机制是位错攀移越过筏状γ′相,[001]和[011]取向单晶合金在蠕变后期的变形机制是<110>超位错剪切筏状γ′相;而[111]取向单晶合金在蠕变期间,应变量较大,有较多位错切入筏状γ′相,使其相邻γ′相之间产生取向差,并形成亚晶结构。在蠕变后期,[011]、[111]取向单晶合金首先在筏形γ′/γ两相界面之间出现裂纹,并沿垂直于应力轴方向逐渐扩展,是其合金在高温蠕变期间的主要断裂机制。
In the paper, by means of measurating of creep curves and microstructure observation, an investigation has been made into the microstructure evolution regularity of the single crystal superalloy with different orientations during creep, and the influence of the crystal orientations on the creep behaviors of the single crystal nickel base superalloy.
Results show that the regular“+”dendritic morphology is displayed in the cross-section of single crystal nickel base superalloy with [001] orientation, and the longer and shorter dendritic are displayed in the cross-section of [011]orientation alloy along the vertical and horizontal directions, and the asymmetry distribution of the dendritic directions along the 60o and 120o angles arrangment is displayed in the cross-section of [111] orientation alloy. After fully heat treated, the microstructure of the superalloys with [001], [011] and [111] orientations consists of that the cubicγ′phase is regularly arranged along the <100> directions, thereinto, the cubicγ′phase in the [011] orientation superalloy is arranged along the direction of 45o angle on (100) crystal plane, and the strip-like raftedγ′phase is formed along [001] orientation during creep. and theγmatrix phase is continuously filled between the raftedγ′phase. The cubicγ′phase in the superalloy with [111] orientation is arranged along the direction of 53o angle on (1?21) plane. During tensile creep, the meah-like raftedγ′phase is formed [111] orientation alloy along the direction of 35o angle on the (?101) plane.
The creep lifetimes of the superalloy with [001], [111] and [011] orientations reduce in turn, under the condition of 1040℃/137MPa. Compared to the [011] and [111] orientations alloys, [001] orientation alloy displays a better creep resistance and longer creep lifetime under the conditions of high and intermediate temperatures. In the range of the experimental temperatures and stresses, the strain rates of the superalloys with [001], [011] and [111] orientations during steaty state creep obey to Dorn rate equation. In the further, the activation energies of the superalloys with [001]、[011] and [111] orientations are calculated to be Q[001] = 469.56kJ/mol, Q[011] = 396.54 kJ/mol and Q[111] = 411.12kJ/mol respectively, and the stress exponents are calculated to be n[001] =4.77、n[011] =4.10 and n[111] =5.56 respectively. The dislocation climbing over theγ′raft phase is thought to be the deformed mechanism of the alloy with [001] orientation during the steady state creep. In the later stage of creep, the deformed mechanism of the alloys with [001] and [011] orientations is <110> super-dislocation shearing into the raftedγ′phase. Significant amount of dislocations appear in the [111] orientation alloy during creep due to the bigger strain, which result in the form of the subgrain as creep goes on.
In the later stage of creep, the microcracks are initiated in the interfaces of the raftedγ′/γphase in the [011] and [111] orientations supperalloy, and propagated on the interfaces of the raftedγ′/γphase along the direction vertical to the applied stress axis as creep goes on, which is thought to be the fracture mechanism of the alloys during creep.
引文
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[2] White R A, Lau Y. Method of repairing directionally solidified and single crystal alloy parts. Metals Handbook, 10th edition, Vol. 1, Properties and Selection. 1990: 981-1006.
[3]吴昌新,孙传棋,李其娟.γ′粒子尺寸对定向凝固高温合金拉伸和持久性能.航空材料学报, 2002, 22(3):1-4.
[4] Mclean M. Directionally solidified materials for high temperature service. London. Superalloys 1983, Warrendale: TMS, 1983: 9-14.
[5]张祖谦,刘志中.涡轮叶片用高温合金发展中的几个问题.国际航空, 1978 (4): 43-48.
[6]师昌绪.中国高温合金四十年.北京:中国科学技术出版社,1996: 1-8.
[7]汤鑫,刘发信,韩梅.高温合金细晶铸造技术研究.材料工程, 1997 (9): 24-25.
[8] Giamei A F, Anton D L. Effect of Rhenium additions on microstructure of a Ni-base superalloy. Metal mater trans, 1985, 16A: 1997-2005.
[9]陈荣章.单晶高温合金发展现状.材料工程, 1995 (9): 3-12.
[10]何明辉,李海燕,聂景旭. DD6单晶合金的高温蠕变损伤研究.燃气涡轮试验与研究, 2002 15(2): 24-28.
[11] Versnyder F L, Guard R W. Directional grain structures for high temperature strength. Scripta Metall Mater, 1960 (52): 485-493.
[12] Gell M, Dupta D N, Sheffler K D. High temperature super conductors with Tc over 30K. Journal of Metals, 1987 (7): 11-15.
[13]田素贵.单晶镍基合金组织演化与蠕变行为及微观特征的研究:[博士学位论文].沈阳:东北大学, 1998.
[14]陈金国.军用航空发动机的发展趋势.航空科学技术, 1994 (5): 9-12.
[15]陈荣章.北京航空材料研究院铸造高温合金发展40年.材料工程, 1998 (10): 3-7.
[16] Cetel A D, Duhl D N. Second-generation nickel-base single crystal superalloy. Symp on Superalloys 1988, Proc. 6th Int.: TMS, 1998: 2-12.
[17] Blavette D, Caron P, Khan T. An atom-probe study of some fine-scale microstructural features in Ni-based single crystal superalloys. Superalloys 1988, Reichman S, et al. eds. The Metallurgical Society, 1988: 3-12.
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