The influence of cooling rate from temperatures above the γ′ solvus on morphology, mismatch and hardness in advanced polycrystalline nickel-base superalloys

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Abstract

The interrelationships between composition, γ′ morphology, hardness and γ–γ′ mismatch were identified for a set of advanced nickel-base polycrystalline superalloys. When cooling from a temperature above the γ′ solvus, i.e. a supersaturated γ matrix, it was found that the cooling rate determines the γ′ morphology and particle distribution that develops. This in turn dictates the hardness of the alloy system and the mismatch in lattice parameters between the γ and γ′ phases. This work shows that the development of these parameters is relatively insensitive to alloy composition. Moreover, the thermal history, especially cooling rate through the γ′ solvus temperature is found to be the critical factor in determining alloy response.

Introduction

Advanced polycrystalline nickel-base superalloys are used extensively in gas turbines for high-pressure rotor discs. They rely upon a combination of solid solution strengthening and precipitation hardening to provide high-temperature mechanical properties. The precipitation strengthening in nickel-base superalloys is provided by the precipitation of large volume fractions of intermetallic phases such as γ′. In general the composition of γ′ is Ni3Al with the aluminium atoms occupying the corners and nickel the faces of the unit cell. In highly alloyed Ni-base superalloys many other elements can substitute for either nickel or aluminium or both so that the composition can be best expressed as (Ni, Co)3(Al, Ti, Ta).

The contribution to strength by the γ′ phase arises due to the precipitates impeding the motion of dislocations, through a variety of mechanisms. These are principally the formation of stress fields arising from coherency strains and anti-phase boundary energies associated with the movement of dislocations through the precipitates [1].

In terms of high temperature mechanical properties, modern nickel-base superalloys rely chiefly on the precipitation of the ordered γ′ phase to provide the main strengthening mechanism. Dislocation slip in the γ and γ′ phase occur on the {1 1 1} glide plane in the 〈1 1 0〉 direction. However, there is a tendency with increasing temperature for dislocations in the γ′ phase to cross-slip on to the {1 0 0} planes where they have a lower anti-phase domain boundary energy. Situations arise where the extended dislocation is then partly on the close-packed plane and partly on the cube plane. Such a dislocation becomes locked, leading to an increase in strength of the γ′ phase at elevated temperatures [2].

The principal mechanism of strengthening in the disordered γ phase is the introduction of solute atoms of different sizes, which create elastic strains in the crystal lattice. These strains are either tensile or compressive, depending on the relative size difference of the solute atom to the lattice. In order to achieve the maximum degree of solid solution strengthening the alloying elements are chosen to provide the largest amount of lattice strain, without precipitation of a new phase. Elements that are commonly used in nickel-base superalloys, which provide varying degrees of solid solution strengthening are tungsten, molybdenum, chromium and cobalt.

The coherent precipitation of an ordered L12 γ′ structure in a disordered fcc γ matrix gives rise to coherency strains between the two phases, due to their different lattice parameters. Both the magnitude and sign of the mismatch are important in determining the microstructure and the interactions that occur between dislocations and the precipitates. The magnitude of the strengthening effect of γ′ is influenced by the mismatch in lattice parameters between the γ′ and the γ matrix, which is given by Eq. (1).δ=2(aγaγ)(aγ+aγ),where δ is the mismatch and a is the lattice parameter of γ′ or γ′, respectively. Small differences in the lattice parameters of either phase can have a large effect on the mismatch.

The distribution of elements within the bulk alloy can have a significant effect upon the lattice parameter of each phase. Elements such as cobalt, aluminium, titanium and tantalum which are all present in γ′ may be present in varying quantities in the γ phase, due to effects of cooling rate for example. Nucleation of γ′ occurs via an ordering transformation and then is followed by diffusion controlled growth. The distribution of γ′ forming elements can have a marked effect on the lattice parameters of both phases. In addition, other elements present in the alloy composition may combine to form phases such as carbides in preference to a γ′ forming element and as such release this element into either the γ or γ′ phase.

The thermal history of the alloys is also important in determining the degree of mismatch. Increased creep lives are being sought in future alloys and this is achieved by increasing grain size. This is typically achieved by solution treatment above the γ′ solvus temperature (∼1150–1200 °C). During cooling from the supersolvus temperature γ′ precipitates out over a range of temperatures, resulting in a range of morphologies, distributions and chemistries. These are then refined further during a lower temperature age to provide optimum mechanical properties and relieve cooling residual stresses. As the chemistry, morphology, volume fraction and precipitate size of the intragranular γ′ changes with cooling rate so does the γ–γ′ mismatch. However, consideration of microstructural stability and the role of precipitate coarsening over time also need to be taken into account.

This work looks at the interrelationships between cooling rate, γ′ morphology, γ–γ′ mismatch and hardness in a range of recently developed polycrystalline nickel-base superalloys, the compositions of which are included in Table 1. Three experimental alloys (alloys UC01–UC03) have been designed to contain 50–55% volume fraction of γ′. This has been achieved by increasing the levels of γ′ forming elements, titanium and tantalum, whilst at the same time reducing solid solution strengthening elements, molybdenum and cobalt. Furthermore, these alloys have increased levels of carbon and boron which are found as carbides and borides on the γ grain boundaries. In the absence of γ′ during solution treatment it is these phases that pin the grain boundaries and thus prevent critical grain growth [3]. The changes in chemistry shown in Table 1, should affect the propensity for grain boundary pinning (carbide/boride formers) and therefore the grain size and the total volume fraction of γ′ (although not the mean size of each population).

Section snippets

Heat treatment

Prior to commencing heat treatments of the alloys, consideration was given to the theoretical phase equilibria using Thermo-Calc with the Super5 database. Thermo-Calc predicts the presence of up to nine different phases over the temperature range 500–1400 °C. Fig. 1 shows the predicted weight fractions for γ, γ′ and liquid phase for Alloy UC01 over such a temperature range. It can be seen that for UC01 the predicted γ′ solvus temperature is in the range 1150–1200 °C. This temperature was found to

Grain size analysis

Previous studies have shown that the experimental alloys under investigation in this work are capable of maintaining a uniform grain size when heat treated above the γ′ solvus temperature [14], [15]. However, there are marked differences in the grain sizes obtained which are attributable to alloy composition. Fig. 3 shows electron backscatter diffraction images of the supersolvus heat-treated microstructures of RR1000 (right) and Alloy UC01 (left). This clearly demonstrates that RR1000 exhibits

Conclusions

The microstructure of RR1000 and three experimental alloys has been studied after supersolvus heat treatments and various cooling rates. Hardness measurements were carried out to determine the influence of the cooling rate on the strength of the material before aging and to relate the observed γ′ distribution to the strength of the material. DSC measurements were undertaken to establish the degree of undercooling as a function of cooling rate of γ′ precipitation. The γ′ distribution was

Acknowledgements

Financial support for this work in the form of an Industrial CASE Award between EPSRC and Rolls-Royce plc. is gratefully acknowledged. For technical assistance at the ESRF, thanks go to the ID31 beamline scientist, Francois Fauth.

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