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Hydrogen embrittlement behavior of high strength rail steels: A comparison between pearlitic and bainitic microstructures
Materials Science & Engineering A 704 (2017) 199–206
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Hydrogen embrittlement behavior of high strength rail steels: A comparison between pearlitic and bainitic microstructures
MARK
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Weijun Hui , Zhibao Xu, Yongjian Zhang, Xiaoli Zhao, Chengwei Shao, Yuqing Weng School of Mechanical, Electronic and Control Engineering, Beijing Jiaotong University, Beijing 100044, PR China
A R T I C L E I N F O
A B S T R A C T
Keywords: Hydrogen embrittlement Microstructures Rail steel Pearlitic steel Bainitic steel
The present study was attempted to evaluate the hydrogen embrittlement (HE) behavior of two high strength rail steels with pearlitic and bainitic microstructures by using slow strain rate test (SSRT) with notched round bar specimens. The results show that the bainitic rail steel with a mixed microstructure of bainite, tempered martensite and M/A constituents exhibits superior mechanical properties in comparison with the pearlitic rail steel. The results of SSRT revealed that the bainitic rail steel is more susceptible to HE than the pearlitic rail steel, which is ascribed mainly to the higher strength level and the microstructural characteristics of the bainitic rail steel. The attempt of re-tempering treatment of the as-received bainitic rail steel exhibits no notable improvement in the resistance to HE at no expense of strength. It is thus suggested that further efforts concerning obtaining low susceptibility to HE besides obtaining excellent combination of strength and toughness should be conducted to guarantee the safety service of the bainitic steel rails.
1. Introduction Rail is one of the important components of railway transportation system which requires being routinely checked and maintained in proper service condition because any premature failure of rail can cause catastrophic accident. Recently, the ever-growing constructions of highspeed railway and the increased axle loads have led to larger wheel/rail contact forces [1]. Therefore, great efforts have been made to enhance the properties of rail materials to improve the performance and to reduce the cost [2–8]. In the last 50 years the railways and rail manufacturers have improved rail performance by increasing its hardness from ~248HB to more than 400HB [6]. Owing to its high values of elastic modulus and superior strength, ductility and wear resistance, nearly all modern rails are made of high C-Mn steels which possess full pearlitic microstructure [4]. The development of pearlitic rail steels has been progressing since the invention of the conventional rail and these efforts are focused primarily on either increasing the C content up to even hypereutectoid level or further refining the pearlitic lamellar structure by either alloying or heat treatment [2,5,9]. However, the strength and hardness level of current pearlitic rail steel has almost reached its theoretical limit and any further improvement becomes extremely difficult [4,10]. Because of their low C content and superior mechanical properties than pearlitic rail steels, bainitic rail steels have attracted significant attentions since bainitic steel “Titan” was first used in the high-speed railway system in
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Corresponding author. E-mail address: [email protected] (W. Hui).
http://dx.doi.org/10.1016/j.msea.2017.08.022 Received 7 March 2017; Received in revised form 5 August 2017; Accepted 5 August 2017 Available online 09 August 2017 0921-5093/ © 2017 Elsevier B.V. All rights reserved.
1980s [3,4,7–12]. This kind of steel is considered as the most potential candidate as the next generation rail steel [11]. It is well known that hydrogen embrittlement (HE), also known as hydrogen-induced delayed fracture (HIDF), has been one of the top issues in the development of high strength steels. The susceptibility to HE generally increases with enhancing steel strength level especially when it reaches over ~1200 MPa [13,14]. As the strength level of bainitic rail steel reaches as high as 1300–1500 MPa grade, a risk to HE susceptibility might exist. In fact, according to the statistical results from railway interests, a series of hydrogen related brittle fracture occurred of the bainitic steel rails and crossings in service, which is one of the reasons that bainitic steel rails and crossings have not been applied globally in the past decades [11,15]. Moreover, rail steels are often adversely affected under severe corrosive environments, for example, near coastal areas. A recent survey in India has shown that approximately one third of all the failures of rails were caused by corrosion [16]. Hydrogen is generally produced on the surface through the corrosion of steels which could enter into the steel and thus cause HE. The presence of high amount of cementite in pearlitic steel renders the structure susceptible to corrosion. Therefore, increasing attentions have also been paid to the HE behavior of high strength pearlitic [16] and bainitic [11,15,17] rail steels. However, the comparison of the HE behavior between high strength pearlitic and bainitic rail steels has not yet been found. In the present study, samples obtained from both pearlitic and
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charging and the results represent the mean of at least three specimens. After SSRT, the notch tensile strengths (σN0 and σNH for the uncharged and hydrogen-charged specimens, respectively), which were defined as the nominal maximum tensile stresses, were obtained. The index of relative susceptibility to HE (HEI) was determined by calculating the relative notch tensile strength loss, which is expressed as:
Table 1 Chemical compositions of the tested rail steels (wt%). Steel
C
Si
Mn
P
S
Cr
Ni
Mo
P B
0.72 0.20
0.25 0.80
1.20 2.00
0.015 0.021
0.011 0.010
— 0.90
— < 0.70
— < 0.50
HEI (%) = (1 − bainitic steel rails used in practice were used to investigate their HE behavior by using slow strain rate test (SSRT). Understanding of such behavior is important for ensuring the safety service of railway transportation and for increasing the lifetime of rails and crossings. It is also beneficial to gain a deeper insight of the relationship between the microstructures and the HE behavior of high strength steels.
σNH ) × 100% σN 0
(1)
TDS was used for the analysis of hydrogen and the tests were carried out within 10 min after completing the hydrogen charging. The specimen was heated from ambient temperature to 800 ℃ at a constant heating rate of 100 ℃/h. 2.3. Microstructural observation and mechanical evaluation
2. Material and methods After standard grinding and polishing, the metallographic specimens were etched in 3% nital solution for microstructural observation by an optical microscope (OP, Zeiss Scope A1) and a scanning electron microscope (SEM, Zeiss EVO 18). A transmission electron microscope (TEM, Hitachi H-800) with energy dispersive X-ray spectrometer (EDS) was used to further study the microstructural characteristics. The operating voltage was 200 kV. Standard chromium trioxide-acetic acid solution was used for the preparation of thin foils in a twin-jet electropolishing apparatus. The volume fraction of retained austenite (RA) was determined via X-rays diffraction (XRD, Rigaku D/MAX 2500) using a Cu Kα radiation operated at 40 kV and 150 mA. Tensile tests were performed on a SUNS 5305 type universal testing machine using a cross-head speed of 1 mm/min at room temperature. The fracture surfaces of the fractured specimens were observed using SEM operated at 20 kV.
2.1. Materials and specimen preparation The samples used in the present study were obtained from steel rails in service. The chemical compositions of the tested rail steels, which were designated as P steel and B steel, are listed in Table 1. Unfortunately the supplier did not provide specific details about the manufacturing process of these tested rails except low-temperature tempering treatment of the B steel rail. Specimens were cut from the head region of rails in the rolling direction as shown in Fig. 1(a). In order to confirm the influence of tempering, one group of the specimens of the B steel was tempered at 280 ℃ for 2 h (designated as B-T steel). Circumferentially notched round bar specimens were used for the SSRT. Fig. 1(b) shows the geometry and dimensions of the notched specimen with notch root radius of 0.15 mm (Kt = 3.2) [13]. Smooth round specimens for tensile test are standard round bars with 5 mm diameter and 25 mm gauge length. Specimens with diameter of 5 mm and length of 25 mm were used to study the hydrogen desorption behavior. Hydrogen was introduced into the SSRT and the thermal desorption spectrometry (TDS) specimens by electrochemical charging in a 0.1 mol/L NaOH aqueous solution at 8 mA/cm2 current density for 72 h.
3. Results 3.1. Microstructure characteristics and mechanical properties Figs. 2–4 present the OP, SEM and TEM microstructures of the tested rail steels, respectively. The P steel has an almost full pearlite microstructure. SEM and TEM observations revealed that there exhibit a large portion of degenerated pearlite (Fig. 3(a) and 4(a)). As seen from Fig. 3(a), the lamellae are generally aligned in the same orientation in one prior austenite grain. The properties of pearlitic steels are mainly determined by the spacing between the ferrite-cementite lamellae, i.e., the strength and the hardness are inversely related to the interlamellar spacing. The average interlamellar spacing determined using intercept method of SEM micrographs is about 300 ± 75 nm for the P steel, which is similar to that for the R260 rail steel
2.2. Measurement of HE susceptibility and hydrogen content To evaluate the hydrogen degradation properties, SSRT was conducted in both the uncharged (as-received condition) and hydrogencharged specimens. SSRT was carried out at room temperature using a WDML-100 kN type uniaxial tensile machine with a constant stroke rate of 0.005 mm/min which is equivalent to a nominal strain rate of 2.1 × 10−6 s−1. The tests were started within 10 min after the hydrogen
Fig. 1. (a) Schematic diagram showing sampling position and orientation in the rail head region and (b) geometry and dimensions (in mm) of the notched specimen with notch root radius of 0.15 mm and Kt = 3.2.
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(a)
(b)
(c)
(d) M
M
Fig. 2. OP micrographs of (a, b) P steel and (c, d) B steel. M refers to martensite.
(a)
(b) thin M/A LP
blocky M/A M
DP
Fig. 3. SEM micrographs of (a) P steel and (b) B steel. LP refers to lamellar pearlite, DP refers to degenerated pearlite, M/A refers to martensite/austenite constituent and M refers to martensite.
stability [8]. Moreover, there are also two types of RA, i.e., filmy and blocky RA, as shown in Fig. 4(d) and (e). The presences of fine ε-carbides within the martensitic laths (Fig. 4(f)) as well as a certain amount of RA reconfirmed that the steel rail had experienced a low-temperature tempering treatment after hot-rolling. The mechanical properties of the tested steels are summarized in Table 2. As expected, the B steel exhibits superior mechanical properties in comparison with the P steel. There is an increase of about 30% in the tensile strength, 85% in the yield strength and 24% in the reduction of area from the pearlitic rail steel to the bainitic rail steel, though there is a little decrease in the total elongation. The higher strength of the B steel could be related to the mixed microstructure of bainite and
(282 ± 11 nm), whereas it is much higher than that for the R370CrHT rail steel (80 ± 3 nm) [18]. As for the B steel, its microstructure is more complex than that of the P steel, which composes of bainite, martensite and martensite/ austenite (M/A) constituents. The volume fraction of martensite using image analysis is about 16.2% and the volume fraction of RA obtained from the XRD analysis is about 5.2% (Fig. 5 and Table 2). As can be seen from Figs. 3(b), 4(b) and 4(c), there exist two types of M/A constituent based on its size and morphology, i.e., large blocky M/A island mainly near the prior austenite grain boundaries and thin M/A strip between the bainitic ferrite (BF) laths. This kind of M/A characteristics was also found in a low C bainitic rail steel and exhibited different thermal 201
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(a)
(b)
blocky M/A LP DP
(d)
(c)
thin M/A BF flimy RA
(f)
(e)
carbide
blocky RA
Fig. 4. TEM micrographs of the tested (a) P steel and (b–f) B steel, showing (a) lamellar pearlite (LP) and degenerated pearlite (DP) of the P steel, (b) blocky M/A, (c) thin M/A, (d) filmy RA, (e) blocky RA and (f) carbide within martensitic laths of the B steel (insert: selected area electron diffraction pattern of the blocky RA in (e)).
lower than that of the P steel. The integration of these curves shows that the hydrogen contents of the B steel both before (C0) and after (CH) hydrogen charging are slightly lower than those of the P steel as shown in Table 3.
martensite, while the presence of a certain amount of RA has great contribution to its ductility [8,19]. Large interlamellar spacing (about 300 ± 75 nm) is detrimental to the mechanical properties of the P steel. Moreover, the presence of degenerated pearlite also decreases the steel's strength and hardness while increases its ductility.
3.3. Hydrogen embrittlement behavior 3.2. Hydrogen absorption behavior Fig. 7 gives the typical SSRT tensile stress versus displacement curves of both the uncharged and hydrogen-charged specimens. It is obvious that the notched specimens exhibited an early fracture before obvious yielding and necking, indicating the significant influence of notch on the tensile behavior of the tested steels. It is well known that the susceptibility of given steel to HE is the interaction of environment (hydrogen), stress and material, and this behavior could be intensified
Fig. 6 shows the hydrogen desorption curves of the tested rail steels. It is obvious that for the uncharged specimens the curves of both steels exhibit slight high-temperature desorption peak at around 435 ℃, while for the hydrogen-charged specimens the curves of both steels exhibit remarkable low-temperature desorption peak at around 155 ℃. It is also shown that the height of the low-temperature peak of the B steel is 202
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fracture of quasi-cleavage and ductile dimples in the crack initiation region, as shown in Fig. 8(c), while the fracture surface of the hydrogen-charged specimen is dominated by brittle intergranular failure with few ductile tearing, as shown in Fig. 8(d). This intergranular nature of failure suggests that the prior austenite grain boundaries of the B steel were severely affected by the enrichment of hydrogen. Compared to that of the P steel, hydrogen has significantly changed the fracture mode of the B steel, as also reflected by the much higher HEI, though both steels possess very close hydrogen contents (Table 3). The fracture mode of the uncharged B steel is similar to that of the lowtemperature tempered bainitic rail steel [8]. This result along with the presence of tempered martensite as shown in Fig. 4(f) indicates again that the tested B steel rail had experienced tempering treatment. 4. Discussion Fig. 5. X-ray diffraction patterns of the tested rail steels.
As the tensile strength of the P steel is only 1050 MPa compared to that of 1360 MPa of the B steel, the HE susceptibility of the P steel is thus significant lower than that of the B steel. It should be noted that since the strength of a steel depends upon both its chemical composition and microstructure, it is rather difficult to independently investigate the effects of strength from microstructure and alloying. The results of a systematic study examined the effects of microstructures, alloying elements, and strength on the susceptibility of fastener grade steels to HE showed that the overriding factor contributing to the HE susceptibility of the steels was strength though microstructural alteration resulted in some improvement in the resistance to HE [23]. It was also found that the degree of susceptibility to HE of the pearlitic especially degenerated pearlitic microstructure is notably lower than those of the bainitic and tempered martensitic microstructures at similar strength level [23,24]. Therefore, it is regarded that both strength and microstructural characteristics play important roles in determining the resistance to HE of the P steel. As mentioned above, there exists as high as ~16 vol% of martensite for the B steel. It is well known that untempered brittle martensite significantly deteriorates the toughness as well as the resistance to HE of high strength steels. Numerous studies have reported that improvement in the resistance of martensite to HE could be achieved by increasing tempering temperature [23,25–27]. It should be noted that tempering usually results in a corresponding loss of strength while it changes the microstructure for conventional low-alloy structural steels such as AISI 4140 and 4340. One of the authors’ previous investigation on a secondary hardening Cr-Mo-V high strength steel confirmed that its resistance to stress corrosion cracking in 3.5% aqueous solution could be significantly enhanced with an increase in tempering temperature even without losing its strength [26]. As the heat treatment of the B steel rail was outsourced, the specific tempering temperature used is not known though it was required that suitable tempering treatment should be performed. The presence of fine ε-carbides within the martensitic laths (Fig. 4(f)) as well as the SSRT fracture mode of the uncharged B steel suggests that the steel rail might have experienced an incorrect lowtemperature tempering treatment after hot-rolling. To confirm this assumption, a group of the B steel was tempered at 280 ℃ for 2 h (the B-T steel) because excellent combination of strength and toughness could be obtained at this temperature [8], and the results are also presented in Table 3. Unfortunately, it is obvious that the B-T steel exhibited no
in the presence of stress risers. As for the notched specimen, local stress concentration will be induced at the notch during SSRT deformation. It has been revealed by numerical simulation [20] and several investigations [21,22] that notch accelerates the diffusion of hydrogen from other part to the notch root, and therefore helps to enhance the mobility of hydrogen in the specimens during loading to make it more effective as an embrittlement agent. Therefore, it is suggested that notched specimen is suitable to reflect the HE susceptibility of the tested high strength rail steels. As seen from Table 3, the notch tensile strength of the B steel without hydrogen charging (σN0) is much higher (by 412 MPa) than that of the P steel mainly due to the much higher strength of the former, whereas the notch tensile strength of the B steel with hydrogen charging (σNH) is much lower (by 451 MPa) than that of the P steel. As a result, the HEI of the B steel (60.6%) is much higher than that of the P steel (26.5%). These results reveal that the B steel exhibits remarkably lower resistance against HE and thus higher susceptibility to HE than that of the P steel. 3.4. Hydrogen embrittlement fracture surface characteristics A detailed fractographic analysis of both the uncharged and hydrogen-charged specimens was performed after SSRT. Fig. 8 shows the fracture surfaces of the SSRT notched specimens. As seen from Fig. 8(a) and (b), both the uncharged and hydrogen-charged specimens of the P steel exhibit primarily cleavage fracture with a few ductile dimples in the crack initiation regions, while there are an increased number of secondary cracks in the hydrogen-charged specimen. Similar result of cleavage type fracture was also observed for pearlitic microstructure either tested in air or tested in cathodically charged salt solution [23]. Despite the similarity in the fracture mode between the uncharged and hydrogen-charged specimens, the crack initiation stress and/or the cohesiveness and energy required to grow the cleavage crack must have been reduced by hydrogen. This result indicates that hydrogen charging did not noticeably change the fracture mode of the P steel except an increase of the amount of secondary cracks, which thus exhibits a relatively lower susceptibility to HE, as also shown in Table 3. As for the B steel, the uncharged specimen exhibits a mixed mode Table 2 Mechanical properties and microstructural parameters of the tested rail steels. Steel
Tensile strength Rm (MPa)
Yield strength RP0.2 (MPa)
Elongation A (%)
Reduction of area Z (%)
Hardness HV5
RA content (vol%)
Martensite content (vol%)
P B
1050 1360
530 980
19.0 14.6
45.5 56.6
264 ± 6 429 ± 4
0 5.2 ± 0.1
– 16.2 ± 2.7
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Fig. 6. Hydrogen desorption rate curves of the uncharged and hydrogen-charged specimens. (a) P steel; (b) B steel.
analysis is required to clarify this issue, and it is important that great efforts should be conducted further to enhance the resistance to HE to guarantee the safety service of the tested bainitic steel rails besides to obtain excellent combination of strength and toughness.
Table 3 Summary of the SSRT and TDS results of the tested steels. Steel
Rm (MPa)
σN0 (MPa)
σNH (MPa)
HEI (%)
C0 (ppm)
CH (ppm)
P B B-T
1050 1360 1340
1799 2211 2251
1323 872 891
26.5 60.6 60.4
0.16 0.12 –
0.75 0.61 –
5. Conclusions (1) Microstructural examinations revealed that the P steel has full pearlitic microstructure with a large portion of degenerated pearlite, while the B steel exhibits a mixed microstructure composing of bainite, tempered martensite and M/A constituents. The B steel exhibits superior mechanical properties in comparison with the P steel. (2) The notch tensile strength of the B steel without hydrogen charging is much higher than that of the P steel, while the notch tensile strength of the B steel with hydrogen pre-charged under the same condition is much lower than that of the P steel. Accordingly, the susceptibility to HE of the former is much higher than that of the latter, which is ascribed mainly to the higher strength level and the microstructural characteristics of the former. (3) The tempering treatment at 280 ℃ for 2 h of the as-received B steel exhibits no notable improvement in the resistance to HE at no expense of strength. It is thus suggested that further investigation mainly concerning chemical composition optimization and suitable tempering treatment after hot-rolling should be conducted to obtain lower susceptibility to HE and thus to guarantee safety service of the bainitic steel rails.
noticeable improvement of the resistance to HE compared to the B steel at no expense of strength. Besides the influence of martensite, the presence of large size blocky RA of the B steel may also deteriorate the resistance to HE for it might transform into fresh untempered martensite at the early beginning of SSRT due to its lower mechanical stability [28,29]. Studies have revealed that unstable blocky RA decomposed while filmy RA retained due to its high thermal and mechanical stability during appropriate tempering treatment [8,30]. Owing to the very high solubility and low diffusivity of hydrogen in it, austenite acts as a sink of hydrogen [31], and thus the hydrogen concentration in austenite is regarded to be much higher than in other phase. From this viewpoint, the hydrogen-enriched transformed martensite is even more detrimental to HE [32]. Moreover, the hydrogen desorption behavior of the B steel is similar to that of the P steel (Fig. 6) and the hydrogen content for the hydrogen-charged B steel is even slightly lower than that for the hydrogen-charged P steel (Table 3). Thus, it is reasonable to suggest that hydrogen content is not the main reason for the higher HE susceptibility of the B steel. Therefore, further detail and systematic
Fig. 7. SSRT curves of the notched specimens of the tested rail steels before and after hydrogen charging. (a) P steel; (b) B steel.
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Hydrogen embrittlement behavior of high strength rail steels: A comparison between pearlitic and bainitic microstructures
MARK
⁎
Weijun Hui , Zhibao Xu, Yongjian Zhang, Xiaoli Zhao, Chengwei Shao, Yuqing Weng School of Mechanical, Electronic and Control Engineering, Beijing Jiaotong University, Beijing 100044, PR China
A R T I C L E I N F O
A B S T R A C T
Keywords: Hydrogen embrittlement Microstructures Rail steel Pearlitic steel Bainitic steel
The present study was attempted to evaluate the hydrogen embrittlement (HE) behavior of two high strength rail steels with pearlitic and bainitic microstructures by using slow strain rate test (SSRT) with notched round bar specimens. The results show that the bainitic rail steel with a mixed microstructure of bainite, tempered martensite and M/A constituents exhibits superior mechanical properties in comparison with the pearlitic rail steel. The results of SSRT revealed that the bainitic rail steel is more susceptible to HE than the pearlitic rail steel, which is ascribed mainly to the higher strength level and the microstructural characteristics of the bainitic rail steel. The attempt of re-tempering treatment of the as-received bainitic rail steel exhibits no notable improvement in the resistance to HE at no expense of strength. It is thus suggested that further efforts concerning obtaining low susceptibility to HE besides obtaining excellent combination of strength and toughness should be conducted to guarantee the safety service of the bainitic steel rails.
1. Introduction Rail is one of the important components of railway transportation system which requires being routinely checked and maintained in proper service condition because any premature failure of rail can cause catastrophic accident. Recently, the ever-growing constructions of highspeed railway and the increased axle loads have led to larger wheel/rail contact forces [1]. Therefore, great efforts have been made to enhance the properties of rail materials to improve the performance and to reduce the cost [2–8]. In the last 50 years the railways and rail manufacturers have improved rail performance by increasing its hardness from ~248HB to more than 400HB [6]. Owing to its high values of elastic modulus and superior strength, ductility and wear resistance, nearly all modern rails are made of high C-Mn steels which possess full pearlitic microstructure [4]. The development of pearlitic rail steels has been progressing since the invention of the conventional rail and these efforts are focused primarily on either increasing the C content up to even hypereutectoid level or further refining the pearlitic lamellar structure by either alloying or heat treatment [2,5,9]. However, the strength and hardness level of current pearlitic rail steel has almost reached its theoretical limit and any further improvement becomes extremely difficult [4,10]. Because of their low C content and superior mechanical properties than pearlitic rail steels, bainitic rail steels have attracted significant attentions since bainitic steel “Titan” was first used in the high-speed railway system in
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Corresponding author. E-mail address: [email protected] (W. Hui).
http://dx.doi.org/10.1016/j.msea.2017.08.022 Received 7 March 2017; Received in revised form 5 August 2017; Accepted 5 August 2017 Available online 09 August 2017 0921-5093/ © 2017 Elsevier B.V. All rights reserved.
1980s [3,4,7–12]. This kind of steel is considered as the most potential candidate as the next generation rail steel [11]. It is well known that hydrogen embrittlement (HE), also known as hydrogen-induced delayed fracture (HIDF), has been one of the top issues in the development of high strength steels. The susceptibility to HE generally increases with enhancing steel strength level especially when it reaches over ~1200 MPa [13,14]. As the strength level of bainitic rail steel reaches as high as 1300–1500 MPa grade, a risk to HE susceptibility might exist. In fact, according to the statistical results from railway interests, a series of hydrogen related brittle fracture occurred of the bainitic steel rails and crossings in service, which is one of the reasons that bainitic steel rails and crossings have not been applied globally in the past decades [11,15]. Moreover, rail steels are often adversely affected under severe corrosive environments, for example, near coastal areas. A recent survey in India has shown that approximately one third of all the failures of rails were caused by corrosion [16]. Hydrogen is generally produced on the surface through the corrosion of steels which could enter into the steel and thus cause HE. The presence of high amount of cementite in pearlitic steel renders the structure susceptible to corrosion. Therefore, increasing attentions have also been paid to the HE behavior of high strength pearlitic [16] and bainitic [11,15,17] rail steels. However, the comparison of the HE behavior between high strength pearlitic and bainitic rail steels has not yet been found. In the present study, samples obtained from both pearlitic and
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charging and the results represent the mean of at least three specimens. After SSRT, the notch tensile strengths (σN0 and σNH for the uncharged and hydrogen-charged specimens, respectively), which were defined as the nominal maximum tensile stresses, were obtained. The index of relative susceptibility to HE (HEI) was determined by calculating the relative notch tensile strength loss, which is expressed as:
Table 1 Chemical compositions of the tested rail steels (wt%). Steel
C
Si
Mn
P
S
Cr
Ni
Mo
P B
0.72 0.20
0.25 0.80
1.20 2.00
0.015 0.021
0.011 0.010
— 0.90
— < 0.70
— < 0.50
HEI (%) = (1 − bainitic steel rails used in practice were used to investigate their HE behavior by using slow strain rate test (SSRT). Understanding of such behavior is important for ensuring the safety service of railway transportation and for increasing the lifetime of rails and crossings. It is also beneficial to gain a deeper insight of the relationship between the microstructures and the HE behavior of high strength steels.
σNH ) × 100% σN 0
(1)
TDS was used for the analysis of hydrogen and the tests were carried out within 10 min after completing the hydrogen charging. The specimen was heated from ambient temperature to 800 ℃ at a constant heating rate of 100 ℃/h. 2.3. Microstructural observation and mechanical evaluation
2. Material and methods After standard grinding and polishing, the metallographic specimens were etched in 3% nital solution for microstructural observation by an optical microscope (OP, Zeiss Scope A1) and a scanning electron microscope (SEM, Zeiss EVO 18). A transmission electron microscope (TEM, Hitachi H-800) with energy dispersive X-ray spectrometer (EDS) was used to further study the microstructural characteristics. The operating voltage was 200 kV. Standard chromium trioxide-acetic acid solution was used for the preparation of thin foils in a twin-jet electropolishing apparatus. The volume fraction of retained austenite (RA) was determined via X-rays diffraction (XRD, Rigaku D/MAX 2500) using a Cu Kα radiation operated at 40 kV and 150 mA. Tensile tests were performed on a SUNS 5305 type universal testing machine using a cross-head speed of 1 mm/min at room temperature. The fracture surfaces of the fractured specimens were observed using SEM operated at 20 kV.
2.1. Materials and specimen preparation The samples used in the present study were obtained from steel rails in service. The chemical compositions of the tested rail steels, which were designated as P steel and B steel, are listed in Table 1. Unfortunately the supplier did not provide specific details about the manufacturing process of these tested rails except low-temperature tempering treatment of the B steel rail. Specimens were cut from the head region of rails in the rolling direction as shown in Fig. 1(a). In order to confirm the influence of tempering, one group of the specimens of the B steel was tempered at 280 ℃ for 2 h (designated as B-T steel). Circumferentially notched round bar specimens were used for the SSRT. Fig. 1(b) shows the geometry and dimensions of the notched specimen with notch root radius of 0.15 mm (Kt = 3.2) [13]. Smooth round specimens for tensile test are standard round bars with 5 mm diameter and 25 mm gauge length. Specimens with diameter of 5 mm and length of 25 mm were used to study the hydrogen desorption behavior. Hydrogen was introduced into the SSRT and the thermal desorption spectrometry (TDS) specimens by electrochemical charging in a 0.1 mol/L NaOH aqueous solution at 8 mA/cm2 current density for 72 h.
3. Results 3.1. Microstructure characteristics and mechanical properties Figs. 2–4 present the OP, SEM and TEM microstructures of the tested rail steels, respectively. The P steel has an almost full pearlite microstructure. SEM and TEM observations revealed that there exhibit a large portion of degenerated pearlite (Fig. 3(a) and 4(a)). As seen from Fig. 3(a), the lamellae are generally aligned in the same orientation in one prior austenite grain. The properties of pearlitic steels are mainly determined by the spacing between the ferrite-cementite lamellae, i.e., the strength and the hardness are inversely related to the interlamellar spacing. The average interlamellar spacing determined using intercept method of SEM micrographs is about 300 ± 75 nm for the P steel, which is similar to that for the R260 rail steel
2.2. Measurement of HE susceptibility and hydrogen content To evaluate the hydrogen degradation properties, SSRT was conducted in both the uncharged (as-received condition) and hydrogencharged specimens. SSRT was carried out at room temperature using a WDML-100 kN type uniaxial tensile machine with a constant stroke rate of 0.005 mm/min which is equivalent to a nominal strain rate of 2.1 × 10−6 s−1. The tests were started within 10 min after the hydrogen
Fig. 1. (a) Schematic diagram showing sampling position and orientation in the rail head region and (b) geometry and dimensions (in mm) of the notched specimen with notch root radius of 0.15 mm and Kt = 3.2.
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(a)
(b)
(c)
(d) M
M
Fig. 2. OP micrographs of (a, b) P steel and (c, d) B steel. M refers to martensite.
(a)
(b) thin M/A LP
blocky M/A M
DP
Fig. 3. SEM micrographs of (a) P steel and (b) B steel. LP refers to lamellar pearlite, DP refers to degenerated pearlite, M/A refers to martensite/austenite constituent and M refers to martensite.
stability [8]. Moreover, there are also two types of RA, i.e., filmy and blocky RA, as shown in Fig. 4(d) and (e). The presences of fine ε-carbides within the martensitic laths (Fig. 4(f)) as well as a certain amount of RA reconfirmed that the steel rail had experienced a low-temperature tempering treatment after hot-rolling. The mechanical properties of the tested steels are summarized in Table 2. As expected, the B steel exhibits superior mechanical properties in comparison with the P steel. There is an increase of about 30% in the tensile strength, 85% in the yield strength and 24% in the reduction of area from the pearlitic rail steel to the bainitic rail steel, though there is a little decrease in the total elongation. The higher strength of the B steel could be related to the mixed microstructure of bainite and
(282 ± 11 nm), whereas it is much higher than that for the R370CrHT rail steel (80 ± 3 nm) [18]. As for the B steel, its microstructure is more complex than that of the P steel, which composes of bainite, martensite and martensite/ austenite (M/A) constituents. The volume fraction of martensite using image analysis is about 16.2% and the volume fraction of RA obtained from the XRD analysis is about 5.2% (Fig. 5 and Table 2). As can be seen from Figs. 3(b), 4(b) and 4(c), there exist two types of M/A constituent based on its size and morphology, i.e., large blocky M/A island mainly near the prior austenite grain boundaries and thin M/A strip between the bainitic ferrite (BF) laths. This kind of M/A characteristics was also found in a low C bainitic rail steel and exhibited different thermal 201
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(a)
(b)
blocky M/A LP DP
(d)
(c)
thin M/A BF flimy RA
(f)
(e)
carbide
blocky RA
Fig. 4. TEM micrographs of the tested (a) P steel and (b–f) B steel, showing (a) lamellar pearlite (LP) and degenerated pearlite (DP) of the P steel, (b) blocky M/A, (c) thin M/A, (d) filmy RA, (e) blocky RA and (f) carbide within martensitic laths of the B steel (insert: selected area electron diffraction pattern of the blocky RA in (e)).
lower than that of the P steel. The integration of these curves shows that the hydrogen contents of the B steel both before (C0) and after (CH) hydrogen charging are slightly lower than those of the P steel as shown in Table 3.
martensite, while the presence of a certain amount of RA has great contribution to its ductility [8,19]. Large interlamellar spacing (about 300 ± 75 nm) is detrimental to the mechanical properties of the P steel. Moreover, the presence of degenerated pearlite also decreases the steel's strength and hardness while increases its ductility.
3.3. Hydrogen embrittlement behavior 3.2. Hydrogen absorption behavior Fig. 7 gives the typical SSRT tensile stress versus displacement curves of both the uncharged and hydrogen-charged specimens. It is obvious that the notched specimens exhibited an early fracture before obvious yielding and necking, indicating the significant influence of notch on the tensile behavior of the tested steels. It is well known that the susceptibility of given steel to HE is the interaction of environment (hydrogen), stress and material, and this behavior could be intensified
Fig. 6 shows the hydrogen desorption curves of the tested rail steels. It is obvious that for the uncharged specimens the curves of both steels exhibit slight high-temperature desorption peak at around 435 ℃, while for the hydrogen-charged specimens the curves of both steels exhibit remarkable low-temperature desorption peak at around 155 ℃. It is also shown that the height of the low-temperature peak of the B steel is 202
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fracture of quasi-cleavage and ductile dimples in the crack initiation region, as shown in Fig. 8(c), while the fracture surface of the hydrogen-charged specimen is dominated by brittle intergranular failure with few ductile tearing, as shown in Fig. 8(d). This intergranular nature of failure suggests that the prior austenite grain boundaries of the B steel were severely affected by the enrichment of hydrogen. Compared to that of the P steel, hydrogen has significantly changed the fracture mode of the B steel, as also reflected by the much higher HEI, though both steels possess very close hydrogen contents (Table 3). The fracture mode of the uncharged B steel is similar to that of the lowtemperature tempered bainitic rail steel [8]. This result along with the presence of tempered martensite as shown in Fig. 4(f) indicates again that the tested B steel rail had experienced tempering treatment. 4. Discussion Fig. 5. X-ray diffraction patterns of the tested rail steels.
As the tensile strength of the P steel is only 1050 MPa compared to that of 1360 MPa of the B steel, the HE susceptibility of the P steel is thus significant lower than that of the B steel. It should be noted that since the strength of a steel depends upon both its chemical composition and microstructure, it is rather difficult to independently investigate the effects of strength from microstructure and alloying. The results of a systematic study examined the effects of microstructures, alloying elements, and strength on the susceptibility of fastener grade steels to HE showed that the overriding factor contributing to the HE susceptibility of the steels was strength though microstructural alteration resulted in some improvement in the resistance to HE [23]. It was also found that the degree of susceptibility to HE of the pearlitic especially degenerated pearlitic microstructure is notably lower than those of the bainitic and tempered martensitic microstructures at similar strength level [23,24]. Therefore, it is regarded that both strength and microstructural characteristics play important roles in determining the resistance to HE of the P steel. As mentioned above, there exists as high as ~16 vol% of martensite for the B steel. It is well known that untempered brittle martensite significantly deteriorates the toughness as well as the resistance to HE of high strength steels. Numerous studies have reported that improvement in the resistance of martensite to HE could be achieved by increasing tempering temperature [23,25–27]. It should be noted that tempering usually results in a corresponding loss of strength while it changes the microstructure for conventional low-alloy structural steels such as AISI 4140 and 4340. One of the authors’ previous investigation on a secondary hardening Cr-Mo-V high strength steel confirmed that its resistance to stress corrosion cracking in 3.5% aqueous solution could be significantly enhanced with an increase in tempering temperature even without losing its strength [26]. As the heat treatment of the B steel rail was outsourced, the specific tempering temperature used is not known though it was required that suitable tempering treatment should be performed. The presence of fine ε-carbides within the martensitic laths (Fig. 4(f)) as well as the SSRT fracture mode of the uncharged B steel suggests that the steel rail might have experienced an incorrect lowtemperature tempering treatment after hot-rolling. To confirm this assumption, a group of the B steel was tempered at 280 ℃ for 2 h (the B-T steel) because excellent combination of strength and toughness could be obtained at this temperature [8], and the results are also presented in Table 3. Unfortunately, it is obvious that the B-T steel exhibited no
in the presence of stress risers. As for the notched specimen, local stress concentration will be induced at the notch during SSRT deformation. It has been revealed by numerical simulation [20] and several investigations [21,22] that notch accelerates the diffusion of hydrogen from other part to the notch root, and therefore helps to enhance the mobility of hydrogen in the specimens during loading to make it more effective as an embrittlement agent. Therefore, it is suggested that notched specimen is suitable to reflect the HE susceptibility of the tested high strength rail steels. As seen from Table 3, the notch tensile strength of the B steel without hydrogen charging (σN0) is much higher (by 412 MPa) than that of the P steel mainly due to the much higher strength of the former, whereas the notch tensile strength of the B steel with hydrogen charging (σNH) is much lower (by 451 MPa) than that of the P steel. As a result, the HEI of the B steel (60.6%) is much higher than that of the P steel (26.5%). These results reveal that the B steel exhibits remarkably lower resistance against HE and thus higher susceptibility to HE than that of the P steel. 3.4. Hydrogen embrittlement fracture surface characteristics A detailed fractographic analysis of both the uncharged and hydrogen-charged specimens was performed after SSRT. Fig. 8 shows the fracture surfaces of the SSRT notched specimens. As seen from Fig. 8(a) and (b), both the uncharged and hydrogen-charged specimens of the P steel exhibit primarily cleavage fracture with a few ductile dimples in the crack initiation regions, while there are an increased number of secondary cracks in the hydrogen-charged specimen. Similar result of cleavage type fracture was also observed for pearlitic microstructure either tested in air or tested in cathodically charged salt solution [23]. Despite the similarity in the fracture mode between the uncharged and hydrogen-charged specimens, the crack initiation stress and/or the cohesiveness and energy required to grow the cleavage crack must have been reduced by hydrogen. This result indicates that hydrogen charging did not noticeably change the fracture mode of the P steel except an increase of the amount of secondary cracks, which thus exhibits a relatively lower susceptibility to HE, as also shown in Table 3. As for the B steel, the uncharged specimen exhibits a mixed mode Table 2 Mechanical properties and microstructural parameters of the tested rail steels. Steel
Tensile strength Rm (MPa)
Yield strength RP0.2 (MPa)
Elongation A (%)
Reduction of area Z (%)
Hardness HV5
RA content (vol%)
Martensite content (vol%)
P B
1050 1360
530 980
19.0 14.6
45.5 56.6
264 ± 6 429 ± 4
0 5.2 ± 0.1
– 16.2 ± 2.7
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Fig. 6. Hydrogen desorption rate curves of the uncharged and hydrogen-charged specimens. (a) P steel; (b) B steel.
analysis is required to clarify this issue, and it is important that great efforts should be conducted further to enhance the resistance to HE to guarantee the safety service of the tested bainitic steel rails besides to obtain excellent combination of strength and toughness.
Table 3 Summary of the SSRT and TDS results of the tested steels. Steel
Rm (MPa)
σN0 (MPa)
σNH (MPa)
HEI (%)
C0 (ppm)
CH (ppm)
P B B-T
1050 1360 1340
1799 2211 2251
1323 872 891
26.5 60.6 60.4
0.16 0.12 –
0.75 0.61 –
5. Conclusions (1) Microstructural examinations revealed that the P steel has full pearlitic microstructure with a large portion of degenerated pearlite, while the B steel exhibits a mixed microstructure composing of bainite, tempered martensite and M/A constituents. The B steel exhibits superior mechanical properties in comparison with the P steel. (2) The notch tensile strength of the B steel without hydrogen charging is much higher than that of the P steel, while the notch tensile strength of the B steel with hydrogen pre-charged under the same condition is much lower than that of the P steel. Accordingly, the susceptibility to HE of the former is much higher than that of the latter, which is ascribed mainly to the higher strength level and the microstructural characteristics of the former. (3) The tempering treatment at 280 ℃ for 2 h of the as-received B steel exhibits no notable improvement in the resistance to HE at no expense of strength. It is thus suggested that further investigation mainly concerning chemical composition optimization and suitable tempering treatment after hot-rolling should be conducted to obtain lower susceptibility to HE and thus to guarantee safety service of the bainitic steel rails.
noticeable improvement of the resistance to HE compared to the B steel at no expense of strength. Besides the influence of martensite, the presence of large size blocky RA of the B steel may also deteriorate the resistance to HE for it might transform into fresh untempered martensite at the early beginning of SSRT due to its lower mechanical stability [28,29]. Studies have revealed that unstable blocky RA decomposed while filmy RA retained due to its high thermal and mechanical stability during appropriate tempering treatment [8,30]. Owing to the very high solubility and low diffusivity of hydrogen in it, austenite acts as a sink of hydrogen [31], and thus the hydrogen concentration in austenite is regarded to be much higher than in other phase. From this viewpoint, the hydrogen-enriched transformed martensite is even more detrimental to HE [32]. Moreover, the hydrogen desorption behavior of the B steel is similar to that of the P steel (Fig. 6) and the hydrogen content for the hydrogen-charged B steel is even slightly lower than that for the hydrogen-charged P steel (Table 3). Thus, it is reasonable to suggest that hydrogen content is not the main reason for the higher HE susceptibility of the B steel. Therefore, further detail and systematic
Fig. 7. SSRT curves of the notched specimens of the tested rail steels before and after hydrogen charging. (a) P steel; (b) B steel.
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