- Email: [email protected]
Thermal hydrogen desorption behavior of cathodically charged Ni3(Si,Ti) alloys
Journal of Alloys and Compounds 364 (2004) 214–220
Thermal hydrogen desorption behavior of cathodically charged Ni3(Si,Ti) alloys T. Nishiue, Y. Kaneno, H. Inoue, T. Takasugi∗ Department of Metallurgy and Materials Science, Graduate School of Engineering, Osaka Prefecture University, 1-1 Gakuen-cho, Sakai, Osaka 599-8531, Japan Received 31 March 2003; accepted 23 April 2003
Abstract Using thermal hydrogen desorption spectroscopy, the effects of lattice defects and boron-doping on the hydrogen desorption behavior in the cathodically charged L12 -type Ni3 (Si,Ti) compounds were investigated. Low- and high-temperature peaks were observed, and attributed to interstitial sites in L12 lattice and grain boundaries, respectively, as the trap sites of hydrogen. The low-temperature peak increased by deformation while the high-temperature peak was insensitive to deformation. The boron-doping to the Ni3 (Si,Ti) specimen has the effect of reducing the desorped hydrogen contents in the measured whole temperatures. As the holding time after hydrogen charging proceeds, the low-temperature peak rapidly decreased while the high-temperature peak remained constant. It is suggested that during holding time, hydrogen atoms in the interstitial sites in the L12 lattice migrate to the outside of the material while those at the grain boundaries remain at the same level. The hydrogen atoms in the interstitial sites in the L12 lattice in the deformed specimen can migrate through dislocations or with the help of excess vacancies more quickly than those in the recrystallized specimen. The results obtained in this study are discussed in relation to the environmental embrittlement of Ni3 (Si,Ti) compounds. © 2003 Elsevier B.V. All rights reserved. Keywords: Intermetallics; Point defects
1. Introduction Many intermetallic compounds including L12 -type compounds have been shown to be suffering from a so-called environmental embrittlement occurring in air at ambient temperature and resulting in reduction of tensile ductility and strength [1–4]. Environmental embrittlement of L12 -type intermetallic compounds including Ni3 Si and related alloys [5–7] has been considered to be caused by hydrogen released from moisture in the air, and to involve microprocesses such as the generation of hydrogen by decomposition of moisture on the alloy surface, absorption of hydrogen into the alloy interior, migration and condensation of hydrogen at grain boundaries in front of a propagating microcrack introduced from the alloy surface [1–4]. Eventually, grain boundary cohesion and associated plastic work are reduced by hydrogen condensation on grain boundary, and conse∗ Corresponding author. Tel.: +81-72-254-9314; fax: +81-72-257-9912. E-mail address: [email protected] (T. Takasugi).
0925-8388/$ – see front matter © 2003 Elsevier B.V. All rights reserved. doi:10.1016/S0925-8388(03)00505-X
quently intergranular fracture occurs, resulting in low ductility. Recent studies reveal that environmental embrittlement of the L12 -type Ni3 (Si,Ti) alloys is dependent on lattice defect structures [8,9] or microstructures [10–14] as well as alloy composition [6,7,15,16]. For instances, prestrained [8], shot-peened [9], fine-grained [14] and duplex [10–12] microstructures are effective in reducing environmental embrittlement of the L12 -type Ni3 (Si,Ti) alloys. Also, some substitutional (such as chromium and iron) [16] and interstitial (such as boron and carbon) [6,7,15] solutes are known to be strongly effective in reducing the moisture-induced embrittlement of the L12 -type Ni3 (Si,Ti) alloys. However, the detailed mechanisms are not fully understood due to lack of understanding of the hydrogen processes such as entry and desorption as well as the trap sites in the materials. Thermal hydrogen desorption spectroscopy (THDS) has been recently applied to many kinds of materials including steels. By this experimental technique, the state of hydrogen or hydrogen trap sites in the materials, and the correlation of the responsible hydrogen in materials with environmental embrittlement can be clarified. This technique is most
T. Nishiue et al. / Journal of Alloys and Compounds 364 (2004) 214–220
advantageous when applied to materials with a low hydrogen diffusion coefficient [17]. In the present study, using THDS, the effects of lattice defect structures (such as vacancies or dislocations introduced by deformation, and grain boundaries) as well as boron-doping on the hydrogen desorption behavior in the L12 -type Ni3 (Si,Ti) alloys, which have very attractive mechanical properties at high temperature as well as ambient temperature [18,19] and therefore is considered to be used as high-temperature structural materials, were investigated. In actual environmental embrittlement of the Ni3 (Si,Ti) alloys, the hydrogen content causing the embrittlement is assumed to be at a very low level. Accordingly, it is difficult to quantitatively evaluate the hydrogen behavior and properties by the THDS technique because the detected hydrogen level is far below the resolution limit. Therefore, cathodically charged Ni3 (Si,Ti) alloys were used in this study.
215
Fig. 1. Typical microstructure of a recrystallized Ni3 (Si,Ti) specimen with a fine grain size of 15 m.
2. Experimental Raw materials used in this study were 99.9 wt.% cobalt, 99.999 wt.% silicon, 99.9 wt.% sponge titanium and 99.9 wt.% boron. The Ni3 (Si,Ti) alloys with a nominal composition of Ni79.5 Si11 Ti9.5 (denoted by atomic pct) without and with 50 mass pct boron were used in this study. Boron-doping was carried out using a nickel–10 wt.% boron master alloy. Alloy button ingots with the dimensions of 15 × 25 × 50 mm were prepared by arc melting in argon gas atmosphere on a copper hearth using a nonconsumable tungsten electrode. Homogenization heat treatment was conducted in vacuum at 1323 K for 2 days, followed by furnace cooling. The alloy button ingots were rolled at 573 K in air and then annealed in vacuum at 1273 K for 5 h. This procedure was repeated several times until the desired thickness was obtained. In the final stage, the rolling was done at room temperature in air, and plates with a thickness of 0.3 mm were obtained as starting materials. Specimens with two kinds of microstructures, i.e. (1) fully annealed (i.e. recrystallized) microstructure, and (2) deformed microstructure were prepared. Specimens with a recrystallized microstructure were prepared by annealing in vacuum at 1273 K, using the starting material. Fig. 1 shows a typical microstructure of the fully annealed specimen observed by an optical microscope. Changing the annealing time, the recrystallized specimens with different grain sizes were prepared to investigate the effect of grain size on the hydrogen desorption behavior. Grain size was measured by a linear intercept method, using an optical microscope. Twin boundaries were excluded from counting. The THDS specimens were cut by an electrodischarge machine (EDM) from the variously processed materials mentioned above. The THDS sample size was about 4 × 15 × 0.3 mm3 . Cathodical hydrogen charging was conducted at room temperature for 6 h in a solution of 0.5 M H2 SO4 +0.05 g/l of NaAsO2 , at a constant current density of 6 mA cm−2 .
After hydrogen charging, the THDS specimens were rinsed in acetone using an ultrasonic cleaner, and then dried, followed by holding in air at room temperature for the desired hours (i.e. zero to 24 h). An X-ray diffraction (XRD) experiment using CuK␣ radiation was carried out to investigate the existence of the hydride after hydrogen charging. The THDS measurement was conducted using a gas chromatograph and was calibrated with a standard mixture of hydrogen and argon gas. High-purity argon was employed as the carrier gas, and the sampling time of the carrier gas to the gas chromatograph was in 5-min intervals. The THDS specimen was heated at a linear heating rate within a quartz tube in a furnace. The standard heating rate was 100 K h−1 . When evaluating the activation energy corresponding to the peak (i.e. the trap site) appearing in the THDS curve, heating rates different from the standard heating rate were also adopted. The hydrogen desorption rate was defined as the amount of hydrogen desorped in 1 min per gram of specimen.
3. Results The Ni3 (Si,Ti) specimens, which were not cathodically charged, did not show any levels in the hydrogen desorption curve at the whole temperatures measured because the desorped hydrogen level is far below the resolution limit of the analyzed system. When the Ni3 (Si,Ti) samples were cathodically charged, the hydrogen desorption level exceeded the resolution limit of the analyzed system and showed certain profiles as a function of temperature. However, XRD did not show any evidence of hydride formation on the specimen surface after hydrogen charging. Fig. 2 shows the hydrogen thermal desorption curves of the recrystallized specimen with a fine grain size (15 m). Here, the THDS measurements were conducted at three different heating rates immediately after hydrogen charging. Describing the result
216
T. Nishiue et al. / Journal of Alloys and Compounds 364 (2004) 214–220 Table 1 Total hydrogen content (massppm) desorped from the Ni3 (Si,Ti) alloys
Fig. 2. Hydrogen thermal desorption curves from the recrystallized Ni3 (Si,Ti) specimen. The THDS measurements were conducted at heating rates of 100, 200 and 400 K h−1 immediately after hydrogen charging.
of the specimen heated at a standard rate (i.e. 100 K h−1 ), there is a large peak at about 370 K and a small peak (or a shoulder) at about 480 K (in the following these two peaks will be described as the low- and high-temperature peaks, respectively). Furthermore, the low-temperature peak may be split into two subpeaks in many cases, and also accompanied by a recognizable shoulder in its higher part. With increasing heating rate, it appears that the hydrogen thermal desorption curve shifts toward higher temperature, and also the split low-temperature peak appears to change to a broad single peak. Fig. 3 shows the hydrogen thermal desorption curves of the recrystallized Ni3 (Si,Ti) specimen without and with boron-doping. The THDS measurements were conducted immediately after hydrogen charging. It is apparent that boron-doping to the recrystallized Ni3 (Si,Ti) specimen re-
Fig. 3. Hydrogen thermal desorption curves of the recrystallized Ni3 (Si,Ti) specimen without and with boron-doping. The THDS measurements were conducted immediately after hydrogen charging.
Alloy
Microstructure
Hydrogen content (massppm) As-charged
6h
24 h
Ni3 (Si,Ti)
Recrystallized Deformed
106 138
67 65
32 34
Boron-doped Ni3 (Si,Ti)
Recrystallized Deformed
48 105
42 26
19 49
sults in a reduced hydrogen content level in the hydrogen thermal desorption curve at almost all temperatures; the measured desorped hydrogen contents were 106 and 48 massppm in the recrystallized Ni3 (Si,Ti) specimen without and with boron-doping, respectively (Table 1). It is interesting that the high-temperature peak and the shoulder of the low-temperature peak become unambiguous by the boron-doping. Fig. 4 shows the hydrogen thermal desorption curve of the deformed Ni3 (Si,Ti) specimen, in comparison with that of the recrystallized Ni3 (Si,Ti) specimen. The prior grain sizes of the two Ni3 (Si,Ti) specimens were the same (i.e. 15 m). The THDS measurements were conducted immediately after hydrogen charging. Basically, the hydrogen thermal desorption curve of the deformed Ni3 (Si,Ti) specimen is similar to that of the recrystallized one. However, the deformation results in an increase in the level of the hydrogen thermal desorption curve at almost all temperatures; the measured desorped hydrogen contents were 106 and 138 massppm in the recrystallized and deformed Ni3 (Si,Ti) specimens, respectively (Table 1). Fig. 5 shows the hydrogen thermal desorption curves of the deformed Ni3 (Si,Ti) specimen without and with boron-doping. The THDS measurements were conducted immediately after hydrogen charging. It is found that the
Fig. 4. Comparison in the hydrogen thermal desorption curves between the recrystallized and deformed Ni3 (Si,Ti) specimens. The THDS measurements were conducted immediately after hydrogen charging.
T. Nishiue et al. / Journal of Alloys and Compounds 364 (2004) 214–220
Fig. 5. Hydrogen thermal desorption curves of the deformed Ni3 (Si,Ti) specimen without and with boron-doping. The THDS measurements were conducted immediately after hydrogen charging.
boron-doping to the deformed Ni3 (Si,Ti) specimen has the effect of reducing the desorped hydrogen content as observed in the recrystallized Ni3 (Si,Ti) specimen (Table 1). Figs. 6 and 7 show changes of the hydrogen thermal desorption curves with holding time at room temperature after hydrogen charging for the recrystallized and deformed Ni3 (Si,Ti) specimens, respectively. In both the Ni3 (Si,Ti) specimens, the levels of the hydrogen desorption curves decrease at whole temperatures with increasing holding time (Table 1). In these figures, it is evident that the low-temperature peak decreases much more rapidly than the high-temperature peak, and also their decreases are much more significant in the deformed Ni3 (Si,Ti) specimen than in the recrystallized one. A further interesting result is that the split low-temperature peaks change into broad single peaks by holding time after hydrogen charging.
Fig. 6. Change of the hydrogen thermal desorption curve with holding time after hydrogen charging for the recrystallized Ni3 (Si,Ti) specimen.
217
Fig. 7. Change of the hydrogen thermal desorption curve with holding time after hydrogen charging for the deformed Ni3 (Si,Ti) specimen.
Figs. 8 and 9 show changes of the hydrogen thermal desorption curves with holding time at room temperature after hydrogen charging for the recrystallized and deformed Ni3 (Si,Ti) specimens with boron-doping, respectively. Similar to the Ni3 (Si,Ti) specimens without boron-doping, the levels of the hydrogen desorption curve decrease at whole temperatures with increasing holding time, regardless of the recrystallized or deformed state (Table 1). Again, the low-temperature peaks decrease much more rapidly than the high-temperature peak, and their decrease is much more significant in the deformed Ni3 (Si,Ti) specimen than in the recrystallized one. Also, the split low-temperature peaks change into broad single peaks by holding time after hydrogen charging. To obtain the activation energies for the trap sites of hydrogen, the hydrogen thermal desorption curves were mea-
Fig. 8. Change of the hydrogen thermal desorption curve with holding time after hydrogen charging for the recrystallized Ni3 (Si,Ti) specimen with boron-doping.
218
T. Nishiue et al. / Journal of Alloys and Compounds 364 (2004) 214–220
Ni3 (Si,Ti) specimen without and with boron-doping were estimated to be about 41 and 35 kJ mol−1 , respectively.
4. Discussion
Fig. 9. Change of the hydrogen thermal desorption curve with holding time after hydrogen charging for the deformed Ni3 (Si,Ti) specimen with boron-doping.
sured at three different heating rates, e.g. as already shown in Fig. 2 for the case of the recrystallized Ni3 (Si,Ti) specimen. The THDS measurements were conducted immediately after hydrogen charging. An increase in the heating rate generally results in an increase in the peak temperature (TC ) in the hydrogen desorption curve. Here, the activation energy, Ea , for the trap sites corresponding to the low-temperature peak can be derived on the basis of the following equation [15]: (Ea α/RT2C ) = Aexp(−Ea /RTC )
(1)
where α is the heating rate, R is the gas constant, TC is a maximum temperature, and A is a constant. The Arrhenius plots of Eq. (1), i.e. ln(α/TC2 ) vs. 1/TC for the trap sites corresponding to the low-temperature peak are shown in Fig. 10. The apparent activation energies of the hydrogen desorption from the trap sites (or soluble position) in the recrystallized
Fig. 10. Arrhenius plot of ln (␣/TC 2 ) vs. 1/TC for the recrystallized Ni3 (Si,Ti) specimen without and with boron-doping.
XRD did not show any evidence of hydride formation on the specimen surface after hydrogen charging. Therefore, the peaks (or the shoulders) observed in the hydrogen thermal desorption curves are attributed to some trap sites where atomic (free) hydrogen atoms are bound. The trap sites of hydrogen corresponding to the low-temperature peak appearing at about 370 K are suggested to be the interstitial sites in L12 lattices. Similar peaks have been observed in the hydrogen desorped curve of the cathodically charged L12 -type Co3 Ti where the low-temperature peak was observed to occur at about 360 K [20]. The low-temperature peak was strongly dependent on deformation. This result means that vacancies or dislocations introduced by deformation contribute to the increase of the trap sites of hydrogen, and consequently enhance the low-temperature peak. It is well known in many metals and alloys that vacancies prefer to bind with hydrogen atoms, in other words, additional hydrogen atoms are introduced by the introduction of excess vacancies [21–23]. On the other hand, the trap sites corresponding to the high-temperature peak appearing at about 480 K are suggested to be ‘grain boundaries’. The high-temperature peak was not so much dependent on grain size although the THDS measurement was conducted on the recrystallized Ni3 (Si,Ti) specimens with two levels of grain sizes (i.e. 15 and 170 m). In the case of the cathodically charged Co3 Ti, the high-temperature peak accompanied with an obvious profile has been observed at about 590 K and clearly depended on grain size [20]. As the interstitial sites in L12 lattices, octahedral sites (O-site) and tetragonal sites (T-sites) are considered. The former sites have larger free volume than the latter sites [24–26]. Furthermore, two kinds of atomic environments are considered as the O-sites; one is called as the OI -site where all (i.e. six) of neighboring constituent atoms are Ni atoms while the other is called as the OII -site where four of neighboring constituent atoms are Ni atoms and two of those are Ti/Si atoms. Therefore, it is possible that the two subpeaks of the low-temperature peak correspond to the OI and OII -sites, which hydrogen atoms independently occupy. Here, it might be anticipated that hydrogen atoms prefer to occupy the OII -sites rather than the OI -sites because hydrogen atoms prefer to bind with Ti atoms than Ni atoms. However, the experimental results that the heights of the two subpeaks of the low-temperature peak are not so different may suggest that there are no preferential O-sites for the trap sites of hydrogen. Therefore, on an assumption that all of the desorped hydrogen atoms (corresponding to the low-temperature peak) occupy the O-sites in the region where hydrogen atoms diffused into during the charging are
T. Nishiue et al. / Journal of Alloys and Compounds 364 (2004) 214–220
219
Table 2 Fraction of the O-sites occupied by hydrogen atoms which were desorbed from trap sites corresponding to the low-temperature peak in the hydrogen thermal desorption curve of the Ni3 (Si, Ti) specimens. The calculation was conducted for the recrystallized and deformed specimens that were as-charged, and subsequently held for 6 h and for 24 h before the THDS measurement, respectively Alloy
Microstructure
Hydrogen content (massppm) As-charged
6h
24 h
Without boron-doping
Recrystallized Deformed
6.5 × 10−5 8.5 × 10−5
4.0 × 10−5 3.3 × 10−5
7.6 × 10−6 1.5 × 10−5
With boron-doping
Recrystallized Deformed
3.1 × 10−5 6.8 × 10−5
2.5 × 10−5 1.4 × 10−5
8.9 × 10−6 2.3 × 10−6
distributed, the values shown in Table 2 are calculated as the fraction of the O-sites occupied by hydrogen atoms. Here, the value of 1.2 × 10−10 m2 s−1 [27] was used as the bulk diffusion data of the hydrogen in Ni3 Al instead of Ni3 (Si,Ti) because the data of the latter compound is not available up to date. Also, as the amount of the desorped hydrogen corresponding to the low-temperature peak, the hydrogen contents desorped at temperatures between 300 and 450 K were taken into calculation. The calculation was conducted for the recrystallized and deformed Ni3 (Si,Ti) specimens without and with boron-doping that were as-charged, and subsequently held for 6 and 24 h, respectively. It is clearly shown that as the holding time after charging increases the concentration of the hydrogen atoms on the O-sites rapidly decreases. It was found in this study that boron-doping has the effect of reducing the hydrogen content desorped from the interstitial trap sites (i.e. the O-sites) in L12 lattices as well as the trap sites on grain boundaries. It is anticipated that boron atoms prefer to occupy the OI -sites rather than the OII -sites because boron atoms prefer to bind to Ni atoms rather than to Ti atoms. However, the fact that reduction of the two subpeaks (comprising the low-temperature peak) by boron-doping occurred to the same extent may suggest that there is no preferential O-site for the trap site of boron. It is speculated that Si atoms substituting for Ti atom sites interferes preferable occupation of boron atoms for the OII sites, and consequently there may be no energetical preference for boron occupation between the OI -sites and OII -sites. Regarding the trap sites of vacancies introduced by deformation, site competition between boron and hydrogen is expected because vacancies are attractive to both atoms. Furthermore, regarding the trap sites of grain boundaries, site competition between boron and hydrogen is expected because both have similar atomic sizes. Indeed, Auger electron spectroscopic (AES) experiment by which boron was shown to be highly enriched (by about a few hundred times) at grain boundaries of the Ni3 (Si,Ti) alloy [28], and also the auto radiographic experiment by which hydrogen was shown to be little detected on grain boundaries of the boron-doped Ni3 (Si,Ti) alloy [29], support this expectation. Thus, it is likely that boron-doping has the effect of reducing the hydrogen contents desorped from the trap sites in the Ni3 (Si,Ti) alloy, regardless of the recrystallized or deformed state.
For the kinetics of hydrogen atoms (i.e. the evolution of the trap sites of hydrogen atoms) in the Ni3 (Si,Ti) specimen, it was found that the hydrogen atoms trapped in the interstitial sites in L12 lattice primarily move the outside from the material interior while the hydrogen atoms trapped at grain boundaries remain the same level. This result suggests that the hydrogen atoms in the O-sites are very diffusible or shallowly trapped while the hydrogen atoms at grain boundaries are deeply trapped. The migration of hydrogen atoms from the O-sites to the outside was much faster in the deformed specimen than in the recrystallized specimen, as understood from the comparison between Figs. 6 and 7, or between Figs. 8 and 9. In the recrystallized specimen, the migration of hydrogen atoms from the O-sites to the outside is mainly due to bulk (lattice) diffusion while in the deformed specimens, that is assumed to be due to dislocation diffusion or bulk diffusion enhanced by excess vacancies. As a result, much faster migration of hydrogen atoms is expected in the deformed specimen than in the recrystallized specimen. It was shown that grain boundaries in Ni3 (Si,Ti) are still effective in trapping hydrogen atoms even after a longer holding time after hydrogen charging. It is accordingly suggested that hydrogen atoms enriched on grain boundaries are species causing grain boundary fracture and low ductility, resulting in the environmental embrittlement in the Ni3 (Si,Ti) alloys [5,6]. Also, it was found that vacancies or dislocations introduced by deformation similarly act as the trap sites of hydrogen atoms and therefore are effective in reducing the environmental embrittlement through suppressing the enrichment of hydrogen atoms on grain boundaries. This result is consistent with the experimental findings that the prestrained [8] or shotpeened [9] L12 -type Ni3 (Si,Ti) alloys are less sensitive to environmental embrittlement. However, since the trap sites (i.e. the O-sites) of hydrogen atoms in L12 lattice appear to be not so energetically deep, hydrogen atoms eventually desorb the outside of the material or move to grain boundaries. Consequently, if the different types of lattice defects, specific solutes or the second-phase dispersions that are energetically deeper as the trap site of hydrogen atoms would be introduced in L12 lattice, it is possible that the environmental embrittlement of L12 -type Ni3 (Si,Ti) alloys could be suppressed. Indeed, it was shown that the
220
T. Nishiue et al. / Journal of Alloys and Compounds 364 (2004) 214–220
introduction of specific second-phase dispersion (and its interface with the matrix) is effective in reducing the environmental embrittlement of the L12 -type Ni3 (Si,Ti) alloys [10–12]. It was shown that boron-doping has the effect of reducing particularly the hydrogen content trapped on grain boundaries. This result is consistent with the experimental result that the boron-doped Ni3 (Si,Ti) alloy is insensitive to environmental embrittlement [6,7,15]. It has been suggested by the present authors that boron atoms enrich on grain boundaries of Ni3 (Si,Ti) and site-compete with hydrogen atoms. Consequently, the hydrogen content at grain boundaries does not reach the critical content above which the intergranular fracture occurs.
5. Conclusions Using THDS, the effects of lattice defects (such as dislocations, vacancies and grain boundaries) and boron-doping on the desorption behavior of hydrogen in cathodically charged Ni3 (Si,Ti) compounds were investigated. The following results were obtained. 1. Low- and high-temperature peaks were observed in the hydrogen thermal desorption curves of the Ni3 (Si,Ti) specimens, and attributed to the interstitial sites in L12 lattice and grain boundaries as the trap sites of hydrogen, respectively. 2. The low-temperature peak increased by deformation while the high-temperature peak was insensitive to deformation. 3. The boron-doping of the Ni3 (Si,Ti) specimen has the effect of reducing the desorped hydrogen contents at all temperatures. 4. As the holding time at room temperature after hydrogen charging increases, the low-temperature peak rapidly decreased while the high-temperature peak remained constant. This change was faster and more significant in the deformed specimen than in the recrystallized specimen, irrespective of boron-doping. 5. It was suggested that hydrogen atoms in the interstitial sites in L12 lattice migrate to the outside of the material while those at grain boundaries remain trapped. The hydrogen atoms in the deformed specimen can migrate through dislocations or with the help of excess vacancies more quickly than those in the recrystallized specimen.
6. The results obtained in this study (THDS method) were discussed in relation to the environmental embrittlement of Ni3 (Si,Ti) compounds.
References [1] G.L. Chen, C.T. Liu, Int. Mater. Rev. 46 (2001) 253. [2] T. Takasugi, in: N. Stoloff, D.J. Duquette, A.F. Giamei (Eds.), Critical Issues in the Development of High Temperature Structural Materials, TMS, USA, 1993, p. 21. [3] C.T. Liu, in: O. Izumi (Eds.), The 6th Int. Symp. Intermetallic Compounds—Structure and Mechanical Properties, JIM, Japan, 1991, p. 703. [4] C.T. Liu, E.P. George, in: C. Deevi, V.N. Sikka, P.J. Maziasz, R.W. Cahn (Eds.), Int. Symp. on Nickel and Iron Aluminide; Processing, Properties and Applications, ASM, USA, 1997, p. 21. [5] T. Takasugi, O. Izumi, M. Yoshida, J. Mater. Sci. 26 (1991) 1173. [6] T. Takasugi, H. Suenaga, O. Izumi, J. Mater. Sci. 26 (1991) 1179. [7] T. Takasugi, M. Yoshida, J. Mater. Sci. 26 (1991) 3032. [8] C.L. Ma, T. Takasugi, S. Hanada, Script. Metall. Mater. 34 (1996) 1633. [9] C.L. Ma, T. Takasugi, S. Hanada, Script. Metall. Mater. 34 (1996) 1131. [10] T. Takasugi, S. Hanada, MRS Symp. Proc. Publ. 552 (1999) KK6.5.1. [11] T. Takasugi, M. Wada, Y. Kaneno, H. Inoue, Mater. Sci. Eng. A 329–331C (2002) 523. [12] T. Takasugi, ISIJ Intern. 43 (2003) 564. [13] T. Takasugi, T. Tsuyumu, Y. Kaneno, H. Inoue, J. Mater. Res. 15 (2000) 1881. [14] Y. Kaneno, M. Wada, H. Inoue, T. Takasugi, Mater. Transaction 42 (2001) 418. [15] C.L. Ma, T. Takasugi, S. Hanada, Mater. Trans. JIM 36 (1995) 30. [16] C.L. Ma, T. Takasugi, S. Hanada, Script. Metall. Mater. 32 (1995) 1025. [17] A. Turnbull, R.B. Hutchings, D.H. Ferriss, Mater. Sci. Eng. A238 (1997) 317. [18] T. Takasugi, M. Nagahsima, O. Izumi, Acta Metall. Mater. 37 (1989) 3425. [19] T. Takasugi, M. Nagahsima, O. Izumi, Acta Metall. Mater. 38 (1990) 747. [20] T. Nishiue, Y. Kaneno, H. Inoue, T. Takasugi, Intermetallics, in press. [21] Y. Hukai, N. Okuma, J. Appl. Phys. 32 (1993) 1256. [22] M. Iwamoto, Y. Hukai, Mater. Trans. JIM 40 (1999) 606. [23] H.K. Birnbaum, C. Buckley, F. Zeides, E. Sioris, P. Rozanak, S. Spooner, J.S. Lin, J. Alloys Comp. 253–254 (1997) 260. [24] H.D. Carstanjen, Phys. Stat. Sol. (a) 59 (1980) 11. [25] L. Yang, R.B. McLellan, Script. Metall. Mater. 42 (1994) 3993. [26] L. Yang, R.B. McLellan, Script. Metall. Mater. 32 (1994) 779. [27] L. Yang, R.B. MacLellan, J. Phys. Chem. Solids 57 (1996) 233. [28] T. Takasugi, K. Hono, S. Suzuki, S. Hanada, T. Sakurai, Script. Metall. Mater. 29 (1993) 1587. [29] H. Saitoh, M. Endoh, T. Aoyama, T. Misawa, T. Ohnishi, C.L. Ma, T. Takasugi, J. Mater. Sci. Lett. 15 (1996) 1297.
Thermal hydrogen desorption behavior of cathodically charged Ni3(Si,Ti) alloys T. Nishiue, Y. Kaneno, H. Inoue, T. Takasugi∗ Department of Metallurgy and Materials Science, Graduate School of Engineering, Osaka Prefecture University, 1-1 Gakuen-cho, Sakai, Osaka 599-8531, Japan Received 31 March 2003; accepted 23 April 2003
Abstract Using thermal hydrogen desorption spectroscopy, the effects of lattice defects and boron-doping on the hydrogen desorption behavior in the cathodically charged L12 -type Ni3 (Si,Ti) compounds were investigated. Low- and high-temperature peaks were observed, and attributed to interstitial sites in L12 lattice and grain boundaries, respectively, as the trap sites of hydrogen. The low-temperature peak increased by deformation while the high-temperature peak was insensitive to deformation. The boron-doping to the Ni3 (Si,Ti) specimen has the effect of reducing the desorped hydrogen contents in the measured whole temperatures. As the holding time after hydrogen charging proceeds, the low-temperature peak rapidly decreased while the high-temperature peak remained constant. It is suggested that during holding time, hydrogen atoms in the interstitial sites in the L12 lattice migrate to the outside of the material while those at the grain boundaries remain at the same level. The hydrogen atoms in the interstitial sites in the L12 lattice in the deformed specimen can migrate through dislocations or with the help of excess vacancies more quickly than those in the recrystallized specimen. The results obtained in this study are discussed in relation to the environmental embrittlement of Ni3 (Si,Ti) compounds. © 2003 Elsevier B.V. All rights reserved. Keywords: Intermetallics; Point defects
1. Introduction Many intermetallic compounds including L12 -type compounds have been shown to be suffering from a so-called environmental embrittlement occurring in air at ambient temperature and resulting in reduction of tensile ductility and strength [1–4]. Environmental embrittlement of L12 -type intermetallic compounds including Ni3 Si and related alloys [5–7] has been considered to be caused by hydrogen released from moisture in the air, and to involve microprocesses such as the generation of hydrogen by decomposition of moisture on the alloy surface, absorption of hydrogen into the alloy interior, migration and condensation of hydrogen at grain boundaries in front of a propagating microcrack introduced from the alloy surface [1–4]. Eventually, grain boundary cohesion and associated plastic work are reduced by hydrogen condensation on grain boundary, and conse∗ Corresponding author. Tel.: +81-72-254-9314; fax: +81-72-257-9912. E-mail address: [email protected] (T. Takasugi).
0925-8388/$ – see front matter © 2003 Elsevier B.V. All rights reserved. doi:10.1016/S0925-8388(03)00505-X
quently intergranular fracture occurs, resulting in low ductility. Recent studies reveal that environmental embrittlement of the L12 -type Ni3 (Si,Ti) alloys is dependent on lattice defect structures [8,9] or microstructures [10–14] as well as alloy composition [6,7,15,16]. For instances, prestrained [8], shot-peened [9], fine-grained [14] and duplex [10–12] microstructures are effective in reducing environmental embrittlement of the L12 -type Ni3 (Si,Ti) alloys. Also, some substitutional (such as chromium and iron) [16] and interstitial (such as boron and carbon) [6,7,15] solutes are known to be strongly effective in reducing the moisture-induced embrittlement of the L12 -type Ni3 (Si,Ti) alloys. However, the detailed mechanisms are not fully understood due to lack of understanding of the hydrogen processes such as entry and desorption as well as the trap sites in the materials. Thermal hydrogen desorption spectroscopy (THDS) has been recently applied to many kinds of materials including steels. By this experimental technique, the state of hydrogen or hydrogen trap sites in the materials, and the correlation of the responsible hydrogen in materials with environmental embrittlement can be clarified. This technique is most
T. Nishiue et al. / Journal of Alloys and Compounds 364 (2004) 214–220
advantageous when applied to materials with a low hydrogen diffusion coefficient [17]. In the present study, using THDS, the effects of lattice defect structures (such as vacancies or dislocations introduced by deformation, and grain boundaries) as well as boron-doping on the hydrogen desorption behavior in the L12 -type Ni3 (Si,Ti) alloys, which have very attractive mechanical properties at high temperature as well as ambient temperature [18,19] and therefore is considered to be used as high-temperature structural materials, were investigated. In actual environmental embrittlement of the Ni3 (Si,Ti) alloys, the hydrogen content causing the embrittlement is assumed to be at a very low level. Accordingly, it is difficult to quantitatively evaluate the hydrogen behavior and properties by the THDS technique because the detected hydrogen level is far below the resolution limit. Therefore, cathodically charged Ni3 (Si,Ti) alloys were used in this study.
215
Fig. 1. Typical microstructure of a recrystallized Ni3 (Si,Ti) specimen with a fine grain size of 15 m.
2. Experimental Raw materials used in this study were 99.9 wt.% cobalt, 99.999 wt.% silicon, 99.9 wt.% sponge titanium and 99.9 wt.% boron. The Ni3 (Si,Ti) alloys with a nominal composition of Ni79.5 Si11 Ti9.5 (denoted by atomic pct) without and with 50 mass pct boron were used in this study. Boron-doping was carried out using a nickel–10 wt.% boron master alloy. Alloy button ingots with the dimensions of 15 × 25 × 50 mm were prepared by arc melting in argon gas atmosphere on a copper hearth using a nonconsumable tungsten electrode. Homogenization heat treatment was conducted in vacuum at 1323 K for 2 days, followed by furnace cooling. The alloy button ingots were rolled at 573 K in air and then annealed in vacuum at 1273 K for 5 h. This procedure was repeated several times until the desired thickness was obtained. In the final stage, the rolling was done at room temperature in air, and plates with a thickness of 0.3 mm were obtained as starting materials. Specimens with two kinds of microstructures, i.e. (1) fully annealed (i.e. recrystallized) microstructure, and (2) deformed microstructure were prepared. Specimens with a recrystallized microstructure were prepared by annealing in vacuum at 1273 K, using the starting material. Fig. 1 shows a typical microstructure of the fully annealed specimen observed by an optical microscope. Changing the annealing time, the recrystallized specimens with different grain sizes were prepared to investigate the effect of grain size on the hydrogen desorption behavior. Grain size was measured by a linear intercept method, using an optical microscope. Twin boundaries were excluded from counting. The THDS specimens were cut by an electrodischarge machine (EDM) from the variously processed materials mentioned above. The THDS sample size was about 4 × 15 × 0.3 mm3 . Cathodical hydrogen charging was conducted at room temperature for 6 h in a solution of 0.5 M H2 SO4 +0.05 g/l of NaAsO2 , at a constant current density of 6 mA cm−2 .
After hydrogen charging, the THDS specimens were rinsed in acetone using an ultrasonic cleaner, and then dried, followed by holding in air at room temperature for the desired hours (i.e. zero to 24 h). An X-ray diffraction (XRD) experiment using CuK␣ radiation was carried out to investigate the existence of the hydride after hydrogen charging. The THDS measurement was conducted using a gas chromatograph and was calibrated with a standard mixture of hydrogen and argon gas. High-purity argon was employed as the carrier gas, and the sampling time of the carrier gas to the gas chromatograph was in 5-min intervals. The THDS specimen was heated at a linear heating rate within a quartz tube in a furnace. The standard heating rate was 100 K h−1 . When evaluating the activation energy corresponding to the peak (i.e. the trap site) appearing in the THDS curve, heating rates different from the standard heating rate were also adopted. The hydrogen desorption rate was defined as the amount of hydrogen desorped in 1 min per gram of specimen.
3. Results The Ni3 (Si,Ti) specimens, which were not cathodically charged, did not show any levels in the hydrogen desorption curve at the whole temperatures measured because the desorped hydrogen level is far below the resolution limit of the analyzed system. When the Ni3 (Si,Ti) samples were cathodically charged, the hydrogen desorption level exceeded the resolution limit of the analyzed system and showed certain profiles as a function of temperature. However, XRD did not show any evidence of hydride formation on the specimen surface after hydrogen charging. Fig. 2 shows the hydrogen thermal desorption curves of the recrystallized specimen with a fine grain size (15 m). Here, the THDS measurements were conducted at three different heating rates immediately after hydrogen charging. Describing the result
216
T. Nishiue et al. / Journal of Alloys and Compounds 364 (2004) 214–220 Table 1 Total hydrogen content (massppm) desorped from the Ni3 (Si,Ti) alloys
Fig. 2. Hydrogen thermal desorption curves from the recrystallized Ni3 (Si,Ti) specimen. The THDS measurements were conducted at heating rates of 100, 200 and 400 K h−1 immediately after hydrogen charging.
of the specimen heated at a standard rate (i.e. 100 K h−1 ), there is a large peak at about 370 K and a small peak (or a shoulder) at about 480 K (in the following these two peaks will be described as the low- and high-temperature peaks, respectively). Furthermore, the low-temperature peak may be split into two subpeaks in many cases, and also accompanied by a recognizable shoulder in its higher part. With increasing heating rate, it appears that the hydrogen thermal desorption curve shifts toward higher temperature, and also the split low-temperature peak appears to change to a broad single peak. Fig. 3 shows the hydrogen thermal desorption curves of the recrystallized Ni3 (Si,Ti) specimen without and with boron-doping. The THDS measurements were conducted immediately after hydrogen charging. It is apparent that boron-doping to the recrystallized Ni3 (Si,Ti) specimen re-
Fig. 3. Hydrogen thermal desorption curves of the recrystallized Ni3 (Si,Ti) specimen without and with boron-doping. The THDS measurements were conducted immediately after hydrogen charging.
Alloy
Microstructure
Hydrogen content (massppm) As-charged
6h
24 h
Ni3 (Si,Ti)
Recrystallized Deformed
106 138
67 65
32 34
Boron-doped Ni3 (Si,Ti)
Recrystallized Deformed
48 105
42 26
19 49
sults in a reduced hydrogen content level in the hydrogen thermal desorption curve at almost all temperatures; the measured desorped hydrogen contents were 106 and 48 massppm in the recrystallized Ni3 (Si,Ti) specimen without and with boron-doping, respectively (Table 1). It is interesting that the high-temperature peak and the shoulder of the low-temperature peak become unambiguous by the boron-doping. Fig. 4 shows the hydrogen thermal desorption curve of the deformed Ni3 (Si,Ti) specimen, in comparison with that of the recrystallized Ni3 (Si,Ti) specimen. The prior grain sizes of the two Ni3 (Si,Ti) specimens were the same (i.e. 15 m). The THDS measurements were conducted immediately after hydrogen charging. Basically, the hydrogen thermal desorption curve of the deformed Ni3 (Si,Ti) specimen is similar to that of the recrystallized one. However, the deformation results in an increase in the level of the hydrogen thermal desorption curve at almost all temperatures; the measured desorped hydrogen contents were 106 and 138 massppm in the recrystallized and deformed Ni3 (Si,Ti) specimens, respectively (Table 1). Fig. 5 shows the hydrogen thermal desorption curves of the deformed Ni3 (Si,Ti) specimen without and with boron-doping. The THDS measurements were conducted immediately after hydrogen charging. It is found that the
Fig. 4. Comparison in the hydrogen thermal desorption curves between the recrystallized and deformed Ni3 (Si,Ti) specimens. The THDS measurements were conducted immediately after hydrogen charging.
T. Nishiue et al. / Journal of Alloys and Compounds 364 (2004) 214–220
Fig. 5. Hydrogen thermal desorption curves of the deformed Ni3 (Si,Ti) specimen without and with boron-doping. The THDS measurements were conducted immediately after hydrogen charging.
boron-doping to the deformed Ni3 (Si,Ti) specimen has the effect of reducing the desorped hydrogen content as observed in the recrystallized Ni3 (Si,Ti) specimen (Table 1). Figs. 6 and 7 show changes of the hydrogen thermal desorption curves with holding time at room temperature after hydrogen charging for the recrystallized and deformed Ni3 (Si,Ti) specimens, respectively. In both the Ni3 (Si,Ti) specimens, the levels of the hydrogen desorption curves decrease at whole temperatures with increasing holding time (Table 1). In these figures, it is evident that the low-temperature peak decreases much more rapidly than the high-temperature peak, and also their decreases are much more significant in the deformed Ni3 (Si,Ti) specimen than in the recrystallized one. A further interesting result is that the split low-temperature peaks change into broad single peaks by holding time after hydrogen charging.
Fig. 6. Change of the hydrogen thermal desorption curve with holding time after hydrogen charging for the recrystallized Ni3 (Si,Ti) specimen.
217
Fig. 7. Change of the hydrogen thermal desorption curve with holding time after hydrogen charging for the deformed Ni3 (Si,Ti) specimen.
Figs. 8 and 9 show changes of the hydrogen thermal desorption curves with holding time at room temperature after hydrogen charging for the recrystallized and deformed Ni3 (Si,Ti) specimens with boron-doping, respectively. Similar to the Ni3 (Si,Ti) specimens without boron-doping, the levels of the hydrogen desorption curve decrease at whole temperatures with increasing holding time, regardless of the recrystallized or deformed state (Table 1). Again, the low-temperature peaks decrease much more rapidly than the high-temperature peak, and their decrease is much more significant in the deformed Ni3 (Si,Ti) specimen than in the recrystallized one. Also, the split low-temperature peaks change into broad single peaks by holding time after hydrogen charging. To obtain the activation energies for the trap sites of hydrogen, the hydrogen thermal desorption curves were mea-
Fig. 8. Change of the hydrogen thermal desorption curve with holding time after hydrogen charging for the recrystallized Ni3 (Si,Ti) specimen with boron-doping.
218
T. Nishiue et al. / Journal of Alloys and Compounds 364 (2004) 214–220
Ni3 (Si,Ti) specimen without and with boron-doping were estimated to be about 41 and 35 kJ mol−1 , respectively.
4. Discussion
Fig. 9. Change of the hydrogen thermal desorption curve with holding time after hydrogen charging for the deformed Ni3 (Si,Ti) specimen with boron-doping.
sured at three different heating rates, e.g. as already shown in Fig. 2 for the case of the recrystallized Ni3 (Si,Ti) specimen. The THDS measurements were conducted immediately after hydrogen charging. An increase in the heating rate generally results in an increase in the peak temperature (TC ) in the hydrogen desorption curve. Here, the activation energy, Ea , for the trap sites corresponding to the low-temperature peak can be derived on the basis of the following equation [15]: (Ea α/RT2C ) = Aexp(−Ea /RTC )
(1)
where α is the heating rate, R is the gas constant, TC is a maximum temperature, and A is a constant. The Arrhenius plots of Eq. (1), i.e. ln(α/TC2 ) vs. 1/TC for the trap sites corresponding to the low-temperature peak are shown in Fig. 10. The apparent activation energies of the hydrogen desorption from the trap sites (or soluble position) in the recrystallized
Fig. 10. Arrhenius plot of ln (␣/TC 2 ) vs. 1/TC for the recrystallized Ni3 (Si,Ti) specimen without and with boron-doping.
XRD did not show any evidence of hydride formation on the specimen surface after hydrogen charging. Therefore, the peaks (or the shoulders) observed in the hydrogen thermal desorption curves are attributed to some trap sites where atomic (free) hydrogen atoms are bound. The trap sites of hydrogen corresponding to the low-temperature peak appearing at about 370 K are suggested to be the interstitial sites in L12 lattices. Similar peaks have been observed in the hydrogen desorped curve of the cathodically charged L12 -type Co3 Ti where the low-temperature peak was observed to occur at about 360 K [20]. The low-temperature peak was strongly dependent on deformation. This result means that vacancies or dislocations introduced by deformation contribute to the increase of the trap sites of hydrogen, and consequently enhance the low-temperature peak. It is well known in many metals and alloys that vacancies prefer to bind with hydrogen atoms, in other words, additional hydrogen atoms are introduced by the introduction of excess vacancies [21–23]. On the other hand, the trap sites corresponding to the high-temperature peak appearing at about 480 K are suggested to be ‘grain boundaries’. The high-temperature peak was not so much dependent on grain size although the THDS measurement was conducted on the recrystallized Ni3 (Si,Ti) specimens with two levels of grain sizes (i.e. 15 and 170 m). In the case of the cathodically charged Co3 Ti, the high-temperature peak accompanied with an obvious profile has been observed at about 590 K and clearly depended on grain size [20]. As the interstitial sites in L12 lattices, octahedral sites (O-site) and tetragonal sites (T-sites) are considered. The former sites have larger free volume than the latter sites [24–26]. Furthermore, two kinds of atomic environments are considered as the O-sites; one is called as the OI -site where all (i.e. six) of neighboring constituent atoms are Ni atoms while the other is called as the OII -site where four of neighboring constituent atoms are Ni atoms and two of those are Ti/Si atoms. Therefore, it is possible that the two subpeaks of the low-temperature peak correspond to the OI and OII -sites, which hydrogen atoms independently occupy. Here, it might be anticipated that hydrogen atoms prefer to occupy the OII -sites rather than the OI -sites because hydrogen atoms prefer to bind with Ti atoms than Ni atoms. However, the experimental results that the heights of the two subpeaks of the low-temperature peak are not so different may suggest that there are no preferential O-sites for the trap sites of hydrogen. Therefore, on an assumption that all of the desorped hydrogen atoms (corresponding to the low-temperature peak) occupy the O-sites in the region where hydrogen atoms diffused into during the charging are
T. Nishiue et al. / Journal of Alloys and Compounds 364 (2004) 214–220
219
Table 2 Fraction of the O-sites occupied by hydrogen atoms which were desorbed from trap sites corresponding to the low-temperature peak in the hydrogen thermal desorption curve of the Ni3 (Si, Ti) specimens. The calculation was conducted for the recrystallized and deformed specimens that were as-charged, and subsequently held for 6 h and for 24 h before the THDS measurement, respectively Alloy
Microstructure
Hydrogen content (massppm) As-charged
6h
24 h
Without boron-doping
Recrystallized Deformed
6.5 × 10−5 8.5 × 10−5
4.0 × 10−5 3.3 × 10−5
7.6 × 10−6 1.5 × 10−5
With boron-doping
Recrystallized Deformed
3.1 × 10−5 6.8 × 10−5
2.5 × 10−5 1.4 × 10−5
8.9 × 10−6 2.3 × 10−6
distributed, the values shown in Table 2 are calculated as the fraction of the O-sites occupied by hydrogen atoms. Here, the value of 1.2 × 10−10 m2 s−1 [27] was used as the bulk diffusion data of the hydrogen in Ni3 Al instead of Ni3 (Si,Ti) because the data of the latter compound is not available up to date. Also, as the amount of the desorped hydrogen corresponding to the low-temperature peak, the hydrogen contents desorped at temperatures between 300 and 450 K were taken into calculation. The calculation was conducted for the recrystallized and deformed Ni3 (Si,Ti) specimens without and with boron-doping that were as-charged, and subsequently held for 6 and 24 h, respectively. It is clearly shown that as the holding time after charging increases the concentration of the hydrogen atoms on the O-sites rapidly decreases. It was found in this study that boron-doping has the effect of reducing the hydrogen content desorped from the interstitial trap sites (i.e. the O-sites) in L12 lattices as well as the trap sites on grain boundaries. It is anticipated that boron atoms prefer to occupy the OI -sites rather than the OII -sites because boron atoms prefer to bind to Ni atoms rather than to Ti atoms. However, the fact that reduction of the two subpeaks (comprising the low-temperature peak) by boron-doping occurred to the same extent may suggest that there is no preferential O-site for the trap site of boron. It is speculated that Si atoms substituting for Ti atom sites interferes preferable occupation of boron atoms for the OII sites, and consequently there may be no energetical preference for boron occupation between the OI -sites and OII -sites. Regarding the trap sites of vacancies introduced by deformation, site competition between boron and hydrogen is expected because vacancies are attractive to both atoms. Furthermore, regarding the trap sites of grain boundaries, site competition between boron and hydrogen is expected because both have similar atomic sizes. Indeed, Auger electron spectroscopic (AES) experiment by which boron was shown to be highly enriched (by about a few hundred times) at grain boundaries of the Ni3 (Si,Ti) alloy [28], and also the auto radiographic experiment by which hydrogen was shown to be little detected on grain boundaries of the boron-doped Ni3 (Si,Ti) alloy [29], support this expectation. Thus, it is likely that boron-doping has the effect of reducing the hydrogen contents desorped from the trap sites in the Ni3 (Si,Ti) alloy, regardless of the recrystallized or deformed state.
For the kinetics of hydrogen atoms (i.e. the evolution of the trap sites of hydrogen atoms) in the Ni3 (Si,Ti) specimen, it was found that the hydrogen atoms trapped in the interstitial sites in L12 lattice primarily move the outside from the material interior while the hydrogen atoms trapped at grain boundaries remain the same level. This result suggests that the hydrogen atoms in the O-sites are very diffusible or shallowly trapped while the hydrogen atoms at grain boundaries are deeply trapped. The migration of hydrogen atoms from the O-sites to the outside was much faster in the deformed specimen than in the recrystallized specimen, as understood from the comparison between Figs. 6 and 7, or between Figs. 8 and 9. In the recrystallized specimen, the migration of hydrogen atoms from the O-sites to the outside is mainly due to bulk (lattice) diffusion while in the deformed specimens, that is assumed to be due to dislocation diffusion or bulk diffusion enhanced by excess vacancies. As a result, much faster migration of hydrogen atoms is expected in the deformed specimen than in the recrystallized specimen. It was shown that grain boundaries in Ni3 (Si,Ti) are still effective in trapping hydrogen atoms even after a longer holding time after hydrogen charging. It is accordingly suggested that hydrogen atoms enriched on grain boundaries are species causing grain boundary fracture and low ductility, resulting in the environmental embrittlement in the Ni3 (Si,Ti) alloys [5,6]. Also, it was found that vacancies or dislocations introduced by deformation similarly act as the trap sites of hydrogen atoms and therefore are effective in reducing the environmental embrittlement through suppressing the enrichment of hydrogen atoms on grain boundaries. This result is consistent with the experimental findings that the prestrained [8] or shotpeened [9] L12 -type Ni3 (Si,Ti) alloys are less sensitive to environmental embrittlement. However, since the trap sites (i.e. the O-sites) of hydrogen atoms in L12 lattice appear to be not so energetically deep, hydrogen atoms eventually desorb the outside of the material or move to grain boundaries. Consequently, if the different types of lattice defects, specific solutes or the second-phase dispersions that are energetically deeper as the trap site of hydrogen atoms would be introduced in L12 lattice, it is possible that the environmental embrittlement of L12 -type Ni3 (Si,Ti) alloys could be suppressed. Indeed, it was shown that the
220
T. Nishiue et al. / Journal of Alloys and Compounds 364 (2004) 214–220
introduction of specific second-phase dispersion (and its interface with the matrix) is effective in reducing the environmental embrittlement of the L12 -type Ni3 (Si,Ti) alloys [10–12]. It was shown that boron-doping has the effect of reducing particularly the hydrogen content trapped on grain boundaries. This result is consistent with the experimental result that the boron-doped Ni3 (Si,Ti) alloy is insensitive to environmental embrittlement [6,7,15]. It has been suggested by the present authors that boron atoms enrich on grain boundaries of Ni3 (Si,Ti) and site-compete with hydrogen atoms. Consequently, the hydrogen content at grain boundaries does not reach the critical content above which the intergranular fracture occurs.
5. Conclusions Using THDS, the effects of lattice defects (such as dislocations, vacancies and grain boundaries) and boron-doping on the desorption behavior of hydrogen in cathodically charged Ni3 (Si,Ti) compounds were investigated. The following results were obtained. 1. Low- and high-temperature peaks were observed in the hydrogen thermal desorption curves of the Ni3 (Si,Ti) specimens, and attributed to the interstitial sites in L12 lattice and grain boundaries as the trap sites of hydrogen, respectively. 2. The low-temperature peak increased by deformation while the high-temperature peak was insensitive to deformation. 3. The boron-doping of the Ni3 (Si,Ti) specimen has the effect of reducing the desorped hydrogen contents at all temperatures. 4. As the holding time at room temperature after hydrogen charging increases, the low-temperature peak rapidly decreased while the high-temperature peak remained constant. This change was faster and more significant in the deformed specimen than in the recrystallized specimen, irrespective of boron-doping. 5. It was suggested that hydrogen atoms in the interstitial sites in L12 lattice migrate to the outside of the material while those at grain boundaries remain trapped. The hydrogen atoms in the deformed specimen can migrate through dislocations or with the help of excess vacancies more quickly than those in the recrystallized specimen.
6. The results obtained in this study (THDS method) were discussed in relation to the environmental embrittlement of Ni3 (Si,Ti) compounds.
References [1] G.L. Chen, C.T. Liu, Int. Mater. Rev. 46 (2001) 253. [2] T. Takasugi, in: N. Stoloff, D.J. Duquette, A.F. Giamei (Eds.), Critical Issues in the Development of High Temperature Structural Materials, TMS, USA, 1993, p. 21. [3] C.T. Liu, in: O. Izumi (Eds.), The 6th Int. Symp. Intermetallic Compounds—Structure and Mechanical Properties, JIM, Japan, 1991, p. 703. [4] C.T. Liu, E.P. George, in: C. Deevi, V.N. Sikka, P.J. Maziasz, R.W. Cahn (Eds.), Int. Symp. on Nickel and Iron Aluminide; Processing, Properties and Applications, ASM, USA, 1997, p. 21. [5] T. Takasugi, O. Izumi, M. Yoshida, J. Mater. Sci. 26 (1991) 1173. [6] T. Takasugi, H. Suenaga, O. Izumi, J. Mater. Sci. 26 (1991) 1179. [7] T. Takasugi, M. Yoshida, J. Mater. Sci. 26 (1991) 3032. [8] C.L. Ma, T. Takasugi, S. Hanada, Script. Metall. Mater. 34 (1996) 1633. [9] C.L. Ma, T. Takasugi, S. Hanada, Script. Metall. Mater. 34 (1996) 1131. [10] T. Takasugi, S. Hanada, MRS Symp. Proc. Publ. 552 (1999) KK6.5.1. [11] T. Takasugi, M. Wada, Y. Kaneno, H. Inoue, Mater. Sci. Eng. A 329–331C (2002) 523. [12] T. Takasugi, ISIJ Intern. 43 (2003) 564. [13] T. Takasugi, T. Tsuyumu, Y. Kaneno, H. Inoue, J. Mater. Res. 15 (2000) 1881. [14] Y. Kaneno, M. Wada, H. Inoue, T. Takasugi, Mater. Transaction 42 (2001) 418. [15] C.L. Ma, T. Takasugi, S. Hanada, Mater. Trans. JIM 36 (1995) 30. [16] C.L. Ma, T. Takasugi, S. Hanada, Script. Metall. Mater. 32 (1995) 1025. [17] A. Turnbull, R.B. Hutchings, D.H. Ferriss, Mater. Sci. Eng. A238 (1997) 317. [18] T. Takasugi, M. Nagahsima, O. Izumi, Acta Metall. Mater. 37 (1989) 3425. [19] T. Takasugi, M. Nagahsima, O. Izumi, Acta Metall. Mater. 38 (1990) 747. [20] T. Nishiue, Y. Kaneno, H. Inoue, T. Takasugi, Intermetallics, in press. [21] Y. Hukai, N. Okuma, J. Appl. Phys. 32 (1993) 1256. [22] M. Iwamoto, Y. Hukai, Mater. Trans. JIM 40 (1999) 606. [23] H.K. Birnbaum, C. Buckley, F. Zeides, E. Sioris, P. Rozanak, S. Spooner, J.S. Lin, J. Alloys Comp. 253–254 (1997) 260. [24] H.D. Carstanjen, Phys. Stat. Sol. (a) 59 (1980) 11. [25] L. Yang, R.B. McLellan, Script. Metall. Mater. 42 (1994) 3993. [26] L. Yang, R.B. McLellan, Script. Metall. Mater. 32 (1994) 779. [27] L. Yang, R.B. MacLellan, J. Phys. Chem. Solids 57 (1996) 233. [28] T. Takasugi, K. Hono, S. Suzuki, S. Hanada, T. Sakurai, Script. Metall. Mater. 29 (1993) 1587. [29] H. Saitoh, M. Endoh, T. Aoyama, T. Misawa, T. Ohnishi, C.L. Ma, T. Takasugi, J. Mater. Sci. Lett. 15 (1996) 1297.