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RETRACTED: Role of hydrogen environment induced hydrogen embrittlement of Ti–8Al–1Mo–2V alloy
Materials Science and Engineering A 494 (2008) 143–146
Contents lists available at ScienceDirect
Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Role of hydrogen environment induced hydrogen embrittlement of Ti–8Al–1Mo–2V alloy Ronnie Higuchi Rusli ∗ School of Graduate Studies, University of Indonesia, Jl. Salemba 4, Jakarta 10430, Indonesia
a r t i c l e
i n f o
Article history: Received 2 March 2008 Received in revised form 2 April 2008 Accepted 7 April 2008 Keywords: Accoustic emission Crack propagation Embrittlement Fractographic Precracked ¨ Widmanstatten structure
a b s t r a c t Structural analysis by mean of metallographic, SEM fractographic and TEM replica technique including acoustics-emission studies have been carried out on Ti–8Al–1Mo–2V alloy specimen tested at room temperature in gaseous hydrogen environment. The result provided evidences of the presence of face centred cubic titanium hydride at the fracture surfaces, with discontinuous nature of crack propagation. The present work confirmed that an essentially continuous path of  phase is necessary for the occurrence of slow crack growth in gaseous hydrogen. Metallographic and fractographic observation leave little doubt that cracks propagates along the ␣– interface rather than through stable ␣ phase. © 2008 Elsevier B.V. All rights reserved.
1. Introduction Several workers [1–5] had reported slow crack growth in ␣– titanium alloys stressed in gaseous hydrogen. The room temperature behavior of Ti–6Al–4V heat treated to produce a variety of microstructures concluded that significant hydrogen gaseous embrittlement requires the presence of continuous paths of the  ¨ type present in Widmanstatten structures [4]. This finding, together with the fact that the hydrogen diffusivity and solubility are greater in  than in the ␣– interface, where by it causes embrittlement of the neighbouring ␣ phase. In this experiment it shows that embrittlement is resulted from repeated formation and rupture of a brittle hydride phase, and Nelson [6] had developed this film-rupture model, however, evidence of the formation of a hydride is meagre. Meyn [5] had ¨ studied Ti–8A1–1Mo–1V with a Widmanstatten structure, and reported that the fracture surfaces produced by hydrogen gaseous embrittlement were covered by a black deposit, and from X-ray diffraction pattern indicated a mixture of bct, fcc, and bcc titanium hydrides. The path of hydrogen gaseous embrittlement is also in some doubt. Nelson et al. [4] suggested, without convincing evidence, that cracking tend to follows the ␣– interface, whereas Meyn [5] had argued that failure involved cleavage of the ␣ on plane 9◦
∗ Tel.: +62 21 7506156; fax: +62 21 7653406. E-mail address: ronnie [email protected]. 0921-5093/$ – see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2008.04.006
from (1 0 1¯ 0). Both groups [4,5] reported the presence of striations or terrace structure on the fracture surface, and considered it as crack-arrest markings produced by discontinuous crack propagation. This study presents a supported evidence of the formation of hydride at the fracture surfaces, which establishes the path of cracking. Acoustic emission data also shows evidences which support the view that crack propagation is indeed discontinuous. A limited photograph of optical, SEM fractograph and TEM replica is presented here, but additional photographic evidence is available elsewhere [7]. 2. Experimental Commercially available Ti–8A1–1Mo–2V alloy was received as 1.2 mm thick sheet in mill annealed condition. Single edge notched tensile specimens with dimension of 150 mm × 20 mm × 1.2 mm, with a 45◦ notch extending to a depth of 3 mm were machined from the sheet, with their long axis perpendicular to the rolling direction. An equiaxed microstructure was obtained by heating for 8 h at 875 ◦ C (in the ␣ +  range) followed by quenching to room temperature. The resulting ␣ grain size was small with approximately 5% ¨  distributed discontinuously. A coarse Widmanstatten structure was produced by annealing the as-received sheet for 4 h at 1100 ◦ C (in the  range), slowly cooling to 875 ◦ C, and then quench to a room temperature. These samples displayed the characteristic of ‘basketweave’ microstructure, which consists of several intermeshing colonies of ␣ separated by essentially continuous layers of .
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R.H. Rusli / Materials Science and Engineering A 494 (2008) 143–146
Specimens were electro-polished and etched before testing, and fatigue precracked in the range of 2500–5000 N intentionally to introduce a sharp crack tip of approximately 0.5 mm in length. Experiments were carried out at room temperature in a tightly sealed stainless-steel chamber, which was evacuated prior to backfilling. Linde ultra high purity hydrogen and Linde high purity argon were employed without additional purification. Gas was flushed through the chamber for 1 h prior to testing and small positive pressure (∼104 N/m2 ) was maintained during the test. Specimens were tested to failure at a crosshead speed of 0.5 mm/min. Acoustic emission was continuously monitored during all tests by means of a differential transducer (Kraut Kramer) of natural frequency 225 kHz and a Yokogawa-Hewlett Packard audio-monitor. In the audio-monitor, the signal from the transducer was mixed with a fixed 225 kHz signal from an adjustable local oscillator to produce a high and low frequency component. The latter was displayed by means of a fibre optics monitor, or an oscilloscope coupled to a bio motion waveform recorders. The audio signal was also simultaneously recorded on a memory card recorder. It should be noted that this technique examined only frequencies in the range of 215–245 kHz and makes no attempt to analyse the complete range of frequencies emitted by the specimen. To simplify the interpretation of the results, the cross head of the tensile machine was stopped during slow crack growth in several tests and the cracks permitted to continue propagating under constant deflection. A replica technique is done using a JEOL 1200 transmission electron microscope (TEM). These studies were conducted directly on the fracture surface and on fragments removed from the fracture surface by conventional extraction techniques. In the former case, the fractured ends of specimens were positioned perpendicular to the electron beam so that patterns were obtained from material protruding from the fracture surface. 3. Results and discussion 3.1. Effect of hydrogen gaseous on mechanical behaviour ¨ Equiaxed and Widmanstatten specimens exhibited approximately the same ultimate strength, failing at loads of ∼19.58 kN. In both cases, overload failure was rapid and no slow crack growth was observed. Testing of the equiaxed material in hydrogen, produced no significant change in maximum load nor did slow crack growth occur during a 48 h exposure at 95% of peak load. Wid¨ manstatten ␣ specimen stressed in hydrogen, failed by slow crack growth and achieving a peak load of approximately 11.0 kN. Cracks
Fig. 1. Optical micrograph showing the intergranular fracture of hydrogen-induced ¨ slow crack growth in Widmanstatten specimens.
were observed to initiate at approximately 9.3 kN, and if the cross lead was stopped after initiation, it is continued to grow at an average velocity of ∼6 × 10−6 m/s until rapid overload failure occurred. 3.2. Crack path and fractographic studies Fracture in argon environment was ductile in both equiaxed and ¨ Widmanstatten specimens, and the fracture surfaces are “dimpled”. The equiaxed specimens tested in hydrogen failed predominantly by ductile rupture, nonetheless, the fracture surfaces exhibited small regions which corresponded to cracking along the ␣– interfaces. In contrast, however, the hydrogen-induced fracture in ¨ Widmanstatten specimens followed the ␣– interface (see Fig. 1), with only a small region of ductile overload failure. Examination of SEM stereospairs of the interfacial fracture surfaces indicated that the surfaces were curved, which is consistent with the path lying in the interface. Furthermore, the fracture surface was free from cleavage steps as shown in Fig. 2 and the crack follows the ␣– interface itself rather than cleaving the adjacent ␣ phase. It is considered that the near {1 0 1¯ 0} cleavage as reported by Meyn [5] probably results from ␣– separation, since the long boundaries of the ␣ plates are nearly parallel to this plane [8]. The striations or “terrace structures”, reported by other workers [5,6] were visible on the interfacial fracture surfaces in the Wid¨ manstatten specimens. These were found to be matching on the opposing surfaces (see Fig. 2), confirming the suggestion that they correspond to crack arrest markings which formed during discontinuous or intermittent propagation. The average crack advance distance of 1 m is similar to that observed by Meyn [5] in this type of alloy. Fig. 2 also illustrates another significant feature of these fracture surfaces, namely the presence of secondary cracks.
¨ Fig. 2. Two fracture surfaces of a hydrogen-induced crack in a Widmanstatten specimen. Notice of the typical crack arrest markings as well as the secondary cracking on one face only.
R.H. Rusli / Materials Science and Engineering A 494 (2008) 143–146
These cracks were observed only on one side of the opposing fracture faces, and occurred along certain striations suggesting that the cracking occurred when the crack was arrested. The confinement of the secondary cracks to one fracture face suggested the possibility that the phase on one side of the interface is embrittled, while the phase on the other side remains ductile. A further evidence that signified these fracture surfaces was the presence of a thin and irregularly shaped flakes like, that loosely attached to the surface. Further tests were carried out in which specimens were rapidly fracture in gaseous hydrogen to produce ductile and dimple fractures and subsequently exposed to the gas for ∼3 h. No flakes were detected in this specimen, which strongly suggest that the flakes were produced during the gaseous embrittlement process and not during subsequent exposure to hydrogen. 3.3. TEM replication studies The flakes described in the preceding section were then removed by means of replication techniques, and subsequently examined with the TEM. A typical flake is shown in Fig. 3. Selected area obtained from the flakes, were found to be consistent with the fcc titanium hydride structure. Extra reflections were occasionally observed, and these probably corresponded to ␣-Ti. Examination of the flakes using dark field indicated that lath like structure existed and as shown in Fig. 3, suggesting that the flakes consist of laths of hydride in ␣-Ti matrix. Limited result of electron-diffraction studies were also carried out directly on the fracture surfaces, and ring patterns were generally obtained which were consistent with the presence of ␣-Ti and the fcc hydride. In no circumstances were a combined of hydride and -Ti diffraction patterns was obtained, which strongly suggesting that fracture occurs by the growth of a thin layer of hydride laths in the ␣ adjacent to the . 3.4. Acoustic-emission analysis Acoustic emission observation during the tensile tests was characterized by discrete emissions with amplitude of five to eight times greater than that of background. These discrete emissions could be classified into two types (see Fig. 4), one having an extremely rapid alternation of less than 1 ms (<1 ms), and the other having considerably longer alternation times, >3 ms. However, the specimens tested in argon displayed only the former, ¨ while the tests of Widmanstatten specimens in hydrogen yielded the former prior to crack initiation and both types during crack growth. It is suggested that discrete emissions generated by deformation twinning are of the former type, while those associated with crack growth are of the latter type. The difference may
145
Fig. 4. Acoustic emission signals observed (a) before and (b) after crack initiation.
be due to the durations of each event. Twinning is essentially instantaneous, whereas each crack advances even it is thought to involve of several processes; the initiation at a localized region of the crack front, the lateral spreading of the crack along the front, and the subsequent plastic deformation associated with the blunting of the crack. It should be noted that both processes are considered to emit a broad spectrum of frequencies, and that only a narrow band, 220 ± 10 kHz, was monitored during our tests.
4. Conclusions This experimental result has confirmed that it is essentially a continuous path of  phase which is necessary for the occurrence of slow crack growth in hydrogen gaseous. Thus, limited amounts of brittle cracking occurred at the ␣– interfaces in equiaxed material, nonetheless, did not significantly influence the overall mechanical behaviour, whereas slow crack growth occurred in the specimens ¨ having the Widmanstatten structure. These results provide strong support for film rupture model ¨ for HGE of the Widmanstatten material [6]. The observations are fully consistent with the view that hydrogen enters and diffuses via the continuous  phase, and causes the precipitation of titanium hydride in the adjacent ␣, probably as laths or platelets [6,7]. The metallographic and fractographic observation leave little doubt to supported that the crack propagates along the ␣– interface itself rather than through the ␣ phase. The occurrence of secondary cracking at one fracture face but not at the matching region on the opposite face is consistent with this conclusion, since the former is thought correspond to hydrides ␣ and the latter to  phase. The loosely attached flakes observed on the fracture surface, result presumably from the linking of closely spaced secondary cracks. These conclusions model of the crack tip during hydrogen gaseous embrittlement, illustrating the hydride formation and cracking which gives rise to the observed fractography [7,8]. Discrete acoustic signals and some noise were found to be emitted during crack propagation. This observation, together with the presence of crack arrest markings on the fracture surfaces, supports the contention that crack propagation is discontinuous.
Acknowledgements
Fig. 3. Dark field micrograph, from a hydrogen-induced fracture surface, with a lath-like substructure clearly can be seen.
I would like to express my appreciation to Dr. K. Takagi and Dr. Y. Yamamura for setting the experimental apparatus. (*) This experimental work done at Sumittomo Steel Industries Ltd., Research Laboratory, Japan and was supported partially by the Japan Institute of Iron and Steel Research Development Administration, contract JAT(11-1)-1198-07.
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R.H. Rusli / Materials Science and Engineering A 494 (2008) 143–146
References [1] R.J. Walter, W.T. Chandler, Effect of high pressure hydrogen on metals, Technical Report no. D8 14.2, ASM, Metals Park, OH, 1968. [2] G.F. Pittinato, S.F. Frederick, Met. Trans. 1 (1970) 3241–3243. [3] D.P. Williams, H.G. Nelson, Met. Trans. 3 (1973) 2107–2113.
[4] [5] [6] [7]
H.G. Nelson, D.P. William, J.G. Stein, Met. Trans. 3 (1972) 469–475. D.A. Meyn, Met. Trans. 3 (1972) 2302–2305. H.G. Nelson, Met. Trans. A 7 (1976) 621–627. KOCH G.H., Intern Report of Stress Corrosion Cracking and Gaseous Hydrogen Embrittlement of Alpha Titanium Alloys, 1996. [8] M.J. Blackburn, Trans. ASM 59 (1996) 876–889.
Contents lists available at ScienceDirect
Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Role of hydrogen environment induced hydrogen embrittlement of Ti–8Al–1Mo–2V alloy Ronnie Higuchi Rusli ∗ School of Graduate Studies, University of Indonesia, Jl. Salemba 4, Jakarta 10430, Indonesia
a r t i c l e
i n f o
Article history: Received 2 March 2008 Received in revised form 2 April 2008 Accepted 7 April 2008 Keywords: Accoustic emission Crack propagation Embrittlement Fractographic Precracked ¨ Widmanstatten structure
a b s t r a c t Structural analysis by mean of metallographic, SEM fractographic and TEM replica technique including acoustics-emission studies have been carried out on Ti–8Al–1Mo–2V alloy specimen tested at room temperature in gaseous hydrogen environment. The result provided evidences of the presence of face centred cubic titanium hydride at the fracture surfaces, with discontinuous nature of crack propagation. The present work confirmed that an essentially continuous path of  phase is necessary for the occurrence of slow crack growth in gaseous hydrogen. Metallographic and fractographic observation leave little doubt that cracks propagates along the ␣– interface rather than through stable ␣ phase. © 2008 Elsevier B.V. All rights reserved.
1. Introduction Several workers [1–5] had reported slow crack growth in ␣– titanium alloys stressed in gaseous hydrogen. The room temperature behavior of Ti–6Al–4V heat treated to produce a variety of microstructures concluded that significant hydrogen gaseous embrittlement requires the presence of continuous paths of the  ¨ type present in Widmanstatten structures [4]. This finding, together with the fact that the hydrogen diffusivity and solubility are greater in  than in the ␣– interface, where by it causes embrittlement of the neighbouring ␣ phase. In this experiment it shows that embrittlement is resulted from repeated formation and rupture of a brittle hydride phase, and Nelson [6] had developed this film-rupture model, however, evidence of the formation of a hydride is meagre. Meyn [5] had ¨ studied Ti–8A1–1Mo–1V with a Widmanstatten structure, and reported that the fracture surfaces produced by hydrogen gaseous embrittlement were covered by a black deposit, and from X-ray diffraction pattern indicated a mixture of bct, fcc, and bcc titanium hydrides. The path of hydrogen gaseous embrittlement is also in some doubt. Nelson et al. [4] suggested, without convincing evidence, that cracking tend to follows the ␣– interface, whereas Meyn [5] had argued that failure involved cleavage of the ␣ on plane 9◦
∗ Tel.: +62 21 7506156; fax: +62 21 7653406. E-mail address: ronnie [email protected]. 0921-5093/$ – see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2008.04.006
from (1 0 1¯ 0). Both groups [4,5] reported the presence of striations or terrace structure on the fracture surface, and considered it as crack-arrest markings produced by discontinuous crack propagation. This study presents a supported evidence of the formation of hydride at the fracture surfaces, which establishes the path of cracking. Acoustic emission data also shows evidences which support the view that crack propagation is indeed discontinuous. A limited photograph of optical, SEM fractograph and TEM replica is presented here, but additional photographic evidence is available elsewhere [7]. 2. Experimental Commercially available Ti–8A1–1Mo–2V alloy was received as 1.2 mm thick sheet in mill annealed condition. Single edge notched tensile specimens with dimension of 150 mm × 20 mm × 1.2 mm, with a 45◦ notch extending to a depth of 3 mm were machined from the sheet, with their long axis perpendicular to the rolling direction. An equiaxed microstructure was obtained by heating for 8 h at 875 ◦ C (in the ␣ +  range) followed by quenching to room temperature. The resulting ␣ grain size was small with approximately 5% ¨  distributed discontinuously. A coarse Widmanstatten structure was produced by annealing the as-received sheet for 4 h at 1100 ◦ C (in the  range), slowly cooling to 875 ◦ C, and then quench to a room temperature. These samples displayed the characteristic of ‘basketweave’ microstructure, which consists of several intermeshing colonies of ␣ separated by essentially continuous layers of .
144
R.H. Rusli / Materials Science and Engineering A 494 (2008) 143–146
Specimens were electro-polished and etched before testing, and fatigue precracked in the range of 2500–5000 N intentionally to introduce a sharp crack tip of approximately 0.5 mm in length. Experiments were carried out at room temperature in a tightly sealed stainless-steel chamber, which was evacuated prior to backfilling. Linde ultra high purity hydrogen and Linde high purity argon were employed without additional purification. Gas was flushed through the chamber for 1 h prior to testing and small positive pressure (∼104 N/m2 ) was maintained during the test. Specimens were tested to failure at a crosshead speed of 0.5 mm/min. Acoustic emission was continuously monitored during all tests by means of a differential transducer (Kraut Kramer) of natural frequency 225 kHz and a Yokogawa-Hewlett Packard audio-monitor. In the audio-monitor, the signal from the transducer was mixed with a fixed 225 kHz signal from an adjustable local oscillator to produce a high and low frequency component. The latter was displayed by means of a fibre optics monitor, or an oscilloscope coupled to a bio motion waveform recorders. The audio signal was also simultaneously recorded on a memory card recorder. It should be noted that this technique examined only frequencies in the range of 215–245 kHz and makes no attempt to analyse the complete range of frequencies emitted by the specimen. To simplify the interpretation of the results, the cross head of the tensile machine was stopped during slow crack growth in several tests and the cracks permitted to continue propagating under constant deflection. A replica technique is done using a JEOL 1200 transmission electron microscope (TEM). These studies were conducted directly on the fracture surface and on fragments removed from the fracture surface by conventional extraction techniques. In the former case, the fractured ends of specimens were positioned perpendicular to the electron beam so that patterns were obtained from material protruding from the fracture surface. 3. Results and discussion 3.1. Effect of hydrogen gaseous on mechanical behaviour ¨ Equiaxed and Widmanstatten specimens exhibited approximately the same ultimate strength, failing at loads of ∼19.58 kN. In both cases, overload failure was rapid and no slow crack growth was observed. Testing of the equiaxed material in hydrogen, produced no significant change in maximum load nor did slow crack growth occur during a 48 h exposure at 95% of peak load. Wid¨ manstatten ␣ specimen stressed in hydrogen, failed by slow crack growth and achieving a peak load of approximately 11.0 kN. Cracks
Fig. 1. Optical micrograph showing the intergranular fracture of hydrogen-induced ¨ slow crack growth in Widmanstatten specimens.
were observed to initiate at approximately 9.3 kN, and if the cross lead was stopped after initiation, it is continued to grow at an average velocity of ∼6 × 10−6 m/s until rapid overload failure occurred. 3.2. Crack path and fractographic studies Fracture in argon environment was ductile in both equiaxed and ¨ Widmanstatten specimens, and the fracture surfaces are “dimpled”. The equiaxed specimens tested in hydrogen failed predominantly by ductile rupture, nonetheless, the fracture surfaces exhibited small regions which corresponded to cracking along the ␣– interfaces. In contrast, however, the hydrogen-induced fracture in ¨ Widmanstatten specimens followed the ␣– interface (see Fig. 1), with only a small region of ductile overload failure. Examination of SEM stereospairs of the interfacial fracture surfaces indicated that the surfaces were curved, which is consistent with the path lying in the interface. Furthermore, the fracture surface was free from cleavage steps as shown in Fig. 2 and the crack follows the ␣– interface itself rather than cleaving the adjacent ␣ phase. It is considered that the near {1 0 1¯ 0} cleavage as reported by Meyn [5] probably results from ␣– separation, since the long boundaries of the ␣ plates are nearly parallel to this plane [8]. The striations or “terrace structures”, reported by other workers [5,6] were visible on the interfacial fracture surfaces in the Wid¨ manstatten specimens. These were found to be matching on the opposing surfaces (see Fig. 2), confirming the suggestion that they correspond to crack arrest markings which formed during discontinuous or intermittent propagation. The average crack advance distance of 1 m is similar to that observed by Meyn [5] in this type of alloy. Fig. 2 also illustrates another significant feature of these fracture surfaces, namely the presence of secondary cracks.
¨ Fig. 2. Two fracture surfaces of a hydrogen-induced crack in a Widmanstatten specimen. Notice of the typical crack arrest markings as well as the secondary cracking on one face only.
R.H. Rusli / Materials Science and Engineering A 494 (2008) 143–146
These cracks were observed only on one side of the opposing fracture faces, and occurred along certain striations suggesting that the cracking occurred when the crack was arrested. The confinement of the secondary cracks to one fracture face suggested the possibility that the phase on one side of the interface is embrittled, while the phase on the other side remains ductile. A further evidence that signified these fracture surfaces was the presence of a thin and irregularly shaped flakes like, that loosely attached to the surface. Further tests were carried out in which specimens were rapidly fracture in gaseous hydrogen to produce ductile and dimple fractures and subsequently exposed to the gas for ∼3 h. No flakes were detected in this specimen, which strongly suggest that the flakes were produced during the gaseous embrittlement process and not during subsequent exposure to hydrogen. 3.3. TEM replication studies The flakes described in the preceding section were then removed by means of replication techniques, and subsequently examined with the TEM. A typical flake is shown in Fig. 3. Selected area obtained from the flakes, were found to be consistent with the fcc titanium hydride structure. Extra reflections were occasionally observed, and these probably corresponded to ␣-Ti. Examination of the flakes using dark field indicated that lath like structure existed and as shown in Fig. 3, suggesting that the flakes consist of laths of hydride in ␣-Ti matrix. Limited result of electron-diffraction studies were also carried out directly on the fracture surfaces, and ring patterns were generally obtained which were consistent with the presence of ␣-Ti and the fcc hydride. In no circumstances were a combined of hydride and -Ti diffraction patterns was obtained, which strongly suggesting that fracture occurs by the growth of a thin layer of hydride laths in the ␣ adjacent to the . 3.4. Acoustic-emission analysis Acoustic emission observation during the tensile tests was characterized by discrete emissions with amplitude of five to eight times greater than that of background. These discrete emissions could be classified into two types (see Fig. 4), one having an extremely rapid alternation of less than 1 ms (<1 ms), and the other having considerably longer alternation times, >3 ms. However, the specimens tested in argon displayed only the former, ¨ while the tests of Widmanstatten specimens in hydrogen yielded the former prior to crack initiation and both types during crack growth. It is suggested that discrete emissions generated by deformation twinning are of the former type, while those associated with crack growth are of the latter type. The difference may
145
Fig. 4. Acoustic emission signals observed (a) before and (b) after crack initiation.
be due to the durations of each event. Twinning is essentially instantaneous, whereas each crack advances even it is thought to involve of several processes; the initiation at a localized region of the crack front, the lateral spreading of the crack along the front, and the subsequent plastic deformation associated with the blunting of the crack. It should be noted that both processes are considered to emit a broad spectrum of frequencies, and that only a narrow band, 220 ± 10 kHz, was monitored during our tests.
4. Conclusions This experimental result has confirmed that it is essentially a continuous path of  phase which is necessary for the occurrence of slow crack growth in hydrogen gaseous. Thus, limited amounts of brittle cracking occurred at the ␣– interfaces in equiaxed material, nonetheless, did not significantly influence the overall mechanical behaviour, whereas slow crack growth occurred in the specimens ¨ having the Widmanstatten structure. These results provide strong support for film rupture model ¨ for HGE of the Widmanstatten material [6]. The observations are fully consistent with the view that hydrogen enters and diffuses via the continuous  phase, and causes the precipitation of titanium hydride in the adjacent ␣, probably as laths or platelets [6,7]. The metallographic and fractographic observation leave little doubt to supported that the crack propagates along the ␣– interface itself rather than through the ␣ phase. The occurrence of secondary cracking at one fracture face but not at the matching region on the opposite face is consistent with this conclusion, since the former is thought correspond to hydrides ␣ and the latter to  phase. The loosely attached flakes observed on the fracture surface, result presumably from the linking of closely spaced secondary cracks. These conclusions model of the crack tip during hydrogen gaseous embrittlement, illustrating the hydride formation and cracking which gives rise to the observed fractography [7,8]. Discrete acoustic signals and some noise were found to be emitted during crack propagation. This observation, together with the presence of crack arrest markings on the fracture surfaces, supports the contention that crack propagation is discontinuous.
Acknowledgements
Fig. 3. Dark field micrograph, from a hydrogen-induced fracture surface, with a lath-like substructure clearly can be seen.
I would like to express my appreciation to Dr. K. Takagi and Dr. Y. Yamamura for setting the experimental apparatus. (*) This experimental work done at Sumittomo Steel Industries Ltd., Research Laboratory, Japan and was supported partially by the Japan Institute of Iron and Steel Research Development Administration, contract JAT(11-1)-1198-07.
146
R.H. Rusli / Materials Science and Engineering A 494 (2008) 143–146
References [1] R.J. Walter, W.T. Chandler, Effect of high pressure hydrogen on metals, Technical Report no. D8 14.2, ASM, Metals Park, OH, 1968. [2] G.F. Pittinato, S.F. Frederick, Met. Trans. 1 (1970) 3241–3243. [3] D.P. Williams, H.G. Nelson, Met. Trans. 3 (1973) 2107–2113.
[4] [5] [6] [7]
H.G. Nelson, D.P. William, J.G. Stein, Met. Trans. 3 (1972) 469–475. D.A. Meyn, Met. Trans. 3 (1972) 2302–2305. H.G. Nelson, Met. Trans. A 7 (1976) 621–627. KOCH G.H., Intern Report of Stress Corrosion Cracking and Gaseous Hydrogen Embrittlement of Alpha Titanium Alloys, 1996. [8] M.J. Blackburn, Trans. ASM 59 (1996) 876–889.