Hydrogen embrittlement

Hydrogen embrittlement

21

Historically, hydrogen embrittlement (HE) has been studied for many years in the failure of oil mining pipelines in the presence of H2S gas, offshore structures in the corrosive sea water environment, nuclear power plants with water containing hydrogen, etc. Due to the importance of this problem in related industries, a huge amount of literature on HE has been published. However, regarding the mechanism of HE, various theories, as numerous as the number of researchers, have been proposed, and a consensus of these researchers has not been well established. The term HE has been used to express the degradation of metals due to hydrogen. The exact mechanism of degradation of mechanical properties of metallic materials due to hydrogen must be elucidated, not only for very high cycle fatigue
related to nonmetallic inclusions, but more importantly to ensure the longterm safety and reliability of the hydrogen economy [1], and also to assess the structural integrity of pipelines for the oil and gas industries, nuclear engineering, and environmental degradation in the offshore structures industry. There are many studies about hydrogen-related fractures under both static stress [2
7] and cyclic stress [8
10]. Ref. [11] treats hydrogen
material interaction problems comprehensively. Regarding fatigue, Ritchie et al. [12
14] investigated crack growth behaviour for a wide range of stress intensity factors and discussed the mechanism of hydrogen-affected near-threshold crack growth properties. Nanninga [15] overviewed crack initiation and fatigue life under exposure to hydrogen. These studies mostly treated the crack growth behaviour of long cracks. Birnbaum and Sofronis [5,6] by an in situ TEM study showed that hydrogen enhanced dislocation mobility under constant stress, and by a calculation method showed that hydrogen affected the interaction between dislocations and solutes such as interstitial. Robertson [7] supported their model by in situ TEM studies and bulk mechanical property tests. The model proposed by the Illinois University group is called the Hydrogen Enhanced Localized Slip (HELP) model. However, it is still unclear how the enhancement of dislocation mobility affects fatigue behaviour. Smith and Stewart [8] showed that crack growth rates in water and hydrogen were higher than in air. Esaklul and Gerberich [9] showed that internal hydrogen caused enhancement of fatigue crack propagation rates and a decrease in threshold stress intensities. Suresh et al. [10] pointed out that the effect of hydrogen on the acceleration of crack growth was caused by less development of oxide-induced crack closure at nearthreshold crack growth rates in dry gaseous hydrogen. However, the intrinsic effect of hydrogen in metal on fatigue crack behaviour has not yet been revealed from the viewpoint of direct observation of the fracture process. Although many simulation papers on HE at the atomic level have been published, even a simple phenomenon Metal Fatigue. DOI: https://doi.org/10.1016/B978-0-12-813876-2.00021-2 © 2019 Elsevier Ltd. All rights reserved.

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Metal Fatigue

in HE introduced in this chapter has not been successfully analysed by such simulations. Most HE simulation models numerically imitate newly obtained experimental phenomena by adjusting numerical constants and have not been applied to solution of real fatigue and fracture design problems. Thus, currently most HE simulations should be regarded as a kind of computer exercise. In this chapter, first the basic mechanism of hydrogen effect on crack initiation and crack tip opening deformation will be explained based on the direct observation of slip bands and crack growth morphologies. Since the mechanics of small cracks (see Section 5.8) must be considered in addition to the hydrogen effect, the effects of hydrogen on fatigue crack growth (FCG) from the viewpoint of a small defect and the frequency effect on various steels will be discussed with emphasis. Without identifying the hydrogen content in materials, we cannot make clear the exact correlation between fatigue mechanisms and hydrogen. Therefore, all the fatigue data explained in this chapter are discussed in terms of hydrogen content in the test materials.

21.1

Effect of hydrogen on loss of ductility in tensile tests

Before we discuss the HE in fatigue, a crucially important mechanism of a hydrogen effect: the loss of ductility in tensile tests of low-carbon steel is explained. As research on HE has been carried out for the last 50 years, it is surprising that even a basic observation, such as that introduced in this section, has not been known. Fig. 21.1 [16] shows transverse sections of the noncharged and hydrogen- charged tensile specimens after tensile fracture. In hydrogen-charged steel, hydrogen concentrates at the lateral (perpendicular to tensile axis) edges of voids where there is lattice expansion due to stress concentration. This enhances slip band concentration, resulting in void growth in the lateral direction. Figs. 21.1c and 21.2 show schematically the mechanism of void growth in noncharged and hydrogen-charged low-strength steels. Although ductility is lost in tensile tests of hydrogen-charged specimens (Fig. 21.1b), it must be noted that the microscopic mechanism of fracture is ductile fracture and not brittle fracture. It will be shown later in this chapter that the basic mechanism of so-called HE in fatigue of most steels is also governed by the mechanism shown in Figs. 21.1 and 21.2.

21.2

Effects of hydrogen charge on the formation of cyclic slip bands in fatigue of annealed carbon steels [17]

It is very difficult to directly capture the behaviour of individual hydrogen atoms in metals. Uyama et al. [17] used a unique experimental method to study the intrinsic

Hydrogen embrittlement

569

Figure 21.1 Mechanism of hydrogen embrittlement in tensile fracture. Development of voids in JIS-SGP (0.078% carbon steel). The mechanism of loss of ductility (decrease in reduction of area) is the lateral growth of ductile voids and microscopic ductile fracture [16].

Figure 21.2 Schematic of void growth in noncharged and hydrogen-charged lowstrength steels.

effect of hydrogen contained in metal on the microscopic fatigue behaviour, in particular the effect of hydrogen on the formation of slip bands and fatigue cracks. In this section, by referring to their work, the evidence of the effect of hydrogen on cyclic slip bands will be shown through observation of the difference of slip morphologies without hydrogen and with hydrogen.

570

Metal Fatigue

21.2.1 Materials, specimens and experimental methods Two annealed carbon steels of 0.47% C steel and 0.45% C steel (equivalent to JIS-S45C) were used for the observation of cyclic slip bands formation without hydrogen and with hydrogen. The basic microstructures of these two steels were ferrite-pearlite with banded microstructures in the longitudinal directions. The Vickers hardness was HV 5 170 for 0.47% C steel and HV 5 185 for 0.45% C steel, respectively (the average of 10 measurements with a measurement load of 2.94 N). Fig. 21.3 shows the shapes and dimensions of three types of fatigue specimens: (a) plain specimen for tension
compression tests, (b) shallow notched specimen for tension
compression tests and (c) shallow notched specimen for rotating bending tests. The notch radius is 5 mm and the notch depths are 0.2 mm for the tension
compression test specimens and 0.3 mm for the rotating bending test specimens. These notches were introduced in order to limit the location of slip bands formation and crack initiation and make observation easy. The fatigue tests were carried out at a stress ratio, R 5
1 at room temperature and in laboratory air. S
N data were also obtained by using plain specimens at a frequency of 20 Hz. The fatigue limit was around 220
230 MPa. Slip band

5 2 R

C 1

.0 8 φ 7 . 2 1 φ 20 41.2

50.4

50.4

142

(a) Plain specimen for tension–compression tests. 5 2 R 0 . 8 φ

C 1

2 . 0

5 R

7 . 2 1 φ

20 50.4

41.2

50.4

142

R65

(b) Shallow notched specimen for tension–compression tests.

φ8

φ 15.0

0.3

R5

8.25 80

33.5 50

80

210

(c) Shallow notched specimen for rotating bending tests.

Figure 21.3 The shapes and dimensions of fatigue test specimens.

Hydrogen embrittlement

571

Hydrogen-charged (0.47%C)

1.0

Hydrogen-charged (0.45%C)

Hydrogen content, CH/ppm

Prefatigue + hydrogen-charged (0.47%C)

0.8 0.6 0.4 0.2 Uncharged: 0.05 ppm 0.0 0.1

1 10 100 Time after hydrogen charge, t/h

1000

Figure 21.4 The variation in hydrogen content with time after hydrogen charging. Hydrogen diffuses out of specimens with time [17].

morphology and fatigue crack behaviour were observed by the replica method on the shallow notched specimens. The test frequency for the shallow notched specimens was 20 Hz for tension
compression tests and 45 Hz for rotating bending tests. Hydrogen was charged by soaking the specimens in an aqueous solution of 20 mass% NH4SCN at 310K for 24 h. The hydrogen content in specimens was measured by thermal desorption spectrometry (TDS). However, as shown in Fig. 21.4, the hydrogen content decreases with time after hydrogen charging due to diffusion out of specimens at room temperature. The details of the variations in hydrogen content with time after hydrogen charging are shown in Ref. [17]. The fatigue tests were started within 4 h after hydrogen charging. Thus, detecting the hydrogen content in specimens is crucially important for exact discussion of HE. Discussion without paying attention to hydrogen content will lead us to misconception of HE.

21.2.2 Effects of hydrogen on slip band morphology and crack initiation near the fatigue limit stress For ease of observation of microscopic fatigue processes, rotating bending tests using the shallow notched specimen were carried out to observe slip bands in more detail. Fig. 21.5 compares the slip bands and fatigue cracks on the specimen surfaces in the uncharged specimen (Nf 5 2.53 3 105) and the hydrogen-charged specimen (Nf 5 2.16 3 105) at a stress amplitude, σa 5 230 MPa. There is a definite difference in the slip bands between the uncharged specimen and the hydrogen-charged specimen. Densely and widely distributed slip bands were formed at most of the ferrite grains in the uncharged specimen. On the other hand, discrete and localised slip bands were nucleated in the hydrogen-charged specimen. In order to investigate the effect of the change in slip morphology shown in Fig. 21.5 on fatigue, the cyclic stress
strain property was measured. Fig. 21.6 shows the difference of cyclic stress
strain properties in terms of hydrogen. It is

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Metal Fatigue

t* = 4 h

t* = 6 h N=5x

104

t* = 7 h N=1x

105

t* = 9 h N = 2 x 105

H-charged (Nf = 2.16 x 105)

Uncharged (Nf = 2.53 x 105)

N = 1 x 104

Axial direction

50 μm

t* shows time after hydrogen charge.

Figure 21.5 Slip bands and fatigue cracks of the hydrogen-charged and uncharged specimens of 0.47% C steel in tension
compression test (σ a 5 230 MPa).

Figure 21.6 Cyclic hysteresis loops of the hydrogen-charged and uncharged specimens of 0.47% C steel.

Hydrogen embrittlement

573

presumed from Figs. 21.5 and 21.6 that the decrease in the cyclic total strain range in the hydrogen-charged specimens was caused by the localisation of slip bands. In the hydrogen-charged specimen, the region where plastic deformation occurred is smaller than that in the uncharged specimen because of localised slip bands. The macroscopic strain is the result of the accumulation of local slip deformation. Therefore, the total amount of the strain corresponding to the slip bands in the hydrogen-charged specimen became smaller than that in the uncharged specimen. These experimental data imply that if the fatigue tests are carried out under a constant strain amplitude, the fatigue life for the hydrogen-charged specimen will be much shorter than the uncharged specimen, because the larger plastic deformation must be supported by a small number of ferrite grains with localised slip. Fig. 21.7 shows micrographs of the slip bands which were taken by a laser microscope using the replica method in the uncharged specimen and the N=0

N = 6 × 105

(a) Type A (Hydrogen-charged specimen, σa = 230 MPa) N=0

N = 6 × 105

(b) Type B (Uncharged specimen, σa = 230 MPa)

N=0

N = 6 × 105

(c) Type C (Hydrogen-charged specimen, σa = 230 MPa) Axial direction Ferrite grain

Type A

Type B

15 μm Slip bands

Type C

Figure 21.7 Micrographs and schematic illustrations of three types of slip bands.

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Metal Fatigue

hydrogen-charged specimen of the 0.45% C steel. Comparing the micrographs at N 5 0 and N 5 6 3 105 in Fig. 21.7, the slip bands were mostly observed in ferrite grains. The morphologies of slip bands can be classified to three types, that is type A, type B and type C: Type A: Slip bands cover a whole ferrite grain and cross each other (Fig. 21.7a); Type B: Slip bands cover a whole ferrite grain in parallel uniformity (Fig. 21.7b); Type C: Discrete localised slip bands exist in ferrite grains in almost the same direction in a zebra pattern (Fig. 21.7c).

In order to evaluate quantitatively the effect of hydrogen on slip bands, types of slip bands in 100 ferrite grains were investigated at the notch root and were classified into type A, B or C. Fig. 21.8 shows micrographs of the notch root in the hydrogen-charged, uncharged and aged specimens. Each ferrite grain was coloured according to the classification of the type of slip bands. Fig. 21.9 shows the number of ferrite grains with the three types of slip band in the uncharged, hydrogen-charged and aged specimen. In the uncharged specimen the type A slip bands were dominant and type C was not observed. On the other hand, in the hydrogen-charged specimen 49% of the slip bands were type C. Thus, it is evident that the formation of the type C slip bands was caused by a synergetic effect of hydrogen and cyclic stress. In the specimen aged for 270 h in air at room temperature after hydrogen charging, the percentage of the type C was 7%, because approximately only 0.2 ppm of hydrogen remained in the specimen 270 h after the hydrogen charge (see Fig. 21.4). Fig. 21.10 shows the number of fatigue cracks and the crack initiation sites after 107 cycles in the hydrogen-charged specimen, the uncharged specimen and the aged specimen at the fatigue limit (230 MPa in the rotating bending test). To obtain these data an area of 0.65 mm2 on the notch root in each specimen was observed with a microscope and cracks having a length longer than two grains were defined as fatigue cracks. Fatigue cracks were predominantly initiated at grain boundaries in all specimens. It follows that persistent slip bands are not necessarily fatigue crack initiation sites for smooth specimens. The percentage of cracks initiating along slip bands in ferrite grains is small in the uncharged specimen. On the other hand, in the hydrogen-charged specimen approximately 20% of fatigue cracks were initiated along slip bands in ferrite grains. Moreover, the total number of fatigue cracks in the hydrogen-charged specimen was more than in the uncharged specimen. It can be interpreted that these cracks are microscopic nonpropagating cracks, as explained in Chapter 1, Mechanism of fatigue in the absence of defects and inclusions. Although some unique differences in crack initiation morphology between noncharged specimens and hydrogen-charged specimens were observed, the fatigue limits in both specimen types did not change within the hydrogen content charged in this experiment. Fig. 21.11 [17] shows the cross-section profiles of type A and type C slip bands measured by an atomic force microscope (AFM). The true configurations of the specimen surface are the reverse of Fig. 21.11, where the specimen surface configuration was obtained by the replica method. It can be seen that a ferrite grain with

Type A Type B Type C No slip bands

100 μm (a)

Hydrogen-charged specimen.

100 μm (b)

Uncharged specimen.

100 μm (c)

Aged specimen.

Figure 21.8 Micrographs of the notch root with 100 ferrite grains coloured according to the type of slip band.

Type A

Type B 4

38

Hydrogencharged

Type C

No slip bands 49

9

60

17

23

56

29

7 8

Uncharged Aged 0

20

40 60 80 Number of ferrite grains

100

Figure 21.9 Number of ferrite grains with the three types of slip band (N 5 6 3 105, σ a 5 230 MPa).

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Metal Fatigue

Grain boundary 102

Hydrogencharged

Ferrite grain 27

2

77 Uncharged 61

5

Aged 0

50 100 Number of fatigue cracks

150

Figure 21.10 Number of fatigue cracks on the specimen surface (N 5 1 3 107, σ a 5 230 MPa).

type A slip bands has steep steps at the grain boundaries. On the other hand, there are discretely distributed steep steps inside a ferrite grain having type C slip bands. It is presumed from the data already show that these steps are likely to become the initiation sites for fatigue cracks. Thus, the change in the slip band morphology is induced by hydrogen contained in ferrite grains and leads to an increase in the number of fatigue cracks emanating from slip bands in ferrite grains in the hydrogencharged specimen. Birnbaum and Sofronis [5] showed with an in situ TEM study that hydrogen enhanced dislocation mobility under constant stress. Enhancement of dislocation mobility due to hydrogen can lead to localised slip bands in ferrite grain under cyclic stress because a certain slip system in ferrite grains can work with priority. Furthermore, hydrogen is attracted to the localised slip bands and then the localised plasticity can become much easier. It is presumed that this is the mechanism of more crack initiations within the ferrite grains in the hydrogen-charged specimen.

21.3

Effects of hydrogen charge on the mechanism of fatigue crack growth of low-strength steels

21.3.1 Fatigue crack growth behaviour of a pipeline steel [18] Fig. 21.12 shows the effect of test frequency on the FCG rate of 10% prestrained pipeline steel. Crack growth rate is accelerated at the test frequency f lower than 0.1 Hz and the acceleration ratio is more than 10 times below f 5 0.01 Hz [18]. Fig. 21.13 compares the difference in slip morphologies near crack tips of hydrogen-charged specimens under different load frequencies [18]. Slip band morphology for noncharged specimens was almost the same regardless of test frequency, as shown in Fig. 21.12a and b for H-charged specimen tested at f 5 10 Hz and 0.1 Hz. The slip bands for H-charged specimen tested at f 5 0.01 Hz are localised only near crack tips. Nevertheless, the crack growth rate of the H-charged specimen is markedly accelerated more than 10 times at low frequency.

Hydrogen embrittlement

577

Figure 21.11 Micrographs, schematic diagrams and cross-section profiles of type A and type C slip bands measured by AFM.

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Metal Fatigue

[(m/cycle)/(MPa ⋅m1/2)2]

((da/dN)/ΔK2, SZWT,max/K2max

10–8

10–9

Hydrogen charged

–10

10

Uncharged

10–11 10–4 10–3 10–2 10–1 100 101 102 Test T frequency f (Hz)

Hydrogen charged Replica (ΔK = 80–81 MPa m1/2) Striation (ΔK = 45–97 MPa m1/2) SZWT,max (Kmax = 138 MPa m1/2) Uncharged Replica (ΔK = 79 – 80 MPa m1/2) Striation (ΔK = 44 – 91 MPa m1/2) SZWT,max (Kmax = 137 MPa m1/2)

⋅

⋅

⋅

⋅

⋅ ⋅

Figure 21.12 Effect of test frequency on FCG rate of 10% prestrained pipeline steel. Crack growth rate was measured by the replica method and striations [18].

Figure 21.13 Slip deformation near crack tips in hydrogen-charged specimens [18]. Residual hydrogen content CH,R 5 1.1 ppm. R-ratio R 5
1.

In order to identify the mechanism in a more definite way, an experiment was planned to observe the difference of the formation of the stretch zone width (SZW) for uncharged specimens and hydrogen-charged specimens [18]. The stretch zone morphology reflects the crack tip opening displacement of fatigue crack with magnification. The maximum stretch zone size (SZWT,max) was measured by tilting the fracture surface as shown in Fig. 21.14b. If crack tip blunting is large, then a large tilting angle is necessary for measurement of SZWT,max. Fig. 21.14c shows the loading pattern to produce a stretch zone during a fatigue test. Fig. 21.15 compares by SEM the values of SZWmax for a noncharged specimen and a hydrogen-charged specimen. The measured tilt angles in noncharged and Hcharged specimens revealed the peculiar crack growth mechanism in the presence of hydrogen as illustrated in Fig. 21.16. From the results of Fig. 21.15, we can illustrate the configuration of stretch zones for both an uncharged specimen and a hydrogen-charged specimen as in Fig. 21.16b, which shows that the crack in a hydrogen-charged specimen grows without significant blunting. This mechanism is very similar to that shown in Figs. 21.1 and 21.2.

Hydrogen embrittlement

579

Figure 21.14 Introduction and measurement of stretch zone at a fatigue crack tip. The maximum stretch zone size (SZWT,max) was measured by tilting the fracture surface. If crack tip blunting is large, then a large tilting angle is necessary for measurement of SZWT,max. The measured tilt angles in uncharged and H-charged specimens revealed the peculiar crack grow mechanism in the presence of hydrogen as shown in Fig. 21.16b.

Figure 21.15 Stretch zone of noncharged specimen (CH,R 5 0 ppm) at Kmax 5 137 MPa m1/ 2 and hydrogen-charged specimen (CH,R 5 1.1 ppm) at Kmax 5 138 MPa m1/2.

From the above-mentioned results and observations, the FCG mechanism under hydrogen effect can be illustrated as shown in Fig. 21.17. These illustrate the effect of hydrogen on the crack closure mechanism during one load cycle. Fig. 21.17 (a-1)!(a-4) shows the crack opening behaviour on the way to the maximum load in the absence of hydrogen. The crack tip opening displacement reaches its

580

Metal Fatigue

SZWT, max SZWT, max (a) Uncharged (b) Hydrogen charged

Figure 21.16 Difference in stretch zone configuration between noncharged and hydrogen-charged specimens. The crack in hydrogen-charged specimen grows without significant blunting [18].

a

P 1

t

a

(a-1)

(b-1)

(a-2)

(b-2)

(a-3)

(b-3)

(a-4)

(b-4)

(a-5)

(b-5)

(a-6)

(b-6)

P

Loading process

2

t P

Plastic zone

Plastic zone

t

3

(c) No hydrogen effect

(d) Hydrogen effect

P

Unloading process

4

t

P t

5 P 6

t

(a) No hydrogen effect

(b) Hydrogen effect

Figure 21.17 Mechanism of FCG acceleration by hydrogen concentration at crack tip under low test frequency. Hydrogen concentration enables successive crack growth without crack tip blunting, with small CTOD [19].

saturated value at a given load level and crack growth ceases as shown in Fig. 21.17 (a-1)!(a-4). However, hydrogen concentrates near the crack tip in the presence of hydrogen. Hydrogen concentration enhances further crack opening by slip, and crack growth continues. Since the corresponding plastic zone at the crack tip does not become large, the plastic zone wake which remains on the fracture surface is shallow. The difference in crack tip blunting mechanism between no Heffect and H-effect is essentially the same as in Figs. 21.1, 21.2 and 21.16. Figs. 21.17c and d are schematic illustrations of plastic zone wakes without and with hydrogen. This phenomenon results both in a decrease in the height of striations (see Fig. 21.24) and a decrease in the crack opening load (decrease in ΔKop and increase in ΔKeff).

21.4

Effect of hydrogen on fatigue behaviour of a Cr
Mo steel SCM435 [20]

Cr
Mo steel JIS SCM435 is a typical steel which is used for hydrogen storage cylinders at hydrogen fuelling stations for fuel cell vehicles (FCVs). Currently, the

Hydrogen embrittlement

581

–5

10

–6

Crack growth rate, da/dN(m/cycle)

10

–7

10

(Hydrogen content:

0.01 ppm) Constant frequency

–8

: f = 20 Hz

10

Frequency switched

: f = 2 Hz : f = 0.02 Hz

10

10

Hydrogen-charged Constant frequency : f = 20 Hz (0.53–0.27 ppm) : f = 2 Hz (0.58–0.49 ppm) : f = 0.2 Hz (0.58–0.49 ppm) Frequency switched : f = 2 Hz (0.58–0.29 ppm) : f = 0.02 Hz (0.58–0.29 ppm)

–9

–10

10

20

30

Schematic image of the mechanism of effect of hydrogen and test frequency on fatigue crack growth.

40 50 60 708090100

Stress intensity factor range, ΔK(MPa m1/2)

Figure 21.18 Effect of hydrogen on FCG behaviour of a Cr
Mo steel SCM435 [20]. Relationship between da/dN and ΔK. Material: SCM435. Hydrogen content indicated by 0.xx!0.yy ppm means that hydrogen content decreased from 0.xx to 0.yy ppm during the fatigue test. ‘Frequency switched’ means that the test frequency was switched between f 5 2 Hz and f 5 0.02 Hz.

maximum hydrogen gas pressure for FCV is 70 MPa. As shown in Fig. 21.18, the crack growth rates under the hydrogen effect are 30 times higher than those for noncharged specimens. The slip bands of the H-charged specimens were localised only at a very narrow area beside the crack line as in the case of pipeline steel (see Figs. 21.13 and 21.17). The reason that crack growth curves of H-charged specimens approach at high ΔK to those of noncharged specimens is that the crack grows faster than does hydrogen diffusion and concentration to the crack tip. These characteristics are commonly observed in H-charged specimens of other steels [1].

21.5

Effect of hydrogen on fatigue behaviour of austenitic stainless steels

21.5.1 Basic parameters: hydrogen content, diffusion coefficient, fatigue crack growth and test frequency It is well known that austenitic stainless steels have higher HE-resistance than other bcc steels. It is also known, as shown in Fig. 21.19, that the hydrogen content of austenitic stainless steels is commonly 2
4 mass ppm in the as-received condition (curve A of Fig. 21.19). Nevertheless, as explained in the previous sections, bcc

Metal Fatigue

Hydrogen desorption intensity (×10–10 A)

582

6 Heating rate: 0.5ºC/s

5 H-charged (cathodic) 3.6 wppm

4 B

3

A

2 1 0

Uncharged 2.6 wppm

C 0

100

200

300 400 500 Temperature (ºC)

NDH-HT 0.4 wppm

600

700

Figure 21.19 Hydrogen thermal desorption spectrum of type 316L. As-received austenitic stainless steels (curve C) always contain 2
3 ppm hydrogen [21].

steels are more sensitive to HE even with one order lower hydrogen content compared with austenitic stainless steels. Fig. 21.20 shows the FCG curves starting from a drilled hole of 100 μm diameter and 100 μm depth in the uncharged and hydrogen-charged specimens of SUS304 (type 304), SUS316 (type 316) and SUS316L (type 316L). The hydrogen content values in this figure were measured immediately after the fatigue test by TDS. For both SUS304 and SUS316, the increased FCG rate due to increased hydrogen content is clear. The crack growth rate in the hydrogen-charged SUS316L was only slightly higher than in the uncharged SUS316L. In the austenitic stainless steels, it can be assumed that hydrogen in a specimen does not redistribute macroscopically during the fatigue test (within 100 h). This is because of the very low hydrogen diffusion coefficient, that is D323KB10215 m2/s (see Table 21.1). Thus, hydrogen loss during the fatigue test is considered to be negligible. Smith and Stewart [8] carried out fatigue tests at frequencies of 0.02
50 Hz and showed that FCG of 2Ni
Cr
Mo
V steel was more accelerated by hydrogen at the lower test frequency. It must be noted that fatigue tests for Fig. 21.20 were carried out at frequencies of 1.2
5 Hz. For reference, the hydrogen diffusion coefficients are summarised in Table 21.1. Fig. 21.21 shows the fatigue cracks emanating from a small hole (same as Fig. 21.20) in the uncharged and hydrogen-charged specimens of SUS304 and SUS316L. The crack growth paths in the hydrogen-charged specimens of SUS304 and SUS316L are less tortuous than the paths of the uncharged specimens. The less tortuous crack growth paths in the hydrogen-charged specimens are observed, not only in austenitic stainless steels, but also in other steels [17]. The relatively linear crack growth path is caused by hydrogen-enhanced slip localisation at fatigue crack tips, and can be one of the indicators for hydrogen effects. Fig. 21.22 shows the more microscopic morphology of slip bands in the vicinity of a crack tip in both the hydrogen-charged specimens and in the uncharged specimens of SUS304 and SUS316L. From this figure it can be observed that the number

Hydrogen embrittlement

Figure 21.20 Influence of hydrogen charging on crack growth from 100 μm hole for austenitic stainless steels type 304, type 316 and type 316L. Hydrogen charging was carried out at 50 C for 672 h by cathodic charging [22].

583

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Metal Fatigue

Table 21.1 Hydrogen diffusion coefficients in specimens. (a) Hydrogen diffusion coefficient at 323K in solution-treated austenitic stainless steels and an annealed ferritic stainless steel (SUS405), and (b) hydrogen diffusion coefficient at 353K in prestrained specimens of SUS304 and SUS316L and a solution-treated specimen of SUS304 [22] (a) Type

Hydrogen diffusion coefficient at 323K (m2/s)

304

1.1 3 10215

316

2.0 3 10215

316L

2.4 3 10216

405

2.8 3 10212 (b)

Type

Hydrogen diffusion coefficient at 353K (m2/s)

Solution-treated 304 (α0 : B3%)

5.5 3 10215

304 prestrained at 29 C (εp 5 0.29, α0 : 1%
3%)

5.5 3 10215

304 prestrained at 270 C (εp 5 0.28, α0 : 65%
69%)

2.4 3 10213

316 L prestrained at 270 C (εp 5 0.30, α0 : 26%
28%)

5.4 3 10214

εp, plastic prestrain; α0 , martensite content.

of slip bands after cyclic loading is greater for the uncharged specimens, whereas more localised slip bands are observed in the hydrogen-charged specimens (see the arrows in Fig. 21.22). Slip bands in the vicinity of the crack tip of the hydrogencharged specimens are more parallel and discrete than those of the uncharged specimens. This is especially clear for the SUS304 specimens where slip bands are distinctly more localised as compared to slip bands in the SUS316L. Localisation of slip bands is assumed to be coupled to hydrogen diffusion and increased hydrogen concentration near the fatigue crack tip. In the SUS304, transformed martensite fractions of up to 50%
70% were observed in the vicinity of the fatigue crack tip. Hydrogen can diffuse through the transformed martensite much faster than through austenite. Therefore, it can be hypothesised that hydrogen in the SUS304 more easily concentrates near the fatigue crack tip than in the SUS316L. Increased hydrogen content in the vicinity of the crack tip enhances further slip localisation. This is consistent with the studies of Birnbaum and his colleagues where flow stress, as estimated from dislocation velocity, decreased as the hydrogen gas pressure

Figure 21.21 Observed fatigue crack paths for SUS304 and SUS316 both with and without hydrogen charging. Hydrogen charging reduces tortuosity.

Figure 21.22 Slip bands in the vicinity of the crack tips of hydrogen-charged and uncharged specimens of SUS304 and SUS316L (SUS304: σ 5 260 MPa, f 5 1.5 Hz, CH 5 6.7 wppm, CU 5 2.2 wppm, SUS316L: σ 5 280 MPa, f 5 1.5 Hz, CH 5 5.1 wppm, CU 5 2.6 wppm) [22].

586

Metal Fatigue

Figure 21.23 Influence of hydrogen and test frequency on crack growth, from 2a 5 200 μm, of (a) SUS304 (σ 5 280 MPa), (b) SUS316L (σ 5 280 MPa) and (c) effect of hydrogen and test frequency on crack growth rate of SUS316L [21].

increased [23
28]. Increased slip localisation at the fatigue crack tip due to hydrogen implies a decrease in the cyclic plastic zone size which is consistent with the results of Katz and Gerberich [29] in which the overload-affected zone size was reduced by hydrogen. All these observations are consistent with the mechanism shown in Fig. 21.17. Figs. 21.23a and b show the relationship between the crack length, 2a, and the number of cycles, N, for type 316L and type 304 specimens. The hydrogen contents of the hydrogen-charged specimens, which were measured immediately after fatigue tests, are also included in the figure. The crack length is defined by adding the initial hole diameter. Fig. 21.23c shows the relationship between the FCG rate, da/dN, and the stress intensity factor range, ΔK, for type 316L. In type 316L tested at a frequency of 1.5 Hz, there was no pronounced difference between the hydrogen-charged specimen and the uncharged specimen. The FCG rates of the hydrogen-charged type 316L, tested at the frequency f 5 0.0015 Hz, were two to three times higher than those of type 316L tested at f 5 1.5 Hz. In particular, in the ΔK range from 8 to 10 MPa m1/2, the increase in the crack growth rate was

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Figure 21.24 [21] Difference in crack growth behaviour between hydrogen-charged specimens and uncharged specimens of type 316L (σ 5 280 MPa): (a) uncharged (f 5 1.5 Hz, 2a 5 743 μm, 2.6 wppm), (b) H-charged (f 5 1.5 Hz, 2a 5 678 μm, 5.1 wppm), (c) uncharged (f 5 0.0015 Hz, 2a 5 841 μm, 2.6 wppm) and (d) H-charged (f 5 0.0015 Hz, 2a 5 684 μm, 3.9 wppm).

remarkable. Surprisingly, the uncharged specimen also exhibited an obviously definite crack growth rate increase as the test frequency decreased from 1.5 Hz to 0.0015 Hz. In type 304, as well as in type 316L, the uncharged specimen exhibited a definite frequency effect. As the test frequency decreased from 1.2 Hz to 0.0015 Hz FCG rates increased (Fig. 21.23b). Fig. 21.24 shows the effect of hydrogen and test frequency on the microscopic morphologies of crack growth. More slips bands become visible for f 5 0.0015 Hz, even for uncharged specimens. It is evident that hydrogen has sufficient time for diffusion under f 5 0.0015 Hz, even with a very slow diffusion rate of D 5 B10215 m2/s.

21.5.2 What happens if nondiffusible hydrogen is removed by a special heat treatment [21]? Even in the uncharged specimens the FCG rate is increased by decreasing the test frequency down to 0.0015 Hz. The only suspected cause for this surprising phenomenon is the effect of so-called nondiffusible hydrogen, of the level of 2
3 wppm, which is unavoidably trapped in the material during the production process. The nondiffusible hydrogen unavoidably contained in solution-treated austenitic stainless steels is different from the diffusible (or reversible) hydrogen, which is charged

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into a material by a hydrogen gas environment, or by electrochemical methods. Historically, this nondiffusible hydrogen has not been suspected in HE studies as a possible cause of HE. However, if this nondiffusible hydrogen can cause HE, it means that the HE of austenitic stainless steels can occur even in the absence of an external hydrogen environment. In order to investigate the influence of so-called nondiffusible hydrogen on FCG, a special heat treatment was applied. This is soaking type 316L at 450 C for 450 h, and type 304 at 300 C for 600 h, in a vacuum of 6.0 3 1023 Pa. In practice this heat treatment was carried out by enclosing specimens, or samples, in an evacuated silica glass tube. During the heat treatment process the partial pressure of hydrogen in the glass tube increases, and this prevents the removal of hydrogen in a specimen to below the equilibrium hydrogen content for the increased partial pressure. Therefore, in order to avoid unwanted increases in the hydrogen partial pressure, a specimen was removed from the glass tube after heat treatment for 150 h, and the same vacuum heat treatment process was then repeated in 150-h increments. Fatigue testing was carried out using specimens which were prepared using the special heat treatment. The specification for this special heat treatment was determined by considering the temperature and time needed to avoid sensitization, and also to ensure hydrogen diffusion out of specimens. This heat treatment is completely different from the socalled ‘baking’ which is conventionally applied to remove hydrogen introduced by processes such as welding and plating. Nondiffusible hydrogen at the level of 2
3 wppm cannot be removed by the so-called baking. This special heat treatment is called nondiffusible hydrogen desorption heat treatment (NDH-HT) [21]. The effect of NDH-HT on FCG is clearly shown in Fig. 21.23, in which the hydrogen content in an ordinarily heat-treated type 316L is 2.6 wppm, and on the other hand in the sample subjected to NDH-HT the content is 0.4 wppm. Thus, NDH-HT removes the hydrogen, which is strongly trapped at the centre (O-site) of an octahedron of the FCC lattice, where the potential energy is much lower than that for hydrogen trapped in the quadratic lattice of a bcc material (T-site) [30]. The hydrogen content after NDH-HT was measured by TDS. In order to check that the material structure had not been sensitised by NDH-HT the microstructure was etched, and this confirmed that chromium carbides did not segregate at grain boundaries. As further evidence of the avoidance of sensitising, the Vickers hardness before and after NDH-HT was found to be largely unchanged (for type 304 HV 5 176 before and 172 after NDH-HT, for type 316L HV 5 157 before and 163 after NDH-HT). Fig. 21.23b shows the FCG behaviour of NDH-HT specimens. Fatigue tests were carried out at a frequency of 1.5 Hz until a crack had grown to 200 μm in length, including the hole diameter (100 μm), and the test frequency was then changed to 0.0015 Hz. Surprisingly, the FCG rate of an NDH-HT specimen was substantially decreased in comparison with a hydrogen-charged specimen, and even in comparison with an uncharged specimen. Further evidence of the effect of NDHHT was confirmed by the measurement of the striation height-spacing ratios (H/s) which were shifted to higher values [21,22]. The above facts are definite proof that

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even the nondiffusible hydrogen contained in ordinary solution-treated austenitic stainless steels influences FCG rates. From the result of this experiment we can identify the following two reasons for the increase in fatigue crack rates, at the very low test frequency of 0.0015 Hz, in solution-treated uncharged specimens, which naturally contain hydrogen at the level of 2
3 wppm. 1. At a very low test frequency, of the order of 0.0015 Hz, there is sufficient time for hydrogen to be transported to fatigue crack tips through lattice defects such as edge dislocations. In this case, it is presumed that the probability of the movement of the hydrogen out of O-sites is increased by the fluctuation of the potential energy due to lattice deformation under very low stress frequencies. 2. As austenite at fatigue crack tips partially transforms to strain-induced martensite, the binding energy between hydrogen and a trapping site is decreased, and the transformed martensite plays the role of the hydrogen diffusion highway [22,31,32] which, eventually, enhances the diffusion and concentration of hydrogen trapped in austenite to fatigue cracks.

For example at 0.0015 Hz, with ΔK 5 10 MPa m1/2 for which the plastic zone size is approximately 20 μm, during one fatigue cycle (11 min) hydrogen can move 0.4 μm in plastically deformed austenite having a slightly higher diffusion rate than that for an undeformed structure (Table 21.1), whereas it can move 3.5 μm through transformed martensite. Thus, it is vitally important to take the effect of loading frequency into consideration for the prediction of fatigue lives and the safety of infrastructures and components which, in actual service, are used in hydrogen environments. Current fatigue data obtained by conventional accelerated tests should therefore not be used for design to ensure the long-term safety of hydrogen energy systems. Therefore, fatigue test methods must be reviewed from the viewpoints both of test frequency and of hydrogen content.

21.5.3 Hydrogen-induced striation formation mechanism The distributions of maximum shear stress and of hydrostatic tensile stress, ahead of the crack tip, under plane strain can be easily calculated from the elastic solution for a crack. In the case where there is no hydrogen, slip from the crack tip occurs in the 75.8 direction, where the shear stress has its maximum under plane strain. Slip in the 75.8 direction causes both crack tip blunting and crack growth at the initial stage of loading. Under a given load level, crack tip blunting occurs as a crack grows and, finally, at the maximum load crack growth is saturated. This mechanism has been well known in previous studies on metal fatigue [33
36]. On the other hand, for the case where hydrogen is present, Sofronis and McMeeking [37] showed, by numerical analysis of hydrogen diffusion near the crack tip, that hydrogen diffuses to, and concentrates at, the region where the hydrostatic tensile stress has its maximum. Tabata and Birmbaum [23] suggested, through TEM observation of the interaction between dislocations and hydrogen, that yield stress decreases as hydrogen pressure increases. Considering their experimental result, it is therefore presumed that yield stress decreases at a region where hydrogen concentrates. As a result, crack tip blunting and crack growth both occur during the whole load cycle.

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Namely, even if crack tip blunting occurs at a given load level that is below the maximum load, further slip takes place at the growing crack tip where hydrogen repeatedly concentrates. This further slip reduces crack tip blunting in the 75.8 direction; both crack tip blunting and crack growth occur in a coupled manner during the whole load cycle. Sofronis et al. proposed a new constitutive equation which includes a parameter for the decrease in yield stress due to the presence of hydrogen [32,38], and also simulated the growth and coalescence of voids due to the presence of hydrogen in a material [39]. Their simulation indicated that, if the hydrostatic stress is larger than the maximum shear stress, then the growth of shallow voids is accelerated by the presence of hydrogen. As shown in Fig. 21.17, the FCG mechanism of ductile materials is based on striations formed by slip at a crack tip. This differs from the static fracture mechanism of bcc metals. However, the diffusion and concentration behaviour of hydrogen near a crack tip, or near a notch root, is similar in both fcc and bcc metals. Furthermore, with decreasing fatigue test frequency, there is sufficient time for hydrogen to diffuse towards crack tips, and a large amount of hydrogen concentrates near crack tips. As a result, a crack continues to grow before the crack tip becomes fully blunt. Bichler and Pippan [36] showed, by 3D analysis of SEM stereophotographs of striations, that the bottom of a striation coincides with the crack tip at the end of the previous loading cycle, as illustrated in Fig. 21.17a. This means that the upper and lower angles of the V-shape at a crack tip do not close during unloading, and that the bottoms of striation are formed at these angles. Murakami et al. [21] conducted 3D analysis of striations in hydrogen-charged specimens, and showed that the peaks and valleys of striations on the upper and lower fracture surfaces correspond to each other. Hence, the basic mechanism of striation formation is the same in both hydrogencharged and uncharged specimens. As has been described in the previous paragraph, a crack grows continuously during loading in the presence of hydrogen, even before the crack opening displacement reaches its maximum value. Consequently, the crack tip shape at the maximum load is sharper in the presence of hydrogen than in its absence. The effect of hydrogen on plastic deformation at a crack tip during unloading is supposed to be reduced. This is because the stress field at the crack tip becomes compressive. It is presumed that the vertical distance between the peak and the valley of a striation becomes small. This is because the crack opening displacement at the maximum load is small in the presence of hydrogen, even though the amount of reverse slip is the same as that in air. This is the possible mechanism for the small ratio of striation height, H, to spacing s, H/s, in the presence of hydrogen [22].

21.5.4 Case study: dispenser hose fatigue failure at a hydrogen station [40] Fig. 21.25 shows a schematic view of the dispenser at the hydrogen station of EXPO 2005 in Nagoya in which a hydrogen leakage incident from the flexible dispenser hose occurred in May, 2005. The dispenser material was type 316L austenitic stainless steel. At the design stage the effect of hydrogen on type 316L was not

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Type 304 [Furukawa and Murakami et al.(2010)]

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Uncharged Type304 (1.5 Hz, 2.2 mass ppm) H-charged Type304 (1.5 Hz, 6.7 mass ppm)

0.5 0.6 0.7 0.80.9 1

2

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1-R Figure 21.25 [40] (a) Incident of dispenser hose at a hydrogen station in 2005. Hydrogen pressure: 35 MPa. Number of hydrogen supplies to buses (number of pressure cycles): 280. (b) Relationship between ratio of striation height H to spacing s, H/s, and stress ratio (1
R). Effect of hydrogen environment on relationship between ratio of striation height H to spacing s in stainless steels. The values of H/s in hydrogen effect are smaller than in air.

considered, because type 316 austenitic stainless steel had been approved for use in high-pressure hydrogen facilities on the basis of existing information that the fatigue properties of type 316L austenitic stainless steel are not affected by hydrogen. However, the incident occurred after an unexpectedly small number of service cycles. By failure analysis [40], the reason for the difference between the predicted and the actual fatigue life was identified as the low load frequency in actual service (8 min for one fuelling cycle) in the presence of hydrogen. The failure analysis of the fracture surface revealed the presence of 270 striations for the 280 service cycles and the evidence of a hydrogen effect from the lower values of the ratio between striation height and spacing, H/s (see Fig. 21.25b). The mechanism can be understood from the frequency effect on FCG, which was explained in detail in the previous sections.

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21.5.5 Hydrogen effect against hydrogen embrittlement [41] As described in the previous section, in the noncharged specimens, many grains were covered with slip bands around a fatigue crack, while in the H-charged specimens, slip bands appeared only in the vicinity of a crack, but, under the same stress, fewer slip bands were observed in grains away from the crack (see Figs. 21.13 and 21.24). This phenomenon shows that hydrogen increased the resistance against crystallographic glide, that is hydrogen hindered dislocation motion, eventually leading to hardening. Nevertheless, FCG rates for the H-charged specimens were significantly increased compared with those for noncharged specimens. What does this contradiction in the results of optical microscopic observations mean? In order to resolve the mystery of HE, Murakami et al. [41] paid particular attention to the quantitative effects of H-content, ranging from a noncharged level to supersaturated levels in the material. From the term ‘HE’, it is natural to assume that the higher the H-content in a material, the lower the strength properties of the material, as shown in the previous section. Murakami et al. [41] obtained a uniform supersaturated high H-content distribution throughout fatigue specimens by exposing types 304 and 316L specimens to gaseous hydrogen at a pressure up to approximately 100 MPa and at a temperature of 553K (280 C). In tests on specimens containing supersaturated H, the appearance of a strong HE was naturally anticipated, but what actually happened was surprising and dramatic. It is expressed as ‘hydrogen effect against HE’. The unprecedented experimental results are presented in Fig. 21.26 [41]. Kirchheim [42] gave an explanation to this effect based on the mechanism of double kinks locked by excess hydrogen. The well-known term ‘hydrogen embrittlement’ (HE) expresses undesirable effects due to hydrogen such as loss of ductility, decreased fracture toughness, and degradation of fatigue properties of metals. However, a dramatic phenomenon was found in which charging a supersaturated level of hydrogen into specimens of austenitic stainless steels of types 304 and 316L drastically improved the FCG Type 304, R = –1, σ = 280 MPa

Crack length 2a (m)

3000

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2000

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0

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Figure 21.26 Effect of H-content on FCG. Hydrogen effect against HE.

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Figure 21.27 Hydrogen atoms pinned by an edge dislocation.

resistance, rather than accelerating FCG rates. Although this mysterious phenomenon has not previously been observed in the history of HE research, its mechanism can be understood as an interaction between hydrogen and dislocations. Hydrogen can play two roles in terms of dislocation mobility: pinning (or dragging) and enhancement of mobility. Competition between these two roles determines whether the resulting phenomenon is damaging or, unexpectedly, desirable. This finding will not only be the crucial key factor to elucidate the mechanism of HE, but also be a trigger to review all existing theories on HE in which hydrogen is regarded as a dangerous culprit. Fig. 21.27a illustrates the pinning effect of hydrogen on a dislocation. The hydrogen trapped by a dislocation increases the critical shear stress for dislocation glide. Hydrogen trapped at a dislocation core, in a dilatational stress field, pins an edge dislocation. The higher the hydrogen concentration, the stronger the pinning effect. It follows that if the hydrogen concentration at dislocation is supersaturated, then the pinning effect should be particularly strong. This pinning effect contributes to increasing the strength of materials. The effect explains why there are fewer grains with slip bands in the case of specimens with high hydrogen content, CH, as shown in Figs. 21.21 and 21.24. Sofronis and Birnbaum [6] showed analytically that hydrogen decreases the interaction energy between dislocations. The analytical model can work both for slip localisation (local softening) and for hardening in the region larger than a grain. Ref. [41] reported variations in the cyclic yield stress, σ0.2,cyclic, and Vickers hardness, HV, as a function of the hydrogen content for type 316L. It has been reported that hydrogen charging increases static flow stress and also increases Vickers hardness with increasing hydrogen content. However, the effect of hydrogen on the cyclic yield stress in terms of hydrogen content is not simple, particularly at low hydrogen content, as for the case of fatigue life (Fig. 21.23). Moriya et al. [43] reported that hydrogen, which is cathodically charged at a certain strain during a tensile test, resulted in softening. However, it must be noted that cathodic charging always creates a hydrogen concentration gradient from a specimen surface to the subsurface, and also that internal dislocations do not trap hydrogen before hydrogen charging. Thus, if we employ cathodic charging, we cannot avoid the trigger effect of surface hydrogen on the start of slip, and also on the subsequent complicated stress
strain curve. In the experiment of Ref. [41], expecting a stronger hydrogen effect, supersaturated hydrogen was charged at high temperature and high pressure into types 304

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Figure 21.28 Macroscopic observation of slip deformation behaviour around fatigue cracks for an uncharged specimen (CH 5 2.2 mass ppm) and a specimen with excessive hydrogen content (CH 5 89.2 mass ppm) [41]. (a) Non-charged type 304 at 1.0 Hz (CH 5 2.2 mass ppm, σ 5 280 MPa, N 5 11,000 and 2a = 0.782 mm). (b) H-charged type 304 at 1.0 Hz (CH = 89.2 mass ppm, σ = 280 MPa, N = 92,000 and 2a = 1.028 mm). Plastic zone H

H

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Softened grain (d) Image of slip localization and slip-blocking effects by surrounding hardened grains near a crack tip

Figure 21.29 Schematic illustration of plastic deformation in the vicinity of a fatigue crack tips. Mechanism of hydrogen effect against HE [41]. (a) Uncharged specimen, (b) an Hcharged specimen, (c) a specimen with supersaturated hydrogen and (d) image of slip localisation and slip-blocking effects by surrounding hardened grains near a crack tip.

and 316L stainless steels, and fatigue tests were carried out at room temperature and 1 Hz. However, surprisingly, FCG rates for these specimens were dramatically decreased, as shown in Fig. 21.26. Fig. 21.28 shows the difference in slip morphologies between an uncharged specimen (CH 5 2.2mass ppm) and an excessively hydrogen-charged specimen (CH 5 89.2mass ppm). It is surprising that almost no slip bands were observed by optical microscopy in excessively hydrogen-charged specimens. The mechanisms of the difference of the slip morphologies are illustrated schematically in Fig. 21.29.

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Summarizing the experimental results of Ref. [41], the following two phenomena require attention. 1. FCG rates are decreased both at extremely low hydrogen content [19,41] and at high, supersaturated, hydrogen content, CHS. Thus, damaging hydrogen effects appear for a hydrogen content, CH, above that of the solution-treated state (2
3 wt ppm), but below the supersaturated value, CHS. In other words, undesirable effects of hydrogen appear over a certain range of CH. 2. Although we do not have precise information on dislocation pinning due to hydrogen in fcc metals, it is presumed from the experimental results of Ref. [41] that dislocation pinning due to hydrogen occurs when hydrogen atoms are trapped in a dislocation core within a dilatational stress field.

Based on the discussion with regard to Figs. 21.28 and 21.29, reconsideration of the experiments of Birnbaum et al. [23,44,45] provides another example of the same interaction effect between hydrogen and dislocations. In their experiments, hydrogen was supplied to thin specimens under a tensile stress within a TEM cell; we must note that at the beginning of a test in the TEM cell, the dislocations in a thin specimen are not initially pinned by hydrogen. The observation of increased dislocation mobility in the internal friction tests of Gavriljuk et al. [46,47] also indicates a softening effect, not by the initially stationary dislocations, but by interaction between hydrogen atoms and mobile dislocations that have been released from pinning by hydrogen. It follows that the basic mechanism of the phenomenon observed by Birnbaum et al. is essentially the same as that observed by Gavriljuk et al. Fig. 21.27b illustrates hydrogen atoms left behind an edge dislocation core after it is released from pinning and moves (Fig. 21.27a). Although these hydrogen atoms move with time toward another dislocation core within a dilatational stress field (molecular dynamic analysis by Kakimoto [48]), the dislocation mobility is enhanced as long as these dislocations are located apart from the latter dislocation (molecular dynamics analysis by Taketomi et al. [49]). Under this condition, a dislocation continues to move at a lower stress than the critical shear stress necessary to release a dislocation that has been pinned by hydrogen. Such a dislocation movement activates Frank
Read sources and generates additional new dislocations; slip is confined to planes and, in the presence of solute hydrogen, cross-slip is prevented. By contrast, in the same specimen, further slip does not originate at dislocations that have not been released from pinning by hydrogen. The chain reaction of hydrogen
dislocation interaction is presumed to be the cause of slip localisation at fatigue crack tips. Sofronis and Birnbaum [6] analysed the problems of the interaction between dislocations and hydrogen atoms distributed at the dilatation area of edge dislocations. In a similar analysis, Chateau et al. [50] also reported a decrease in interaction energy between co-planar dislocations. The HELP model, which emphasises screening stress from obstacles by hydrogen atmospheres, may be one of the important basic models to explain the interaction between hydrogen and mobile dislocations. However, FCG rates dramatically change depending on hydrogen content, for example much decreased crack growth rates for the cases of almost

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Metal Fatigue

zero H-content and supersaturated H-content and increased crack growth rates for the case of a few ppm H-content (Fig. 21.23). In order to explain, without contradiction, the experimental results of the present study (supersaturated hydrogen), and also the results of Ref. [19](almost no hydrogen), it is necessary to consider the way the HELP model is applied to fatigue phenomena in relation to hydrogen content and FCG mechanisms. If we interpret the HELP model as acting inside the fatigue process zone at crack tip but that the zone outside the process zone does not satisfy the critical condition for slip, then the phenomena can be understood without contradiction. The supersaturated hydrogen in the aforementioned mechanism influences the size of the plastic zone at a fatigue crack tip, as shown in Fig. 21.17. As explained with regard to Fig. 21.17, in the presence of hydrogen, a growing fatigue crack has a smaller plastic zone size and a smaller crack tip opening displacement than in the absence of a hydrogen effect. In the supersaturated hydrogen condition, it is natural to presume that the plastic zone size at a fatigue crack tip is strongly limited by additional constraint produced by the increased yield stress, outside the plastic zone, due to dislocation pinning by hydrogen. Consequently, slip displacement at a crack tip is limited by a zone that encircles the plastic zone. Cross-slip at the crack tip is also restricted for the same reason. Hence, planarity of slip is maintained regardless of variations in SFE. It must be noted that these mechanisms, that is slip localisation inside the plastic zone at a fatigue crack tip, are consistent with the HELP model. With respect to the hydrogen
dislocation interaction, the effect of the frequency, f, of fatigue loading is also important. Although the degree of the frequency effect naturally depends on the hydrogen content, CH, a general tendency is for an increase in FCG rate with decreasing f, but there is an upper limit to the effect of hydrogen on FCG acceleration (Figs. 21.12 and 21.18, and Ref. [41]). The reason for the existence of an upper limit is as follows. With decreasing f, hydrogen diffuses into, and concentrates within, the zone of the hydrostatic tensile stress field at a crack tip. In this diffusion process, in which hydrogen travels before being trapped by dislocations, hydrogen atoms enhance dislocation mobility, producing both slip concentration and localisation. If f is decreased to very low values, hydrogen atoms are likely to be trapped at dislocation cores and, hence, hinder dislocation motion by a pinning effect. Thus, the decrease in f causes the phenomenon of slip blocking and eventually retards FCG rate acceleration. Although essentially the same phenomenon occurs in bcc metals [20], the diffusion rate of hydrogen in bcc metals is four orders higher and, accordingly, the lower content of stable hydrogen in bcc metals, compared with fcc metals, results in a very different frequency effect. The different effects of external hydrogen and of internal hydrogen are of scientific interest and are also of practical importance. Here, external hydrogen is defined as the hydrogen that diffuses into a material, through a specimen surface, from H2 gas outside the specimen, during a fatigue test. Internal hydrogen is defined as the hydrogen that is charged into a specimen before mechanical testing. It is known that, in general, external hydrogen diffuses into a material following Sieverts’ law. It can be seen from Fig. 21.26 that when a specimen contains

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Figure 21.30 Martensitic transformation near crack tip in austenitic stainless steel type 304.

supersaturated internal hydrogen, external hydrogen due to 0.7 MPa hydrogen gas has little influence. There are two possible reasons for this result. One is that supersaturated internal hydrogen has a stronger effect of blocking the extension of the plastic zone at the crack tip than does the hydrogen-enhanced slip localisation effect caused by the external hydrogen. The other, more probable reason, is that the hydrogen content in the material is already supersaturated before fatigue testing in H2 gas at 0.7 MPa and room temperature (CHS 5 5.1 wt ppm). It is therefore thermophysically very difficult for external hydrogen to diffuse into the material. On the other hand, FCG rates for an uncharged specimen tested in 0.7 MPa H2 gas, at f 5 0.01 Hz, are much higher than those for an uncharged specimen tested in air. This is because the hydrogen content, CH, of the uncharged specimen is much lower than the hydrogen content for the supersaturated condition, and external hydrogen can enter through the new, fresh surfaces of a crack as hydrogen atoms, H. Fig. 21.30 shows EBSD observations of the strain-induced martensite produced in the vicinity of a fatigue crack, where there is severe cyclic deformation as FCG proceeds. Fig. 21.30a shows strain-induced martensite in an uncharged specimen (CH 5 2.2 wt ppm) tested in air at f 5 1 Hz. With little hydrogen effect, martensite was produced extensively, reflecting severe cyclic plastic deformation, including

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Figure 21.31 Slip deformation behaviour near fatigue cracks in type 304: (a) uncharged type 304 at 1.0 Hz (CH 5 2.2 wt ppm, σ 5 280 MPa, N 5 13,200 and 2a 5 3.019 mm), (b) H-charged type 304 at 1.0 Hz (CH 5 47.2 wt ppm, σ 5 280 MPa, N 5 18,400, and 2a 5 2.871 mm) and (c) H-charged type 304 at 1.0 Hz (CH 5 89.2 wt ppm, σ 5 280 MPa, N 5 96,200, and 2a 5 2.926 mm) [41].

cross-slip. Fig. 21.30b shows the small amount of strain-induced martensite in an uncharged specimen (CH of the internal hydrogen 5 2.2 wt ppm) tested in 0.7 MPa H2 gas at f 5 0.01 Hz. This condition definitely brought about slip localisation due to a hydrogen effect. Although in this case the martensitic transformation is limited to a small area in the vicinity of the crack, the FCG rates are higher than those of an uncharged specimen tested in air at f 5 1 Hz. Fig. 21.30c shows martensitic transformation limited to an extremely small area due to the effect of the supersaturated hydrogen (CH 5 109.3 wt ppm). This increased the flow stress outside the plastic zone and blocked the extension of the slip zone, resulting in very low FCG rates. All these observations are consistent with slip morphologies in the vicinity of cracks, as shown in Fig. 21.31. As reported in Ref. [41], the fracture surfaces of the hydrogen-supersaturated specimens are less covered with clear striations than those of the specimen with a low hydrogen content. This implies that an increase in the

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hydrogen content confines the plasticity to a smaller zone at the crack tip so that the formation of striations becomes difficult. The data introduced in this section verify the possibility of increasing fatigue strength by precharging hydrogen to a supersaturated level. Thus, the hydrogen effect is not always damaging and we can term a nondamaging hydrogen effect as ‘hydrogen effect against HE’. This phenomenon is unique to fcc austenitic stainless steels with the supersaturated hydrogen tested under cyclic loading. In bcc metals, it is difficult to keep the supersaturated hydrogen within the lattice, since the hydrogen diffusivity is high and hydrogen escapes from the material quickly. It is also difficult to perform a fatigue test using such a hydrogen-supersaturated specimen for bcc metals. Actually, the supersaturated hydrogen would not be generated in the service conditions used for components of bcc metals. Although the lattice of the austenitic stainless steels would not be supersaturated with hydrogen in service, it is possible to achieve the hydrogen supersaturation by artificial high-pressure hydrogen gas charging at a temperature higher than in service. Even after taking specimens out of an autoclave at a high pressure and at a medium temperature, the supersaturated hydrogen hardly diffuses from the austenitic stainless steels over a long period. This is due to the difference in the lattice structure between fcc and bcc and lower diffusivity in fcc than bcc by four orders of magnitude.

21.6

Hydrogen embrittlement of other materials

21.6.1 High-strength steels HE of low-strength steels, medium-strength steels and austenitic stainless steels includes common aspects such as slip localisation at crack tips and load frequency effect due to hydrogen. However, in high-strength steels other mechanisms become preferentially dominant compared to the common aspects for low-strength steels. Although most researches for high-strength steels were carried out for specimens having long cracks and the results are outside the scope of this book, learning the essential mechanisms will be useful for the problem of small cracks and defects. A transition of fatigue fracture surface morphology from transgranular (TG) to intergranular (IG) cracking is often observed for high-strength steels in the presence of hydrogen [51,52] and also in low-strength steels. Such hydrogen-induced intergranular cracking is explained in terms of the decrease in cohesion strength of the grain boundary by hydrogen [53
56]. Therefore, it is presumed that the mechanism of HE for high-strength steels with IG cracking is not identical to that for lowstrength steels. Kameda and McMahon [54] and Novak et al. [57] explained IG cracking in high-strength steels by a decohesion model based on impingement of dislocation pile-up on the carbide boundary. Liu [58] discussed a mixture of general mechanisms such as cleavage of grain boundaries and interface decohesion at grain-boundary carbides for IG cracking. It is presumed that the effects of hydrogen, cyclic frequency, time dependency, microstructures of high-strength steels, and applied stress levels on FCG rate acceleration should be mutually coupled.

600

Metal Fatigue

Figure 21.32 A model of hydrogen-induced intergranular fracture in steel [57]. Hydrogen-induced slip bands impinge a carbide at a grain boundary resulting in intergranular cracking. Precharged by 100 MPa hydrogen gas

Precharged by 100 MPa hydrogen gas R = 0.1 0.2 Hz, C = 2.04 wppm, ΔK increase

R = 0.1 0.2 Hz, C

= 2.04 wppm, ΔK increase

0.2 Hz, C

= 1.58 wppm, ΔK increase

0.2 Hz, C

0.2 Hz, C

= 1.33 wppm, ΔK decrease

0.2 Hz, C

H,R H,R H,R

2 Hz, C 2 Hz, C

H,R

= 1.51 wppm, ΔK increase

H,R

= 1.33 wppm, ΔK increase

H,R

Uncharged R = 0.1 0.2 Hz, ΔK decrease

H,R

= 1.58 wppm, ΔK increase

H,R

= 1.33 wppm, ΔK decrease

2 Hz, C

= 1.51 wppm, ΔK increase

2 Hz, C

= 1.33 wppm, ΔK increase

H,R H,R

Uncharged R = 0.1 0.2 Hz, ΔK decrease

20 Hz, C

= 1.17 wppm, ΔK increase

20 Hz, ΔK increase

20 Hz, C

= 1.17 wppm, ΔK increase

20 Hz, ΔK increase

20 Hz, C

= 1.15 wppm, ΔK decrease

20 Hz, ΔK decrease R = 0.5

20 Hz, C

= 1.15 wppm, ΔK decrease

20 Hz, ΔK decrease R = 0.5

= 1.54 wppm, ΔK decrease

20 Hz, ΔK decrease

= 1.54 wppm, ΔK decrease

20 Hz, ΔK decrease

H,R H,R

R = 0.5 20 Hz, C

H,R

H,R H,R

R = 0.5 20 Hz, C

H,R

10–3

10–3 H-charged 0.2 Hz, R = 0.1 ( , , )

–4

10–4

10

H-charged 20 Hz, R = 0.5 ( )

10–6

H-charged 20 Hz, R = 0.1 ( , )

10–7 Uncharged, R = 0.1 0.2 Hz ( ) 20 Hz ( , )

10–8

10–9

10–5

H-charged 2 Hz, R = 0.1 ( , ) da/dt (m/s)

da/dN (m/cycle)

10–5

10–6

10–7

10–8

10–9

Uncharged 20 Hz, R = 0.5 ( )

10–10

10–10 2

4

6

8 10

30

ΔK (MPa •m1/2)

(a) Relationship between da/dN and ΔK

2

4

6

8 10

30

Kmax (MPa •m1/2)

(b) Relationship between da/dt and Kmax

Figure 21.33 Effects of hydrogen and frequency on FCG rates [59].

Novak et al. [57] proposed a model of the initiation of grain boundary crack for high-strength steels in the presence of hydrogen by the model of impingement of hydrogen-induced slip bands to carbides existing along the grain boundary (Fig. 21.32). Yamabe et al. [59] investigated the frequency effect on FCG of a high-strength steel charged in 100 MPa H2 gas. Fig. 21.33 shows the high acceleration of FCG

Hydrogen embrittlement

601

Figure 21.34 FE-TEM micrographs at the cross-section of the IG fracture surface for the 1/2 H-charged specimen (f 5 20 Hz, R 5 0.1, ΔK 5 11.6 MPa m , CH,R 5 1.17 mass ppm)[59].

(a)

(b) Packet boundary Deformation twin Block boundary

(c)

Block boundary Deformation twin

Block boundary Deformation twin

Crack

Crack IG crack

Carbide

Carbide

Prior-austenite grain boundary

Figure 21.35 Schematic images of formations of (a) IG crack, (b) carbide boundary crack and (c) carbide cracking due to stress concentration by deformation twins [59].

rate observed at low frequency. If the data are plotted in terms of da/dt as shown Fig. 21.33b, then it can be understood that the crack growth rate is time-dependent. By the detailed observation of fracture surfaces, secondary grain boundary cracks ahead of main cracks were observed. Secondary cracks are classified into three types. It was made clear by TEM observation as in Fig. 21.34 that these grain boundary cracks were produced by impingement of hydrogen-induced twins onto grain boundaries. Fig. 21.35 illustrates the mechanism for three typical secondary crack initiation modes caused by deformation twins, that is IG cracks (crack A in Fig. 21.35a) [60], carbide boundary cracks (crack B in Fig. 21.35b), and carbide cracking (crack C in Fig. 21.35c). Kameda and McMahon [54] and Novak et al. [57] proposed a model for hydrogen-induced intergranular fracture in steel under static bending by assuming a dislocation pile-up impinging on a carbide at a grain boundary and leading to decohesion at a carbide
matrix interface. However, the secondary cracks observed in Ref. [59] exist outside the plastic zone ahead of the primary crack, and therefore we cannot attribute the initiation of the secondary cracks to pile-up of dislocations impinging onto carbides. It is known that the formation of deformation twins is enhanced by hydrogen in some steels and alloys [61
65], though the role of deformation twins in the initiation of cracks, especially in fatigue, has not been

602

Metal Fatigue

discussed. Ref. [59] showed experimental evidence that the decohesion of grain boundaries and carbide boundaries caused by deformation twins outside the plastic zone ahead of the primary crack can be another mechanism for IG fracture. It is natural to assume that these hydrogen-induced deformation twins cause an array of small IG cracks, as illustrated in Fig. 21.35a. In this model, subsequently, due to the stress concentration between these small IG cracks, an overall IG fracture surface is formed by coalescence of these small IG cracks. It is presumed that the formation of tear ridges at a fracture surface is a consequence of plastic deformation near a grain boundary, caused by the stress concentration between small IG cracks. These tear ridges are observed on only one side of the IG fracture surfaces due to the difference in ease of plastic deformation between two mating grains having different orientation.

21.6.2 Aluminium alloys Since there are many kinds of aluminium alloy, fatigue of aluminium alloys related to HE cannot be summarised by a simple model. Moreover, it is not clear if the HE mechanism in an environment of humidity and water is the same as that in hydrogen gas. Hydrogen content measurement in Al alloys is very difficult, even with TDS. The hydrogen content measured by a simple sample with TDS very often gives values almost 10 times higher value than the true value, because hydrogen quantity trapped by the sample surface is one order higher than the internal hydrogen [66]. If hydrogen content is measured in such a condition, erroneous conclusions can be derived in the interpretation of experimental results. This implies that the HE phenomenon in humidity or water is very complicated. Actually, most discussions of HE of Al alloys have been made without identifying the hydrogen content. When SIMS is applied to identify hydrogen in materials, so-called background hydrogen in SIMS chamber and hydrogen trapped on specimen surface mostly as the form of H2O disturbs the exact measurement of internally trapped hydrogen. Therefore, as only one example, the fatigue of A6016-T6 in high-pressure hydrogen gas for long cracks is introduced for reference in this section. A6016-T6 is a typical Al alloy used for the liner of high-pressure hydrogen CFRP container of FCV. Fig. 21.36a [67] shows the FCG rates of A6061-T6 specimens tested under different environments, 90 MPa H2 gas (purity 99.99%), 90 MPa N2 gas (purity 99.999%), air and deionized water, and test frequency at f 5 1 Hz or 0.001 Hz. For ΔK 5 15 MPa m1/2 there are no definite differences in the da/dN data under the different conditions. Fig. 21.37 [67] shows the morphologies of the fracture surfaces observed in the tests. Except for the cases of f 5 1 Hz in 90 MPa H2 gas (purity 99.99%) and 90 MPa N2 gas (purity 99.999%), clear striations were observed and the average striation spacing s was s 5 0.64 μm, which coincides with the average macroscopic crack growth rate da/dNC6 3 1027 m/cycle of Fig. 21.36b. On the other hand, the fracture morphologies in 90 MPa H2 gas and 90 MPa N2 gas showed so-called unclear striation-like patterns. Even though the environments of these

Hydrogen embrittlement

603

(a) da/dN–ΔK

Fatigue crack growth rate da/dN (m/cycle)

10−6

A6061−T6 R = 0.1

10−7

10−8

In H 2 (90 MPa, 1 Hz) In H 2 (90 MPa, 10 Hz) In H 2 (40 MPa, 1 Hz) In air (30 Hz) In deionized water (1 Hz)

−9

10

1

10

50

Stress intensity factor range ΔK (MPa࣭m1/2)

(b) Effect of test frequency f on da/dN

Fatigue crack growth rate da/dN (m/cycle)

A6061−T6

1/2

R = 0.1, ΔK = 15 MPa࣭m

10

−6

10

−7

H2 N2 Air Deionized water 90 MPa 0.1 MPa −4

10

10

−3

−2

−1

0

10 10 10 Frequency f (Hz)

10

1

10

2

Figure 21.36 (a) FCG rates of A6061-T6 specimens tested under different environments, 90 MPa and 40 MPa H2 gas (purity 99.99%), air and deionized water, and test frequency at f 5 1 Hz or 0.001 Hz. (b) Effect of test frequency for ΔK 5 15 MPa m1/2 [67].

fatigue tests were well controlled, interpretation of the different fracture surface morphologies is not easy [67]. Therefore, when we discuss the HE mechanism of various Al alloys in various environments, we need to use data in which the test environment is well controlled. Otherwise, we cannot identify the key factors which control the HE crack growth behaviour, such as intergranular cracking, transgranular cracking, precipitation elements, grain size and other microstructural factors.

604

Metal Fatigue

Figure 21.37 The morphologies of the fracture surfaces observed in the tests under different environments [67].

References [1] Y. Murakami, S. Matsuoka, Y. Kondo, S. Nishimura, Mechanism of Hydrogn Embrittlement and Guide for Fatigue Design, Yokendo, Tokyo, 2012. [2] A.R. Troiano, The role of hydrogen and other interstitials in the mechanical behavior of metals, Trans. ASM 52 (1960) 54
80. [3] R.A. Oriani, P.H. Josephic, Equilibrium aspects of hydrogen-induced cracking of steels, Acta Metall. 22 (1974) 1065
1074. [4] C.D. Beachem, A new model for hydrogen-assisted cracking (hydrogen ‘embrittlement’), Metall. Trans. 3 (1972) 437
451. [5] H.K. Birnbaum, P. Sofronis, Hydrogen-enhanced localized plasticity
a mechanism for hydrogen-related fracture, Mater. Sci. Eng. A176 (1994) 191
202. [6] P. Sofronis, H.K. Birnbaum, Mechanics of the hydrogen-dislocation-impurity interactions
I. Increasing shear modulus, J. Mech. Phys. Solids 43 (1995) 49
90. [7] I.M. Robertson, The effect of hydrogen on dislocation dynamics, Eng. Fract. Mech. 68 (2001) 671
692. [8] P. Smith, A.T. Stewart, Effect of aqueous and hydrogen environments on fatigue crack growth in 2Ni-Cr-Mo-V rotor steel, Metal Sci. 13 (1979) 429
435. [9] K.A. Esaklul, W.W. Gerberich, On the influence of internal hydrogen on fatigue thresholds of HSLA steel, Scripta Metal. 17 (1983) 1079
1082. [10] S. Suresh, G.F. Zamiski, R.O. Ritchie, Oxide-induced crack closure: an explanation for near-threshold corrosion fatigue crack growth behavior, Metal. Trans. 12A (1981) 1435
1443. [11] Gaseous hydrogen embrittlement of materials in energy technologies, in: R.P. Gangloff, B.P. Somerday (Eds.), The Problem, Its Characterization and Defects on Particular Alloy Classes, Vo. 1, Woodhead Publishing Limited, 2012. [12] R.O. Ritchie, S. Suresh, C.M. Moss, Near-threshold fatigue crack growth in 21/4Cr 1Mo pressure vessel steel in air and hydrogen, Trans. ASME, J. Eng. Mater. Tech. 102 (1980) 293
309. [13] R.O. Ritchie, Near-threshold fatigue crack propagation in steels, Int. Metals Rev. 20 (1979) 205
230.

Hydrogen embrittlement

605

[14] Y. Murakami, R.O. Ritchie, Effects of hydrogen on fatigue-crack propagation in steels, 379-416, in: R.P. Gangloff, B.P. Somerday (Eds.), Gaseous Hydrogen Embrittlement of Materials in Energy Technologies, Vol. 1: The Problem, Its Characterization and Defects on Particular Alloy Classes, Woodhead Publishing Limited, 2012. [15] N.E. Nanninga, Fatigue crack initiation and fatigue life of metals exposed to hydrogen, 347-378, in: R.P. Gangloff, B.P. Somerday (Eds.), Gaseous Hydrogen Embrittlement of Materials in Energy Technologies, Vol. 1: The Problem, Its Characterization and Defects on Particular Alloy Classes, Woodhead Publishing Limited, 2012. [16] T. Matsuo, N. Homma, S. Matsuoka, Y. Murakami, Effect of hydrogen and prestrain on tensile properties of carbon steel SGP (0.078 C-0.012 Si-0.35 Mn, mass%) for 0.1 MPa hydrogen pipelines, Trans. JSME A 74 (2008) 1164
1173. [17] H. Uyama, M. Nakashima, K. Morishige, Y. Mine, Y. Murakami, Effects of hydrogen charge on microscopic fatigue behaviour of annealed carbon steels, Fatigue Frac. Eng. Mater. Struc. 29-12 (2006) 1066
1074. [18] S. Matsuoka, N. Tsutsumi, Y. Murakami, Effects of hydrogen on fatigue crack growth and stretch zone of 0.08% C low carbon steel pipe, Trans. JSME A 74 (2008) 1528
1537. [19] T. Kanezaki, Effects of Hydrogen on the Strength Properties of Stainless Steels (PhD Dissertation), Kyushu University, 2007. And also Ref. [21]. [20] H. Tanaka, N. Honma, S. Matsuoka, Y. Murakami, Effect of hydrogen and frequency on fatigue behavior of SCM435 Steel for storage cylinder of hydrogen station, Trans. JSME A 73 (2007) 1358
1365. [21] Y. Murakami, T. Kanezaki, Y. Mine, S. Matsuoka, Hydrogen embrittlement mechanism in fatigue of austenitic stainless steels, Metall. Mater. Trans. A 39 (2008) 1327
1339. And also Ref. [19]. [22] T. Kanezaki, C. Narazaki, Y. Mine, S. Matsuoka, Y. Murakami, Effects of hydrogen on fatigue crack growth behavior of austenitic stainless steels, Int. J. Hydrogen Energy 33 (2008) 2604
2619. [23] T. Tabata, H.K. Birnbaum, Direct observations of hydrogen enhanced crack propagation in iron, Scr. Metall. 18 (1984) 231
236. [24] G.M. Bond, I.M. Robertson, H.K. Birnbaum, The influence of hydrogen on deformation and fracture processes in highs trength aluminum alloys, Acta Metall. 35 (1987) 2289
2296. [25] D.S. Shih, I.M. Robertson, H.K. Birnbaum, Hydrogen embrittlement of a titanium: in situ TEM studies, Acta Metall. 36 (1988) 111
124. [26] P. Rozenak, D. Eliezer, Effects of metallurgical variables on hydrogen embrittlement in AISI type 316, 321 and 347 stainless steels, Mater. Sci. Eng. 61 (1983) 31
41. [27] H.K. Birnbaum, P. Sofronis, Hydrogen-enhanced localized plasticity—a mechanism for hydrogen related fracture, Mater. Sci. Eng. A176 (1994) 191
202. [28] P.J. Ferreira, I.M. Robertson, H.K. Birnbaum, Hydrogen effects on the interaction between dislocations, Acta Mater. 46 (1998) 1749
1757. [29] Y. Katz, W.W. Gerberich, Fatigue transients and hydrogen induced cracking in AISI 304 metastable austenitic stainless steel, in: N.R. Moody, A.W. Thompson (Eds.), Proceedings of the 4th International Conference on the Effect of Hydrogen on the Behavior of Materials: Hydrogen Effects on Material Behavior, TMS, Warrendale, PA, 1990, pp. 469
479. [30] Y. Fukai, The Metal-Hydrogen System: Basic Bulk Properties, Springer, 1993. [31] Y. Murakami, Effects of hydrogen on metal fatigue, in: Proceedings of the International Hydrogen Energy Development Forum, Fukuoka, Japan, 2007, pp. 96
105.

606

[32] [33] [34] [35] [36]

[37] [38] [39] [40] [41] [42] [43] [44] [45] [46] [47] [48] [49] [50] [51] [52]

[53]

[54] [55] [56] [57]

[58] [59]

Metal Fatigue

P. Sofronis, Y. Liang, N. Aravas, Eur. J. Mech. A/Solids 20 (2001) 857. C. Laird, ASTM STP 415 (1967) 131. R.M.N. Pelloux, Trans. ASM 62 (1969) 281. P. Neumann, Acta Metall. 22 (1974) 1155. C.H. Bichler, R. Pippan, Direct observation of the formation of striation, in: J.H. Beynon, M.W. Brown, T.C. Lindley, B. Tomkins (Eds.), Engineering Against Fatigue: Proc Int Conf on Fatigue Sheffield, March 1997, Balkema, Rotterdam, 1999, pp. 211
218. P. Sofronis, R.M. McMeeking, J. Mech. Phys. Solids 37 (1989) 317. Y. Liang, P. Sofronis, N. Aravas, Acya Mater. 51 (2003) 2727. D.C. Ahn, P. Sofronis, R.H. Dodds, Jr., Int. J. Hydrogen Energy (available on line). Y. Murakami, T. Kanezaki, Y. Fukushima, H. Tanaka, J. Tomuro, K. Kuboyama, M. Matsue, Y. Ito, H. Ando, Trans. Japan Soc. Mech. Eng. A 75 (93) (2009) 92
102. Y. Murakami, T. Kanezaki, Y. Mine, Hydrogen effect against hydrogen embrittlement, Metall. Mater. Trans. A 41 (2010) 2548
2562. R. Kirchheim, Solid solution softening and hardening by mobile solute atoms with special focus on hydrogen, Scripta Mater. 67 (2012) 767
770. S. Moriya, H. Matsui, H. Kimura, Mater. Sci. Eng. 40 (1979) 217. I.M. Robertson, H.K. Birnbaum, Acta Metall. 34 (1986) 353
366. P. Rozenak, I.M. Robertson, H.K. Birnbaum, Acta Metall. Mater. 38 (1990) 2031
2040. V.G. Gavriljuk, V.N. Shivanyuk, J. Foct, Acta Mater. 51 (2003) 1293. V.G. Gavriljuk, V.N. Shivanyuk, B.D. Shanina, Acta Mater. 53 (2005) 5017. K. Kakimoto, Molecular dynamic analysis of hydrogen trap by dislocation, the paper presented at an informal meeting at Kyushu University, 2008. S. Taketomi, R. Matsumoto, N. Miyazaki, J. Mater. Sci. 43 (2008) 1166. J.P. Chateau, D. Delafosse, T. Magnin, Acta Mater. 50 (2002) 1507. J. Toplosky, R.O. Ritchie, On the influence of gaseous hydrogen in decelerating fatigue crack growth rates in ultrahigh strength steels, Scr. Metall. 15 (1981) 905
908. Y. Kondo, Y.,M. Kubota, K. Shimada, Hydrogen enhanced crack propagation of SCM440H low-alloy steel under long-term varying load, Eng. Fract. Mech. 77 (2010) 1963
1974. K. Yoshino, C.J. McMahon Jr, The cooperative relation between temper embrittlement and hydrogen embrittlement in a high strength steel, Metall. Trans. 5B (1974) 363
370. J. Kameda, C.J. McMahon Jr, Solute segregation and brittle fracture in an alloy steel, Metall. Trans. 11A (1980) 91
101. J. Kameda, C.J. McMahon Jr., Soltue segregation and hydrogen-induced intergranular fracture in an alloy steel, Metall. Trans. 14A (1983) 903
911. C.J. McMahon Jr., Hydrogen-induced intergranular fracture of steels, Eng. Fract. Mech. 68 (2001) 773
788. P. Novak, R. Yuan, B.P. Somerday, P. Sofronis, R.O. Ritchie, A statistical, physicalbased, micro-mechanical model of hydrogen-induced intergranular fracture in steel, J. Mech. Phys. Solids 58 (2010) 206
226. H.W. Liu, Hydrogen assisted intergranular cracking in steels, Eng. Fract. Mech. 78 (2011) 2563
2571. J. Yamabe, T. Matsumoto, S. Matsuoka, Y. Murakami, A new mechanism in hydrogenenhanced fatigue crack growth behavior of a 1900-MPa-class high-strength steel, Int. J. Fract. 177 (2012) 141
162.

Hydrogen embrittlement

607

[60] A. Gilbert, A.,G.T. Hahn, C.N. Reid, B.A. Wilcox, Twin-induced grain boundary cracking in b.c.c. metals, Acta Metall. 12 (1964) 754
755. [61] J.M. Rigsbee, R.B. Benson, A TEM investigation of hydrogen-induced deformation twinning and associated martensitic phases in 304-type stainless steel, J. Mater. Sci. Lett. 12 (1977) 406
409. [62] C. Hwang, I.M. Bernstein, Hydrogen induced slip and twinning in iron alloys, Scr. Metall. 16 (1982) 85
90. [63] K. Jagannadham, R.W. Armstrong, J.P. Hirth, Deformation twinning in high-hydrogensolubility refractory alloy crystals, Phil. Mag. 68A (1993) 419
451. [64] T.D. Le, I.M. Bernstein, S. Mahajan, Effects of hydrogen on micro-twinning in a FeTi-C alloy, Acta Metall. Mater. 41 (1993) 3363
3379. [65] E.G. Astafurova, G.G. Zakharova, H.J. Maier, Hydrogen-induced twinning in ,111. Hadfield steel single cystals, Scr. Mater. 63 (2010) 1189
1192. [66] J. Yamabe, T. Awane, Y. Murakami, Hydrogen trapped at intermetallic particles in aluminum alloy 6061-T6 exposed to high-pressure hydrogen gas and the reason for high resistance against hydrogen embrittlement, Int. J. Hydrogen Energy 42 (2017) 24560
24568. [67] H. Itoga, S. Watanabe, Y. Fukushima, S. Matsuoka, Y. Murakami, Fatigue crack growth of aluminum alloy A6061-T6 in high pressure hydrogen gas and failure analysis on 35 MPa compressed hydrogen tanks VH3 for fuel cell vehicles, Trans. JSME (A) 78 (2012) 442
457.