DOI 10.1007/s11041-015-9861-y Metal Science and Heat Treatment, Vol. 57, Nos. 3 – 4, July, 2015 (Russian Original Nos. 3 – 4, March – April, 2015)
UDC 669.788:620.178.2:669.018.298
EFFECT OF VARIOUS FACTORS ON HYDROGEN EMBRITTLEMENT OF STRUCTURAL STEELS V. G. Khanzhin,1 V. Yu. Turilina,1 S. O. Rogachev,1 A. V. Nikitin,1 and V. A. Belov1 Translated from Metallovedenie i Termicheskaya Obrabotka Metallov, No. 4, pp. 19 – 26, April, 2015.
Results of studies of hydrogen embrittlement of structural steels of different strength are presented. The effect of various factors on delayed hydrogen fracture is analyzed using the results of investigations by the methods of acoustic emission, metallography and fractography of bolts under the conditions of tension with bending after galvanic hydrogen charging.
Key words: hydrogen embrittlement, galvanic hydrogen charging, acoustic emission, delayed fracture.
The grain boundary fracture due to hydrogen may by intensified in the presence of segregations of fine dispersed particles over boundaries [6, 7]. All the three successive stages of DF, i.e., nucleation of a crack, its slow (steady) growth, and then accelerated (unsteady) growth, may be controlled by hydrogen as a result of competition between the “chemical” damage due to the interaction between the hydrogen and the metal and the “mechanical” damage due to the action of applied stress on the components of the structure of the metal. The processes of hydrogen embrittlement of structural heat-hardenable steels of different strength and the possibilities of its elimination have been considered in [8 – 11]. To obtain objective experimental data on the mechanisms and kinetics of hydrogen embrittlement and on effective control of the resistance of metals to this kind of fracture we should resort to complex methods of study of the state of the material and to results of evaluation of the action of various factors on the hydrogen embrittlement of steels and on the delayed fracture of articles from them. The present work is devoted to these topics.
INTRODUCTION Solution of the problem of hydrogen embrittlement of steels and of delayed fracture (DF) of steel parts, especially of critical ones, related to the former remains an important task of the recent science of materials. The susceptibility of steels to hydrogen embrittlement depends on many, often interconnected, factors (the chemical composition, the structure, the strength level, the internal stresses), which determine the capacity of the metal to absorb hydrogen and the mechanism and kinetics of its delayed fracture [1 – 7]. Atomic hydrogen can be present in structural steels due to electrolysis of water in the environment of dissolved industrial gases. If the dissolved atomic hydrogen cannot leave the metal, it may be accumulated in structural traps but now in a molecular (gaseous) form and create high mechanical stresses in the volume surrounding the trap due to its pressure. Traps for atomic hydrogen in a steel may be inclusions, particles of second phases, and regions of high tensile stresses on phase boundaries, where the concentration of hydrogen is elevated due to ascending diffusion. Formation, growth and merging of pores with hydrogen constitute a typical mechanism of hydrogen embrittlement in steels with low and medium strength and relatively low resistance to plastic deformation. In high-strength steels with high resistance to deformations voids with hydrogen often give rise to sharp microcracks, which causes intragrain and intergrain brittle fracture. 1
METHODS OF STUDY We studied structural heat-hardenable steels of type 35KhGM (states 1 and 2 ) and 33KhMF (state 3 ), the chemical compositions and the mechanical properties of which are presented in Table 1. To test the steels we cut cylindrical specimens 5 mm in diameter with functional part 100 mm long. Different strengths were obtained in the steels by oil
NITU “MISiS,” Moscow, Russia (e-mail:
[email protected]).
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1 4
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Fig. 1. Scheme of testing of a flat specimen by the method of fourpoint bending with galvanic charging with hydrogen and measurement of the acoustic emission: 1 ) flat specimen; 2 ) round rests of the loading device; 3 ) electrolytic cell with platinum anode; 4 ) electrodes; 5 ) acoustic emission sensor.
quenching and tempering at different temperatures, namely, at 550°C for 1 h (states 1 and 3 ) and at 400°C for 1 h (state 2 ). In all the cases the cooling was conducted in oil. The structure of the steels in states 1 and 3 was represented by tempered troostite; in state 2 it was represented by troostite-martensite. The structure of steel 33KhM (state 3 ) contained a great number of particles of carbides and carbonitrides. The size of the austenite grains in all the cases studied was 12 – 14 mm. The difference in the strength and ductility of steel 35KhGM in states 1 and 2 was determined primarily by the strength of the matrix rather than by the precipitated carbide and carbonitride particles; the density of the distribution of the latter for states 1 and 2 did not differ substantially and amounted to (2.8 ± 1.1) ´ 108 and (2.7 ± 0.5) ´ 108 mm – 2, respectively. For steel 33KhMF (state 3 ), on the contrary, a high density was provided by a great number of fine particles in the structure, the density of which was an order of magnitude higher (1.0 ± 0.5) ´ 109 mm – 2 than in steel 35KhGM in states 1 and 2 [10]. The specimens were charged with hydrogen by a specially developed method that involved loading by extension to a stress s = 0.8s0.2 and local (in an electrochemical cell) galvanic hydrogen charging of the specimens for 4 – 168 h at a specified load [8]. The hydrogen-charged specimens were
TABLE 1. Mechanical Properties of Steels of Types 35KhGM and 33KhMF under Tension Steel
State
HV, kgf/mm2
s0.2 , MPa
sr , MPa
eu , %
d, %
y, %
35KhGM
1
350 ± 2
940
1050
8.3
11.5
63
35KhGM
2
1460
4.2
7.8
54
33KhMF
3
500 ± 2 1360 510 ± 2 1440
1550
4.4
8.1
53
Notations: eu is the uniform strain; d is the elongation; y is the contraction.
5
Fig. 2. Scheme of testing of bolts for delayed fracture under tension with bending and simultaneous galvanic charging with hydrogen: 1 ) tubes for feeding the electrolytic solution; 2 ) loading nut; 3 ) bolt; 4 ) electrodes; 5 ) air-tight electrochemical cell; 6 ) acoustic emission sensor.
loaded to s = 0.2s0.2 and then extended until failure at a speed of 6 ´ 10 – 4 sec – 1 at 20°C in an Instron LX150 testing machine with simultaneous recording of the acoustic emission (AE) [8]. Thus, we analyzed the hydrogen-induced fracture under the conditions of uniform uniaxial loading. Delayed fracture under galvanic hydrogen charging under the conditions of bending was studied by loading of flat specimens in a specially designed device for creating fourpoint bending [11] (Fig. 1). The bending gave rise to tensile stresses on the external (upper) surface of the specimen and compressive stresses on the opposite (internal) surface. In this way we created a nonuniform stress-state state over the section of the flat specimen. Delayed hydrogen fracture was studied for hexagonal bolts M9 with 1.0-mm thread pitch produced from a steel of type 35KhGM (state 1 ) and from a stronger steel of type 33KhMF (state 3 ). A complex stress state was created in the specimen by deforming it through a diagonally cut nut in a special device [11] (Fig. 2). Local galvanic charging with hydrogen was performed in an electrochemical cell mounted tightly on the bolt (Fig. 2). The resistance of the bolt to delayed fracture under hydrogen charging was evaluated from the results of long-term (for up to 700 h) tests. The kinetics of the delayed fracture of the hydrogen-charged bolt under a load was detected using the method of acoustic emission described in [8, 10]. The studied states of the steels were also compared with respect to their capacity to absorb hydrogen by the method of [9]. The solubility of hydrogen in the metal was evaluated in terms of the volume of the gas emitted from the specimen after hydrogen charging in an electrochemical cell for 15 h. We observed real-time emission of hydrogen bubbles through a layer of glycerin with simultaneous video shooting of the process. The volume of the emitted hydrogen was determined automatically from the total volume of bubbles when we analyzed the images with the help of a specially designed computer software [9]. In accordance with the method of
Effect of Various Factors on Hydrogen Embrittlement of Structural Steels
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Fig. 3. Stress-strain diagrams superimposed on AE diagrams of steels: a, b ) 35KhGM in state 1; c, d ) 35KhGM is sate 2; e, f ) steel 33KhGM is state 3; a, c, e) prior to charging with hydrogen; b, d, f ) after charging with hydrogen for 50 h.
500
500
30
20
20 10 0
[10] we observed the distribution of hydrogen microbubbles in the structure of the steels. RESULTS AND DISCUSSION Hydrogen-Induced Fracture under Tension To evaluate the ductility of the steels after hydrogen charging for different periods of time the specimens were subjected to tensile deformation until failure. The obtained typical stress-strain diagrams of the hydrogen-charged specimens were superimposed on the diagrams of acoustic emission (Fig. 3). The stress-strain diagram and the AE data were used to analyze the kinetics of fracture under tension for
0.04
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0
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e
specimens without hydrogen (Fig. 3a, c, and e) and after charging with hydrogen (Fig. 3b, d, and f ). For all the states the maximum load corresponded to the stop of uniform deformation and localization of strain in the neck. The earlier localization of strain in the necks of the specimens of steel 33KhMF without hydrogen in the stronger state 3 was caused by formation of microcracks in the stage of uniform deformation detected from the AE pulses with a high amplitude of 10 – 30 dB (Fig. 3e ). After the extension of hydrogen-charged specimens in all the states studied high-amplitude AE signals from surface cracks were detected even in the elastic range of the stress– strain diagram.
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The specimens of 35KhGM in state 1 before and after 50-h hydrogen charging had the same fracture kinetics. The geometry of the fractures was also the same, i.e., had an “asterisk” pattern. Only after hydrogen charging for 168 h under a load the uniform strain decreased noticeably (by 1.5%) and the contraction in the neck y = 46 – 50% (without the action of hydrogen y = 60 – 63%). Long-term charging with hydrogen caused “layering” in a part of tensile specimens, when large and flat axial cracks separated the specimen into 2 – 3 sectors. The short period of crack nucleation under hydrogen charging at a load and the growth of brittle cracks over grain boundaries decreased markedly the ductility margin of steel 35KhGM with higher strength (state 2 ). All the extended specimens broke by cleavage. For example, 5-h galvanic charging with hydrogen produced uniform strain eu » 1 – 1.5% in a specimen with surface creaks. AE signals with an amplitude of 10 – 55 dB accompanied the growth of a surface crack (Fig. 3d ) from the start of the tension to a critical length amounting to 0.2 – 0.3 mm. After 8 – 15 h of hydrogen charging the specimens of 35KhGM in state 2 failed by cleavage upon tension already in the elastic range. The measured volumes of emitted hydrogen showed that its solubility in steel 35KhGM in state 2 was 40 times higher than in state 1 at the same conditions of hydrogen charging and measurement times. However, the dissolved hydrogen may not affect noticeably the fracture resistance of steels. Under uniaxial extension of specimens locally charged with hydrogen at a constant load fracture starts similarly whatever the level of their strength, i.e., surface cracks nucleate in the hydrogen-charged zone (Fig. 4). A neck forms in the same zone. In steel 35KhGM with lower strength (state 1 ) fracture starts by detachment of the “bottom” in the long neck; this occurs either in the zone of maximum tensile stresses on the axis of the specimen or at the tip of the deepest surface crack after long-term (over 100 h) charging with hydrogen. The “bottom” crack opens by merging of pores under lumped elongation in the neck and creates a ductile relief, which is the coarser the longer the time of the preliminary charging of the specimen. Lumped strain also creates grain elongation and tangential stresses required for opening of longitudinal cracks in the neck upon extension of hydro-
b
Fig. 4. Cracks in steel 33KhMF (state 3 ) after charging with hydrogen for 22 h under load: a) on the surface of the specimen in the zone of hydrogen charging; b ) on a longitudinal lap.
gen-charged specimens of steel 35KhGM in plastic state 1. Growth in the hydrogen concentration in the steels changes the fracture pattern from a “symmetric” asterisk to an “asymmetric” one, and then to brittle “separation” [9]. In steels 35KhGM and 33KhMF with higher strength (states 2 and 3 ) with hydrogen the fracture upon tension is controlled fully by the surface hydrogen cracks. If the length of the crack induced by hydrogen charging under a load is less than the critical one, the chain of the events under the tension includes growth of the surface grain-boundary cracks already under elastic deformation of the specimen, shortterm uniform deformation, loss of the flow stability and a short lumped deformation, and breakage of the specimen from the “critical” crack with deformation of the fracture surface in the form of flat-bottom dimples. The kinetics of delayed fracture under a load in hydrogen charging is characterized in the first turn by the incubation period (the time to crack nucleation on the surface of the specimen). The incubation period in steel 35KhGM in state 1 was 8 – 10 h from the start of the test, when the AE signals reflected formation of surface cracks 0.06 – 0.08 mm deep. Their subsequent slow growth (at a speed of 5.2 ´ 10 – 7 mm/sec) did not cause failure of the specimens even after the maximum (168 h) time of hydrogen charging under load. The incubation period for steel 33KhMF with higher strength (state 3 ) was only 35 – 50 min. The specimens of steel 33KhMF with the same content of hydrogen as in steel 35KhGM in state 1 fractured in the brittle manner in 18 – 24 h (Fig. 5). Independently of the strength level and of the presence of particles of second phases in the structure delayed hydrogen embrittlement of the steels started similarly, i.e., with nucleation of surface cracks in the zone of local hydrogen charging at a constant load. Under subsequent tensile deformation of the cylindrical hydrogen-charged specimens a neck formed in the same zone. The ductility of the less strong steel 35KhGM in state 1 decreased noticeably only upon hydrogen charging at a load (for over 100 h). In this case the fracture geometry in the specimen also changed from a deep surface crack. In the stronger steel 33KhMF (state 3 ) the rate of growth of the
Effect of Various Factors on Hydrogen Embrittlement of Structural Steels SVp , conv. units 2000
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surface crack was an order of magnitude higher. The much higher (by factor of 27) capacity of steel 33KhMF to saturation with hydrogen and the lower resistance to delayed hydrogen embrittlement than in steel 35KhGM in state 1 were connected with the higher (by an order of magnitude) density of particles of second phases in the structure of steel 33KhMF. Delayed Fracture under Bending Nonuniform distribution of stresses under delayed fracture creates local stresses and, as a consequence, the conditions for steady delayed crack growth. On the contrary, in the state with uniform strain distribution the susceptibility to delayed fracture may decrease or disappear with time [12, 13]. The effect of hydrogen on the plastic deformation of steels is explainable by its high mobility. When the hydrogen is distributed nonuniformly over the cross section, the concentration expansion creates macroscopic internal stresses. Being redistributed to equilibrium in the stress field, hydrogen accumulates in the zone of tensile stresses, for example, at the tip of a crack. The concentration stresses are added to the residual stresses in the structure upon saturation of the metal with hydrogen and reach the yield point in local regions; the flow is then localized in narrow shear bands creating a wavy macrobrittle surface under delayed hydrogen fracture (DHF). A typical diagram of peak AE amplitudes (Vp ) and a curve of added AE amplitudes (SVp ) in the tests for DHF performed in the mode of bending of flat specimens are presented in Fig. 5.
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In the process of hydrogen charging under pure bending of flat specimens, just like in extension of round specimens, the fracture under DHF starts similarly in all the steels with formation of microcracks at moment t1 detected by the acoustic emission (Fig. 5). Analyzing the fracture surface we may conclude that the cracks 10 – 100 mm in size in steel 35KhGM (state 1 ) have nucleated both on the surface of the specimen in the hydrogen-charged zone and at a distance h ~ 250 – 300 mm from the surface. Such nuclei cracks are absent under the surface of the specimen in the case of fracture of the stronger steel 33KhMF (state 3 ). The difference in the resistance to nucleation of hydrogen cracks under bending determines the difference in the incubation periods of delayed fracture. According to the AE data, the incubation period of DHF in bending tests of specimens of steel 35KhGM in states 1 and 2 is t1 ~ 60 – 120 min and ~ 15 – 40 min, respectively. After a 10-h hold under load with hydrogen charging the specimens were unloaded and broken. The morphology of the hydrogen macrocracks was used to estimate the difference in the resistance of steel 35KhGM in state 1 and steel 33KhMF in state 3 to hydrogen embrittlement under bending. In the stronger steel 35KhGM (state 3 ) the surface bore a straight crack 5 – 6 mm long, which corresponded to the zone of hydrogen charging of the meal in an electrochemical cell; the fracture surface due to breaking of the specimen was plane. On the surface of flat specimens of plastic steel 35KhGM (state 1 ) after the test we observed short (2 – 3 mm) curved cracks; the considerable plastic strain of the section of the specimen not damaged by DHF created an oblique fracture after the breakage. Analysis of panoramic fracture images of flat surfaces has shown three zones of hydrogen crack propagation typical for delayed fracture, namely, (1 ) a zone of crack nucleation, (2 ) a zone of its steady growth, and (3 ) a zone of final breakage. The rate of nucleation of a hydrogen crack v1 was evaluated in terms of the ratio of the depth of the zone of nucleation a1 to the time of crack nucleation Dt1 = t2 – t1 (Fig. 5a ). At Dt1 ~ 300 – 400 min the rate of nucleation in steel 35KhGM in state 1 was v1 ~ (1.3 – 1.7) ´ 10 – 5 mm/sec. In same steel in state 2 the rate of nucleation of a hydrogen crack under DHF was of the same order of magnitude, i.e., v1 ~ (1.2 – 1.5) ´ 10 – 5 mm/sec (Dt1 ~ 160 – 200 min). In steels 35KhGM (state 1 ) and 33KhMF (state 3 ) the average stress intensity factors active at the front of a semi-elliptical hydrogen crack upon its transition to steady growth after the nucleation did not differ significantly, i.e., KI = 16.0 ± 0.7 MPa × m1/2 (eu = 750 MPa) for a crack in steel 1 and KI = 17.0 ± 0.6 MPa × m1/2 (eu = 1150 MPa) for a crack in steel 3 [10]. Transition to a steady crack growth under DHF (t2 , Fig. 5a ) is accompanied by growth of the amplitudes (Vp ~ 20 – 40 dB, Fig. 5b ) and intensity of the AE. The dominantly intergrain and intragrain mechanism of fracture of the
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metal under DHF is preserved in the stage of steady growth of the crack, while the fraction of ductile bridges between brittle facets increases. When a specimen is tested for t0 = 600 min, the time of steady growth of the crack (Dt2 = t0 – t2 ) amounts to Dt2 ~ 350 – 400 min for steel 35KhGM (state 1 ) and to Dt2 ~ 120 – 150 min for steel 33KhMF (state 3 ). The rate of steady growth of a hydrogen crack is much higher than in the stage of its nucleation. For a “narrow” crack in steel 35KhGM (state 1 ) this rate is higher than in steel 33KhMF, i.e. (8.6 – 16) ´ 10 – 5 mm/sec and (4.9 – 5.4) ´ 10 – 5 mm/sec, respectively. Thus, the time of formation of a surface crack in steel 1 under bending of a hydrogen-charged flat specimen is twice higher than in steel 2. Delayed Fracture of Bolts Hydrogen brittleness is a very important cause of early fracture of parts from high-strength steels. The greatest dan-
600
80 700 t, h
Fig. 6. Kinetic curves of AE and typical flaws on the surface of bolts from steels 33KhMF in state 3 (a) and 35KhGM in state 1 (b ): Vp ) peak AE amplitudes; SVp ) sum of AE amplitudes.
ger of hydrogen charging is the development of susceptibility to delayed fracture (DF) in parts, especially those containing stress concentrators. This often happens in operation of bolts and studs serving under conditions of tension with bending, because the nonuniformity of stresses over the cross section of a part produces a nonuniform concentration of hydrogen [6]. Fracture starts with formation of shallow surface cracks in the region of tensile stresses. A change in the hydrogen content in the metal changes the duration of the incubation period of DF, i.e., the time before the appearance of the first crack (from several minutes to several hours for highstrength steels). AE diagrams and typical flaws in testing of bolts are presented in Fig. 6. After 700 h of testing grain boundary microcracks (Fig. 7a) and circular hydrogen cracks (Fig. 7b and c) are observable on the surface of bolts from the stronger steel 33KhMF (state 3) in the region of local hydrogen charging. Accumulation of damage in the metal of the bolts in the zones of hydrogen charging is accompanied by AE for
Effect of Various Factors on Hydrogen Embrittlement of Structural Steels
à
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b
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c
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Fig. 7. Cracks on the surface (a, b ) and in the body (c) of a bolt from steel 33KhMF in state 3 after 200-h (a) and 700-h (b, c) galvanic charging with hydrogen.
both steels. For example, for steel 35KhGM in states 1 and 2 the bolts started to fracture under hydrogen charging at a load after a short incubation period, i.e., 30 – 40 h and 50 h, respectively. For the bolts from steel 33KhMF (state 3 ) active accumulation of damage occurred after testing for 50 – 300 h (Fig. 6a ). Analysis of the fracture surfaces has shown formation of first single surface grain-boundary microcracks and then of their nets (Fig. 7a ). After 300 h of testing the surface of the bolt exhibited chains of shallow round sources of fracture on process marks, i.e., the metal was eroded. Shallow (up to 200 – 500 mm) circular cracks formed on these sources after 600 – 700 h of testing. The kinetics of hydrogen embrittlement of the bolts from steel 1 was different; the kinetic curves had no long regions with enhanced activity of AE. The intensity and the level of the peak AE amplitudes was considerably lower (Fig. 6b ). The surface of the bolts contained usual sources of local corrosion, i.e., chains of round regions of damage (erosion of the metal), pits or short and wide flaws. Through cracks causing failure of bolts were not detected even after 700 h of hydrogen charging. CONCLUSIONS The main factor determining the difference in the mechanism and kinetics of the development of delayed hydrogen embrittlement and in the fracture resistance of steels is the level of their strength and toughness, which affects both the stage of nucleation of hydrogen cracks, the possibility of their propagation, and the kinetics of growth to a critical size. The effect of particles in the structure of the steels on the processes of hydrogen embrittlement is only an additional factor. Bending of the specimen during charging with hydrogen creates nonuniform distribution of stresses over its cross section and accelerates considerably the process of delayed hy-
drogen fracture (DHF) of the metal. In steel 35KhGM of medium strength (state 1 ) DHF under bending starts after a comparatively short incubation period (5 – 10 times shorter than the DHF under tension). In a stronger steel 33KhMF (state 3 ) the incubation period is only 15 – 40 min (in state 1 the incubation period is 60 – 120 min). At an insignificant difference in the stress intensity factors the transition to steady fracture at the front of hydrogen cracks in the tough steel 35KHGM (state 1 ) under bending starts much later than in steel 33KhMF in state 3. Comparative analysis of the resistance of bolts from steels of different strength to hydrogen embrittlement under the conditions of tension with bending and galvanic hydrogen charging has made it possible to determine the kinetics and the mechanism of hydrogen fracture in the metal right in the article, when the state of the surface of the article becomes a very important factor changing the conditions of dissociation of hydrogen and the diffusion of its atoms into the metal. Growth in the strength of the material changes the duration of the incubation period and the intensity of damage accumulation (from formation of many surface cracks to growth of a through crack). The probability of accelerated (catastrophic) fracture is not high. REFERENCES 1. L. S. Moroz and B. B. Chechulin, Hydrogen Brittleness of Metals [in Russian], Metallurgiya, Moscow (1967), 255 p. 2. S. M. Beloglazov, Hydrogen Charging of Steels in Electrical Processes [in Russian], Izd. LGU, Leningrad (1975), 412 p. 3. W. W. Gerberich and Y. T. Chen, “Hydrogen-controlled cracking — an approach to threshold stress intensity,” Metall. Trans., A13(2), 305 – 311 (1975). 4. J. P. Hirth and H. H. Johnson, “Hydrogen problems in energy related technology,” Corrosion, 32(1), 3 – 25 (1976).
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5. R. V. Hertzberg, Deformation and Fracture Mechanics of Structural Materials [in Russian], Metallurgiya, Moscow (1989), 576 p. 6. M. A. Shtremel, Strength of Alloys, Part II, Deformation [in Russian], Izd. Dom MISiS, Moscow (1997), 527 p. 7. G. N. Kasatkin, Hydrogen in Structural Steels [in Russian], Intermet Engineering, Moscow (2003), 336 p. 8. V. G. Khanzhin, S. A. Nikulin, V. Yu. Turilina, et al., “Hydrogen embrittlement of steels. I. Analysis of the kinetics of hydrogen embrittlement by measuring acoustic emission,” Deform. Razrush. Mater., No. 8, 35 – 41 (2011). 9. V. G. Khanzhin, S. A. Nikulin, V. A. Belov, et al., “Hydrogen embrittlement of steels. II, Effect of strength,” Deform. Razrush. Mater., No. 1, 40 – 47 (2012).
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10. V. G. Khanzhin, S. A. Nikulin, V. A. Belov, et al., “Hydrogen embrittlement of steels. III, Effect of particles,” Deform. Razrush. Mater., No. 3, 34 – 40 (2012). 11. V. G. Khanzhin, S. A. Nikulin, O. V. Khanzhin, et al., “Hydrogen embrittlement of steels. IV, Delayed fracture upon bending,” Deform. Razrush. Mater., No. 4, 45 – 48 (2012). 12. G. A. Filipov and V. I. Sarrak, “Local distribution of hydrogen and internal microstresses in the structure of hardened steel,” Fiz. Met. Metalloved., 49(1), 121 – 125 (1980). 13. M. Garet, A. M. Brass, C. Haut, and F. Gutterez-Solana, “Hydrogen trapping on non metallic inclusions in Cr – Mo low alloy steels,” Corros. Sci., 39(6), 1073 – 1086 (1978).
