Creep of Ti–5Al–5Mo–5V–1Fe–1Cr alloy with equiaxed and lamellar microstructures

Materials Science & Engineering A 651 (2016) 37–44

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Creep of Ti–5Al–5Mo–5V–1Fe–1Cr alloy with equiaxed and lamellar microstructures Xian Nie a, Huiqun Liu a,b,n, Xiaozhou Zhou a, Danqing Yi a,b, Baiyun Huang a, Zhan Hu a, Yanfei Xu c, Qi Yang d, Dingchun Wang d, Qi Gao d a

School of Materials Science and Engineering, Central South University, Changsha, Hunan 410083, PR China Key Lab of Nonferrous Materials, Ministry of Education, Central South University, Changsha, Hunan 410083, PR China c School of Mechanical Engineering, Xiangtan University, Xiangtan, Hunan 411105, PR China d BAOTI Group Ltd, BaoJi, Shanxi 721014, PR China b

art ic l e i nf o

a b s t r a c t

Article history: Received 3 September 2015 Received in revised form 16 October 2015 Accepted 24 October 2015 Available online 26 October 2015

The influence of microstructure on creep behavior of Ti–5Al–5Mo–5V–1Fe–1Cr (TC18) alloy at 350– 500 °C and tensile stress of 200–400 MPa was investigated by creep testing and transmission electron microscope (TEM). The creep mechanism was discussed by microstructure evolution. The results showed that, the alloy with lamellar structure showed better creep resistance than that of equiaxed structure. At 350–400 °C, the stress exponent n of samples with both structure was in the range of 1.2–1.9, the creep ¯ 〉 dislocations gliding on { 1010} prism plane, {0002} basal plane and { 1011} was controlled by a/3〈 1120 pyramidal plane. The creep activation energy of samples with both structure increased from 119 to 364 kJ/mol, with the increase of tensile stress from 200 to 400 MPa. At 450 °C, the stress exponent n of samples with both structure was about 3.0. At 500 °C, n is 4.3–4.8 under lower tensile stress, and 8.2–8.5 under higher tensile stress. At 450 °C and 500 °C, there was a dramatical rise of total creep strain, the jogged dislocation glided on {0002} basal plane was dominated, and the creep was controlled by dislocation climb. & 2015 Elsevier B.V. All rights reserved.

Keywords: Near-β titanium alloy Creep Microstructure Dislocation glide Creep mechanism

1. Introduction Near-β titanium alloys have been used as structural component in aircraft, such as under-carriage, and compressor disk in engine in 350–400 °C [1,2], Ti–5Al–5Mo–5V–1Fe–1Cr (TC18) is a typical near-β titanium alloy, with high strength-to-density ratio, high ductility, deep hardenability, excellent fatigue/crack propagation and corrosion resistance [3–6]. It is well known that for high temperature titanium alloy, creep resistance is one of the main criteria because it affects the service life and safety reliability of alloy. Although most of the structural titanium alloys are not easy to be failure induced by creep, Hooke's law (ε ¼ s/E) shows that titanium alloys with lower modulus than other structural materials can produce large elastic deformation by creep, after a long time service at intermediate temperature, the accumulation of elastic deformation probably induce plastic deformation and ultimate failure. TC18 alloy used at 350–400 °C will inevitably face with creep problem during long time service. n Corresponding author at: School of Materials Science and Engineering, Central South University, Changsha, Hunan 410083, PR China. E-mail address: [email protected] (H. Liu).

http://dx.doi.org/10.1016/j.msea.2015.10.092 0921-5093/& 2015 Elsevier B.V. All rights reserved.

The high temperature creep behaviors of some typical titanium alloys, such as Ti64 (Ti–6Al–4V), Ti6242 (Ti–6Al–2Sn–4Zr–2Mo), Ti6242s (Ti–6Al–2Sn–4Zr–2Mo–0.1Si), TC11 (Ti–6.5Al–1.5Zr– 3.5Mo–0.3Si) [7–12], have been extensively investigated. Previous studies have shown that microstructure is an important variable affecting creep strain [12–17]. For several typical microstructures of titanium alloy, some previous studies showed that basketweave and lamellar structures are more creep resistant than bimodel and equiaxed structures [12–15]. However, Gu and Zeng [12,16] found that TC11 alloy with bi-modal structure showed better creep resistance than fine acicular basket-weave structure. This result attributed to that very fine lamellar α phase in basketweave structure decreased creep resistance. Furthermore, some works [12,17] reported that creep rates increased with volume fraction of primary α and the width of α lath in near α-Ti alloys. The creep mechanism of titanium alloys is also strongly dependant on their microstructure. The creep mechanism discussed in literatures for metallic alloys includes dislocation gliding and climbing [18,19], grain boundary or interface sliding [17], viscous dislocation gliding either dragged by jogs or solute atoms [20] and diffusional creep [12]. Many researches [12,17,21,22] suggested that creep behavior of titanium alloy was controlled by deformation of α phase. Viswanathan et al. [20,23] reported that the creep

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¯ 〉 dislocation gliding on basal plane was dominated by a-type 〈1120 in α phase with few 〈a þc〉 type dislocation at α/β interface. Furthermore, some works [12,17,22,24] reported that dislocation networks and subgrains formed after creep test. The creep of TC11 alloy with bi-modal microstructure is controlled by dislocation climbing and changes to a jogged-screw creep model in basketweave structure [12]. However, the influence of microstructure on creep behavior of TC18 alloy has not been studied in detail up to now. Work on similar metastable near β-titanium alloy is also limited. In our previous work [25], the creep behavior of TC18 alloy with typical equiaxed microstructure was investigated. We found that there was a steep increase of creep strain above 400 °C and 300 MPa. This is probably the critical condition dominating creep property of TC18 alloy. The effects of modified microstructure as well as applied stress and temperature on creep behavior of TC18 alloy deserves further studying. In this paper, the microstructure of TC18 alloy was tailored to equiaxed and lamellar state through different heat treatments. Our aim is to understand the relationship between microstructure and creep properties of TC18 alloy at stress–temperature regime relevant to aircraft engine applications. It is also intended to discuss creep mechanism under different conditions.

Route II

Route I

Fig. 1. Heat treatment process of TC18 alloy bar.

Table 1 Heat treatment schemes and observed microstructure of TC18 samples. Heat treatment

Microstructure

830 °C/2 h/FC to 750 °C/2 h/AC þ600 °C/8 h/AC(Route I) 890 °C/2 h/FC to 750 °C/2 h/AC þ600 °C/8 h/AC(Route II)

Equiaxed Lamellar

2. Experimental Forged TC18 alloy bar of 35 mm in diameter provided by BAOTI Group Ltd. with following composition: 5.16Al, 5.08Mo, 5.08V, 0.95Fe, 0.95Cr, balance Ti, in wt%, was used as raw material. The β/ α þ β transus temperature (Tβ/α þ β) of this alloy is about 865 °C. In order to obtain desired microstructure, the as-received bar was solution treated at 830 °C (below Tβ/α þ β) and 890 °C (above Tβ/α þ β) for 2 h followed by cooling in furnace (FC) to 750 °C and held for 2 h prior to air cooling (AC). Then the bar was aged at 600 °C for 8 h, before air cooling again. Fig. 1 shows details of heat treatment process of TC18 alloy bar. Microstructures of heat treated alloy samples were observed by optical microscopy. The schemes of heat treatment and corresponding microstructure were summarized in Table 1. In order to investigate the influence of microstructure on creep behavior of TC18 alloy, specimens were cut and machined from the bars after heat treatment. Fig. 2 shows the size of cylindrical specimens, with diameter of 5 mm and gage length of 25 mm. Gage section of specimen was mechanically polished, ultrasonically cleaned in acetone and dried before testing. Creep tests were performed on a 30 kN capacity creep machine (RDL30) in air. Creep strain was continuously monitored using an extensometer. Temperature was monitored using a chromel–alumel thermocouple and maintained within 7 3 °C. Specimens were heated to desired test temperatures in electric resistance heating furnace at heating rate of  10 °C/min. Before creep testing, the specimens were kept at test temperature for 30 min to stabilize the temperature over the gage section. Creep testing was performed over 350–500 °C at applied stress levels ranging from 200 MPa to 400 MPa. The specimen was quenched in water immediately after creep testing interrupted at 100 h. The samples for microstructure examination were cut from plastic deformation area. Microstructural analysis of tested samples was performed by transmission electron microscopy (TEM, Tecnai G220, 200 kV). Ion beam thinning technique was used to prepare TEM samples.

Fig. 2. Geometry and dimensions of creep testing specimen (in mm).

3. Results and discussion 3.1. Microstructure of as-heat-treated samples Microstructure of TC18 samples can be controlled through processing as well as heat treatment. This alloy exhibits two important types of microstructure after different heat-treatments, as shown in Fig. 3. After heat treatment Route I and II (in Fig. 1 and Table 1), an equiaxed structure and lamellar structure were obtained, respectively. The microstructure in Fig. 3(a) is constituted of a large number of short rod-like primary α phase distributing within equiaxed β grains. After slow cooling from β to α þ β twophase region, nucleation and growth of lamellar like α-phase growing from β-grain boundaries are presented in Fig. 3(b). The resulting lamellar structure is like basket-weave or Widmanstatten microstructure.

X. Nie et al. / Materials Science & Engineering A 651 (2016) 37–44

39

Fig. 3. Optical microstructures (in TD direction) of TC18 alloy samples (a) with equiaxed structure after heat treatment Route I and (b) with lamellar structure after heat treatment Route II.

3.2. Creep behavior of samples with equiaxed and lamellar structure The creep curves and data of TC18 alloy samples with equiaxed structure obtained at 350–500 °C under different stresses are shown in Fig. 4 and Table 2, respectively. From Fig. 4, we see steadystate creep rate increased with testing temperature and applied stress. Total true creep strain increased slowly at lower temperature and/or stress, and it increased significantly at 450 °C and over 300 MPa. Experimental creep curves at 350 °C and 400 °C appear only the first two stages, namely the primary creep stage and the steady-state creep stage. When temperature increased to 450 °C and 500 °C, the tertiary stage of accelerated creep appeared. Creep is a thermal activation process and very sensitive to temperature, and occurs by dislocation movement, combining effects of atom thermal motion and stress. At the same temperature, the stress is the main factor affecting the creep, the main obstacle of dislocation motion is generated by long-range stress field, which is overcome through the internal shear stress. Therefore, at high external stress is, dislocations will overcome the hindrance of local obstacle with ease, and increase creep deformation. Fig. 5 shows the primary creep strain and total strain vs. temperature under applied stress of 200 MPa and 300 MPa. At 350 °C and 400 °C, primary creep strain and total strain increases slowly, the total strain increase from 0.351% to 0.501% under 300 MPa. However, the creep increase dramatically when temperature is over 400 °C, with total strain reached 1.428% (450 °C) and 9.56% (500 °C) at 300 MPa. At 200 MPa, the total strain and primary strain shows similar variation as those at 300 MPa. The creep curves of TC18 alloy samples with equiaxed and lamellar structures at 350–500 °C under 200 MPa and 300 MPa are illustrated in Fig. 6. The creep properties with lamellar structure are shown in Table 3. Compared with equiaxed structure, on the whole, samples with lamellar structure have lower total creep strain and steady-state creep rate, especially at temperature over 450 °C. Thus, it suggests that the creep resistance of lamellar structure is better than that of equiaxed structure. 3.3. Creep activation energy and stress exponent The creep of TC18 alloy sample consists of three stages, i.e. primary creep, steady state creep and trinary creep. Creep can be regarded as a process where as two phenomena are concurrent: strain hardening due to long range dislocation interaction and strain recovery due to thermal activation of short range dislocation movement. At primary stage, strain recovery gradually increases with creep deformation and creep strain rate gradually decreases until steady state stage. At trinary stage, the creep strain rate

Fig. 4. Creep strain versus time of TC18 samples with equiaxed structure at 350– 500 °C under applied stress of 200–400 MPa.

grows quickly until fracture. For these three stages, steady state stage occupied most of the procedure. Therefore, we need to discuss the mechanism of steady state creep regime. Knowledge of activation energy Qc and stress exponent n is necessary to establish the creep mechanisms. The temperature and stress dependence of steady-state creep strain rate is generally expressed in terms of Arrhenius-like equation [26]:

ϵ̇=Aσ n exp( −

Qc ) RT

(1) 1

where ϵ̇ is steady-state strain rate (s ), A is material constant, s is flow stress, R is gas constant, T is absolute temperature (K) and n is stress exponent, Qc is apparent activation energy for creep (kJ/ mol). In order to understand the mechanism of creep, the kinetic parameters n and Qc can be deduced from Eq. (1).

Q c = − R(

n=(

∂lnϵ )c ∂(1/T )

∂ln ϵ )T ∂ln σ

(2)

(3)

Fig. 7 shows Arrhenius plots of steady-state creep strain rate versus reciprocal of absolute temperature. The slope of the plot

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Table 2 Creep strain and steady-state creep rate of TC18 samples with equiaxed structure at different conditions. Creep conditions

Primary creep strain (%)

Total strain (%)

Steady state creep rate (s  1)

Creep time (h)

350 °C, 350 °C, 350 °C, 400 °C, 400 °C, 400 °C, 450 °C, 450 °C, 450 °C, 500 °C, 500 °C, 500 °C,

0.21 0.33 0.35 0.25 0.39 0.52 0.38 0.72 1.2 0.85 1.35 1.5

0.225 0.351 0.47 0.295 0.501 0.7 0.669 1.428 2.46 1.82 9.56 15.4 Fracture

2.91  10  10 4.77  10  10 1.37  10  9 1.95  10  9 3.71  10  9 6.85  10  9 5.50  10  9 1.89  10  8 4.48  10  8 2.85  10  8 1.98  10  7 2.26  10  6

100 100 100 100 100 100 100 100 100 100 100 11.5

200 MPa 300 MPa 400 MPa 200 MPa 300 MPa 400 MPa 200 MPa 300 MPa 400 MPa 200 MPa 300 MPa 400 MPa

Fig. 5. Primary creep strain and total strain vs. temperature of TC18 samples with equiaxed structure.

corresponds to apparent activation energy Qc of creep. The value of activation energy ranges from 119 kJ/mol to 364 kJ/mol for equiaxed structure, and from 124 kJ/mol to 306 kJ/mol for lamellar structure. The sample with equiaxed structure shows higher activation energy than that with lamellar structure under 300 MPa and 400 MPa, while shows a somewhat lower value at low stress of 200 MPa. In addition, two distinct regions are observed at 400 MPa between 400 °C and 500 °C for samples of both equiaxed and lamellar structures. Above 450 °C, the activation energy of samples with equiaxed and lamellar structure are 364 kJ/mol and 306 kJ/mol respectively. Below 450 °C, the activation energy becomes 152 kJ/mol and 138 kJ/mol respectively. Fig. 8 shows the creep data of samples with equiaxed and lamellar structure plotted as steady-state creep rate versus stress and reciprocal of temperature in double natural logarithmic scale. The slope of fitted curve corresponds to apparent stress exponent n. The stress exponents of sample with equiaxed structure fitted from the slopes in Fig. 8(a) are n ¼1.2, 1.8 and 3.0 for 350 °C, 400 °C and 450 °C respectively. At 500 °C, the change of n value is from 4.8 to 8.5 with increased applied stress. The change of activation energy and stress exponent indicated that the power low breakdown occurred at 500 °C and 300 MPa. From Fig. 8, it reveals that the stress exponent n is increased with temperature. The stress exponents obtained from equiaxed structure at all tested temperatures are consistent with dislocation creep. According to Fig. 8 (b), the stress exponents of samples with lamellar structure are similar with those of equiaxed structure. This indicates that change of microstructure has no significant influence on creep mechanism.

Fig. 6. Creep strain versus time of TC18 samples with equiaxed and lamellar structures at 350–500 °C under applied stress (a) 200 MPa and (b) 300 MPa.

3.4. Microstructure of samples after creep Fig. 9 shows typical TEM microstructure of TC18 alloy samples with equiaxed structure after creep testing under 300 MPa at different temperatures. Numerous dislocations were observed in these samples in α phase, while the dislocation density in β phase was very low. At 350 °C (Fig. 9a), a large number of short rod-like α existed within β matrix. It can be seen from Fig. 9(b), there are three orientations of slipping systems activated within α phase of

X. Nie et al. / Materials Science & Engineering A 651 (2016) 37–44

Table 3 Creep strain and steady-state creep rate of TC18 samples with lamellar structure at different conditions. Creep conditions Primary creep strain (%)

Total strain (%)

Steady state creep rate (s  1)

350 °C, 400 °C, 450 °C, 500 °C, 350 °C, 400 °C, 450 °C, 500 °C, 350 °C, 400 °C, 450 °C, 500 °C,

0.228 0.31 0.564 1.43 0.325 0.476 1.232 7.98 0.416 0.65 2.2 14 Fracture

3.56  10  10 2.12  10  9 4.87  10  9 1.90  10  8 4.48  10  10 3.24  10  9 1.47  10  8 1.07  10  7 1.07  10  9 5.62  10  9 3.78  10  8 1.86  10  6

0.21 0.23 0.35 0.82 0.305 0.33 0.7 1.2 0.47 0.52 1.1 1.35

-12

.5

-14

n=8 n=4.

-16

lnε (s )

200 MPa 200 MPa 200 MPa 200 MPa 300 MPa 300 MPa 300 MPa 300 MPa 400 MPa 400 MPa 400 MPa 400 MPa

41

8 n=3.0

-18

n=1.8 -20

n=1.2 -22 5.2

5.3

5.4

5.5

5.6

ln

-12 200 MPa 300 MPa 400 MPa

Q 64

=3

-14

5.7

-14

/m Q=

157

Q=

-20

15

kJ/

119

2k

mo

l kJ/m

J/m

lnε (s )

lnε (s )

ol

-18

ol

ol

n=4.

3 n=3.1

-18

n=1.9 -20

-22 1.25

n=1.3 1.30

1.35

1.40

1.45

1.50

1.55

1.60

1.65

1000/T (K )

5.4

5.5

5.6

=3 06 kJ /m ol Q=

-18

126

Q=

-20

Q= kJ/m

124

138

ol

kJ/m

kJ/

mo

l

ol

-22 1.30

1.35

1.40

1.45

1.50

1.55

1.60

5.7

5.8

5.9

6.0

(MPa)

Fig. 8. Double natural logarithmic plots of steady-state creep rate versus stress of samples with (a) equiaxed and (b) lamellar structure.

200 MPa 300 MPa 400 MPa

Q

lnε (s )

5.3

ln

-16

1.25

-22 5.2

-12

-14

6.0

8.2

-16

Q=

5.9

n=

kJ

-16

5.8

(MPa)

1.65

1000/T (K ) Fig. 7. Plots of double natural logarithmic steady-state creep rate versus reciprocal temperature of TC18 samples with (a) equiaxed and (b) lamellar structure.

HCP structure, which are defined as dislocation 1, 2 and 3, respectively. It can be determined by means of diffraction contrast → analysis according to invisibility criterion g ⋅b = 0 [27] where ¯ 〉, dislocations 1, 2 and 3 are distinguished as the type of a/3 〈1120 gliding on three different planes respectively. Dislocation 1, 2 and 3 glides on the {1010} prism plane, {0002} basal plane and {1011}

pyramidal plane, respectively. At 400 °C, as shown in Fig. 9(c), numerous dislocations piled up at α cell boundary suggesting a network rearrangement by gliding and climbing of edge segment. At 450 °C, the dislocation in interior of α phase contains many super jogs (indicated by rings in Fig. 9d) as a result of dislocation interactions, cross slipping or climbing. These structures are similar to those reported by Gollapudi et al. [28] and Viswanathan [20,29]. Under 300 MPa at 500 °C, cell structure formed by dislocation gliding and climbing (Fig. 9e), thus large numbers of subgrains with low-angle boundaries were observed in Fig. 9(f). On the whole, similar dislocation structures are observed for la¯ 〉 dismellar structure in Fig. 10. However, the a-type a/3 〈 1120 locations gliding on {0002} basal plane appears to be dominant (Fig. 10a and b) for lamellar structure creeping at 400 °C under 300 MPa. This may be the consequence of Burgers orientation relationship between α and β phase. The subgrains formed by entangled dislocations which were also observed in Fig. 10(c) for lamellar structure after creeping at 500 °C under 300 MPa. Therefore, at higher temperature (450–500 °C), crept structures for samples with both equiaxed and lamellar structure are mainly composed of dislocation slipping, climbing and grain boundary sliding.

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β α

β

1

α

2

α

200 nm

α

3

α 0.2 μm

α [2 1 1 0]

α 200 nm

jog

200 nm Fig. 9. TEM micrographs of dislocation substructures of TC18 samples with equiaxed structure after creep testing at 300 MPa for different temperatures: (a) 350 °C, ¯ 〉α axis. (b) enlarged α phase in (a), (c) 400 °C, (d) 450 °C, (e) 500 °C, (f) enlarged subgrain boundary morphology in (e). Beam direction is parallel to 〈 1120

X. Nie et al. / Materials Science & Engineering A 651 (2016) 37–44

43

0.2 μm

Fig. 10. TEM micrographs of dislocation substructures of TC18 samples with lamellar structure after creep testing at different conditions: (a) 400 °C, 300 MPa, (b) the ¯ 〉α zone axis of (a); (c) 500 °C, 300 MPa. corresponding electron diffraction pattern along the 〈 1120

3.5. Creep mechanism Reported self-diffusion activation energy values of α-Ti are in the range of 150–240 kJ/mol, and that of β-Ti 153 kJ/mol [29–32]. However, Koppers et al. [33] reviewed that the amount and nature of impurities or solution elements, particularly those fast diffusing impurities Fe and Ni, have dramatic effect on Ti self-diffusion and solute diffusion. In high purity α-Ti, activation energies of 303 and 329 kJ/mol were found for Ti self-diffusion and Al solute diffusion, respectively. The apparent activation energies obtained in present work are 119–364 kJ/mol from 350 to 500 °C. These values lie in the range of activation energy reported for self-diffusion of α-Ti and β-Ti. It is indicated that the creep mechanism in the temperature range is related to diffusion processes. The difference in stress exponent n indicates different creep mechanism. The stress exponents of present study revealed that the creep process is controlled by dislocation movement. At 350– 400 °C, stress exponents is 1 on o2, the creep is controlled by dislocation sliding. It has been proposed that creep in titanium at lower temperature is controlled by thermally activation instead of interstitial solute obstacles. In this temperature region, dislocation

sliding resistance is too large leading to low dislocation slide rate. Strain hardening rate due to dislocation sliding is slower than the strain recovery rate. The steady-state creep rate is determined by the dislocation sliding rate. At 450 °C and 500 °C, the stress exponents are n ¼3 and 4.8, respectively. Creep is controlled by dislocation climbing. These mechanisms are consistent with crept structures showed by TEM observation (see Figs. 9 and 10). At high temperature (350 °C and 450 °C), dislocation sliding becomes easy due to decrease of slipping resistance (see Figs. 9c, d and 10a, c), and reach to steady state with larger creep strain. With the temperature increasing to 500 °C under 300 MPa and 400 MPa, dislocation density increases obviously and intends to form subgrain boundary (see Fig. 9e and f) as well as stress exponent increasing drastically. Strain recovery rate induced by dislocation climbing is slower than that of strain hardening. The steady-state creep rate is determined by dislocation climbing rate. Furthermore, the dislocation climb rate increased due to the increase of diffusion rate and strain energy. These two reasons caused the significant increase of steady state creep rate and creep strain. The lamellar α and β phase in samples with lamellar structure is semi-coherent which meets Burgers orientation relationship,

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X. Nie et al. / Materials Science & Engineering A 651 (2016) 37–44

with small diffusion rate and dislocation climbing rate. However, the primary α and neighboring β phase with equiaxed structure has non-coherent relationship, which has high rate of dislocation climbing, leading to higher rate of creeping [13]. Therefore, for most Ti alloys, many studies [13,14] showed that the samples with lamellar and basket-weave structures have higher creep resistance at high-temperature followed by bimodal and equiaxed structures. In the present study, at low temperature the creep behavior difference between samples with different microstructure is not obvious because of the diffusion rates are relatively low. Therefore, the alloy with lamellar microstructure has better creep resistant than that of alloy with equiaxed structure.

4. Conclusions The high temperature tensile creep behaviors of TC18 samples with equiaxed and lamellar structure were compared at 350– 500 °C under 200–400 MPa. The major conclusions are summarized as follows. (1) The samples with lamellar structure showed better creep resistance than those with equiaxed structure. The total creep strain and steady creep rate increased with temperature and applied stress, while a dramatical rise of creep strain appeared between 400 °C and 450 °C. (2) The creep activation energy of samples with both microstructure increased from 119 to 364 kJ/mol, at 200–400 MPa. At 350 and 400 °C, the stress exponent n of samples with both microstructure was in the range of 1.2–1.9. The stress exponent n of samples with both microstructure was about 3.0 at 450 °C, 4.3–4.8 at 500 °C under lower applied stress, and 8.2–8.5 at 500 °C under higher applied stress. (3) At lower temperatures (350 and 400 °C), the creep of samples ¯ 〉 dislocawith equiaxed structure was controlled by a/3〈 1120 tions gliding on { 1010} prism plane, {0002} basal plane and ¯ 〉 {1011} pyramidal plane. Those for lamellar structure, a/3〈1120 dislocations glided on {0002} basal plane. At higher temperature (450 and 500 °C), for both microstructure, jogged dislocation gliding on {0002} basal plane became dominated, with the creep controlled by dislocation climbing.

Acknowledgment The authors wish to thank BAOTI Group Ltd for providing

forged TC18 alloy bar. This support of the National Natural Science Foundation of China under Grant no. 51401175 is gratefully acknowledged. The Key Projects in the National Science and Technology (No. 2014BAC03B05) and the partial financial support from the National Basic Research Program of China (Sub-contract no. 2014CB644001-2) were also acknowledged.

References [1] R.R. Boyer, Mater. Sci. Eng. A 213 (1996) 103–114. [2] D. Banerjee, J.C. Williams, Acta Mater. 61 (2013) 844–879. [3] O.M. Ivasishin, P.E. Markovsky, Y.V. Matviychuk, S.L. Semiatin, C.H. Ward, S. Fox, J. Alloy. Compd. 457 (2008) 296–309. [4] O.M. Ivasishin, P.E. Markovsky, S.L. Semiatin, C.H. Ward, Mater. Sci. Eng. A 405 (2005) 296–305. [5] O.M. Ivasishin, P.E. Markovsky, Y.V. Matviychuk, Metall. Mater. Trans. A 34 (2003) 147–158. [6] S. Nag, R. Banerjee, R. Srinivasan, J.Y. Hwang, M. Harper, H.L. Fraser, Acta Mater. 57 (2009) 2136–2147. [7] M.J.R. Barboza, C. Moura Neto, C.R.M. Silva, Mater. Sci. Eng. A 369 (2004) 201–209. [8] M.J.R. Barboza, E.A.C. Perez, M.M. Medeiros, D.A.P. Reis, M.C.A. Nono, F.P. Neto, C.R.M. Silva, Mater. Sci. Eng. A 428 (2006) 319–326. [9] M. Es-Souni, Mater. Charact. 45 (2000) 153–164. [10] J. Koike, K. Maruyama, Mater. Sci. Eng. A 263 (1999) 155–159. [11] X. Li, T. Sugui, B. Xianyu, C. Liqing, Mater. Sci. Eng. A 559 (2013) 401–406. [12] Y. Gu, F. Zeng, Y. Qi, C. Xia, X. Xiong, Mater. Sci. Eng. A 575 (2013) 74–85. [13] H. Mishra, P. Ghosal, T.K. Nandy, P.K. Sagar, Mater. Sci. Eng. A 399 (2005) 222–231. [14] M.A. Morris, T. Lipe, Intermetallics 5 (1997) 329–337. [15] H. Mishra, D. Satyanarayana, T. Nandy, P. Sagar, Scr. Mater. 59 (2008) 591–594. [16] L.Y. Zeng, G.J. Yang, Q. Hong, Y.Q. Zhao, Mater. Heat. Treat. 32 (2011) 81–85. [17] L. Ponsonnet, C. Quesne, R. Penelle, Mater. Sci. Eng. A 262 (1999) 50–63. [18] R.W. Hayes, P.L. Martin, Acta Met. Mater. 43 (1995) 2761–2772. [19] J.N. Wang, A.J. Schwartz, T.G. Nieh, D. Clemens, Mater. Sci. Eng. A 206 (1996) 63–70. [20] G.B. Viswanathan, S. Karthikeyan, R.W. Haye, M.J. Mills, Acta Mater. 50 (2002) 4965–4980. [21] R.W. Hayes, G.B. Viswanathan, M.J. Mills, Acta Mater. 50 (2002) 4953–4963. [22] M. Es-Souni, Mater. Charact. 46 (2001) 365–379. [23] G.B. Viswanathan, R.W. Hayes, M.J. Mill, Mater. Sci. Eng. A 319–321 (2001) 706–710. [24] W.H. Miller, R.T.S.E. Chen, Metall. Trans. A 18 (1987) 1451. [25] X.A. Nie, Z. Hu, H.Q. Liu, D.Q. Yi, T.Y. Chen, B.F. Wang, Q. Gao, D.C. Wang, Mater. Sci. Eng. A 613 (2014) 306–316. [26] A.M. Brown, M.F. Ashby, Scr. Met. 14 (1980) 1297–1302. [27] X. Li, T. Sugui, B. Xianyu, C. Liqing, Mater. Sci. Eng. A 529 (2011) 452–458. [28] S. Gollapudi, I. Charit, K.L. Murty, Acta Mater. 56 (2008) 2406–2419. [29] S. Karthikeyan, G.B. Viswanathan, P.I. Gouma, V.K. Vasudevan, Y.W. Kim, M. J. Mills, Mater. Sci. Eng. A 329–331 (2002) 621–630. [30] R.P.R.G. Malakondaiah, P.R. Rao, Acta Met. 29 (1981) 1263–1275. [31] I. Weiss, S.L. Semiatin, Mater. Sci. Eng. A 243 (1998) 46–65. [32] F. Dyment, C.M. Libanati, J. Mater. Sci. 3 (1968) 349–359. [33] M. Koppers, C.H.R. Herzig, M. Friesel, Y. Mishin, Acta Mater. 45 (1997) 4181–4191.