2024 Volume 65 Issue 2 Pages 177-183
To investigate the effect of the microstructure adjacent to the grain boundaries on the mechanical properties and hydrogen embrittlement susceptibilities in the Al–Cu base 2219 alloy, alloy specimens were solution-treated and then aged at 100°C, 130°C, and the usual aging temperature of 190°C, to control the alloy microstructure in the vicinity of grain boundaries. The slow strain rate technique was conducted on the specimens in humid air and dry nitrogen gas environments to evaluate the effect of environmental hydrogen on them. Transmission electron microscopy was used to measure the grain boundary precipitate size and precipitate-free zone width of the specimens. Thermal desorption analysis was conducted on the gauge sections of the fractured specimens to evaluate the trapping sites and amount of hydrogen desorbed. The specimens aged below 190°C had finer grain boundary precipitates than those aged at 190°C. The test environment did not affect the specimen strength under any of the aging conditions 100°C, 130°C, and 190°C. Some samples had intergranular fractures on their entire fracture surfaces, irrespective of the test environment. The thermal desorption analysis results showed no significant difference between the hydrogen emission spectrum and the amount of hydrogen released within each temperature range. Thus, hydrogen embrittlement does not occur in the 2219 alloy, irrespective of the characteristics of its microstructure adjacent to the grain boundaries.
Hydrogen embrittlement (HE), the degradation of mechanical properties of a material due to its hydrogen content, is a disadvantage of metallic materials. HE in aluminum alloys, such as Al–Zn–Mg(–Cu) alloys, is a well-known problem. In Al–Zn–Mg(–Cu) alloys, HE exhibits the following characteristics: (1) Stress corrosion cracking (SCC) of Al–Zn–Mg(–Cu) alloys is considered to be caused by crack initiation due to the dissolution of solute atomic defects (anodic dissolution) and crack propagation due to HE.1–7) (2) Hydrogen is generated through the reaction between water vapor and the unoxidized aluminum surfaces of the alloys during slow-strain-rate deformation.8,9) (3) The hydrogen produced by the reaction leads to a significant decrease in the ductility of the alloys in humid air (HA) environments, while no such decrease in the ductility of the alloys occurs in high-pressure hydrogen gas atmospheres.10–18) (4) Hydrogen embrittlement susceptibility (HES) is low for solution-treated alloys, is maximum for underaged alloys, and is low after peak-aged and overaged alloys in that order.16,19) (5) The fracture surfaces caused by HE can exhibit either intergranular fracture (IGF) or quasi-cleavage fracture modes.11,16,17) In the case of smooth IGFs, elongation is low, and this type of fracture is often observed in HE. In contrast, when IGFs occur but elongation is high, dimples form on the IGF surface, distinguishing it from the fracture surface caused by HE.10,11,16,20,21) (6) Coarsening the size of grain boundary precipitates (GBPs) or causing grain boundary waviness due to strain-induced grain boundary migration reduces HES.16)
By contrast, Al–Cu(–Mg) alloys have a low resistance to SCC but are known for their resistance to HE. The HE characteristics of Al–Cu(–Mg) alloys can be listed as indicated below: (1) Crack initiation and propagation during SCC are both caused by anodic dissolution.1,2) (2) HES is low in solution-treated, underaged, peak-aged, and overaged alloys.10,22) (3) Hydrogen has no significant effect on the fracture surfaces of the alloys.22,23) (4) The Al–Cu(–Mg) alloy has a lower HES than the Al–Zn–Mg(–Cu) alloy when the strength levels of the two alloys are the same.10) Whether hydrogen can penetrate Al–Cu(–Mg) alloys and if so why HE does not occur in them have not been clarified so far. Only general aging conditions have been specifically evaluated for Al–Cu(–Mg) alloys so far in terms of their HES. To enable their comparison with Al–Zn–Mg(–Cu) alloys and discuss the HE mechanism of Al–Cu alloys in detail, their low-temperature aging in the presence of fine microstructures adjacent to GBs should be evaluated.
The following mechanisms have been employed to explain HE in aluminum alloys: hydrogen-enhanced decohesion (HEDE), hydrogen-enhanced localized plasticity (HELP), and hydrogen-enhanced strain-induced vacancy (HESIV) mechanisms.8) However, no unified view of the HE mechanism of aluminum alloys has yet been expressed. Thus, although many empirical findings on HE in aluminum alloys have been presented, only few efforts have so far been made to theoretically link those findings. In particular, many reports explaining why alloys such as Al–Zn–Mg(–Cu) alloy are more prone to HE than alloys such as Al–Cu(–Mg) alloy have been presented. However, only a few reports elucidate the reasons why alloys such as Al–Cu(–Mg) alloy, are not prone to HE. Understanding the HE resistance mechanism in the Al–Cu(–Mg) alloys would expand the design and use of aluminum alloys, as well as other alloys like steel and magnesium alloys.
In this study, we focused on the microstructure adjacent to the GBs in the Al–Cu(–Mg) alloys, which can be used as a HE parameter. We use the slow strain rate technique (SSRT) to evaluate the change in the mechanical properties of the alloy due to hydrogen, and thermal desorption analysis (TDA) to determine the change in the level of hydrogen penetration and number of hydrogen trapping sites due to differences in the test environment and microstructure adjacent to the GBs, and quantified the effects of the GBs and microstructural factors, such as dislocations, on the HE in the alloy. The objective of the study was to determine the effect of the microstructure adjacent to the GBs on the HE of aluminum alloys.
The chemical composition of the 2219 alloy used in this study is listed in Table 1. The samples with a thickness of 10 mm were first solution-treated at 535°C for 3.6 ks (1 h) and then water-quenched. After they were solution-treated, 90% cold rolling was conducted to obtain 1 mm thick plates required for tensile specimen preparation. The plates were cold-rolled with a 10% rolling reduction per 1 pass and water-cooled after each pass. The tensile test specimens were cut out from the 1 mm thick plates employing electrical discharge machining. The longitudinal directions of the tensile test specimens were aligned with the rolling direction. The specimens were thereafter re-solution-treated at 535°C for 3.6 ks and water-quenched. They were aged at 100°C for 864 ks (10 d), 130°C for 86.4 (24 h), 864 (10 d), and 3456 ks (40 d), and at 190°C for 93.6 (26 h) and 129.6 ks (36 h). Electric furnaces (FT-105, FT-01P, FT-01X, FULL-TECH, and NHK-170, Nitto Kagaku) were used for the heat treatment of the specimens. The aging conditions were determined using general T6 data (190°C), preliminary experiments, and previous reports (100 and 130°C).24) For underaging, peak aging, and overaging, a temperature of 130°C was selected as the criterion, while 100°C was selected specifically for underaging. As solution-treated (As ST) and aged specimens were used for the tensile tests at various. Prior to the tensile test, wet polishing (#400–800), and alkaline washing with a 10% NaOH solution for 60 s, and dismutation with a 10% HNO3 solution for 60 s were conducted. SSRT was conducted on the test specimens at an initial strain rate of 1.67 × 10−6 s−1. The environmental conditions used were HA and dry nitrogen gas (DNG). The relative humidity was ≥95% in HA and ≤5% in DNG. The test temperature was 25°C. Scanning electron microscopy (SEM, S-3400N, Hitachi High-Tech), and field emission scanning electron microscopy (FE-SEM, SU5000, Hitachi High-Tech) were used for fracture surface observation. The Vickers hardness test was conducted with a 10 N for 15 s. The metallographic structures of the specimens were observed using an optical microscope. Transmission electron microscopy (TEM, JEM-2010, JEOL) was used to evaluate the specimen microstructures adjacent to the GBs. The TEM samples were mechanically polished to #80–2000, punched to a diameter of 3 mm, electropolished, and cleaned with ethanol. The nitric acid:methanol:water volume ratio used in the electropolishing solution was 4:5:1. The TDA was performed in each test environment after the tensile test to assess the amount of hydrogen that entered the specimens during their tensile deformation. Hydrogen desorption was analyzed using gas chromatography (SGHA-P2, NISSHA FIS) with argon as the carrier gas. The gas flow rate was maintained at 20 cc/min. The specimen was heated from 25 to 540°C at a heating rate of 100°C/h. Before TDA, the specimens were wet-polished (#400–1500) and cleaned using acetone. The desorption rate and the amount of hydrogen in a specimen were normalized per unit mass of the specimen measured before the TDA. The SSRT allowed for the continuous rupture of the dense oxide film and exposure of the unoxidized aluminum surface. During the SSRT conducted with monotonic loading at slow strain rates, the constant reaction of the unoxidized aluminum surface with the surrounding water vapor produced atomic hydrogen which continuously penetrated the aluminum alloy. The chemical equation expressing the formation of atomic hydrogen from the unoxidized aluminum surface and water vapor is as follows:
| \begin{equation} \text{2Al(s)} + (3 + X)\text{H$_{2}$O(g)} \to \text{Al$_{2}$O$_{3}$} \cdot \text{$X$H$_{2}$O(s)} + \text{6H} \end{equation} | (1) |
The optical micrographs of the specimens after they were aged at 130°C for 86.4 ks are shown in Fig. 1. An equiaxed grain structure and ∼2 µm second-phase particles could be seen in the micrographs. The second phase particles were reported to have come from a Al20Cu2Mn3 compound.24) The average grain size of the specimens was almost 16 µm, and their subsequent aging had not affected it. Thus, the effect of grain size in each specimen could be disregarded in the hardness tests, tensile tests, and TDA.
Optical micrograph of the specimen aged at 130°C for 86.4 ks.
The Vickers hardness values of the specimens at the different aging conditions are shown in Fig. 2. The hardness of the specimens aged at 100°C and 130°C tends to be superior to the that of specimens aged at 190°C.
Vickers hardness values of the as solution-treated specimen and other specimens aged under different conditions.
The stress vs. strain curves obtained from the SSRTs at an initial strain rate of 1.67 × 10−6 s−1 in the HA and DNG environments are shown in Fig. 3. The specimens aged at low temperatures, such as 100°C and 130°C, showed higher strength and ductility than those aged at 190°C. The shape of the stress–strain curve depended on the aging conditions. The as-ST specimen (Fig. 3(a)) and the specimen aged at 130°C for 86.4 ks (Fig. 3(c)) showed a gradual increase in their stress and high ductility. In the specimen aged at 100°C for 864 ks (Fig. 3(b)) and that aged at 130°C for 864 ks (Fig. 3(d)), the stress reached its maximum value earlier than in the specimens of Figs. 3(a) and (b), and their local deformation after tensile strength is less than that in the specimens of Figs. 3(a) and (c). In the specimens of Fig. 3(e), which was aged at 130°C for 3456 ks, the stress reached its maximum value at a smaller strain than in the specimens of Figs. 3(c) and (d). In the specimen of Fig. 3(f), which was aged at 190°C for 93.6 ks and the specimen of Fig. 3(g), which was aged at 190°C for 129.6 ks, the stress reached its maximum value earlier than in the specimen of Fig. 3(e), and the stress in it started to decrease after reaching its maximum value, which is significant, indicating large local deformation. The 0.2% proof stress and ultimate tensile strength (UTS) exhibited almost no difference between their values in the HA and DNG environments. In some specimens, the elongation to failure was slightly greater in the DNG environment than in the HA environment. This result is consistent with the results of previous studies that showed a slight decrease in the ductility of the specimens in an HA environment.10,22)
Stress vs. strain curves of the specimens in humid air and dry nitrogen gas environments. (a) As-ST, (b) 100°C–864 ks, (c) 130°C–86.4 ks, (d) 130°C–864 ks, (e) 130°C–8640 ks, (f) 190°C–93.6 ks, and (g) 190°C–129.6 ks specimens.
The relationship between the HES index and 0.2% proof stress is shown in Fig. 4. The HES index, I(δ), is defined as follows:
| \begin{equation} \text{I}(\delta) = (\delta_{\text{DNG}} - \delta_{\text{HA}})/\delta_{\text{DNG}} \end{equation} | (2) |
where δHA and δDNG are the elongation to failure in the HA and DNG environments, respectively. If the HES index was greater than zero, the ductility in the HA environment would have been lower than that in the DNG environment. Therefore, the higher the HES index, the higher the HE caused by environmental hydrogen. Figure 4 shows that the HES index of the specimens is low under all aging conditions and strength levels. Al–Cu(–Mg) alloys subjected to general aging conditions have been found to possess a low HES, which is confirmed by the results of this study.10,22) The as-ST and underaged Al–Zn–Mg(–Cu) alloys are known to be more susceptible to HE than the peak-aged alloys.1,2) Because the HES index of the Al–Zn–Mg(–Cu) alloys is in the range of 0.5–0.9, the HES of the Al–Cu alloy specimens used in the study is lower than that of the Al–Zn–Mg(–Cu) alloys. This result indicates that the tensile properties of the Al–Cu alloys are not significantly degraded by environmental hydrogen, regardless of the differences in their strengthening and deformation behaviors subsequent to their yielding.
Relationship between the hydrogen embrittlement susceptibility index and 0.2% proof stress of the specimens.
The bright-field TEM image and selected area electron diffraction (SAED) patterns of the microstructure in the vicinity of the GBs of the specimens aged at 130°C for 86.4 ks, 130°C for 864 ks, and 190°C for 93.6 ks are shown in Fig. 5. Typical aging conditions were selected for the TEM specimens based on the differences in their mechanical behaviors, as shown in Fig. 3. The precipitation process of the Al–Cu alloy is as follows:25,26)
| \begin{align} & \alpha - \text{Al} \to \text{Supersaturated solid solution} \to\\ & \quad \text{GP(1) zone} \to \text{GP(2) zone (or $\theta''$)} \to \theta' \to \theta \end{align} | (3) |
The precipitation phases of the specimens can be categorized into three zones based on their aging conditions. In the specimens aged at 130°C for 86.4 ks, Guinier-Preston (GP)(1) zone was formed. In the specimens aged at 100°C for 864 ks, and at 130°C for 864 and 3456 ks, GP(2) zone (or θ′′) and θ′ were formed. In the specimens aged at 190°C for 93.6 and 129.6 ks, θ′ was formed. The GBPs consisted of the respective θ phases.24–26)
Bright-field transmission electron microscopy image and selected area electron diffraction patterns of the microstructures of the specimens in the vicinity of their grain boundaries. (a) and (b) 130°C–86.4 ks, (c) and (d) 130°C–864 ks, and (e) and (f) 190°C–93.6 ks specimens.
The average size of the GBPs and the width of the PFZ are shown in Fig. 6. Based on Fig. 5, the average size of the GBPs was defined in the direction of the GBs, while the PFZ width was defined in the direction perpendicular to the GBs. The specimens aged at 130°C have a finer microstructure adjacent to the GBs than the specimens aged at 190°C. The finest microstructure was observed in the 130°C–86.4 ks specimen, and the average size of its GBP was 76 nm. By contrast, the average size of the GBPs in the 190°C–93.6 ks specimen was 270 nm. The smaller the GBP, the narrower the PFZ width. According to Itoh et al., the average size of the GBPs in the 7475 alloy T6 temper, a commercially available Al–Zn–Mg(–Cu) alloy, is 15 nm.27,28) This result shows that the size of the GBPs and the width of the PFZ in the Al–Cu alloys aged at 130°C are closer to those of the Al–Zn–Mg(–Cu) alloy than to those of the Al–Cu alloys aged at 190°C.
Relationship between the average size of the grain boundary precipitates of the specimens and their precipitate-free zone widths.
The results of this study indicate that unlike the test environments, the aging conditions affect the mechanical properties of Al–Cu alloys. Thus, the effect of the microstructure adjacent to the GBs on the fracture behavior of the alloys would be considered when discussing the effect of the aging conditions on them. The TDA results would be considered when discussing why environmental hydrogen did not affect the mechanical properties of the alloys.
The secondary electron SEM images of the fracture surfaces of the specimens in the HA and DNG environments, obtained after conducting the SSRT are shown in Fig. 7. Transgranular ductile fractures were observed in the as-ST specimens. IGFs with dimples could be observed on the entire surfaces of the 100°C–864 ks and 130°C–864 ks specimens. In the other specimens, the fracture surfaces were a mixture of transgranular ductile fractures and IGFs with dimples. The finer the microstructure adjacent to the GBs, the smaller the dimples on the IGFs.21) An IGF with dimples is macroscopically brittle but microscopically ductile, and will not always deteriorate.21) The size of the dimples almost corresponds to the size of the GBPs. Larger and deeper dimples were also observed in the specimens, which indicates that a larger local deformation has occurred before fracturing, as shown in the stress–strain curves of Fig. 3. The fracture morphology did not change with the change in the test environment. Hence, environmental hydrogen will not promote void formation, growth, or coalescence near GBs.
Fracture surface observations of the specimens in humid air and dry nitrogen gas environments. (a) and (b) 100°C–864 ks, (c) and (d) 130°C–86.4 ks, and (e) and (f) 130°C–864 ks specimens.
The hydrogen desorption rates and spectra of different temperature ranges obtained from the TDA performed on the parallel sections of the specimens fractured after conducting the tensile test in the HA and DNG environments are shown in Figs. 8 and 9. A model of the hydrogen trapping sites is shown in Fig. 8. The main trapping site of hydrogen in the aluminum present in the specimens is in the free surfaces of their interiors, such as voids, blisters, and pores.28,29) However, trapping sites are also present in vacancies, dislocations, GBs, and phase interfaces.28,29) In Fig. 8, hydrogen emission peaks are present at 90°C, 300°C, 400°C, and 500°C. At 90°C, 300°C, and 400°C, hydrogen trapping sites are presumed to be in surface oxide films, and GBs, dislocations, respectively and in free surfaces, such as blisters and pores, at temperatures above 500°C.28–31) The aging condition affects the amount of hydrogen desorbed at 90°C. The precipitation microstructure could have some effect on the hydrogen desorption behavior; however, the reason for which could not be identified in the study. No difference in the amount of hydrogen trapped at the trapping sites within the two test environments could be observed. The amounts of hydrogen trapped in the surface oxide films, GBs, and dislocations are not affected by the test environment. According to Itoh et al., no differences could be discerned in the hydrogen desorbed from the 7075 alloy specimens subjected to 7.0% tensile deformation in HA and DNG environments.19,33) The amount of hydrogen absorbed by the specimens from a humid environment during their tensile deformation is much smaller than the amount of hydrogen within it. Therefore, we could not conclude whether hydrogen was absorbed or not by the specimens used in the study during their tensile deformation in the HA environment. The increased amount of hydrogen emitted by the specimens at temperatures above 450°C and the difference between that and the amount of hydrogen emitted between 400°C and 450°C is presented in Fig. 9; however, the hydrogen emitted at temperatures above 450°C will not be related to the tensile deformation at room temperature (25°C) because it is strongly trapped.
Hydrogen desorption rates of the specimens after they were fractured in humid air and dry nitrogen gas environments, calculated using thermal desorption analysis and the model of hydrogen trapping sites and hydrogen desorption temperature of the specimens. (a) As-ST, (b) 100°C–864 ks, (c) 130°C–86.4 ks, (d) 130°C–864 ks, (e) 130°C–3456 ks, (f) 190°C–93.6 ks, and (g) 190°C–129.6 ks specimens.
Hydrogen desorption amounts of the specimens after they were fractured in humid air and dry nitrogen gas environments, calculated using thermal desorption analysis. (a) As-ST, (b) 100°C–864 ks, (c) 130°C–86.4 ks, (d) 130°C–864 ks, (e) 130°C–3456 ks, (f) 190°C–93.6 ks, and (g) 190°C–129.6 ks specimens.
The effect of the microstructure adjacent to the GBs on the mechanical properties and HES of Al–Cu alloys was investigated. The TEM observations confirmed that the average size of the GBPs and the width of the PFZ were more refined in the specimens aged at 130°C than those aged at 190°C. The SSRT confirmed that the strength and ductility of the alloys increased during low-temperature aging. A comparison between the strength and ductility values of the specimens aged in HA and DNG environments revealed that the mechanical properties of the specimens were similar in both environments. The HES of the specimens remained low under all aging conditions. Fracture surface observations showed IGFs with dimples were present on the entire or part of the surfaces of the specimens aged in both HA and DNG environments. However, the fracture morphologies of the specimens were not affected by the test environment. The TDA results showed no significant difference between the hydrogen desorption rate and spectra in the two different test environments.
The authors are grateful to (a) JST, the establishment of university fellowships toward the creation of science technology innovation, Grant Number JPMJFS2105, and (b) Japan Aluminum Association, Reiwa 4th Aluminum Research Grant Program. The first author is also grateful to his JST fellowship program mentor, Associate Professor Keitaro Horikawa, at Osaka University.
