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Visualization of trapped hydrogen along grain boundaries and its quantitative contribution to hydrogen-induced intergranular fracture in pure nickel
Materialia 53 (2019) 100478
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Visualization of trapped hydrogen along grain boundaries and its quantitative contribution to hydrogen-induced intergranular fracture in pure nickel Kentaro Wada a,b,∗, Junichiro Yamabe b,c, Hisao Matsunaga b,d,e,f a
Graduate School of Engineering, Kyushu University, 744 Moto-oka, Nishi-ku, Fukuoka 819-0395, Japan AIST–Kyushu University Hydrogen Materials Laboratory (HydoMate), National Institute of Advanced Industrial Science and Technology (AIST), 744 Moto-oka, Nishi-ku, Fukuoka 819-0395, Japan c Department of Mechanical Engineering, Fukuoka University, 8-19-1 Nanakuma, Jonan-ku, Fukuoka 814-0180, Japan d Department of Mechanical Engineering, Kyushu University, 744 Moto-oka, Nishi-ku, Fukuoka 819-0395, Japan e Research Center for Hydrogen Industrial Use and Storage (HYDROGENIUS), Kyushu University, 744 Moto-oka, Nishi-ku, Fukuoka 819-0395, Japan f International Institute for Carbon-Neutral Energy Research (WPI-I2CNER), Kyushu University, 744 Moto-oka, Nishi-ku, Fukuoka 819-0395, Japan b
a r t i c l e
i n f o
Keywords: Nickel Hydrogen embrittlement Secondary ion mass spectroscopy (SIMS) Grain boundary segregation Tension test Thermal desorption analysis (TDA)
a b s t r a c t The phenomenon of hydrogen-trapping and its quantitative contribution to hydrogen-induced intergranular (IG) fracture were studied using a combination of thermal desorption analysis, secondary ion mass spectrometry and slow strain rate tensile tests. Hydrogen was trapped along grain boundaries (GBs) with a binding energy of ≈20 kJ/mol, accompanied by IG sulfur. The true fracture stress and fracture surface morphology were strongly dependent on the concentration of trapped hydrogen, leading to the conclusion that the hydrogen-induced IG fracture of pure Ni is controlled by the concentration of hydrogen trapped along GBs, and not by the concentration of lattice hydrogen.
Introduction In order to increase the reliability of components that are exposed to hydrogen gas, a mechanism-based understanding of hydrogen embrittlement (HE) is mandatory. To date, many basic HE mechanisms have been proposed, including hydrogen-enhanced decohesion (HEDE) [1,2], hydrogen-enhanced stabilized lattice defects [3–6] and hydrogenenhanced localized plasticity (HELP) [7–9]. At the same time, numerous efforts have been undertaken to reasonably explain the HE phenomena in various materials, based on experimental and observational results [1–7,9–12]. As one of the commonly-researched model metals, several studies have been conducted on pure Ni [1,4,10–15], owing to its catastrophic HE sensitivity and accompanying intergranular (IG) fracture. For instance, Martin et al. documented that IG cracking can result from the HEDE, with the assistance of the HELP [11]. As another mechanism, Robertson et al. stated that hydrogen accumulated around GBs locally softens the material via the HELP, unaided by the HEDE, resulting in locally ductile fracture with IG features [10]. While controversy reigns over the specific HE mechanism for pure Ni, most major studies postulate either hydrogen accumulation along
∗
GBs via thermodynamic equilibrium [12,13,15], or transportation by moving dislocations [11]. For example, Martin et al. stated that moving dislocations transport hydrogen atoms towards GBs so that the hydrogen concentration reaches at the critical value to cause GB separation via HEDE. Conversely, Harris et al. conducted the slow strain rate tensile (SSRT) tests with H-charged pure Ni at RT and 77 K, indicating the degradation of ductility even at 77 K, at which hydrogen transportation by moving dislocation cannot take place [12]. Also, the authors have recently obtained a similar result to that by Harris et al., showing that pure Ni was still embrittled by hydrogen at 77 K [16]. These results suggest that initially-segregated hydrogen via the thermodynamic equilibrium along GBs plays a major role in the course of HE [12,16]. Indeed, using a secondary ion mass spectrometry (SIMS) analysis, Fukushima and Birnbaum detected highly-concentrated deuterium (not hydrogen) along GBs [17]. Similarly, Oudriss et al. discovered hydrogen segregation at low-energy grain boundaries, using the SIMS measurement [18]. However, no existing research has ever quantitatively identified the contribution of segregated hydrogen to hydrogen-induced IG fracture. The present study focuses on the accumulation of hydrogen along GBs and its contribution to hydrogen-induced IG fracture quantitatively. The hydrogen-trapping phenomenon in thermally H-charged pure Ni
Corresponding author at: Graduate School of Engineering, Kyushu University, 744 Moto-oka, Nishi-ku, Fukuoka 819-0395, Japan. E-mail address: [email protected] (K. Wada).
https://doi.org/10.1016/j.mtla.2019.100478 Received 12 September 2019; Accepted 18 September 2019 Available online 20 September 2019 2589-1529/© 2019 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
K. Wada, J. Yamabe and H. Matsunaga
was visualized using thermal desorption analysis (TDA) and SIMS. Subsequently, using SSRT tests, the quantitative contribution of the trapped hydrogen to the IG fracture was duly investigated. Experimental Vacuum induction melting was used to produce an ingot of pure Ni, which contained impurities of 0.001 C, 0.020 O, 0.0006 N, <0.001 P and <0.001 S (wt%). The ingot was hot-rolled at 1423 K so as to obtain homogenous microstructures and equiaxed grains, with an average grain size of 200 μm. Cylindrical samples were then machined with both diameter and length of 5 mm for the TDA, and with a diameter of 8 mm and a length of 3 mm for the SIMS analysis. SSRT specimens with a gagediameter of 4 mm and a gage-length of 30 mm were also produced from the same ingot, ensuring that the tensile direction lies parallel to the rolling direction. The gage-sections of the specimens were later mirrorpolished with 1-μm diamond paste. All samples and specimens were annealed at 923 K for two hours in vacuum, to ensure the removal of any residual strain potentially introduced during the machining and polishing processes. Some of these samples and specimens were thermally charged in hydrogen gas at pressures of 0.7, 11, 50 and 100 MPa, at a temperature of 543 K for 200 h. Maintaining a constant heating rate of 100 K/h, the TDA samples were heated up to 1073 K and the desorbed hydrogen was accordingly measured by means of TDA (JTF-20W, J-SCIENCE LAB Co., Ltd., Japan). Hydrogen distribution was visualized via SIMS (IMS–7f, CAMECA, AMETEC, France). Cesium ion with an intensity of 60 nA and an acceleration voltage of 15 kV was used as the primary ion, which was rastered over a square area of 300 μm. To avoid any false signals of hydrogen associated with surface contaminations, cold-trapping and silicon-sputtering methods were applied [19–22]. The SSRT tests were conducted in laboratory air at room temperature, with a constant deformation rate of 0.09 mm/min, corresponding to an approximate initial strain rate of
Materialia 53 (2019) 100478
5 × 10−5 s−1 . After the SSRT tests, ductility was evaluated through a reduction of area (RA). Using a two-dimensional optical micrometer, minimum specimen diameters were measured during the deformation process in order to calculate the true stress and the maximum true stress was defined as the true fracture stress (TFS). Results and discussion Fig. 1(a–d) shows the TDA profiles of H-charged samples. Deconvoluted red and green Gaussian curves are represented in the figure. The sum of these curves is denoted by the broken blue line, being in good agreement with the experimental data indicated by the black line. Despite the different hydrogen-gas pressures, the H-charged samples displayed two major peaks at around 600 K (Peak 1) and 750 K (Peak 2) except for the peak 2 in 0.7-MPa H-charged sample with lower temperature (683 K). This reason is not clear; however, may be attributed to lower hydrogen concentration at the corresponding trap-site. Fig. 1(f) presents the hydrogen concentrations of Peak 1, cH 1 , and Peak 2, cH 2 , as functions of the hydrogen-gas fugacity expressed as follows [23]: 𝑓 = 𝑝H exp
𝑏 ⋅ 𝑝H 𝑅𝑇
(1)
where, pH is the hydrogen-gas pressure, b is a constant, which represents the finite volume of the hydrogen molecules (1.584 × 10−5 m3 /mol), R is the universal gas constant (8.314 J K−1 mol−1 ) and T is the absolute temperature (K). As illustrated by the red plots, the values of cH 1 demonstrated a linear correlation with the square root of fugacity, f 1/2 , following the Sieverts’ Law derived from the thermodynamic equilibrium [24]. The proportional constant, defined as the solubility of hydrogen at a given temperature, was determined to be 4.97 wt ppm/MPa1/2 . In contrast, the values of cH 2 remained almost constant regardless of the fugacity, leading to the assertion that Peak 2 hydrogen corresponds to some sort of trapped one. Based on the equilibrium hydrogen-trapping theory, the trap-site occupancy, 𝜃 x , can be quantified by the following
Fig. 1. TDA profiles of non-deformed pure Ni that was H-charged in (a) 100 MPa-, (b) 50 MPa-, (c) 11 MPa- and (d) 0.7 MPa-hydrogen gas, as well as H-free, nondeformed pure Ni (e). The profiles were deconvoluted into two Gaussian peaks, designated as Peak 1 and Peak 2, with the hydrogen contents, cH , of each peak provided in (f).
K. Wada, J. Yamabe and H. Matsunaga
Materialia 53 (2019) 100478
equation [25]: 𝐸 𝜃X = 𝜃L exp b 1 − 𝜃X 𝑅𝑇
(2)
where, 𝜃 L is the lattice-site occupancy (𝜃 L << 1) and Eb is the binding energy of a trap-site (kJ/mol). Based on the simple assumption that one trap-site can trap one hydrogen atom and that the O-site is the only possible location, 𝜃 L is equivalent to the lattice-hydrogen concentration in at. ppm. In the determination of trap-site occupancy, it was presumed that an equilibrium state was achieved in the specimens since, in this study, they were aged at the room temperature (T ≈ 300 K) for more than three hours before starting tensile testing after the H-charging. In fact, Lassila et al. reported that solute hydrogen in pure Ni can reach an equilibrium state in approximately 103 s at 300 K after a high-temperature H-charging [13]. Fig. 2(a) presents the trap-site occupancy as a function of binding energy for each H-charging pressure, where the measured values of cH 1 were used as 𝜃 L . In order to determine the binding energy of the trapsite, the relationship between 𝜃 x and cH 2 were obtained as a function of Eb by using Eq. (2). Fig. 2(b) shows the 𝜃 x ‒ cH 2 relation in the case of Eb = 20.6 kJ/mol, where the solid line represents the regression. This Eb value was optimized so that 𝜃 x becomes proportional to cH 2 , i.e., a coefficient of determination, R2 , was maximized (cf. Fig. 2(c)). The R2 is defined by the following equation: ∑𝑛 2 (𝜃X,𝑖 − 𝜃̂X,𝑖 ) 𝑅2 = 1 − ∑𝑖=1 (3) 𝑛 ̄ 2 𝑖=1 (𝜃X,𝑖 − 𝜃X,𝑖 ) where, n is the number of data, 𝜃 x, i is the trap-site occupancy obtained from experiment in data number of i, 𝜃̄x is the mean value of trap-site occupancies and 𝜃̂x, 𝑖 is the trap-site occupancy calculated from a regression line in data number of i. Above result suggests that the second peak corresponds to hydrogen trapped at an unidentified trap-site with a binding energy of ≈ 20 kJ/mol. In the case of pure Ni, several potential trap-sites have been proposed in the existing literature, including the actual GBs (Eb = 21– 29 kJ/mol) [26–29], vacancies and their clusters (23–50 kJ/mol) [30– 32], dislocation cores (32 kJ/mol) [29], as well as interstitial sulfur (3– 4 kJ/mol) [33]. Among these trap-sites, the trap-energy determined in this study corresponded to that of the actual GBs or vacancies. Fig. 3(a–f) showcase the two-dimensional mappings of SIMS signals for 1 H and 32 S, with GB overlap, as taken via electron backscatter diffraction (EBSD) pattern analysis (indicated by red and white lines). Red lines represent ‘special’ (3 ≤ Σ ≤ 29) boundaries whereas white ones represent random boundaries [34]. As was already reported [17,18], the 1 H signals were spotted along GBs (indicated by black arrows in Fig. 3(a–c)) in many cases, thus clearly confirming that the Peak 2 hydrogen detected in the TDA measurement was trapped there. In several previous researches, it has been reported that the ‘special’ boundaries preferentially trap the hydrogen [18,34]. Indeed, hydrogen-trapping phenomena varied depending on GBs also in the present study; however, no clear difference in the trapping capability between ‘special’ and random boundaries was detected (Fig. 3(a–c)). Additionally, two-dimensional SIMS mapping offered visual confirmation that some of the hydrogen trapped along GBs existed at the same location as sulfur (indicated by white arrows in Fig. 3(a–f)). A segregant such as IG sulfur has been proposed in past studies to behave as a trap-site [14,35], and a similar accumulation of deuterium around IG sulfur was also previously reported [17]. Thus, although its binding energy has not yet been clarified, IG sulfur undeniably assists with hydrogen accumulation along GBs, with the trapping of hydrogen around the segregated sulfur [14,35]. Fig. 4(a) and (b) demonstrates the stress‒stroke curves of H-free and H-charged specimens, as well as the TFS values as functions of hydrogengas fugacity. Based on the data, two notable facts can be immediately ascertained. First of all, flow stresses increased with each increase in fugacity. Secondly, in the specimens H-charged at 11, 50 and 100 MPa, the RA and TFS values were nearly constant (RA of 0.05–0.07 and TFS of
Fig. 2. (a) Trap-site occupancy as calculated by Eq. (2). (b) Relationship between trap-site occupancy and Peak 2 hydrogen content for the binding energy Eb of 20.6 kJ/mol. (c) Coefficient of determination, R2 , as a function of binding energy.
228–265 MPa, respectively) and the fracture surfaces were covered only with IG facets (Fig. 4(d) and (e)). On the other hand, the 0.7-MPa Hcharged specimen failed after plastic instability, with an RA of 0.57 and TFS of 448 MPa. Furthermore, the IG fracture surface was discovered in the limited region of the specimen (Fig. 4(c)). Specifically, the degree of HE in pure Ni was not dependent on the lattice-hydrogen concentration
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Materialia 53 (2019) 100478
Fig. 3. Two-dimensional mappings of 1 H (a–c) and 32 S (d–f), as analyzed via SIMS. Red and white lines represent ‘special’ (3 ≤ Σ ≤ 29) and random grain boundaries examined using the EBSD technique, respectively. In (a–f), white and black arrows indicate the hydrogen trapped at IG sulfur and the hydrogen trapped along GB. These mappings should be compared with (g–i) showing inversed pole figure maps around the observation areas indicated by red rectangles.
Fig. 4. (a) Stress–stroke curves of H-free and H-charged pure Ni. (b) Peak 2 hydrogen content and true fracture stress as functions of hydrogen-gas fugacity. (c–e) Fracture surface morphologies of H-charged pure Ni in 0.7 MPa (c), 50 MPa (d) and 100 MPa (e) hydrogen gas.
when gas pressure exceeded 11 MPa; the magnitude of HE was rapidly mitigated with the decrease in gas pressure from 11 to 0.7 MPa. These tendencies correlate with the concentration of trapped hydrogen, cH 2 (Fig. 1(f) and (b)), leading to the assertion that the extent of degradation is attributed to the hydrogen trapped along GBs, as associated with the Peak 2 hydrogen. This assertion is well-consistent with the conclusion of our previous research, demonstrating that initially-segregated hydrogen dominates the hydrogen-induced ductility loss [16].
In addition, based on the SIMS visualization of the IG sulfur along GBs, the impact of segregated sulfur on hydrogen-induced degradation must be considered. The segregated sulfur along GBs was reported to play two possible roles in the enhancement of hydrogen-induced IG fracture [14,35]. Firstly, as sulfur itself is well-known IG segregant to decrease GB strength [36,37], the IG sulfur possibly increases the propensity for IG fracture, in tandem with the hydrogen-induced IG cracking. Secondly, as obtained in the present SIMS observation, the IG sulfur
K. Wada, J. Yamabe and H. Matsunaga
may attract the hydrogen atoms [14,17,35], thereby making it easier to fulfill the critical hydrogen concentration required at GBs in order to drive the hydrogen-induced IG cracking. Consequently, through a combination of these two roles, it has been inferred that segregated sulfur and hydrogen synergistically weaken the GB strength [14,35]. Conclusions The present paper highlighted the hydrogen-trapping phenomenon and its quantitative contribution to the hydrogen-induced IG fracture in pure Ni. The TDA profiles revealed the existence of a trap-site with a binding energy of ≈ 20 kJ/mol, corresponding to that of the GBs or vacancies reported in existing literature. Visualization via SIMS measurement clearly indicated that this hydrogen was trapped along GBs, primarily accompanied by IG sulfur. Furthermore, from the viewpoints of the RA and TFS obtained in the SSRT tests, the degree of HE was dependent on the trapped hydrogen. Therefore, it can be concluded that the HE of pure Ni is dominated by the hydrogen existing along GBs and not by the lattice-hydrogen concentration. Declaration of Competing Interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. Acknowledgments A part of this study (EBSD analysis) was conducted with support from the Advanced Characterization Platform of the Nanotechnology Platform Japan, sponsored by the Ministry of Education, Culture, Sports, Science and Technology (MEXT) of Japan. The authors also wish to thank the staff of the Cameca Division – Ametek Co., Ltd., and Kobe Material Testing Laboratory Co., Ltd., for providing technical support and helpful information in relation to the SIMS analysis. Funding This research did not receive any specific grant from funding agencies in the public, commercial, or not-for-profit sectors. References [1] A.R. Troiano, The role of hydrogen and other interstitials in the mechanical behavior of metals, Trans. ASM 52 (1960) 54–80. [2] R.A. Oriani, P.H. Josephic, Equilibrium aspects of hydrogen-induced cracking of steels, Acta Metall. 22 (1974) 1065–1074, doi:10.1016/0001-6160(74)90061-3. [3] M. Nagumo, Hydrogen related failure of steels – a new aspect, Mater. Sci. Technol. 20 (2004) 940–950, doi:10.1179/026708304225019687. [4] S.K. Lawrence, Y. Yagodzinskyy, H. Hänninen, E. Korhonen, F. Tuomisto, Z.D. Harris, B.P. Somerday, Effects of grain size and deformation temperature on hydrogenenhanced vacancy formation in Ni alloys, Acta Mater. 128 (2017) 218–226, doi:10.1016/j.actamat.2017.02.016. [5] R. Kirchheim, Reducing grain boundary, dislocation line and vacancy formation energies by solute segregation. I. Theoretical background, Acta Mater. 55 (2007) 5129– 5138, doi:10.1016/j.actamat.2007.05.047. [6] R. Kirchheim, Reducing grain boundary, dislocation line and vacancy formation energies by solute segregation. II. Experimental evidence and consequences, Acta Mater. 55 (2007) 5139–5148, doi:10.1016/j.actamat.2007.05.033. [7] I.M. Robertson, P. Sofronis, A. Nagao, M.L. Martin, S. Wang, D.W. Gross, K.E. Nygren, Hydrogen Embrittlement Understood, Metall. Mater. Trans. B 46 (2015) 1085– 1103, doi:10.1007/s11663-015-0325-y. [8] I.M. Robertson, The effect of hydrogen on dislocation dynamics, Eng. Fract. Mech. 68 (2001) 671–692, doi:10.1016/S0013-7944(01)00011-X. [9] H.K. Birnbaum, P. Sofronis, Hydrogen-enhanced localized plasticity-a mechanism for hydrogen-related fracture, Mater. Sci. Eng. A 176 (1994) 191–202, doi:10.1016/0921-5093(94)90975-X. [10] I.M. Robertson, T. Tabata, W. Wei, F. Heubaum, H.K. Birnbaum, Hydrogen embrittlement and grain boundary fracture, Scr. Metall. 18 (1984) 841–846, doi:10.1016/0036-9748(84)90407-1.
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Full Length Article
Visualization of trapped hydrogen along grain boundaries and its quantitative contribution to hydrogen-induced intergranular fracture in pure nickel Kentaro Wada a,b,∗, Junichiro Yamabe b,c, Hisao Matsunaga b,d,e,f a
Graduate School of Engineering, Kyushu University, 744 Moto-oka, Nishi-ku, Fukuoka 819-0395, Japan AIST–Kyushu University Hydrogen Materials Laboratory (HydoMate), National Institute of Advanced Industrial Science and Technology (AIST), 744 Moto-oka, Nishi-ku, Fukuoka 819-0395, Japan c Department of Mechanical Engineering, Fukuoka University, 8-19-1 Nanakuma, Jonan-ku, Fukuoka 814-0180, Japan d Department of Mechanical Engineering, Kyushu University, 744 Moto-oka, Nishi-ku, Fukuoka 819-0395, Japan e Research Center for Hydrogen Industrial Use and Storage (HYDROGENIUS), Kyushu University, 744 Moto-oka, Nishi-ku, Fukuoka 819-0395, Japan f International Institute for Carbon-Neutral Energy Research (WPI-I2CNER), Kyushu University, 744 Moto-oka, Nishi-ku, Fukuoka 819-0395, Japan b
a r t i c l e
i n f o
Keywords: Nickel Hydrogen embrittlement Secondary ion mass spectroscopy (SIMS) Grain boundary segregation Tension test Thermal desorption analysis (TDA)
a b s t r a c t The phenomenon of hydrogen-trapping and its quantitative contribution to hydrogen-induced intergranular (IG) fracture were studied using a combination of thermal desorption analysis, secondary ion mass spectrometry and slow strain rate tensile tests. Hydrogen was trapped along grain boundaries (GBs) with a binding energy of ≈20 kJ/mol, accompanied by IG sulfur. The true fracture stress and fracture surface morphology were strongly dependent on the concentration of trapped hydrogen, leading to the conclusion that the hydrogen-induced IG fracture of pure Ni is controlled by the concentration of hydrogen trapped along GBs, and not by the concentration of lattice hydrogen.
Introduction In order to increase the reliability of components that are exposed to hydrogen gas, a mechanism-based understanding of hydrogen embrittlement (HE) is mandatory. To date, many basic HE mechanisms have been proposed, including hydrogen-enhanced decohesion (HEDE) [1,2], hydrogen-enhanced stabilized lattice defects [3–6] and hydrogenenhanced localized plasticity (HELP) [7–9]. At the same time, numerous efforts have been undertaken to reasonably explain the HE phenomena in various materials, based on experimental and observational results [1–7,9–12]. As one of the commonly-researched model metals, several studies have been conducted on pure Ni [1,4,10–15], owing to its catastrophic HE sensitivity and accompanying intergranular (IG) fracture. For instance, Martin et al. documented that IG cracking can result from the HEDE, with the assistance of the HELP [11]. As another mechanism, Robertson et al. stated that hydrogen accumulated around GBs locally softens the material via the HELP, unaided by the HEDE, resulting in locally ductile fracture with IG features [10]. While controversy reigns over the specific HE mechanism for pure Ni, most major studies postulate either hydrogen accumulation along
∗
GBs via thermodynamic equilibrium [12,13,15], or transportation by moving dislocations [11]. For example, Martin et al. stated that moving dislocations transport hydrogen atoms towards GBs so that the hydrogen concentration reaches at the critical value to cause GB separation via HEDE. Conversely, Harris et al. conducted the slow strain rate tensile (SSRT) tests with H-charged pure Ni at RT and 77 K, indicating the degradation of ductility even at 77 K, at which hydrogen transportation by moving dislocation cannot take place [12]. Also, the authors have recently obtained a similar result to that by Harris et al., showing that pure Ni was still embrittled by hydrogen at 77 K [16]. These results suggest that initially-segregated hydrogen via the thermodynamic equilibrium along GBs plays a major role in the course of HE [12,16]. Indeed, using a secondary ion mass spectrometry (SIMS) analysis, Fukushima and Birnbaum detected highly-concentrated deuterium (not hydrogen) along GBs [17]. Similarly, Oudriss et al. discovered hydrogen segregation at low-energy grain boundaries, using the SIMS measurement [18]. However, no existing research has ever quantitatively identified the contribution of segregated hydrogen to hydrogen-induced IG fracture. The present study focuses on the accumulation of hydrogen along GBs and its contribution to hydrogen-induced IG fracture quantitatively. The hydrogen-trapping phenomenon in thermally H-charged pure Ni
Corresponding author at: Graduate School of Engineering, Kyushu University, 744 Moto-oka, Nishi-ku, Fukuoka 819-0395, Japan. E-mail address: [email protected] (K. Wada).
https://doi.org/10.1016/j.mtla.2019.100478 Received 12 September 2019; Accepted 18 September 2019 Available online 20 September 2019 2589-1529/© 2019 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
K. Wada, J. Yamabe and H. Matsunaga
was visualized using thermal desorption analysis (TDA) and SIMS. Subsequently, using SSRT tests, the quantitative contribution of the trapped hydrogen to the IG fracture was duly investigated. Experimental Vacuum induction melting was used to produce an ingot of pure Ni, which contained impurities of 0.001 C, 0.020 O, 0.0006 N, <0.001 P and <0.001 S (wt%). The ingot was hot-rolled at 1423 K so as to obtain homogenous microstructures and equiaxed grains, with an average grain size of 200 μm. Cylindrical samples were then machined with both diameter and length of 5 mm for the TDA, and with a diameter of 8 mm and a length of 3 mm for the SIMS analysis. SSRT specimens with a gagediameter of 4 mm and a gage-length of 30 mm were also produced from the same ingot, ensuring that the tensile direction lies parallel to the rolling direction. The gage-sections of the specimens were later mirrorpolished with 1-μm diamond paste. All samples and specimens were annealed at 923 K for two hours in vacuum, to ensure the removal of any residual strain potentially introduced during the machining and polishing processes. Some of these samples and specimens were thermally charged in hydrogen gas at pressures of 0.7, 11, 50 and 100 MPa, at a temperature of 543 K for 200 h. Maintaining a constant heating rate of 100 K/h, the TDA samples were heated up to 1073 K and the desorbed hydrogen was accordingly measured by means of TDA (JTF-20W, J-SCIENCE LAB Co., Ltd., Japan). Hydrogen distribution was visualized via SIMS (IMS–7f, CAMECA, AMETEC, France). Cesium ion with an intensity of 60 nA and an acceleration voltage of 15 kV was used as the primary ion, which was rastered over a square area of 300 μm. To avoid any false signals of hydrogen associated with surface contaminations, cold-trapping and silicon-sputtering methods were applied [19–22]. The SSRT tests were conducted in laboratory air at room temperature, with a constant deformation rate of 0.09 mm/min, corresponding to an approximate initial strain rate of
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5 × 10−5 s−1 . After the SSRT tests, ductility was evaluated through a reduction of area (RA). Using a two-dimensional optical micrometer, minimum specimen diameters were measured during the deformation process in order to calculate the true stress and the maximum true stress was defined as the true fracture stress (TFS). Results and discussion Fig. 1(a–d) shows the TDA profiles of H-charged samples. Deconvoluted red and green Gaussian curves are represented in the figure. The sum of these curves is denoted by the broken blue line, being in good agreement with the experimental data indicated by the black line. Despite the different hydrogen-gas pressures, the H-charged samples displayed two major peaks at around 600 K (Peak 1) and 750 K (Peak 2) except for the peak 2 in 0.7-MPa H-charged sample with lower temperature (683 K). This reason is not clear; however, may be attributed to lower hydrogen concentration at the corresponding trap-site. Fig. 1(f) presents the hydrogen concentrations of Peak 1, cH 1 , and Peak 2, cH 2 , as functions of the hydrogen-gas fugacity expressed as follows [23]: 𝑓 = 𝑝H exp
𝑏 ⋅ 𝑝H 𝑅𝑇
(1)
where, pH is the hydrogen-gas pressure, b is a constant, which represents the finite volume of the hydrogen molecules (1.584 × 10−5 m3 /mol), R is the universal gas constant (8.314 J K−1 mol−1 ) and T is the absolute temperature (K). As illustrated by the red plots, the values of cH 1 demonstrated a linear correlation with the square root of fugacity, f 1/2 , following the Sieverts’ Law derived from the thermodynamic equilibrium [24]. The proportional constant, defined as the solubility of hydrogen at a given temperature, was determined to be 4.97 wt ppm/MPa1/2 . In contrast, the values of cH 2 remained almost constant regardless of the fugacity, leading to the assertion that Peak 2 hydrogen corresponds to some sort of trapped one. Based on the equilibrium hydrogen-trapping theory, the trap-site occupancy, 𝜃 x , can be quantified by the following
Fig. 1. TDA profiles of non-deformed pure Ni that was H-charged in (a) 100 MPa-, (b) 50 MPa-, (c) 11 MPa- and (d) 0.7 MPa-hydrogen gas, as well as H-free, nondeformed pure Ni (e). The profiles were deconvoluted into two Gaussian peaks, designated as Peak 1 and Peak 2, with the hydrogen contents, cH , of each peak provided in (f).
K. Wada, J. Yamabe and H. Matsunaga
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equation [25]: 𝐸 𝜃X = 𝜃L exp b 1 − 𝜃X 𝑅𝑇
(2)
where, 𝜃 L is the lattice-site occupancy (𝜃 L << 1) and Eb is the binding energy of a trap-site (kJ/mol). Based on the simple assumption that one trap-site can trap one hydrogen atom and that the O-site is the only possible location, 𝜃 L is equivalent to the lattice-hydrogen concentration in at. ppm. In the determination of trap-site occupancy, it was presumed that an equilibrium state was achieved in the specimens since, in this study, they were aged at the room temperature (T ≈ 300 K) for more than three hours before starting tensile testing after the H-charging. In fact, Lassila et al. reported that solute hydrogen in pure Ni can reach an equilibrium state in approximately 103 s at 300 K after a high-temperature H-charging [13]. Fig. 2(a) presents the trap-site occupancy as a function of binding energy for each H-charging pressure, where the measured values of cH 1 were used as 𝜃 L . In order to determine the binding energy of the trapsite, the relationship between 𝜃 x and cH 2 were obtained as a function of Eb by using Eq. (2). Fig. 2(b) shows the 𝜃 x ‒ cH 2 relation in the case of Eb = 20.6 kJ/mol, where the solid line represents the regression. This Eb value was optimized so that 𝜃 x becomes proportional to cH 2 , i.e., a coefficient of determination, R2 , was maximized (cf. Fig. 2(c)). The R2 is defined by the following equation: ∑𝑛 2 (𝜃X,𝑖 − 𝜃̂X,𝑖 ) 𝑅2 = 1 − ∑𝑖=1 (3) 𝑛 ̄ 2 𝑖=1 (𝜃X,𝑖 − 𝜃X,𝑖 ) where, n is the number of data, 𝜃 x, i is the trap-site occupancy obtained from experiment in data number of i, 𝜃̄x is the mean value of trap-site occupancies and 𝜃̂x, 𝑖 is the trap-site occupancy calculated from a regression line in data number of i. Above result suggests that the second peak corresponds to hydrogen trapped at an unidentified trap-site with a binding energy of ≈ 20 kJ/mol. In the case of pure Ni, several potential trap-sites have been proposed in the existing literature, including the actual GBs (Eb = 21– 29 kJ/mol) [26–29], vacancies and their clusters (23–50 kJ/mol) [30– 32], dislocation cores (32 kJ/mol) [29], as well as interstitial sulfur (3– 4 kJ/mol) [33]. Among these trap-sites, the trap-energy determined in this study corresponded to that of the actual GBs or vacancies. Fig. 3(a–f) showcase the two-dimensional mappings of SIMS signals for 1 H and 32 S, with GB overlap, as taken via electron backscatter diffraction (EBSD) pattern analysis (indicated by red and white lines). Red lines represent ‘special’ (3 ≤ Σ ≤ 29) boundaries whereas white ones represent random boundaries [34]. As was already reported [17,18], the 1 H signals were spotted along GBs (indicated by black arrows in Fig. 3(a–c)) in many cases, thus clearly confirming that the Peak 2 hydrogen detected in the TDA measurement was trapped there. In several previous researches, it has been reported that the ‘special’ boundaries preferentially trap the hydrogen [18,34]. Indeed, hydrogen-trapping phenomena varied depending on GBs also in the present study; however, no clear difference in the trapping capability between ‘special’ and random boundaries was detected (Fig. 3(a–c)). Additionally, two-dimensional SIMS mapping offered visual confirmation that some of the hydrogen trapped along GBs existed at the same location as sulfur (indicated by white arrows in Fig. 3(a–f)). A segregant such as IG sulfur has been proposed in past studies to behave as a trap-site [14,35], and a similar accumulation of deuterium around IG sulfur was also previously reported [17]. Thus, although its binding energy has not yet been clarified, IG sulfur undeniably assists with hydrogen accumulation along GBs, with the trapping of hydrogen around the segregated sulfur [14,35]. Fig. 4(a) and (b) demonstrates the stress‒stroke curves of H-free and H-charged specimens, as well as the TFS values as functions of hydrogengas fugacity. Based on the data, two notable facts can be immediately ascertained. First of all, flow stresses increased with each increase in fugacity. Secondly, in the specimens H-charged at 11, 50 and 100 MPa, the RA and TFS values were nearly constant (RA of 0.05–0.07 and TFS of
Fig. 2. (a) Trap-site occupancy as calculated by Eq. (2). (b) Relationship between trap-site occupancy and Peak 2 hydrogen content for the binding energy Eb of 20.6 kJ/mol. (c) Coefficient of determination, R2 , as a function of binding energy.
228–265 MPa, respectively) and the fracture surfaces were covered only with IG facets (Fig. 4(d) and (e)). On the other hand, the 0.7-MPa Hcharged specimen failed after plastic instability, with an RA of 0.57 and TFS of 448 MPa. Furthermore, the IG fracture surface was discovered in the limited region of the specimen (Fig. 4(c)). Specifically, the degree of HE in pure Ni was not dependent on the lattice-hydrogen concentration
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Fig. 3. Two-dimensional mappings of 1 H (a–c) and 32 S (d–f), as analyzed via SIMS. Red and white lines represent ‘special’ (3 ≤ Σ ≤ 29) and random grain boundaries examined using the EBSD technique, respectively. In (a–f), white and black arrows indicate the hydrogen trapped at IG sulfur and the hydrogen trapped along GB. These mappings should be compared with (g–i) showing inversed pole figure maps around the observation areas indicated by red rectangles.
Fig. 4. (a) Stress–stroke curves of H-free and H-charged pure Ni. (b) Peak 2 hydrogen content and true fracture stress as functions of hydrogen-gas fugacity. (c–e) Fracture surface morphologies of H-charged pure Ni in 0.7 MPa (c), 50 MPa (d) and 100 MPa (e) hydrogen gas.
when gas pressure exceeded 11 MPa; the magnitude of HE was rapidly mitigated with the decrease in gas pressure from 11 to 0.7 MPa. These tendencies correlate with the concentration of trapped hydrogen, cH 2 (Fig. 1(f) and (b)), leading to the assertion that the extent of degradation is attributed to the hydrogen trapped along GBs, as associated with the Peak 2 hydrogen. This assertion is well-consistent with the conclusion of our previous research, demonstrating that initially-segregated hydrogen dominates the hydrogen-induced ductility loss [16].
In addition, based on the SIMS visualization of the IG sulfur along GBs, the impact of segregated sulfur on hydrogen-induced degradation must be considered. The segregated sulfur along GBs was reported to play two possible roles in the enhancement of hydrogen-induced IG fracture [14,35]. Firstly, as sulfur itself is well-known IG segregant to decrease GB strength [36,37], the IG sulfur possibly increases the propensity for IG fracture, in tandem with the hydrogen-induced IG cracking. Secondly, as obtained in the present SIMS observation, the IG sulfur
K. Wada, J. Yamabe and H. Matsunaga
may attract the hydrogen atoms [14,17,35], thereby making it easier to fulfill the critical hydrogen concentration required at GBs in order to drive the hydrogen-induced IG cracking. Consequently, through a combination of these two roles, it has been inferred that segregated sulfur and hydrogen synergistically weaken the GB strength [14,35]. Conclusions The present paper highlighted the hydrogen-trapping phenomenon and its quantitative contribution to the hydrogen-induced IG fracture in pure Ni. The TDA profiles revealed the existence of a trap-site with a binding energy of ≈ 20 kJ/mol, corresponding to that of the GBs or vacancies reported in existing literature. Visualization via SIMS measurement clearly indicated that this hydrogen was trapped along GBs, primarily accompanied by IG sulfur. Furthermore, from the viewpoints of the RA and TFS obtained in the SSRT tests, the degree of HE was dependent on the trapped hydrogen. Therefore, it can be concluded that the HE of pure Ni is dominated by the hydrogen existing along GBs and not by the lattice-hydrogen concentration. Declaration of Competing Interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. Acknowledgments A part of this study (EBSD analysis) was conducted with support from the Advanced Characterization Platform of the Nanotechnology Platform Japan, sponsored by the Ministry of Education, Culture, Sports, Science and Technology (MEXT) of Japan. The authors also wish to thank the staff of the Cameca Division – Ametek Co., Ltd., and Kobe Material Testing Laboratory Co., Ltd., for providing technical support and helpful information in relation to the SIMS analysis. Funding This research did not receive any specific grant from funding agencies in the public, commercial, or not-for-profit sectors. References [1] A.R. Troiano, The role of hydrogen and other interstitials in the mechanical behavior of metals, Trans. ASM 52 (1960) 54–80. [2] R.A. Oriani, P.H. Josephic, Equilibrium aspects of hydrogen-induced cracking of steels, Acta Metall. 22 (1974) 1065–1074, doi:10.1016/0001-6160(74)90061-3. [3] M. Nagumo, Hydrogen related failure of steels – a new aspect, Mater. Sci. Technol. 20 (2004) 940–950, doi:10.1179/026708304225019687. [4] S.K. Lawrence, Y. Yagodzinskyy, H. Hänninen, E. Korhonen, F. Tuomisto, Z.D. Harris, B.P. Somerday, Effects of grain size and deformation temperature on hydrogenenhanced vacancy formation in Ni alloys, Acta Mater. 128 (2017) 218–226, doi:10.1016/j.actamat.2017.02.016. [5] R. Kirchheim, Reducing grain boundary, dislocation line and vacancy formation energies by solute segregation. I. Theoretical background, Acta Mater. 55 (2007) 5129– 5138, doi:10.1016/j.actamat.2007.05.047. [6] R. Kirchheim, Reducing grain boundary, dislocation line and vacancy formation energies by solute segregation. II. Experimental evidence and consequences, Acta Mater. 55 (2007) 5139–5148, doi:10.1016/j.actamat.2007.05.033. [7] I.M. Robertson, P. Sofronis, A. Nagao, M.L. Martin, S. Wang, D.W. Gross, K.E. Nygren, Hydrogen Embrittlement Understood, Metall. Mater. Trans. B 46 (2015) 1085– 1103, doi:10.1007/s11663-015-0325-y. [8] I.M. Robertson, The effect of hydrogen on dislocation dynamics, Eng. Fract. Mech. 68 (2001) 671–692, doi:10.1016/S0013-7944(01)00011-X. [9] H.K. Birnbaum, P. Sofronis, Hydrogen-enhanced localized plasticity-a mechanism for hydrogen-related fracture, Mater. Sci. Eng. A 176 (1994) 191–202, doi:10.1016/0921-5093(94)90975-X. [10] I.M. Robertson, T. Tabata, W. Wei, F. Heubaum, H.K. Birnbaum, Hydrogen embrittlement and grain boundary fracture, Scr. Metall. 18 (1984) 841–846, doi:10.1016/0036-9748(84)90407-1.
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