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Fatigue damage and its interaction with hydrogen in martensitic steels
Materials Science and Engineering A348 (2003) 192 /200 www.elsevier.com/locate/msea
Fatigue damage and its interaction with hydrogen in martensitic steels M. Nagumo a,b,*, H. Shimura a, T. Chaya a, H. Hayashi c, I. Ochiai c a
Department of Materials Science and Engineering, Waseda University, Okubo 3-4-1, Shinjuku, Tokyo 169-8555, Japan b Laboratory for Materials Science and Technology, Waseda University, Tokyo, Japan c R&D Section, Suzuki Metal Co., Ltd, Higashi-Narashino 7-5-1, Narashino, Chiba 275-8577, Japan Received 19 June 2002; received in revised form 19 September 2002
Abstract Fatigue damage involving creation of point defects has been revealed by means of hydrogen thermal desorption analysis, utilizing hydrogen as a probe of defects. Two types of high-strength martensitic steels were subjected to a rotational bending test for various cycles. The amount of absorbed hydrogen decreased or was almost constant in the early stage of the fatigue process and then began to increase after a substantial number of fatigue cycles. The increase in the final stage was due to the creation of point defects, presumably vacancies, while the decrease in the early stage was ascribed to changes in microstructural constitutions. Hydrogenprecharging drastically reduced the fatigue life, but the effect of hydrogen on the fatigue limit was not significant. The fracture surface of hydrogen-precharged specimens was smooth without cleavage-like crystallographic features. The defect density in fractured specimens in the presence of hydrogen was higher than in specimens fatigued for a similar number of cycles without hydrogen. Involvement of hydrogen in fatigue damage in the crack nucleation and growth stages is discussed. # 2002 Elsevier Science B.V. All rights reserved. Keywords: Martensitic steel; Fatigue; Fatigue damage; Vacancy; Hydrogen thermal desorption; Hydrogen embrittlement
1. Introduction The fatigue process initiates with irreversible changes in microstructural constitutions in the incubation period preceding crack initiation. Nucleation of microscopic defects then takes place and leads to the formation of a macroscopic flaw. Crack growth then follows, leading to eventual failure. The term fatigue damage in a broad sense refers to such irreversible changes in specimens, and plastic strain is the primary factor that induces such structural changes. Besides surface defects, heterogeneities of microstructures such as second phase particles and grain boundaries are the sites of stress and strain concentration. Formation of slip bands induces strain localization, and persistent slip bands (PSBs) have been observed mostly in fcc crystals [1]. The PSBs are sites involved with crack nucleation, and the process leading to the formation of PSBs is important not only for their
* Corresponding author. Tel.: /81-3-5286-3316; fax: /81-3-32002567 E-mail address: [email protected] (M. Nagumo).
configuration as a stress concentrator, but also for creation of defects. The characteristic dislocation veinstructures of PSBs are walls of dislocation dipoles [2], and annihilation of the dipoles creates a high density of vacancies [3]. One problem in steels is whether microstructural changes like PSBs take place, but the development of slip bands may result in substantial defect creation in localized regions. Once a crack is formed, a small-scale yielded-zone is associated with the crack tip, wherein strain localization becomes prominent. Fatigue resistance is strongly affected by the surrounding environment, and several mechanisms have been proposed as documented in review articles [4,5]. The operating mechanism may differ by materials and environmental conditions. In the case of a corrosive environment, one model [6] has ascribed the reduction of fatigue resistance to the role of hydrogen liberated in corrosion reactions. Accelerated transport of hydrogen by dislocation motion [7] and a similarity with stress corrosion cracking [5] have been discussed as the roles of hydrogen embrittlement in fatigue failure. In a gaseous hydrogen environment, an accelerated crack growth rate and formation of sub-critical cleavage of the ferrite
0921-5093/02/$ - see front matter # 2002 Elsevier Science B.V. All rights reserved. PII: S 0 9 2 1 - 5 0 9 3 ( 0 2 ) 0 0 7 4 5 - 1
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matrix were reported in a duplex stainless steel [8]. A recent topic concerning high-strength steels is fatigue failure after a very high number of cycles [9], which appears in two stages or with the disappearance of the fatigue limit. Murakami et al. have revealed [10] that a small amount of hydrogen plays a role associated with the evolution of an optically dark area around nonmetallic inclusions. Involvement of hydrogen in environmental fatigue is fairly evident, but the function of hydrogen and its relation to fatigue damage are issues that are still not clearly understood. The mechanism of hydrogen-related failure has been discussed for decades. Most theories [11 /13] have presumed the function of hydrogen in solid solution, but the solid solubility of hydrogen in iron is extremely low, say, 10 8 under one atmospheric pressure of hydrogen gas, according to Sievert’s Law [14], and most hydrogen is in trapped states at various defects. When the fugacity of hydrogen is high, molecular hydrogen is ready to precipitate at the interface between second phase particles and the matrix, but incipient cracks associated with hydrogen gas precipitation are hardly observed under mild environments or at low hydrogen concentrations. A point so far overlooked is the role of weakly trapped hydrogen. When vacancies are defects acting as traps, the binding energy with hydrogen reduces the effective formation energy of vacancies and drastically increases the vacancy density [15,16]. Then, if vacancies are created as one aspect of fatigue damage in the presence of hydrogen, interactions between vacancies and hydrogen are reasonably assumed to enhance the accumulation of fatigue damage. In a previous study [17], we have demonstrated an enhanced susceptibility to delayed fracture of a highstrength martensitic steel when the specimen had been subjected to a pre-fatigue treatment. The detrimental effect of pre-fatigue disappeared when the pre-fatigued specimens were subsequently annealed at temperatures as low as 250 8C. The observation implied that point defects created in the course of fatigue, presumably vacancies, were responsible for the enhanced susceptibility to delayed fracture. This enhanced susceptibility is in accord with our model of hydrogen-related failure [18] that states that strain-induced vacancies and their aggregates lead to failure, being assisted by hydrogen that increases vacancy density. The detection of vacancies is a difficult task particularly in structures with high densities of defects such as dislocations, grain boundaries, impurities and precipitates. A method we have developed [19] is to utilize hydrogen as a probe of defects by means of hydrogen thermal desorption analysis (TDA), which has showed the point defect nature of strain-induced defects in a ferritic steel. Since most hydrogen in steel is in trapped states at various defects, the thermal desorption characteristics give information about not only the density of
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defects but also the binding energies of hydrogen with defects [20,21]. There are two types of rate-determining process in the formation of a thermal desorption peak. One is the thermal dissociation of strong hydrogen/ defect couples [20], and the other is the diffusion of lattice hydrogen interacting with weakly trapped states under a local equilibrium [21]. In the latter case, a TDA peak is formed as a result of two competing processes on heating, i.e. an increasing diffusion rate and decreasing total hydrogen content. The presence of traps with a higher binding energy and a higher density shifts the desorption peak to higher temperatures. This method has been used in the present study to detect the defects created during fatigue. The purpose of the present study is firstly to reveal the fatigue damage in high-strength martensitic steels associated with microstructural changes, paying particular attention to the creation of point defects, and secondly to show the effect of hydrogen on fatigue resistance.
2. Experimental procedures Two types of martensitic steels in 5-mm-diameter round bars were used in the present study. One was a Si /Cr steel, JIS SWOSC-V, oil-quenched after being austenitized at 910 8C for 15 min followed by tempering at 460 8C for 9 min to a tensile strength of 1849 MPa. The other was a JIS G3137 steel for a pre-stressed concrete reinforcement (PC) bar, which was induction heat-treated to a tensile strength of 1493 MPa. The maximum temperatures for austenitizing and tempering were 1010 and 525 8C, respectively. The chemical compositions are shown in Table 1. The fatigue test was conducted with a rotational bending machine under constant stress amplitude with a rotational speed of 5000 cpm. Fig. 1 shows the S /N curves of the two steels in terms of the stress ratio, Ds/sB, vs. cycles to failure. Ds is the applied stress amplitude and sB is the ultimate tensile stress. After obtaining the S /N curves, specimens subjected to various fatigue cycles were prepared at two stress amplitudes, one at a level of 0.33 Ds /sB and the other close to the fatigue limit, 0.31 Ds/sB (580 MPa) and 0.29 Ds/sB (430 MPa) for the Si /Cr steel and the PC bar, respectively. The stress levels are also indicated in Fig. 1 by dashed lines. In order to discriminate hydrogen-trap sites, some fatigued specimens were subsequently annealed at 200 8C for 1 h to annihilate some of the defects, and attention was paid to the annihilation of point defects. In order to detect defects created during fatigue, hydrogen was introduced to fatigued-specimens as a probe and was subjected to TDA. Prior to hydrogencharging, the specimen surface was polished with #320 SiC-grit paper and cleaned with acetone. Hydrogencharging was conducted by cathodic electrolysis using a
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Table 1 Chemical compositions of the tested steels (mass %)
Si /Cr PC-bar
C
Si
Mn
P
S
Cu
Cr
Al
0.57 0.316
1.42 1.64
0.65 0.75
0.012 0.009
0.006 0.004
0.01 0.05
0.67 /
/ 0.023
3. Experimental results
Fig. 1. Stress amplitude vs. number of cycles to failure of the tested steels. The stress amplitude, Ds , is expressed as the ratio to the tensile strength, sB. The stress levels for sampling during fatigue are indicated by dashed lines.
3% NaCl/3 g l 1 NH4SCN aqueous solution at a current density of 0.75 mA cm 2 for 12 h, which exceeded the period for attaining the saturated hydrogen concentration. In order to examine the binding strength of hydrogen with defects, some hydrogen-charged specimens were kept at 30 8C for 30 h for degassing ‘diffusive’ hydrogen before TDA. The hydrogen content was measured by means of TDA using a gas-chromatograph with a linear heating rate of 200 8C h 1 up to 400 8C. The sample gas was analyzed at 5-min intervals using Ar as the carrier gas. Since the precision of TDA is sensitive to how specimens are handled, the time to the start of TDA after hydrogen charging was fixed at 15 min so as to keep the same condition, except in the cases where ‘non-diffusive’ hydrogen was measured. The effect of hydrogen on fatigue resistance was examined using hydrogen-precharged Si /Cr steel specimens. Hydrogen was charged to the specimens prior to the fatigue test for 12 h using the same solution and current density as described above. The fatigue test was started within 15 min after the completion of hydrogencharging. Diffusing out of hydrogen during the fatigue test was anticipated, but its effect on the results was assumed in the analysis to be insignificant when the fatigue life was less than a few hours.
Defect creation and alteration of microstructural constitutions during fatigue were examined using hydrogen as a probe. The hydrogen content, i.e. integrated desorption peak intensity in TDA up to 400 8C, is shown in Fig. 2 for specimens subjected to various fatigue cycles. The quantity represents the hydrogen absorption capacity of the sample. The hydrogen content in the Si /Cr steel was substantial in the asheat-treated, non-deformed specimen, and was unchanged in the early stage of fatigue. With increasing fatigue cycles, the hydrogen content decreased significantly in specimens tested in long fatigue cycles at the applied stress near the fatigue limit. The decrease then turned to an increase in the final stage of fatigue. In the
Fig. 2. Integrated intensity of hydrogen desorption curve, i.e. hydrogen content, of (a) Si /Cr steel and (b) PC bar fatigued for various cycles. Arrows show hydrogen-content levels in the as-heat-treated samples.
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PC bar, the hydrogen content increased from that in the as-heat-treated specimen soon after the application of fatigue cycles, but the following changes were quite similar to those seen for the Si /Cr steel. A small change of hydrogen content in the TDA data must be carefully examined with respect to the precision of the analysis. Important are reproducibility and consistency, based on careful handling procedures. Besides the common behavior observed for the two steels, the present result reproduced the finding of a separately conducted preliminary investigation [22]. The amount of absorbed hydrogen was reduced nearly 90% when the specimens were kept for 30 h at 30 8C after hydrogen-charging, maintaining non-diffusive hydrogen at a level of 0.2 ppm irrespective of fatigue cycles. In these specimens, most traps of absorbed hydrogen were defects having weak binding with hydrogen. The origin of the change in the hydrogen content with fatigue cycles was examined by annealing the fatigued specimens at 200 8C for 1 h. Fig. 3 shows the hydrogen content in specimens with/without annealing, together with the difference between the hydrogen content of the two specimens. In the Si/Cr steel, Fig. 3(a), annealing
Fig. 3. Integrated peak intensity of hydrogen desorption curve of (a) Si /Cr steel and (b) PC bar specimens. The solid and open marks are for as-fatigued and subsequently annealed specimens, respectively. The (') and (^) marks denote the difference of the upper two curves. Arrows show hydrogen-content levels in the as-heat-treated samples.
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hardly affected the hydrogen absorbing capacity until the fatigue cycle when the hydrogen desorption began to increase, but thereafter a remarkable reduction by annealing was observed at both applied stress levels. The reduction by annealing implies the creation of point defects in the final stage of fatigue. On the other hand, in the PC bar, Fig. 3(b), annealing reduced the increased total hydrogen content to the original level of nonloaded specimens. With the progress of fatigue, the reduction by annealing was almost unchanged or slightly decreased and finally turned to an increase in a manner similar to the Si/Cr steel. The hydrogen content is an important quantity relevant to defect density, but the TDA curve profile gives additional information about the binding strength of hydrogen with defects. Fig. 4 shows examples of TDA curves before and after annealing for the Si/Cr steel and PC bar. The difference, dHan, between the curves of asfatigued and subsequently annealed specimens is shown in Figs. 5 and 6 and corresponds to the decrease of desorption by annealing. The change in the dHan profile with increasing fatigue cycles suggests alterations in annealed out defects that are concerned with desorption. Inevitable scatters are present in dHan because of the precision of the original TDA curves, and distinct
Fig. 4. Examples of hydrogen TDA curves of fatigued specimens with/ without subsequent annealing. (a) Si /Cr steel tested at 610 MPa for 1 /106 cycles and (b) PC bar tested at 490 MPa for 2 /105 cycles.
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Fig. 5. Difference between TDA curves of as-fatigued and subsequently annealed Si /Cr steel specimens; tested at (a) 610 MPa and (b) 580 MPa for various cycles.
features are to be noted. In the Si/Cr steel tested at 610 MPa, Fig. 5(a), a slight decrease of desorption by annealing took place initially on the higher temperature side of the desorption peak, and then the decrease appeared on the lower temperature side around 100 8C as the fatigue cycles increased. Similar behavior was also observed when this steel was tested at 580 MPa near the fatigue limit, Fig. 5(b). In the PC bar tested at 490 MPa, Fig. 6(a), the decreases took place initially in the middle part of the peak and then on the higher temperature side when the total hydrogen content began to increase at 4 /105 cycles. At the applied stress of 430 MPa near the fatigue limit, Fig. 6(b), the initial decrease in the middle part of the desorption peak and the decrease on the low temperature side in the late stage appeared in a similar way as in Fig. 5 and Fig. 6(a). The effect of hydrogen on fatigue resistance was examined with the Si /Cr steel. The S /N curve of hydrogen-precharged specimens is shown in Fig. 7 together with the curve of non-charged specimens in Fig. 1. The fatigue life was drastically reduced in hydrogen-precharged specimens at higher stress amplitudes than the fatigue limit. Since the charged-hydrogen was ready to diffuse out during the fatigue test, the effect of hydrogen would be diminished as the fatigue cycles
Fig. 6. Difference between TDA curves of as-fatigued and subsequently annealed PC bar specimens; tested at (a) 490 MPa and (b) 430 MPa for various cycles.
Fig. 7. S /N curves of the Si /Cr steel with/without hydrogenprecharging. The arrows indicate no failure.
increased. However, in Fig. 7, the stress level at which the fatigue life was noticeably extended was only slightly altered by hydrogen-precharging, suggesting that the presence of hydrogen did not significantly affect the fatigue limit. The defect density in specimens fatigue-fractured in the presence of hydrogen was evaluated in terms of the absorption capacity of hydrogen recharged to the fractured specimens after degassing pre-charged hydro-
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Fig. 8. Hydrogen absorption capacity in hydrogen-precharged and fatigue-fractured Si /Cr steel specimens, compared with specimens fatigued without hydrogen-precharging. The fatigue cycle applied to each specimen was respectively 6 /106 or 5.5 /106 for the noncharged or pre-charged specimens at the stress level of 580 MPa, and 2 /105 or 8 /104 for the non-charged or pre-charged specimens at the stress level of 610 MPa.
gen by keeping them at room temperature for more than 50 h. Fig. 8 shows the hydrogen content in fractured Si/ Cr steel specimens compared with specimens subjected to exceeding fatigue cycles without hydrogen-precharging. The fatigue cycle applied to each specimen was respectively 6 /106 or 5.5 /106 for the non-charged or pre-charged specimens at the stress level of 580 MPa, and 2 /105 or 8 /104 for the non-charged or precharged specimens at the stress level of 610 MPa. The presence of hydrogen during fatigue prominently increased the hydrogen content, which implies that the earlier failure in the presence of hydrogen resulted from an increased density of defects. The fractographic features near the crack initiation site of the Si /Cr steel specimens tested at 610 MPa with/ without hydrogen-precharging are shown in Fig. 9(a and b), respectively. When hydrogen was present, the fracture surface was fairly flat with some irregular streaks. The flat area extended across grains with a diameter of about 10 mm, and no crystallographic units characterized by cleavage or intergranular cracking were observed. The extent of the flat area was about half of the whole fracture surface, and the remaining area showed dimple-type fracture. In the hydrogen-precharged specimen, the flat area occasionally mixed with the dimples. Involvement of hydrogen in the crack growth stage is evident and is associated with microstructural changes that precede the crack advance.
4. Discussion The change in hydrogen content with the progress of fatigue shown in Fig. 2 indicates the change of interest in the defect density detected using hydrogen as a probe. While the extent differed by steels and stress levels, near
Fig. 9. Fractograph of a Si /Cr steel specimen subjected to a fatigue test at a stress amplitude of 610 MPa (a) with and (b) without hydrogen-precharging.
constancy or a decrease in the early stage followed by turning to an increase was observed in almost all cases. In a rotational bending test, a strain gradient exists across the diameter of a specimen, and strain localization is expected in the area near the surface. Since TDA gives information averaged over an entire specimen, the change in TDA would be more pronounced if we could detect hydrogen only in the near-surface area. Martensite in steel is a structure with a high density of various lattice defects that act as traps of hydrogen. Most of the hydrogen absorbed in as-heat-treated specimens was diffusive at room temperature, and the defects of interest that acted as traps were ones weakly coupled with hydrogen. The increase in hydrogen content in the final stage was due to the creation of point defects, presumably vacancies, which are annealed out at 200 8C as shown in Fig. 3. The annealing out of defects that trap the strain-induced increment of hydrogen is quite similar to
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previous results reported for tensile straining of ferritic [19] and martensitic [23,24] steels. The turn to the increase took place at cycles corresponding to about 50% of the fatigue life, suggesting that the creation of point defects is concerned with the crack growth stage rather than the incubation period. On the other hand, the preceding decrease in hydrogen content, remarkable in long fatigue cycles when specimens were tested at a stress level near the fatigue limit, was not associated with the creation of point defects in the Si /Cr steel, Fig. 3(a). The decrease might be ascribed to changes in microstructural constitutions that compose the major part of hydrogen traps and associated desorption. Dislocations in as-heat-treated martensite are a kind of trap for diffusive hydrogen, and their unstable configurations are apt to be affected by cyclic stressing. The decrease in hydrogen content in the early stage likely resulted from reconfiguration or annihilation of dislocations, reducing the dislocation density [25]. In the PC bar, Fig. 3(b), the decrease in the early stage was more prominent than in the Si /Cr steel and was associated with point defect creation. Since the PC bar has been tempered at temperatures higher than 200 8C, the creation of point defects in the early stage must have taken place soon after the application of fatigue cycles. The creation of point defects in the early stage, contrary to the Si /Cr steel, might be due to the difference in dislocation-configurations. The PC bar was induction heat-treated, and dislocation-configurations might be less stable than that in the Si /Cr steel. The strain-induced creation of point defect is likely to originate in interactions between dislocations, such as annihilation of dislocation dipoles [3] or dragging of intersection jogs [26]. The stability of dislocation substructures might be different between induction and furnace heat-treated steels, but a detailed discussion is beyond the scope of the present study. A consequence of the above considerations is that the creation process of point defects differs between the early and final stages of fatigue. A relevant feature might be the change in TDA curve profiles between the as-fatigued and subsequently annealed specimens as shown in Figs. 5 and 6. When the TDA curve is determined by the effective diffusion constant of hydrogen [27], the curve profile changes when the trapping defects are altered. Annealing in the final stage of fatigue reduced the desorption on the low temperature side of the peak around 100 8C, both in the Si/Cr steel, Fig. 5, and the PC bar subjected to long fatigue cycles, Fig. 6(b). The reduction was due to annealing out of point defects created after a substantial number of fatigue cycles, presumably after evolution of considerable strain localization in areas such as extended slip bands or crack front. The decrease of hydrogen desorption on the lower temperature side implies that the annealed out defects have a low binding energy with
hydrogen. On the other hand, annealing out of point defects, created soon after the application of fatigue cycles in the PC bar, appeared in the middle part of the desorption curve, Fig. 6. The implication is that the binding energy with hydrogen of point defects, presumably vacancies, created in the early stage is slightly higher than that of defects created in the final stage, presumably by coupling with impurity atoms or clustering [28]. Since extended strain localization was not expected in the early stage, reconfiguration and annihilation of dislocations in a wide region might be a possible mechanism of point defect creation. As for the involvement of hydrogen in fatigue failure, Fig. 7 shows a remarkable effect of hydrogen on reducing the fatigue life, while the decrease of the fatigue limit was not significant. Miller suggested [29] that the fatigue limit is characterized by the maximum non-propagating crack size, further crack growth being permitted by the increased stress level. If the controlling process of the fatigue limit is the crack growth overcoming a strong barrier to crack propagation, the implication of Fig. 7, which shows a prominent effect of hydrogen on fatigue life, is that hydrogen drastically reduces the resistance to the once-started crack growth. The decrease in crack growth resistance in the presence of hydrogen was observed previously in a three-point bending test of ferritic steels [30], in which a finite element analysis showed prominent strain localization in front of the advancing crack associated with an increased density of nucleation voids. The fractographic feature shown in Fig. 9(a) suggests microstructural changes associated with the crack advance. The increased defect density in the fractured hydrogen-precharged specimens shown in Fig. 8 also supports the idea that hydrogen is a key factor in reducing fatigue life. However, fatigue failure in the presence of hydrogen occurred much earlier than the stage in which crack growth presumably took place as inferred from the point defect creation in Fig. 3. The relation of hydrogen to crack nucleation may also be involved in the reduction of fatigue life. At the stress level above the fatigue limit, an incipient crack is ready to grow overcoming various barriers. If hydrogen favors the nucleation of an incipient crack, the nucleation as well as the growth of a fatigue crack is expected to reduce the fatigue life. Fatigue crack initiation associated with the formation of well-developed PSBs has been reported [31], but the results in Fig. 7 suggest that hydrogen also affected premature alterations of the substructures in the incubation period, leading to the crack nucleation. In Fig. 7, the effect of hydrogen in reducing the fatigue life was more prominent at lower stress levels. The applied stress amplitude in the present experiment was as low as about one third of the ultimate tensile strength, implying strain localization in very limited areas. In such situa-
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tions, the extent of plastic deformation that interacts with hydrogen must depend on the stress level, resulting in a sensitive effect of hydrogen on the crack nucleation and growth process through interactions with plastic deformation. The effect of hydrogen in increasing vacancy density has been well recognized [15,16]. One mechanism proposed by Essmann concerning the formation of PSBs is the annihilation of dislocation dipoles [3]. It is not definite if his proposed model can be applied to bcc alloys, but vacancy formation must originate in interactions of moving dislocations, which may take place preferentially in an area of high dislocation density. The PSBs and crack front area are sites of substantial strain localization, and the effect of hydrogen in increasing vacancy density could be prominent once a crack is formed as a strain concentrator. However, hydrogen might accelerate the formation of substructures like PSBs in the incubation period before fatigue crack initiation through synergistic interactions between hydrogen and dislocation movement. The decrease in hydrogen content preceding the final stage of increasing hydrogen content suggests reconfigurations of dislocations and the involvement of hydrogen in the process. Interactions between hydrogen and vacancies are expected to affect fatigue resistance in both the crack initiation and growth stages. Vacancy formation during fatigue is consistent with a previous result [17] that showed an increased susceptibility to delayed fracture in the same Si /Cr steel bar that had been given pre-fatigue treatment by a rotational bending fatigue test at a stress amplitude of 640 MPa and a rotational speed of 5000 rpm for 8 /104 cycles, close to the fatigue life, and 4 /104 cycles. Since the increased susceptibility recovered to the original level by subsequently annealing the pre-fatigued specimens at 200 8C, the effect was ascribed to an interaction of hydrogen with vacancies created during fatigue. In that case, if vacancy creation in the pre-fatigued specimens was associated with a definite crack, the onceformed crack would remain after annealing, still favoring the increase in delayed fracture susceptibility. Then, it is reasonable to consider that fatigue damage associated with vacancy creation is not confined to the area in front of an existing fatigue crack.
5. Summary and conclusions Defects created during rotational bending fatigue in high-strength martensitic steels have been detected utilizing hydrogen as a probe. The amount of absorbed hydrogen decreased or was almost constant in the early stage of fatigue, and then turned to an increase after a substantial number of fatigue cycles. Subsequent annealing of fatigued specimens reduced the increased hydro-
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gen in the final stage, suggesting the creation of point defects, presumably vacancies, in that stage. On the other hand, the decrease in the early stage was ascribed to changes in microstructural constitutions, reconfigurations or annihilations of dislocations. In the induction heat-treated PC bar, point defect creation also took place in the early stage, presumably associated with reconfiguration process of dislocations. Hydrogen-precharging drastically reduced fatigue life, but the effect of hydrogen on the fatigue limit was not significant. The defect density in fractured specimens in the presence of hydrogen was higher than in specimens fatigued for similar cycles without hydrogen, indicating a hydrogenassisted increase in defect density. The fracture surface of hydrogen-precharged specimens was smooth without cleavage-like crystallographic features. The reduction of fatigue life at applied stress levels above the fatigue limit and significant creation of point defects in the final stage of the fatigue process imply that hydrogen reduces crack growth resistance coupled with point defects. However, very early fracture in the presence of hydrogen and changes in microstructural constitutions in the early stage suggest that hydrogen is also involved in crack initiation through interactions with substructural changes in the incubation period.
Acknowledgements The present work has been conducted under a contract with the Ministry of Education, Culture, Sports, Science and Technology of Japan. The authors also wish to thank the Netsuren Co. Ltd for supplying the PC steel bars.
References [1] C. Laird, in: R.W. Cahn, P.H. Haasen (Eds.), Physical Metallurgy, fourth ed., Elsevier Science, Amsterdam, 1996, p. 2293. [2] L.P. Karjalainen, Mater. Sci. Eng. 38 (1979) 145. [3] U. Essmann, H. Mughrabi, Phil. Mag. 40 (1979) 731. [4] R.P. Wei, in: J. Fong (Ed.), Fatigue Mechanism, STP 675, ASTM, Philadelphia, 1979, p. 826. [5] D.J. Duquette, Hydrogen Embrittlement and Stress Corrosion Cracking, ASM, 1980, p. 250. [6] R.J. Jacko, D.H. Duquette, in: I.M. Bernstein, A.W. Thompson (Eds.), Hydrogen Effects in Materials, AIME, Warrendale, PA, 1981, p. 477. [7] H.H. Johnson, J.P. Hirth, Metall. Trans. A 7A (1976) 1543. [8] T.J. Marrow, C.A. Hippsley, J.E. King, Acta Metall. Mater. 39 (1991) 1367. [9] T. Sakai, M. Takeda, K. Shiozawa, Y. Ochi, M. Nakajima, T. Nakamura, N. Oguma, in: X.-R. Wu, Z.G. Wang (Eds.), Fatigue’99, vol. I, Higher Education Press, Beijing, 1999, p. 573. [10] Y. Murakami, H. Konishi, K. Takai, Y. Murakami, Tetsu-toHagane 86 (2000) 777. [11] R.A. Oriani, P.H. Josephic, Acta Metall. 22 (1974) 1065. [12] H.K. Birnbaum, P. Sofronis, Mater. Sci. Eng. A 176 (1994) 191.
200 [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23]
M. Nagumo et al. / Materials Science and Engineering A348 (2003) 192 /200 C.D. Beachem, Metall. Trans. 3 (1972) 437. J.P. Hirth, Metall. Trans. A 11A (1980) 861. R.B. McLellan, L. Yang, Acta Metall. Mater. 43 (1995) 2463. Y. Fukai, Y. Kurokawa, H. Hiraoka, J. Jpn Inst. Metals 61 (1997) 663. M. Nagumo, S. Sekiguchi, H. Hayashi, T. Takai, Mater. Sci. Eng. A, in press. M. Nagumo, M. Nakamura, K. Takai, Metallur. Mater. Trans. A 32A (2001) 339. M. Nagumo, K. Ohta, H. Saitoh, Scripta Mater. 40 (1999) 313. W.Y. Choo, J.Y. Lee, Metall. Trans. A 13A (1982) 135. K. Ono, M. Meshii, Acta Metall. Mater. 40 (1992) 1357. H. Shimura, T. Chaya, M. Nagumo, H. Hayashi, I. Ochiai, CAMP ISIJ 15 (2002) 572. S. Suzuki, N. Ishii, Y. Tsuchida, Te´tsu-to-Hagane´ 80 (1994) 855.
[24] M. Aoki, H. Saitoh, M. Mori, Y. Ishida, M. Nagumo, J. Jpn Inst. Metals 58 (1994) 1141. [25] K. Rahka, P. Nenonen, Z.-G. Wang, C. Laird, in: P.O. Kettunen, T.K. Lepisto¨, M.E. Lehtonen (Eds.), Strength of Metals and Alloys, vol. 2, Pergamon Press, Oxford, 1989, p. 767. [26] A.M. Cuitino, M. Ortiz, Acta Mater. 44 (1996) 427. [27] R.A. Oriani, Acta Metall. 18 (1970) 147. [28] A. Vehanen, P. Hautojarvi, J. Johnson, J. Yli-Kauppila, Phys. Rev. B 25 (1982) 762. [29] K.J. Miller, in: A.S. Argon (Ed.), Topics in Fracture and Fatigue, Springer, New York, 1992, p. 309. [30] M. Nagumo, H. Yoshida, Y. Shimomura, T. Kadokura, Mater. Trans. 42 (2001) 132. [31] A. Hunsche, P. Neumann, Acta Metall. 34 (1986) 207.
Fatigue damage and its interaction with hydrogen in martensitic steels M. Nagumo a,b,*, H. Shimura a, T. Chaya a, H. Hayashi c, I. Ochiai c a
Department of Materials Science and Engineering, Waseda University, Okubo 3-4-1, Shinjuku, Tokyo 169-8555, Japan b Laboratory for Materials Science and Technology, Waseda University, Tokyo, Japan c R&D Section, Suzuki Metal Co., Ltd, Higashi-Narashino 7-5-1, Narashino, Chiba 275-8577, Japan Received 19 June 2002; received in revised form 19 September 2002
Abstract Fatigue damage involving creation of point defects has been revealed by means of hydrogen thermal desorption analysis, utilizing hydrogen as a probe of defects. Two types of high-strength martensitic steels were subjected to a rotational bending test for various cycles. The amount of absorbed hydrogen decreased or was almost constant in the early stage of the fatigue process and then began to increase after a substantial number of fatigue cycles. The increase in the final stage was due to the creation of point defects, presumably vacancies, while the decrease in the early stage was ascribed to changes in microstructural constitutions. Hydrogenprecharging drastically reduced the fatigue life, but the effect of hydrogen on the fatigue limit was not significant. The fracture surface of hydrogen-precharged specimens was smooth without cleavage-like crystallographic features. The defect density in fractured specimens in the presence of hydrogen was higher than in specimens fatigued for a similar number of cycles without hydrogen. Involvement of hydrogen in fatigue damage in the crack nucleation and growth stages is discussed. # 2002 Elsevier Science B.V. All rights reserved. Keywords: Martensitic steel; Fatigue; Fatigue damage; Vacancy; Hydrogen thermal desorption; Hydrogen embrittlement
1. Introduction The fatigue process initiates with irreversible changes in microstructural constitutions in the incubation period preceding crack initiation. Nucleation of microscopic defects then takes place and leads to the formation of a macroscopic flaw. Crack growth then follows, leading to eventual failure. The term fatigue damage in a broad sense refers to such irreversible changes in specimens, and plastic strain is the primary factor that induces such structural changes. Besides surface defects, heterogeneities of microstructures such as second phase particles and grain boundaries are the sites of stress and strain concentration. Formation of slip bands induces strain localization, and persistent slip bands (PSBs) have been observed mostly in fcc crystals [1]. The PSBs are sites involved with crack nucleation, and the process leading to the formation of PSBs is important not only for their
* Corresponding author. Tel.: /81-3-5286-3316; fax: /81-3-32002567 E-mail address: [email protected] (M. Nagumo).
configuration as a stress concentrator, but also for creation of defects. The characteristic dislocation veinstructures of PSBs are walls of dislocation dipoles [2], and annihilation of the dipoles creates a high density of vacancies [3]. One problem in steels is whether microstructural changes like PSBs take place, but the development of slip bands may result in substantial defect creation in localized regions. Once a crack is formed, a small-scale yielded-zone is associated with the crack tip, wherein strain localization becomes prominent. Fatigue resistance is strongly affected by the surrounding environment, and several mechanisms have been proposed as documented in review articles [4,5]. The operating mechanism may differ by materials and environmental conditions. In the case of a corrosive environment, one model [6] has ascribed the reduction of fatigue resistance to the role of hydrogen liberated in corrosion reactions. Accelerated transport of hydrogen by dislocation motion [7] and a similarity with stress corrosion cracking [5] have been discussed as the roles of hydrogen embrittlement in fatigue failure. In a gaseous hydrogen environment, an accelerated crack growth rate and formation of sub-critical cleavage of the ferrite
0921-5093/02/$ - see front matter # 2002 Elsevier Science B.V. All rights reserved. PII: S 0 9 2 1 - 5 0 9 3 ( 0 2 ) 0 0 7 4 5 - 1
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matrix were reported in a duplex stainless steel [8]. A recent topic concerning high-strength steels is fatigue failure after a very high number of cycles [9], which appears in two stages or with the disappearance of the fatigue limit. Murakami et al. have revealed [10] that a small amount of hydrogen plays a role associated with the evolution of an optically dark area around nonmetallic inclusions. Involvement of hydrogen in environmental fatigue is fairly evident, but the function of hydrogen and its relation to fatigue damage are issues that are still not clearly understood. The mechanism of hydrogen-related failure has been discussed for decades. Most theories [11 /13] have presumed the function of hydrogen in solid solution, but the solid solubility of hydrogen in iron is extremely low, say, 10 8 under one atmospheric pressure of hydrogen gas, according to Sievert’s Law [14], and most hydrogen is in trapped states at various defects. When the fugacity of hydrogen is high, molecular hydrogen is ready to precipitate at the interface between second phase particles and the matrix, but incipient cracks associated with hydrogen gas precipitation are hardly observed under mild environments or at low hydrogen concentrations. A point so far overlooked is the role of weakly trapped hydrogen. When vacancies are defects acting as traps, the binding energy with hydrogen reduces the effective formation energy of vacancies and drastically increases the vacancy density [15,16]. Then, if vacancies are created as one aspect of fatigue damage in the presence of hydrogen, interactions between vacancies and hydrogen are reasonably assumed to enhance the accumulation of fatigue damage. In a previous study [17], we have demonstrated an enhanced susceptibility to delayed fracture of a highstrength martensitic steel when the specimen had been subjected to a pre-fatigue treatment. The detrimental effect of pre-fatigue disappeared when the pre-fatigued specimens were subsequently annealed at temperatures as low as 250 8C. The observation implied that point defects created in the course of fatigue, presumably vacancies, were responsible for the enhanced susceptibility to delayed fracture. This enhanced susceptibility is in accord with our model of hydrogen-related failure [18] that states that strain-induced vacancies and their aggregates lead to failure, being assisted by hydrogen that increases vacancy density. The detection of vacancies is a difficult task particularly in structures with high densities of defects such as dislocations, grain boundaries, impurities and precipitates. A method we have developed [19] is to utilize hydrogen as a probe of defects by means of hydrogen thermal desorption analysis (TDA), which has showed the point defect nature of strain-induced defects in a ferritic steel. Since most hydrogen in steel is in trapped states at various defects, the thermal desorption characteristics give information about not only the density of
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defects but also the binding energies of hydrogen with defects [20,21]. There are two types of rate-determining process in the formation of a thermal desorption peak. One is the thermal dissociation of strong hydrogen/ defect couples [20], and the other is the diffusion of lattice hydrogen interacting with weakly trapped states under a local equilibrium [21]. In the latter case, a TDA peak is formed as a result of two competing processes on heating, i.e. an increasing diffusion rate and decreasing total hydrogen content. The presence of traps with a higher binding energy and a higher density shifts the desorption peak to higher temperatures. This method has been used in the present study to detect the defects created during fatigue. The purpose of the present study is firstly to reveal the fatigue damage in high-strength martensitic steels associated with microstructural changes, paying particular attention to the creation of point defects, and secondly to show the effect of hydrogen on fatigue resistance.
2. Experimental procedures Two types of martensitic steels in 5-mm-diameter round bars were used in the present study. One was a Si /Cr steel, JIS SWOSC-V, oil-quenched after being austenitized at 910 8C for 15 min followed by tempering at 460 8C for 9 min to a tensile strength of 1849 MPa. The other was a JIS G3137 steel for a pre-stressed concrete reinforcement (PC) bar, which was induction heat-treated to a tensile strength of 1493 MPa. The maximum temperatures for austenitizing and tempering were 1010 and 525 8C, respectively. The chemical compositions are shown in Table 1. The fatigue test was conducted with a rotational bending machine under constant stress amplitude with a rotational speed of 5000 cpm. Fig. 1 shows the S /N curves of the two steels in terms of the stress ratio, Ds/sB, vs. cycles to failure. Ds is the applied stress amplitude and sB is the ultimate tensile stress. After obtaining the S /N curves, specimens subjected to various fatigue cycles were prepared at two stress amplitudes, one at a level of 0.33 Ds /sB and the other close to the fatigue limit, 0.31 Ds/sB (580 MPa) and 0.29 Ds/sB (430 MPa) for the Si /Cr steel and the PC bar, respectively. The stress levels are also indicated in Fig. 1 by dashed lines. In order to discriminate hydrogen-trap sites, some fatigued specimens were subsequently annealed at 200 8C for 1 h to annihilate some of the defects, and attention was paid to the annihilation of point defects. In order to detect defects created during fatigue, hydrogen was introduced to fatigued-specimens as a probe and was subjected to TDA. Prior to hydrogencharging, the specimen surface was polished with #320 SiC-grit paper and cleaned with acetone. Hydrogencharging was conducted by cathodic electrolysis using a
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Table 1 Chemical compositions of the tested steels (mass %)
Si /Cr PC-bar
C
Si
Mn
P
S
Cu
Cr
Al
0.57 0.316
1.42 1.64
0.65 0.75
0.012 0.009
0.006 0.004
0.01 0.05
0.67 /
/ 0.023
3. Experimental results
Fig. 1. Stress amplitude vs. number of cycles to failure of the tested steels. The stress amplitude, Ds , is expressed as the ratio to the tensile strength, sB. The stress levels for sampling during fatigue are indicated by dashed lines.
3% NaCl/3 g l 1 NH4SCN aqueous solution at a current density of 0.75 mA cm 2 for 12 h, which exceeded the period for attaining the saturated hydrogen concentration. In order to examine the binding strength of hydrogen with defects, some hydrogen-charged specimens were kept at 30 8C for 30 h for degassing ‘diffusive’ hydrogen before TDA. The hydrogen content was measured by means of TDA using a gas-chromatograph with a linear heating rate of 200 8C h 1 up to 400 8C. The sample gas was analyzed at 5-min intervals using Ar as the carrier gas. Since the precision of TDA is sensitive to how specimens are handled, the time to the start of TDA after hydrogen charging was fixed at 15 min so as to keep the same condition, except in the cases where ‘non-diffusive’ hydrogen was measured. The effect of hydrogen on fatigue resistance was examined using hydrogen-precharged Si /Cr steel specimens. Hydrogen was charged to the specimens prior to the fatigue test for 12 h using the same solution and current density as described above. The fatigue test was started within 15 min after the completion of hydrogencharging. Diffusing out of hydrogen during the fatigue test was anticipated, but its effect on the results was assumed in the analysis to be insignificant when the fatigue life was less than a few hours.
Defect creation and alteration of microstructural constitutions during fatigue were examined using hydrogen as a probe. The hydrogen content, i.e. integrated desorption peak intensity in TDA up to 400 8C, is shown in Fig. 2 for specimens subjected to various fatigue cycles. The quantity represents the hydrogen absorption capacity of the sample. The hydrogen content in the Si /Cr steel was substantial in the asheat-treated, non-deformed specimen, and was unchanged in the early stage of fatigue. With increasing fatigue cycles, the hydrogen content decreased significantly in specimens tested in long fatigue cycles at the applied stress near the fatigue limit. The decrease then turned to an increase in the final stage of fatigue. In the
Fig. 2. Integrated intensity of hydrogen desorption curve, i.e. hydrogen content, of (a) Si /Cr steel and (b) PC bar fatigued for various cycles. Arrows show hydrogen-content levels in the as-heat-treated samples.
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PC bar, the hydrogen content increased from that in the as-heat-treated specimen soon after the application of fatigue cycles, but the following changes were quite similar to those seen for the Si /Cr steel. A small change of hydrogen content in the TDA data must be carefully examined with respect to the precision of the analysis. Important are reproducibility and consistency, based on careful handling procedures. Besides the common behavior observed for the two steels, the present result reproduced the finding of a separately conducted preliminary investigation [22]. The amount of absorbed hydrogen was reduced nearly 90% when the specimens were kept for 30 h at 30 8C after hydrogen-charging, maintaining non-diffusive hydrogen at a level of 0.2 ppm irrespective of fatigue cycles. In these specimens, most traps of absorbed hydrogen were defects having weak binding with hydrogen. The origin of the change in the hydrogen content with fatigue cycles was examined by annealing the fatigued specimens at 200 8C for 1 h. Fig. 3 shows the hydrogen content in specimens with/without annealing, together with the difference between the hydrogen content of the two specimens. In the Si/Cr steel, Fig. 3(a), annealing
Fig. 3. Integrated peak intensity of hydrogen desorption curve of (a) Si /Cr steel and (b) PC bar specimens. The solid and open marks are for as-fatigued and subsequently annealed specimens, respectively. The (') and (^) marks denote the difference of the upper two curves. Arrows show hydrogen-content levels in the as-heat-treated samples.
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hardly affected the hydrogen absorbing capacity until the fatigue cycle when the hydrogen desorption began to increase, but thereafter a remarkable reduction by annealing was observed at both applied stress levels. The reduction by annealing implies the creation of point defects in the final stage of fatigue. On the other hand, in the PC bar, Fig. 3(b), annealing reduced the increased total hydrogen content to the original level of nonloaded specimens. With the progress of fatigue, the reduction by annealing was almost unchanged or slightly decreased and finally turned to an increase in a manner similar to the Si/Cr steel. The hydrogen content is an important quantity relevant to defect density, but the TDA curve profile gives additional information about the binding strength of hydrogen with defects. Fig. 4 shows examples of TDA curves before and after annealing for the Si/Cr steel and PC bar. The difference, dHan, between the curves of asfatigued and subsequently annealed specimens is shown in Figs. 5 and 6 and corresponds to the decrease of desorption by annealing. The change in the dHan profile with increasing fatigue cycles suggests alterations in annealed out defects that are concerned with desorption. Inevitable scatters are present in dHan because of the precision of the original TDA curves, and distinct
Fig. 4. Examples of hydrogen TDA curves of fatigued specimens with/ without subsequent annealing. (a) Si /Cr steel tested at 610 MPa for 1 /106 cycles and (b) PC bar tested at 490 MPa for 2 /105 cycles.
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Fig. 5. Difference between TDA curves of as-fatigued and subsequently annealed Si /Cr steel specimens; tested at (a) 610 MPa and (b) 580 MPa for various cycles.
features are to be noted. In the Si/Cr steel tested at 610 MPa, Fig. 5(a), a slight decrease of desorption by annealing took place initially on the higher temperature side of the desorption peak, and then the decrease appeared on the lower temperature side around 100 8C as the fatigue cycles increased. Similar behavior was also observed when this steel was tested at 580 MPa near the fatigue limit, Fig. 5(b). In the PC bar tested at 490 MPa, Fig. 6(a), the decreases took place initially in the middle part of the peak and then on the higher temperature side when the total hydrogen content began to increase at 4 /105 cycles. At the applied stress of 430 MPa near the fatigue limit, Fig. 6(b), the initial decrease in the middle part of the desorption peak and the decrease on the low temperature side in the late stage appeared in a similar way as in Fig. 5 and Fig. 6(a). The effect of hydrogen on fatigue resistance was examined with the Si /Cr steel. The S /N curve of hydrogen-precharged specimens is shown in Fig. 7 together with the curve of non-charged specimens in Fig. 1. The fatigue life was drastically reduced in hydrogen-precharged specimens at higher stress amplitudes than the fatigue limit. Since the charged-hydrogen was ready to diffuse out during the fatigue test, the effect of hydrogen would be diminished as the fatigue cycles
Fig. 6. Difference between TDA curves of as-fatigued and subsequently annealed PC bar specimens; tested at (a) 490 MPa and (b) 430 MPa for various cycles.
Fig. 7. S /N curves of the Si /Cr steel with/without hydrogenprecharging. The arrows indicate no failure.
increased. However, in Fig. 7, the stress level at which the fatigue life was noticeably extended was only slightly altered by hydrogen-precharging, suggesting that the presence of hydrogen did not significantly affect the fatigue limit. The defect density in specimens fatigue-fractured in the presence of hydrogen was evaluated in terms of the absorption capacity of hydrogen recharged to the fractured specimens after degassing pre-charged hydro-
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Fig. 8. Hydrogen absorption capacity in hydrogen-precharged and fatigue-fractured Si /Cr steel specimens, compared with specimens fatigued without hydrogen-precharging. The fatigue cycle applied to each specimen was respectively 6 /106 or 5.5 /106 for the noncharged or pre-charged specimens at the stress level of 580 MPa, and 2 /105 or 8 /104 for the non-charged or pre-charged specimens at the stress level of 610 MPa.
gen by keeping them at room temperature for more than 50 h. Fig. 8 shows the hydrogen content in fractured Si/ Cr steel specimens compared with specimens subjected to exceeding fatigue cycles without hydrogen-precharging. The fatigue cycle applied to each specimen was respectively 6 /106 or 5.5 /106 for the non-charged or pre-charged specimens at the stress level of 580 MPa, and 2 /105 or 8 /104 for the non-charged or precharged specimens at the stress level of 610 MPa. The presence of hydrogen during fatigue prominently increased the hydrogen content, which implies that the earlier failure in the presence of hydrogen resulted from an increased density of defects. The fractographic features near the crack initiation site of the Si /Cr steel specimens tested at 610 MPa with/ without hydrogen-precharging are shown in Fig. 9(a and b), respectively. When hydrogen was present, the fracture surface was fairly flat with some irregular streaks. The flat area extended across grains with a diameter of about 10 mm, and no crystallographic units characterized by cleavage or intergranular cracking were observed. The extent of the flat area was about half of the whole fracture surface, and the remaining area showed dimple-type fracture. In the hydrogen-precharged specimen, the flat area occasionally mixed with the dimples. Involvement of hydrogen in the crack growth stage is evident and is associated with microstructural changes that precede the crack advance.
4. Discussion The change in hydrogen content with the progress of fatigue shown in Fig. 2 indicates the change of interest in the defect density detected using hydrogen as a probe. While the extent differed by steels and stress levels, near
Fig. 9. Fractograph of a Si /Cr steel specimen subjected to a fatigue test at a stress amplitude of 610 MPa (a) with and (b) without hydrogen-precharging.
constancy or a decrease in the early stage followed by turning to an increase was observed in almost all cases. In a rotational bending test, a strain gradient exists across the diameter of a specimen, and strain localization is expected in the area near the surface. Since TDA gives information averaged over an entire specimen, the change in TDA would be more pronounced if we could detect hydrogen only in the near-surface area. Martensite in steel is a structure with a high density of various lattice defects that act as traps of hydrogen. Most of the hydrogen absorbed in as-heat-treated specimens was diffusive at room temperature, and the defects of interest that acted as traps were ones weakly coupled with hydrogen. The increase in hydrogen content in the final stage was due to the creation of point defects, presumably vacancies, which are annealed out at 200 8C as shown in Fig. 3. The annealing out of defects that trap the strain-induced increment of hydrogen is quite similar to
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previous results reported for tensile straining of ferritic [19] and martensitic [23,24] steels. The turn to the increase took place at cycles corresponding to about 50% of the fatigue life, suggesting that the creation of point defects is concerned with the crack growth stage rather than the incubation period. On the other hand, the preceding decrease in hydrogen content, remarkable in long fatigue cycles when specimens were tested at a stress level near the fatigue limit, was not associated with the creation of point defects in the Si /Cr steel, Fig. 3(a). The decrease might be ascribed to changes in microstructural constitutions that compose the major part of hydrogen traps and associated desorption. Dislocations in as-heat-treated martensite are a kind of trap for diffusive hydrogen, and their unstable configurations are apt to be affected by cyclic stressing. The decrease in hydrogen content in the early stage likely resulted from reconfiguration or annihilation of dislocations, reducing the dislocation density [25]. In the PC bar, Fig. 3(b), the decrease in the early stage was more prominent than in the Si /Cr steel and was associated with point defect creation. Since the PC bar has been tempered at temperatures higher than 200 8C, the creation of point defects in the early stage must have taken place soon after the application of fatigue cycles. The creation of point defects in the early stage, contrary to the Si /Cr steel, might be due to the difference in dislocation-configurations. The PC bar was induction heat-treated, and dislocation-configurations might be less stable than that in the Si /Cr steel. The strain-induced creation of point defect is likely to originate in interactions between dislocations, such as annihilation of dislocation dipoles [3] or dragging of intersection jogs [26]. The stability of dislocation substructures might be different between induction and furnace heat-treated steels, but a detailed discussion is beyond the scope of the present study. A consequence of the above considerations is that the creation process of point defects differs between the early and final stages of fatigue. A relevant feature might be the change in TDA curve profiles between the as-fatigued and subsequently annealed specimens as shown in Figs. 5 and 6. When the TDA curve is determined by the effective diffusion constant of hydrogen [27], the curve profile changes when the trapping defects are altered. Annealing in the final stage of fatigue reduced the desorption on the low temperature side of the peak around 100 8C, both in the Si/Cr steel, Fig. 5, and the PC bar subjected to long fatigue cycles, Fig. 6(b). The reduction was due to annealing out of point defects created after a substantial number of fatigue cycles, presumably after evolution of considerable strain localization in areas such as extended slip bands or crack front. The decrease of hydrogen desorption on the lower temperature side implies that the annealed out defects have a low binding energy with
hydrogen. On the other hand, annealing out of point defects, created soon after the application of fatigue cycles in the PC bar, appeared in the middle part of the desorption curve, Fig. 6. The implication is that the binding energy with hydrogen of point defects, presumably vacancies, created in the early stage is slightly higher than that of defects created in the final stage, presumably by coupling with impurity atoms or clustering [28]. Since extended strain localization was not expected in the early stage, reconfiguration and annihilation of dislocations in a wide region might be a possible mechanism of point defect creation. As for the involvement of hydrogen in fatigue failure, Fig. 7 shows a remarkable effect of hydrogen on reducing the fatigue life, while the decrease of the fatigue limit was not significant. Miller suggested [29] that the fatigue limit is characterized by the maximum non-propagating crack size, further crack growth being permitted by the increased stress level. If the controlling process of the fatigue limit is the crack growth overcoming a strong barrier to crack propagation, the implication of Fig. 7, which shows a prominent effect of hydrogen on fatigue life, is that hydrogen drastically reduces the resistance to the once-started crack growth. The decrease in crack growth resistance in the presence of hydrogen was observed previously in a three-point bending test of ferritic steels [30], in which a finite element analysis showed prominent strain localization in front of the advancing crack associated with an increased density of nucleation voids. The fractographic feature shown in Fig. 9(a) suggests microstructural changes associated with the crack advance. The increased defect density in the fractured hydrogen-precharged specimens shown in Fig. 8 also supports the idea that hydrogen is a key factor in reducing fatigue life. However, fatigue failure in the presence of hydrogen occurred much earlier than the stage in which crack growth presumably took place as inferred from the point defect creation in Fig. 3. The relation of hydrogen to crack nucleation may also be involved in the reduction of fatigue life. At the stress level above the fatigue limit, an incipient crack is ready to grow overcoming various barriers. If hydrogen favors the nucleation of an incipient crack, the nucleation as well as the growth of a fatigue crack is expected to reduce the fatigue life. Fatigue crack initiation associated with the formation of well-developed PSBs has been reported [31], but the results in Fig. 7 suggest that hydrogen also affected premature alterations of the substructures in the incubation period, leading to the crack nucleation. In Fig. 7, the effect of hydrogen in reducing the fatigue life was more prominent at lower stress levels. The applied stress amplitude in the present experiment was as low as about one third of the ultimate tensile strength, implying strain localization in very limited areas. In such situa-
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tions, the extent of plastic deformation that interacts with hydrogen must depend on the stress level, resulting in a sensitive effect of hydrogen on the crack nucleation and growth process through interactions with plastic deformation. The effect of hydrogen in increasing vacancy density has been well recognized [15,16]. One mechanism proposed by Essmann concerning the formation of PSBs is the annihilation of dislocation dipoles [3]. It is not definite if his proposed model can be applied to bcc alloys, but vacancy formation must originate in interactions of moving dislocations, which may take place preferentially in an area of high dislocation density. The PSBs and crack front area are sites of substantial strain localization, and the effect of hydrogen in increasing vacancy density could be prominent once a crack is formed as a strain concentrator. However, hydrogen might accelerate the formation of substructures like PSBs in the incubation period before fatigue crack initiation through synergistic interactions between hydrogen and dislocation movement. The decrease in hydrogen content preceding the final stage of increasing hydrogen content suggests reconfigurations of dislocations and the involvement of hydrogen in the process. Interactions between hydrogen and vacancies are expected to affect fatigue resistance in both the crack initiation and growth stages. Vacancy formation during fatigue is consistent with a previous result [17] that showed an increased susceptibility to delayed fracture in the same Si /Cr steel bar that had been given pre-fatigue treatment by a rotational bending fatigue test at a stress amplitude of 640 MPa and a rotational speed of 5000 rpm for 8 /104 cycles, close to the fatigue life, and 4 /104 cycles. Since the increased susceptibility recovered to the original level by subsequently annealing the pre-fatigued specimens at 200 8C, the effect was ascribed to an interaction of hydrogen with vacancies created during fatigue. In that case, if vacancy creation in the pre-fatigued specimens was associated with a definite crack, the onceformed crack would remain after annealing, still favoring the increase in delayed fracture susceptibility. Then, it is reasonable to consider that fatigue damage associated with vacancy creation is not confined to the area in front of an existing fatigue crack.
5. Summary and conclusions Defects created during rotational bending fatigue in high-strength martensitic steels have been detected utilizing hydrogen as a probe. The amount of absorbed hydrogen decreased or was almost constant in the early stage of fatigue, and then turned to an increase after a substantial number of fatigue cycles. Subsequent annealing of fatigued specimens reduced the increased hydro-
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gen in the final stage, suggesting the creation of point defects, presumably vacancies, in that stage. On the other hand, the decrease in the early stage was ascribed to changes in microstructural constitutions, reconfigurations or annihilations of dislocations. In the induction heat-treated PC bar, point defect creation also took place in the early stage, presumably associated with reconfiguration process of dislocations. Hydrogen-precharging drastically reduced fatigue life, but the effect of hydrogen on the fatigue limit was not significant. The defect density in fractured specimens in the presence of hydrogen was higher than in specimens fatigued for similar cycles without hydrogen, indicating a hydrogenassisted increase in defect density. The fracture surface of hydrogen-precharged specimens was smooth without cleavage-like crystallographic features. The reduction of fatigue life at applied stress levels above the fatigue limit and significant creation of point defects in the final stage of the fatigue process imply that hydrogen reduces crack growth resistance coupled with point defects. However, very early fracture in the presence of hydrogen and changes in microstructural constitutions in the early stage suggest that hydrogen is also involved in crack initiation through interactions with substructural changes in the incubation period.
Acknowledgements The present work has been conducted under a contract with the Ministry of Education, Culture, Sports, Science and Technology of Japan. The authors also wish to thank the Netsuren Co. Ltd for supplying the PC steel bars.
References [1] C. Laird, in: R.W. Cahn, P.H. Haasen (Eds.), Physical Metallurgy, fourth ed., Elsevier Science, Amsterdam, 1996, p. 2293. [2] L.P. Karjalainen, Mater. Sci. Eng. 38 (1979) 145. [3] U. Essmann, H. Mughrabi, Phil. Mag. 40 (1979) 731. [4] R.P. Wei, in: J. Fong (Ed.), Fatigue Mechanism, STP 675, ASTM, Philadelphia, 1979, p. 826. [5] D.J. Duquette, Hydrogen Embrittlement and Stress Corrosion Cracking, ASM, 1980, p. 250. [6] R.J. Jacko, D.H. Duquette, in: I.M. Bernstein, A.W. Thompson (Eds.), Hydrogen Effects in Materials, AIME, Warrendale, PA, 1981, p. 477. [7] H.H. Johnson, J.P. Hirth, Metall. Trans. A 7A (1976) 1543. [8] T.J. Marrow, C.A. Hippsley, J.E. King, Acta Metall. Mater. 39 (1991) 1367. [9] T. Sakai, M. Takeda, K. Shiozawa, Y. Ochi, M. Nakajima, T. Nakamura, N. Oguma, in: X.-R. Wu, Z.G. Wang (Eds.), Fatigue’99, vol. I, Higher Education Press, Beijing, 1999, p. 573. [10] Y. Murakami, H. Konishi, K. Takai, Y. Murakami, Tetsu-toHagane 86 (2000) 777. [11] R.A. Oriani, P.H. Josephic, Acta Metall. 22 (1974) 1065. [12] H.K. Birnbaum, P. Sofronis, Mater. Sci. Eng. A 176 (1994) 191.
200 [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23]
M. Nagumo et al. / Materials Science and Engineering A348 (2003) 192 /200 C.D. Beachem, Metall. Trans. 3 (1972) 437. J.P. Hirth, Metall. Trans. A 11A (1980) 861. R.B. McLellan, L. Yang, Acta Metall. Mater. 43 (1995) 2463. Y. Fukai, Y. Kurokawa, H. Hiraoka, J. Jpn Inst. Metals 61 (1997) 663. M. Nagumo, S. Sekiguchi, H. Hayashi, T. Takai, Mater. Sci. Eng. A, in press. M. Nagumo, M. Nakamura, K. Takai, Metallur. Mater. Trans. A 32A (2001) 339. M. Nagumo, K. Ohta, H. Saitoh, Scripta Mater. 40 (1999) 313. W.Y. Choo, J.Y. Lee, Metall. Trans. A 13A (1982) 135. K. Ono, M. Meshii, Acta Metall. Mater. 40 (1992) 1357. H. Shimura, T. Chaya, M. Nagumo, H. Hayashi, I. Ochiai, CAMP ISIJ 15 (2002) 572. S. Suzuki, N. Ishii, Y. Tsuchida, Te´tsu-to-Hagane´ 80 (1994) 855.
[24] M. Aoki, H. Saitoh, M. Mori, Y. Ishida, M. Nagumo, J. Jpn Inst. Metals 58 (1994) 1141. [25] K. Rahka, P. Nenonen, Z.-G. Wang, C. Laird, in: P.O. Kettunen, T.K. Lepisto¨, M.E. Lehtonen (Eds.), Strength of Metals and Alloys, vol. 2, Pergamon Press, Oxford, 1989, p. 767. [26] A.M. Cuitino, M. Ortiz, Acta Mater. 44 (1996) 427. [27] R.A. Oriani, Acta Metall. 18 (1970) 147. [28] A. Vehanen, P. Hautojarvi, J. Johnson, J. Yli-Kauppila, Phys. Rev. B 25 (1982) 762. [29] K.J. Miller, in: A.S. Argon (Ed.), Topics in Fracture and Fatigue, Springer, New York, 1992, p. 309. [30] M. Nagumo, H. Yoshida, Y. Shimomura, T. Kadokura, Mater. Trans. 42 (2001) 132. [31] A. Hunsche, P. Neumann, Acta Metall. 34 (1986) 207.