Creep behavior and microstructure evolution of directionally solidified superalloy IC10 at high temperature

ZHAO Xi-hong(赵希宏)^1,2, SUN Yue-jun(孙跃军)¹, HUANG Zhao-hui(黄朝辉)²,

GONG Sheng-kai(宫声凯)¹, XU Hui-bin(徐惠彬)¹

1. School of Materials Science and Engineering, Beijing University of Aeronautics and Astronautics,

Beijing 100083, China;

2. Beijing Institute of Aeronautical Materials, Beijing 100095, China

Received 20 April 2006; accepted 30 June 2006

Abstract: The creep behavior of directionally solidified (DS) superalloy IC10 was investigated under 192 MPa and 218 MPa at 980 ℃. Under the testing conditions, the marked creep characteristic of the superalloy is that the creep curve has a short primary and secondary stage and a long tertiary stage. Another creep characteristic is that the superalloy has excellent plasticity at high temperature. To study the creep behavior, the microstructure was observed by SEM and TEM. Different from other microstructure of Ni-base superalloys, superalloy IC10 forms incompletely rafted γ' phase during the creep processes. To understand the creep deformation mechanism of superalloy IC10, the movement of dislocations was analyzed. The results show that the dislocations moving in the γ matrix and climbing over the γ' precipitates is the main deformation mechanism under the experimental conditions.

Key words: superalloy IC10; creep; microstructure evolution; directional solidification; dislocation; plasticity

1 Introduction

The evolution of efficient gas turbine engine technology has brought about increased interest in high-temperature operating materials[1]. Ni-base superalloy has excellent high-temperature strength and hot corrosion resistance with the addition of refractory elements such as W, Cr, Ta, Mo, Re, Ru, and has been used as gas turbine blade material. The high-temperature strength of Ni-base supralloy depends strongly on the precipitation of the γ' phase[2].

It is well known that the γ' precipitates have the long-range ordered L1₂ structure based on Ni₃Al, and they are embedded in the fcc matrix phase of the γ phase (Ni solid solution). The directional coarsening of γ', which is called rafting, takes place in the creep of Ni-base superalloys at higher temperature. Some observations imply that rafting reduces the creep resistance at temperature below 1 000 ℃ and at high temperature[3,4]. The coarsening of γ' had been studied extensively[3-5]. Therefore, the creep behavior and structure stability must be considered by investigators and manufactures. In this paper, the creep behavior of a new DS superalloy IC10 at high temperature was investigated. The aim of this investigation is to show creep behavior and microstructure evolution of the superalloy IC10 at high-temperature and offer useful information for the manufactures.

2 Experimental

Rods were cut from a standard heat-treated superalloy along the [001] orientation. Test pieces were machined from the rods having (27±0.1) mm gauge length and (5.00±0.02) mm diameter. Tensile creep tests were performed using RD₂-3 creep testing machine. Temperature vibration was controlled within ±1 ℃ and the fluctuations along the gauge length were kept within ±3 ℃ during the whole experiments. Disk samples for microstructure observation were cut from the gauge length of the specimens deformed under different conditions along the horizontal and vertical directions. The microstructures were observed by using HITACHI S-3500N scanning electron microscope (SEM) and H-800 transmission electron microscope (TEM).

3.1 Creep deformation

For most of the Ni-base superalloys, the typical creep curve is composed of a short primary and tertiary stage and a long secondary stage[6, 7]. However, in the present study, the creep curve shape of the DS superalloy IC10 is quite different. As shown in Fig.1, after tested under 192 MPa and 218 MPa at 980 ℃, the creep curve shows a short primary and secondary stage then a long tertiary stage. Another creep characteristic of the tested superalloy is its excellent plasticity at high temperature with the maximum strain of about 34%.

Fig.1 Creep curves of superalloy tested under 218 MPa and 192 MPa at 980 ℃

3.2 Structure evolution

Fig.2 shows the microstructure of the tested superalloy with standard heat treatment. The microstructure consists of γ matrix, spherical γ' precipitates, γ-γ' eutectics and a small amount of carbides as shown in Fig.2(a). Carbides are mainly MC, in which M is substituted for Ta and Hf. The γ' precipitates distribute in the γ phase uniformly, as shown in Fig.2(b).

Fig.3 shows the changes of microstructure of the superalloy after tested for 4 and 30 h at 980 ℃. After crept 4 h, the γ' precipitates became coarser but kept spherical shape always compared with which under standard heat-treatment. No rafted structure was observed for the sample tested at the first stage of creep (tested for 4 h), as shown in Fig.3(a). Incompletely rafted γ' phase images perpendicular to the stress axis were observed at the secondary stage (tested for 30 h), accompanying with the existence of small spherical γ' precipitates (Fig.3(b)).

Fig.4 shows the microstructure is of the superalloy after failure at 980 ℃. There are no visible changes between the microstructure near the rupture area and far away from the rupture area. The majority of γ' precipita-

Fig.2 Initial microstructures of DS superalloy before creep testing

Fig.3 Development of γ' morphology at 980 ℃ under 218 MPa: (a) Interrupted for 4 h; (b) Interrupted for 30 h

tes grew normal to the applied stress, however, there are partial γ' precipitates grew along different orientations. No perfect rafted γ' was observed. Usually, γ' phase in Ni-base superalloy rafts rapidly during the first stage of creep, and the time which formed γ' phase becomes shorter

Fig.4 Morphologies of γ' phase of superalloy after failure, tested under 218 MPa at 980 ℃: (a) 0.2 mm from fracture; (b) 2 mm from fracture; (c) 6 mm from fracture; (d) 10 mm from fracture

with increasing temperature and stress. The γ' coarsen along the directions normal to the applied stress when the γ'/γ lattice mismatch δ is negative[8], where δ=(α_γ'-α_γ)/α_γ', in which α_γ' and α_γ  are the lattice parameters of γ' precipitates and γ matrix. The kinetics of the γ' coarsening depends on the thermodynamic driving force. The driving force mainly arises from the decrease of the interfacial energy. When the interfacial energy decreases, the smaller precipitates dissolve and the larger ones grow, and the kinetics are generally controlled by volume diffusion in which the solute is transferred through the matrix from the shrinking particles to the growing ones. Such coarsening behavior has been theoretically predicted by LIFSHITZ and SLYOZOV[9], and also independently by WAGNER[10]. Considering that in the present study, the temperature and applied stress are large enough to raft γ' phase[11],  the incomplete rafted γ' phase must be due to the insufficient lattice mismatch which has been reported by QIU[12].

Fig.5 shows the microstructure changes at 980 ℃ under different applied stresses. Under 218 MPa, the volume fraction of γ' phase became larger and the γ channel became narrower than that under 192 MPa. With the increasing of the applied stress, the alloy elements, which form γ' phase, diffused from γ matrix to γ' phase easily. For this season, the number of γ' precipitates increased and the partial γ matrix adjacent to γ' phase disappeared.

Fig.5 Microstructures of superalloy after failure at 980 ℃ under different applied stresses: (All specimens were sectioned parallelly to applied stress axis) (a) Under 192 MPa; (b) Under 218 MPa

3.3 Creep deformation mechanism

Generally, the creep mechanism changes with applied stress at high-temperature. Under low stress, because the dislocations are unable to cut through the spherical γ' particles, deformation is based on the movement of dislocations in the matrix. When the creep comes to the secondary stage, the lattice misfit can promote the formation of rafted γ'/γ structure and also result in dense γ'/γ interfacial dislocation networks[13-15]. The formation of fine and continuous γ' phase provides an ideal structure for creep resistance because the mobile dislocations will be hindered from passing these γ' phase and they will be forced to shear γ' phase[16]. All these make the superalloy has a short primary and a long secondary stages. Fig.6 shows the TEM microstructure of the specimen. It is shown clearly that dislocations migrate mainly within the γ matrix in creep processes. There is no evidence that dislocations cut through γ' particles because both SF (stack fault) and APB (anti- phase boundary) can not be observed in the γ' precipitates. This figure also shows that dislocations network formatted at the γ-γ' interfaces with little dislocations in the γ' precipitates.

Fig.6 TEM microstructure of superalloy IC10, crept at 980 ℃ and 218 MPa interrupted for 30 h showing dislocation climb mechanism

Therefore, at medium stresses (192-218 MPa) the dislocations are unable to cut through or bow between the γ' particles. Deformation is based now on the movement of dislocations in the matrix and the change of the γ' morphology. In this region the dislocations are forced to overcome the particles by climb.

3.4 Creep failure

Fig.7 shows the cracks morphologies near the fracture. During the creep process, cracks originated at the edge of carbides which distributed at the grain boundary, then extended along the grain boundary normal to the stress axis, at last been hindered by the dendritical boundary. According to this failure process, the extended velocity of cracks can be decreased largely.

Fig.7 Crack morphologies in tested superalloy after creep failure under 218 MPa at 980 ℃

For superalloy IC10, insufficient rafted γ' phase made the dislocations move easily in the γ matrix, so the superalloy has a short secondary and a long tertiary stages. The low extended velocity of cracks can elongate the tertiary stage also. As a result, the superalloy has a excellent plasticity at high temperature.

4 Conclusions

The creep behavior of DS superalloy IC10 was investigated at 980 ℃ under 192 MPa and 218 MPa. The typical creep curve showed shorter primary and secondary stages and a longer tertiary stage. There is no perfect rafted γ' phase observed during the creep processes.  Dislocations moving in the γ matrix and climbing over the γ' precipitates is the main deformation mechanism under the experimental conditions.

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(Edited by LONG Huai-zhong)

Corresponding author: XU Hui-bin; Tel: +86-10-82317117; Fax: +86-10-82338200; E-mail: xuhb@buaa.edu.cn