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Hydrogen induced softening mechanism in near alpha titanium alloy
Journal of Alloys and Compounds 541 (2012) 60–64
Contents lists available at SciVerse ScienceDirect
Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jalcom
Hydrogen induced softening mechanism in near alpha titanium alloy Zong Yingying, Huang Shuhui, Feng Yingjuan, Shan Debin ⇑ School of Materials Science and Engineering, Harbin Institute of Technology, Harbin 150001, PR China
a r t i c l e
i n f o
Article history: Received 29 April 2012 Received in revised form 2 June 2012 Accepted 15 June 2012 Available online 28 June 2012 Keywords: Near alpha titanium alloy Hydrogen Softening mechanism
a b s t r a c t Hydrogen induced softening mechanism in near alpha titanium alloy Ti600 (Ti-6.0Al-2.5Sn-4.0Zr-0.5Mo0.4Si-0.1Y) is studied in this paper. The results show that, the hydrogenated alloy has a lower flow stress than the unhydrogenated alloy at the same temperature, and 0.3 wt.% is the optimum hydrogen content for Ti600 deformed at 850 °C. Compared with unhydrogenated Ti600 deformed at 950 °C, deformation resistance decreases by one third. A new method based on thermogravimetry analysis for studying dehydrogenation process of titanium alloys is proposed. The results demonstrate that with the addition of 0.3 wt.% hydrogen the phase transformation temperature decreases by around 100 °C. The reasons of hydrogen-induced softening are studied by XRD and TEM methods, which are attributed to the effects of hydrogen on the phase transformation, dislocation mobility, dynamic recovery and recrystallization during high temperature deformation process. Ó 2012 Elsevier B.V. All rights reserved.
1. Introduction Near alpha titanium alloy, an ideal material for aero engine, can be used nearly 600 °C and has been researched widely [1–3]. Near alpha titanium alloys have excellent high temperature mechanical properties, especially creep resistance. On the one hand, these properties make them can be used at high temperature; on the other hand, these properties also make their deformation more difficult than the other titanium alloys [3–5]. Thermohydrogen processing of titanium alloy, which hydrogen is used as a temporary alloying element, is a new technology for improving the hot workability of titanium alloy through changing microstructure [6,7]. Of course, hydrogen must be removed by vacuum heat treatment after deformation, in case of hydrogen embrittlement during service. Some reports show that [8–10] thermohydrogen technology can decrease the forging load, the flow stress by around 70% or the forming temperature by 100–150 °C. The reasons why thermohydrogen technology decreases the hot deformation resistance and temperature of titanium alloys are reported from the following aspects [11–14]. (1) Hydrogen addition makes the b phase transformation temperature decrease, so b phase volume fraction increases at the same temperature. While b phase has more slip systems than a phase, so it can be deformed more easily. (2) Diffusion of hydrogen in titanium alloy makes dislocation be activated and move ⇑ Corresponding author. Tel./fax: +86 451 86418732. E-mail address: [email protected] (S. Debin). 0925-8388/$ - see front matter Ó 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.jallcom.2012.06.099
easily, and then plastic deformation resistance decreases. (3) Hydrogen addition accelerates the element diffusion in titanium alloy. Especially for Al, the diffusion coefficient is increased tenfold both in a phase and b phase. With the increasing of elements diffusion, b phase diffusion creep and a phase grain boundary sliding are strengthened, which promotes dynamic recovery and recrystallization, and then plasticity of titanium alloy is enhanced. The purpose of this work is to research the effects of hydrogen addition on the high temperature deformation and microstructure evolution in Ti600, which can provide evidence for the application of thermohydrogen processing to hot deformation of near alpha titanium alloy. 2. Materials and methods Ti600, IMI834 and Ti1100 are representatives of near alpha titanium alloy. The chemical compositions (wt.%) and mechanical properties of the three near alpha titanium alloys are shown in Tables 1 and 2. Tensile strength, creep resistance and high temperature endurance of Ti600 are improved by yttrium addition [2]. The a + b ? b transition temperature of Ti600 alloy is determined to be about 1010 °C. Specimens for hydrogenation and hot compression tests are cylinders of 8 mm in diameter and 12 mm in height. The hydrogenated process of specimens is as follows: making vacuum ? adding argon ? heating to 750 °C ? adding hydrogen ? insulation for 2 h ? furnace cooling to room temperature, according to phase diagram and experiments. The hydrogen content of specimens is controlled through equilibrium partial pressure, and then measured by Sartorius BT125D electronic balance with an accuracy of 0.01 mg. Isothermal hot compression tests are carried out on an Instron 5500 machine with an initial strain rate of 0.01 s1 at the deformation temperatures of 750, 800, 850, 900 and 950 °C. Glass lubricant is coated on the surface of the specimens. The specimens are held at compression temperature for 5 min and subsequently deformed until deformation extent reaches 50%. After compression, the specimens
Z. Yingying et al. / Journal of Alloys and Compounds 541 (2012) 60–64 Table 1 Chemical composition (wt.%) of near alpha titanium alloys. Alloy
Al
Sn
Zr
Mo
Si
Ti600 IMI834 Ti1100
6.0 5.8 6.0
2.5 4.0 2.75
4.0 3.5 4.0
0.5 0.5 0.4
0.4 0.35 0.4
Nb
Y
Ti
0.1
Bal. Bal. Bal.
0.7
Table 2 Mechanical properties of near alpha titanium alloys [2]. Alloy
Ti600 IMI834 Ti1100
Tensile at 20 °C/600 °C
Creep with 600 °C/100 h/150 MPa
rb (Mpa)
rs (Mpa)
d (%)
w (%)
ee (%)
1068/745 1070/680 960/630
1050/615 960/550 860/530
11/16 14/15 11/14
13/31 20/50 18/30
0.03 0.1 0.1
are quenched immediately in water. Thermogravimetry analysis (TGA) is performed by using a Mettler Toledo TGA/SDTA851 in the temperature range 70–1150 °C in an argon atmosphere with a heating rate of 10 °C/min. Microstructures are observed by Philips-CM12 transmission electron microscopy (TEM).
3. Results and discussion 3.1. Deformation behavior The true stress–strain curves pertaining to the unhydrogenated and hydrogenated alloys are presented in Fig. 1. The conclusion is apparent: (1) under all deformation conditions, flow stress of the hydrogenated samples is lower than the
Fig. 1. True stress–strain curves of the alloys deformed with 0.01 s1: (a) unhydrogenated Ti600, (b) 0.3 wt.% hydrogenated Ti600.
61
unhydrogenated. (2) When the deformation temperature is lower than 800 °C, the true stress–strain curves of the unhydrogenated alloy are very flat meaning that dynamic recovery is the main softening mechanism; the true stress–strain curves of the hydrogenated alloy decrease sharply meaning that dynamic recrystallization is the main softening mechanism. In a phase region, hydrogen addition can promote dynamic recrystallization of a phase during hot deformation process, which is the reason of the above differences. (3) When the deformation temperature is higher than 850 °C, the true stress–strain curves of the unhydrogenated alloy exhibit dynamic recrystallization characteristic, which flow stress decreases rapidly after the peak stress; and the true stress–strain curves of the hydrogenated alloy exhibit dynamic recovery characteristic, which flow stress stabilizes at a certain level. Since hydrogen addition has decreased the phase transformation temperature, in a + b phase region at the same temperature, there is more b phase in hydrogenated alloy than in unhydrogenated. So dynamic recrystallization of a phase is the main softening mechanism of the unhydrogenated alloy, and dynamic recovery of b phase is main softening mechanism of hydrogenated alloy. The effect of hydrogen at high temperature deformation of Ti600 alloy is researched by lots of hot compression tests, including hydrogen content of 0, 0.1, 0.2, 0.3, 0.4, 0.5, 1.0 wt.% and temperature of 750, 800, 850, 900, 950 °C at 0.01 s1. The relationship between the flow stress and hydrogen content is presented in Fig. 2. The flow stress during the hot deformation process is influenced by many factors such as strain, strain rate and temperature etc. This paper focuses on the effect of hydrogen content and temperature on the flow stress. It is shown that: (1) Hydrogen addition can decrease the flow stress significantly, and then compared with the unhydrogenated alloy, the maximum degree of reduction is by around 70%. (2) The flow stress decreases with the rising of temperature, and the unhydrogenated alloy is more sensitive to deformation temperature than hydrogenated alloy. When the deformation temperature increases from 750 to 950 °C, for the unhydrogenated alloy, the reduction of the flow stress is about 300 MPa, while for the hydrogenated alloy, the reduction of the flow stress ranges from about 100 to 250 MPa. The above appearance can be explained as follows: on the one hand, with the rising of temperature, the increase of b phase volume fraction is the main reason for the decrease of the flow stress. BCC b phase has more slip systems and can be deformed more easily than a HCP a phase, so plasticity of b phase is better than a phase. In the low deformation temperature region, the increasing rate of b phase volume fraction becomes faster with the rising of
Fig. 2. Effects of hydrogen content on the flow stress of Ti600 alloys deformed at 0.01 s1.
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Z. Yingying et al. / Journal of Alloys and Compounds 541 (2012) 60–64
Fig. 5. XRD patterns of the alloys after quenching with deformed at 0.01 s1, 850 °C: (a) unhydrogenated, (b) 0.3 wt.% hydrogenated.
Fig. 3. (a) TGA curves of unhydrogenated and hydrogenated Ti600, (b) quality loss curve of hydrogenated Ti600.
with the increasing of hydrogen addition, Young’s modulus and shear modulus of a phase decrease, on the contrary, the two modulus of b phase increase. According to this result, the theory of hydrogen addition causing a phase softening and b phase hardening is proposed. It can be used to explain that optimal hydrogen content exists for decreasing deformation resistance for a certain titanium alloy at a certain temperature range. Hydrogen addition softens a phase and increases the amount of b phase, which makes the deformation resistance decrease. But hydrogen addition hardens b phase at the same time, which makes deformation resistance increase. So with the increasing of hydrogen content, deformation resistance reduces firstly, but the decreasing rate slows down gradually. When the two converse effects reach balance, the hydrogen content is optimum. As can be seen from Fig. 2, when the deformation temperature is less than 800 °C (Lower deformation temperature and deformation resistance are equally important for the formation of titanium alloys), 0.3 wt.% is an obviously optimum hydrogen content for Ti600. So this work is to research the effects of 0.3 wt.% hydrogen addition on the high temperature deformation characteristics, dynamic recovery, recrystallization and microstructure evolution in Ti600. 3.2. Microstructure evolution
Fig. 4. The equilibrium phase transition of 0.3 wt.% hydrogenated and unhydrogenated Ti600.
temperature, which leads to the decreasing rate of the flow stress becoming faster too. On the other hand, when the alloys are deformed in the lower part of a + b phase region, the amount of a phase is larger than b phase, so the main softening mechanism is dynamic recrystallization of a phase. When the alloys are deformed in the upper part of a + b phase region, the amount of b phase is more than a phase, so the main softening mechanism is dynamic recovery of b phase. Above all, the sensitivity of stress to deformation temperature decreases with the rising of temperature. Senkov [15,16] researched the elastic module of hydrogenated titanium alloys by using laser ultrasonic. The result is obvious:
Because of oxidation and dehydrogenation, the phase transition of hydrogenated titanium alloy cannot be directly measured by TGA or DTA. In this paper, a new method based on TGA for studying dehydrogenation process of titanium alloys is proposed, and the major process is as follows. The TGA curves of unhydrogenated and hydrogenated Ti600 are shown in Fig. 3, and they are obtained by using the following equations:
mC ¼
M ST 1 100% MS
mL ¼ mCunh mCh
ð1Þ ð2Þ
where MS is initial quality of the samples, MST is changing quality of the samples during heating process, mC is percentage change in quality of the samples during heating process, and mL is percentage loss in quality of the hydrogenated sample during heating process. The bond energy of b phase is smaller than a phase [16], so b phase can be oxidized more easily, meanwhile diffusion velocity of oxygen atoms in b phase is faster than in a phase. Then at the beginning of the a ? b phase transition, the increasing rate
Z. Yingying et al. / Journal of Alloys and Compounds 541 (2012) 60–64
63
Fig. 6. TEM micrographs of Ti600 alloys deformed at 0.01 s1: (a) dislocation cells of unhydrogenated Ti600 at 850 °C, (b) DRX of unhydrogenated Ti600 at 950 °C, (c) the lamellar microstructure of 0.3 wt.% hydrogenated Ti600 at 850 °C, (d) titanium hydride d phase of 0.3 wt.% hydrogenated Ti600 at 850 °C, (e) SAED pattern of the FCC d titanium hydride with twins substructure at 850 °C, (f) the lamellar microstructure of 0.3 wt.% hydrogenated Ti600 at 950 °C, (g) titanium hydride d phase of 0.3 wt.% hydrogenated Ti600 at 950 °C.
of quality is accelerated because of oxidation. At the end of the a ? b phase transition, the increasing rate of quality reaches a
maximum. Above all, the a + b phase region of unhydrogenated Ti600 alloy ranges from about 790 to 1010 °C.
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Z. Yingying et al. / Journal of Alloys and Compounds 541 (2012) 60–64
The increasing quality of unhydrogenated Ti600 is caused by oxidation, while the change of quality is the result of oxidation and dehydrogenated for hydrogenated Ti600. So the difference of the mCunh and mCh is the percentage loss in quality, which caused by dehydrogenation, as shown in Fig. 3b. The transition of aH ? a + H2 occurs at the first stage, and then quality loss goes slowly because of low hydrogen solubility in aH. At the second stage, aH ? bH and b ?b + H2 occur about 700 °C, and then quality loss goes quickly because of high hydrogen solubility in bH. The maximum of quality loss comes at 900 °C meaning that aH ? bH has finished. At the third stage, oxidation is the main reason of quality change for the two samples, so the difference tends to zero quickly. Above all, the a + b phase region of 0.3 wt.% hydrogenated Ti600 alloy lies approximately between 700 and 900 °C. Based on Fig. 3, equilibrium phase transition of unhydrogenated and hydrogenated Ti600 is described in Fig. 4. The a + b phase region of unhydrogenated Ti600 ranges from 790 to 1010 °C. The XRD pattern of the alloys after quenching with deformed at a phase region of 850 °C is shown in Fig. 5a. There is only a phase as illustrated in Fig. 5a. It suggests b phase cannot be retained to room temperature by quenching from a + b phase region, and b ? a phase transition goes completely. It is a characteristic of near alpha titanium alloy. The a + b phases region of hydrogenated Ti600 lies approximately from 700 to 900 °C. The XRD pattern of the alloys after quenching with deformed at a + aH + b + bH phase region of 850 °C is shown in Fig. 5b with a, a0 , d and b phases. bH ? a0 (a) + d, bH ? b+d and b ? a phase transitions occur during the process of quenching. The metastable b phase of hydrogenated Ti600 only comes from bH ? b + d phase transition. It is hard to distinguish a0 phase and a by XRD, so the above conclusions concerned with a0 phase and a phase must be proved by TEM. Fig. 6 displays the effects of hydrogen and temperature on the TEM micrographs of the Ti600 alloys. After quenching, lamellar microstructure in the deformed hydrogenated alloy increases obviously, and lamellar microstructure comes from the bH phase quenching at high temperature. a phase dynamic recovery and recrystallization of unhydrogenated Ti600 alloy are shown in Fig. 6a and b, including dislocation cells and recrystallization grains. That suggests dynamic recovery of a phase is the main softening mechanism at 850 °C, and dynamic recrystallization of a phase is the main softening mechanism at 950 °C. The lamellar titanium hydride a phase of 0.3 wt.% hydrogenated Ti600 alloy, which comes from the quenching of bH phase at high temperature, is shown in Fig. 6c. The substructure of the lamellar titanium hydride phase is twin, as shown in Fig. 6d. Fig. 6e is the SAED pattern of the FCC a titanium hydride, which reveals the orientation rela 1 1 1). All of tionship of adjacent twinned regions to be (1 1 1)k( the samples are held at deformation temperature for 15 min to ensure the phase transition completely. Compared Fig. 6c with Fig. 6f, the width of lath titanium hydride phase expands with the rising of deformation temperature. With a comparison between Fig. 6d and g, it can be seen that the width of the twin substructure becomes bigger with the rising of deformation temperature. The lamellar microstructure consists of martensite a0 , metastable b and titanium hydride d. The dislocation substructure of martensite a0 and metastable b is caused by the deformation. The twin substructure of titanium hydride d is caused by bH ? a + d transformation during quenching process. The difference of the specific volume between the new phase d and the parent phase bH is enormous, and the
new phase d is limited in the growing process, which results in generating twins. The b phase of unhydrogenated Ti600 cannot be retained to room temperature by quenching, so it cannot be strengthened by heat treatment. Because of hydrogen addition, bH phase and d phase are retained by quenching, then the near alpha titanium alloy changes to a + b titanium alloy, which can be strengthened by heat treatment. Of course, the hydrogen will be removed by vacuum dehydrogenation heat treatment in case of Hydrogen embrittlement. 4. Conclusions (1) Hydrogen addition causes a phase softening and decreases b phase transformation temperature, but hydrogen causes b phase hardening. When the two contrary effects reach equilibrium, the hydrogen content is optimum. 0.3 wt.% is an optimal hydrogen content for Ti600 deformed at 850 °C. Compared with unhydrogenated Ti600 deformed at 950 °C, deformation temperature decreases by 100 °C and deformation resistance decreases by one third. Because of hydrogen addition, the main softening mechanism changes from dynamic recrystallization of a phase to dynamic recovery of b phase. (2) A new method based on TGA for studying phase transition of hydrogenated titanium alloy is proposed. With the study of weight gain by oxidation and weight loss by dehydrogenation, the a + b phase region of unhydrogenated and 0.3 wt.% hydrogenated Ti600 alloy are about from 790 to 1010 °C and from 700 to 900 °C. And the optimum temperature for dehydrogenation of titanium alloy is above 700 °C, which is the turning point of dehydrogenation rate. (3) The lamellar microstructure is gained through quenching after high temperature deformation in 0.3 wt.% hydrogenated Ti600, which consists of martensite a0 , metastable b and titanium hydride d, and the width of the lamellar become larger with the rising of deformation temperature.
Acknowledgments This work was supported by the National Natural Science Foundation of China (No. 50805037). References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16]
Q. Hong, Rare Metal Mat. Eng. 34 (2005) 1334–1337. Y.Q. Zhao, Titanium industry progress 1 (2001) 33–39. S.W. Xin, Q. Hong, Trans. Nonferrous Met. Soc. China 20 (2010) 2142–2147. P. Wanjara, M. Jahazi, H. Monajati, S. Yue, J.P. Immarigeon, Mater. Sci. Eng. A 396 (2005) 50–60. X.Z. Cai, D. Eylon, Rare Metal Mat. Eng. 23 (1994) 3–7. F.H. Froes, O.N. Senkov, J.I. Qazi, Int. Mater. Rev. 49 (2004) 227–245. H.L. Hou, Trans. Nonferrous Met. Soc. China 13 (2003) 533–549. Y.Y. Zong, D.B. Shan, Y. Lv, Int. J. Hydrogen Energy 32 (2007) 3936–3940. Y.Y. Zong, D.B. Shan, Y. Lv, Scr. Mater. 58 (2008) 449–453. L.S. Luo, Y.Q. Su, J.J. Guo, H.Z. Fu, J. Alloys Compd. 425 (2006) 140–144. T.K. Zhu, M.Q. Li, J. Alloys Compd. 481 (2009) 480–485. J.W. Zhao, H. Ding, H.L. Hou, Z.Q. Li, J. Alloys Compd. 491 (2010) 673–678. L.I. Anisimova, Met. Sci. Heat Treat. 34 (1992) 143–147. Z.G. Sun, H.L. Hou, W.L. Zhou, Y.Q. Wang, Z.Q. Li, J. Alloys Compd. 476 (2009) 550–555. O.N. Senkov, J.J. Jonas, Metall. Mater. Trans. A 27 (1996) 1303–1312. O.N. Senkov, M. Dubois, J.J. Jonas, Metall. Mater. Trans. A 27 (1996) 3963–3970.
Contents lists available at SciVerse ScienceDirect
Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jalcom
Hydrogen induced softening mechanism in near alpha titanium alloy Zong Yingying, Huang Shuhui, Feng Yingjuan, Shan Debin ⇑ School of Materials Science and Engineering, Harbin Institute of Technology, Harbin 150001, PR China
a r t i c l e
i n f o
Article history: Received 29 April 2012 Received in revised form 2 June 2012 Accepted 15 June 2012 Available online 28 June 2012 Keywords: Near alpha titanium alloy Hydrogen Softening mechanism
a b s t r a c t Hydrogen induced softening mechanism in near alpha titanium alloy Ti600 (Ti-6.0Al-2.5Sn-4.0Zr-0.5Mo0.4Si-0.1Y) is studied in this paper. The results show that, the hydrogenated alloy has a lower flow stress than the unhydrogenated alloy at the same temperature, and 0.3 wt.% is the optimum hydrogen content for Ti600 deformed at 850 °C. Compared with unhydrogenated Ti600 deformed at 950 °C, deformation resistance decreases by one third. A new method based on thermogravimetry analysis for studying dehydrogenation process of titanium alloys is proposed. The results demonstrate that with the addition of 0.3 wt.% hydrogen the phase transformation temperature decreases by around 100 °C. The reasons of hydrogen-induced softening are studied by XRD and TEM methods, which are attributed to the effects of hydrogen on the phase transformation, dislocation mobility, dynamic recovery and recrystallization during high temperature deformation process. Ó 2012 Elsevier B.V. All rights reserved.
1. Introduction Near alpha titanium alloy, an ideal material for aero engine, can be used nearly 600 °C and has been researched widely [1–3]. Near alpha titanium alloys have excellent high temperature mechanical properties, especially creep resistance. On the one hand, these properties make them can be used at high temperature; on the other hand, these properties also make their deformation more difficult than the other titanium alloys [3–5]. Thermohydrogen processing of titanium alloy, which hydrogen is used as a temporary alloying element, is a new technology for improving the hot workability of titanium alloy through changing microstructure [6,7]. Of course, hydrogen must be removed by vacuum heat treatment after deformation, in case of hydrogen embrittlement during service. Some reports show that [8–10] thermohydrogen technology can decrease the forging load, the flow stress by around 70% or the forming temperature by 100–150 °C. The reasons why thermohydrogen technology decreases the hot deformation resistance and temperature of titanium alloys are reported from the following aspects [11–14]. (1) Hydrogen addition makes the b phase transformation temperature decrease, so b phase volume fraction increases at the same temperature. While b phase has more slip systems than a phase, so it can be deformed more easily. (2) Diffusion of hydrogen in titanium alloy makes dislocation be activated and move ⇑ Corresponding author. Tel./fax: +86 451 86418732. E-mail address: [email protected] (S. Debin). 0925-8388/$ - see front matter Ó 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.jallcom.2012.06.099
easily, and then plastic deformation resistance decreases. (3) Hydrogen addition accelerates the element diffusion in titanium alloy. Especially for Al, the diffusion coefficient is increased tenfold both in a phase and b phase. With the increasing of elements diffusion, b phase diffusion creep and a phase grain boundary sliding are strengthened, which promotes dynamic recovery and recrystallization, and then plasticity of titanium alloy is enhanced. The purpose of this work is to research the effects of hydrogen addition on the high temperature deformation and microstructure evolution in Ti600, which can provide evidence for the application of thermohydrogen processing to hot deformation of near alpha titanium alloy. 2. Materials and methods Ti600, IMI834 and Ti1100 are representatives of near alpha titanium alloy. The chemical compositions (wt.%) and mechanical properties of the three near alpha titanium alloys are shown in Tables 1 and 2. Tensile strength, creep resistance and high temperature endurance of Ti600 are improved by yttrium addition [2]. The a + b ? b transition temperature of Ti600 alloy is determined to be about 1010 °C. Specimens for hydrogenation and hot compression tests are cylinders of 8 mm in diameter and 12 mm in height. The hydrogenated process of specimens is as follows: making vacuum ? adding argon ? heating to 750 °C ? adding hydrogen ? insulation for 2 h ? furnace cooling to room temperature, according to phase diagram and experiments. The hydrogen content of specimens is controlled through equilibrium partial pressure, and then measured by Sartorius BT125D electronic balance with an accuracy of 0.01 mg. Isothermal hot compression tests are carried out on an Instron 5500 machine with an initial strain rate of 0.01 s1 at the deformation temperatures of 750, 800, 850, 900 and 950 °C. Glass lubricant is coated on the surface of the specimens. The specimens are held at compression temperature for 5 min and subsequently deformed until deformation extent reaches 50%. After compression, the specimens
Z. Yingying et al. / Journal of Alloys and Compounds 541 (2012) 60–64 Table 1 Chemical composition (wt.%) of near alpha titanium alloys. Alloy
Al
Sn
Zr
Mo
Si
Ti600 IMI834 Ti1100
6.0 5.8 6.0
2.5 4.0 2.75
4.0 3.5 4.0
0.5 0.5 0.4
0.4 0.35 0.4
Nb
Y
Ti
0.1
Bal. Bal. Bal.
0.7
Table 2 Mechanical properties of near alpha titanium alloys [2]. Alloy
Ti600 IMI834 Ti1100
Tensile at 20 °C/600 °C
Creep with 600 °C/100 h/150 MPa
rb (Mpa)
rs (Mpa)
d (%)
w (%)
ee (%)
1068/745 1070/680 960/630
1050/615 960/550 860/530
11/16 14/15 11/14
13/31 20/50 18/30
0.03 0.1 0.1
are quenched immediately in water. Thermogravimetry analysis (TGA) is performed by using a Mettler Toledo TGA/SDTA851 in the temperature range 70–1150 °C in an argon atmosphere with a heating rate of 10 °C/min. Microstructures are observed by Philips-CM12 transmission electron microscopy (TEM).
3. Results and discussion 3.1. Deformation behavior The true stress–strain curves pertaining to the unhydrogenated and hydrogenated alloys are presented in Fig. 1. The conclusion is apparent: (1) under all deformation conditions, flow stress of the hydrogenated samples is lower than the
Fig. 1. True stress–strain curves of the alloys deformed with 0.01 s1: (a) unhydrogenated Ti600, (b) 0.3 wt.% hydrogenated Ti600.
61
unhydrogenated. (2) When the deformation temperature is lower than 800 °C, the true stress–strain curves of the unhydrogenated alloy are very flat meaning that dynamic recovery is the main softening mechanism; the true stress–strain curves of the hydrogenated alloy decrease sharply meaning that dynamic recrystallization is the main softening mechanism. In a phase region, hydrogen addition can promote dynamic recrystallization of a phase during hot deformation process, which is the reason of the above differences. (3) When the deformation temperature is higher than 850 °C, the true stress–strain curves of the unhydrogenated alloy exhibit dynamic recrystallization characteristic, which flow stress decreases rapidly after the peak stress; and the true stress–strain curves of the hydrogenated alloy exhibit dynamic recovery characteristic, which flow stress stabilizes at a certain level. Since hydrogen addition has decreased the phase transformation temperature, in a + b phase region at the same temperature, there is more b phase in hydrogenated alloy than in unhydrogenated. So dynamic recrystallization of a phase is the main softening mechanism of the unhydrogenated alloy, and dynamic recovery of b phase is main softening mechanism of hydrogenated alloy. The effect of hydrogen at high temperature deformation of Ti600 alloy is researched by lots of hot compression tests, including hydrogen content of 0, 0.1, 0.2, 0.3, 0.4, 0.5, 1.0 wt.% and temperature of 750, 800, 850, 900, 950 °C at 0.01 s1. The relationship between the flow stress and hydrogen content is presented in Fig. 2. The flow stress during the hot deformation process is influenced by many factors such as strain, strain rate and temperature etc. This paper focuses on the effect of hydrogen content and temperature on the flow stress. It is shown that: (1) Hydrogen addition can decrease the flow stress significantly, and then compared with the unhydrogenated alloy, the maximum degree of reduction is by around 70%. (2) The flow stress decreases with the rising of temperature, and the unhydrogenated alloy is more sensitive to deformation temperature than hydrogenated alloy. When the deformation temperature increases from 750 to 950 °C, for the unhydrogenated alloy, the reduction of the flow stress is about 300 MPa, while for the hydrogenated alloy, the reduction of the flow stress ranges from about 100 to 250 MPa. The above appearance can be explained as follows: on the one hand, with the rising of temperature, the increase of b phase volume fraction is the main reason for the decrease of the flow stress. BCC b phase has more slip systems and can be deformed more easily than a HCP a phase, so plasticity of b phase is better than a phase. In the low deformation temperature region, the increasing rate of b phase volume fraction becomes faster with the rising of
Fig. 2. Effects of hydrogen content on the flow stress of Ti600 alloys deformed at 0.01 s1.
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Z. Yingying et al. / Journal of Alloys and Compounds 541 (2012) 60–64
Fig. 5. XRD patterns of the alloys after quenching with deformed at 0.01 s1, 850 °C: (a) unhydrogenated, (b) 0.3 wt.% hydrogenated.
Fig. 3. (a) TGA curves of unhydrogenated and hydrogenated Ti600, (b) quality loss curve of hydrogenated Ti600.
with the increasing of hydrogen addition, Young’s modulus and shear modulus of a phase decrease, on the contrary, the two modulus of b phase increase. According to this result, the theory of hydrogen addition causing a phase softening and b phase hardening is proposed. It can be used to explain that optimal hydrogen content exists for decreasing deformation resistance for a certain titanium alloy at a certain temperature range. Hydrogen addition softens a phase and increases the amount of b phase, which makes the deformation resistance decrease. But hydrogen addition hardens b phase at the same time, which makes deformation resistance increase. So with the increasing of hydrogen content, deformation resistance reduces firstly, but the decreasing rate slows down gradually. When the two converse effects reach balance, the hydrogen content is optimum. As can be seen from Fig. 2, when the deformation temperature is less than 800 °C (Lower deformation temperature and deformation resistance are equally important for the formation of titanium alloys), 0.3 wt.% is an obviously optimum hydrogen content for Ti600. So this work is to research the effects of 0.3 wt.% hydrogen addition on the high temperature deformation characteristics, dynamic recovery, recrystallization and microstructure evolution in Ti600. 3.2. Microstructure evolution
Fig. 4. The equilibrium phase transition of 0.3 wt.% hydrogenated and unhydrogenated Ti600.
temperature, which leads to the decreasing rate of the flow stress becoming faster too. On the other hand, when the alloys are deformed in the lower part of a + b phase region, the amount of a phase is larger than b phase, so the main softening mechanism is dynamic recrystallization of a phase. When the alloys are deformed in the upper part of a + b phase region, the amount of b phase is more than a phase, so the main softening mechanism is dynamic recovery of b phase. Above all, the sensitivity of stress to deformation temperature decreases with the rising of temperature. Senkov [15,16] researched the elastic module of hydrogenated titanium alloys by using laser ultrasonic. The result is obvious:
Because of oxidation and dehydrogenation, the phase transition of hydrogenated titanium alloy cannot be directly measured by TGA or DTA. In this paper, a new method based on TGA for studying dehydrogenation process of titanium alloys is proposed, and the major process is as follows. The TGA curves of unhydrogenated and hydrogenated Ti600 are shown in Fig. 3, and they are obtained by using the following equations:
mC ¼
M ST 1 100% MS
mL ¼ mCunh mCh
ð1Þ ð2Þ
where MS is initial quality of the samples, MST is changing quality of the samples during heating process, mC is percentage change in quality of the samples during heating process, and mL is percentage loss in quality of the hydrogenated sample during heating process. The bond energy of b phase is smaller than a phase [16], so b phase can be oxidized more easily, meanwhile diffusion velocity of oxygen atoms in b phase is faster than in a phase. Then at the beginning of the a ? b phase transition, the increasing rate
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Fig. 6. TEM micrographs of Ti600 alloys deformed at 0.01 s1: (a) dislocation cells of unhydrogenated Ti600 at 850 °C, (b) DRX of unhydrogenated Ti600 at 950 °C, (c) the lamellar microstructure of 0.3 wt.% hydrogenated Ti600 at 850 °C, (d) titanium hydride d phase of 0.3 wt.% hydrogenated Ti600 at 850 °C, (e) SAED pattern of the FCC d titanium hydride with twins substructure at 850 °C, (f) the lamellar microstructure of 0.3 wt.% hydrogenated Ti600 at 950 °C, (g) titanium hydride d phase of 0.3 wt.% hydrogenated Ti600 at 950 °C.
of quality is accelerated because of oxidation. At the end of the a ? b phase transition, the increasing rate of quality reaches a
maximum. Above all, the a + b phase region of unhydrogenated Ti600 alloy ranges from about 790 to 1010 °C.
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The increasing quality of unhydrogenated Ti600 is caused by oxidation, while the change of quality is the result of oxidation and dehydrogenated for hydrogenated Ti600. So the difference of the mCunh and mCh is the percentage loss in quality, which caused by dehydrogenation, as shown in Fig. 3b. The transition of aH ? a + H2 occurs at the first stage, and then quality loss goes slowly because of low hydrogen solubility in aH. At the second stage, aH ? bH and b ?b + H2 occur about 700 °C, and then quality loss goes quickly because of high hydrogen solubility in bH. The maximum of quality loss comes at 900 °C meaning that aH ? bH has finished. At the third stage, oxidation is the main reason of quality change for the two samples, so the difference tends to zero quickly. Above all, the a + b phase region of 0.3 wt.% hydrogenated Ti600 alloy lies approximately between 700 and 900 °C. Based on Fig. 3, equilibrium phase transition of unhydrogenated and hydrogenated Ti600 is described in Fig. 4. The a + b phase region of unhydrogenated Ti600 ranges from 790 to 1010 °C. The XRD pattern of the alloys after quenching with deformed at a phase region of 850 °C is shown in Fig. 5a. There is only a phase as illustrated in Fig. 5a. It suggests b phase cannot be retained to room temperature by quenching from a + b phase region, and b ? a phase transition goes completely. It is a characteristic of near alpha titanium alloy. The a + b phases region of hydrogenated Ti600 lies approximately from 700 to 900 °C. The XRD pattern of the alloys after quenching with deformed at a + aH + b + bH phase region of 850 °C is shown in Fig. 5b with a, a0 , d and b phases. bH ? a0 (a) + d, bH ? b+d and b ? a phase transitions occur during the process of quenching. The metastable b phase of hydrogenated Ti600 only comes from bH ? b + d phase transition. It is hard to distinguish a0 phase and a by XRD, so the above conclusions concerned with a0 phase and a phase must be proved by TEM. Fig. 6 displays the effects of hydrogen and temperature on the TEM micrographs of the Ti600 alloys. After quenching, lamellar microstructure in the deformed hydrogenated alloy increases obviously, and lamellar microstructure comes from the bH phase quenching at high temperature. a phase dynamic recovery and recrystallization of unhydrogenated Ti600 alloy are shown in Fig. 6a and b, including dislocation cells and recrystallization grains. That suggests dynamic recovery of a phase is the main softening mechanism at 850 °C, and dynamic recrystallization of a phase is the main softening mechanism at 950 °C. The lamellar titanium hydride a phase of 0.3 wt.% hydrogenated Ti600 alloy, which comes from the quenching of bH phase at high temperature, is shown in Fig. 6c. The substructure of the lamellar titanium hydride phase is twin, as shown in Fig. 6d. Fig. 6e is the SAED pattern of the FCC a titanium hydride, which reveals the orientation rela 1 1 1). All of tionship of adjacent twinned regions to be (1 1 1)k( the samples are held at deformation temperature for 15 min to ensure the phase transition completely. Compared Fig. 6c with Fig. 6f, the width of lath titanium hydride phase expands with the rising of deformation temperature. With a comparison between Fig. 6d and g, it can be seen that the width of the twin substructure becomes bigger with the rising of deformation temperature. The lamellar microstructure consists of martensite a0 , metastable b and titanium hydride d. The dislocation substructure of martensite a0 and metastable b is caused by the deformation. The twin substructure of titanium hydride d is caused by bH ? a + d transformation during quenching process. The difference of the specific volume between the new phase d and the parent phase bH is enormous, and the
new phase d is limited in the growing process, which results in generating twins. The b phase of unhydrogenated Ti600 cannot be retained to room temperature by quenching, so it cannot be strengthened by heat treatment. Because of hydrogen addition, bH phase and d phase are retained by quenching, then the near alpha titanium alloy changes to a + b titanium alloy, which can be strengthened by heat treatment. Of course, the hydrogen will be removed by vacuum dehydrogenation heat treatment in case of Hydrogen embrittlement. 4. Conclusions (1) Hydrogen addition causes a phase softening and decreases b phase transformation temperature, but hydrogen causes b phase hardening. When the two contrary effects reach equilibrium, the hydrogen content is optimum. 0.3 wt.% is an optimal hydrogen content for Ti600 deformed at 850 °C. Compared with unhydrogenated Ti600 deformed at 950 °C, deformation temperature decreases by 100 °C and deformation resistance decreases by one third. Because of hydrogen addition, the main softening mechanism changes from dynamic recrystallization of a phase to dynamic recovery of b phase. (2) A new method based on TGA for studying phase transition of hydrogenated titanium alloy is proposed. With the study of weight gain by oxidation and weight loss by dehydrogenation, the a + b phase region of unhydrogenated and 0.3 wt.% hydrogenated Ti600 alloy are about from 790 to 1010 °C and from 700 to 900 °C. And the optimum temperature for dehydrogenation of titanium alloy is above 700 °C, which is the turning point of dehydrogenation rate. (3) The lamellar microstructure is gained through quenching after high temperature deformation in 0.3 wt.% hydrogenated Ti600, which consists of martensite a0 , metastable b and titanium hydride d, and the width of the lamellar become larger with the rising of deformation temperature.
Acknowledgments This work was supported by the National Natural Science Foundation of China (No. 50805037). References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16]
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