Microstructure and mechanical behavior of a novel Co20Ni20Fe20Al20Ti20 alloy fabricated by mechanical alloying and spark plasma sintering

Author’s Accepted Manuscript Microstructure and mechanical behavior of a novel Co20Ni20Fe20Al20Ti20 alloy fabricated by mechanical alloying and spark plasma sintering Zhiqiang Fu, Weiping Chen, Haiming Wen, Sam Morgan, Fei Chen, Baolong Zheng, Yizhang Zhou, Lianmeng Zhang, Enrique J. Lavernia www.elsevier.com/locate/msea

PII: DOI: Reference:

S0921-5093(15)30204-5 http://dx.doi.org/10.1016/j.msea.2015.07.052 MSA32590

To appear in: Materials Science & Engineering A Received date: 28 June 2015 Accepted date: 18 July 2015 Cite this article as: Zhiqiang Fu, Weiping Chen, Haiming Wen, Sam Morgan, Fei Chen, Baolong Zheng, Yizhang Zhou, Lianmeng Zhang and Enrique J. Lavernia, Microstructure and mechanical behavior of a novel Co20Ni20Fe20Al20Ti20 alloy fabricated by mechanical alloying and spark plasma sintering, Materials Science & Engineering A, http://dx.doi.org/10.1016/j.msea.2015.07.052 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. 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.

1

Microstructure and mechanical behavior of a novel

2

Co20Ni20Fe20Al20Ti20 alloy fabricated by mechanical alloying and

3

spark plasma sintering

4

Zhiqiang Fua,b*, Weiping Chena, Haiming Wenc, Sam Morganb, Fei Chend, Baolong Zhengb,

5

Yizhang Zhoub, Lianmeng Zhangd, Enrique J. Laverniab† a

6

School of Mechanical and Automotive Engineering, South China University of Technology, Guangzhou, Guangdong 510640, China

7 b

8

Department of Chemical Engineering and Materials Science, University of California at Davis, Davis, CA 95616, USA

9 c

10 11

d

Characterization Department, Idaho National Laboratory, Idaho Falls, ID 83415, USA

State Key Lab of Advanced Technology for Materials Synthesis and Processing, Wuhan University of Technology, Wuhan 430070, China

12 13 14 15

Abstract

16

A novel equiatomic Co20Ni20Fe20Al20Ti20 (at.%) alloy was designed and synthesized to study

17

the effect of high atomic concentrations of Al and Ti elements on the microstructure, phase

18

composition and mechanical behavior of high-entropy alloys (HEAs) fabricated by mechanical

19

alloying (MA) and spark plasma sintering (SPS). Following the MA process, the

20

Co20Ni20Fe20Al20Ti20 alloy was composed of a primary body-centered cubic (BCC) supersaturated

*

Corresponding author. E-mail address: [email protected] (Z. Fu)

†

Corresponding author. E-mail address:[email protected] (E.J. Lavernia) 1

21

solid solution and a face-centered cubic (FCC) supersaturated solid solution. However, following

22

SPS, a primary FCC solid-solution phase, a BCC solid-solution phase and a trace amount of Al3Ti

23

intermetallics were observed. Transmission electron microscopy (TEM) results confirmed the

24

presence of the FCC solid-solution phase, the BCC (B2-type) solid-solution phase and Al3Ti

25

intermetallics in the bulk alloy. The FCC and B2-type phases are ultrafine-grained, and Al3Ti

26

intermetallics is nano/ultrafine-grained. Our results suggest that consideration of a single existing

27

empirical design criterion is inadequate to explain phase formation in the Co20Ni20Fe20Al20Ti20 alloy.

28

Solid-solution strengthening, grain-boundary strengthening, twin-boundary strengthening, the

29

presence of the strong B2-type BCC phase, and precipitate strengthening due to the presence of a

30

trace amount of Al3Ti are responsible for the ultra-high compressive of ~2988 MPa and hardness of

31

~704 Hv. The strain-to-failure of ~5.8% with visible ductility is dominated by the FCC

32

solid-solution phase.

33

Keywords: High-entropy alloys; intermetallics; alloy design; microstructure; solid solution.

34

1. Introduction

35

Traditional alloy systems are based on one principal metallic element with small concentrations

36

of additional elements added to tailor microstructures and properties, which inherently limits the

37

number of alloy systems achievable due to the limited number of metallic elements in the periodic

38

table. More recently, a novel alloy design strategy involving the concept of multiple principal

39

elements with equal or nearly equal molar fractions of component elements, has been proposed and

40

the resulting materials have been termed high-entropy alloys (HEAs) [1]. Each constituent element

41

with a concentration between 5 and 35 at.% is considered as a principal element, and the presence

42

of five or more principal elements results in the fact that “solvent” or “solute” atoms are not distinct 2

43

[1]. A high entropy decreases the Gibbs free energy of mixing, Gmix , in HEA systems especially at

44

elevated temperatures based on the equation Gmix  H mix  TSmix [1,2,3]. As a result,

45

unexpectedly stable body- or/and face-centered cubic (BCC or/and FCC) or even hexagonal

46

close-packed (HCP) solid-solutions are formed [1,3,4].

47

HEAs

can

attain

expectedly high

solid-solution

strengthening

derived

from

the

48

multiple-principal elements which create strong lattice distortion and therefore high strength and

49

hardness are achievable in this novel class of alloys [1,3]. Elements possessing larger atomic radii

50

compared to other component elements, such as Al and Ti, are usually introduced into HEA systems

51

as a strategy to increase the effect of solid-solution strengthening [5-9]. However, Al and Ti have

52

good affinity with a variety of elements that are usually used in HEA systems, and accordingly high

53

atomic concentrations of them results in intermetallics and/or other complex phases, further

54

downgrading ductility and comprehensive mechanical properties [8-10]. Inspection of the literature

55

reveals that few investigations have focused on the HEA systems exhibiting high atomic

56

concentrations of Al and Ti. However, introducing a small amount of intermetallics into metals or

57

alloys would facilitate a reasonable reinforced strength [11,12]. Additionally, previous studies have

58

shown that equiatomic FCC structured CoNiFe alloy has good ductility [13,14]. Hence, on the basis

59

of the arguments described above, we designed a novel equiatomic Co20Ni20Fe20Al20Ti20

60

(CoNiFeAlTi) alloy, to involve a slightly higher concentrations of both Al and Ti into the

61

equiatomic CoNiFe alloy, with the aim of obtaining high solid-solution strengthening and a small

62

amount of intermetallic phases.

63

Inspection of the available literature reveals that casting and mechanical alloying (MA)

64

followed by sintering are the two most common processing routes for bulk HEAs [3,15]. The 3

65

references [15,16] show that typical defects of casting processes, such as phase segregation and

66

inhomogeneous microstructures, result in degradation of the mechanical properties of HEAs. The

67

combination of MA and sintering exhibit better homogeneity of microstructure compared to that of

68

casting and can avoid the potential segregation and inhomogeneous microstructures in HEAs

69

[17,18]. Furthermore, spark plasma sintering (SPS) can consolidate powders within short sintering

70

time, thereby achieving ultra-fine or even nanocrystalline grain sizes in the bulk alloys after

71

consolidation [19,20].

72

Thus, in view of the above discussion, we designed a novel Co20Ni20Fe20Al20Ti20 alloy in this

73

study to provide insight into the effect of high atomic concentrations of Al and Ti elements on the

74

alloying behavior, phase composition and evolution, microstructure and mechanical properties of

75

HEAs. The Co20Ni20Fe20Al20Ti20 alloy was fabricated via mechanical alloying followed by spark

76

plasma sintering.

77

2. Experimental procedures

78

The Co20Ni20Fe20Al20Ti20 alloy powders were processed by mechanical alloying of elemental

79

powders of Co, Ni, Fe, Al and Ti with particle sizes of ≤ 45 µm (-325 mesh) in a high-energy

80

planetary ball mill using the following procedure. First, elemental powders of greater than 99.7

81

wt.% purity were placed in stainless steel vials with tungsten carbide balls, without a process

82

control agent (PCA). A ball-to-powder weight ratio of 10:1 and a high-purity argon atmosphere

83

were applied during the whole MA process. Then, the powders were subjected to 4 hours wet

84

milling with high-purity ethanol as a PCA, followed after by 45 h of dry milling at 300 rpm. Before

85

analysis and sintering, the ethanol was removed by drying via evaporation in a vacuum oven for a

86

minimum of 24 hours. Bulk samples were prepared by SPS of the dried Co20Ni20Fe20Al20Ti20 alloy 4

87

powders followed by passing through a 75 mm sieve, using a Dr. Sinter 825 apparatus (Sumitomo

88

Coal Mining Co. Ltd., Japan). SPS was conducted at 1273 K for 8 min with a heating rate of ~90 K

89

min–1. During the entire SPS process, a constant pressure of 30 MPa was applied while maintaining

90

a vacuum pressure < 8 Pa. Specimens used for the subsequent testing were sectioned from the

91

sintered bulk samples by electrical discharge machining.

92

X-ray diffraction (XRD) studies of the powders and bulk samples were performed on a Bruker

93

D8 ADVANCE X-ray diffractometer with a Cu Kα radiation. A NETZSCH STA 449C differential

94

scanning calorimeter (DSC) was used to perform the thermal analysis under a high purity Ar

95

atmosphere, by performing from room temperature to 1323 K with a heating rate of 10 K min–1.

96

Phase composition of each bulk sample and their detailed chemical compositions were carried out

97

by a transmission electron microscope (TEM) with selected area electron diffraction (SAED),

98

performed on a JEOL JEM-2100 (Tokyo, Japan) operated at 200 kV with an energy dispersive

99

spectrometer (EDS). Specimens for TEM observations were prepared by mechanical thinning

100

followed by ion milling at ambient temperature. Uniaxial compression tests of cylindrical

101

specimens with the dimensions of Ø3 mm×4.5 mm were performed at ambient temperature in an

102

Instron 5500 testing system at an engineering strain rate of 1×10-3 s-1. Three nominally identical

103

specimens of each sample were compressed to obtain average value. Vickers hardness of bulk

104

samples were measured at minimum of 10 measurements for each specimen with a HVS-1000

105

digital micro-hardness tester with a load of 2942 mN.

106

3. Results and discussion

107

3.1 Alloying behavior

108

The X-ray diffraction (XRD) patterns of the mechanically alloyed (MA’ed) powders of the 5

109

Co20Ni20Fe20Al20Ti20 alloy with different milling times are presented in Fig. 1(a). Peaks

110

corresponding to elemental Co, Ni, Fe, Al and Ti are evident in the XRD pattern before milling.

111

After 6 h of milling, diffraction peaks associated with most component elements exhibit dramatic

112

decreases in intensity. As milling time increases to 15 h, diffraction peaks corresponding to Ni, Fe,

113

and Ti are still observed, whereas those corresponding to Al and Co are absent, indicating that Al

114

and Co have been dissolved. After 30 h of milling, no elemental components can be identified, and

115

peaks corresponding to both BCC and FCC solid-solution phases are evident, with the BCC phase

116

being the main phase. The alloying rates of the component elements are inversely correlated with

117

their melting points and ductility. It ascribes an element having a lower melting point exhibiting a

118

higher intrinsic diffusion coefficient in comparison with that having a higher melting point [17, 21].

119

On the other hand, a brittle element could be crushed more easily than a ductile element when they

120

have similar melting points, therefore resulting in acceleration of alloying rate [21]. Accordingly the

121

anticipated alloying rates in the equiatomic Co20Ni20Fe20Al20Ti20 alloy is: Al→Co→Ni→Fe→Ti,

122

which is consistent with the observations in previous studies [17,21]. Further milling after 30 h was

123

performed to ensure complete formation of a solid-solution and to refine its grain size. Moreover, 4

124

h of wet milling was carried out after 45 h of dry milling to achieve smaller sizes of powder

125

particles, which facilitate the subsequent sintering of the powders. The XRD patterns of the 45 h

126

and 49 h milled powders are similar to that of the 30 h powders, however, peaks are broadened,

127

attributable to grain refinement and an increment in lattice strain [17]. Fig. 1(b) displays the pattern

128

of a slow san of the 49 h milled powders in the range of 2θ=40-48°, and the fitting result verifies

129

that the asymmetrically overlapped peak contains two peaks corresponding to a main BCC

130

solid-solution phase and an FCC solid-solution phase, respectively. 6

131

3.2 Microstructure, phase composition and evolution

132

The XRD pattern of the bulk alloy after SPS is presented in Fig. 2(a) with that of the 49 h

133

milled powders for comparison. The bulk alloy consists of a primary FCC solid-solution phase, a

134

BCC solid-solution phase and a trace amount of Al3Ti intermetallic phase. In contrast, the primary

135

phase is the BCC solid-solution phase and there is no Al3Ti intermetallic phase in the 49 h milled

136

powders. These observations reveal phase transformation during SPS. A differential scanning

137

calorimetry (DSC) curve of the 49 h milled powders in a temperature range of room temperature to

138

1323 K is shown in Fig. 2(b). Two endothermic peaks at ~724.85 K and ~1062.75 K are visible,

139

suggesting that phase transformations may occur at these two temperatures. Based on the

140

endothermic peaks in Figs. 2(a) and (b), it is postulated that the energy absorption at ~724.85 K is

141

attributable to the phase formation of Al3Ti, and that at ~1062.75 K corresponds to a phase

142

evolution that results in the change in volume fractions of the FCC and BCC phases. Because of the

143

solubility extension generated during the non-equilibrium MA process [22], supersaturated

144

solid-solutions are formed in the milled powders, which consist of a primary BCC phase and an

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FCC phase. These metastable supersaturated solid-solution phases transform to equilibrium phases

146

during sintering or heating, leading to the above-mentioned phase evolution. The excessive amount

147

of energy stored in the milled powders in the form of grain boundaries with significant volume

148

fraction and dislocations with high density may reduce the activation energy for phase evolution

149

and accordingly facilitate its occurrence during sintering or heating [17,22].

150

Fig. 3(a) shows a bright-field transmission electron microscope (TEM) image of the bulk

151

Co20Ni20Fe20Al20Ti20 HEA after SPS at 1273 K for 8 min. The selected-area electron diffraction

152

(SAED) pattern of grain A is presented in Fig. 3(b), corresponding to that of an FCC structure along 7

153

[011] zone axis. Similarly, the SAED pattern of grain B is displayed in Fig. 3(c), corresponding to

154

that of an ordered BCC structure (B2-type structure) along [001] zone axis. In addition, the arrows

155

in Fig. 3(a) indicate Al3Ti intermetallics, which are also confirmed by chemical composition

156

analysis using EDS/TEM as listed in Table 1, i.e., ~71.03±1.38 (Al), 23.13±1.27 (Ti), 2.30±0.72

157

(Fe), 2.24±0.66 (Co) and 1.30±0.57 (Ni) at.%. As shown in Table 1, the chemical composition of

158

the FCC phase obtained by EDS/TEM is ~5.39±0.61 (Al), 8.92±0.73 (Ti), 31.69±0.90 (Fe),

159

26.85±0.34 (Ni) and 27.15±1.41 (Co) at.%. In contrast, the chemical composition of the B2-type

160

BCC phase is ~19.98±0.27 (Al), 17.37±2.56 (Ti), 12.31±3.22 (Fe), 25.71±0.42 (Ni) and 24.63±0.62

161

(Co) at.%. That is, the FCC phase is a Ni-Co-Fe-based solid-solution with some Al and Ti present,

162

whereas the B2-type BCC phase is a Ni-Co-Al-Ti-based solid-solution with some Fe present. This

163

result is consistent with previous reports that higher Al and Ti contents are favorable to formation of

164

BCC phases, and lower Al and Ti contents benefit formation of FCC phases in HEAs [3,15]. In

165

addition, ordered BCC phases, i.e., B2-type BCC phases, are observed in many Al- and

166

Ti-containing HEAs because of Al and Ti displaying strong bonding with 3d transition metals, such

167

as Cr, Fe, Ni, Co, etc. [6,8,15]. It is obvious that the FCC and B2-type BCC phases are ultra-fine in

168

grain size, and Al3Ti intermetallic phase is nanocrystalline or ultra-fine in grain size.

169

In addition, twins with nanoscale lamella thickness are observed in some FCC grains as shown

170

in Fig 4(a). Fig. 4(b) shows the SAED pattern of grain C along [011]M (matrix) and [011]T (twin)

171

zone axes, confirming the presence of twins. Previous studies have shown that twins are readily

172

formed in FCC phases of HEAs with low stacking fault energy [23,24].

_ _

173

Due to the lack of complete phase diagrams, it is challenging to design HEAs with desirable

174

phases, microstructures and properties, resulting in many attempts to develop empirical design 8

175

criteria by analyzing the available data on HEA systems. Although the available empirical criteria

176

for HEAs are not very accurate, researchers have relied heavily on them to design HEAs, and in

177

particular, the criterion proposed by Yang et al. has been applied in many HEAs [25]. In early

178

studies, Zhang et al. proposed a criterion for solid-solution formation in HEAs, that is, atomic

179

radius differences (δ), enthalpy of mixing ( H mix ) and entropy of mixing ( S mix ) should meet the

180

following

181

-2.685δ-2.54< H mix <-1.28δ+5.44 KJ mol–1. The parameters δ, H mix and S mix are defined as

182

follows [25,26],

conditions

simultaneously

δ

183

n

 c (1  r / r ) i 1

H mix 

184

S mix

[26]:

i

2

i

, r  i 1 ci ri n

n

 ij cic j 

i 1, i  j

JK-1mol-1,

>13.38

δ<4.6%

and

(1)

n

 4H

i 1, i  j

mix ij i

c cj

(2)

n

S mix   R (ci ln ci )

185

(3)

i 1



n

c  1 , ri is the atomic radius of the ith component. ci and c j are the atomic

186

Where

187

concentrations of the ith and jth components, respectively. R is the gas constant equal to 8.314

188

JK-1mol-1; and H ijmix represents the enthalpy of mixing the ith and jth components, which are from

189

the calculations on the basis of Miedema macroscopic model for ideal binary liquid alloys [25,27].

190

Subsequently, this criterion was amended as follows [28]: 12< S mix <17.5 JK-1mol-1, δ<6.4% and

191

-20< H mix <5 KJ mol–1. In addition, a less restrictive criterion, i.e., 11≤ S mix ≤19.5 JK-1mol-1,

192

δ≤8.5%, -22≤ H mix ≤7 KJ mol–1, was proposed by Guo et al. [29]. However, it is challenging to

193

determine the threshold values of these parameters.

194

i 1 i

In subsequent work, Yang et al. proposed that solid-solutions are formed when satisfying

9

TmΔS mix ≥1.1 ( Tm is the mole averaged melting point) and δ ≤ ΔH mix

195

simultaneously the conditions of Ω 

196

6.6% [25]. Recently, another criterion γ 

197

formation of solid-solutions in HEAs [30]. Small and Larg e indicate the solid angles around the

198

smallest and largest atoms, described as follows [30],

199

S m a l l1 

Small <1.175, was proposed by Wang et al. for the L arg e

2 2 (rL a g  r er )  r

, L arg e  1 

2 (rL a g  r er )

(rSmall  r )2  r 2 (rSmall  r ) 2

(4)

200

Here rSmall and rLagre are the atomic radii of the smallest and largest component atoms. In addition,

201

average valence electron concentration VEC=  ci (VEC )i , where ci and (VEC)i are the mole

n

i 1

202

fraction and valence electron concentration of the ith element, respectively, is regarded as a useful

203

parameter for predicting the crystal structures of solid-solution phases in HEAs [29,31]. According

204

to the proposition of Guo et al., single FCC and single BCC solid-solution phases would be formed

205

when VEC ≥ 8 and VEC < 6.87, respectively, and when 6.87 ≤ VEC < 8, both FCC and BCC

206

solid-solution phases would be formed [31].

207

According to the results presented above, the bulk Co20Ni20Fe20Al20Ti20 alloy after SPS

208

consists of a primary FCC solid-solution phase, a B2-type BCC solid-solution phase and a trace

209

amount of Al3Ti. As given in Table 2, the estimated values of H mix , S mix , δ , Ω, VEC and γ for

210

the Co20Ni20Fe20Al20Ti20 alloy are -26.40 kJ mol-1, 13.38 JK-1mol-1, 6.85%, 0.83, 6.8 and 1.018,

211

respectively. For this alloy, the entropy of mixing S mix =13.38 JK-1mol-1 is high, facilitating the

212

formation of solid-solutions [28,29]. Also the value of

213

Co20Ni20Fe20Al20Ti20 alloy, which satisfies the criterion proposed by Wang et al. for formation of

214

solid-solution phases in HEAs [30]. However, a trace amount of Al3Ti was identified in the bulk 10

γ

is 1.108 <1.175 for the

215

Co20Ni20Fe20Al20Ti20 alloy in addition to the two solid-solution phases, indicating that these two

216

empirical criteria involving the parameters of S mix and γ do not apply to this alloy. In addition,

217

the values of δ and Ω for this alloy are 6.85% and 0.83, respectively, which do not satisfy the

218

conditions of Ω≥1.1 and δ ≤ 6.6% proposed by Yang et al. for formation of solid-solution phases in

219

HEAs [25]. According to the criterion of Yang et al., intermetallics or other complicated phases are

220

anticipated to be formed in the Co20Ni20Fe20Al20Ti20 alloy, which is consistent with our

221

experimental results. Therefore, this criterion applies to the Co20Ni20Fe20Al20Ti20 alloy. Furthermore,

222

the value of VEC is 6.8 for the Co20Ni20Fe20Al20Ti20 alloy, very close to the threshold value of 6.87

223

for the formation of a single BCC solid-solution phase (VEC < 6.87) or a mixture of BCC+FCC

224

solid-solution phases (6.87 ≤ VEC < 8) [31]. One can assume that for VEC values close to the

225

threshold value, the formation of single BCC solid-solution phase or the formation of BCC+FCC

226

solid-solution phases are both possible, since the threshold value of 6.87 may not be accurate or

227

may be different for different alloys. Evidently, the consideration of a single existing empirical

228

design criterion is not adequate for predicting or explaining phase formation in the

229

Co20Ni20Fe20Al20Ti20 alloy prepared by MA and SPS. To achieve more adequate predication of

230

formation of solid-solution phases in HEAs, all of the existing criteria should be considered.

231

Table 3 shows atomic radii and enthalpy of mixing in binary equiatomic alloys calculated by

232

Miedema’s approach [27,29]. Ni, Co and Fe have similar atomic radii and electronegativity, which

233

is in favor of formation of solid solutions [27,29,32]. Therefore, the formation of an FCC

234

solid-solution phase rich in Ni-Co-Fe is anticipated, consistent with previous studies [13,14]. Al and

235

Ti exhibit a significantly negative enthalpy of mixing (-30 kJ mol-1) and possess larger atomic radii

236

than Fe, Ni and Co, and thus affinity is more preferred between these two elements. Moreover, the 11

237

enthalpies of mixing for different atom-pairs among Al, Ti, Co and Ni are significantly negative

238

ranging from -19 to -35 kJ mol-1. In particular, Ti and Ni have the most negative enthalpy of mixing

239

of -35 kJ mol-1 [27,32]. It appears that these elements have similar affinity leading to similar ability

240

to attract each other, that is, a trend of “mixing” which is referred to as “mixing entropy effect” or

241

“high entropy effect” [1,15], and accordingly resulting in the formation of the B2-type BCC

242

solid-solution phase enriched in Ni-Co-Al-Ti. Additionally, besides the formation of FCC and BCC

243

solid-solution phases, residual Al and Ti formed a trace amount of Al3Ti intermetallics, because of

244

their good affinity mentioned above.

245

3.3 Mechanical behavior

246

Fig. 5 displays the engineering stress-strain curve of the bulk Co20Ni20Fe20Al20Ti20 alloy under

247

compression at room temperature. The bulk Co20Ni20Fe20Al20Ti20 alloy exhibits an ultra-high

248

compressive strength (σmax) of ~2988 MPa, an ultra-high hardness of ~704 Hv and a strain-to-failure

249

(εf) of ~5.8%, respectively. Table 4 summarizes mechanical properties of the SPS’ed

250

Co20Ni20Fe20Al20Ti20 alloy and some typical SPS’ed HEAs at room temperature. Note that values of

251

εf for some alloys were re-estimated according to their engineering stress-strain curves under

252

compression at room temperature, which is attributed to incorrectly elastic deformation behaviors in

253

these HEAs. It is obvious that the SPS’ed Co20Ni20Fe20Al20Ti20 alloy exhibits the highest

254

compressive strength among the typical SPS’ed HEAs given in Table 4. In HEA systems, B2-type

255

BCC phases exhibits significantly higher strength and less ductility in comparison with FCC

256

structured phases [3,15]. Hence, the B2-type BCC contributes to strengthening the bulk alloy while

257

the FCC phase primarily contributes to ductility. Although Al3Ti is formed after SPS, the

258

engineering stress-strain curve of the bulk alloy shows a visible stage of plasticity originating from 12

259

the primary FCC solid-solution phase [15], and its plasticity is comparable with some typical

260

SPS’ed HEAs as given in Table 4 [7,17,33,36-39], Co0.5NiFeCrTi0.5, CoCrFeNiAl, AlFeNiCoCrMn,

261

for example. This is attributed to the bulk Co20Ni20Fe20Al20Ti20 alloy possessing a mostly ductile

262

FCC phase, and low volume fractions of B2-type BCC phase and Al3Ti intermetallics. Therefore,

263

the ultra-high strength and hardness are dominantly attributed to solid-solution strengthening caused

264

by Al and Ti with larger atomic radii than Fe, Ni and Co [3,17], the presence of the strong B2-type

265

BCC phase [3,15], grain-boundary strengthening resulting from the ultra-fine FCC and B2-type

266

BCC grains [34,35], and twin-boundary strengthening originating from nanoscale twins [35], as

267

well as precipitate strengthening due to the presence of a trace amount of nanocrystalline or

268

ultra-fine Al3Ti precipitates [11,34,35].

269

4. Conclusions

270

Bulk Co20Ni20Fe20Al20Ti20 alloy was fabricated by MA and SPS. After MA, a primary BCC

271

solid-solution phase and an FCC solid-solution phase were formed. In contrast, a primary FCC

272

solid-solution phase with a BCC solid-solution phase and a trace amount of Al3Ti intermetallic

273

phase was formed after SPS. TEM results of the bulk Co20Ni20Fe20Al20Ti20 alloy confirmed the FCC,

274

BCC (B2-type) and Al3Ti phases. The FCC and B2-type BCC phases were ultra-fine in size,

275

whereas Al3Ti intermetallic phase was nanocrystalline or ultra-fine in size. The compressive

276

strength, strain-to-failure and hardness of the bulk Co20Ni20Fe20Al20Ti20 alloy are ~2988 MPa, 5.8%

277

and 704 Hv, respectively. Ultra-high strength and hardness are attributed to the presence of strong

278

B2-type BCC phase, solid-solution strengthening, grain-boundary strengthening, twin-boundary

279

strengthening, and precipitate strengthening due to the presence of a trace amount of Al3Ti

280

intermetallic phase. 13

281

Acknowledgements

282

The authors acknowledge the financial support from Fundamental Research Funds for the

283

Central Universities, SCUT (2013ZZ014), from Specialized Research Fund for the Doctoral

284

Program of Higher Education (20130172120027), from the China Scholarship Council (CSC), and

285

from the US Army Research Office (W911NF-14-1-0627). H.M. Wen utilized his private time to

286

perform related work.

287

14

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Figure Captions:

342

Fig. 1. (a) XRD patterns of mechanically milled powders with different milling times, (b) XRD

343

pattern of a slow scan in the range of 2θ=40-48° for the 49 h milled powders, and the fitting

344

result showing a main BCC phase and an FCC phase.

345

Fig. 2. (a) XRD pattern of the bulk alloy after SPS, with the pattern of the 49 h milled powders for

346

comparison, (b) DSC curve of the 49 h milled powders from room temperature to 1323 K, with

347

a heating rate of 10 K min–1.

348

Fig. 3. (a) Bright-field TEM image of the bulk alloy, (b) SAED pattern of grain A corresponding to

349

that of an FCC structure along [011] zone axis, (c) SAED pattern of grain B corresponding to

350

that of a B2-type BCC structure along [001] zone axis. The arrows indicate Al3Ti

351

intermetallics.

352

Fig. 4. (a) Bright-field TEM image showing nanoscale twins in FCC grain C, (b) SAED pattern of _ _

353 354

grain C along [011]M (matrix) and [011]T (twin) zone axes. Fig. 5. Engineering stress-strain curve of consolidated alloy under compression at room

355

temperature.

356

Table Captions:

357

Table 1. Chemical composition (in at.%) analysis results of the detailed phases by EDS/TEM.

358

Table 2. The values of parameters, ΔHmix, ΔSmix, Tm, δ, Ω, e/a, VEC for the alloys investigated in the

359

present study. 17

360

Table 3. Enthalpy of mixing ( H ijmix , kJ mol–1) in binary equiatomic alloys calculated by Miedema’s approach [27,29].

361

362

Table 4. Mechanical properties of the SPS’ed Co20Ni20Fe20Al20Ti20 alloy and some typical SPS’ed HEAs at room temperature.

363 364

(a)

(b)

365 366

Fig. 1. (a) XRD patterns of mechanically milled powders with different milling times, (b) XRD

367

pattern of a slow scan in the range of 2θ=40-48° for the 49 h milled powders, and the fitting result

368

showing a main BCC phase and an FCC phase.

369 370 371

18

(a)

(b)

SPS

49h MA 372 373

Fig. 2. (a) XRD pattern of the bulk alloy after SPS, with the pattern of the 49 h milled powders for

374

comparison, (b) DSC curve of the 49 h milled powders from room temperature to 1323 K, with a

375

heating rate of 10 K min–1.

376 377

A B

378 379

Fig. 3. (a) Bright-field TEM image of the bulk alloy, (b) SAED pattern of grain A corresponding to

380

that of an FCC structure along [011] zone axis, (c) SAED pattern of grain B corresponding to that of

381

a B2-type BCC structure along [001] zone axis. The arrows indicate Al3Ti intermetallics.

382 383 19

C

384 385

Fig. 4. (a) Bright-field TEM image showing nanoscale twins in FCC grain C, (b) SAED pattern of

386

grain C along [011]M (matrix) and [011]T (twin) zone axes.

_ _

387

388 389

Fig. 5. Engineering stress-strain curve of consolidated alloy under compression at room temperature.

390 391 392

393

394

Table 1. Chemical composition (in at.%) analysis results of the detailed phases by EDS/TEM. Phases

Al

Ti

Fe

Co

Ni

Nominal

20

20

20

20

20

20

composition FCC BCC Particles

5.39±0.61 19.98±0.27 71.03±1.38

8.92±0.73 17.37±2.56 23.13±1.27

31.69±0.90 12.31±3.22 2.30±0.72

26.85±0.34 25.71±0.42 2.24±0.66

27.15±1.41 24.63±0.62 1.30±0.57

395

396

397

Table 2. The values of parameters, ΔHmix, ΔSmix, Tm, δ, Ω, e/a, VEC for the alloys investigated in the

398

present study.

Alloys

ΔHmix (kJ mol-1)

ΔSmix (JK-1mol-1)

Tm (K)

δ (%)

Ω

e/a

VEC

γ

Al20Ti20Co20Ni20Fe20

-26.40

13.38

1634.95

6.85

0.83

2.2

6.8

1.018

399

400

401

Table 3. Enthalpy of mixing ( H ijmix , kJ mol–1) in binary equiatomic alloys calculated by Miedema’s

402

approach [27,29]. Element (atomic radii, Å)

Al

Fe

Ni

Co

Ti

Al (1.43) Fe (1.27) Ni (1.25) Co(1.26) Ti (1.46)

— -11 -22 -19 -30

-11 — -2 -1 -17

-22 -2 — 0 -35

-19 -1 0 — -28

-30 -17 -35 -28 —

403 404

405

Table 4. Mechanical properties of the SPS’ed Co20Ni20Fe20Al20Ti20 alloy and some typical SPS’ed

406

HEAs at room temperature. Alloys

σmax

εf 21

Hardness

Refs.

Co20Ni20Fe20Al20Ti20 Co0.5NiFeCrTi0.5 CoNiFeCrAl0.6Ti0.4 AlFeNiCoCrMn Co0.3NiFeCrAl0.7 Al0.5CrFeNiCo0.3C0.2 CoCrFeNiAl FeNiCoCrMn 407 408



(MPa)

(%)

(Hv)

2988 2690 2520 2142 2635 2131 1907 1987

5.8

704 846 573 662 624 617 625 646



~2.0 ~7.0 ~6.0 8.1 3.0 ~6.5 ~8.0

This work [7] [17] [33] [36] [37] [38] [39]

Indicates that values of εf for these alloys were re-estimated according to their engineering stress-strain curves

under compression at room temperature.

409

410

22