ɤ-TiAl alloy : Tensile creep deformation behaviour and creep life at 832

Creep deformation in single phase ɤ-TiAl alloy manufactured using different processing techniques has been an extensively studied topic since the late 1970s. The present work revisits the original work on understanding the tensile creep deformation behaviour of wrought singlephase ɤ-TiAl alloy by Hayes and Martin [1] and is aimed to develop an understanding of steady state creep. Besides, it is also aimed to investigate the creep life for stress levels of 69.4 and 103.4 MPa at 832 C using Monkman-Grant [2] approach.


Introduction
A series of reports on the room temperature tensile deformation behaviour of ɤ-TiAl alloys (L10 structure) show that the near-ɤ, two phase compositions having ~ 48 at.% Al possess very high toughness (ultimate tensile strength~ 844-1010 MPa and tensile ductility~ 3-4.6% at room temperature) [1], [3]- [10]. Extensive creep deformation studies have been carried out on a number of two phase near-ɤ-TiAl alloys produced by various processing routes. In addition, compression creep based investigations have been performed on a number of single phase ɤ-TiAl alloys and a number of literatures have reported that minimum strain rates during creep testing at different regimes of temperature and stress, hugely influence the grain size of single phase ɤ-TiAl alloys (during creep testing) [11]- [24]. Minimum strain rate of such intermetallic-based alloys may be defined in terms of Mukherjee-Bird-Dorn (MBD) equation [19], [25], [26].
The work of Hayes and Martin [1] discusses an analysis of the minimum strain rate deformation of a wrought single phase ɤ -TiAl alloy within temperature range of 760-1000 ⁰ C and stress range of 32-345 MPa. The work by Hayes and Martin [1] also predicts the main mechanism for creep rupture at different temperatures, in a given stress range, using Larson-Miller (L-M) [27] and Monkman-Grant (M-G) plots [2] and extensive microstructural characterisation using Optical microscope (OM) and the Transmission Electron Microscope (TEM). Besides, a recent report (from the author) on the tensile creep deformation behaviour of ɤ-TiAl has determined the stress exponent and creep activation energies between 760 ⁰ C and 900 ⁰ C at 69.4 and 103.4 MPa [28]. Moreover, based on creep activation energies and stress exponents, it has been reported that there is a transition from dislocation-glide to dislocation-climb controlled creep at very low creep stress levels (~ 66.68 MPa) [28] and that there is no steady-state creep observed for both interrupted and uninterrupted creep tests at 832 ⁰ C [1]. However, there is

Figs. 1(a and b)
show the variation of ε̇ with ε (in both linear and logarithmic scales respectively) for three cases viz. uninterrupted creep testing and interrupted creep testing with termination strains of 0.18 and 0.5%. Fig. 1(b . 1) indicating that the material (used in the present work) does not exhibit resistance to creep deformation for the aforementioned deformation condition. Besides there is a limited regime of tertiary creep observed for creep tests terminated a 0.18% and 0.5% strain ( Fig. 1). In addition, the material shows a dominant tertiary creep regime for the case of uninterrupted tensile creep testing (Fig. 1). Secondary (or steady-state creep) is not observed in any case (Fig. 1). In other words, the interrupted creep tests lead to a higher rate of vacancy formation at grain boundaries (GBs) normal to the tensile stress subsequently followed by void growth and coalescence leading to an intergranular fracture, as mentioned in Refs. [30]- [32].
Explanation of the above tendencies during interrupted and uninterrupted tensile creep test (at 832 °C ) is subject to extensive microstructural investigations which is beyond the scope of the present discussion. For uninterrupted creep testing, there is a dominant tertiary creep regime (Fig. 1). Moreover, it has been reported that the extent of tertiary creep hugely influences the creep life of the material [31]. To this end, creep life determination has been performed for samples (with uninterrupted creep testing) at 832 °C . of the M-G curve (in Fig. 2) is calculated as -1.33 which suggests that the mechanism of creep rupture is power-law breakdown at 832 °C (using the criteria mentioned in Ref. [14], [16], [33]- [37]). Hayes and Martin [1] and Saha [28] have reported that ε̇ decreases with increasing stress levels (from 69.4-103.4 MPa) at 832 °C . Moreover, Saha [28] has reported that the creep mechanisms operating at different stress levels (69.4 and 103.4 MPa) at 832 °C are independent of each other and hence, are not sequential. Hence, combining the trend observed between ε̇ with tr ( Fig. 2) with the aforementioned reports (Refs. [1], [28]), it may be inferred that M-G based approach to determine the creep life also predicts that tr decreases with increasing stress levels (from 69.4 to 103.4 MPa) at 832 °C . been normalised with rupture strain (εr) determined from Fig. 1. A justification of the above normalisation (for M-G plots) has been provided in Ref. [27].
Based on the Ashby's model [13], [14], [38], [39], the constitutive equation for NH (or lattice diffusion creep) is given as: where, ̇ is the strain rate, is the lattice diffusivity, b is the burgers vector, d is the grain size, k is the Boltzmann constant, σ is the creep stress and G is the shear modulus of the material.
Similarly, using Ashby's model, the constitutive equation for Coble creep (or GB diffusion creep) is given as: where, is the GB width and is the activation energy for GB diffusion. Both and are highly temperature-dependent [14]. However, the magnitude of is higher than that of at lower homologous temperatures whereas the magnitude of is higher than that of at higher homologous temperatures [40], [41]. This is because GBs offer a higher pathway for diffusion at lower temperatures when compared with that of the lattice [42]. Moreover, a comparison of the equations (3) and (

Conclusions
Based on the present work, the following may be concluded: •