Research Article
Microstructure and dislocation structure evolution during creep life of Ni-based single crystal superalloys

https://doi.org/10.1016/j.jmst.2019.11.028Get rights and content

Abstract

The high performance of Ni single crystal superalloys during high temperature low stress creep service, is intrinsically determined by the combined effects of microstructural evolution and the dislocation behaviour. In the field of the evolution of dislocation network, two main recovery mechanism based on dislocation migration dominate the process. One is superdislocations shearing into γ’ rafts through a two-superpartials-assisted approach. Another is the compact dislocations migrating along γ/γ′ interface. These two mechanisms are similarly climb-rate-controlled process. In this work, a model for the minimum creep rate based on thermodynamic and kinetic calculations and using an existing detailed dislocation dynamics model has been built by taking the dislocation migration behaviours as well as the rafted microstructure into consideration, which can well reproduce the ([100] tensile) creep properties of existing Ni superalloy grades, without the need to make the dislocation parameter values composition dependent.

Introduction

Ni-based single crystal superalloys have been widely used for the blades and other loaded structures of aero-engines and gas turbines due to their superior mechanical properties, in particular their excellent creep resistance at high temperatures [1]. Their outstanding creep resistance, not only origins from the absence of grain boundaries but is largely determined by the unique microstructure characterised by the presence of a high-volume fraction of the long-range ordered L12 γ′ phase, which appears as cubes coherently embedded in a face-centered cubic solid solution γ matrix. In general, the size, volume fraction and morphology of γ′ precipitates mainly determine the mechanical properties of Ni-superalloys [2,3]. In the as-produced condition the cuboidal γ′-precipitates have a size of around 0.4um size and they are separated by γ-channels with around 0.1um size. The typical precipitate volume fraction at room temperature is 50% or higher [1].

When exposed to their typical use conditions, a relatively high temperature (>950 °C) and a modest stress (<250 MPa), there is directional coarsening of γ′ precipitates during the early creep stage, which is so-called “rafting” stage [[4], [5], [6], [7]]. During this time, the initially adjacent cuboidal γ′ particles coalesce and form platelets that turn into plate-like or rod-like structures. This lamellar γ/γ’ rafted microstructure will remain more or less unchanged during the long stable creep stage, until the γ’ gradually interconnects and becomes the matrix phase surrounding isolated γ phase islands [8,9]. This process is known as the ‘topological phase inversion’, which has been considered as the microstructural indicator marking the transition from quasi-stationary creep to accelerated creep. This inverted microstructure is maintained during the accelerated creep stage but rapidly loses it regularity, the morphology evolution of the phases during the entire creep life is presented in Fig. 1 [10].

On the other hand, the creep response not only depends on the microstructure evolution but also on the changes in the dynamics and topology of the dislocations and dislocation networks [11]. At the beginning of creep loading, the deformation is governed by the dislocation glide and dislocation multiplication in the γ channels. Soon thereafter the mobile dislocations start to accumulate and become rearranged at the γ/γ’ interface, while the formation of lamellar rafts takes place, leading to the widely observed formation of dislocation networks on the  γ/γ’ interface [[11], [12], [13], [14], [15]]. Analogous to the lamellar microstructure, the dislocation network will remain stable until the end of the stable creep stage (stage II creep) when the network begins to degrade by huge amounts of dislocations cutting into the rafted γ′ through the interface. Ultimately this chaotic dislocation multiplication process leads to rupture.

Summing up, the creep properties of Ni single crystal superalloys are strongly dependent on the combined effect of the microstructure evolution and the dislocation behaviour. Therefore, a good understanding and some quantification of microstructure and dislocation evolution during creep is crucial to better comprehend the mechanical properties of Ni superalloys. In this work, the current work on the dependence of mechanical properties in Ni single crystals on the microstructural and dislocation behaviour during isothermal creep tests is reviewed. The experimental observations and the corresponding models and simulations are combined to show their mutual interaction. A simple model has been built in which the governing dislocation dynamics equation is thermodynamically coupled to the chemical composition of some commercial superalloys to predict their minimum creep rate at the loading conditions.

Section snippets

Initial microstructure

The γ′ precipitates in Ni superalloys undergo a succession of morphology changes from spheres to cubes during the heat treatments preceding the actual use phase. At the very beginning, the γ’ precipitates nucleate as spheres to minimize the surface area [16]. As the particles grow, the misfit strain energy induced by the lattice and modulus mismatch between the γ  and the γ’ phase increases, and the precipitates become cuboids as the reduction in strain energy more than compensates for the

Dislocations in initial microstructure

For the commercial Ni single crystal superalloys with the typical microstructure of aligned cubic γ’ embedded in γ matrix, the well-organized γ/γ’ coherent initial microstructure is obtained by a multi-step solution and aging treatment followed by a slow air-cooling process. The density of dislocations in the initial microstructure will be at a sufficiently low level after the high-temperature heat treatment and the as-processed γ/γ′ microstructure can be approximately considered as a

The dependence of minimum creep rate on interfacial dislocation density

As shown in Fig. 4, the minimum creep rate of commercial Ni single crystal grades increases linearly with the spacing of the interfacial dislocations [11,57,65]. This phenomenological relationship can give an intuitive first-order connection between the creep properties and the dislocation properties. However, the parameter dislocation spacing is highly dependent on the accuracy of experimental observation. Moreover, the application of this relationship to quantify the creep behaviour of

Conclusions

  • (1)

    The high performance of Ni single crystal superalloys during high temperature low stress creep service, is primarily controlled by the combined effects of microstructural evolution, namely the formation of rafting lamellae, and the dislocation behaviour, i.e., the well-arranged dislocation network located on the γ/γ′ interface.

  • (2)

    During the secondary creep stage which takes longest time of creep life, two main recovery mechanism based on dislocation migration dominate the process. One is

Acknowledgements

This work was financially supported by the National Natural Science Foundation of China (No. 51722101) and the Key Research and Development Project (No. 2017YFB0703001).

References (94)

  • T. Murakumo et al.

    Acta Mater.

    (2004)
  • L. Shui et al.

    Mater. Sci. Eng. A

    (2007)
  • R.A. Ricks et al.

    Acta Metall.

    (1983)
  • W. Johnson et al.

    Acta Metall.

    (1988)
  • S. Socrate

    D.M. Parks. Acta Metall. Mater.

    (1993)
  • A. Epishin et al.

    Acta Mater.

    (2001)
  • J.V. Goerler et al.

    Acta Mater.

    (2017)
  • X. Tan et al.

    Mater. Sci. Eng. A

    (2011)
  • T. Gabb et al.

    Mater. Sci. Eng. A

    (1989)
  • R. Ricks et al.

    Acta Metall.

    (1983)
  • Z. Zhu et al.

    Acta Mater.

    (2015)
  • M. Ignat et al.

    Acta Metall. Mater.

    (1993)
  • M. Kamaraj et al.

    Mater. Sci. Eng. A

    (2001)
  • K. Serin et al.

    Mater. Sci. Eng. A

    (2004)
  • A. Epishin et al.

    Acta Mater.

    (2000)
  • A. Epishin et al.

    Mater. Sci. Eng. A

    (2009)
  • C. Rae et al.

    Acta Mater.

    (2007)
  • R. Reed et al.

    Acta Mater.

    (1999)
  • B. Fedelich et al.

    Mater. Sci. Eng. A

    (2009)
  • T.M. Pollock et al.

    Acta Metall. Mater.

    (1992)
  • A. Pineau

    Acta Metall.

    (1976)
  • D. Arrell et al.

    Scr. Metall. Mater.

    (1994)
  • J.Y. Buffiere et al.

    Acta Metall. Mater.

    (1995)
  • N. Matan et al.

    Acta Mater.

    (1999)
  • N. Ratel et al.

    Acta Mater.

    (2006)
  • J. Svoboda et al.

    Acta Mater.

    (1998)
  • B. Fedelich et al.

    Comput. Mater. Sci.

    (2012)
  • T. Tinga et al.

    Comput. Mater. Sci.

    (2009)
  • M. Gururajan et al.

    Acta Mater.

    (2007)
  • D. Li et al.

    Acta Mater.

    (1998)
  • G. Boussinot et al.

    Acta Mater.

    (2010)
  • N. Zhou et al.

    Acta Mater.

    (2007)
  • N. Zhou et al.

    Acta Mater.

    (2014)
  • L.T. Mushongera et al.

    Acta Mater.

    (2015)
  • Y.H. Wen et al.

    Acta Mater.

    (2010)
  • J. Zhang et al.

    Scr. Mater.

    (2003)
  • T. Link et al.

    Acta Mater.

    (2000)
  • R.R. Keller et al.

    Scr. Metall. Mater.

    (1993)
  • H. Gabrisch et al.

    Acta Mater.

    (2000)
  • G. Eggeler et al.

    Acta Mater.

    (1997)
  • J. Zhang et al.

    Acta Mater.

    (2005)
  • J. Zhang et al.

    Scr. Mater.

    (2003)
  • Z. Luo et al.

    Mater. Sci. Eng. A

    (2003)
  • R. Srinivasan et al.

    Acta Mater.

    (2000)
  • J.B. le Graverend et al.

    Int. J. Plasticity

    (2014)
  • H. Yang et al.

    Comput. Mater. Sci.

    (2015)
  • Z. Zhu et al.

    Acta Mater.

    (2012)
  • Cited by (29)

    • Subsurface damage in laser-assisted machining titanium alloys

      2023, International Journal of Mechanical Sciences
    View all citing articles on Scopus
    View full text