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Hydrogen-assisted degradation of some non-ferrous metals and alloys
Journal of Materials Processing Technology 109 (2001) 206±214
Hydrogen-assisted degradation of some non-ferrous metals and alloys A. Zielinski Department of Materials Science and Engineering, Technical University of Gdansk, G. Narutowicza 11/12, 80-952 Gdansk, Poland
Abstract The hydrogen degradation of some non-ferrous metals and alloys, and its dependence on microstructure and environment is reviewed. The hydrogen-enhanced localised plasticity, formation and decomposition of brittle hydrides, and appearance of hydrogen and methane bubbles are discussed as possible mechanisms of hydrogen embrittlement. # 2001 Elsevier Science B.V. All rights reserved. Keywords: Hydrogen; Embrittlement; Copper; Aluminium; Nickel; Titanium; Zirconium; Niobium; Vanadium; Intermetallics
1. Introduction Non-ferrous metals and alloys are industrially relevant materials for many applications including hydrogen evolution and absorption. The hydrogen entry into metal or alloy may result in degradation of its mechanical properties to the extent dependent on the environment, stress state, alloy microstructure, etc. The present paper reviews the recent results obtained for this groups of metallic materials, especially on degradation of their microstructure and mechanical properties, conditions which are critical for hydrogen embrittlement to occur, and possible atomic mechanisms of hydrogen delayed cracking. 2. Copper and its alloys Hydrogen solubility in bulk copper is extremely low, about 10ÿ2 ppm, and no hydride formation has been reported for hydrogen charged copper. The properties of bulk copper are in¯uenced by the hydrogen presence: the tensile yield stress increased between 77 and 200 K after high temperature gas hydrogen charging, and hardness increased 3±5% following cathodic charging [1]. At industrial applications the high supersaturation of copper with hydrogen is not expected except electrolytic plating of copper layers in which the hydrogen content may reach even 200 to 1000 ppm [2]. Such high hydrogen content was associated with loss in plasticity of copper layer [2,3]. The decrease in plasticity became reversible only in a small part [2]. The precipitation of hydrogen resulting in bubbles and voids was observed following the gas charging in bulk copper [4] and in copper layer [2,3,5]. Among three types
of voids the smallest originated from hydrogen agglomeration, and the larger, present along the grain boundaries, were pressurised gas bubbles with dislocations punched out of the void [2]. The appearance of voids was postulated as a cause of the loss in plasticity. Thermal charging of Cu±Al alloys increased the microhardness of annealed material. The effect was accompanied by an increase in dislocation density [6]. Effect of hydrogen on mechanical behaviour of a deformation processed Cu±20%Nb composite was investigated in [7]. No substantial degradation, except a small decrease in uniform elongation, was observed in the composite (on the contrary to high susceptibility of niobium). The alumina dispersed copper alloys were investigated after their cathodic and gas charging. The loss in ductility reached about 22% but the increase in Al content had a bene®cial effect on the plasticity. The increase in annealing temperature of cold worked specimens decreased the degree of embrittlement due to decrease in number of reversible hydrogen trapping sites [8]. 3. Aluminium and its alloys Pure aluminium was shown, on the contrary to some previous reports, to be susceptible to the hydrogen presence. The specimens cathodically charged at 300 Amÿ2 were tested at strain rate 10ÿ4 sÿ1. All samples failed in a completely ductile manner. The cathodic charging did cause differences in the stress±strain behaviour: increase in the yield stress, slight increase in the ultimate tensile stress, and decrease in the strain to failure (Fig. 1). The increase in microhardness from 46 to 67 MPa was noticed [9,10].
0924-0136/01/$ ± see front matter # 2001 Elsevier Science B.V. All rights reserved. PII: S 0 9 2 4 - 0 1 3 6 ( 0 0 ) 0 0 7 9 7 - 4
A. Zielinski / Journal of Materials Processing Technology 109 (2001) 206±214
Fig. 1. Effect of hydrogen on failure strain of 99.99% aluminium [11].
In [11] the critical slip shear stress was reduced after hydrogen charging and the atomic binding energy of aluminium decreased about 11%. Water vapour accelerated the fatigue crack propagation in pure aluminium single crystals under mode II or I±II loading [12]. The susceptibility to hydrogen-related degradation of aluminium alloys was observed to depend on microstructure of an aluminium alloy and test environment. The 2090 Al±Li±Cu±Zr alloy after cathodic charging showed the linear decrease in fracture strain following increasing hydrogen content [13]. The hydrogen charging of 8090 Al±Li±Cu±Mg alloys increased the dislocation density and caused the intergranular failure, especially in the overaged condition [14] but did not affect the bending fatigue strength [15]. In [16] the hydrogen charging resulted in ductility loss (Fig. 2), increasing with increasing ageing time. The testing of the 2024 Al±Cu±Mg alloy disclosed severe degradation in overall plastic elongation. The embrittlement was observed after the onset of necking [17]. Reverse torsional fatigue tests of the alloy exposed to hydrogen and humid air showed that hydrogen absorption promoted multiple crack initiation and altered the fracture mode [18].
Fig. 2. Effect of charging time on the ductility loss of an Al±Li±Cu±Mg± Zr alloy [17].
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The 2124 Al±Cu±Mg alloy (2024 alloy with low Fe and Si) showed no signi®cant loss in ductility for any temper following hydrogen charging but lost about 25% RA when simultaneously strained and charged with hydrogen [19]. The fracture mode was not altered by hydrogen. The other fatigue testing of precharged specimens revealed an increase in fatigue life [20]. The Al±8%Mg alloy was observed to decrease its Young modulus, 0.2% proof stress and tensile stress following hydrogen charging [21]. An exposure of Al±Zn±Mg alloys to water vapour was observed to result in their embrittlement [22±28]. The crack velocity was strongly stress dependent at low stress intensities while at intermediate and high stress intensities there was no effect of stress on the crack growth rate [25]. The cracks were intergranular with a brittle appearance. The crack propagation rate increased with level of humidity, with no effect observed in dry gases, including dry hydrogen [25]. The loss in ductility was observed in tensile tests [23]. The solute hydrogen enhanced dislocation mobility and reduced the ¯ow stress. The cathodic polarisation disclosed similar embrittlement dependent on the alloy and its microstructure. The Al±Zn± Mg alloy was embrittled for all microstructures, and the Al± Zn±Mg±Cu alloy was resistant in the peak strength and overaged condition [29±31]. The effect of cathodic and gaseous hydrogen on fatigue properties was investigated in [22,28,32,33]. Experiments made on Al±Zn±Mg±Cu alloy showed the decrease in fatigue resistance [22,32] (Fig. 3) and in increase in the degree of transgranular crack initiation and propagation [32]. The hydrogen embrittlement was explained by different models,dependingonthealloy.IntheAl±Lialloythehydrogen degradation was caused by precipitation of internal LiH and LiAlH4 hydrides [34,35] which brittle fracture resulted in crack propagation [15]. In pure aluminium the enhanced dislocation mobility was observed [9] in the presence of hydrogen. Similar hydrogen-enhanced localised plasticity seems responsible for the hydrogen embrittlement in the
Fig. 3. Resolved shear stress versus cycles to failure for single crystals of Al±5.5Zn±2.5Mg±1.5Cu in air, and 0.5 M NaCl solution at corrosion and cathodic potential [33].
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other aluminium alloys, like Al±Zn±Mg and Al±Zn±Mg±Cu [36], although formation of magnesium hydride [31,37,38] was also postulated to be a cause of the embrittlement. The appearance of the bubbles decreased the plasticity loss [23]. 4. Nickel and its alloys Numerous studies [39±45] showed that solute hydrogen can cause intergranular fracture of nickel and thereby reduce the ductility. The most characteristic feature is a change in fracture mode from ductile predominantly transgranular microvoid coalescence-type to intergranular fracture with increasing cathodic potential [44] and amount of hydrogen [39,40,42±44]. The intergranular fracture initially occurred as isolated microcracks on the surface exposed to hydrogen. The mode transition was explained by different hydrogen trapping: by grain boundaries, continuous incoherent precipitates or discrete incoherent precipitates [42]. A critical stage was the link-up of microcracks into a macrocrack which propagated, resulting in failure [45]. The critical hydrogen content at the grain boundary was between 0.04 and 0.10 (Fig. 4) [39]. The susceptibility to cracking was remarkably dependent on the presence of impurities within the grain boundary area: carbon decreased the extent of hydrogen embrittlement and sulphur increased the amount of degradation (Fig. 5) [42±44]. The fracture mode was also related to the presence of sulphur: at low sulphur content (0.03 monolayers at the grain boundary) nickel failed in microvoid coalescence-type manner, and at high sulphur content (0.20 monolayers) in brittle intergranular manner [44]. The effect of impurities was explained by their in¯uence on hydrogen concentration which was required for intergranular embrittlement; segregated carbon increased the critical hydrogen concentration and grain boundary diffusivity, and sulphur decreased this value [41]. The hydrogen cracking was postulated to advance rather by the hydrogen-enhanced plasticity than by decohesion or hydride formation [46,47].
Fig. 4. The percent of intergranular fracture versus the grain boundary hydrogen concentration [40].
Fig. 5. Percent intergranular fracture versus grain boundary composition of nickel for several cathodic test potentials [45].
Hydrogen embrittlement of the austenitic Ni±Cr±Fe alloys was often observed following cathodic charging [48±55]. Tensile testing of the charged alloy with Ni content varying between 29 and 59% showed loss in elongation and ultimate tensile stress, with surface cracks appearing on the surface during subsequent ageing of the alloy [48]. The simultaneous charging with hydrogen and deformation in tension of the Inconel 600 76Ni±16Cr±8Fe alloy decreased markedly the fracture strain and changed the fracture mode to completely intergranular [49,50]. The exposure of the Alloy 600 in steam caused hydrogeninduced intergranular stress corrosion cracking [51±53]. The appearance of 30 and 50 nm bubbles along the grain boundaries ®lled in with high pressure methane was observed. The cracks were initiated by the nucleation of high density bubbles on the grain boundary under the combined action of the applied stress and high pressure methane formed from carbon in solution reacting with hydrogen injected by corrosion [52]. The necessary condition for IGSCC to occur was a presence of water/steam; no cracks were formed in dry hydrogen of 27.3 MPa pressure at 673 K [52]. The susceptibility was related to the stability of NiO in this environment [54]. For Fe±Ni±Co superalloys hydrogen was observed to signi®cantly reduce tensile true fracture strain (Fig. 6), slow crack growth thresholds and fracture toughness values [55± 57]. The decrease in fracture toughness was accompanied by a change in fracture mode from microvoid coalescence in the uncharged specimens to principally slip band fracture with some twin bands and ductile intergranular fracture in the hydrogen charged samples. The fracture initiated at matrix carbides [55]. The testing of Ni±Cu Monel alloys showed some loss in tensile strength and RA value [58,59]. The embrittlement decreased with increasing strain rate and grain size. Fracture in air was completely ductile and that in hydrogen brittle and intergranular. The gaseous hydrogen embrittlement was observed for amorphous metallic glass Fe40Ni38Mo4B18 [60]. The fracture
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Fig. 7. Ultimate tensile strength as a function of hydrogen concentration in Ti±10V±2Fe±3Al alloy [73].
Fig. 6. Effect of hydrogen on tensile true fracture strain of the IN903 alloy [56].
toughness of notched specimens was signi®cantly reduced at slower strain rates. The degree of embrittlement was similar to that of ferritic or unstable austenitic stainless steels. 5. Titanium alloys Titanium alloys are very complex materials in which two matrix phases, hcp a-phase more susceptible to hydrogen embrittlement with lower hydrogen solubility and diffusivity, and bcc b-phase with opposite properties are present. As much as four different hydrides may appear in titanium alloys: d fcc TiH1.5, e fct TiH2, g fct hydride and strain induced bcc hydride. Therefore, the hydrogen-assisted cracking is strongly related to the alloy microstructure [61]. Commercially pure titanium is very resistant to the embrittlement caused by hydrogen when tested in the form of ®ne-grained specimens at low-to-moderate strain rates in unaxial tension but it becomes susceptible in the presence of a notch, at low temperatures or high strain rates, or large grain sizes. The last effect was reported to be a consequence of an enhancement of both void nucleation and void link-up at large grain sizes/biaxial stresses [62,63]. Testing on oxygen-strengthened titanium demonstrated no effect of hydrogen on the fracture toughness but pronounced effect on impact resistance [64,65]. In testing under sustained load the room temperature rupture times were observed to decrease [66]. The testing of Ti±6Al±4V alloy containing 80 to 720 ppm of hydrogen showed that hydrogen promoted creep of this alloy at room temperature, markedly increasing both creep strain and rate in the primary stage [67]. There was an evidence of softening of the alloy due to hydrogen. The accelerated creep was also observed for the Ti±5Al±2.5Sn alloy [68]. The mechanism of cracking was often suggested to result from formation and decomposition of brittle hydride phases.
In the hcp. a-Ti±4%Al alloy the investigations in highvoltage electron microscope showed that in gaseous hydrogen environment at room temperature two fracture mechanisms could operate, depending on stress intensity: at low stress intensity the cracks propagated by repeated formation and cleavage fracture of hydrides, and at high stress intensity the fracture mode transition occurred at crack propagation rate which exceeded that at which hydride could form in front of the crack, and cracking was possible by hydrogen-enhanced localised plasticity process [61]. The hydride cracking was postulated for the other a and a b-Ti alloys [69]. The hydrogen embrittlement was observed to occur also in the b alloys, like the Ti±Mo± Nb±Al alloy [70] and Ti±V±Cr±Al±Sn alloy [71], and Ti±V± Fe±Al (Figs. 7 and 8) [72], well below hydrogen concentration required to hydride the b-phase. In these alloys the effect of hydrogen is through two primary mechanisms: hydrogen embrittlement due to the effects other than hydride formation, and stabilisation of the b-phase [72,73]. In the alloys, in which the hydride phase was formed, the variation of tensile stress, yield stress and ductility was shown to be proportional to an amount of hydride (Fig. 9) [74]. The effect of microstructure was substantial as shown for the Ti±8Al±1Mo±1V alloy undergoing different heat treatment: in the near-a alloys the cleavage-like fracture
Fig. 8. Reduction in area as a function of hydrogen concentration in Ti±10V±2Fe±3Al alloy [73].
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Fig. 9. Variation of ultimate tensile strength and yield stress with the hydrogen concentration, i.e. amount of hydride [75].
occurred, and in the a-b alloys an alternating extensive a cleavage and ductile rupture of the b ligaments became active [69,75]. For the Ti±6Al±6V±2Sn alloy fracture was characterised by extensive cleavage of the a grains separated by ductile rupture of the b ligaments from threshold to near failure under plane strain conditions, and the material fractured in the plane stress regions failed in a ductile mode [76]. The mechanism of cracking was related also on test environment. In gaseous hydrogen the hydrogen diffusion occurred via b-phase to a-phase, then hydride phase precipitated and crack advanced along a-b interface. During hydrogen charging or exposure in water solution the hydrogen diffusion, hydride formation and fracture took place in the a-phase [77,78]. 6. Zirconium alloys The ductility of the Zr±2.5Nb Zircaloy±2 alloy was shown to decrease with increasing hydrogen content (Fig. 10) [79]. The cracks initiated at hydride blisters which grew on the surface as a result of high pressure. The thermo-migration of
hydrogen to colder points and possible formation of solid zirconium hydrides were the critical steps in initiation of fracture. The critical applied stress was necessary to initiate delayed hydride cracking. The threshold for cracking initiation from a blister ranged between 10.7 and 15.4 MPa1/2 [80]. The amount of plastic strain needed to induce fracture in hydrides decreased as the average hydride length and the axiality of the stress increased. Regardless of the stress state, when hydride platelets had lengths larger than about 50 to 100 mm, little plastic deformation was needed to initiate hydride fracture [81]. The hydrogen embrittlement was a consequence of void nucleation due to the strain induced hydrogen fracture and of the subsequent void growth and void link-up [79]. The mechanism of hydride fracture appeared to involve the slip-induced nucleation of cracks which grew to a critical size under suf®ciently high normal stress. The critical size of a hydride was estimated at 2 to 4 mm in length [82]. The hydrogen degradation induced by brittle fracture of hydrides was also observed in Zr±1.5Sn Zircaloy±4 alloy [83]. The transition from ductile to brittle behaviour was suggested to be prevented by using the ®ne microstructures with elongated grain in the loading direction. The delayed hydride cracking was proved to be very sensitive to prior heat treatment [84]. The maximum crack growth rate was produced in a microstructure consisting of a continuous grain boundary phase b, and the crack velocity decreased progressively when this phase was decomposed and broken up. Rapid quenching which produced ®nely dispersed hydrides resulted in higher crack growth rates. The model of the delayed hydride cracking was proposed which can be applicable for the other metal-hydrogen system in which hydride may appear. In such systems hydrides must be present in a bulk, and hydrogen must be able to ¯ow to the crack tip from these bulk hydrides. The bulk hydrides must be in a dissolving mode, while the crack tip must be in a growing or nucleating mode. The delayed hydrogen cracking was assumed to be caused by preferential and repeated accumulation, and fracture of hydrides at stress raisers such as cracks, causing the cracks to slowly increase over a period of time. The crack tip hydride was said to grow to the critical length as a necessary condition for such type of cracking [85]. 7. Niobium alloys
Fig. 10. The in¯uence of hydrogen on the reduction of area at fracture (circles) and the true strain at fracture (triangles). The tensile tests made either parallel to the rolling direction (black symbols) or transverse to the rolling direction (light symbols) [80].
The Va group transition metals, like niobium, vanadium and tantalum were usually ductile down to 77 K but became brittle at low temperature when they contained even a small amount of hydrogen. For low hydrogen concentration one single minimum and for high hydrogen content the two peaks were observed in relation of tensile elongation on temperature (Fig. 11) [86,87]. No appreciable effect of hydrogen concentration up to 130 ppm on the yield stress was noticed [86].
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Fig. 11. The uniform strain as a function of temperature in single crystals Nb±H alloys at 80 ppm wt. H (D), 130 ppm (- - -) and 230 ppm (&).
In fatigue testing, the decrease in stress intensity range, DK, was observed [88]. The hydrogen amount needed to initiate the embrittlement was below critical value necessary to form a hydride. Further increase in hydrogen content had little effect on the fatigue mechanical properties but appeared to change the brittle fracture path. Various models were proposed for the hydrogen embrittlement. The cracking propagation was recently suggested to take place as a repeated process in which hydride would precipitate at an advancing crack tip with the aid of a local stress and fracture in a brittle manner, followed by the crack advancement [89]; the other model proposed the dislocation-hydrogen interaction as an important factor in hydrogen degradation [86,88]. The nature of two step ductile-brittle transition for bulk specimens was explained in several ways. They included: thermal hysteresis of precipitation and dissolution of hydride, decohesion or hydrogen-enhanced localised plasticity at higher temperature and hydride cracking at lower temperature, or ordering of hydrogen within the hydride phase that would lead to partial recovery of ductility between two transition temperatures [86,88].
Fig. 12. Effect of carbon and combined carbon and hydrogen on the temperature dependence of ductility in vanadium [92]. Circles are for pure carbon (black 0.012%C, light 0.20%C), and squares for 0.21%C and 0.27%H (black) or 0.50%H (light).
about 0.6% hydrogen had a moderate strengthening effect on the yield stress but the low temperature decrease in plasticity in suf®ciently hydrogen charged specimens was observed [93]. Similar behaviour was reported for the V±Ti alloys [94,95]. Hydrogen embrittlement was observed as decrease in plasticity (RA) in both the hydride forming and some nonhydride forming V±Ti alloys. The transgranular initiation of cracks for the hydride forming alloy and intergranular for alloys which did not form hydrides was found [95]. For the V±Cr±Ti alloys the hydrogen embrittlement was noticed in both hydride and non-hydride forming alloys. The ductile-to-brittle transition was present in these alloys as a consequence of segregation of sulphur and titanium to the grain boundaries (internal embrittlement). However, the addition of less than 1000 ppm of hydrogen raised the transition temperature from about 210 K to above room temperature (Fig. 13) [96].
8. Vanadium and its alloys The hydrogen entry into vanadium caused a small strengthening effect (increase in yield stress) and, at hydrogen content above a value critical for formation of hydride phase, transition from ductile to brittle behaviour and deep decrease in RA value in a low temperature region (Fig. 12) [89±91]. Hydrogen embrittlement in vanadium and its alloys was resulted from repetition of formation and cleavage fracture of a hydride particle at the crack tip [92]. However, in the testing of effects of oxygen and nitrogen on mechanical behaviour of hydrogenated V, Nb and Ta such simple model appeared inadequate for explaining the mechanical behaviour of the Va group metals [89±91]. For the vanadium-niobium alloys, the hydrogen degradation was similar to that of pure vanadium: an addition of
Fig. 13. Temperature dependence of the ductility (reduction in area) for V±15Cr±5Ti with 0.10%H (black circles) and without hydrogen (light circles) [97].
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Addition of hydrogen to V±Cr alloys caused appreciable strengthening but it also resulted in ductile brittle transition behaviour. Failure was caused by the initiation of surface cracks, transgranular in non-hydride forming alloy, and intergranular in hydride forming alloy [97]. 9. Intermetallics Many ordered intermetallics, like Ni3Al, Ni3(Al,Ti), Ni3Si, Ni3Fe, Co3Ti have been shown to exhibit environmental embrittlement when tested in air or hydrogen at room temperature. This behaviour was suggested to be caused by atomic hydrogen released from the reaction of moisture in the air with a reactive element. Ni3Al [98±102], NiAl [103], Ni3(Al,Mn) [104,105] were shown to become brittle in laboratory air, after pre-charging and testing in air, and during simultaneous charging and straining. An addition of boron (at least 200 ppm) reduced the susceptibility in air but not in presence of cathodic charging [104,106±110]. Ni3Fe was severely embrittled by hydrogen. The degradation was observed in both ordered conditions. Under cathodic pre-charging conditions diffusion of hydrogen occurred along the grain boundaries and led to the intergranular fracture [111]. The both tensile and plastic properties became lower following hydrogen charging (Fig. 14) [111]. The in¯uence of test environment on the tensile properties and fracture behaviour of L12 ordered (Co,Fe)3V alloys was studied [113±115] (Fig. 15). The elongation at room temperature decreased in the sequence of oxygen, vacuum, air and distilled water [115]. The loss in elongation was accompanied by a change in a fracture mode from transgranular to intergranular [115]. The addition of 200 ppm of boron was shown to completely eliminate the environmental embrittlement [112].
Fig. 14. Engineering stress±strain curves for ordered Ni3Fe alloy tested under several conditions [112].
Fig. 15. Comparison of room temperature environmental embrittlement in two (Co,Fe)3V alloys [114].
10. Mechanisms of hydrogen degradation The above reviewed results show that hydrogen degradation of non-ferrous metals and alloys may often occur, even if the hydrogen solubility in a number of materials (e.g. copper and aluminium) is low and hydrogen moves slowly in the lattice. The process of degradation may involve different forms of hydrogen: lattice hydrogen, hydrogen trapped by dislocations, hydrogen bound in hydrides, hydrogen clusters and molecules, and hydrogen compounds (methane). The atomic process of degradation depends on the hydrogen content, structure and microstructure of the metal. The occurrence of hydride-related degradation depends on the stability of this form: if the hydride is very unstable (Al±Zn± Mg alloys) or very stable (nickel), degradation proceeds by the other mechanism. The degradation through the formation and local plastic deformation of the hydrogen bubbles may occur if such bubbles appear at very low hydrogen content, below the limit at which the embrittlement associated with interstitial hydrogen or by hydride decomposition becomes important. The main mechanism of hydrogen cracking seems in any case to be associated with hydrogen-dislocation interaction. The author has shown that such interaction occurs even in high purity copper [116,117] and nickel [117±120]. If neither hydride nor hydrogen bubbles appear, the hydrogen-dislocation interaction and hydrogen-enhanced localised plasticity seems the most likely mechanism of the crack advancement. However, even for the other mechanisms the hydrogendislocation interaction is an important factor. The fast hydrogen transport by moving dislocations seems necessary to explain the fast crack advancement in many metal-hydrogen systems, e.g. in Ni. As shown [120], the hydrogendislocation binding enthalpy is relatively low, 0.11 eV, but even at such value at room temperature and high hydrogen content the amount of hydrogen trapped by dislocations is substantial. When the hydrogen entry occurs at the crack tip and crack walls the high hydrogen content may be reached in this area even if the content of total hydrogen is low.
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11. Conclusions Hydrogen presence in non-ferrous metals and alloys can result in severe degradation of their plasticity and occurrence of hydrogen-induced delayed cracking, depending on environment, stress intensity, chemical composition and microstructure. The cracks propagate by different mechanisms: hydrogen-enhanced localised plasticity, precipitation and decomposition of brittle hydrides, appearance and plastic degradation of hydrogen bubbles. The occurrence of hydrogen-dislocation interaction which causes the fast hydrogen transport by moving dislocation and/or limited plastic deformation or localised softening seems important for any type of degradation. Even if the hydrogen-dislocation binding enthalpy and hydrogen solubility are relatively small in a number of metal-hydrogen systems the hydrogen-dislocation interaction may appear and hydrogen content in a crack area reach high value. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27]
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Hydrogen-assisted degradation of some non-ferrous metals and alloys A. Zielinski Department of Materials Science and Engineering, Technical University of Gdansk, G. Narutowicza 11/12, 80-952 Gdansk, Poland
Abstract The hydrogen degradation of some non-ferrous metals and alloys, and its dependence on microstructure and environment is reviewed. The hydrogen-enhanced localised plasticity, formation and decomposition of brittle hydrides, and appearance of hydrogen and methane bubbles are discussed as possible mechanisms of hydrogen embrittlement. # 2001 Elsevier Science B.V. All rights reserved. Keywords: Hydrogen; Embrittlement; Copper; Aluminium; Nickel; Titanium; Zirconium; Niobium; Vanadium; Intermetallics
1. Introduction Non-ferrous metals and alloys are industrially relevant materials for many applications including hydrogen evolution and absorption. The hydrogen entry into metal or alloy may result in degradation of its mechanical properties to the extent dependent on the environment, stress state, alloy microstructure, etc. The present paper reviews the recent results obtained for this groups of metallic materials, especially on degradation of their microstructure and mechanical properties, conditions which are critical for hydrogen embrittlement to occur, and possible atomic mechanisms of hydrogen delayed cracking. 2. Copper and its alloys Hydrogen solubility in bulk copper is extremely low, about 10ÿ2 ppm, and no hydride formation has been reported for hydrogen charged copper. The properties of bulk copper are in¯uenced by the hydrogen presence: the tensile yield stress increased between 77 and 200 K after high temperature gas hydrogen charging, and hardness increased 3±5% following cathodic charging [1]. At industrial applications the high supersaturation of copper with hydrogen is not expected except electrolytic plating of copper layers in which the hydrogen content may reach even 200 to 1000 ppm [2]. Such high hydrogen content was associated with loss in plasticity of copper layer [2,3]. The decrease in plasticity became reversible only in a small part [2]. The precipitation of hydrogen resulting in bubbles and voids was observed following the gas charging in bulk copper [4] and in copper layer [2,3,5]. Among three types
of voids the smallest originated from hydrogen agglomeration, and the larger, present along the grain boundaries, were pressurised gas bubbles with dislocations punched out of the void [2]. The appearance of voids was postulated as a cause of the loss in plasticity. Thermal charging of Cu±Al alloys increased the microhardness of annealed material. The effect was accompanied by an increase in dislocation density [6]. Effect of hydrogen on mechanical behaviour of a deformation processed Cu±20%Nb composite was investigated in [7]. No substantial degradation, except a small decrease in uniform elongation, was observed in the composite (on the contrary to high susceptibility of niobium). The alumina dispersed copper alloys were investigated after their cathodic and gas charging. The loss in ductility reached about 22% but the increase in Al content had a bene®cial effect on the plasticity. The increase in annealing temperature of cold worked specimens decreased the degree of embrittlement due to decrease in number of reversible hydrogen trapping sites [8]. 3. Aluminium and its alloys Pure aluminium was shown, on the contrary to some previous reports, to be susceptible to the hydrogen presence. The specimens cathodically charged at 300 Amÿ2 were tested at strain rate 10ÿ4 sÿ1. All samples failed in a completely ductile manner. The cathodic charging did cause differences in the stress±strain behaviour: increase in the yield stress, slight increase in the ultimate tensile stress, and decrease in the strain to failure (Fig. 1). The increase in microhardness from 46 to 67 MPa was noticed [9,10].
0924-0136/01/$ ± see front matter # 2001 Elsevier Science B.V. All rights reserved. PII: S 0 9 2 4 - 0 1 3 6 ( 0 0 ) 0 0 7 9 7 - 4
A. Zielinski / Journal of Materials Processing Technology 109 (2001) 206±214
Fig. 1. Effect of hydrogen on failure strain of 99.99% aluminium [11].
In [11] the critical slip shear stress was reduced after hydrogen charging and the atomic binding energy of aluminium decreased about 11%. Water vapour accelerated the fatigue crack propagation in pure aluminium single crystals under mode II or I±II loading [12]. The susceptibility to hydrogen-related degradation of aluminium alloys was observed to depend on microstructure of an aluminium alloy and test environment. The 2090 Al±Li±Cu±Zr alloy after cathodic charging showed the linear decrease in fracture strain following increasing hydrogen content [13]. The hydrogen charging of 8090 Al±Li±Cu±Mg alloys increased the dislocation density and caused the intergranular failure, especially in the overaged condition [14] but did not affect the bending fatigue strength [15]. In [16] the hydrogen charging resulted in ductility loss (Fig. 2), increasing with increasing ageing time. The testing of the 2024 Al±Cu±Mg alloy disclosed severe degradation in overall plastic elongation. The embrittlement was observed after the onset of necking [17]. Reverse torsional fatigue tests of the alloy exposed to hydrogen and humid air showed that hydrogen absorption promoted multiple crack initiation and altered the fracture mode [18].
Fig. 2. Effect of charging time on the ductility loss of an Al±Li±Cu±Mg± Zr alloy [17].
207
The 2124 Al±Cu±Mg alloy (2024 alloy with low Fe and Si) showed no signi®cant loss in ductility for any temper following hydrogen charging but lost about 25% RA when simultaneously strained and charged with hydrogen [19]. The fracture mode was not altered by hydrogen. The other fatigue testing of precharged specimens revealed an increase in fatigue life [20]. The Al±8%Mg alloy was observed to decrease its Young modulus, 0.2% proof stress and tensile stress following hydrogen charging [21]. An exposure of Al±Zn±Mg alloys to water vapour was observed to result in their embrittlement [22±28]. The crack velocity was strongly stress dependent at low stress intensities while at intermediate and high stress intensities there was no effect of stress on the crack growth rate [25]. The cracks were intergranular with a brittle appearance. The crack propagation rate increased with level of humidity, with no effect observed in dry gases, including dry hydrogen [25]. The loss in ductility was observed in tensile tests [23]. The solute hydrogen enhanced dislocation mobility and reduced the ¯ow stress. The cathodic polarisation disclosed similar embrittlement dependent on the alloy and its microstructure. The Al±Zn± Mg alloy was embrittled for all microstructures, and the Al± Zn±Mg±Cu alloy was resistant in the peak strength and overaged condition [29±31]. The effect of cathodic and gaseous hydrogen on fatigue properties was investigated in [22,28,32,33]. Experiments made on Al±Zn±Mg±Cu alloy showed the decrease in fatigue resistance [22,32] (Fig. 3) and in increase in the degree of transgranular crack initiation and propagation [32]. The hydrogen embrittlement was explained by different models,dependingonthealloy.IntheAl±Lialloythehydrogen degradation was caused by precipitation of internal LiH and LiAlH4 hydrides [34,35] which brittle fracture resulted in crack propagation [15]. In pure aluminium the enhanced dislocation mobility was observed [9] in the presence of hydrogen. Similar hydrogen-enhanced localised plasticity seems responsible for the hydrogen embrittlement in the
Fig. 3. Resolved shear stress versus cycles to failure for single crystals of Al±5.5Zn±2.5Mg±1.5Cu in air, and 0.5 M NaCl solution at corrosion and cathodic potential [33].
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other aluminium alloys, like Al±Zn±Mg and Al±Zn±Mg±Cu [36], although formation of magnesium hydride [31,37,38] was also postulated to be a cause of the embrittlement. The appearance of the bubbles decreased the plasticity loss [23]. 4. Nickel and its alloys Numerous studies [39±45] showed that solute hydrogen can cause intergranular fracture of nickel and thereby reduce the ductility. The most characteristic feature is a change in fracture mode from ductile predominantly transgranular microvoid coalescence-type to intergranular fracture with increasing cathodic potential [44] and amount of hydrogen [39,40,42±44]. The intergranular fracture initially occurred as isolated microcracks on the surface exposed to hydrogen. The mode transition was explained by different hydrogen trapping: by grain boundaries, continuous incoherent precipitates or discrete incoherent precipitates [42]. A critical stage was the link-up of microcracks into a macrocrack which propagated, resulting in failure [45]. The critical hydrogen content at the grain boundary was between 0.04 and 0.10 (Fig. 4) [39]. The susceptibility to cracking was remarkably dependent on the presence of impurities within the grain boundary area: carbon decreased the extent of hydrogen embrittlement and sulphur increased the amount of degradation (Fig. 5) [42±44]. The fracture mode was also related to the presence of sulphur: at low sulphur content (0.03 monolayers at the grain boundary) nickel failed in microvoid coalescence-type manner, and at high sulphur content (0.20 monolayers) in brittle intergranular manner [44]. The effect of impurities was explained by their in¯uence on hydrogen concentration which was required for intergranular embrittlement; segregated carbon increased the critical hydrogen concentration and grain boundary diffusivity, and sulphur decreased this value [41]. The hydrogen cracking was postulated to advance rather by the hydrogen-enhanced plasticity than by decohesion or hydride formation [46,47].
Fig. 4. The percent of intergranular fracture versus the grain boundary hydrogen concentration [40].
Fig. 5. Percent intergranular fracture versus grain boundary composition of nickel for several cathodic test potentials [45].
Hydrogen embrittlement of the austenitic Ni±Cr±Fe alloys was often observed following cathodic charging [48±55]. Tensile testing of the charged alloy with Ni content varying between 29 and 59% showed loss in elongation and ultimate tensile stress, with surface cracks appearing on the surface during subsequent ageing of the alloy [48]. The simultaneous charging with hydrogen and deformation in tension of the Inconel 600 76Ni±16Cr±8Fe alloy decreased markedly the fracture strain and changed the fracture mode to completely intergranular [49,50]. The exposure of the Alloy 600 in steam caused hydrogeninduced intergranular stress corrosion cracking [51±53]. The appearance of 30 and 50 nm bubbles along the grain boundaries ®lled in with high pressure methane was observed. The cracks were initiated by the nucleation of high density bubbles on the grain boundary under the combined action of the applied stress and high pressure methane formed from carbon in solution reacting with hydrogen injected by corrosion [52]. The necessary condition for IGSCC to occur was a presence of water/steam; no cracks were formed in dry hydrogen of 27.3 MPa pressure at 673 K [52]. The susceptibility was related to the stability of NiO in this environment [54]. For Fe±Ni±Co superalloys hydrogen was observed to signi®cantly reduce tensile true fracture strain (Fig. 6), slow crack growth thresholds and fracture toughness values [55± 57]. The decrease in fracture toughness was accompanied by a change in fracture mode from microvoid coalescence in the uncharged specimens to principally slip band fracture with some twin bands and ductile intergranular fracture in the hydrogen charged samples. The fracture initiated at matrix carbides [55]. The testing of Ni±Cu Monel alloys showed some loss in tensile strength and RA value [58,59]. The embrittlement decreased with increasing strain rate and grain size. Fracture in air was completely ductile and that in hydrogen brittle and intergranular. The gaseous hydrogen embrittlement was observed for amorphous metallic glass Fe40Ni38Mo4B18 [60]. The fracture
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209
Fig. 7. Ultimate tensile strength as a function of hydrogen concentration in Ti±10V±2Fe±3Al alloy [73].
Fig. 6. Effect of hydrogen on tensile true fracture strain of the IN903 alloy [56].
toughness of notched specimens was signi®cantly reduced at slower strain rates. The degree of embrittlement was similar to that of ferritic or unstable austenitic stainless steels. 5. Titanium alloys Titanium alloys are very complex materials in which two matrix phases, hcp a-phase more susceptible to hydrogen embrittlement with lower hydrogen solubility and diffusivity, and bcc b-phase with opposite properties are present. As much as four different hydrides may appear in titanium alloys: d fcc TiH1.5, e fct TiH2, g fct hydride and strain induced bcc hydride. Therefore, the hydrogen-assisted cracking is strongly related to the alloy microstructure [61]. Commercially pure titanium is very resistant to the embrittlement caused by hydrogen when tested in the form of ®ne-grained specimens at low-to-moderate strain rates in unaxial tension but it becomes susceptible in the presence of a notch, at low temperatures or high strain rates, or large grain sizes. The last effect was reported to be a consequence of an enhancement of both void nucleation and void link-up at large grain sizes/biaxial stresses [62,63]. Testing on oxygen-strengthened titanium demonstrated no effect of hydrogen on the fracture toughness but pronounced effect on impact resistance [64,65]. In testing under sustained load the room temperature rupture times were observed to decrease [66]. The testing of Ti±6Al±4V alloy containing 80 to 720 ppm of hydrogen showed that hydrogen promoted creep of this alloy at room temperature, markedly increasing both creep strain and rate in the primary stage [67]. There was an evidence of softening of the alloy due to hydrogen. The accelerated creep was also observed for the Ti±5Al±2.5Sn alloy [68]. The mechanism of cracking was often suggested to result from formation and decomposition of brittle hydride phases.
In the hcp. a-Ti±4%Al alloy the investigations in highvoltage electron microscope showed that in gaseous hydrogen environment at room temperature two fracture mechanisms could operate, depending on stress intensity: at low stress intensity the cracks propagated by repeated formation and cleavage fracture of hydrides, and at high stress intensity the fracture mode transition occurred at crack propagation rate which exceeded that at which hydride could form in front of the crack, and cracking was possible by hydrogen-enhanced localised plasticity process [61]. The hydride cracking was postulated for the other a and a b-Ti alloys [69]. The hydrogen embrittlement was observed to occur also in the b alloys, like the Ti±Mo± Nb±Al alloy [70] and Ti±V±Cr±Al±Sn alloy [71], and Ti±V± Fe±Al (Figs. 7 and 8) [72], well below hydrogen concentration required to hydride the b-phase. In these alloys the effect of hydrogen is through two primary mechanisms: hydrogen embrittlement due to the effects other than hydride formation, and stabilisation of the b-phase [72,73]. In the alloys, in which the hydride phase was formed, the variation of tensile stress, yield stress and ductility was shown to be proportional to an amount of hydride (Fig. 9) [74]. The effect of microstructure was substantial as shown for the Ti±8Al±1Mo±1V alloy undergoing different heat treatment: in the near-a alloys the cleavage-like fracture
Fig. 8. Reduction in area as a function of hydrogen concentration in Ti±10V±2Fe±3Al alloy [73].
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Fig. 9. Variation of ultimate tensile strength and yield stress with the hydrogen concentration, i.e. amount of hydride [75].
occurred, and in the a-b alloys an alternating extensive a cleavage and ductile rupture of the b ligaments became active [69,75]. For the Ti±6Al±6V±2Sn alloy fracture was characterised by extensive cleavage of the a grains separated by ductile rupture of the b ligaments from threshold to near failure under plane strain conditions, and the material fractured in the plane stress regions failed in a ductile mode [76]. The mechanism of cracking was related also on test environment. In gaseous hydrogen the hydrogen diffusion occurred via b-phase to a-phase, then hydride phase precipitated and crack advanced along a-b interface. During hydrogen charging or exposure in water solution the hydrogen diffusion, hydride formation and fracture took place in the a-phase [77,78]. 6. Zirconium alloys The ductility of the Zr±2.5Nb Zircaloy±2 alloy was shown to decrease with increasing hydrogen content (Fig. 10) [79]. The cracks initiated at hydride blisters which grew on the surface as a result of high pressure. The thermo-migration of
hydrogen to colder points and possible formation of solid zirconium hydrides were the critical steps in initiation of fracture. The critical applied stress was necessary to initiate delayed hydride cracking. The threshold for cracking initiation from a blister ranged between 10.7 and 15.4 MPa1/2 [80]. The amount of plastic strain needed to induce fracture in hydrides decreased as the average hydride length and the axiality of the stress increased. Regardless of the stress state, when hydride platelets had lengths larger than about 50 to 100 mm, little plastic deformation was needed to initiate hydride fracture [81]. The hydrogen embrittlement was a consequence of void nucleation due to the strain induced hydrogen fracture and of the subsequent void growth and void link-up [79]. The mechanism of hydride fracture appeared to involve the slip-induced nucleation of cracks which grew to a critical size under suf®ciently high normal stress. The critical size of a hydride was estimated at 2 to 4 mm in length [82]. The hydrogen degradation induced by brittle fracture of hydrides was also observed in Zr±1.5Sn Zircaloy±4 alloy [83]. The transition from ductile to brittle behaviour was suggested to be prevented by using the ®ne microstructures with elongated grain in the loading direction. The delayed hydride cracking was proved to be very sensitive to prior heat treatment [84]. The maximum crack growth rate was produced in a microstructure consisting of a continuous grain boundary phase b, and the crack velocity decreased progressively when this phase was decomposed and broken up. Rapid quenching which produced ®nely dispersed hydrides resulted in higher crack growth rates. The model of the delayed hydride cracking was proposed which can be applicable for the other metal-hydrogen system in which hydride may appear. In such systems hydrides must be present in a bulk, and hydrogen must be able to ¯ow to the crack tip from these bulk hydrides. The bulk hydrides must be in a dissolving mode, while the crack tip must be in a growing or nucleating mode. The delayed hydrogen cracking was assumed to be caused by preferential and repeated accumulation, and fracture of hydrides at stress raisers such as cracks, causing the cracks to slowly increase over a period of time. The crack tip hydride was said to grow to the critical length as a necessary condition for such type of cracking [85]. 7. Niobium alloys
Fig. 10. The in¯uence of hydrogen on the reduction of area at fracture (circles) and the true strain at fracture (triangles). The tensile tests made either parallel to the rolling direction (black symbols) or transverse to the rolling direction (light symbols) [80].
The Va group transition metals, like niobium, vanadium and tantalum were usually ductile down to 77 K but became brittle at low temperature when they contained even a small amount of hydrogen. For low hydrogen concentration one single minimum and for high hydrogen content the two peaks were observed in relation of tensile elongation on temperature (Fig. 11) [86,87]. No appreciable effect of hydrogen concentration up to 130 ppm on the yield stress was noticed [86].
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211
Fig. 11. The uniform strain as a function of temperature in single crystals Nb±H alloys at 80 ppm wt. H (D), 130 ppm (- - -) and 230 ppm (&).
In fatigue testing, the decrease in stress intensity range, DK, was observed [88]. The hydrogen amount needed to initiate the embrittlement was below critical value necessary to form a hydride. Further increase in hydrogen content had little effect on the fatigue mechanical properties but appeared to change the brittle fracture path. Various models were proposed for the hydrogen embrittlement. The cracking propagation was recently suggested to take place as a repeated process in which hydride would precipitate at an advancing crack tip with the aid of a local stress and fracture in a brittle manner, followed by the crack advancement [89]; the other model proposed the dislocation-hydrogen interaction as an important factor in hydrogen degradation [86,88]. The nature of two step ductile-brittle transition for bulk specimens was explained in several ways. They included: thermal hysteresis of precipitation and dissolution of hydride, decohesion or hydrogen-enhanced localised plasticity at higher temperature and hydride cracking at lower temperature, or ordering of hydrogen within the hydride phase that would lead to partial recovery of ductility between two transition temperatures [86,88].
Fig. 12. Effect of carbon and combined carbon and hydrogen on the temperature dependence of ductility in vanadium [92]. Circles are for pure carbon (black 0.012%C, light 0.20%C), and squares for 0.21%C and 0.27%H (black) or 0.50%H (light).
about 0.6% hydrogen had a moderate strengthening effect on the yield stress but the low temperature decrease in plasticity in suf®ciently hydrogen charged specimens was observed [93]. Similar behaviour was reported for the V±Ti alloys [94,95]. Hydrogen embrittlement was observed as decrease in plasticity (RA) in both the hydride forming and some nonhydride forming V±Ti alloys. The transgranular initiation of cracks for the hydride forming alloy and intergranular for alloys which did not form hydrides was found [95]. For the V±Cr±Ti alloys the hydrogen embrittlement was noticed in both hydride and non-hydride forming alloys. The ductile-to-brittle transition was present in these alloys as a consequence of segregation of sulphur and titanium to the grain boundaries (internal embrittlement). However, the addition of less than 1000 ppm of hydrogen raised the transition temperature from about 210 K to above room temperature (Fig. 13) [96].
8. Vanadium and its alloys The hydrogen entry into vanadium caused a small strengthening effect (increase in yield stress) and, at hydrogen content above a value critical for formation of hydride phase, transition from ductile to brittle behaviour and deep decrease in RA value in a low temperature region (Fig. 12) [89±91]. Hydrogen embrittlement in vanadium and its alloys was resulted from repetition of formation and cleavage fracture of a hydride particle at the crack tip [92]. However, in the testing of effects of oxygen and nitrogen on mechanical behaviour of hydrogenated V, Nb and Ta such simple model appeared inadequate for explaining the mechanical behaviour of the Va group metals [89±91]. For the vanadium-niobium alloys, the hydrogen degradation was similar to that of pure vanadium: an addition of
Fig. 13. Temperature dependence of the ductility (reduction in area) for V±15Cr±5Ti with 0.10%H (black circles) and without hydrogen (light circles) [97].
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Addition of hydrogen to V±Cr alloys caused appreciable strengthening but it also resulted in ductile brittle transition behaviour. Failure was caused by the initiation of surface cracks, transgranular in non-hydride forming alloy, and intergranular in hydride forming alloy [97]. 9. Intermetallics Many ordered intermetallics, like Ni3Al, Ni3(Al,Ti), Ni3Si, Ni3Fe, Co3Ti have been shown to exhibit environmental embrittlement when tested in air or hydrogen at room temperature. This behaviour was suggested to be caused by atomic hydrogen released from the reaction of moisture in the air with a reactive element. Ni3Al [98±102], NiAl [103], Ni3(Al,Mn) [104,105] were shown to become brittle in laboratory air, after pre-charging and testing in air, and during simultaneous charging and straining. An addition of boron (at least 200 ppm) reduced the susceptibility in air but not in presence of cathodic charging [104,106±110]. Ni3Fe was severely embrittled by hydrogen. The degradation was observed in both ordered conditions. Under cathodic pre-charging conditions diffusion of hydrogen occurred along the grain boundaries and led to the intergranular fracture [111]. The both tensile and plastic properties became lower following hydrogen charging (Fig. 14) [111]. The in¯uence of test environment on the tensile properties and fracture behaviour of L12 ordered (Co,Fe)3V alloys was studied [113±115] (Fig. 15). The elongation at room temperature decreased in the sequence of oxygen, vacuum, air and distilled water [115]. The loss in elongation was accompanied by a change in a fracture mode from transgranular to intergranular [115]. The addition of 200 ppm of boron was shown to completely eliminate the environmental embrittlement [112].
Fig. 14. Engineering stress±strain curves for ordered Ni3Fe alloy tested under several conditions [112].
Fig. 15. Comparison of room temperature environmental embrittlement in two (Co,Fe)3V alloys [114].
10. Mechanisms of hydrogen degradation The above reviewed results show that hydrogen degradation of non-ferrous metals and alloys may often occur, even if the hydrogen solubility in a number of materials (e.g. copper and aluminium) is low and hydrogen moves slowly in the lattice. The process of degradation may involve different forms of hydrogen: lattice hydrogen, hydrogen trapped by dislocations, hydrogen bound in hydrides, hydrogen clusters and molecules, and hydrogen compounds (methane). The atomic process of degradation depends on the hydrogen content, structure and microstructure of the metal. The occurrence of hydride-related degradation depends on the stability of this form: if the hydride is very unstable (Al±Zn± Mg alloys) or very stable (nickel), degradation proceeds by the other mechanism. The degradation through the formation and local plastic deformation of the hydrogen bubbles may occur if such bubbles appear at very low hydrogen content, below the limit at which the embrittlement associated with interstitial hydrogen or by hydride decomposition becomes important. The main mechanism of hydrogen cracking seems in any case to be associated with hydrogen-dislocation interaction. The author has shown that such interaction occurs even in high purity copper [116,117] and nickel [117±120]. If neither hydride nor hydrogen bubbles appear, the hydrogen-dislocation interaction and hydrogen-enhanced localised plasticity seems the most likely mechanism of the crack advancement. However, even for the other mechanisms the hydrogendislocation interaction is an important factor. The fast hydrogen transport by moving dislocations seems necessary to explain the fast crack advancement in many metal-hydrogen systems, e.g. in Ni. As shown [120], the hydrogendislocation binding enthalpy is relatively low, 0.11 eV, but even at such value at room temperature and high hydrogen content the amount of hydrogen trapped by dislocations is substantial. When the hydrogen entry occurs at the crack tip and crack walls the high hydrogen content may be reached in this area even if the content of total hydrogen is low.
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11. Conclusions Hydrogen presence in non-ferrous metals and alloys can result in severe degradation of their plasticity and occurrence of hydrogen-induced delayed cracking, depending on environment, stress intensity, chemical composition and microstructure. The cracks propagate by different mechanisms: hydrogen-enhanced localised plasticity, precipitation and decomposition of brittle hydrides, appearance and plastic degradation of hydrogen bubbles. The occurrence of hydrogen-dislocation interaction which causes the fast hydrogen transport by moving dislocation and/or limited plastic deformation or localised softening seems important for any type of degradation. Even if the hydrogen-dislocation binding enthalpy and hydrogen solubility are relatively small in a number of metal-hydrogen systems the hydrogen-dislocation interaction may appear and hydrogen content in a crack area reach high value. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27]
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