Al-Ni包晶合金定向凝固组织演化及小平面包晶相生长机制
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摘要
本文以Al-Ni包晶合金(L+Al_3Ni2Al_3Ni：Al_3Ni相为符合化学计量比的无任何固溶度的有序金属间化合物)为主要研究对象，通过差示扫描量热法(DSC)及Bridgman法定向凝固实验，研究了Al_3Ni相从液相直接析出时的生长特性，包晶Al_3Ni相的形核机制，定向凝固过程中Al-Ni包晶合金的组织演化规律及包晶Al_3Ni相的生长机制。
采用Al_3Ni相为初生相的Al-6at%Ni过共晶合金，通过Bridgman法定向凝固实验，研究了Al_3Ni相从液相直接析出时的凝固特性，发现在较低凝固速度(V=5m/s，G=30K/mm)下，Al_3Ni初生相具有尖锐的棱角，呈现典型的小平面相生长特征，且形貌复杂。采用Al_3Ni为包晶相的Al-Ni包晶合金，综合DSC实验及组织观察，研究了Al_3Ni相作为包晶相时的形核与生长机制。发现当外推冷却速度为零时，相比于固溶体型包晶相，无任何固溶度的包晶Al_3Ni相具有较大的形核过冷度。组织观察表明，连续凝固过程中，包晶Al_3Ni相以初生Al_3Ni2相为衬底形核并生长，Al_3Ni包晶相具有尖锐的棱角，并且包晶相层由多个晶粒组成。
就理论而言，定向凝固过程中，晶体生长起始于平面状的初始固/液界面，但是定向凝固实验过程中，初始固/液界面的准备由试样在定向凝固炉中熔化及后续的热稳定化处理(保温)组成。Al-18at%Ni(凝固区间为320℃)、Cu-13at%Sn(凝固区间为110℃)包晶合金在定向凝固初始固/液界面的准备过程中，位于完全液相区及未熔固相区之间，形成由液、固两相组成的糊状区，该糊状区与完全液相区之间的界面对应于初始固/液界面。随热稳定化处理时间的增加，糊状区内液相体积分数逐渐减少，定向凝固初始固/液界面向低温区移动。然而，即使经历了长达2h的热稳定化处理后，初始固/液界面仍不是平界面。本文建立了温度梯度作用下糊状区内的溶质扩散模型，对上述实验现象进行了合理的解释。
研究了Al-18at%Ni、Cu-13at%Sn包晶合金定向凝固初始固/液界面形貌对后续凝固组织的影响。当初始固/液界面为平界面，定向凝固工艺条件(温度梯度G、抽拉速度V)满足固相平面状生长时，后续定向凝固过程中生长界面为平界面；当初始固/液界面为非平界面时，即使定向凝固工艺条件满足固相平面状生长，后续定向凝固过程中生长界面仍为非平界面。在定向凝固工艺条件不满足固相平面状生长条件时，初始固/液界面为平界面与否对后续定向凝固过程中生长界面形貌以及定向凝固组织演化的影响将不明显，不会改变定向凝固包晶合金中各相的凝固顺序及生长机制。
依据相图推测，近平衡凝固时，Al-Ni包晶合金的凝固顺序应为初生Al_3Ni2相领先于包晶Al_3Ni相自液相析出。然而，Al-18at%Ni合金定向凝固过程中，发现凝固速度为1m/s、8m/s时，初生Al_3Ni2相分别呈胞状和枝晶生长，随凝固距离的增加，固/液界面处的领先相由初生Al_3Ni2相转变为包晶Al_3Ni相。根据溶质守恒和菲克扩散定律，建立了无任何固溶度的金属间化合物以平界面定向生长时的溶质再分配模型，发现当合金初始成分小于金属间化合物成分时，随凝固距离的增加，固/液界面处液相中溶质浓度线性递减，并利用这一模型揭示了上述现象发生的机理。
在Al-25at%Ni包晶合金定向凝固过程中，当凝固速度处于5m/s~500m/s范围内，初生Al_3Ni2相将以枝晶形貌领先于包晶Al_3Ni相从液相析出，温度降至包晶反应温度以下时，Al_3Ni相以Al_3Ni2相为基底形核并长大。当凝固速度为5、10、20m/s时，包晶反应界面前后初生Al_3Ni2相体积分数明显减少。当凝固速度高于20m/s时，包晶反应界面前后初生相Al_3Ni2体积分数不大，包晶Al_3Ni相的生长机制为自液相直接凝固。定量测量了Al-25at%Ni合金在凝固速度为20m/s条件下定向凝固时试样对应不同温度的横截面上各相体积分数。上述实验现象与利用经典包晶凝固理论模型计算的Al-Ni包晶合金中包晶反应以及包晶转变对初生Al_3Ni2相的消耗量有很大差异。
对定向凝固Al-Ni包晶合金组织演化的研究发现，在定向凝固过程中，初生Al_3Ni2相二次枝晶臂上端形成大量的包晶Al_3Ni相，而初生Al_3Ni2相二次枝晶臂下端处并未发现包晶Al_3Ni相。分析表明，在定向凝固包晶合金枝晶生长过程中，温度梯度的存在会导致温度梯度区域熔化效应(TGZM：Temperature Gradient Zone Melting)，并且当外界强制冷却不足以抑制TGZM效应引起的重熔现象时，TGZM效应将导致初生Al_3Ni2相与包晶Al_3Ni相的重熔和凝固，这促使一种特殊的分离式包晶反应机制的发生，进而导致定向凝固过程中大量初生Al_3Ni2相发生溶解。
In this study, Al-Ni peritectic alloys exhibiting peritectic reactionL+Al_3Ni2Al_3Ni in which the peritectic Al_3Ni phase is an intermetallic compoundwith nil solubility, have been chosen for investigation. Growth of Al_3Ni phasedirectly from the melt, nucleation of Al_3Ni phase as a peritectic phase,microstructure evolution and phase selection of Al-Ni peritectic alloys duringdirectional solidification, and growth mechanism of Al_3Ni phase as a peritecticphase have been investigated using differential scanning calorimetry (DSC)analysis and a Bridgman-type directional solidification apparatus.
Al-6at%Ni alloy has been selected to investigate the growth mechanism ofAl_3Ni phase as the primary phase precipitating from the melt. During directionalsolidification, at a growth rate of5m/s and a temperature gradient of30K/mm,Al_3Ni phase precipitates from the melt directly, exhibiting faceted growth interfaceand complex3-D morphology. Nucleation and growth of Al_3Ni as a peritectic phaseduring continuous solidification process of Al-Ni peritectic alloys have beenanalyzed. The DSC analysis results show that when extrapolating the cooling ratesto0℃/min, the degenerated nucleation temperature of solid solution type peritecticphase is equal to the equilibrium peritectic temperature, while a great amount ofundercooling has been observed for Al_3Ni phase. The microstructure analysis showsthat the primary Al_3Ni2phase can act as a sound heterogeneous nucleation substrate.The Al_3Ni phase layer envelopping the primary Al_3Ni2phase exhibits facetedgrowth interface and is composed of several grains.
The preparation of the initial solid/liquid interface during directionalsolidification on which the growth starts is a critical step, which consists of meltingand thermal stabilization under a temperature gradient. Experiments on Al-Ni andCu-Sn peritectic alloys consisting of melting followed by thermal stabilizationranging from0to2h have been carried out in a Bridgman-type furnace. In thedirectional melting process, due to the temperature gradient imposed on the rod, amushy zone is created between the complete liquid zone and non-molten solid zoneparallel to the temperature gradient. With the thermal stabilization time increase,the volume fraction of the liquid phase in the mushy zone decreases, and the initialsolid/liquid interface moves downward parallel to the temperature gradient. After2h thermal stabilization, the initial solid-liquid interface is corrugated, whichdeparts from the theorectical prediction. A solute diffusion model under atemperature gradient has been built up to explain the above experimental observations.
The consequence of the initial solid/liquid interface morphology on themicrostructure evolution during subsequent directional solidification process of Al-Ni and Cu-Sn alloys has been investigated by changing the thermal stabilizationduration. It is found that if the initial solid/liquid interface is non-planar, even whenthe experimental parameters such as temperature gradient and growth velocitymaintaining planar growth interface are met, the growth interface is non-planar.However, if the initial solid/liquid interface is planar, a planar growth interface canbe obtained. When the experimental conditions maitaining planar growth interfaceare not met, whether the initial solid/liquid is planar or not doesn’t have obviousinfluence on the subsequent microstructure evolution during directionalsolidification.
Systematic directional solidification experiments have been performed on Al-Ni peritectic alloys. During directional solidification of Al-18at%Ni alloy, withpulling rates of1m/s and8m/s, with freezing distance increase, the precipitatingsolid phase from the liquid at the quenching solid/liquid interface transforms fromprimary Al_3Ni2phase to peritectic Al_3Ni phase. Based on the principle of soluteconservation and Fick’s diffusion law, solute redistribution behavior ofintermetallic compound with nil solubility range during planar directional growthhas been analyzed, which has been used to explain the experimental observations indirectionally solidified Al-18at%Ni alloy.
During directional solidification of Al-25at%Ni alloy, with pulling rate rangingfrom5to500m/s, primary Al_3Ni2phase precipitates from the liquid firstexhibiting dendrite morphology, and peritectic Al_3Ni phase forms at a locationbelow the peritectic interface. At growth rate of5,10, and20m/s, a great volumefraction of primary Al_3Ni2phase is consumed below the peritectic reactiontemperature. With growth rate beyond20m/s, the change of Al_3Ni2volumefraction changes a little below the peritectic reaction temperature, which indicatesthe growth mechanism of the peritectic Al_3Ni phase is direct solidification from theliquid. Volume fractions of each phase in the transverse sections with differenttemperatures have been evaluated for the sample at pulling rate of20m/s. PrimaryAl_3Ni2phase dissolution is much greater than that should be consumed by peritecticreaction and transformation which is calculated based on traditional peritecticsolidification theory.
During dendritic solidification of Al-Ni peritectic alloys under a temperaturegradient, it is observed that a thick peritectic layer forms on the front edge of thesecondary dendrite arm of the primary Al_3Ni2phase, while there is almost no peritectic Al_3Ni phase on the back edge of the secondary dendrite arm. Thisobservation is explained satisfactorily by a new version of secondary dendrite armmigration caused by temperature gradient zone melting (TGZM) during peritecticsolidification, which involves both primary and peritectic phases. A divorcedperitectic reaction model caused by TGZM in directional solidification process hasbeen proposed, which can satisfactorily explain the fast primary Al_3Ni2dissolutionduring directional solidification.
采用Al_3Ni相为初生相的Al-6at%Ni过共晶合金，通过Bridgman法定向凝固实验，研究了Al_3Ni相从液相直接析出时的凝固特性，发现在较低凝固速度(V=5m/s，G=30K/mm)下，Al_3Ni初生相具有尖锐的棱角，呈现典型的小平面相生长特征，且形貌复杂。采用Al_3Ni为包晶相的Al-Ni包晶合金，综合DSC实验及组织观察，研究了Al_3Ni相作为包晶相时的形核与生长机制。发现当外推冷却速度为零时，相比于固溶体型包晶相，无任何固溶度的包晶Al_3Ni相具有较大的形核过冷度。组织观察表明，连续凝固过程中，包晶Al_3Ni相以初生Al_3Ni2相为衬底形核并生长，Al_3Ni包晶相具有尖锐的棱角，并且包晶相层由多个晶粒组成。
就理论而言，定向凝固过程中，晶体生长起始于平面状的初始固/液界面，但是定向凝固实验过程中，初始固/液界面的准备由试样在定向凝固炉中熔化及后续的热稳定化处理(保温)组成。Al-18at%Ni(凝固区间为320℃)、Cu-13at%Sn(凝固区间为110℃)包晶合金在定向凝固初始固/液界面的准备过程中，位于完全液相区及未熔固相区之间，形成由液、固两相组成的糊状区，该糊状区与完全液相区之间的界面对应于初始固/液界面。随热稳定化处理时间的增加，糊状区内液相体积分数逐渐减少，定向凝固初始固/液界面向低温区移动。然而，即使经历了长达2h的热稳定化处理后，初始固/液界面仍不是平界面。本文建立了温度梯度作用下糊状区内的溶质扩散模型，对上述实验现象进行了合理的解释。
研究了Al-18at%Ni、Cu-13at%Sn包晶合金定向凝固初始固/液界面形貌对后续凝固组织的影响。当初始固/液界面为平界面，定向凝固工艺条件(温度梯度G、抽拉速度V)满足固相平面状生长时，后续定向凝固过程中生长界面为平界面；当初始固/液界面为非平界面时，即使定向凝固工艺条件满足固相平面状生长，后续定向凝固过程中生长界面仍为非平界面。在定向凝固工艺条件不满足固相平面状生长条件时，初始固/液界面为平界面与否对后续定向凝固过程中生长界面形貌以及定向凝固组织演化的影响将不明显，不会改变定向凝固包晶合金中各相的凝固顺序及生长机制。
依据相图推测，近平衡凝固时，Al-Ni包晶合金的凝固顺序应为初生Al_3Ni2相领先于包晶Al_3Ni相自液相析出。然而，Al-18at%Ni合金定向凝固过程中，发现凝固速度为1m/s、8m/s时，初生Al_3Ni2相分别呈胞状和枝晶生长，随凝固距离的增加，固/液界面处的领先相由初生Al_3Ni2相转变为包晶Al_3Ni相。根据溶质守恒和菲克扩散定律，建立了无任何固溶度的金属间化合物以平界面定向生长时的溶质再分配模型，发现当合金初始成分小于金属间化合物成分时，随凝固距离的增加，固/液界面处液相中溶质浓度线性递减，并利用这一模型揭示了上述现象发生的机理。
在Al-25at%Ni包晶合金定向凝固过程中，当凝固速度处于5m/s~500m/s范围内，初生Al_3Ni2相将以枝晶形貌领先于包晶Al_3Ni相从液相析出，温度降至包晶反应温度以下时，Al_3Ni相以Al_3Ni2相为基底形核并长大。当凝固速度为5、10、20m/s时，包晶反应界面前后初生Al_3Ni2相体积分数明显减少。当凝固速度高于20m/s时，包晶反应界面前后初生相Al_3Ni2体积分数不大，包晶Al_3Ni相的生长机制为自液相直接凝固。定量测量了Al-25at%Ni合金在凝固速度为20m/s条件下定向凝固时试样对应不同温度的横截面上各相体积分数。上述实验现象与利用经典包晶凝固理论模型计算的Al-Ni包晶合金中包晶反应以及包晶转变对初生Al_3Ni2相的消耗量有很大差异。
对定向凝固Al-Ni包晶合金组织演化的研究发现，在定向凝固过程中，初生Al_3Ni2相二次枝晶臂上端形成大量的包晶Al_3Ni相，而初生Al_3Ni2相二次枝晶臂下端处并未发现包晶Al_3Ni相。分析表明，在定向凝固包晶合金枝晶生长过程中，温度梯度的存在会导致温度梯度区域熔化效应(TGZM：Temperature Gradient Zone Melting)，并且当外界强制冷却不足以抑制TGZM效应引起的重熔现象时，TGZM效应将导致初生Al_3Ni2相与包晶Al_3Ni相的重熔和凝固，这促使一种特殊的分离式包晶反应机制的发生，进而导致定向凝固过程中大量初生Al_3Ni2相发生溶解。
In this study, Al-Ni peritectic alloys exhibiting peritectic reactionL+Al_3Ni2Al_3Ni in which the peritectic Al_3Ni phase is an intermetallic compoundwith nil solubility, have been chosen for investigation. Growth of Al_3Ni phasedirectly from the melt, nucleation of Al_3Ni phase as a peritectic phase,microstructure evolution and phase selection of Al-Ni peritectic alloys duringdirectional solidification, and growth mechanism of Al_3Ni phase as a peritecticphase have been investigated using differential scanning calorimetry (DSC)analysis and a Bridgman-type directional solidification apparatus.
Al-6at%Ni alloy has been selected to investigate the growth mechanism ofAl_3Ni phase as the primary phase precipitating from the melt. During directionalsolidification, at a growth rate of5m/s and a temperature gradient of30K/mm,Al_3Ni phase precipitates from the melt directly, exhibiting faceted growth interfaceand complex3-D morphology. Nucleation and growth of Al_3Ni as a peritectic phaseduring continuous solidification process of Al-Ni peritectic alloys have beenanalyzed. The DSC analysis results show that when extrapolating the cooling ratesto0℃/min, the degenerated nucleation temperature of solid solution type peritecticphase is equal to the equilibrium peritectic temperature, while a great amount ofundercooling has been observed for Al_3Ni phase. The microstructure analysis showsthat the primary Al_3Ni2phase can act as a sound heterogeneous nucleation substrate.The Al_3Ni phase layer envelopping the primary Al_3Ni2phase exhibits facetedgrowth interface and is composed of several grains.
The preparation of the initial solid/liquid interface during directionalsolidification on which the growth starts is a critical step, which consists of meltingand thermal stabilization under a temperature gradient. Experiments on Al-Ni andCu-Sn peritectic alloys consisting of melting followed by thermal stabilizationranging from0to2h have been carried out in a Bridgman-type furnace. In thedirectional melting process, due to the temperature gradient imposed on the rod, amushy zone is created between the complete liquid zone and non-molten solid zoneparallel to the temperature gradient. With the thermal stabilization time increase,the volume fraction of the liquid phase in the mushy zone decreases, and the initialsolid/liquid interface moves downward parallel to the temperature gradient. After2h thermal stabilization, the initial solid-liquid interface is corrugated, whichdeparts from the theorectical prediction. A solute diffusion model under atemperature gradient has been built up to explain the above experimental observations.
The consequence of the initial solid/liquid interface morphology on themicrostructure evolution during subsequent directional solidification process of Al-Ni and Cu-Sn alloys has been investigated by changing the thermal stabilizationduration. It is found that if the initial solid/liquid interface is non-planar, even whenthe experimental parameters such as temperature gradient and growth velocitymaintaining planar growth interface are met, the growth interface is non-planar.However, if the initial solid/liquid interface is planar, a planar growth interface canbe obtained. When the experimental conditions maitaining planar growth interfaceare not met, whether the initial solid/liquid is planar or not doesn’t have obviousinfluence on the subsequent microstructure evolution during directionalsolidification.
Systematic directional solidification experiments have been performed on Al-Ni peritectic alloys. During directional solidification of Al-18at%Ni alloy, withpulling rates of1m/s and8m/s, with freezing distance increase, the precipitatingsolid phase from the liquid at the quenching solid/liquid interface transforms fromprimary Al_3Ni2phase to peritectic Al_3Ni phase. Based on the principle of soluteconservation and Fick’s diffusion law, solute redistribution behavior ofintermetallic compound with nil solubility range during planar directional growthhas been analyzed, which has been used to explain the experimental observations indirectionally solidified Al-18at%Ni alloy.
During directional solidification of Al-25at%Ni alloy, with pulling rate rangingfrom5to500m/s, primary Al_3Ni2phase precipitates from the liquid firstexhibiting dendrite morphology, and peritectic Al_3Ni phase forms at a locationbelow the peritectic interface. At growth rate of5,10, and20m/s, a great volumefraction of primary Al_3Ni2phase is consumed below the peritectic reactiontemperature. With growth rate beyond20m/s, the change of Al_3Ni2volumefraction changes a little below the peritectic reaction temperature, which indicatesthe growth mechanism of the peritectic Al_3Ni phase is direct solidification from theliquid. Volume fractions of each phase in the transverse sections with differenttemperatures have been evaluated for the sample at pulling rate of20m/s. PrimaryAl_3Ni2phase dissolution is much greater than that should be consumed by peritecticreaction and transformation which is calculated based on traditional peritecticsolidification theory.
During dendritic solidification of Al-Ni peritectic alloys under a temperaturegradient, it is observed that a thick peritectic layer forms on the front edge of thesecondary dendrite arm of the primary Al_3Ni2phase, while there is almost no peritectic Al_3Ni phase on the back edge of the secondary dendrite arm. Thisobservation is explained satisfactorily by a new version of secondary dendrite armmigration caused by temperature gradient zone melting (TGZM) during peritecticsolidification, which involves both primary and peritectic phases. A divorcedperitectic reaction model caused by TGZM in directional solidification process hasbeen proposed, which can satisfactorily explain the fast primary Al_3Ni2dissolutionduring directional solidification.
引文
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