- Email: [email protected]
Effects of cyclic thermo-hydrogen processing on microstructural and mechanical properties of Ti6Al4V alloy at room temperature
Journal Pre-proof Effects of cyclic thermo-hydrogen processing on microstructural and mechanical properties of Ti6Al4V alloy at room temperature Baoguo Yuan, Xiaoxue Zhang, Yujie Wang, Qiang Chen, Yuanyuan Wan, Yubin Zheng, Zhihui Xing, Hong Zhan PII:
S0042-207X(19)32269-9
DOI:
https://doi.org/10.1016/j.vacuum.2019.109015
Reference:
VAC 109015
To appear in:
Vacuum
Received Date: 8 August 2019 Revised Date:
13 October 2019
Accepted Date: 14 October 2019
Please cite this article as: Yuan B, Zhang X, Wang Y, Chen Q, Wan Y, Zheng Y, Xing Z, Zhan H, Effects of cyclic thermo-hydrogen processing on microstructural and mechanical properties of Ti6Al4V alloy at room temperature, Vacuum, https://doi.org/10.1016/j.vacuum.2019.109015. 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 Elsevier Ltd. All rights reserved.
1
Effects of cyclic thermo-hydrogen processing on microstructural
2
and mechanical properties of Ti6Al4V alloy at room temperature
3
Baoguo Yuana, Xiaoxue Zhangb, Yujie Wanga, Qiang Chena,c,*, Yuanyuan Wanc, Yubin Zhenga, Zhihui Xingc,
4
Hong Zhanc
5
a
6
b
7
c
8
* Corresponding author. E-mail address:[email protected], [email protected] (Q. Chen)
School of Materials Science and Engineering, Hefei University of Technology, Hefei 230009, PR China School of Mechanical Engineering, Anhui Sanlian University, Hefei 230601, PR China
Southwest Technology and Engineering Research Institute, Chongqing 400039, PR China
9 10
Abstract: Compression tests were performed at room temperature to investigate the effects of cyclic
11
thermo-hydrogen processing (CTHP) on microstructural and mechanical properties of Ti6Al4V alloy. It was
12
found that the amount of β-phase in the microstructure of the alloy increased significantly under CTHP.
13
Moreover, the ultimate compression of the alloy increased first and then started to decrease with the increasing
14
number of CTHP cycles. The ultimate compression of the CTHP 2-treated Ti6Al4V alloy was improved by
15
22% as compared to Ti6Al4V alloy without hydrogen and reached 42%. Furthermore, with the increasing
16
number of CTHP cycles, yield strength and yield ratio decreased first and then increased slightly; however, no
17
conspicuous change in compressive strength was noticed.
18
Keywords: Titanium alloy; Cyclic thermo-hydrogen processing; Plasticity; Microstructure
19
1. Introduction
20
Titanium (Ti) and its alloys, due to their low density, high specific strength, good corrosion resistance, and
21
biochemical compatibility, are widely used in different engineering fields including chemical, aviation,
22
aerospace, energy, marine, and biomedical [1-3]. Ti alloy has become an important metal in many industries
23
along with magnesium alloy, steel, aluminum alloy and composites [4-6]. However, their poor plasticity at
24
room temperature, high yield ratios and high deformation resistance hinder the cold forming of α+β-type Ti
25
alloys [7-9]. 1
26
Fortunately, the incorporation of hydrogen into Ti alloys has been reported to be an effective approach to
27
improve the plasticity of Ti alloys at room temperature [10-13], which can be utilized to achieve cold forming of
28
Ti alloys. Sun et al. [11] investigated the effect of hydrogen on the cold deformation behaviors of Ti6Al4V
29
alloy. They found the plasticity of Ti6Al4V alloy increased with the increase of hydrogen content, and
30
improved by one time compared with that of the unhydrogenated alloy when the Ti6Al4V alloy contained 0.9
31
wt.% H. Ma et al. [12] studied the effect of hydrogen content on room-temperature compressive properties of
32
Ti44Al6Nb alloy. They found fracture strain increased first and then decreased when hydrogen content
33
increased. The Ti44Al6Nb alloy exhibited higher plasticity after addition of 0.42 at.% H, which was enhanced
34
by 31%. In addition, hydrogen can also increase ductility and reduce flow stress at high temperatures [14-17],
35
and can improve superplastic behavior [18-22] and diffusion bonding [23-26] of Ti alloys . The process of
36
hydrogen addition (as a temporary alloying agent) into Ti alloys in order to enhance their microstructural and
37
mechanical properties is called thermo-hydrogen processing (THP) [27-30]. Vacuum annealing of THP-treated
38
Ti alloys is further executed in order to resist hydrogen embrittlement. However, in the traditional THP, the
39
improvements in plasticity of Ti alloys (at room temperature) are not high enough to meet the demands of actual
40
production.
41
In the present study, cyclic thermo-hydrogen processing (CTHP) was employed to improve the plasticity of
42
Ti alloys at room temperature. The effects of CTHP on microstructural and mechanical properties of Ti6Al4V
43
alloy were investigated in detail. Yuan et al. [31] propounded that hydrogenated Ti6Al4V alloy manifested
44
excellent plasticity at room temperature for hydrogen contents of 0.6~0.8 wt%. Hence, in the current work, the
45
hydrogenation temperature was set to at 750°C, which is below the eutectoid transformation temperature of
46
Ti6Al4V alloy, and the hydrogen content was considered as ~0.65 wt%.
47
2. Materials and Methods
48
Ti6Al4V alloy (consisting of α-phase and β-phase) was the main material in the present experiment.
49
Chemical composition of the alloy in wt.% was as follows: 6.22 Al, 4.20 V, 0.15 Fe, 0.01 C, 0.01 N and
50
balance Ti. The CTHP-treated cylindrical specimens (diameter of 4 mm and height of 6 mm) were divided into
51
three groups: CTHP 1 (hydrogenation → quenching), CTHP 2 (hydrogenation → dehydrogenation →
52
hydrogenation → quenching), and CTHP 3 (hydrogenation → dehydrogenation → hydrogenation →
53
dehydrogenation → hydrogenation → quenching). The specimens were first hydrogenated in a tube furnace at 2
54
750°C for 2 h under a hydrogen pressure of 0.0163 MPa, and then air-cooled to room temperature. The furnace
55
tube chamber with specimens was vacuumed before introducing hydrogen. The hydrogenated specimens were
56
further dehydrogenated by vacuum annealing at 750°C for 10 h followed by furnace cooling to room
57
temperature. The hydrogenated specimens were then vacuum-sealed in a quartz tube, then heated in a resistance
58
furnace at 850°C for 30 min, and immediately quenched into water. In contrast, the original specimen without
59
hydrogen was first placed in a tube furnace at 750°C for 2 h under vacuum and air-cooled to room temperature,
60
subsequently, heated in a resistance furnace at 850°C for 30 min after vacuum-sealed in a quartz tube, and
61
finally quenched into water. This process was termed as CTHP 0. The Ti6Al4V alloy samples before and after
62
CTHP were shown in Fig.1. It can be seen that the Ti6Al4V alloy samples were not oxidized during the process
63
of CTHP.
64 65
Fig. 1. Photos of Ti6Al4V alloy samples before (a) and after (b) CTHP
66
Uniaxial compression tests were performed in an axial/torsional test system (MTS 809) at room temperature
67
at a constant compressive velocity of 0.5 mm/min. White vaseline lubricant was coated on both end faces of the
68
specimens in order to reduce friction between specimens and compressive bars. The displacement-load data
69
were recorded by a computer connected to the MTS 809 system. Microstructural characteristics of the
70
specimens were observed by an optical microscope (OM; Olympus BHM-2UM). OM specimens were first cut
71
axially on an electric discharging machine, grounded with emery papers, polished on a polishing machine, and
72
finally, etched in a mixed solution consisting of 1 ml hydrofluoric acid, 1 ml nitric acid and 8 ml water. Grain
73
sizes of specimens were determined from OM images using Image-pro Plus software. X-ray diffraction (XRD)
74
was carried out on a Panalytical X-Pert PRO X-ray diffractometer equipped with a Cu-Kα radiation source (λ =
75
1.5406 Å) under 40 kV and 40 mA (with a scanning parameter of 3 °/min). Thin foils for transmission electron
76
microscopy (TEM; JEOL JEM-2100F) were electropolished by a twin-jet electropolisher (Struers Tenupol-5) 3
77
in a solution bath of 60% CH3OH + 34% C4H9OH + 6% HClO4 (volume ratio) at -25°C~-30°C under 20 V and
78
30~40 mA. Compression fracture morphologies were investigated by a scanning electron microscope (SEM;
79
JSM-6490LV) at 20 kV.
80
3. Results and Discussion
81
3.1 Microstructures
82
Fig. 2 displays the OM micrographs of Ti6Al4V alloys treated by different CTHP procedures. It is clear from
83
Fig. 2a that Ti6Al4V alloy without hydrogen was composed of α-phase and β-phase, and significant changes in
84
its microstructure were observed under CTHP. Both tiny acicular α' martensite phase with a hexagonal
85
closed-packed (HCP) structure and wider acicular α'' martensite phase with an orthorhombic structure appeared
86
in the microstructures of the CTHP-treated alloys (as identified by TEM provided subsequently), which was
87
because that the quenching temperature was higher than the transformation temperature of β-phase in the
88
CTHP-treated alloys, and consequently, the transformations of β → α' and β → α'' occurred, which has been
89
reported previously by Qazi et al. [32]. In addition, the amount of α'' martensite of the CTHP 2-treated alloy
90
increased obviously as compared to the CTHP 1-treated alloy, and increased slightly as compared to the CTHP
91
3-treated alloy, which may be caused by the increasing of stability of β-phase. Moreover, the grain size of the
92
CTHP 2-treated alloy decreased by 0.8% as compared to the CTHP 1-treated alloy, while the grain size of the
93
CTHP 3-treated alloy increased by 25.0% as compared to the CTHP 2-treated alloy, as shown in Fig. 3, and the
94
CTHP 2-treated alloy possessed a more uniform microstructure. The reason may be that orthorhombic α''
95
martensite needles were thicker and their relative orientations were different from HCP α' martensite needles
96
due to different crystal structures, thus resulting in a more homogenous microstructure [32]. As recrystallization
97
occurred during dehydrogenation, the addition of hydrogen increased the nucleation rate of recrystallization and
98
facilitated the grain refinement process. During secondary recrystallization, when Ti6Al4V alloy was
99
dehydrogenated for the second time, anomalous grain growth occurred; hence, the CTHP 3-treated alloy had a
100
coarser microstructure as compared to the CTHP 2-treated alloy.
4
101
102 103
Fig. 2. OM micrographs of Ti6Al4V alloys treated by different CTHP procedures: (a) CTHP 0, (b) CTHP 1, (c)
104
CTHP 2, (d) CTHP 3
105 106
Fig. 3. Grain size of Ti6Al4V alloys treated by different CTHP procedures
107
Fig. 4 presents the XRD patterns of Ti6Al4V alloys treated by different CTHP procedures. It was found that
108
the microstructure of Ti6Al4V alloy without hydrogen consisted of a large amount of α-phase and a small 5
109
quantity of β-phase. However, after CTHP, the relative intensities of β-phase diffraction peaks increased
110
significantly, which reveals that the phase transformation of α → β happened when the alloy was hydrogenated
111
at 750°C (hydrogen is a stabilizing element of β-phase). Moreover, the addition of hydrogen decreased the
112
transformation temperature of β-phase and increased the amount of β-phase after quenching. In addition,
113
face-centered cubic (FCC) δ-hydrides appeared in the CTHP-treated alloys after hydrogenation, which is
114
confirmed by TEM analysis subsequently. Furthermore, due to the lattice expansion of β-phase in the solution
115
of hydrogen atoms, diffraction peaks of β-phase moved to lower angles, whereas no conspicuous shifts in
116
diffraction peaks of α-phase were detected because of the low solubility of hydrogen in α-phase. Moreover, both
117
α' martensite phase and α'' martensite phase appeared in the microstructures of the CTHP-treated alloys, and the
118
relative intensity of α'' martensite phase peak was found to be the highest for the CTHP 2-treated alloy due to the
119
phase transformations of β → α'' and β → α' during quenching. As the lattice constants of the CTHP-treated
120
alloys increased with the addition of hydrogen, their diffraction peaks became wider. In addition, diffraction
121
peaks of α' and α'' martensite phases were overlapped with that of α-phase. With the increasing number of
122
CTHP cycles, the relative intensity of β-phase increased first and then decreased. The relative intensity of
123
β-phase was found to be the highest in the CTHP 2-treated alloy, which signifies that the CTHP 2-treated alloy
124
yielded the highest amount of β-phase.
125
6
126 127
Fig. 4. XRD patterns of Ti6Al4V alloys treated by different CTHP procedures: (a) CTHP 0, (b) CTHP 1, (c)
128
CTHP 2, (d) CTHP 3
129
Fig. 5 displays the TEM images of Ti6Al4V alloys treated by different CTHP procedures. After CTHP, HCP
130
α' martensite, orthorhombic α'' martensite and FCC δ-hydride phase appeared. Formation of both HCP α'
131
martensite and orthorhombic α'' martensite was typical in hydrogenated Ti6Al4V alloys, which had been
132
reported previously by Qazi et al. [32]. α'' martensite needles were thicker than α' martensite needles, as shown
133
in Fig. 5b and Fig. 2. The simultaneous presence of HCP α' martensite and orthorhombic α'' martensite in the
134
hydrogenated titanium alloys had been explained by Qazi et al. [32] for three reasons: (a) a lower solubility of
135
hydrogen in α' martensite than in α'' martensite; (b) the existence of a minimum critical hydrogen concentration
136
required to form α'' martensite; and (c) high hydrogen diffusivity at temperatures between Ms and Mf, where Ms
137
and Mf are the martensite start and finish temperatures, respectively. In addition, δ-hydrides were lamellar and
138
parallel to each other, and most of them originated at the grain (phase) boundary due to hydrogenation [33].
139
When H2 molecules reacted with Ti6Al4V alloy, physical adsorption occurred on the surface and H atoms were
140
generated. Grain (phase) boundaries possessed large amounts of defects and high energy, thus providing
141
channels for the diffusion of H atoms [33]. Therefore, H atoms diffused preferentially in short distances along
142
grain (phase) boundaries, and hydrogen concentration reached its saturation point in a short time. Hence, due to
143
their high hydrogen concentrations, grain (phase) boundaries satisfied the required compositional fluctuation
144
and energy fluctuation for the nucleation of δ-hydrides. When hydrogen exceeded its saturated solid solubility,
145
it reacted with Ti to generate Ti-H compound [34-38].
7
146
147
148 149
Fig. 5. TEM micrographs of Ti6Al4V alloys treated by different CTHP procedures: (a) δ-hydride laths and its
150
SAED pattern in the CTHP 1-treated Ti6Al4V alloy, (b) α' and α'' matrix structure and its SAED pattern in
151
the CTHP 1-treated Ti6Al4V alloy, (c) α' matrix structure and its SAED pattern in the CTHP 2-treated
152
Ti6Al4V alloy, (d) δ-hydride laths and its SAED pattern in the CTHP 2-treated Ti6Al4V alloy, (e) α' matrix
153
structure and its SAED pattern in the CTHP 3-treated Ti6Al4V alloy, (f) δ-hydride laths and its SAED pattern
154
in the CTHP 3-treated Ti6Al4V alloy.
155
3.2 Mechanical properties
8
156
Compression tests of Ti6Al4V alloys treated under different CTHP conditions were performed at room
157
temperature. The difference between strain corresponding to compressive strength and strain corresponding to
158
yield strength is defined as ultimate compression. As shown in Fig. 6, it is noticeable that ultimate compression
159
increased first and then started to decrease slightly with the increasing number of CTHP cycles. The ultimate
160
compression of the CTHP 2-treated Ti6Al4V alloy was improved by 22% as compared to Ti6Al4V alloy
161
without hydrogen and reached 42% (highest among all CTHP-treated alloys). The CTHP 3-treated alloy
162
manifested a slight decrease in ultimate compression as compared to the CTHP 2-treated alloy. Reasons for the
163
improvement in ultimate compression of CTHP-treated alloys are that the amounts of softer β-phase and α''
164
martensite phase increased under CTHP, and the CTHP 2-treated Ti6Al4V alloy yielded the highest amounts of
165
softer β-phase and α'' martensite phase, and possessed the minimum grain size and a more uniform
166
microstructure among the CTHP-treated alloys. Besides the phase evolution, the slight decrease in ultimate
167
compression of CTHP 3-treated alloy was caused by the increase of grain size.
168 169
Fig. 6. Ultimate compression of Ti6Al4V alloys treated by different CTHP procedures
170
Fig. 7 depicts yield strengths (σ0.2) and compressive strengths (σb) of Ti6Al4V alloys treated by different
171
CTHP procedures. With the increasing number of CTHP cycles, yield strength decreased first and then
172
increased slightly. The CTHP 2-treated alloy manifested the lowest yield strength (decreased by 11% as
173
compared to Ti6Al4V alloy without hydrogen) among all CTHP-treated alloys; however, no significant change
174
in compressive strength was observed, it happened because the CTHP-treated alloys contained more β-phase
175
and α'' martensite phase as compared to Ti6Al4V alloy without hydrogen. The existence of softer β-phase and 9
176
α'' martensite phase can reduce the compressive strength of Ti6Al4V alloy, whereas the presence of hydrogen in
177
both solid solution and δ-hydride can improve the compressive strength of Ti6Al4V alloy. Therefore,
178
compressive strengths of the CTHP-treated Ti6Al4V alloys were affected by softening induced by β-phase and
179
α'' martensite phase as well as by strengthening induced by hydrogen existed in both solid solution and
180
δ-hydride. Moreover, with the increasing number of CTHP cycles, the yield ratio (the ratio of yield strength to
181
compressive strength) decreased first and then increased slightly (Fig. 8). The CTHP 2-treated Ti6Al4V alloy
182
also manifested the lowest yield ratio (decreased by 11% as compared to Ti6Al4V alloy without hydrogen)
183
among all CTHP-treated alloys. However, lower yield ratio is beneficial for Ti alloys to deform plastically at
184
room temperature.
185 186
Fig. 7. Yield strength and compressive strength of Ti6Al4V alloys treated by different CTHP procedures
10
187 188
Fig. 8. Yield ratio of Ti6Al4V alloys treated by different CTHP procedures
189
3.3 Fracture morphologies
190
Macroscopic fracture surfaces of Ti6Al4V alloys (both without hydrogen and treated by CTHP) were found
191
to be aligned at a 45˚ angle with the axes of cylindrical specimens, which indicates that fracture under
192
compression occurred in a shear mode. Fig. 9 exhibits the fracture morphologies of Ti6Al4V alloys treated by
193
different CTHP processes. Two types of fracture characteristics, relatively smooth surface and ductile dimples,
194
were noticed on fracture surfaces of Ti6Al4V alloys (both without hydrogen and treated by CTHP). Therefore,
195
it can be inferred that the dominant fracture mode of the alloys was ductile fracture. With the increasing number
196
of CTHP cycles, areas of ductile dimples increased first and then started to decrease. The CTHP 2-treated
197
Ti6Al4V alloy yielded the largest ductile dimple area and, thus, manifested the best plasticity, which is
198
corresponding to the results of ultimate compression of Ti6Al4V alloys treated under different CTHP processes
199
(Fig. 6).
11
(a)
(b) Smooth surface
Smooth surface
Ductile dimple
Ductile dimple 200
(c)
Ductile dimple
(d)
Smooth surface
Ductile dimple
Smooth surface 201 202
Fig. 9. Fracture surfaces of Ti6Al4V alloys treated by different CTHP procedures: (a) CTHP 0, (b) CTHP 1, (c)
203
CTHP 2, (d) CTHP 3
204
4. Conclusions
205
The ultimate compression of Ti6Al4V alloy increased first and then started to decrease with the increasing
206
number of CTHP cycles. The ultimate compression of the CTHP 2-treated Ti6Al4V alloy was improved by
207
22% as compared to Ti6Al4V alloy without hydrogen and reached 42% (highest among all CTHP-treated
208
alloys). Furthermore, with the increasing number of CTHP cycles, yield strength and yield ratio decreased first
209
and then increased slightly. The CTHP 2-treated Ti6Al4V alloy manifested the lowest yield strength (decreased
210
by 11% as compared to Ti6Al4V alloy without hydrogen) among all CTHP-treated alloys; however, no
211
significant change in compressive strength was noticed. Furthermore, the CTHP 2-treated Ti6Al4V alloy also
212
manifested the lowest yield ratio (decreased by 11% as compared to Ti6Al4V alloy without hydrogen) among
213
all CTHP-treated alloys. In addition, relatively smooth surface and ductile dimples were observed on fracture
214
surfaces of Ti6Al4V alloys (both without hydrogen and treated by CTHP). With the increasing number of
215
CTHP cycles, areas of ductile dimples increased first and then started to decrease, and the CTHP 2-treated 12
216
Ti6Al4V alloy yielded the largest ductile dimple area. Reasons for the improvement in ultimate compression of
217
CTHP-treated alloys are that the amounts of softer β-phase and α'' martensite phase increased under CTHP, and
218
the CTHP 2-treated Ti6Al4V alloy yielded the highest amounts of softer β-phase and α'' martensite phase, and
219
possessed the minimum grain size and a more uniform microstructure among the CTHP-treated alloys. In
220
addition, compressive strengths of the CTHP-treated Ti6Al4V alloys were affected by softening induced by
221
β-phase and α'' martensite phase as well as by strengthening induced by hydrogen existed in both solid solution
222
and δ-hydride.
223
Acknowledgements
224
This work was financially supported by the National Natural Science Foundation of China (No.51875157)
225
and the University Natural Science Research Project of Anhui Province (No.KJ2019A0894).
226
References
227
1.
Y.C. Lin, Y. Tang, X.Y. Zhang, C. Chen, H. Yang, K.C. Zhou, Effects of solution temperature and
228
cooling rate on microstructure and micro-hardness of a hot compressed Ti-6Al-4V alloy, Vacuum
229
159(2019)191-199.
230
2.
Y. Geng, X. Mei, K. Wang, X. Dong, X. Yan, Z. Fan, W. Duan, W. Wang, Effect of Laser Shock
231
Peening on Residual Stress, Microstructure and Hot Corrosion Behavior of Damage-Tolerant TC21
232
Titanium Alloy, J. Mater. Eng. Perform. 27(2018)4703-4713.
233
3.
Y.Q. Jiang, Y.C. Lin, X.Y. Zhang, C. Chen, Q.W. Wang, G.D. Pang, Isothermal tensile deformation
234
behaviors and fracture mechanism of Ti-5Al-5Mo-5V-1Cr-1Fe alloy in beta phase field, Vacuum
235
156 (2018) 187-197.
236
4.
G. Chen, X. Chang, J. Zhang, Y. Jin, C. Sun, Q. Chen, Z. Zhao, Microstructures and Mechanical
237
Properties of In-Situ Al3Ti/2024 Aluminum Matrix Composites Fabricated by Ultrasonic Treatment
238
and
239
https://doi.org/10.1007/s12540-019-00396-y.
240
5.
Subsequent
Squeeze
Casting,
Met.
Mater.
Int.
(2019).
Y. Wang, F. Li, X.W. Li, W.B. Fang, Unusual texture formation and mechanical property in AZ31
241
magnesium alloy sheets processed by CVCDE, J. Mater. Process. Technol. 275 (2020).
242
https://doi.org/10.1016/j.jmatprotec.2019.116360
13
243
6.
Y. Meng, J.Y. Zhang, H. Zhan, Q. Chen, J. Zhou, S. Sugiyama, J. Yanagimoto, Effects of
244
subsequent treatments on the microstructure and mechanical properties of SKD11 tool steel samples
245
processed
246
https://doi.org/10.1016/j.msea.2019.138070
247
7.
248 249
by
multi-stage
thixoforging,
Mater.
Sci.
Eng.,
A
762
(2019).
X. Li, H. Sun, P. Zhang, W. Fang, The effect of strain on dynamic recrystallization of PM Ti–45Al– 10Nb intermetallics during isothermal forging, Intermetallics 55(2014)90-94.
8.
Y. Yu, W. Zhang, W. Dong, J. Yang, Y. Feng, Effects of pre-sintering on microstructure and
250
properties of TiBw/Ti6Al4V composites fabricated by hot extrusion with steel cup, Mater. Sci. Eng.,
251
A 638(2015)38-45.
252
9.
Y.C. Lin, X.Y. Jiang, C.J. Shuai, C.Y. Zhao, D.G. He, M.S. Chen, C. Chen, Effects of initial
253
microstructures on hot tensile deformation behaviors and fracture characteristics of Ti-6Al-4V alloy,
254
Mater. Sci. Eng., A 711(2018)293-302.
255 256
10. B.A. Kolachev, Reversible hydrogen alloying of titanium-alloys, Met. Sci. Heat Treat. 35(1993)586-591.
257
11. Z.G. Sun, G.Q. Chen, Y.Q. Wang, W.L. Zhou, H.L. Hou, Investigations on cold deformation
258
mechanisms of the hydrogenated Ti-6Al-4V alloys, Mater. Sci. Eng., A 527(2010)1003-1007.
259
12. T.F. Ma, R.R. Chen, D.S. Zheng, J.J. Guo, H.S. Ding, Y.Q. Su, H.Z. Fu, Hydrogen enhanced the
260
room-temperature compressive properties of Ti44Al6Nb alloy, Mater. Lett. 213(2018)170-173.
261
13. B.G. Yuan, Y.B. Zheng, L.Q. Gong, Q. Chen, M. Lv, Effect of hydrogen content on microstructures
262 263 264
and room-temperature compressive properties of TC21 alloy, Mater. Des. 94(2016)330-337. 14. Y.Y. Zong, D.B. Shan, M. Xu, Y. Lv, Flow softening and microstructural evolution of TC11 titanium alloy during hot deformation, J. Mater. Process. Technol. 209(2009)1988-1994.
265
15. X. Wang, L. Wang, L. Luo, X. Liu, Y. Tang, X. Li, R. Chen, Y. Su, J. Guo, H. Fu, Hot deformation
266
behavior and dynamic recrystallization of melt hydrogenated Ti-6Al-4V alloy, J. Alloys Compd.
267
728(2017)709-718.
268 269 270 271
16. T. Ma, R. Chen, D. Zheng, J. Guo, H. Ding, Y. Su, H. Fu, Hydrogen-induced softening of Ti-44Al-6Nb-1Cr-2V alloy during hot deformation, Int. J. Hydrogen Energy 42(2017)8329-8337. 17. B. Shao, S.X. Wan, D.B. Shan, B. Guo, Y.Y. Zong, 2017. Hydrogen-Induced Improvement of the Cylindrical Drawing Properties of a Ti-22Al-25Nb Alloy. Adv. Eng. Mater. 19, 201600621.
14
272 273 274 275 276 277
18. M.A. Murzinova, G.A. Salishchev, D.D. Afonichev, Superplasticity of hydrogen-containing VT6 titanium alloy with a submicrocrystalline structure, Phys. Met. Metall. 104(2007)195-202. 19. X. Zhang, Y. Zhao, W. Zeng, Effect of hydrogen on the superplasticity of Ti600 alloy, Int. J. Hydrogen Energy 35(2010)4354-4360. 20. X.L. Wang, Y.Q. Zhao, H.L. Hou, Y.Q. Wang, Effect of hydrogen content on superplastic forming/diffusion bonding of TC21 alloys, J. Alloys Compd. 503(2010)151-154.
278
21. X. Li, G. Jia, F. Qu, H. Wu, J. Chen, Ultrafine Grain Refinement and Superplasticity of Ti-55 Alloy
279
Obtained by Hydrogen Absorption and Desorption, J. Mater. Eng. Perform. 27(2018)3472-3477.
280
22. G.P. Grabovetskaya, E.N. Stepanova, I.V. Ratochka, I.P. Mishin, O.V. Zabudchenko, Effect of
281
Hydrogen on the Development of Superplastic Deformation in the Submicrocrystalline Ti–6Al–4V
282
Alloy, Mater. Sci. Forum 838-839(2016)344-349.
283 284
23. Z. Wang, C. Li, J. Qi, J. Feng, J. Cao, Characterization of hydrogenated niobium interlayer and its application in TiAl/Ti2AlNb diffusion bonding, Int. J. Hydrogen Energy 44(2019)6929-6937.
285
24. L.X. Zhang, Z. Sun, J.M. Shi, X.F. Ye, Z.Y. Yang, J.C. Feng, Low-temperature direct diffusion
286
bonding of hydrogenated TC4 alloy and GH3128 superalloy, Int. J. Hydrogen Energy
287
44(2019)3906-3916.
288 289
25. J.W. Wang, H.R. Gong, Adsorption and diffusion of hydrogen on Ti, Al, and TiAl surfaces, Int. J. Hydrogen Energy 39(2014)6068-6075.
290
26. J.C. Feng, H. Liu, P. He, J. Cao, Effects of hydrogen on diffusion bonding of hydrogenated Ti6Al4V
291
alloy containing 0.3 wt% hydrogen at fast heating rate, Int. J. Hydrogen Energy 32(2007)3054-3058.
292
27. A.A. Ilyin, S.V. Skvortsova, A.M. Mamonov, G.V. Permyakova, D.A. Kurnikov, Effect of
293
thermohydrogen treatment on the structure and properties of titanium alloy castings, Met. Sci. Heat
294
Treat. 44(2002)185-189.
295 296 297 298 299 300
28. F.H. Froes, O.N. Senkov, J.O. Qazi, Hydrogen as a temporary alloying element in titanium alloys: thermohydrogen processing, Int. Mater. Rev. 49(2004)227-245. 29. C.P. Liang, H.R. Gong, Fundamental influence of hydrogen on various properties of alpha-titanium, Int. J. Hydrogen Energy 35(2010)3812-3816. 30. L.C. Campanelli, P.S.C.P. da Silva, A.M. Jorge, C. Bolfarini, Effect of hydrogen on the fatigue behavior of the near-β Ti-5Al-5Mo-5V-3Cr alloy, Scr. Mater. 132(2017)39-43.
15
301 302 303 304 305 306 307 308 309 310 311 312
31. B.G. Yuan, C.F. Li, H.P. Yu, D.L. Sun, Influence of hydrogen content on tensile and compressive properties of Ti-6Al-4V alloy at room temperature, Mater. Sci. Eng., A 527(2010)4185-4190. 32. J.I. Qazi, O.N. Senkov, J. Rahim, F.H. Froes, Kinetics of martensite decomposition in Ti-6Al-4V-xH alloys, Mater. Sci. Eng., A 359(2003)137-149. 33. X. Wang, Y. Zhao, W. Zeng, H. Hou, Y. Wang, Kinetics of Hydrogen Absorption of Ti600 Alloy, Rare Met. Mater. Eng. 38(2009)1219-1222. 34. T.K. Zhu, M.Q. Li, Effect of 0.770 wt%H addition on the microstructure of Ti-6Al-4V alloy and mechanism of [delta] hydride formation, J. Alloys Compd. 481(2009)480-485. 35. C.Y. Yu, C.C. Shen, T.P. Perng, Microstructure of Ti-6Al-4V processed by hydrogenation, Scr. Mater. 55(2006)1023-1026. 36. J. Zhao, H. Ding, Y. Zhong, C.S. Lee, Effect of thermo hydrogen treatment on lattice defects and microstructure refinement of Ti6Al4V alloy, Int. J. Hydrogen Energy 35(2010)6448-6454.
313
37. N. Pushilina, A. Panin, M. Syrtanov, E. Kashkarov, V. Kudiiarov, O. Perevalova, R. Laptev, A. Lider,
314
A. Koptyug, 2018. Hydrogen-Induced Phase Transformation and Microstructure Evolution for
315
Ti-6Al-4V Parts Produced by Electron Beam Melting. Met. 8, 301.
316
38. R. de Araujo-Silva, A.M. Neves, L.E.R. Vega, M.R.M. Triques, D.R. Leiva, C.S. Kiminami, T.T.
317
Ishikawa, A.M. Jorge Junior, W.I. Botta, Synthesis of beta-Ti-Nb alloys from elemental powders by
318
high-energy ball milling and their hydrogenation features, Int. J. Hydrogen Energy
319
43(2018)18382-18391.
320
16
321
Figure captions:
322
Fig. 1. Photos of Ti6Al4V alloy samples before (a) and after (b) CTHP
323
Fig. 2. OM micrographs of Ti6Al4V alloys treated by different CTHP procedures: (a) CTHP 0, (b) CTHP 1, (c)
324
CTHP 2, (d) CTHP 3
325
Fig. 3. Grain size of Ti6Al4V alloys treated by different CTHP procedures
326
Fig. 4. XRD patterns of Ti6Al4V alloys treated by different CTHP procedures: (a) CTHP 0, (b) CTHP 1, (c)
327
CTHP 2, (d) CTHP 3
328
Fig. 5. TEM micrographs of Ti6Al4V alloys treated by different CTHP procedures: (a) δ-hydride laths and its
329
SAED pattern in the CTHP 1-treated Ti6Al4V alloy, (b) α' and α'' matrix structure and its SAED pattern in
330
the CTHP 1-treated Ti6Al4V alloy, (c) α' matrix structure and its SAED pattern in the CTHP 2-treated
331
Ti6Al4V alloy, (d) δ-hydride laths and its SAED pattern in the CTHP 2-treated Ti6Al4V alloy, (e) α' matrix
332
structure and its SAED pattern in the CTHP 3-treated Ti6Al4V alloy, (f) δ-hydride laths and its SAED pattern
333
in the CTHP 3-treated Ti6Al4V alloy.
334
Fig. 6. Ultimate compression of Ti6Al4V alloys treated by different CTHP procedures
335
Fig. 7. Yield strength and compressive strength of Ti6Al4V alloys treated by different CTHP procedures
336
Fig. 8. Yield ratio of Ti6Al4V alloys treated by different CTHP procedures
337
Fig. 9. Fracture surfaces of Ti6Al4V alloys treated by different CTHP procedures: (a) CTHP 0, (b) CTHP 1, (c)
338
CTHP 2, (d) CTHP 3
17
Fig. 1. Photos of Ti6Al4V alloy samples before (a) and after (b) CTHP
Fig. 2. OM micrographs of Ti6Al4V alloys treated by different CTHP procedures: (a) CTHP 0, (b) CTHP 1, (c) CTHP 2, (d) CTHP 3
Fig. 3. Grain size of Ti6Al4V alloys treated by different CTHP procedures
Fig. 4. XRD patterns of Ti6Al4V alloys treated by different CTHP procedures: (a) CTHP 0, (b) CTHP 1, (c) CTHP 2, (d) CTHP 3
Fig. 5. TEM micrographs of Ti6Al4V alloys treated by different CTHP procedures: (a) δ-hydride laths and its SAED pattern in the CTHP 1-treated Ti6Al4V alloy, (b) α' and α'' matrix structure and its SAED pattern in the CTHP 1-treated Ti6Al4V alloy, (c) α' matrix structure and its SAED pattern in the
CTHP 2-treated Ti6Al4V alloy, (d) δ-hydride laths and its SAED pattern in the CTHP 2-treated Ti6Al4V alloy, (e) α' matrix structure and its SAED pattern in the CTHP 3-treated Ti6Al4V alloy, (f) δ-hydride laths and its SAED pattern in the CTHP 3-treated Ti6Al4V alloy.
Fig. 6. Ultimate compression of Ti6Al4V alloys treated by different CTHP procedures
Fig. 7. Yield strength and compressive strength of Ti6Al4V alloys treated by different CTHP procedures
Fig. 8. Yield ratio of Ti6Al4V alloys treated by different CTHP procedures
(a)
Smooth surface
(b)
Ductile dimple
(c)
Ductile dimple
Smooth surface
Smooth surface
Ductile dimple
(d)
Smooth surface
Ductile dimple
Fig. 9. Fracture surfaces of Ti6Al4V alloys treated by different CTHP procedures: (a) CTHP 0, (b) CTHP 1, (c) CTHP 2, (d) CTHP 3
Declarations of interest: none
Highlights Plasticity increased first and then decreased under CTHP. Yield ratio decreased first and then increased slightly under CTHP. The amounts of softer β and α'' phases increased after treatment of CTHP. Area of ductile dimple increased first and then decreased under CTHP.
S0042-207X(19)32269-9
DOI:
https://doi.org/10.1016/j.vacuum.2019.109015
Reference:
VAC 109015
To appear in:
Vacuum
Received Date: 8 August 2019 Revised Date:
13 October 2019
Accepted Date: 14 October 2019
Please cite this article as: Yuan B, Zhang X, Wang Y, Chen Q, Wan Y, Zheng Y, Xing Z, Zhan H, Effects of cyclic thermo-hydrogen processing on microstructural and mechanical properties of Ti6Al4V alloy at room temperature, Vacuum, https://doi.org/10.1016/j.vacuum.2019.109015. 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 Elsevier Ltd. All rights reserved.
1
Effects of cyclic thermo-hydrogen processing on microstructural
2
and mechanical properties of Ti6Al4V alloy at room temperature
3
Baoguo Yuana, Xiaoxue Zhangb, Yujie Wanga, Qiang Chena,c,*, Yuanyuan Wanc, Yubin Zhenga, Zhihui Xingc,
4
Hong Zhanc
5
a
6
b
7
c
8
* Corresponding author. E-mail address:[email protected], [email protected] (Q. Chen)
School of Materials Science and Engineering, Hefei University of Technology, Hefei 230009, PR China School of Mechanical Engineering, Anhui Sanlian University, Hefei 230601, PR China
Southwest Technology and Engineering Research Institute, Chongqing 400039, PR China
9 10
Abstract: Compression tests were performed at room temperature to investigate the effects of cyclic
11
thermo-hydrogen processing (CTHP) on microstructural and mechanical properties of Ti6Al4V alloy. It was
12
found that the amount of β-phase in the microstructure of the alloy increased significantly under CTHP.
13
Moreover, the ultimate compression of the alloy increased first and then started to decrease with the increasing
14
number of CTHP cycles. The ultimate compression of the CTHP 2-treated Ti6Al4V alloy was improved by
15
22% as compared to Ti6Al4V alloy without hydrogen and reached 42%. Furthermore, with the increasing
16
number of CTHP cycles, yield strength and yield ratio decreased first and then increased slightly; however, no
17
conspicuous change in compressive strength was noticed.
18
Keywords: Titanium alloy; Cyclic thermo-hydrogen processing; Plasticity; Microstructure
19
1. Introduction
20
Titanium (Ti) and its alloys, due to their low density, high specific strength, good corrosion resistance, and
21
biochemical compatibility, are widely used in different engineering fields including chemical, aviation,
22
aerospace, energy, marine, and biomedical [1-3]. Ti alloy has become an important metal in many industries
23
along with magnesium alloy, steel, aluminum alloy and composites [4-6]. However, their poor plasticity at
24
room temperature, high yield ratios and high deformation resistance hinder the cold forming of α+β-type Ti
25
alloys [7-9]. 1
26
Fortunately, the incorporation of hydrogen into Ti alloys has been reported to be an effective approach to
27
improve the plasticity of Ti alloys at room temperature [10-13], which can be utilized to achieve cold forming of
28
Ti alloys. Sun et al. [11] investigated the effect of hydrogen on the cold deformation behaviors of Ti6Al4V
29
alloy. They found the plasticity of Ti6Al4V alloy increased with the increase of hydrogen content, and
30
improved by one time compared with that of the unhydrogenated alloy when the Ti6Al4V alloy contained 0.9
31
wt.% H. Ma et al. [12] studied the effect of hydrogen content on room-temperature compressive properties of
32
Ti44Al6Nb alloy. They found fracture strain increased first and then decreased when hydrogen content
33
increased. The Ti44Al6Nb alloy exhibited higher plasticity after addition of 0.42 at.% H, which was enhanced
34
by 31%. In addition, hydrogen can also increase ductility and reduce flow stress at high temperatures [14-17],
35
and can improve superplastic behavior [18-22] and diffusion bonding [23-26] of Ti alloys . The process of
36
hydrogen addition (as a temporary alloying agent) into Ti alloys in order to enhance their microstructural and
37
mechanical properties is called thermo-hydrogen processing (THP) [27-30]. Vacuum annealing of THP-treated
38
Ti alloys is further executed in order to resist hydrogen embrittlement. However, in the traditional THP, the
39
improvements in plasticity of Ti alloys (at room temperature) are not high enough to meet the demands of actual
40
production.
41
In the present study, cyclic thermo-hydrogen processing (CTHP) was employed to improve the plasticity of
42
Ti alloys at room temperature. The effects of CTHP on microstructural and mechanical properties of Ti6Al4V
43
alloy were investigated in detail. Yuan et al. [31] propounded that hydrogenated Ti6Al4V alloy manifested
44
excellent plasticity at room temperature for hydrogen contents of 0.6~0.8 wt%. Hence, in the current work, the
45
hydrogenation temperature was set to at 750°C, which is below the eutectoid transformation temperature of
46
Ti6Al4V alloy, and the hydrogen content was considered as ~0.65 wt%.
47
2. Materials and Methods
48
Ti6Al4V alloy (consisting of α-phase and β-phase) was the main material in the present experiment.
49
Chemical composition of the alloy in wt.% was as follows: 6.22 Al, 4.20 V, 0.15 Fe, 0.01 C, 0.01 N and
50
balance Ti. The CTHP-treated cylindrical specimens (diameter of 4 mm and height of 6 mm) were divided into
51
three groups: CTHP 1 (hydrogenation → quenching), CTHP 2 (hydrogenation → dehydrogenation →
52
hydrogenation → quenching), and CTHP 3 (hydrogenation → dehydrogenation → hydrogenation →
53
dehydrogenation → hydrogenation → quenching). The specimens were first hydrogenated in a tube furnace at 2
54
750°C for 2 h under a hydrogen pressure of 0.0163 MPa, and then air-cooled to room temperature. The furnace
55
tube chamber with specimens was vacuumed before introducing hydrogen. The hydrogenated specimens were
56
further dehydrogenated by vacuum annealing at 750°C for 10 h followed by furnace cooling to room
57
temperature. The hydrogenated specimens were then vacuum-sealed in a quartz tube, then heated in a resistance
58
furnace at 850°C for 30 min, and immediately quenched into water. In contrast, the original specimen without
59
hydrogen was first placed in a tube furnace at 750°C for 2 h under vacuum and air-cooled to room temperature,
60
subsequently, heated in a resistance furnace at 850°C for 30 min after vacuum-sealed in a quartz tube, and
61
finally quenched into water. This process was termed as CTHP 0. The Ti6Al4V alloy samples before and after
62
CTHP were shown in Fig.1. It can be seen that the Ti6Al4V alloy samples were not oxidized during the process
63
of CTHP.
64 65
Fig. 1. Photos of Ti6Al4V alloy samples before (a) and after (b) CTHP
66
Uniaxial compression tests were performed in an axial/torsional test system (MTS 809) at room temperature
67
at a constant compressive velocity of 0.5 mm/min. White vaseline lubricant was coated on both end faces of the
68
specimens in order to reduce friction between specimens and compressive bars. The displacement-load data
69
were recorded by a computer connected to the MTS 809 system. Microstructural characteristics of the
70
specimens were observed by an optical microscope (OM; Olympus BHM-2UM). OM specimens were first cut
71
axially on an electric discharging machine, grounded with emery papers, polished on a polishing machine, and
72
finally, etched in a mixed solution consisting of 1 ml hydrofluoric acid, 1 ml nitric acid and 8 ml water. Grain
73
sizes of specimens were determined from OM images using Image-pro Plus software. X-ray diffraction (XRD)
74
was carried out on a Panalytical X-Pert PRO X-ray diffractometer equipped with a Cu-Kα radiation source (λ =
75
1.5406 Å) under 40 kV and 40 mA (with a scanning parameter of 3 °/min). Thin foils for transmission electron
76
microscopy (TEM; JEOL JEM-2100F) were electropolished by a twin-jet electropolisher (Struers Tenupol-5) 3
77
in a solution bath of 60% CH3OH + 34% C4H9OH + 6% HClO4 (volume ratio) at -25°C~-30°C under 20 V and
78
30~40 mA. Compression fracture morphologies were investigated by a scanning electron microscope (SEM;
79
JSM-6490LV) at 20 kV.
80
3. Results and Discussion
81
3.1 Microstructures
82
Fig. 2 displays the OM micrographs of Ti6Al4V alloys treated by different CTHP procedures. It is clear from
83
Fig. 2a that Ti6Al4V alloy without hydrogen was composed of α-phase and β-phase, and significant changes in
84
its microstructure were observed under CTHP. Both tiny acicular α' martensite phase with a hexagonal
85
closed-packed (HCP) structure and wider acicular α'' martensite phase with an orthorhombic structure appeared
86
in the microstructures of the CTHP-treated alloys (as identified by TEM provided subsequently), which was
87
because that the quenching temperature was higher than the transformation temperature of β-phase in the
88
CTHP-treated alloys, and consequently, the transformations of β → α' and β → α'' occurred, which has been
89
reported previously by Qazi et al. [32]. In addition, the amount of α'' martensite of the CTHP 2-treated alloy
90
increased obviously as compared to the CTHP 1-treated alloy, and increased slightly as compared to the CTHP
91
3-treated alloy, which may be caused by the increasing of stability of β-phase. Moreover, the grain size of the
92
CTHP 2-treated alloy decreased by 0.8% as compared to the CTHP 1-treated alloy, while the grain size of the
93
CTHP 3-treated alloy increased by 25.0% as compared to the CTHP 2-treated alloy, as shown in Fig. 3, and the
94
CTHP 2-treated alloy possessed a more uniform microstructure. The reason may be that orthorhombic α''
95
martensite needles were thicker and their relative orientations were different from HCP α' martensite needles
96
due to different crystal structures, thus resulting in a more homogenous microstructure [32]. As recrystallization
97
occurred during dehydrogenation, the addition of hydrogen increased the nucleation rate of recrystallization and
98
facilitated the grain refinement process. During secondary recrystallization, when Ti6Al4V alloy was
99
dehydrogenated for the second time, anomalous grain growth occurred; hence, the CTHP 3-treated alloy had a
100
coarser microstructure as compared to the CTHP 2-treated alloy.
4
101
102 103
Fig. 2. OM micrographs of Ti6Al4V alloys treated by different CTHP procedures: (a) CTHP 0, (b) CTHP 1, (c)
104
CTHP 2, (d) CTHP 3
105 106
Fig. 3. Grain size of Ti6Al4V alloys treated by different CTHP procedures
107
Fig. 4 presents the XRD patterns of Ti6Al4V alloys treated by different CTHP procedures. It was found that
108
the microstructure of Ti6Al4V alloy without hydrogen consisted of a large amount of α-phase and a small 5
109
quantity of β-phase. However, after CTHP, the relative intensities of β-phase diffraction peaks increased
110
significantly, which reveals that the phase transformation of α → β happened when the alloy was hydrogenated
111
at 750°C (hydrogen is a stabilizing element of β-phase). Moreover, the addition of hydrogen decreased the
112
transformation temperature of β-phase and increased the amount of β-phase after quenching. In addition,
113
face-centered cubic (FCC) δ-hydrides appeared in the CTHP-treated alloys after hydrogenation, which is
114
confirmed by TEM analysis subsequently. Furthermore, due to the lattice expansion of β-phase in the solution
115
of hydrogen atoms, diffraction peaks of β-phase moved to lower angles, whereas no conspicuous shifts in
116
diffraction peaks of α-phase were detected because of the low solubility of hydrogen in α-phase. Moreover, both
117
α' martensite phase and α'' martensite phase appeared in the microstructures of the CTHP-treated alloys, and the
118
relative intensity of α'' martensite phase peak was found to be the highest for the CTHP 2-treated alloy due to the
119
phase transformations of β → α'' and β → α' during quenching. As the lattice constants of the CTHP-treated
120
alloys increased with the addition of hydrogen, their diffraction peaks became wider. In addition, diffraction
121
peaks of α' and α'' martensite phases were overlapped with that of α-phase. With the increasing number of
122
CTHP cycles, the relative intensity of β-phase increased first and then decreased. The relative intensity of
123
β-phase was found to be the highest in the CTHP 2-treated alloy, which signifies that the CTHP 2-treated alloy
124
yielded the highest amount of β-phase.
125
6
126 127
Fig. 4. XRD patterns of Ti6Al4V alloys treated by different CTHP procedures: (a) CTHP 0, (b) CTHP 1, (c)
128
CTHP 2, (d) CTHP 3
129
Fig. 5 displays the TEM images of Ti6Al4V alloys treated by different CTHP procedures. After CTHP, HCP
130
α' martensite, orthorhombic α'' martensite and FCC δ-hydride phase appeared. Formation of both HCP α'
131
martensite and orthorhombic α'' martensite was typical in hydrogenated Ti6Al4V alloys, which had been
132
reported previously by Qazi et al. [32]. α'' martensite needles were thicker than α' martensite needles, as shown
133
in Fig. 5b and Fig. 2. The simultaneous presence of HCP α' martensite and orthorhombic α'' martensite in the
134
hydrogenated titanium alloys had been explained by Qazi et al. [32] for three reasons: (a) a lower solubility of
135
hydrogen in α' martensite than in α'' martensite; (b) the existence of a minimum critical hydrogen concentration
136
required to form α'' martensite; and (c) high hydrogen diffusivity at temperatures between Ms and Mf, where Ms
137
and Mf are the martensite start and finish temperatures, respectively. In addition, δ-hydrides were lamellar and
138
parallel to each other, and most of them originated at the grain (phase) boundary due to hydrogenation [33].
139
When H2 molecules reacted with Ti6Al4V alloy, physical adsorption occurred on the surface and H atoms were
140
generated. Grain (phase) boundaries possessed large amounts of defects and high energy, thus providing
141
channels for the diffusion of H atoms [33]. Therefore, H atoms diffused preferentially in short distances along
142
grain (phase) boundaries, and hydrogen concentration reached its saturation point in a short time. Hence, due to
143
their high hydrogen concentrations, grain (phase) boundaries satisfied the required compositional fluctuation
144
and energy fluctuation for the nucleation of δ-hydrides. When hydrogen exceeded its saturated solid solubility,
145
it reacted with Ti to generate Ti-H compound [34-38].
7
146
147
148 149
Fig. 5. TEM micrographs of Ti6Al4V alloys treated by different CTHP procedures: (a) δ-hydride laths and its
150
SAED pattern in the CTHP 1-treated Ti6Al4V alloy, (b) α' and α'' matrix structure and its SAED pattern in
151
the CTHP 1-treated Ti6Al4V alloy, (c) α' matrix structure and its SAED pattern in the CTHP 2-treated
152
Ti6Al4V alloy, (d) δ-hydride laths and its SAED pattern in the CTHP 2-treated Ti6Al4V alloy, (e) α' matrix
153
structure and its SAED pattern in the CTHP 3-treated Ti6Al4V alloy, (f) δ-hydride laths and its SAED pattern
154
in the CTHP 3-treated Ti6Al4V alloy.
155
3.2 Mechanical properties
8
156
Compression tests of Ti6Al4V alloys treated under different CTHP conditions were performed at room
157
temperature. The difference between strain corresponding to compressive strength and strain corresponding to
158
yield strength is defined as ultimate compression. As shown in Fig. 6, it is noticeable that ultimate compression
159
increased first and then started to decrease slightly with the increasing number of CTHP cycles. The ultimate
160
compression of the CTHP 2-treated Ti6Al4V alloy was improved by 22% as compared to Ti6Al4V alloy
161
without hydrogen and reached 42% (highest among all CTHP-treated alloys). The CTHP 3-treated alloy
162
manifested a slight decrease in ultimate compression as compared to the CTHP 2-treated alloy. Reasons for the
163
improvement in ultimate compression of CTHP-treated alloys are that the amounts of softer β-phase and α''
164
martensite phase increased under CTHP, and the CTHP 2-treated Ti6Al4V alloy yielded the highest amounts of
165
softer β-phase and α'' martensite phase, and possessed the minimum grain size and a more uniform
166
microstructure among the CTHP-treated alloys. Besides the phase evolution, the slight decrease in ultimate
167
compression of CTHP 3-treated alloy was caused by the increase of grain size.
168 169
Fig. 6. Ultimate compression of Ti6Al4V alloys treated by different CTHP procedures
170
Fig. 7 depicts yield strengths (σ0.2) and compressive strengths (σb) of Ti6Al4V alloys treated by different
171
CTHP procedures. With the increasing number of CTHP cycles, yield strength decreased first and then
172
increased slightly. The CTHP 2-treated alloy manifested the lowest yield strength (decreased by 11% as
173
compared to Ti6Al4V alloy without hydrogen) among all CTHP-treated alloys; however, no significant change
174
in compressive strength was observed, it happened because the CTHP-treated alloys contained more β-phase
175
and α'' martensite phase as compared to Ti6Al4V alloy without hydrogen. The existence of softer β-phase and 9
176
α'' martensite phase can reduce the compressive strength of Ti6Al4V alloy, whereas the presence of hydrogen in
177
both solid solution and δ-hydride can improve the compressive strength of Ti6Al4V alloy. Therefore,
178
compressive strengths of the CTHP-treated Ti6Al4V alloys were affected by softening induced by β-phase and
179
α'' martensite phase as well as by strengthening induced by hydrogen existed in both solid solution and
180
δ-hydride. Moreover, with the increasing number of CTHP cycles, the yield ratio (the ratio of yield strength to
181
compressive strength) decreased first and then increased slightly (Fig. 8). The CTHP 2-treated Ti6Al4V alloy
182
also manifested the lowest yield ratio (decreased by 11% as compared to Ti6Al4V alloy without hydrogen)
183
among all CTHP-treated alloys. However, lower yield ratio is beneficial for Ti alloys to deform plastically at
184
room temperature.
185 186
Fig. 7. Yield strength and compressive strength of Ti6Al4V alloys treated by different CTHP procedures
10
187 188
Fig. 8. Yield ratio of Ti6Al4V alloys treated by different CTHP procedures
189
3.3 Fracture morphologies
190
Macroscopic fracture surfaces of Ti6Al4V alloys (both without hydrogen and treated by CTHP) were found
191
to be aligned at a 45˚ angle with the axes of cylindrical specimens, which indicates that fracture under
192
compression occurred in a shear mode. Fig. 9 exhibits the fracture morphologies of Ti6Al4V alloys treated by
193
different CTHP processes. Two types of fracture characteristics, relatively smooth surface and ductile dimples,
194
were noticed on fracture surfaces of Ti6Al4V alloys (both without hydrogen and treated by CTHP). Therefore,
195
it can be inferred that the dominant fracture mode of the alloys was ductile fracture. With the increasing number
196
of CTHP cycles, areas of ductile dimples increased first and then started to decrease. The CTHP 2-treated
197
Ti6Al4V alloy yielded the largest ductile dimple area and, thus, manifested the best plasticity, which is
198
corresponding to the results of ultimate compression of Ti6Al4V alloys treated under different CTHP processes
199
(Fig. 6).
11
(a)
(b) Smooth surface
Smooth surface
Ductile dimple
Ductile dimple 200
(c)
Ductile dimple
(d)
Smooth surface
Ductile dimple
Smooth surface 201 202
Fig. 9. Fracture surfaces of Ti6Al4V alloys treated by different CTHP procedures: (a) CTHP 0, (b) CTHP 1, (c)
203
CTHP 2, (d) CTHP 3
204
4. Conclusions
205
The ultimate compression of Ti6Al4V alloy increased first and then started to decrease with the increasing
206
number of CTHP cycles. The ultimate compression of the CTHP 2-treated Ti6Al4V alloy was improved by
207
22% as compared to Ti6Al4V alloy without hydrogen and reached 42% (highest among all CTHP-treated
208
alloys). Furthermore, with the increasing number of CTHP cycles, yield strength and yield ratio decreased first
209
and then increased slightly. The CTHP 2-treated Ti6Al4V alloy manifested the lowest yield strength (decreased
210
by 11% as compared to Ti6Al4V alloy without hydrogen) among all CTHP-treated alloys; however, no
211
significant change in compressive strength was noticed. Furthermore, the CTHP 2-treated Ti6Al4V alloy also
212
manifested the lowest yield ratio (decreased by 11% as compared to Ti6Al4V alloy without hydrogen) among
213
all CTHP-treated alloys. In addition, relatively smooth surface and ductile dimples were observed on fracture
214
surfaces of Ti6Al4V alloys (both without hydrogen and treated by CTHP). With the increasing number of
215
CTHP cycles, areas of ductile dimples increased first and then started to decrease, and the CTHP 2-treated 12
216
Ti6Al4V alloy yielded the largest ductile dimple area. Reasons for the improvement in ultimate compression of
217
CTHP-treated alloys are that the amounts of softer β-phase and α'' martensite phase increased under CTHP, and
218
the CTHP 2-treated Ti6Al4V alloy yielded the highest amounts of softer β-phase and α'' martensite phase, and
219
possessed the minimum grain size and a more uniform microstructure among the CTHP-treated alloys. In
220
addition, compressive strengths of the CTHP-treated Ti6Al4V alloys were affected by softening induced by
221
β-phase and α'' martensite phase as well as by strengthening induced by hydrogen existed in both solid solution
222
and δ-hydride.
223
Acknowledgements
224
This work was financially supported by the National Natural Science Foundation of China (No.51875157)
225
and the University Natural Science Research Project of Anhui Province (No.KJ2019A0894).
226
References
227
1.
Y.C. Lin, Y. Tang, X.Y. Zhang, C. Chen, H. Yang, K.C. Zhou, Effects of solution temperature and
228
cooling rate on microstructure and micro-hardness of a hot compressed Ti-6Al-4V alloy, Vacuum
229
159(2019)191-199.
230
2.
Y. Geng, X. Mei, K. Wang, X. Dong, X. Yan, Z. Fan, W. Duan, W. Wang, Effect of Laser Shock
231
Peening on Residual Stress, Microstructure and Hot Corrosion Behavior of Damage-Tolerant TC21
232
Titanium Alloy, J. Mater. Eng. Perform. 27(2018)4703-4713.
233
3.
Y.Q. Jiang, Y.C. Lin, X.Y. Zhang, C. Chen, Q.W. Wang, G.D. Pang, Isothermal tensile deformation
234
behaviors and fracture mechanism of Ti-5Al-5Mo-5V-1Cr-1Fe alloy in beta phase field, Vacuum
235
156 (2018) 187-197.
236
4.
G. Chen, X. Chang, J. Zhang, Y. Jin, C. Sun, Q. Chen, Z. Zhao, Microstructures and Mechanical
237
Properties of In-Situ Al3Ti/2024 Aluminum Matrix Composites Fabricated by Ultrasonic Treatment
238
and
239
https://doi.org/10.1007/s12540-019-00396-y.
240
5.
Subsequent
Squeeze
Casting,
Met.
Mater.
Int.
(2019).
Y. Wang, F. Li, X.W. Li, W.B. Fang, Unusual texture formation and mechanical property in AZ31
241
magnesium alloy sheets processed by CVCDE, J. Mater. Process. Technol. 275 (2020).
242
https://doi.org/10.1016/j.jmatprotec.2019.116360
13
243
6.
Y. Meng, J.Y. Zhang, H. Zhan, Q. Chen, J. Zhou, S. Sugiyama, J. Yanagimoto, Effects of
244
subsequent treatments on the microstructure and mechanical properties of SKD11 tool steel samples
245
processed
246
https://doi.org/10.1016/j.msea.2019.138070
247
7.
248 249
by
multi-stage
thixoforging,
Mater.
Sci.
Eng.,
A
762
(2019).
X. Li, H. Sun, P. Zhang, W. Fang, The effect of strain on dynamic recrystallization of PM Ti–45Al– 10Nb intermetallics during isothermal forging, Intermetallics 55(2014)90-94.
8.
Y. Yu, W. Zhang, W. Dong, J. Yang, Y. Feng, Effects of pre-sintering on microstructure and
250
properties of TiBw/Ti6Al4V composites fabricated by hot extrusion with steel cup, Mater. Sci. Eng.,
251
A 638(2015)38-45.
252
9.
Y.C. Lin, X.Y. Jiang, C.J. Shuai, C.Y. Zhao, D.G. He, M.S. Chen, C. Chen, Effects of initial
253
microstructures on hot tensile deformation behaviors and fracture characteristics of Ti-6Al-4V alloy,
254
Mater. Sci. Eng., A 711(2018)293-302.
255 256
10. B.A. Kolachev, Reversible hydrogen alloying of titanium-alloys, Met. Sci. Heat Treat. 35(1993)586-591.
257
11. Z.G. Sun, G.Q. Chen, Y.Q. Wang, W.L. Zhou, H.L. Hou, Investigations on cold deformation
258
mechanisms of the hydrogenated Ti-6Al-4V alloys, Mater. Sci. Eng., A 527(2010)1003-1007.
259
12. T.F. Ma, R.R. Chen, D.S. Zheng, J.J. Guo, H.S. Ding, Y.Q. Su, H.Z. Fu, Hydrogen enhanced the
260
room-temperature compressive properties of Ti44Al6Nb alloy, Mater. Lett. 213(2018)170-173.
261
13. B.G. Yuan, Y.B. Zheng, L.Q. Gong, Q. Chen, M. Lv, Effect of hydrogen content on microstructures
262 263 264
and room-temperature compressive properties of TC21 alloy, Mater. Des. 94(2016)330-337. 14. Y.Y. Zong, D.B. Shan, M. Xu, Y. Lv, Flow softening and microstructural evolution of TC11 titanium alloy during hot deformation, J. Mater. Process. Technol. 209(2009)1988-1994.
265
15. X. Wang, L. Wang, L. Luo, X. Liu, Y. Tang, X. Li, R. Chen, Y. Su, J. Guo, H. Fu, Hot deformation
266
behavior and dynamic recrystallization of melt hydrogenated Ti-6Al-4V alloy, J. Alloys Compd.
267
728(2017)709-718.
268 269 270 271
16. T. Ma, R. Chen, D. Zheng, J. Guo, H. Ding, Y. Su, H. Fu, Hydrogen-induced softening of Ti-44Al-6Nb-1Cr-2V alloy during hot deformation, Int. J. Hydrogen Energy 42(2017)8329-8337. 17. B. Shao, S.X. Wan, D.B. Shan, B. Guo, Y.Y. Zong, 2017. Hydrogen-Induced Improvement of the Cylindrical Drawing Properties of a Ti-22Al-25Nb Alloy. Adv. Eng. Mater. 19, 201600621.
14
272 273 274 275 276 277
18. M.A. Murzinova, G.A. Salishchev, D.D. Afonichev, Superplasticity of hydrogen-containing VT6 titanium alloy with a submicrocrystalline structure, Phys. Met. Metall. 104(2007)195-202. 19. X. Zhang, Y. Zhao, W. Zeng, Effect of hydrogen on the superplasticity of Ti600 alloy, Int. J. Hydrogen Energy 35(2010)4354-4360. 20. X.L. Wang, Y.Q. Zhao, H.L. Hou, Y.Q. Wang, Effect of hydrogen content on superplastic forming/diffusion bonding of TC21 alloys, J. Alloys Compd. 503(2010)151-154.
278
21. X. Li, G. Jia, F. Qu, H. Wu, J. Chen, Ultrafine Grain Refinement and Superplasticity of Ti-55 Alloy
279
Obtained by Hydrogen Absorption and Desorption, J. Mater. Eng. Perform. 27(2018)3472-3477.
280
22. G.P. Grabovetskaya, E.N. Stepanova, I.V. Ratochka, I.P. Mishin, O.V. Zabudchenko, Effect of
281
Hydrogen on the Development of Superplastic Deformation in the Submicrocrystalline Ti–6Al–4V
282
Alloy, Mater. Sci. Forum 838-839(2016)344-349.
283 284
23. Z. Wang, C. Li, J. Qi, J. Feng, J. Cao, Characterization of hydrogenated niobium interlayer and its application in TiAl/Ti2AlNb diffusion bonding, Int. J. Hydrogen Energy 44(2019)6929-6937.
285
24. L.X. Zhang, Z. Sun, J.M. Shi, X.F. Ye, Z.Y. Yang, J.C. Feng, Low-temperature direct diffusion
286
bonding of hydrogenated TC4 alloy and GH3128 superalloy, Int. J. Hydrogen Energy
287
44(2019)3906-3916.
288 289
25. J.W. Wang, H.R. Gong, Adsorption and diffusion of hydrogen on Ti, Al, and TiAl surfaces, Int. J. Hydrogen Energy 39(2014)6068-6075.
290
26. J.C. Feng, H. Liu, P. He, J. Cao, Effects of hydrogen on diffusion bonding of hydrogenated Ti6Al4V
291
alloy containing 0.3 wt% hydrogen at fast heating rate, Int. J. Hydrogen Energy 32(2007)3054-3058.
292
27. A.A. Ilyin, S.V. Skvortsova, A.M. Mamonov, G.V. Permyakova, D.A. Kurnikov, Effect of
293
thermohydrogen treatment on the structure and properties of titanium alloy castings, Met. Sci. Heat
294
Treat. 44(2002)185-189.
295 296 297 298 299 300
28. F.H. Froes, O.N. Senkov, J.O. Qazi, Hydrogen as a temporary alloying element in titanium alloys: thermohydrogen processing, Int. Mater. Rev. 49(2004)227-245. 29. C.P. Liang, H.R. Gong, Fundamental influence of hydrogen on various properties of alpha-titanium, Int. J. Hydrogen Energy 35(2010)3812-3816. 30. L.C. Campanelli, P.S.C.P. da Silva, A.M. Jorge, C. Bolfarini, Effect of hydrogen on the fatigue behavior of the near-β Ti-5Al-5Mo-5V-3Cr alloy, Scr. Mater. 132(2017)39-43.
15
301 302 303 304 305 306 307 308 309 310 311 312
31. B.G. Yuan, C.F. Li, H.P. Yu, D.L. Sun, Influence of hydrogen content on tensile and compressive properties of Ti-6Al-4V alloy at room temperature, Mater. Sci. Eng., A 527(2010)4185-4190. 32. J.I. Qazi, O.N. Senkov, J. Rahim, F.H. Froes, Kinetics of martensite decomposition in Ti-6Al-4V-xH alloys, Mater. Sci. Eng., A 359(2003)137-149. 33. X. Wang, Y. Zhao, W. Zeng, H. Hou, Y. Wang, Kinetics of Hydrogen Absorption of Ti600 Alloy, Rare Met. Mater. Eng. 38(2009)1219-1222. 34. T.K. Zhu, M.Q. Li, Effect of 0.770 wt%H addition on the microstructure of Ti-6Al-4V alloy and mechanism of [delta] hydride formation, J. Alloys Compd. 481(2009)480-485. 35. C.Y. Yu, C.C. Shen, T.P. Perng, Microstructure of Ti-6Al-4V processed by hydrogenation, Scr. Mater. 55(2006)1023-1026. 36. J. Zhao, H. Ding, Y. Zhong, C.S. Lee, Effect of thermo hydrogen treatment on lattice defects and microstructure refinement of Ti6Al4V alloy, Int. J. Hydrogen Energy 35(2010)6448-6454.
313
37. N. Pushilina, A. Panin, M. Syrtanov, E. Kashkarov, V. Kudiiarov, O. Perevalova, R. Laptev, A. Lider,
314
A. Koptyug, 2018. Hydrogen-Induced Phase Transformation and Microstructure Evolution for
315
Ti-6Al-4V Parts Produced by Electron Beam Melting. Met. 8, 301.
316
38. R. de Araujo-Silva, A.M. Neves, L.E.R. Vega, M.R.M. Triques, D.R. Leiva, C.S. Kiminami, T.T.
317
Ishikawa, A.M. Jorge Junior, W.I. Botta, Synthesis of beta-Ti-Nb alloys from elemental powders by
318
high-energy ball milling and their hydrogenation features, Int. J. Hydrogen Energy
319
43(2018)18382-18391.
320
16
321
Figure captions:
322
Fig. 1. Photos of Ti6Al4V alloy samples before (a) and after (b) CTHP
323
Fig. 2. OM micrographs of Ti6Al4V alloys treated by different CTHP procedures: (a) CTHP 0, (b) CTHP 1, (c)
324
CTHP 2, (d) CTHP 3
325
Fig. 3. Grain size of Ti6Al4V alloys treated by different CTHP procedures
326
Fig. 4. XRD patterns of Ti6Al4V alloys treated by different CTHP procedures: (a) CTHP 0, (b) CTHP 1, (c)
327
CTHP 2, (d) CTHP 3
328
Fig. 5. TEM micrographs of Ti6Al4V alloys treated by different CTHP procedures: (a) δ-hydride laths and its
329
SAED pattern in the CTHP 1-treated Ti6Al4V alloy, (b) α' and α'' matrix structure and its SAED pattern in
330
the CTHP 1-treated Ti6Al4V alloy, (c) α' matrix structure and its SAED pattern in the CTHP 2-treated
331
Ti6Al4V alloy, (d) δ-hydride laths and its SAED pattern in the CTHP 2-treated Ti6Al4V alloy, (e) α' matrix
332
structure and its SAED pattern in the CTHP 3-treated Ti6Al4V alloy, (f) δ-hydride laths and its SAED pattern
333
in the CTHP 3-treated Ti6Al4V alloy.
334
Fig. 6. Ultimate compression of Ti6Al4V alloys treated by different CTHP procedures
335
Fig. 7. Yield strength and compressive strength of Ti6Al4V alloys treated by different CTHP procedures
336
Fig. 8. Yield ratio of Ti6Al4V alloys treated by different CTHP procedures
337
Fig. 9. Fracture surfaces of Ti6Al4V alloys treated by different CTHP procedures: (a) CTHP 0, (b) CTHP 1, (c)
338
CTHP 2, (d) CTHP 3
17
Fig. 1. Photos of Ti6Al4V alloy samples before (a) and after (b) CTHP
Fig. 2. OM micrographs of Ti6Al4V alloys treated by different CTHP procedures: (a) CTHP 0, (b) CTHP 1, (c) CTHP 2, (d) CTHP 3
Fig. 3. Grain size of Ti6Al4V alloys treated by different CTHP procedures
Fig. 4. XRD patterns of Ti6Al4V alloys treated by different CTHP procedures: (a) CTHP 0, (b) CTHP 1, (c) CTHP 2, (d) CTHP 3
Fig. 5. TEM micrographs of Ti6Al4V alloys treated by different CTHP procedures: (a) δ-hydride laths and its SAED pattern in the CTHP 1-treated Ti6Al4V alloy, (b) α' and α'' matrix structure and its SAED pattern in the CTHP 1-treated Ti6Al4V alloy, (c) α' matrix structure and its SAED pattern in the
CTHP 2-treated Ti6Al4V alloy, (d) δ-hydride laths and its SAED pattern in the CTHP 2-treated Ti6Al4V alloy, (e) α' matrix structure and its SAED pattern in the CTHP 3-treated Ti6Al4V alloy, (f) δ-hydride laths and its SAED pattern in the CTHP 3-treated Ti6Al4V alloy.
Fig. 6. Ultimate compression of Ti6Al4V alloys treated by different CTHP procedures
Fig. 7. Yield strength and compressive strength of Ti6Al4V alloys treated by different CTHP procedures
Fig. 8. Yield ratio of Ti6Al4V alloys treated by different CTHP procedures
(a)
Smooth surface
(b)
Ductile dimple
(c)
Ductile dimple
Smooth surface
Smooth surface
Ductile dimple
(d)
Smooth surface
Ductile dimple
Fig. 9. Fracture surfaces of Ti6Al4V alloys treated by different CTHP procedures: (a) CTHP 0, (b) CTHP 1, (c) CTHP 2, (d) CTHP 3
Declarations of interest: none
Highlights Plasticity increased first and then decreased under CTHP. Yield ratio decreased first and then increased slightly under CTHP. The amounts of softer β and α'' phases increased after treatment of CTHP. Area of ductile dimple increased first and then decreased under CTHP.