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Hydrogen embrittlement of the Ni-base Alloy 600 correlated with hydrogen transport by dislocations
Materials Science and Engineering A262 (1999) 173 – 183
Hydrogen embrittlement of the Ni-base Alloy 600 correlated with hydrogen transport by dislocations F. Lecoester a,1, J. Cheˆne a,*, D. Noel b a
Laboratoire de Me´tallurgie Structurale, URA CNRS n° 1107, Bt413, 91405 Orsay Cedex, France b Electricite´ de France, Les Renardie`res, route de Sens, 77250 Moret-sur-Loing, France Received 18 June 1998; received in revised form 2 September 1998
Abstract Hydrogen transport by dislocations is one of the mechanisms reported in the literature to be responsible for the hydrogen embrittlement (H.E.) of some metals even if it needs to be associated with one or several cracking mechanisms. This study brings new evidence of dislocation sweeping of hydrogen and some arguments in favor of its contribution to the H.E. mechanism of Alloy 600 (Ni–Cr–Fe). A study of tritium desorption assisted by the plastic deformation has been conducted by coupling a b counting technique and a tensile test. The experimental results obtained with Alloy 600 strongly support the mechanism of hydrogen transport by dislocations. The H.E. characteristics of Alloy 600 have been measured on smooth tensile specimens in relationship with the microstructure and the testing conditions. The influence of different parameters including the prestrain level, the strain rate and the temperature supports the role of hydrogen transport in the embrittlement process and the existence of local hydrogen enrichment associated with dislocations pile-ups. © 1999 Elsevier Science S.A. All rights reserved. Keywords: Hydrogen embrittlement; Ni-base alloy 600; Dislocation sweeping of hydrogen; Tritium desorption
1. Introduction Dislocation sweeping of hydrogen in dilute solution in metals has received the support of several experimental results and some simulations. Enhanced hydrogen or tritium diffusion and penetration depths into nickel and nickel base alloys strained while exposed to hydrogen support a mechanism of hydrogen transport assisted by dislocations [1 – 7]. Tritium enrichment on deformation products (dislocations, slip bands, o and a% martensite) formed in fcc iron and nickel base alloys after straining of tritiated sample has been evidenced by tritium autoradiogaphy [7,8]. Hydrogen distribution in the stress field of edge dislocations and its consequence on the dislocations mobility has been recently simulated [9,10]. The free energy minimisation of the metal–hydrogen system leads to hydrogen segregation with atmospheres formation in the hydrostatic tensile stress * Corresponding author. Tel.: + 33-1-69157020; fax: +33-169157833; e-mail: [email protected]. 1 Present address: Laboratoire d’Etude des Mate´riaux Hors Equilibre, Bt 415, Universite´ Paris -Sud, 91405 Orsay Cedex, France.
field and to H depletion in compressive ones. At low strain rates and convenient temperatures, hydrogen mobility may be consistent with dislocation velocity so that atmospheres remain tied to moving dislocations and can be transported [5]. The present paper reports some experimental results concerning the interactions of hydrogen with the Nibase Alloy 600 which are interpreted on the basis of a transport of hydrogen by dislocations. These results refer more precisely to the accelerated desorption of tritium from tensile samples during straining at 293 K and to different characteristics of the embrittlement of Alloy 600 by internal hydrogen.
2. Experimental procedure
2.1. Materials Two industrial heats of Alloy 600 have been studied (noted respectively A and B). Their composition are reported in Table 1 and the different microstructures are shown in Fig. 1. The respective average grain sizes
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Table 1 Chemical composition of Alloy 600 (heats A and B)
A B
C
S
P
Si
Mn
Ni
Cr
Fe
Ti
Cu
Co
Al
0.067 0.058
0.003 0.001
0.006 0.007
0.33 0.45
0.28 0.81
bal. bal.
16.2 16.05
8.2 8.8
0.23 0.29
0.03 0.02
0.02 0.04
0.19 0.24
of heats A and B in the as received state are about 30 and 20 mm. The microstructure of heat A is characterized by the presence of numerous intergranular and intragranular precipitates of chromium carbides. Part of the intragranular carbides are localized on the former austenitic grain boundaries (Fig. 1a). Intergranular and intragranular carbide precipitates are also observed in heat B (Fig. 1f). Their particular repartition in the latter microstructure is illustrated by the presence of bands with a large density of chromium carbides.
2.2. Thermomechanical treatments The different heat treatments and the main characteristics of the corresponding microstructures are reported in Table 2. The as received heat A was solutionized (1 h at 1423 K) and quenched in liquid nitrogen in order to dissolve chromium carbide precipitates and to avoid its precipitation during cooling. Very few intergranular and intragranular precipitates are present in the heat treated alloy (Fig. 1b) and most of the carbon atoms are retained in solid solution. The average grain size of microstructure 2 is about 400 mm. With a lower quenching rate (air cooling), intergranular precipitation occurs (Fig. 1c). An additional annealing for 16 h at 973 K leads to a dense precipitation of intergranular and intragranular chromium carbides (Fig. 1d). The average grain size of microstructures 3 and 4 is comparable (400 mm) to microstructure 2. A 30% cold working has also been performed on the recrystallized microstructure (heat A). It favors a grain elongation as indicated in Fig. 1e.
2.3. Techniques 2.3.1. Hydrogen and tritium charging Assuming that the diffusion rate of hydrogen and its isotopes is very low in fcc structures at 298 K (D= 9.2 × 10 − 11 cm2s − 1 in Alloy 600) [11], hydrogen and tritium were introduced at 423 K by cathodic charging in an eutectic mixture of molten salts (NaHSO4, H2O 53.5%-KHSO4 46,5%) [12]. The electrolytic cell was fitted with a reference electrode Ag/Ag + , a platinum anticathode, and the sample was used as a cathode. A cathodic potential of −1000 mV Ag/Ag + was applied to the sample. The introduction of tritium was performed by substituting part of the water contained in the molten salts by tritiated water. Isotopic ratios Ri =
[3H]/[3H]+ [1H] in the range of 10 − 4 to 10 − 5 may be obtained in this way for the hydrogen species introduced in the samples [12].
2.3.2. Tritium desorption out of Alloy 600 during deformation The study of tritium desorption during deformation was performed with an original equipment by coupling a b counting technique and a tensile test [13]. The procedure was the following: a small tritiated tensile specimen was fixed on the jaws of a tensile machine, and then introduced into a flask containing the liquid scintillator (Fig. 2). The flask was introduced between the two photomultipliers of a b counting apparatus in such a way that the system is hermetically sealed against light. Thus, tritium desorption out of the tensile specimen may be quantified with or without simultaneous straining of the specimen. As hydrogen (and tritium) introduced in the alloy by cathodic charging at 423 K is in supersaturated solid solution, it may desorb from the tensile sample when cathodic charging is interrupted and the temperature brought back to 293 K. The tritium desorption may occur until a thermodynamic equilibrium is achieved between the tritium activity in the alloy and in the environment. The b counting technique allowed the temporal evolution of the tritium activity in the liquid scintillation to be recorded with a very high sensitivity. Thus, the quantity of tritium atoms emanating from Alloy 600 has been evaluated using the tritium decay law: dN = N/t, dt
t= 5.8× 108 s
where N and dN/dt represent the number of tritium atoms contained in the scintillator and the number of desintegrations s − 1. The efficiency of the counting procedure may be alterated by the tensile device introduced in the flask. This efficiency loss was taken into account for the determination of the cumulative amount of tritium which desorbs out of the tensile sample.
2.3.3. Hydrogen embrittlement of Alloy 600 Numerous hydrogen embrittlement tests were performed in previous studies of Alloy 600 [14–19]. In most cases mechanical tests were performed on tensile samples previously hydrogenated by cathodic or
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Fig. 1. Typical microstructures of Alloy 600 (heat A: a–e and B: f): (a) as received, (b) recrystallized for 1 h at 1423 K and quenched in liquid nitrogen, (c) recrystallized for 1 h at 1423 K and air cooled, (d) recrystallized for 1 h at 1423 K and annealed 16 h at 973 K, (e) recrystallized for 1 h at 1423 K and cold worked 30%, (f) as received.
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Table 2 Heat treatments and main characteristics of the different microstructures of Alloy 600 Heat
Heat and mechanical treatments
A
As received
A
Recrystallized 1 liquid nitrogen Recrystallized 1 Recrystallized 1 at 973 K Recrystallized 1 (30%) As received
A A A B
Grain size (mm) Chromium carbides precipitation
30 h at 1423 K and quenched in
400
Inter and intragranular. Intra on former grain 1 boundaries No precipitation 2
h at 1423 K and air cooled h at 1423 K and annealed 16 h
400 400
Intergranular precipitation Intensive inter and intragranular precipitation
3 4
h at 1423 K and cold rolled
400
Intergranular precipitation
5
Inter and intragranular bands with intensive precipitation
6
20
gaseous charging. However, the hydrogen mobility in fcc alloys is very weak so that very often the samples were not fully hydrogenated. In some cases, the hydrogen concentration in the bulk of the sample was almost negligible so that the tensile test was performed on a composite material with large amounts of hydrogen in a small volume located near the surface and free of hydrogen in the bulk. The hydrogen diffusion coefficient in industrial Alloy 600 was measured at different temperatures [18–22]. All these data allow the following Arrhenius law for hydrogen diffusion in Alloy 600 in the temperature range from 298 to 770 K to be defined [22]:
D =2.4× 10 − 3exp
Microstructure reference
−42300 (cm2s − 1), RT
R =8.31 J mol − 1 K − 1 In order to get tensile samples with a quite homogeneous distribution of hydrogen up to the bulk, the hydrogen charging procedure was the following:
Fig. 2. Schematic view of the b counting of a tritiated sample during deformation.
1. small prismatic tensile specimens (8 mm long, 2 mm wide and 0.5 mm thick) were used, 2. they were cathodically charged for 5 h at 423 K in the eutectic mixture of molten salts as indicated above. Assuming that the coverage rate of surface sites was almost constant during charging and a value of D= 1.4× 10 − 8 cm2 s − 1 for the hydrogen diffusion coefficient at 423 K [22], it can be shown that the tensile samples were hydrogenated throughout the bulk. The diffusion was computed by assimilating the tensile samples to a plane sheet [23]. The computed diffusion profile obtained in such conditions is shown in Fig. 3. Just after charging the tensile samples are cleaned and strained to rupture in air with different strain rates (from 3.6 ×10 − 6 to 9 × 10 − 4 s − 1). The complete time required for a tensile test ranges between 10 s to 46 h for these experimental conditions.
Fig. 3. Theoretical diffusion profile of hydrogen in a 0.5 mm thick tensile sample after 5 h charging at 423 K (D= 1.4 ×10 − 8 cm2 s − 1, the concentration C is normalized with respect to the concentration on the surface C0).
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Fig. 4. Tritium desorption at 293 K from a tensile specimen of Alloy 600 (heat A, microstructure 2): effect of the deformation (do/dt= 4× 10 − 4s − 1) on the desorption rate. (a) tritium activity as a function of time, (b) applied stress as a function of time.
The hydrogen embrittlement sensitivity Ie of each microstructure was obtained by measuring the relative homogeneous elongation loss Ie =(Eh −Eh0)/Eh0 where
Fig. 5. Typical tensile curves of hydrogenated and non hydrogenated Alloy 600 (293 K, do/dt=4× 10 − 4 s − 1).
Eh and Eh0 are the homogeneous elongations of the hydrogenated and the uncharged tensile samples. The mechanical properties were measured at different temperature (from 293 to 673 K). For the tests performed above room temperature the procedure was the following: axisymetric tensile specimens (diameter: 2 mm) were cathodically charged for 5 h at 423 K; they were cleaned and coated by PVD with a gold layer 50 nm thick. This coating aimed at reducing the hydrogen desorption during the tensile test at high temperature; the temperature rising rate of the furnace was optimized to minimize the exposure time of the sample at high temperature without straining.
2.3.4. Hydrogen concentration measurement The total amount of hydrogen introduced by cathodic charging in Alloy 600 was measured by extraction at high temperature (1753 K) using a HMAT 2205 Strohlein analyser. The method consists in melting 0.1
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Table 3 Hydrogen embrittlement sensitivity of different microstructures of Alloy 600 (heat A)a Microstructures
As received
Recrystallized 1 h, 1423 K+quenched 77 K
Recrystallized 1 h, 1423 K+annealed 16 h, 973 K
H.E. sensitivity Ie Hydrogen (weight ppm)
73% 24
11% 51
77% 44
a
Comparison with the hydrogen concentration in the tensile samples
to 0.5 g samples in pure argon as a carrier gas. The amount of hydrogen is expressed in weight ppm.
3. Results
3.1. Tritium desorption during straining The recrystallized microstructure 2 (Table 2) which exhibits few obstacles to dislocations motion was chosen for this study. The tensile specimens were first tritiated for 1 h at 423 K in molten salts and stored in liquid nitrogen until testing. The tritiated surface of the sample was about 100 mm2. The b counting started right after immersion of the tensile specimen under a moderate static tensile stress (100 MPa) in the liquid scintillator at 293 K. After the specimen had been submitted to this static stress in the elastic range for 100 min, it was strained (do/dt =4× 10 − 4 s − 1). The tensile test was interrupted at 30% homogeneous deformation. The corresponding macroscopic tensile stress reached a maximum value of 440 MPa and then slightly decreased due to relaxation phenomena. The b counting was continued for about 35 min after interruption of the tensile test and before unloading of the specimen. The evolutions of the tritium activity in the flask and of the imposed load as a function of time, are compared in Fig. 4. The results clearly demonstrate that a continuous plastic deformation assists tritium desorption. Indeed, a strong enhancement of tritium desorption is observed in the plastic deformation regime. A similar effect is hardly seen in the elastic range which is very narrow. The mean tritium desorption flow from the sample under a static stress of 100 MPa is about 3 ×108 atoms cm − 2s − 1. At the onset of the plastic flow the tritium desorption flow from the pulled sample is multiplied almost tenfold (F = 2.4 ×109 atoms cm − 2s − 1). The tritium desorption rate decreases after the arrest of the tensile test but an enhanced desorption is observed for about 10 min after the straining interruption. At the end of the test the tritium desorption rate from the 30% deformed sample is lower than before deformation.
3.2. Characteristics of the hydrogen embrittlement of Alloy 600 3.2.1. Role of the microstructure on the hydrogen embrittlement sensiti6ity of Alloy 600 In order to identify the role of the microstructure on the embrittlement, tensile tests were performed on three different microstructures (1, 2 and 4) of the same heat (A) for the same hydrogen charging conditions. Typical tensile curves obtained at 293 K and a strain rate of 4 ×10 − 4 s − 1 are schematically presented in Fig. 5: the fracture of hydrogenated samples always occurred after significant plastic deformation. The effect of strain rate on the elongation loss at this temperature will be discussed below. The apparent decrease of the Ludwig consolidation coefficient [24] originates from both initiation of microcracks and a softening effect of hydrogen on the plasticity of Alloy 600 [9,10,25]. The role of the deformation level on the crack initiation was not systematically investigated but hydrogen is considered to be fully responsible for the softening observed at low deformation levels as no crack initiation was observed on the surface of the specimens below a 5% deformation. The intergranular character of the fracture increases with the H.E. sensitivity [24]. The hydrogen embrittlement sensitivity of each microstructure is reported in Table 3 together with the amount of hydrogen introduced in 1 mm thick samples. The recrystallized microstructure with copious intergranular and intragranular precipitation of chromium carbides exhibits the largest H.E. sensitivity. In contrast the recrystallized microstructure free of chromium carbide precipitates is the least sensitive to H.E.. The absence of any correlation between the H.E. sensitivity of Alloy 600 and the hydrogen mean concentration in the samples suggests that the H.E. sensitivity intrinsically depends on the microstructure of the alloy. 3.2.2. Influence of carbides precipitation The very different H.E. sensitivity of microstructures 2 and 4 (Table 3) may be attributed to the intergranular precipitation of chromium carbides [19,24]. Recent observations have shown that the H.E. sensitivity of Alloy 600 increases with the density of intergranular precipitates [22]. Thus, the fracture surface of microstructure 3 (heavy intergranular precipitation) is essentially inter-
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Fig. 6. Fracture surface of hydrogenated samples (5 h, 423 K) of Alloy 600: (a) heat A recrystallized for 1 h at 1423 K and and air cooled, (b) heat A recrystallized and quenched in liquid nitrogen.
granular in the presence of hydrogen (Fig. 6a) whereas it is mainly ductile in the case of microstructure 2 free of intergranular precipitates (Fig. 6b). This strong influence of intergranular carbides will be discussed below.
3.2.3. Influence of a tensile predeformation The hydrogen embrittlement tests are particularly reproducible with microstructure 4 which has been used to study the influence of a monotonic plastic predeformation on the hydrogen embrittlement. Tensile specimens (0.8 mm thick) were first submitted to various degrees of tensile deformation (from 0 to 5%) and thinned to 0.5 mm by mechanical polishing. They were then hydrogenated for 5 h at 423 K in molten salts. Taking into account the different gauge lengths of the predeformed specimens, different elongation rates were imposed in order to keep a constant strain rate do/dt =4 × 10 − 4 s − 1 for each test. The results reported on Table 4 show that the H.E. sensitivity is reduced when predeformation of 5% is applied to the samples. Similar results were obtained with microstructure 3 [22]. The fracture mode also strongly depends on the predeformation (Fig. 7). Whereas the fracture is fully intergranular without predeformation (Fig. 7a) it exhibits some traces of ductility when predeformed to 2% and, in spite of intergranular decohesion, is mainly ductile above a predeformation of 5% (Fig. 7b). Table 4 Influence of a predeformation on the hydrogen embrittlement sensitivity of Alloy 600 (heat A, recrystallized 1 h at 1423 K and annealed 16 h at 973 K) Deformation rate (%)
oh = 0
oh = 2
oh = 3
oh = 4
oh =5
Ie (%)
77
64
70
60
51
The predeformation is not considered to impede the hydrogen absorption in the tensile specimens during cathodic charging at 423 K, reducing in this way the H.E sensitivity. Taking into account the small thickness of the samples (0.5 mm) hydrogen is expected to be homogeneously distributed. This is supported by the absence of a change in the fracture mode between the surface and the bulk of the samples (Fig. 7b). The beneficial effect of a predeformation, presumably associated with an increase of the dislocation density, will be discussed below.
3.2.4. Influence of the strain rate Tensile tests were performed at different strain rates with microstructures 1, 2 and 6 (Fig. 8). The hydrogen embrittlement tends to disappear at high strain rate, in agreement with previous results [15]. Moreover, a dependence of the fracture mode on the strain rate is noticed (Figs. 8 and 9). For heat A it is intergranular at low strain rates and mixed (coexistence of brittle intergranular and ductile) at high strain rates contrary to heat B in the as received conditions which is mixed at low strain rates and ductile at high deformation rates. For the microstructure 1, the dependence of the H.E. sensitivity with the strain rate (Fig. 8) is associated with a transition in the rupture mode between do/dt =4× 10 − 4 s − 1 and do/dt = 8×10 − 4 s − 1. The transition range of microstructure 6 which is less sensitive is shifted towards lower strain rates (do/dt =1×10 − 4 s − 1; do/dt = 4×10 − 4 s − 1). In the case of microstructure 2 which is particularly resistant to H. E., the transition range is much lower. It must be emphasized that the fracture mode does not only depend on the macroscopic distribution of hydrogen in the sample since a ductile fracture is also observed in the edge of the fracture surface where the hydrogen concentration is expected to be high.
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Fig. 7. Role of a predeformation on the fracture surface of hydrogenated Alloy 600 (heat A, recrystallized for 1 h at 1423 K and annealed for 16 h at 973 K): (a) without predeformation (b) with 5% predeformation.
3.2.5. Influence of the testing temperature The role of the temperature of the tensile test was investigated with the as received heat A (microstructure 1) which exhibits a high H.E. sensitivity (Table 3). The results are shown in Fig. 10. As expected, the H.E. sensitivity decreases when the testing temperature increases. However a significant H.E. is still observed at 523 K whereas 473 K was considered in a previous study as the higher temperature for H.E. to occur [14]. Taking into account the diameter of the specimens (2 mm) and the charging time at 423 K it can be shown [22] that the specimens were not fully hydrogenated throughout the bulk. This is in good agreement with the fracture obtained at 473 K shown in Fig. 11a which exhibits a brittle intergranular zone located on the edge of the fracture surface. At 523 K, the elongation loss is small but the necking is
Fig. 8. Influence of the strain rate on the hydrogen embrittlement sensitivity of Alloy 600 at 293 K: (a) heat A, as received, (b) heat A, recrystallized 1 h at 1423 K, (c) heat B, as received.
also very small indicating that some hydrogen embrittlement effect still exists (Fig. 11b). Moreover the previous comments together with a possible hydrogen desorption from the sample during the tests at high temperature are in favor of the existence of a hydrogen embrittlement effect in this alloy above a temperature of 523 K.
4. Discussion
4.1. Tritium assisted desorption during tensile straining The enhanced tritium desorption measured under continuous deformation can be explained by a mechanism of hydrogen transport by dislocations. Hydrogen sweeping by dislocations is presumably responsible for the increase of the desorption flow. The relatively large partial molar volume of hydrogen in the alloy favors its segregation after cathodic charging into the tensile stress fields [26] such as those generated by edge dislocations, interstitial or substitutionnal impurity atoms or mutual elastic interactions between defects. When the hydrogenated tensile sample is submitted to a continuous deformation part of the hydrogen population trapped by dislocations may be dragged by moving dislocations, depending on their mobility which is a function of the strain rate (5). The hydrogen atmospheres can be tied to moving dislocations only if the dislocation speed is compatible with the hydrogen mobility. In the present experiment performed at 293 K the relatively low strain rate (4×10 − 4 s − 1) favors hydrogen dragging. Dislocations emerging at the surface of the tensile specimen can liberate their tritium atmosphere in the liquid scintillator leading to an enhanced tritium activity. This effect is anticipated to be less pronounced when numerous defects reducing both
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Fig. 9. Role of the strain rate on the fracture mode at 293 K of hydrogenated tensile samples of Alloy 600 (heat A): (a) brittle intergranular (do/dt= 4 × 10 − 5 s − 1). (b) mixed (do/dt=8× 10 − 4 s − 1).
hydrogen and dislocations motion are present in the microstructure. This justifies the choice of a recrystallized microstructure for this experiment. An enhanced desorption flow of tritium is still measured for about 10 min after the interruption of the tensile test. This is presumably related to the stress relaxation associated with the imposed deformation and may also be a consequence of tritium transport by dislocations moving at a low velocity during the relaxation process. Alternatively, a slow kinetics of the tritium transfer at the interface metal-scintillator could contribute to this phenomenon which needs further investigation. The possibility of hydrogen transport enhanced by mobile dislocations at room temperature [9,10] together with the fact that H.E. sensitivity of the alloy is still observed above 523 K (Section 3.2), suggest that dislocation assisted transport during the creep of alloy 600 may contribute to the stress-corrosion cracking of this alloy at 630 K [27,28].
4.2. Hydrogen embrittlement The dependence of the H.E. sensitivity of Alloy 600 on the different parameters investigated in this study account for a contribution of the hydrogen sweeping phenomenon to the embrittlement process. (i) A predeformation leads to a reduction of the H.E. sensitivity because the dislocation network generated by the predeformation may impede the hydrogen transport by mobile dislocations. (ii) The intergranular precipitation of chromium carbides increases the H.E. sensitivity of the alloy in agreement with a mechanism of hydrogen transport and enrichment at grain boundaries during straining. The latter is presumably a consequence of preferential formation of dislocation pile-ups at the carbide enriched grain boundaries. This strain-induced hydrogen enrichment near the grain boundaries together with a possible preexisting trapping of hydrogen on the intergranular carbides is expected to favor intergranular fracture. (iii) The hydrogen embrittlement sensitivity of Alloy 600 decreases with increasing temperature of the tensile test. In a similar way, the temperature dependence of the hydrogen trapping on dislocations can be expressed to a first approximation by the following relationship [29]: c= c0/(1− c0) exp(Hb/kT)
Fig. 10. Influence of the testing temperature on the hydrogen embrittlement sensitivity of Alloy 600 (heat A, as received, do/dt= 4 × 10 − 4 s − 1).
where c, occupancy fraction of the dislocation; c0, occupancy fraction of non pertubed lattice sites; Hb, binding enthalpy of hydrogen to the dislocation. One can then anticipate a decrease of the amount of hydrogen trapped and transported by mobile dislocations when the testing temperature increases and as a consequence a reduced embrittlement effect at high temperature.
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Fig. 11. Role of the testing temperature on the fracture surface of hydrogenated Alloy 600 (heat A, as received, do/dt=4 × 10 − 4 s − 1): (a) T= 473 K (b) T =523 K.
(iv) The H.E. sensitivity of Alloy 600 increases at low strain rates together with the amount of hydrogen transported by dislocations [5]. These various effects of the different parameters on the H.E sensitivity of Alloy 600 can all be explained by a predominant effect of dislocation sweeping. This accelerated transport would lead to local overconcentration of hydrogen responsible for the crack initiation, whatever the cracking mechanism. The hydrogen transport by dislocations appears consequently as a necessary step for the embrittlement to occur. The hydrogen transport allows also to explain the change in fracture mode with strain rate. At high strain rate, the dislocation velocity is too high for significant hydrogen sweeping and the local hydrogen concentration is not sufficient to initiate a crack. At low strain rate, microcracks associated with the local hydrogen overconcentration may nucleate and will favor the initiation of a macroscopic crack. Once this crack has developed extensively, its propagation will depend on the local deformation rate at the crack tip. For moderate deformation rate, a local hydrogen enrichment on obstacles interacting with dislocations in the plastic zone away from the crack tip may favor a hydrogen-induced decohesion during further straining. This is expected to occur especially in the vicinity of grain boundaries acting as preferential sites for crack propagation and favoring intergranular brittle fracture. At high local deformation rates the absence of significant hydrogen transport to defects acting as a barrier to dislocation motion (grain boundaries) will favor a transgranular cracking with a large ductility. The frac-
ture mode is ductile but the presence of hydrogen may affect the dimple size.
5. Conclusion An accelerated desorption of tritium from Alloy 600 has been measured at 293 K with tritiated tensile specimen, by b counting during straining. The tritium desorption flow from the sample is ten times larger during straining for the present experimental conditions. These results bring strong arguments in favor of a mechanism of hydrogen transport by dislocations. The role of different metallurgical parameters on the hydrogen embrittlement (H.E.) sensitivity of Alloy 600 has been investigated with precharged tensile samples: the H.E. sensitivity at 293 K strongly depends on the microstructure of the alloy. The intergranular precipitation of chromium carbides (Cr7C3) increases the H.E. sensitivity whereas it decreases with a predeformation of the alloy; the H.E. sensitivity of any microstructures depends on the testing conditions: strain rate and temperature. The fracture mode may change from fully brittle intergranular at low strain rate to fully ductile at high strain rate with a mixed mode in an intermediate range of strain rate. The change in fracture mode as a function of the strain rate depends on the microstructure, the more sensitive the microstructure the higher the strain rate corresponding to the brittle/ductile transition. The H.E. sensitivity decreases at high temperature.
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However a significant embrittlement has been evidenced up to 523 K in the as received Alloy 600. The role of these different parameters can be rationalized in term of a determinent influence of hydrogen transport by mobile dislocations leading to a local hydrogen enrichment in the dislocation pile ups which favor crack initiation. This accelerated transport mechanism is considered to be the step which may control the hydrogen embrittlement process of alloy 600, whatever the cracking mechanism.
[12]
[13] [14] [15] [16]
[17] [18]
Acknowledgements Financial support provided by EDF, Direction des Etudes et Recherches is gratefully acknowledged.
[19]
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alliages de nickel. Rapport EDF-DER n° HT-45/PV G 330-A, 1988. P. Lacombe, M. Aucouturier, J. Cheˆne, in: R. Gibala, R.F. Hehemann (Eds.), Hydrogen Embrittlement and Stress Corrosion Cracking, A.S.M., Metals Park, 1984, p. 79. J. Cheˆne, L. Grisy, (Submitted for publication). R. Lagneborg, J. Iron Steel Inst. 207 (1969) 363. M. Cornet, C. Bertrand, M. Da Cunha Belo, Metall. Trans. 13A (1982) 141. Y. Shehu, P. Menut, J. Cheˆne, M. Aucouturier, in: P. Azou, Proceedings of the Third International Conference on Hydrogen and Materials, Paris, 1982, pp. 899. Z. Szklarska-Smialowska, Z. Xia, R.R. Valbuena, Corrosion NACE 50 (9) (1994) 676. K. Mummert, H.J. Engelmann, S. Schwartz, M. Uhlemann, in: A.W. Thompson, N.R. Moody (Eds.), Hydrogen effect in materials, T.M.S., Warrendale, 1996, p. 679. I. Lenartova, Fragilisation par l’hydroge`ne et corrosion sous contrainte d’alliages de nickel et d’un acier inoxydable utilise´s dans les ge´ne´rateurs de vapeur: influence de la microstructure, Thesis Ecole Centrale Paris, 1996. R. M. Latanision, M. Kurkela, Corrosion 39 (1983) 174. A. Turnbull, R.G. Ballinger, J.S. Hwang, M.M. Morra, M. Psaila-Dombrowski, R.M. Gates, Met. Trans. 23 (1992) 3231. F. Lecoester, Interactions hydroge`ne-alliage 600 (NC15Fe): application au cas des tubes de ge´ne´rateur de vapeur des re´acteurs a` eau pressurise´e, soumis au milieu primaire, Thesis Universite´ Paris-Sud, 1997. J. Crank, The mathematics of diffusion, 2nd ed., Oxford University Press, London, 1975, p. 50. F. Lecoester, A.M. Brass, J. Cheˆne, D. Noel, in: T. Magnin, Corrosion-Deformation Interactions, The Institute of Materials, London, 1997, p. 473. C. D. Beachem, Met. Trans. 3A (1972) 437. J. Cheˆne, A.M. Brass, in: D. Desjardins, R. Oltra (Eds.), Corrosion Sous Contrainte, Les Editions de Physique, Les Ulis, 1992, p. 159. T. M. Angeliu, D.J. Paraventi, G.S. Was, Corrosion 51 (11) (1995) 837. J.F. Kergaravat, Influence de la microstructure sur la sensibilite´ a` la corrosion sous contrainte des alliages 600 et 690 en milieu primaire des des re´acteurs a` eau sous pression, Thesis, Institut National Polytechnique de Grenoble, 1996. J.P. Hirth, Met. Trans. 11 (1980) 861.
.
Hydrogen embrittlement of the Ni-base Alloy 600 correlated with hydrogen transport by dislocations F. Lecoester a,1, J. Cheˆne a,*, D. Noel b a
Laboratoire de Me´tallurgie Structurale, URA CNRS n° 1107, Bt413, 91405 Orsay Cedex, France b Electricite´ de France, Les Renardie`res, route de Sens, 77250 Moret-sur-Loing, France Received 18 June 1998; received in revised form 2 September 1998
Abstract Hydrogen transport by dislocations is one of the mechanisms reported in the literature to be responsible for the hydrogen embrittlement (H.E.) of some metals even if it needs to be associated with one or several cracking mechanisms. This study brings new evidence of dislocation sweeping of hydrogen and some arguments in favor of its contribution to the H.E. mechanism of Alloy 600 (Ni–Cr–Fe). A study of tritium desorption assisted by the plastic deformation has been conducted by coupling a b counting technique and a tensile test. The experimental results obtained with Alloy 600 strongly support the mechanism of hydrogen transport by dislocations. The H.E. characteristics of Alloy 600 have been measured on smooth tensile specimens in relationship with the microstructure and the testing conditions. The influence of different parameters including the prestrain level, the strain rate and the temperature supports the role of hydrogen transport in the embrittlement process and the existence of local hydrogen enrichment associated with dislocations pile-ups. © 1999 Elsevier Science S.A. All rights reserved. Keywords: Hydrogen embrittlement; Ni-base alloy 600; Dislocation sweeping of hydrogen; Tritium desorption
1. Introduction Dislocation sweeping of hydrogen in dilute solution in metals has received the support of several experimental results and some simulations. Enhanced hydrogen or tritium diffusion and penetration depths into nickel and nickel base alloys strained while exposed to hydrogen support a mechanism of hydrogen transport assisted by dislocations [1 – 7]. Tritium enrichment on deformation products (dislocations, slip bands, o and a% martensite) formed in fcc iron and nickel base alloys after straining of tritiated sample has been evidenced by tritium autoradiogaphy [7,8]. Hydrogen distribution in the stress field of edge dislocations and its consequence on the dislocations mobility has been recently simulated [9,10]. The free energy minimisation of the metal–hydrogen system leads to hydrogen segregation with atmospheres formation in the hydrostatic tensile stress * Corresponding author. Tel.: + 33-1-69157020; fax: +33-169157833; e-mail: [email protected]. 1 Present address: Laboratoire d’Etude des Mate´riaux Hors Equilibre, Bt 415, Universite´ Paris -Sud, 91405 Orsay Cedex, France.
field and to H depletion in compressive ones. At low strain rates and convenient temperatures, hydrogen mobility may be consistent with dislocation velocity so that atmospheres remain tied to moving dislocations and can be transported [5]. The present paper reports some experimental results concerning the interactions of hydrogen with the Nibase Alloy 600 which are interpreted on the basis of a transport of hydrogen by dislocations. These results refer more precisely to the accelerated desorption of tritium from tensile samples during straining at 293 K and to different characteristics of the embrittlement of Alloy 600 by internal hydrogen.
2. Experimental procedure
2.1. Materials Two industrial heats of Alloy 600 have been studied (noted respectively A and B). Their composition are reported in Table 1 and the different microstructures are shown in Fig. 1. The respective average grain sizes
0921-5093/99/$ - see front matter © 1999 Elsevier Science S.A. All rights reserved. PII: S 0 9 2 1 - 5 0 9 3 ( 9 8 ) 0 1 0 0 6 - 5
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Table 1 Chemical composition of Alloy 600 (heats A and B)
A B
C
S
P
Si
Mn
Ni
Cr
Fe
Ti
Cu
Co
Al
0.067 0.058
0.003 0.001
0.006 0.007
0.33 0.45
0.28 0.81
bal. bal.
16.2 16.05
8.2 8.8
0.23 0.29
0.03 0.02
0.02 0.04
0.19 0.24
of heats A and B in the as received state are about 30 and 20 mm. The microstructure of heat A is characterized by the presence of numerous intergranular and intragranular precipitates of chromium carbides. Part of the intragranular carbides are localized on the former austenitic grain boundaries (Fig. 1a). Intergranular and intragranular carbide precipitates are also observed in heat B (Fig. 1f). Their particular repartition in the latter microstructure is illustrated by the presence of bands with a large density of chromium carbides.
2.2. Thermomechanical treatments The different heat treatments and the main characteristics of the corresponding microstructures are reported in Table 2. The as received heat A was solutionized (1 h at 1423 K) and quenched in liquid nitrogen in order to dissolve chromium carbide precipitates and to avoid its precipitation during cooling. Very few intergranular and intragranular precipitates are present in the heat treated alloy (Fig. 1b) and most of the carbon atoms are retained in solid solution. The average grain size of microstructure 2 is about 400 mm. With a lower quenching rate (air cooling), intergranular precipitation occurs (Fig. 1c). An additional annealing for 16 h at 973 K leads to a dense precipitation of intergranular and intragranular chromium carbides (Fig. 1d). The average grain size of microstructures 3 and 4 is comparable (400 mm) to microstructure 2. A 30% cold working has also been performed on the recrystallized microstructure (heat A). It favors a grain elongation as indicated in Fig. 1e.
2.3. Techniques 2.3.1. Hydrogen and tritium charging Assuming that the diffusion rate of hydrogen and its isotopes is very low in fcc structures at 298 K (D= 9.2 × 10 − 11 cm2s − 1 in Alloy 600) [11], hydrogen and tritium were introduced at 423 K by cathodic charging in an eutectic mixture of molten salts (NaHSO4, H2O 53.5%-KHSO4 46,5%) [12]. The electrolytic cell was fitted with a reference electrode Ag/Ag + , a platinum anticathode, and the sample was used as a cathode. A cathodic potential of −1000 mV Ag/Ag + was applied to the sample. The introduction of tritium was performed by substituting part of the water contained in the molten salts by tritiated water. Isotopic ratios Ri =
[3H]/[3H]+ [1H] in the range of 10 − 4 to 10 − 5 may be obtained in this way for the hydrogen species introduced in the samples [12].
2.3.2. Tritium desorption out of Alloy 600 during deformation The study of tritium desorption during deformation was performed with an original equipment by coupling a b counting technique and a tensile test [13]. The procedure was the following: a small tritiated tensile specimen was fixed on the jaws of a tensile machine, and then introduced into a flask containing the liquid scintillator (Fig. 2). The flask was introduced between the two photomultipliers of a b counting apparatus in such a way that the system is hermetically sealed against light. Thus, tritium desorption out of the tensile specimen may be quantified with or without simultaneous straining of the specimen. As hydrogen (and tritium) introduced in the alloy by cathodic charging at 423 K is in supersaturated solid solution, it may desorb from the tensile sample when cathodic charging is interrupted and the temperature brought back to 293 K. The tritium desorption may occur until a thermodynamic equilibrium is achieved between the tritium activity in the alloy and in the environment. The b counting technique allowed the temporal evolution of the tritium activity in the liquid scintillation to be recorded with a very high sensitivity. Thus, the quantity of tritium atoms emanating from Alloy 600 has been evaluated using the tritium decay law: dN = N/t, dt
t= 5.8× 108 s
where N and dN/dt represent the number of tritium atoms contained in the scintillator and the number of desintegrations s − 1. The efficiency of the counting procedure may be alterated by the tensile device introduced in the flask. This efficiency loss was taken into account for the determination of the cumulative amount of tritium which desorbs out of the tensile sample.
2.3.3. Hydrogen embrittlement of Alloy 600 Numerous hydrogen embrittlement tests were performed in previous studies of Alloy 600 [14–19]. In most cases mechanical tests were performed on tensile samples previously hydrogenated by cathodic or
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Fig. 1. Typical microstructures of Alloy 600 (heat A: a–e and B: f): (a) as received, (b) recrystallized for 1 h at 1423 K and quenched in liquid nitrogen, (c) recrystallized for 1 h at 1423 K and air cooled, (d) recrystallized for 1 h at 1423 K and annealed 16 h at 973 K, (e) recrystallized for 1 h at 1423 K and cold worked 30%, (f) as received.
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Table 2 Heat treatments and main characteristics of the different microstructures of Alloy 600 Heat
Heat and mechanical treatments
A
As received
A
Recrystallized 1 liquid nitrogen Recrystallized 1 Recrystallized 1 at 973 K Recrystallized 1 (30%) As received
A A A B
Grain size (mm) Chromium carbides precipitation
30 h at 1423 K and quenched in
400
Inter and intragranular. Intra on former grain 1 boundaries No precipitation 2
h at 1423 K and air cooled h at 1423 K and annealed 16 h
400 400
Intergranular precipitation Intensive inter and intragranular precipitation
3 4
h at 1423 K and cold rolled
400
Intergranular precipitation
5
Inter and intragranular bands with intensive precipitation
6
20
gaseous charging. However, the hydrogen mobility in fcc alloys is very weak so that very often the samples were not fully hydrogenated. In some cases, the hydrogen concentration in the bulk of the sample was almost negligible so that the tensile test was performed on a composite material with large amounts of hydrogen in a small volume located near the surface and free of hydrogen in the bulk. The hydrogen diffusion coefficient in industrial Alloy 600 was measured at different temperatures [18–22]. All these data allow the following Arrhenius law for hydrogen diffusion in Alloy 600 in the temperature range from 298 to 770 K to be defined [22]:
D =2.4× 10 − 3exp
Microstructure reference
−42300 (cm2s − 1), RT
R =8.31 J mol − 1 K − 1 In order to get tensile samples with a quite homogeneous distribution of hydrogen up to the bulk, the hydrogen charging procedure was the following:
Fig. 2. Schematic view of the b counting of a tritiated sample during deformation.
1. small prismatic tensile specimens (8 mm long, 2 mm wide and 0.5 mm thick) were used, 2. they were cathodically charged for 5 h at 423 K in the eutectic mixture of molten salts as indicated above. Assuming that the coverage rate of surface sites was almost constant during charging and a value of D= 1.4× 10 − 8 cm2 s − 1 for the hydrogen diffusion coefficient at 423 K [22], it can be shown that the tensile samples were hydrogenated throughout the bulk. The diffusion was computed by assimilating the tensile samples to a plane sheet [23]. The computed diffusion profile obtained in such conditions is shown in Fig. 3. Just after charging the tensile samples are cleaned and strained to rupture in air with different strain rates (from 3.6 ×10 − 6 to 9 × 10 − 4 s − 1). The complete time required for a tensile test ranges between 10 s to 46 h for these experimental conditions.
Fig. 3. Theoretical diffusion profile of hydrogen in a 0.5 mm thick tensile sample after 5 h charging at 423 K (D= 1.4 ×10 − 8 cm2 s − 1, the concentration C is normalized with respect to the concentration on the surface C0).
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Fig. 4. Tritium desorption at 293 K from a tensile specimen of Alloy 600 (heat A, microstructure 2): effect of the deformation (do/dt= 4× 10 − 4s − 1) on the desorption rate. (a) tritium activity as a function of time, (b) applied stress as a function of time.
The hydrogen embrittlement sensitivity Ie of each microstructure was obtained by measuring the relative homogeneous elongation loss Ie =(Eh −Eh0)/Eh0 where
Fig. 5. Typical tensile curves of hydrogenated and non hydrogenated Alloy 600 (293 K, do/dt=4× 10 − 4 s − 1).
Eh and Eh0 are the homogeneous elongations of the hydrogenated and the uncharged tensile samples. The mechanical properties were measured at different temperature (from 293 to 673 K). For the tests performed above room temperature the procedure was the following: axisymetric tensile specimens (diameter: 2 mm) were cathodically charged for 5 h at 423 K; they were cleaned and coated by PVD with a gold layer 50 nm thick. This coating aimed at reducing the hydrogen desorption during the tensile test at high temperature; the temperature rising rate of the furnace was optimized to minimize the exposure time of the sample at high temperature without straining.
2.3.4. Hydrogen concentration measurement The total amount of hydrogen introduced by cathodic charging in Alloy 600 was measured by extraction at high temperature (1753 K) using a HMAT 2205 Strohlein analyser. The method consists in melting 0.1
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Table 3 Hydrogen embrittlement sensitivity of different microstructures of Alloy 600 (heat A)a Microstructures
As received
Recrystallized 1 h, 1423 K+quenched 77 K
Recrystallized 1 h, 1423 K+annealed 16 h, 973 K
H.E. sensitivity Ie Hydrogen (weight ppm)
73% 24
11% 51
77% 44
a
Comparison with the hydrogen concentration in the tensile samples
to 0.5 g samples in pure argon as a carrier gas. The amount of hydrogen is expressed in weight ppm.
3. Results
3.1. Tritium desorption during straining The recrystallized microstructure 2 (Table 2) which exhibits few obstacles to dislocations motion was chosen for this study. The tensile specimens were first tritiated for 1 h at 423 K in molten salts and stored in liquid nitrogen until testing. The tritiated surface of the sample was about 100 mm2. The b counting started right after immersion of the tensile specimen under a moderate static tensile stress (100 MPa) in the liquid scintillator at 293 K. After the specimen had been submitted to this static stress in the elastic range for 100 min, it was strained (do/dt =4× 10 − 4 s − 1). The tensile test was interrupted at 30% homogeneous deformation. The corresponding macroscopic tensile stress reached a maximum value of 440 MPa and then slightly decreased due to relaxation phenomena. The b counting was continued for about 35 min after interruption of the tensile test and before unloading of the specimen. The evolutions of the tritium activity in the flask and of the imposed load as a function of time, are compared in Fig. 4. The results clearly demonstrate that a continuous plastic deformation assists tritium desorption. Indeed, a strong enhancement of tritium desorption is observed in the plastic deformation regime. A similar effect is hardly seen in the elastic range which is very narrow. The mean tritium desorption flow from the sample under a static stress of 100 MPa is about 3 ×108 atoms cm − 2s − 1. At the onset of the plastic flow the tritium desorption flow from the pulled sample is multiplied almost tenfold (F = 2.4 ×109 atoms cm − 2s − 1). The tritium desorption rate decreases after the arrest of the tensile test but an enhanced desorption is observed for about 10 min after the straining interruption. At the end of the test the tritium desorption rate from the 30% deformed sample is lower than before deformation.
3.2. Characteristics of the hydrogen embrittlement of Alloy 600 3.2.1. Role of the microstructure on the hydrogen embrittlement sensiti6ity of Alloy 600 In order to identify the role of the microstructure on the embrittlement, tensile tests were performed on three different microstructures (1, 2 and 4) of the same heat (A) for the same hydrogen charging conditions. Typical tensile curves obtained at 293 K and a strain rate of 4 ×10 − 4 s − 1 are schematically presented in Fig. 5: the fracture of hydrogenated samples always occurred after significant plastic deformation. The effect of strain rate on the elongation loss at this temperature will be discussed below. The apparent decrease of the Ludwig consolidation coefficient [24] originates from both initiation of microcracks and a softening effect of hydrogen on the plasticity of Alloy 600 [9,10,25]. The role of the deformation level on the crack initiation was not systematically investigated but hydrogen is considered to be fully responsible for the softening observed at low deformation levels as no crack initiation was observed on the surface of the specimens below a 5% deformation. The intergranular character of the fracture increases with the H.E. sensitivity [24]. The hydrogen embrittlement sensitivity of each microstructure is reported in Table 3 together with the amount of hydrogen introduced in 1 mm thick samples. The recrystallized microstructure with copious intergranular and intragranular precipitation of chromium carbides exhibits the largest H.E. sensitivity. In contrast the recrystallized microstructure free of chromium carbide precipitates is the least sensitive to H.E.. The absence of any correlation between the H.E. sensitivity of Alloy 600 and the hydrogen mean concentration in the samples suggests that the H.E. sensitivity intrinsically depends on the microstructure of the alloy. 3.2.2. Influence of carbides precipitation The very different H.E. sensitivity of microstructures 2 and 4 (Table 3) may be attributed to the intergranular precipitation of chromium carbides [19,24]. Recent observations have shown that the H.E. sensitivity of Alloy 600 increases with the density of intergranular precipitates [22]. Thus, the fracture surface of microstructure 3 (heavy intergranular precipitation) is essentially inter-
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Fig. 6. Fracture surface of hydrogenated samples (5 h, 423 K) of Alloy 600: (a) heat A recrystallized for 1 h at 1423 K and and air cooled, (b) heat A recrystallized and quenched in liquid nitrogen.
granular in the presence of hydrogen (Fig. 6a) whereas it is mainly ductile in the case of microstructure 2 free of intergranular precipitates (Fig. 6b). This strong influence of intergranular carbides will be discussed below.
3.2.3. Influence of a tensile predeformation The hydrogen embrittlement tests are particularly reproducible with microstructure 4 which has been used to study the influence of a monotonic plastic predeformation on the hydrogen embrittlement. Tensile specimens (0.8 mm thick) were first submitted to various degrees of tensile deformation (from 0 to 5%) and thinned to 0.5 mm by mechanical polishing. They were then hydrogenated for 5 h at 423 K in molten salts. Taking into account the different gauge lengths of the predeformed specimens, different elongation rates were imposed in order to keep a constant strain rate do/dt =4 × 10 − 4 s − 1 for each test. The results reported on Table 4 show that the H.E. sensitivity is reduced when predeformation of 5% is applied to the samples. Similar results were obtained with microstructure 3 [22]. The fracture mode also strongly depends on the predeformation (Fig. 7). Whereas the fracture is fully intergranular without predeformation (Fig. 7a) it exhibits some traces of ductility when predeformed to 2% and, in spite of intergranular decohesion, is mainly ductile above a predeformation of 5% (Fig. 7b). Table 4 Influence of a predeformation on the hydrogen embrittlement sensitivity of Alloy 600 (heat A, recrystallized 1 h at 1423 K and annealed 16 h at 973 K) Deformation rate (%)
oh = 0
oh = 2
oh = 3
oh = 4
oh =5
Ie (%)
77
64
70
60
51
The predeformation is not considered to impede the hydrogen absorption in the tensile specimens during cathodic charging at 423 K, reducing in this way the H.E sensitivity. Taking into account the small thickness of the samples (0.5 mm) hydrogen is expected to be homogeneously distributed. This is supported by the absence of a change in the fracture mode between the surface and the bulk of the samples (Fig. 7b). The beneficial effect of a predeformation, presumably associated with an increase of the dislocation density, will be discussed below.
3.2.4. Influence of the strain rate Tensile tests were performed at different strain rates with microstructures 1, 2 and 6 (Fig. 8). The hydrogen embrittlement tends to disappear at high strain rate, in agreement with previous results [15]. Moreover, a dependence of the fracture mode on the strain rate is noticed (Figs. 8 and 9). For heat A it is intergranular at low strain rates and mixed (coexistence of brittle intergranular and ductile) at high strain rates contrary to heat B in the as received conditions which is mixed at low strain rates and ductile at high deformation rates. For the microstructure 1, the dependence of the H.E. sensitivity with the strain rate (Fig. 8) is associated with a transition in the rupture mode between do/dt =4× 10 − 4 s − 1 and do/dt = 8×10 − 4 s − 1. The transition range of microstructure 6 which is less sensitive is shifted towards lower strain rates (do/dt =1×10 − 4 s − 1; do/dt = 4×10 − 4 s − 1). In the case of microstructure 2 which is particularly resistant to H. E., the transition range is much lower. It must be emphasized that the fracture mode does not only depend on the macroscopic distribution of hydrogen in the sample since a ductile fracture is also observed in the edge of the fracture surface where the hydrogen concentration is expected to be high.
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Fig. 7. Role of a predeformation on the fracture surface of hydrogenated Alloy 600 (heat A, recrystallized for 1 h at 1423 K and annealed for 16 h at 973 K): (a) without predeformation (b) with 5% predeformation.
3.2.5. Influence of the testing temperature The role of the temperature of the tensile test was investigated with the as received heat A (microstructure 1) which exhibits a high H.E. sensitivity (Table 3). The results are shown in Fig. 10. As expected, the H.E. sensitivity decreases when the testing temperature increases. However a significant H.E. is still observed at 523 K whereas 473 K was considered in a previous study as the higher temperature for H.E. to occur [14]. Taking into account the diameter of the specimens (2 mm) and the charging time at 423 K it can be shown [22] that the specimens were not fully hydrogenated throughout the bulk. This is in good agreement with the fracture obtained at 473 K shown in Fig. 11a which exhibits a brittle intergranular zone located on the edge of the fracture surface. At 523 K, the elongation loss is small but the necking is
Fig. 8. Influence of the strain rate on the hydrogen embrittlement sensitivity of Alloy 600 at 293 K: (a) heat A, as received, (b) heat A, recrystallized 1 h at 1423 K, (c) heat B, as received.
also very small indicating that some hydrogen embrittlement effect still exists (Fig. 11b). Moreover the previous comments together with a possible hydrogen desorption from the sample during the tests at high temperature are in favor of the existence of a hydrogen embrittlement effect in this alloy above a temperature of 523 K.
4. Discussion
4.1. Tritium assisted desorption during tensile straining The enhanced tritium desorption measured under continuous deformation can be explained by a mechanism of hydrogen transport by dislocations. Hydrogen sweeping by dislocations is presumably responsible for the increase of the desorption flow. The relatively large partial molar volume of hydrogen in the alloy favors its segregation after cathodic charging into the tensile stress fields [26] such as those generated by edge dislocations, interstitial or substitutionnal impurity atoms or mutual elastic interactions between defects. When the hydrogenated tensile sample is submitted to a continuous deformation part of the hydrogen population trapped by dislocations may be dragged by moving dislocations, depending on their mobility which is a function of the strain rate (5). The hydrogen atmospheres can be tied to moving dislocations only if the dislocation speed is compatible with the hydrogen mobility. In the present experiment performed at 293 K the relatively low strain rate (4×10 − 4 s − 1) favors hydrogen dragging. Dislocations emerging at the surface of the tensile specimen can liberate their tritium atmosphere in the liquid scintillator leading to an enhanced tritium activity. This effect is anticipated to be less pronounced when numerous defects reducing both
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Fig. 9. Role of the strain rate on the fracture mode at 293 K of hydrogenated tensile samples of Alloy 600 (heat A): (a) brittle intergranular (do/dt= 4 × 10 − 5 s − 1). (b) mixed (do/dt=8× 10 − 4 s − 1).
hydrogen and dislocations motion are present in the microstructure. This justifies the choice of a recrystallized microstructure for this experiment. An enhanced desorption flow of tritium is still measured for about 10 min after the interruption of the tensile test. This is presumably related to the stress relaxation associated with the imposed deformation and may also be a consequence of tritium transport by dislocations moving at a low velocity during the relaxation process. Alternatively, a slow kinetics of the tritium transfer at the interface metal-scintillator could contribute to this phenomenon which needs further investigation. The possibility of hydrogen transport enhanced by mobile dislocations at room temperature [9,10] together with the fact that H.E. sensitivity of the alloy is still observed above 523 K (Section 3.2), suggest that dislocation assisted transport during the creep of alloy 600 may contribute to the stress-corrosion cracking of this alloy at 630 K [27,28].
4.2. Hydrogen embrittlement The dependence of the H.E. sensitivity of Alloy 600 on the different parameters investigated in this study account for a contribution of the hydrogen sweeping phenomenon to the embrittlement process. (i) A predeformation leads to a reduction of the H.E. sensitivity because the dislocation network generated by the predeformation may impede the hydrogen transport by mobile dislocations. (ii) The intergranular precipitation of chromium carbides increases the H.E. sensitivity of the alloy in agreement with a mechanism of hydrogen transport and enrichment at grain boundaries during straining. The latter is presumably a consequence of preferential formation of dislocation pile-ups at the carbide enriched grain boundaries. This strain-induced hydrogen enrichment near the grain boundaries together with a possible preexisting trapping of hydrogen on the intergranular carbides is expected to favor intergranular fracture. (iii) The hydrogen embrittlement sensitivity of Alloy 600 decreases with increasing temperature of the tensile test. In a similar way, the temperature dependence of the hydrogen trapping on dislocations can be expressed to a first approximation by the following relationship [29]: c= c0/(1− c0) exp(Hb/kT)
Fig. 10. Influence of the testing temperature on the hydrogen embrittlement sensitivity of Alloy 600 (heat A, as received, do/dt= 4 × 10 − 4 s − 1).
where c, occupancy fraction of the dislocation; c0, occupancy fraction of non pertubed lattice sites; Hb, binding enthalpy of hydrogen to the dislocation. One can then anticipate a decrease of the amount of hydrogen trapped and transported by mobile dislocations when the testing temperature increases and as a consequence a reduced embrittlement effect at high temperature.
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Fig. 11. Role of the testing temperature on the fracture surface of hydrogenated Alloy 600 (heat A, as received, do/dt=4 × 10 − 4 s − 1): (a) T= 473 K (b) T =523 K.
(iv) The H.E. sensitivity of Alloy 600 increases at low strain rates together with the amount of hydrogen transported by dislocations [5]. These various effects of the different parameters on the H.E sensitivity of Alloy 600 can all be explained by a predominant effect of dislocation sweeping. This accelerated transport would lead to local overconcentration of hydrogen responsible for the crack initiation, whatever the cracking mechanism. The hydrogen transport by dislocations appears consequently as a necessary step for the embrittlement to occur. The hydrogen transport allows also to explain the change in fracture mode with strain rate. At high strain rate, the dislocation velocity is too high for significant hydrogen sweeping and the local hydrogen concentration is not sufficient to initiate a crack. At low strain rate, microcracks associated with the local hydrogen overconcentration may nucleate and will favor the initiation of a macroscopic crack. Once this crack has developed extensively, its propagation will depend on the local deformation rate at the crack tip. For moderate deformation rate, a local hydrogen enrichment on obstacles interacting with dislocations in the plastic zone away from the crack tip may favor a hydrogen-induced decohesion during further straining. This is expected to occur especially in the vicinity of grain boundaries acting as preferential sites for crack propagation and favoring intergranular brittle fracture. At high local deformation rates the absence of significant hydrogen transport to defects acting as a barrier to dislocation motion (grain boundaries) will favor a transgranular cracking with a large ductility. The frac-
ture mode is ductile but the presence of hydrogen may affect the dimple size.
5. Conclusion An accelerated desorption of tritium from Alloy 600 has been measured at 293 K with tritiated tensile specimen, by b counting during straining. The tritium desorption flow from the sample is ten times larger during straining for the present experimental conditions. These results bring strong arguments in favor of a mechanism of hydrogen transport by dislocations. The role of different metallurgical parameters on the hydrogen embrittlement (H.E.) sensitivity of Alloy 600 has been investigated with precharged tensile samples: the H.E. sensitivity at 293 K strongly depends on the microstructure of the alloy. The intergranular precipitation of chromium carbides (Cr7C3) increases the H.E. sensitivity whereas it decreases with a predeformation of the alloy; the H.E. sensitivity of any microstructures depends on the testing conditions: strain rate and temperature. The fracture mode may change from fully brittle intergranular at low strain rate to fully ductile at high strain rate with a mixed mode in an intermediate range of strain rate. The change in fracture mode as a function of the strain rate depends on the microstructure, the more sensitive the microstructure the higher the strain rate corresponding to the brittle/ductile transition. The H.E. sensitivity decreases at high temperature.
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However a significant embrittlement has been evidenced up to 523 K in the as received Alloy 600. The role of these different parameters can be rationalized in term of a determinent influence of hydrogen transport by mobile dislocations leading to a local hydrogen enrichment in the dislocation pile ups which favor crack initiation. This accelerated transport mechanism is considered to be the step which may control the hydrogen embrittlement process of alloy 600, whatever the cracking mechanism.
[12]
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Acknowledgements Financial support provided by EDF, Direction des Etudes et Recherches is gratefully acknowledged.
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