Solidification rate driven microstructural stability and its effect on the creep property of a polycrystalline nickel-based superalloy K465

Journal Pre-proof Solidification rate driven microstructural stability and its effect on the creep property of a polycrystalline nickel-based superalloy K465 Xiaotong Guo, Stoichko Antonov, Fan Lu, Weiwei Zheng, Xiaofei Yuan, Jonathan Cormier, Qiang Feng PII:

S0921-5093(19)31316-4

DOI:

https://doi.org/10.1016/j.msea.2019.138530

Reference:

MSA 138530

To appear in:

Materials Science & Engineering A

Received Date: 8 August 2019 Revised Date:

3 October 2019

Accepted Date: 8 October 2019

Please cite this article as: X. Guo, S. Antonov, F. Lu, W. Zheng, X. Yuan, J. Cormier, Q. Feng, Solidification rate driven microstructural stability and its effect on the creep property of a polycrystalline nickel-based superalloy K465, Materials Science & Engineering A (2019), doi: https://doi.org/10.1016/ j.msea.2019.138530. This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. Please note that, during the production process, errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain. © 2019 Published by Elsevier B.V.

1

Solidification rate driven microstructural stability and its effect

2

on the creep property of a polycrystalline nickel-based

3

superalloy K465

4 5

Xiaotong Guoa, Stoichko Antonova,*, Fan Lua, Weiwei Zhenga, Xiaofei Yuana,

6

Jonathan Cormierb, Qiang Fenga,*

7 8

a

9

Laboratory for Advanced Metals and Materials, University of Science and Technology

Beijing Advanced Innovation Center for Materials Genome Engineering, State Key

10

Beijing, Beijing 100083, China

11

b

12

ISAE-ENSMA, BP40109, 86961 Futuroscope - Chasseneuil Cedex, France

13

*

Institut Pprime, UPR CNRS 3346, Department of Physics and Mechanics of Materials,

Corresponding author: [email protected]; [email protected]

14

15

Abstract

16

Different solidification rates may cause a significant difference in microstructure

17

and creep property in different locations of turbine blades in aircraft engines. In this

18

study, the effect of solidification rate on microstructural stability and creep property

19

was revealed in a polycrystalline nickel-based superalloy K465. The K465 alloy was

20

cast as a turbine blade, solid bars and hollow tubes. A larger cross section size caused

21

a slower solidification rate in the blade shank and bar than that of the blade airfoil and

22

tube. Microstructural characteristics and corresponding stress rupture properties under

23

975 °C/225 MPa were investigated after thermal exposure at 900 °C for 300 h to 1000

24

h. Plate-like µ phase formed only in the interdendritic regions of the blade shank and

25

the bar, but not in the blade airfoil and the tube after thermal exposure. The

26

precipitation of µ phase was mainly responsible for the much worse stress rupture

27

property of the bar in comparison with the tube. The microsegregation degree, 1

28

chemical composition of γ matrix and precipitates including γ’ phase and various

29

carbides, and dislocation configurations were examined. Compared to the tube, a

30

slower solidification rate caused a higher degree of microsegregation, coarser γ’

31

precipitates and carbides, as well as a much higher dislocation density in the bar after

32

standard solution treatment. The formation of µ phase was stress-induced and

33

attributed to the remaining dislocations in the interdendritic regions. A longer solution

34

treatment was suggested to effectively suppress the formation of µ phase in the bar

35

and blade shank for practical applications. These results provide a guidance for the

36

manufacturing and evaluation of microstructural degradation of turbine blades made

37

from conventionally cast polycrystalline nickel-based superalloys.

38

Keywords: Turbine blade; Microsegregation; Microstructure; µ phase; Dislocations

39 40

1. Introduction

41

Polycrystalline cast nickel-based superalloys are widely used for aero-engines

42

turbine blades [1, 2], which are subjected to complex elevated temperature and stress

43

couplings during service [1, 3]. At present, blades are generally designed with

44

complicated geometrical configurations to achieve the requirements for sustaining

45

such high loads and temperatures. This inevitably results in the variation of

46

cross-section size at different locations, such as the shank and airfoil tip. The size

47

difference can result in a variation of thermal mass, which have an influence on

48

casting parameters such as solidification rate, and ultimately affect the resulting

49

microstructure. Zheng and Cai [4] reported that the γ’ solvus temperature in the

50

thickest area was nearly 30 °C higher than that in the thinnest area for Udimet 500

51

cast turbine blade. Pal et al. [5] showed that topologically close-packed (TCP) phases

52

were prone to precipitate only in some specific locations of a serviced René N5

53

turbine blade. Such studies show that the casting section size can significantly affect

54

the local microstructural stability of a blade after casting [1, 6], solution-treatment [7],

55

during service and even after rejuvenation [5]. 2

56

Generally, solid bars are used to evaluate the microstructure and mechanical

57

properties of a given alloy. However, thin-wall specimens such as hollow tubes are

58

considered to be more suitable for mimicking the creep behavior of the hollow airfoil

59

of turbine blades [6, 8-10]. K465 superalloy (namely M963 superalloy or ЖС6У

60

superalloy) is a conventionally cast nickel-based superalloy that has been widely used

61

in turbine blades and vanes of aircraft engines. Until now, most creep property studies

62

of K465 superalloy utilize solid bar samples, which may not be the representative of

63

the blade airfoil. Such studies generally report microstructural degradation, such as

64

coarsening and dissolution of γ’ precipitates, decomposition of MC carbides, as well

65

as the formation of M6C and M23C6 carbides occurring during thermal exposure in the

66

850-1050 °C temperature range [11-14]. Additionally, plate-like µ phase precipitation,

67

which causes significant creep resistance degradation of K465 alloy, has been

68

reported during thermal exposure at 900 °C of bulk samples [11, 12]. However, no µ

69

phase has been observed and reported by microstructural evolution studies, utilizing

70

the airfoils of serviced K465 turbine blades [15].

71

The addition of refractory elements, such as W and Mo, promotes the

72

precipitation of TCP phases such as µ phase in nickel-based superalloys [2]. TCP

73

phases are always detrimental to the creep durability of superalloys due to their brittle

74

nature, and alloy design typically aims at suppressing/limiting their formation [16].

75

Generally, µ phase precipitates from the supersaturated matrix during solidification

76

[17] or under elevated temperature and/or stress [18, 19], and thus most studies on the

77

precipitation behavior of µ phase focus on the effect of alloy composition [17, 19-22].

78

However, limited studies have explored the precipitation behavior of µ phase in

79

different locations of the same turbine blade.

80

The objective of this study is to explain the solidification rate related

81

microstructural differences produced in different locations of the same blade during

82

casting, and explore the effect of such difference on the microstructural evolution and

83

creep behavior of a K465 superalloy blade. Additionally, it aims at elucidating the 3

84

reason for the observation of µ phase in solid bars and its absence in airfoils produced

85

from the same material after thermal exposure at 900 °C. This study will be beneficial

86

for the optimization of solidification conditions and thermomechanical processing

87

schemes of aero-engine components.

88

2. Experimental

89

The stock material was vacuum induction melted (VIM) and cast by Beijing

90

Institute of Aeronautical Materials into turbine blades, solid bar (representing the

91

blade shank) or hollow tube (representing the blade airfoil) shapes using the same

92

investment casting technology. The cross-section width of the blade shank was ~12

93

mm, while the airfoil thickness was ~2 mm. The cylindrical solid bars were 100 mm

94

long with a corresponding diameter of 14 mm. The hollow tubes (cast to size, rather

95

than machined) were also 100 mm long with an intermediate gauge length of 36 mm,

96

11 mm in outer diameter and 2 mm in wall thickness. The 2 mm thickness was

97

purposely selected to match the thickness of the airfoil. Subsequently, the cast blades,

98

bars and tubes were subjected to the standard solution heat treatment at 1210 °C for 4

99

h followed by air cooling [23]. Fig. 1(a) shows the sketch of a K465 turbine blade,

100

and the blade samples were sectioned at the shank and leading edge in the middle of

101

airfoil. The blade, bars and tubes were then thermally exposed at 900 °C for 300-1000

102

h followed by air cooling. This temperature reasonably represents the normal service

103

temperature of the airfoil. To evaluate the effect of microstructural evolution on creep

104

resistance after thermal exposure, the stress rupture lives of the bar and tube were

105

compared under 975 °C/225 MPa in air. The stress rupture tests were only conducted

106

on the tubes under 975 °C/225 MPa, considering that the stress rupture properties of

107

the bars have been reported in Ref. [12]. The sketch of the stress rupture specimens of

108

the tube is shown in Fig. 1(b). Two to four specimens were tested for each heat

109

treatment condition for statistical significance.

110 4

111 112

(a)

(b)

113

Fig. 1 (a) Sketch of a K465 turbine blade, where the red lines indicate the observation

114

positions of the shank and airfoil; (b) sketch of the stress rupture specimen of the tube

115

of K465 superalloy.

116

The chemical composition of all materials was measured by NCS Testing

117

Technology Co., Ltd, and is listed in Table 1. The compositions of the dendrite cores

118

and interdendritic regions were characterized by a JXA-8100 electron probe

119

microanalyzer (EPMA) for ten locations randomly selected for each specimen, where

120

the testing spot didn’t include interdendritic carbides or γ/γ’ eutectic pools.

121 122

Table 1 Measured alloy compositions of the turbine blade, bar and tube of K465

123

superalloy after the standard solution treatment (wt%). Ni

W

Co

Cr

Al

Ti

Mo

Nb

C

Blade

61.77

9.90

9.50

8.66

5.10

2.52

1.44

0.95

0.16

Bar

60.07

10.40

10.01

8.42

5.65

2.73

1.51

1.06

0.15

Tube

60.81

9.96

9.88

8.67

5.38

2.59

1.55

1.00

0.16

124

Samples for microstructural characterization were sectioned and prepared using

125

the standard metallographic techniques. The methods used for the etching of γ’

126

precipitates, carbides and µ phase have already been described in Ref. [11]. All

127

microstructures were observed using ZEISS AXIO Imager A2m optical microscope 5

128

(OM) and a ZEISS SUPRA 55 field-emission scanning electron microscope (FE-SEM)

129

operating at an accelerating voltage of 15 kV. The area fractions of γ’ precipitates

130

were quantified using the standard point count method according to Chinese standard

131

GB/T 15749-2008. Three to five representative locations were considered for the

132

quantification of γ’ precipitates.

133

For dislocation observation and phase identification, the foils for transmission

134

electron microscopy (TEM) were prepared by manually grinding 3 mm disks to a

135

thickness of 60 µm, followed by electrochemical thinning in a twin-jet polisher with a

136

solution of 90 vol.% ethanol and 10 vol.% perchloric acid. A FEI TECNAI F30 TEM

137

operating at 300 kV was used to observe dislocations and identify the carbides

138

through selective area diffraction (SAD). Site-specific samples for TEM-EDS analysis

139

were prepared by focused ion beam (FIB) using a FEI Helios Nanolab 600i. A JEOL

140

JEM-2100F TEM equipped with Oxford Instruments energy dispersive X-ray

141

spectroscopy (EDS) detector was used to obtain the chemical composition of γ matrix,

142

γ’ and µ phases. In addition, to obtain a statistical mass fraction and composition of

143

carbides, the physicochemical phase analysis method was adopted and conducted by

144

NCS Testing Technology Co., Ltd.

145

3. Results

146

3.1 Microstructure and evolution of the turbine blade

147

Fig. 2 shows the typical microstructure of K465 turbine blade. OM images in Figs.

148

2(a) and (b) show the dendrite morphology of blade shank and airfoil after the standard

149

solution treatment. The shank exhibited much larger dendrites, and the secondary

150

dendrite arm spacing (λ2) of the shank (150.6 µm) was ~3 times larger than that of the

151

airfoil (56.6 µm). SEM-BSE images of the interdendritic microstructures in the shank

152

and airfoil are shown in Figs. 2(c)-(f) at a higher magnification. Both locations

153

contained script-like gray-contrast MC carbides and white-contrast M6C carbides in

154

the interdendritic regions after the standard solution treatment, while the carbides and 6

155 156

157 158

(a)

(b)

(c)

(d)

(e)

(f)

159 160

161 162 163

Fig. 2 Typical microstructure of K465 turbine blade: OM images of dendrite

164

morphologies in the shank (a) and airfoil (b) after the standard solution treatment;

165

SEM-BSE images of interdendritic microstructures in the shank (c) and airfoil (d)

166

after the standard solution treatment, and in the shank (e) and airfoil (f) after thermal

167

exposure at 900 °C for 1000 h. 7

168

γ’ precipitates in the shank were significantly larger than those in the airfoil (Figs. 2(c)

169

and (d)). Extremely distinct interdendritic microstructures were observed in the shank

170

and airfoil after thermal exposure at 900 °C for 1000 h (Figs. 2(e) and (f)). A

171

needle-like µ phase had formed in the shank apart from carbides such as MC and M6C

172

carbides (Fig. 2(e)), while the airfoil was dominated by carbides especially M23C6

173

carbides (Fig. 2(f)). This distinct interdendritic microstructural evolution at different

174

locations of the same blade prompted a further investigation into the involved

175

mechanisms and the effect on the stress rupture life.

176

3.2 Microstructure and evolution of the solid bar and hollow tube

177

3.2.1 Microstructure after the standard solution treatment

178

Precipitation of µ phase in solid bars at 900 °C has also been reported by different

179

groups in Refs. [12, 13], suggesting the microstructural similarity between the blade

180

shank and a solid bar at this specific temperature. Considering the obvious

181

microstructural differences in the shank and airfoil (Figs. 2(e) and (f)), a cast-to-size

182

hollow tube with the same thickness as the blade airfoil was chosen to simulate the

183

airfoil microstructure.

184

Figs. 3(a) and (b) are OM images showing the dendrite morphologies and grain

185

structures (inset) in the bar and tube after the standard solution treatment. The grains

186

and dendrites in the bar were much larger than those in the tube, and the value of

187

secondary dendrite arm spacing in the bar and tube was around 96.6 and 47.7 µm,

188

respectively. Figs. 3(c) and (d) present the interdendritic microstructures in the bar and

189

tube, script-like MC and M6C carbides also distributed in the interdendritic regions,

190

while the carbides in the bar were much coarser than those in the tube. The general

191

microstructure of the solid bar was similar to that of the blade shank (Figs. 2(c) and

192

3(c)), while that of the hollow tube can be considered as the representative of that of

193

the blade airfoil (Figs. 2(d) and 3(d)). Figs. 3(e) and (f) reveal γ’ morphologies in the

194

dendrite core regions of the bar and tube, cuboidal γ’ precipitates were observed in

195

both the bar and tube, and they were more regular than those in interdendritic regions. 8

196

197 198

(a)

(b)

(c)

(d)

(e)

(f)

199 200

201 202 203

Fig. 3 Typical microstructure of the bar and tube of K465 superalloy after the

204

standard solution treatment: OM images of dendrite morphologies and grain structures

205

(inset) in the bar (a) and tube (b); SEM-BSE images of interdendritic microstructures

206

in the bar (c) and tube (d); SEM-SE images of γ’ morphologies in the dendrite core

207

regions of the bar (e) and tube (f).

208 9

209

Multiple small γ’ precipitates distributed in the γ channel formed during the cooling

210

process of the standard solution treatment. The small γ’ precipitates were not of concern

211

due to their rapid dissolution during thermal exposure. The quantitative

212

characterization of γ’ precipitates was thus performed in the dendrite core regions.

213

Table 2 lists the area fractions (Af) of γ’ precipitates in the dendrite core regions of the

214

bar and tube after the standard solution treatment. The Af of γ’ precipitates in the bar

215

and tube was similar and quantified as 61.7% and 63.2%, respectively.

216 217

Table 2 Area fractions of γ’ precipitates in the dendrite core regions of the bar and

218

tube of K465 superalloy after the standard solution treatment and after thermal

219

exposure at 900 °C for various times (%). Heat treatment condition

Bar

Tube

Standard solution treatment

61.7±1.4

63.2±2.2

900 °C/300 h

67.4±1.8

66.1±1.7

900 °C/500 h

67.4±2.4

66.8±0.6

900 °C/1000 h

67.8±1.2

66.1±1.5

220 221 222

To further evaluate the assumption that the blade shank and airfoil can be

223

simulated by a solid bar and a hollow tube, respectively. The dendrite compositions of

224

the turbine blade, bar and tube after the standard solution treatment were determined

225

via EPMA, and are summarized in Table 3. The degree of microsegregation was

226

evaluated using the microsegregation coefficient, k’, i.e., the ratio of the concentration

227

(in wt%) of an element in the dendrite core to the that in the interdendritic region [24],

228

and the results are shown in Table 3 and Fig. 4. A k’ value larger than 1 indicates that

229

the element segregates to the dendrite core region, while a k’ lower than 1 indicates

230

segregation to the interdendritic region. The results illustrated that the solid solution

231

strengthening elements such as W, Co, Cr and Mo segregated to the dendrite core

232

region, while γ’-forming elements, Al and Ti, segregated to the interdendritic region 10

233

(Table 3 and Fig. 4), in agreement with the typical segregation behavior during

234

casting [25]. The elemental segregation behavior in the blade shank and bar was

235

similar and more severe than that of the blade airfoil and tube, which were also

236

similar to each other (Fig. 4). Fig. 5 reveals the measured compositions of γ matrix and

237

γ’ phase in the dendrite core and interdendritic regions of the bar and tube after the

238

standard solution treatment using TEM-EDS. The results indicated that Co, Cr, W and

239

Mo segregated to the γ matrix, while Al and Ti segregated to the γ’ phase. The

240

segregation behavior of Nb was not obvious due to its lower content. More

241

importantly, the compositions of the γ matrix and γ’ phase were similar in all the

242

investigated regions of the bar and tube.

243 244

Table 3 Dendrite compositions of the turbine blade, bar and tube of K465 superalloy

245

after the standard solution treatment determined by EPMA (wt%), as well as

246

microsegregation coefficients, k’, the ratio of the concentration (in wt%) of an element

247

in the dendrite core to that in the interdendritic region.

Blade

Region

Ni

W

Co

Cr

Al

Ti

Mo

Nb

Dendrite core

59.85

11.38

10.60

7.97

5.14

1.94

1.38

0.54

Interdendritic region

62.93

7.13

10.03

7.32

5.86

2.49

1.16

0.68

k’

0.95

1.60

1.06

1.09

0.88

0.78

1.19

0.79

Dendrite core

60.19

9.22

10.56

9.99

5.11

2.16

1.24

0.64

Interdendritic region

61.66

9.02

10.43

9.51

5.47

2.35

1.21

0.66

k’

0.98

1.02

1.01

1.05

0.93

0.92

1.02

0.97

Dendrite core

60.08

9.71

10.54

9.90

5.34

2.24

1.56

0.63

Interdendritic region

63.01

7.86

9.76

8.53

6.05

2.80

1.28

0.70

k’

0.95

1.24

1.08

1.16

0.88

0.80

1.22

0.90

Dendrite core

60.24

9.35

10.63

10.15

5.27

2.29

1.34

0.63

Interdendritic region

61.50

9.11

10.19

9.32

5.52

2.43

1.25

0.68

k’

0.98

1.03

1.04

1.09

0.95

0.94

1.07

0.93

shank

Blade airfoil

Bar

Tube

248 11

249 250 251

252 253

Fig. 4 Microsegregation coefficients, k’, of the turbine blade shank and airfoil, bar

254

and tube of K465 superalloy after the standard solution treatment.

255 256

257 258

(a)

(b)

259

Fig. 5 Measured compositions of γ matrix (a) and γ’ phase (b) in the dendrite core and

260

interdendritic regions of the bar and tube of K465 superalloy after the standard

261

solution treatment using TEM-EDS (at. %).

12

262 263

Figs. 6(a)-(f) are TEM images showing the typical dislocation configurations in

264

the bar and tube after the standard solution treatment. Very limited dislocations existed

265

in the dendrite core regions of the bar or tube, as shown in Figs. 6(a) and (b). The

266

parallel fringes along the γ’ precipitates were moiré fringes due to the inclined γ/γ’

267

interface, as marked with a small white arrow in Figs. 6(a)-(d). Figs. 6(c)-(f) show the

268

dislocation morphologies in the interdendritic regions of the bar and tube, where denser

269

dislocations can be observed compared to the dendrite core regions (Figs. 6(a) and (b)).

270

Multiple single dislocations distributed in the γ matrix of the interdendritic region in

271

the bar, along with several dislocation loops (Fig. 6(c)), while only several single

272

dislocations were present in the interdendritic region of the tube (Fig. 6 (d)). Figs. 6(e)

273

and (f) illustrate the morphologies of interdendritic carbides, where MC carbides were

274

surrounded by M6C carbides and γ’ film. Dislocations accumulated at the interface of

275

carbides and γ’ film, while a higher dislocation density was observed along the bigger

276

carbides in the bar. Overall, limited literatures reported dislocation configurations in the

277

solution treated microstructure, while Fig. 6 confirms that a higher dislocation density

278

occurred in the interdendritic regions of the bar in comparison with the tube of K465

279

superalloy after the standard solution treatment.

280

3.2.2 Microstructural evolution during thermal exposure

281

Fig. 7 represents the SEM-SE images of γ’ precipitates in the dendrite core

282

regions of the bar and tube after thermal exposure at 900 °C for 1000 h, and the area

283

fractions of γ’ precipitates are listed in Table 2. Coarsening and coalescence of γ’

284

precipitates occurred during thermal exposure, maintaining a somewhat cuboidal

285

morphology. The size of γ’ precipitates in the bar was still larger than that in the tube

286

(Figs. 7(a) and (b)), arising from a larger original size (Figs. 3(e) and (f)). The Af of γ’

287

precipitates in the dendrite core regions of the bar and tube remained constant at

288

various thermal exposure times after an initial increase of 6% and 3%, respectively

289

(Table 2). 13

290 291

(a)

(b)

(c)

(d)

(e)

(f)

292 293

294 295 296

Fig. 6 TEM images of typical dislocation configurations in K465 superalloy after the

297

standard solution treatment: the dendrite core regions of the bar (a) and tube (b); the

298

interdendritic regions of the bar (c) and tube (d) and the interdendritic carbides in the

299

bar (e) and tube (f).

14

300 301

(a)

(b)

302

Fig. 7 SEM-SE images of γ’ precipitates in the dendrite core regions of the bar (a) and

303

tube (b) of K465 superalloy after thermal exposure at 900 °C for 1000 h.

304

Figs. 8 (a)-(d) are SEM-BSE images showing the typical morphologies of the

305

interdendritic regions in the bar after thermal exposure at 900 °C for various times.

306

No obvious microstructural evolution of the carbides was observed after thermal

307

exposure for 300 h, while some needle-like phase precipitated (Fig. 8(a)). The

308

needle-like phase also formed on MC carbides and extended into the γ matrix, as

309

depicted in Fig. 8(b). Upon 500 h thermal exposure, the size of MC carbides

310

decreased and the area fraction was noticeably lower (Fig. 8(c)), indicating the

311

degeneration of MC carbides. Meanwhile, some M23C6 carbides were observed in the

312

vicinity of MC and M6C carbides. Fig. 8(d) shows that more M6C and M23C6 carbides

313

formed after thermal exposure for 1000 h, and some needle-like phase was observed

314

to grow from M6C carbides. Fig. 8(e) illustrates the three-dimensional morphologies

315

of needle-like phase after deep etching, the needle-like phase was found to possess a

316

thin plate-like morphology with multiple plates intersecting at each other. Fig. 8(f)

317

shows the TEM morphology and corresponding SAD pattern (inset) of the plate-like

318

phase, which was identified as µ phase with a rhombohedral structure. The

319

composition of µ phase in the bar after thermal exposure at 900 °C for 1000 was

320

determined by TEM-EDS and summarized in Table 4. The µ phase was rich in W, Co,

321

Cr and Mo.

15

322 323

(a)

(b)

(c)

(d)

(e)

(f)

324 325

326 327 328

Fig. 8 SEM-BSE images of typical morphologies of the interdendritic regions in the

329

bar of K465 superalloy after thermal exposure at 900 °C for 300 h with a lower

330

magnification (a) and a higher magnification (b), for 500 h (c) and 1000 h (d); (e)

331

SEM-SE image showing the three-dimensional morphologies of plate-like µ phase in

332

the bar after thermal exposure at 900 °C for 1000 h; (f) TEM image and the

333

corresponding TEM-SAD pattern (inset) of µ phase in the bar after thermal exposure

334

at 900 °C for 1000 h. 16

335

Table 4 Chemical compositions of carbides in the bar and tube of K465 superalloy

336

after the standard solution treatment and after thermal exposure at 900 °C for 1000 h,

337

measured by using physicochemical phase analysis (at. %), as well as the composition

338

of µ phase in the bar after thermal exposure at 900 °C for 1000 h determined by

339

TEM-EDS (at. %). Condition

Standard

W

Ni

Cr

Co

Mo

Ti

Nb

C*

MC

9.29

0

1.14

0

2.32

25.68

11.57

50.00

M 6C

30.60

15.31

11.57

6.00

6.78

13.04

2.41

14.29

MC

9.62

0

1.18

0

2.45

25.14

11.61

50.00

M 6C

31.34

14.41

11.99

5.83

7.40

12.52

2.22

14.29

MC

8.27

0

0.81

0

2.27

26.20

12.45

50.00

M 6C

30.26

12.22

15.16

13.57

7.68

3.58

3.24

14.29

M23C6

4.47

2.90

67.48

1.56

2.91

0

0

20.69

µ

29.32

20.33

23.25

20.27

6.83

0

0

0

MC

8.22

0

1.13

0

2.35

25.44

12.87

50.00

M 6C

24.22

11.95

22.48

14.32

7.21

2.75

2.79

14.29

M23C6

5.00

3.52

66.24

1.74

2.81

0

0

20.69

Bar

solution treatment

Tube

Bar

900 °C/1000 h

Tube

340

*Note: The value of C is the theoretical value due to the inaccuracy in measuring light

341

elements using physicochemical phase analysis.

342

SEM images of typical morphologies in the interdendritic regions of the tube

343

after thermal exposure at 900 °C are shown in Fig. 9. M23C6 carbides formed after

344

thermal exposure for 300 h (Fig. 9(a)). Similar to the solid bar (Figs. 8(a)-(d)), MC

345

carbides degenerated with the prolonged exposure in favor of M6C and M23C6

346

carbides (Figs. 9(a) and (b)). However, the M23C6 carbide fraction in the tube was

347

much higher after thermal exposure even for 1000 h in comparison with the bar (Figs.

348

8(d) and 9(b)). Fig. 9(c) is a SEM-SE image depicting the three-dimensional

349

morphologies of carbides, indicating that multiple fine granular M23C6 carbides

350

formed on coarsen MC carbides. Both the M6C and M23C6 carbides in the tube have a

351

face-centered cubic structure through the analysis of TEM-SAD patterns (not shown

352

here), and are consistent with previous reports of this alloy [12, 14, 26]. 17

353 354

(a)

(b)

355 356

(c)

357

Fig. 9 SEM-BSE images of typical morphologies in the interdendritic regions of the

358

tube of K465 superalloy after thermal exposure at 900 °C for 300 h (a) and 1000 h (b);

359

(c) SEM-SE images showing the three-dimensional morphologies of carbides in the

360

tube after thermal exposure at 900 °C for 1000 h.

361 362

Fig. 10 shows the mass fraction of various carbides in the bar and tube after the

363

standard solution treatment and after thermal exposure at 900 °C for 1000 h. The mass

364

fraction of MC and M6C carbides in the bar (1.27% and 0.41%, respectively) were

365

similar to the tube (1.23% MC and 0.27% M6C) after the standard solution treatment.

366

After thermal exposure at 900 °C for 1000 h, the mass fraction of MC carbides

367

decreased to 0.75% and 0.67% in the bar and tube, respectively, with evident increase

368

in the precipitation of M6C and M23C6 carbides. However, more M23C6 carbides

369

precipitated in the tube (nearly twice that of the bar) consistent with the

370

microstructural observations (Figs. 8(d) and 9(b)). 18

371

The compositions of carbides in the bar and tube after the standard solution

372

treatment and after thermal exposure at 900 °C for 1000 h were measured through

373

physicochemical phase analysis, and are shown in Table 4. The results show that the

374

same type of carbides in the bar and tube showed no obvious difference in chemical

375

composition. MC carbides were rich in Ti and Nb, while M23C6 carbides were

376

composed of Cr, and W. Both µ phase and M6C carbide were rich in W, Co, Cr and

377

Mo, except that M6C carbide also contained C.

378

379 380

Fig. 10 Mass fraction of various carbides in the bar and tube of K465 superalloy after

381

the standard solution treatment and after thermal exposure at 900 °C for 1000 h.

382 383

3.3 Stress rupture property

384

Fig. 11 depicts the curves of stress ruputre lives in the bar and tube under

385

975 °C/225 MPa as a function of thermal exposure time at 900 °C. The data of stress

386

rupture property in the bar was reproduced from the results of Ref. [12]. The average

387

stress rupture life of the bar and tube after the standard solution treatment was 52 and

388

63 h, respectively, and upon thermal exposure, the rupture lives of both conditions

389

decreased. Interestingly, after thermal exposure at 900 °C, the stress rupture property

390

of the bar worsened to a much higher degree compared to the tube. The rupture life of

391

the bar drastically decreased with increasing the thermal exposure time, and was only

392

10 h after thermal exposure for 1000 h. Meanwhile, the rupture life of the tube 19

393

initially showed a slightly decreasing trend, but remained around 50 h even after

394

thermal exposure at 900 °C for 1000 h.

395

396 397

Fig. 11 The curves of stress rupture lives as a function of thermal exposure time at

398

900°C in the bar [12] and tube of K465 superalloy under 975°C/225 MPa.

399 400

To further investigate the strain effect on the precipitation behavior of µ phase,

401

the fracture microstructures at the longitudinal sections close to the fracture surface of

402

stress rupture specimens of the tube were examined and are shown in Fig. 12. No µ

403

phase exists in the fracture microstructure of the tube after the standard solution

404

treatment (Fig. 12(a)). However, µ phase formed in the fracture microstructure of the

405

tube after thermal exposure (Figs. 12(b) and (c)), but the quantity of µ phase in the

406

tube was significantly less than that of the fractured bar for equivalent thermal

407

exposure conditions.

408

409 410

(a)

(b) 20

411 412

(c)

413

Fig. 12 SEM-BSE images of the fracture microstructure at the longitudinal sections

414

close to the fracture surface of stress rupture specimens of the tube of K465

415

superalloy after the standard solution treatment (a) and after thermal exposure at

416

900 °C for 300 h (b) and 1000 h (c).

417

4. Discussion

418

The complexity and intricate design of turbine blades can lead to location specific

419

microstructure during the casting and the subsequent heat treatment, making it

420

difficult to predict the microstructural evolution and creep property of the turbine

421

blade as a whole. In the present study, the effect of the solidification rate on the

422

microstructural evolution and creep durability of a K465 superalloy turbine blade was

423

simulated and evaluated using solid bar and hollow tube samples. A solid bar and

424

hollow tube can be used to simulate the microstructure of the blade shank and airfoil,

425

respectively, based on the present study (Figs. 2, 3 and 4). The following discussion

426

will mainly focus on the solid bar and hollow tube for simplicity.

427

4.1 Effect of solidification rate on microstructure after the standard

428

solution treatment

429

During the casting of conventional nickel-based superalloys, γ dendrites form and

430

grow from the liquid, followed by MC carbides, sometimes M6C carbides and γ/γ’

431

eutectic solidify at the last [27]. Due to the dendritic solidification, segregation of

432

alloying elements is inevitable because of the difference in weight and elemental 21

433

diffusion rates in Ni. Increasing the solidification rate typically decrease the size of γ

434

dendrites and precipitates, as well as the degree of microsegregation [28], due to the

435

increased solid nucleation rate compared to dendrite growth kinetics [29]. During the

436

final stage of the casting process, shrinkage is also unavoidable due to the volumetric

437

contraction associated with the liquid to solid transition and further cooling. A higher

438

degree of microsegregation could cause a larger mismatch of thermal expansion

439

coefficients between dendrite core and interdendritic regions due to chemistry

440

difference and result in a higher stress concentration at dendrite core / interdendritic

441

region interfaces. Dislocations thus have to be generated to alleviate these stresses. As

442

the dendrite core regions are typically rich in the refractory alloying elements, which

443

strengthen the matrix, the interdendritic regions could be deformed relatively easily

444

and therefore dislocations typically form in this region during solidification and

445

cooling [30, 31]. Additionally, precipitates have significantly different compositions

446

and thermal expansion coefficients compared with the surrounding area, therefore

447

dislocations also form at their interface to accommodate the interfacial stresses. The

448

subsequent solution treatment aims at homogenizing the bulk composition and

449

annihilating the dislocations. However, dislocations sometimes cannot be completely

450

annihilated after solution treatment, especially in the interdendritic regions, as

451

reported by Nörtershäuser et al. in Ref. [30].

452

In the present study, Figs. 3(a) and (b) show that the λ2 value in the solid bar was

453

2 times that of the hollow tube. Typically, the dendrite size is described by the

454

secondary dendrite arm spacing (SDAS or λ2), which satisfies the following equation

455

[29]:

456

λ2=KV-1/3

(1)

457

where K value is a parameter related to the local chemical composition and can be

458

considered as a constant, and V is the local solidification rate. This indicates that the

459

solidification rate of the tube was about 8 times that of the bar during solidification.

460

Compared with the tube, the much slower solidification rate resulted in a larger 22

461

dendrite size in the bar (Figs. 3(a) and (b)), along with a higher degree of

462

microsegregation (Fig. 4), coarser γ’ precipitates and carbides (Figs. 3(c)-(f)). It can

463

be inferred that more dislocations existed in the bar than in the tube after the casting.

464

The standard solution treatment for K465 alloy (1210 °C/4h, air cooling) is only a

465

semi-solution treatment in order to decrease the cost and production times. During

466

this solution treatment, the γ’ precipitates in the dendrite core regions were completely

467

dissolved, however, some γ’ precipitates and carbides still remained in the

468

interdendritic regions due to higher microstructural stability. This resulted in

469

remaining dislocations in the interdendritic regions of the bar and tube after the

470

standard solution treatment (Figs. 6(c)-(f)). Due to the smaller degree of

471

microsegregation (Fig. 4), finer γ’ precipitates and carbides (Figs. 3(c)-(f)), the

472

interdendritic region in the tube had a much lower dislocation density than that in the

473

bar (Figs. 6(c)-(f)).

474

4.2 Effect of solidification rate on microstructural evolution

475

4.2.1 Microstructural evolution of γ’ precipitates, carbides and µ phase

476

The γ’ phase is the primary strengthening phase in Ni-based superalloys, hence

477

good γ’ phase stability is highly desired. During thermal exposure, local chemical

478

heterogeneity of alloying elements and minimization of the surface area cause the

479

coarsening and coalescence of γ’ precipitates [32]. This process is exacerbated by

480

internal chemical gradients and misfit strain [33]. The volume fraction of γ’ phase

481

plays a significant role on the mechanical properties of superalloys during thermal

482

exposure [34], and in general, a higher γ’ volume fraction typically means a better

483

stress rupture property [35].

484

In this study, Fig. 5 and Table 2 show that the composition of γ’ precipitates and

485

area fractions of γ’ precipitates in the dendrite core regions were similar in the bar and

486

tube after the standard solution treatment. During thermal exposure at 900 °C, the γ’

487

precipitates in the bar and tube coarsened and started coalesce, but kept a similar 23

488

morphology and area fraction (Fig. 7 and Table 2). Compared to the dendrite core

489

regions, the γ’ phase in the interdendritic regions had a higher area fraction due to the

490

enrichment of γ’ forming elements (Table 3), and the precipitates exhibited more

491

irregular morphology. Nevertheless, no obvious difference was identified between the

492

bar and tube, suggesting that the solidification rate had a negligible effect on the

493

formation and microstructural degradation of γ’ precipitates in the dendrite core and

494

interdendritic regions of K465 superalloy.

495

The main microstructural differences were observed for the carbide degradation

496

and µ phase formation in the interdendritic regions of the samples with different

497

geometries, as indicated in Figs. 8 and 9. The degradation of carbides during thermal

498

exposure is easy to understand, as MC carbides decompose to form M6C and/or

499

M23C6 carbides by releasing C into the γ matrix, where it reacts with Cr to form

500

M23C6 carbides, or with W and Mo to form M6C carbides [36]. As the diffusion rate of

501

Cr is higher than those of W, Co, and Mo, more M23C6 carbides precipitated compared

502

to M6C carbides at 900 °C (Fig. 10) [11, 37].

503

The formation of µ phase in the bar at 900 °C has also been previously reported

504

by Yuan et al. [11, 12]. It is generally accepted that the precipitation of µ phase in the

505

γ matrix can be attributed to two factors: the enrichment of µ phase formers (W, Mo

506

and Cr) [17, 38] and stress-induced formation [19, 20]. In the present study, it is

507

surprising that µ phase precipitated only in the bar, but not in the tube during thermal

508

exposure at 900 °C (Figs. 8 and 9). The forming elements of µ phase typically

509

segregated to the dendrite core regions (Fig. 4), hence it is expected that this phase

510

would form in these regions first, prior to the precipitation in the interdendritic

511

regions, but this was not the case in this study (Fig. 8). Due to the higher fraction of γ’

512

forming elements in the interdendritic regions, the γ forming elements in the matrix at

513

this location might be concentrated, and then drive the precipitation of µ phase.

514

However, this was ruled out by the measurement of the local precipitate and matrix 24

515

chemistries (Fig. 5), where no such enrichment was observed. This implies that the

516

precipitation of µ phase in this case may be stress-induced.

517

4.2.2 Formation mechanism of µ phase

518

Kontis et al. [39] reported that Co and Cr segregated to the dislocations during

519

high temperature deformation, and Wu et al. [25] found that the dislocations were

520

generally rich in γ formers and lack in γ’ formers. This may imply that the dislocations

521

in the interdendritic regions of the bar were rich in Co and Cr, and some W and Mo (γ

522

and µ phase formers). The enrichment would result in a high driving force for the µ

523

phase nucleation [19], while the general chemistry of the alloy, high W and Mo

524

contents, would provide the driving force for the growth during the thermal exposure.

525

Interestingly, some µ precipitates were observed to connect to the MC or M6C

526

carbides (Figs. 8(b) and (d)). As discussed previously and shown in Fig. 6(e), there is

527

an accumulation of dislocations between carbides and the surrounding area, and since

528

MC carbides are rich in Ti and Nb, this could result in local enrichment of W, Co, Cr,

529

and Mo in the vicinity of MC carbides [22, 40]. This further supports the nucleation

530

of µ phase driven by dislocation segregation. With the transformation of MC to M6C

531

carbides during thermal exposure, the µ phase can also nucleate on M6C carbides

532

because of the high similarity in crystal structure and composition of µ phase and

533

M6C carbide (Fig. 8(f) and Table 4) [41, 42]. Overall, multiple dislocations in the

534

interdendritic regions contributed to the formation of µ particles, once µ precipitates

535

have formed in the interdendritic regions, the internal chemical gradients contribute to

536

µ phase growth toward dendrite cores due to the microsegregation (Fig. 8(a)).

537

In the tube, the initial dislocation density was very low after the standard heat

538

treatment (Figs. 6(b) and (d)), which didn’t provide a strong enough driving force for

539

µ phase nucleation. As the MC carbides were finer than those in the bar, the

540

dislocation accumulation at the interface of MC carbides and γ matrix was also lower 25

541

(Fig. 6(f)). Therefore, MC carbides in the tube favored the formation of M6C and

542

M23C6 carbides over µ phase during thermal exposure at 900 °C (Fig. 9).

543

The stress-induced formation of µ phase is further confirmed by the stress rupture

544

specimens of the tube. As significant amount of dislocations accumulated during the

545

creep, µ phase can form during the testing of the thermally exposed samples (Figs.

546

12(b) and (c)). This is consistent with other studies observing accelerated TCP phase

547

formation with stress [20]. For the stress rupture test of the tube after the standard

548

solution treatment, no µ phase formed since the MC carbides had not started

549

decomposing. These results indicate that although the formation of µ phase is

550

stress-induced, some degree of microstructural degeneration is still needed.

551

Our results reveal that the formation of µ phase was mainly controlled by the

552

amount of remaining dislocations and degree of microsegregation in the sample after

553

the standard solution treatment. As the homogenization during solution treatment is

554

elemental diffusion controlled, a larger λ2 value leads to longer elemental diffusion

555

distances, which require more time to eliminate the microsegregation [43]. Hence

556

increasing the solution treatment time could annihilate dislocations and reduce the

557

degree of microsegregation more effectively, and thus suppress the formation of µ

558

phase. In order to validate this assumption, we solution treated the solid bar sample to

559

a much longer time for simulating the blade shank. Fig. 13 is a SEM-BSE image of

560

the typical microstructure of the interdendritic regions in the bar after solution

561

treatment at 1210 °C for additional 24 h and subsequently thermally exposed at 900 °C

562

for 500 h. It’s observed that no µ phase formed in the interdendritic region. This means

563

that prolonging the solution treatment time is one effective way to suppress the initial

564

formation of µ phase for practical engineering application, through annihilation of

565

dislocations and reducing the degree the microsegregation inherited from the

566

solidification of the alloy.

567 26

568 569

Fig. 13 SEM-BSE image of the typical morphologies of the interdendritic regions in

570

the bar of K465 superalloy after solution treatment at 1210 °C for additional 24 h and

571

subsequently thermally exposed at 900 °C for 500 h.

572 573

4.3 Effect of solidification rate on the stress rupture property

574

Generally, the creep properties of hollow tubes are worse than those of solid bars

575

with a much higher diameter [6, 8, 44, 45], which has been attributed to

576

microstructural changes [9, 46], more pronounced effect of number of grains [6, 10]

577

and oxidation [47, 48]. In the present study, the stress rupture life of the bar was

578

significantly less than that of the tube after thermal exposure at 900 °C (Fig. 11),

579

which means that a thin-wall related debit of properties cannot be considered for this

580

creep rupture condition. Although the larger casting geometry size caused a coarser

581

microstructure in the bar, the grain number in the cross sections of the bar and tube

582

was close and was basically between 5-10 (Figs. 3(a) and (b)). The differences in the

583

original microstructure was negligible, considering the stress rupture lives of the bars

584

and tubes was close after the standard solution treatment. Furthermore, no serious

585

oxidation occurred during the creep rupture tests, as all tests lasted less than 65 h (Fig.

586

11). Therefore, the difference in creep property between the bar and tube was mainly

587

caused by the difference in microstructural degradation during thermal exposure.

588

During thermal exposure at 900 °C, the progressive degradation of γ’ precipitates

589

and carbides can explain the gradually decrease in stress rupture lives of the tube with 27

590

increasing thermal exposure time. Microstructural characterizations showed that the

591

degradation degree of γ’ phase and MC carbides were close in the tube and bar after

592

thermal exposure at 900 °C (Figs. 7 and 10, Table 2). This means that the additional

593

deterioration of the creep rupture lives of the bar could be attributed to the large

594

amount of µ phase that formed in the interdendritic regions of the bar after thermal

595

exposure at 900 °C. The µ phase was rich in W, Co and Mo (Table 4), and its

596

formation depleted these strengthening elements in the matrix. In addition, multiple

597

cracks initiated at the interface of µ phase and γ matrix, based on the fracture

598

morphology [12], leading to a reduction in the load bearing area and deterioration of

599

creep resistance. Although µ phase also precipitated in the tube during the stress

600

rupture tests (Figs. 12(b) and (c)), its quantity was siginifcantly lower than that of the

601

bar, and its damaging effect was less severe.

602

In summary, this study used a solid bar and hollow tube to simulate the

603

microstructural evolution of a blade shank and airfoil, respectively. A full solution

604

treatment suppressed the formation of µ phase during thermal exposure. This study is

605

of high relevance to industry as it shows that the solidification rate during casting can

606

lead to striking differences in microstructural stability and mechanical properties, and

607

the optimization of the standard solution treatment need to be considered to ensure the

608

long-term safe service of K465 turbine blades.

609

5. Conclusions

610

The effect of solidification rate on microstructural stability and the stress rupture

611

property in K465 superalloy was revealed, combined with a systematic investigaton of

612

chemical

613

configurations of a hollow turbine blade, solid bars and hollow tubes made of K465

614

superalloy. The conclusions can be summarized as follows:

compositions,

microstructural

characterizations

and

dislocation

615

(1) For K465 superalloy turbine blade, the shank exhibited larger dendrites, a

616

higher degree of microsegregation and coarser carbides compared to the airfoil after 28

617

the standard solution treatment. A significant amount of µ phase formed only in the

618

shank, while much M23C6 carbides precipitated in the airfoil after thermal exposure at

619

900 °C.

620

(2) Considering the microsegregation behavior, standard solution-treated

621

microstructure and thermally-exposed microstructure at 900 °C, the blade shank and

622

bar were equivalent, while the blade airfoil and tube were similar. This suggests that

623

the bar and tube can simulate the microstructure of blade shank and airfoil,

624

respectively.

625

(3) A slower solidification rate in the bar was primarily responsible for the larger

626

dendrites, higher degree of microsegregation and coarser carides in comparison with

627

the tube. Dislocations remained in the interdendriitc regions of the bar and tube after

628

the standard solution treatment, while the bar contained much denser dislocations.

629

(4) The high dislocation density and significant microsegregation degree of the bar

630

were mainly responsible for the nucleation and growth of µ phase during thermal

631

exposure at 900 °C, respectively. µ phase could also nucleate on MC and M6C carbides

632

and grow towards the W and Mo rich dendrite cores. Meanwhile, the low dislocation

633

density in the tube suppressed the formation of µ phase and instead promoted the

634

formation of M6C and M23C6 carbides.

635

(5) The precipitation of µ phase caused a substantial decrease in stress rupture life

636

of the bar compared to the tube. A longer solution treatment successfully suppressed

637

the precipitation of µ phase in the bar during thermal exposure at 900 °C by

638

annihilating dislocations and reducing the degree of microsegregation more

639

effectively.

640

6. Acknowledgements

641

The authors are grateful to Yunrong Zheng and Paraskevas Kontis for their

642

professional advices and help. The financial supports provided by the National Key

643

Research and Development Program of China (Grant No. 2016YFB0701403), and 29

644

National High Technology Research and Development Program of China (Grant No.:

645

2012AA03A513) as well as the 111 Project (No. B170003).

646

References

647 648 649 650

[1] M.J. Donachie, S.J. Donachie, Superalloys: a technical guide (2nd edition), ASM international, Ohio, 2002, pp. 79-90. [2] T.M. Pollock, S. Tin, Nickel-based superalloys for advanced turbine engines: chemistry, microstructure and properties, J. Propul. Power. 22(2) (2006) 361-374.

651

[3] R.C. Reed, The superalloys: fundamentals and applications, Cambridge university

652

press, Cambridge, 2008, pp. 90-98.

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Declaration of interest statement We declare that we do not have any commercial or associative interest that represents a conflict of interest in connection with the work submitted.