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 ]

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 pos