The long and large tertiary stage, that dominates the creep curve shape of ?' reinforced nickel base superalloys for temperatures/stresses relevant to high temperature components, has been often modeled supposing a single strain softening mechanism is operative. For example the accelerating tertiary creep has been described by a linear dependence of strain rate on strain[1-2]: ? . = ? 0(1+C?) (1) where ? . and ? correspond, respectively, to the instantaneous strain rate and the accumulated creep strain, ? 0 represents the creep strain rate extrapolated to ? = 0, and C is a parameter of proportionality between creep damage and the strain, W = C?. The relationship in Eq. 1 has been physically justified in [3] supposing the softening in nickel base superalloys is due to the accumulation of mobile dislocations that are proportional to the creep strain. Time softening, due to the ripening of the particles, is not considered in Eq. 1. Single crystal nickel base superalloys, with a large fraction of hardening cuboidal ?', if creep tested under tensile load along <001> crystalline direction at high temperature, can produce a lamellar or rafted ?/?' pattern perpendicular to the loading axis. This ?' coalescence process can happen only at high temperatures, typically at T=900°C depending on the alloy, and very early, i.e. within the first 1-3% of creep life, for T>1000°C. In fact, for tests performed around 900°C, the cuboidal microstructure is generally present for a considerable part of the creep test: the raft development can start in correspondence of the minimum creep rate and ends well inside the tertiary creep and consistently only a slight influence of the microstructure instability can appear on the creep behaviour at this temperature. At higher temperatures, instead, (1050-1100°C) the cuboid microstructure disappears soon, the cuboid ?' develops into lamellae in the early primary creep, and the raft structure is present during almost the whole creep test. To extrapolate the creep behaviour at such experimental conditions from data obtained at lower temperatures, the microstructure instability must be taken into account, since the raft development can strongly influence the dislocation mobility when the dislocations cannot easily cut the long rafted ?', particularly at low stress values typical of the creep tests performed at such high temperatures. The purpose of the present paper is to study the influence of the ?' morphology evolution on the creep strain rate in the 900-1050°C temperature range mostly important for single crystal components in high performance gas turbines. The rafts formation and their evolution, and their effect on the creep behaviour, have been studied in superalloy SMP 14 - developed by CSIR, Pretoria, RSA and supplied by Ross &Catherall Ltd Shefield UK - at 900, 950, 1000 and 1050°C and with applied stresses in the range 135 - 425 MPa to produce times to rupture between 300 and 3000 h. The nominal composition of the SMP14 is compared in Table 1 with the well-established CMSX-4 and the third generation alloy TMS75, designed for gas turbine single crystal blades/vanes. The heat treatment, a wide microstructure characterization and a mechanical comparison of SMP14 with CMSX-4 can be found in [6]. Figs. 2 and 3 show the experimental creep curves as ? . vs. ? and log? . vs. time. In particular the Figs. 2 show that Eq. 1 can well describe the rate of strain accumulation for almost the whole experimental tertiary creep. The plots of Figs 2 and 3 are equivalent, i.e. experimental points that show a linear relationship in a plot vs , must also display a linear relationship in a plot log vs time. The latter plot expands the initial portion of the creep curve, and it appears to be more convenient to observe possible time dependent effects occurring during the early creep strain. In fact, the juxtaposition of the same creep curves in Figs. 2 and in Figs. 3 reveals new unusual features hidden in the ? . vs. ? plots. The experimental results show that, in addition to the strain softening damage described by Eq. 1, other microstructure dependent mechanisms influence the creep strain accumulation. In particular the analysis of Figs 3, combined with the microstructure observations of creep interrupted tests showed that: - stage A: In correspondence of the building up of the raft microstructure, the thickness of the horizontal ? channels significantly increases producing a reduction of the Orowan resistance of the dislocation mobility. Such microstructure evolution damage, added to the damage described by Eq.1, causes the fastest acceleration stage; - stage B: after raft formation at high stress, the ? channels continue to grow moderately. During the stage B, where most of the creep strain is accumulated in a relatively short time, the strain acceleration can be described by Eq. 1; - stage C: after raft formation at low stress, a further regime of strain accumulation, stage C, appears between stages A and B; it can be due to two concurring mechanisms, i.e.: i) The increment of dislocation density in the ?/?' interface relaxes the misfit stresses, reducing the local stress in the ? channels and making more difficult the movement of dislocations. This process can partially explain the rapid decrement of the strain rate after the raft development in the tests at the highest temperatures. ii) For material with completely developed rafted structure, the ? phase cannot easily flow around the ?' phase because the ? lamellae are essentially unconnected. Supposing the dislocation activity is mainly restricted into the ? phase, during creep the phase can plastically deform, while the ?' phase basically can deform only elastically. Essentially the material can be considered as a metal matrix composite with the two phases deforming simultaneously. During creep the stress is off-loaded from the "soft" matrix regions (?), while the stress in the "hard" reinforcing zones (?') is amplified. In the rafted material, the decreasing of the stage C slope, becoming negative at applied stress values below 200 MPa, reflects the reducing of the local stress in the ? channel. In principle the creep rate should progressively decrease, with increasing creep strain, as the local stress in the ? phase approaches the Orowan and solid solution resistance. Actually the reduction of the creep strain rate with the strain ends when, in absence of some other damage mechanisms, a recovery process prevents a further stress redistribution between the two phases. The simplest recovery process consists in the cutting of the ?' phase by the interface dislocations; this recovery process could be operative after the accumulation of plastic strain in the phase has built up a sufficient internal stress, that in conjunction with the external applied stress, allows the cutting of the ?' phase. As a conclusion, the analysis of the experimental creep curves of the SMP 14 has shown: oThe tertiary creep dominates the creep curves. oIn the explored stress/temperature field, a rafted microstructure is developed during creep. oAt the highest explored temperatures the raft development is very rapid and takes only a few percent of the creep life, while at 900°C rafts are developed only in late tertiary creep. During tertiary creep, different stages can be distinguished: oa strongly creep accelerating stage, that occurs during the raft development; oa stage that produces a reduction of the strain rate for the highest temperatures/lowest stresses explored: this stage has been attributed to a redistribution of the local stresses between the ? and ?' phases and it disappears in the highest stresses tests; oa further stage where the strain accelerates again, with the strain rate following a linear strain softening relationship in the 900 - 1000°C interval. oThe fastest and shortest final creep accelerating stage, leading to fracture.
Le curve di creep di superleghe a base nichel monocristalline rinforzate dalla precipitazione della fase ?' sono spesso dominate dallo stadio accelerante/terziario dovuto all'accumulazione di un danno interno non direttamente relazionato a meccanismi di frattura, ma piuttosto a una variazione della densità e/o mobilità delle dislocazioni mobili che può essere correlata con la deformazione da creep accumulata. Un'attenta esamina dello stadio accelerante ottenuto sulla superlega SMP14 mostra differenti regimi di accumulazione della deformazione in funzione dalla sollecitazione e temperatura applicate. Tale comportamento sperimentale può essere razionalizzato dall'evoluzione, durante il creep, della morfologia della fase rinforzante ?'.
Softening and hardening mechanisms due to rafting development in a SX nickel base superalloy : [ Meccanismi di addolcimento ed incrudimento dovuti allo sviluppo di raft in una superlega monocristallina a base di nichel ]
Maldini M;Angella G;Lupinc V
2011
Abstract
The long and large tertiary stage, that dominates the creep curve shape of ?' reinforced nickel base superalloys for temperatures/stresses relevant to high temperature components, has been often modeled supposing a single strain softening mechanism is operative. For example the accelerating tertiary creep has been described by a linear dependence of strain rate on strain[1-2]: ? . = ? 0(1+C?) (1) where ? . and ? correspond, respectively, to the instantaneous strain rate and the accumulated creep strain, ? 0 represents the creep strain rate extrapolated to ? = 0, and C is a parameter of proportionality between creep damage and the strain, W = C?. The relationship in Eq. 1 has been physically justified in [3] supposing the softening in nickel base superalloys is due to the accumulation of mobile dislocations that are proportional to the creep strain. Time softening, due to the ripening of the particles, is not considered in Eq. 1. Single crystal nickel base superalloys, with a large fraction of hardening cuboidal ?', if creep tested under tensile load along <001> crystalline direction at high temperature, can produce a lamellar or rafted ?/?' pattern perpendicular to the loading axis. This ?' coalescence process can happen only at high temperatures, typically at T=900°C depending on the alloy, and very early, i.e. within the first 1-3% of creep life, for T>1000°C. In fact, for tests performed around 900°C, the cuboidal microstructure is generally present for a considerable part of the creep test: the raft development can start in correspondence of the minimum creep rate and ends well inside the tertiary creep and consistently only a slight influence of the microstructure instability can appear on the creep behaviour at this temperature. At higher temperatures, instead, (1050-1100°C) the cuboid microstructure disappears soon, the cuboid ?' develops into lamellae in the early primary creep, and the raft structure is present during almost the whole creep test. To extrapolate the creep behaviour at such experimental conditions from data obtained at lower temperatures, the microstructure instability must be taken into account, since the raft development can strongly influence the dislocation mobility when the dislocations cannot easily cut the long rafted ?', particularly at low stress values typical of the creep tests performed at such high temperatures. The purpose of the present paper is to study the influence of the ?' morphology evolution on the creep strain rate in the 900-1050°C temperature range mostly important for single crystal components in high performance gas turbines. The rafts formation and their evolution, and their effect on the creep behaviour, have been studied in superalloy SMP 14 - developed by CSIR, Pretoria, RSA and supplied by Ross &Catherall Ltd Shefield UK - at 900, 950, 1000 and 1050°C and with applied stresses in the range 135 - 425 MPa to produce times to rupture between 300 and 3000 h. The nominal composition of the SMP14 is compared in Table 1 with the well-established CMSX-4 and the third generation alloy TMS75, designed for gas turbine single crystal blades/vanes. The heat treatment, a wide microstructure characterization and a mechanical comparison of SMP14 with CMSX-4 can be found in [6]. Figs. 2 and 3 show the experimental creep curves as ? . vs. ? and log? . vs. time. In particular the Figs. 2 show that Eq. 1 can well describe the rate of strain accumulation for almost the whole experimental tertiary creep. The plots of Figs 2 and 3 are equivalent, i.e. experimental points that show a linear relationship in a plot vs , must also display a linear relationship in a plot log vs time. The latter plot expands the initial portion of the creep curve, and it appears to be more convenient to observe possible time dependent effects occurring during the early creep strain. In fact, the juxtaposition of the same creep curves in Figs. 2 and in Figs. 3 reveals new unusual features hidden in the ? . vs. ? plots. The experimental results show that, in addition to the strain softening damage described by Eq. 1, other microstructure dependent mechanisms influence the creep strain accumulation. In particular the analysis of Figs 3, combined with the microstructure observations of creep interrupted tests showed that: - stage A: In correspondence of the building up of the raft microstructure, the thickness of the horizontal ? channels significantly increases producing a reduction of the Orowan resistance of the dislocation mobility. Such microstructure evolution damage, added to the damage described by Eq.1, causes the fastest acceleration stage; - stage B: after raft formation at high stress, the ? channels continue to grow moderately. During the stage B, where most of the creep strain is accumulated in a relatively short time, the strain acceleration can be described by Eq. 1; - stage C: after raft formation at low stress, a further regime of strain accumulation, stage C, appears between stages A and B; it can be due to two concurring mechanisms, i.e.: i) The increment of dislocation density in the ?/?' interface relaxes the misfit stresses, reducing the local stress in the ? channels and making more difficult the movement of dislocations. This process can partially explain the rapid decrement of the strain rate after the raft development in the tests at the highest temperatures. ii) For material with completely developed rafted structure, the ? phase cannot easily flow around the ?' phase because the ? lamellae are essentially unconnected. Supposing the dislocation activity is mainly restricted into the ? phase, during creep the phase can plastically deform, while the ?' phase basically can deform only elastically. Essentially the material can be considered as a metal matrix composite with the two phases deforming simultaneously. During creep the stress is off-loaded from the "soft" matrix regions (?), while the stress in the "hard" reinforcing zones (?') is amplified. In the rafted material, the decreasing of the stage C slope, becoming negative at applied stress values below 200 MPa, reflects the reducing of the local stress in the ? channel. In principle the creep rate should progressively decrease, with increasing creep strain, as the local stress in the ? phase approaches the Orowan and solid solution resistance. Actually the reduction of the creep strain rate with the strain ends when, in absence of some other damage mechanisms, a recovery process prevents a further stress redistribution between the two phases. The simplest recovery process consists in the cutting of the ?' phase by the interface dislocations; this recovery process could be operative after the accumulation of plastic strain in the phase has built up a sufficient internal stress, that in conjunction with the external applied stress, allows the cutting of the ?' phase. As a conclusion, the analysis of the experimental creep curves of the SMP 14 has shown: oThe tertiary creep dominates the creep curves. oIn the explored stress/temperature field, a rafted microstructure is developed during creep. oAt the highest explored temperatures the raft development is very rapid and takes only a few percent of the creep life, while at 900°C rafts are developed only in late tertiary creep. During tertiary creep, different stages can be distinguished: oa strongly creep accelerating stage, that occurs during the raft development; oa stage that produces a reduction of the strain rate for the highest temperatures/lowest stresses explored: this stage has been attributed to a redistribution of the local stresses between the ? and ?' phases and it disappears in the highest stresses tests; oa further stage where the strain accelerates again, with the strain rate following a linear strain softening relationship in the 900 - 1000°C interval. oThe fastest and shortest final creep accelerating stage, leading to fracture.I documenti in IRIS sono protetti da copyright e tutti i diritti sono riservati, salvo diversa indicazione.
