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Improved hydrogen embrittlement resistance after quenching–tempering treatment for a Cr-Mo-V high strength steel
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Improved hydrogen embrittlement resistance after quenchingetempering treatment for a Cr-Mo-V high strength steel Yafei Wang*, Songyan Hu, Yun Li, Guangxu Cheng School of Chemical Engineering and Technology, Xi’an Jiaotong University, Xi’an 710049, China
highlights Decomposition, precipitation and coarsening of carbides confirmed after QT treatment. Beneficial effect of QT-treatment on HE resistance. Decreased reversible hydrogen traps responsible for improved HE resistance.
article info
abstract
Article history:
The effect of quenching-tempering (QT) treatment on the hydrogen embrittlement (HE)
Received 14 August 2019
resistance of a reactor pressure vessel steel was studied. Decomposition of M3C/VC car-
Received in revised form
bides and precipitation of M7C3 carbides were confirmed by transmission electron micro-
17 September 2019
scopy and atom probe tomography observations. Tensile tests showed that HE sensitivity
Accepted 19 September 2019
decreased to a negligible level after QT treatment. The improvement of HE resistance was
Available online 11 October 2019
mainly attributed to the decreased number of M3C carbides which act as the reversible trapping sites for hydrogen. This was supported by the decreased concentration of
Keywords:
reversible hydrogen as measured by thermal desorption spectroscopy. The amount of
Steel
irreversible hydrogen (probably trapped at VC carbides) also decreased, which is however
Heat treatment
not considered responsible for the HE improvement.
Carbides
© 2019 Hydrogen Energy Publications LLC. Published by Elsevier Ltd. All rights reserved.
Hydrogen embrittlement
Introduction Quenching-tempering (QT) treatment is an important heat treatment procedure during the fabrication of hydrogenation reactors, which operate at elevated temperature and high hydrogen partial pressure. Moreover, high temperature hydrogen attack (HTHA) is the main failure mode for these equipment and components that work in a high temperature hydrogen environment. Significant research efforts have
been devoted in the past decades on the study of HTHA mechanism and prevention strategies. The Nelson curves have been extensively used in the industry to determine the operation limit of steels for hydrogen service at elevated temperatures and pressures, according to the American Petroleum Institute (API) Recommended Practice 941 [1]. When the temperature and pressure are controlled below the operation limit, the HTHA is not considered a problem, although the exact values of operation limit for some steels have been revised from time to time based on the experiences
* Corresponding author. E-mail address: [email protected] (Y. Wang). https://doi.org/10.1016/j.ijhydene.2019.09.142 0360-3199/© 2019 Hydrogen Energy Publications LLC. Published by Elsevier Ltd. All rights reserved.
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from some cases of accidental failures (below the operation limit) due to HTHA [2]. It is well-recognized that the carbides in steel can reduce the amount of carbon that is available to react with hydrogen to form methane, which is the key pathway of HTHA. Therefore, the carbide-rich 2.25Cr-1Mo-0.25 V steel has been widely accepted as one of the optimal materials for many components working under high temperature hydrogen environment. According to the Nelson curves, the temperature limit for 2.25Cr-1Mo-0.25 V steel is 510 C when the hydrogen partial pressure is higher than 12.5 MPa, and the temperature limit increases linearly with the further decreasing of hydrogen pressure. The latest version of Nelson curve for 2.25Cr-1Mo-0.25 V steel is based on the data obtained from 10000 þ hours of laboratory tests. Assuming that this curve is accurate enough, the HTHA is not a major concern for hydrogenation reactors made of this material, considering that many reactors operate even far below the temperature limit. In this way, some other failure modes such as hydrogen embrittlement (HE) or hydrogen induced cracking (HIC) may become the important factors that need to be considered for the safe operation of reactors. On the other hand, some cases of failures have been reported during the fabrication of reactors, mostly during welding or heat treatments [3e5]. In these cases, the adsorbed hydrogen, either from the high temperature hydrogen gas in service or from the water molecules in atmosphere during fabrication processes, can play a critical role. The interaction of hydrogen with local microstructures may become the controlling factor that determines the final failures. Therefore, the HE of 2.25Cr-1Mo and 2.25Cr1Mo-0.25 V steel has been reported in some studies [6e9]. However, the hydrogen enhanced delayed fracture is influenced by many factors that take place in different scales, which make the failure mechanism complicated. Thus it is usually difficult to determine the underlying controlling factors by using a single method or methods in similar scales, and subtle change in microstructures can result in significantly different macroscale behaviors [10]. The 2.25Cr-1Mo-0.25 V steel is a carbide-rich high strength low alloy steel, and the interaction of hydrogen with carbides is crucial to understand the hydrogen enhanced fracture because the carbides have been widely acknowledged as typical hydrogen trapping sites. The hydrogen trapping at carbides has been directly evidenced by atom probe tomography (APT) observation [11e13]. Notably, the trapped hydrogen at carbides are typically but not necessarily “irreversible hydrogen”, while hydrogen atoms trapped at dislocations and grain boundaries are typically “reversible hydrogen” [14]. Many literature studies showed that the HE is mainly attributed to the reversible hydrogen rather than irreversible hydrogen [15e17]. Thus the addition of carbides, which leads to the decrease of the fraction of diffusive hydrogen, has been considered as an effective method to mitigate HE [18]. However, hydrogen trapping ability of carbides is affected by the structures, sizes and distributions [19]. Both detrimental and beneficial effect of carbides on HE has been reported [20e23]. For the 2.25Cr-1Mo and 2.25Cr-1Mo-0.25 V steel, up to six types of carbides have been reported including MC, M2C, M7C3, M3C, M23C6 and M6C [24e27]. Additionally,
these carbides are sensitive to high temperature exposure, during which complicated precipitation, decomposition and coarsening processes may occur. The 2.25Cr-1Mo-0.25 V steel heavy pressure vessels undergo multiple heat treatments, including preheating before hot forming, post-weld heat treatment, quenching and tempering, during their fabrication process. However, the effect of heat treatment on the carbide evolution and HE resistance has not been well understood. In the present study, the effect of quenching-tempering in a head-fabrication heat treatment procedure on the HE resistance of 2.25Cr-1Mo-0.25 V steel was investigated. The variations of microstructures, mechanical properties and HE prior to and after heat treatment were studied. The correlations among the carbide precipitation, hydrogen trapping and macroscale ductility loss were clarified. The present study may contribute to the better understanding of hydrogen trapping behavior of carbides and its role in HE mechanism, which may help explain some cases of failures during the complicated fabrication procedures and starts and shutdowns of reactors, which take place at low temperature.
Experimental Materials and heat treatment The material used in this study was the 2.25Cr-1Mo-0.25 V steel plate, with chemical compositionas follows: (wt.%): C 0.15, Cr 2.30, Mo 0.98, V 0.30, Mn 0.54, Si 0.10, P 0.009, S 0.01, Cu 0.02, Ni 0.05, Al 0.05 and Fe balance. The steel plate with a thickness of 98 mm was manufactured by the ArcelorMittal Company and provided by Lanzhou LS Heavy Equipment Co., Ltd. The heavy pressure vessels, for example hydrogenation reactors, mainly consist of two parts: a shell body and two heads. The fabrication of heads includes four steps: preheating for high temperature stamping (940 C for 100 min) þ hot stamping þ quenching (930 C for 150 min, water cooling) þ tempering (720 C for 200 min, air cooling). Previous study showed that the effect of high temperature plastic deformation on HE was mainly attributed to the heat treatment effect instead of plastic deformation [28]. Therefore, the hot stamping procedure was ignored and the combination of preheating (stage 1) þ quenching (stage 2) þ tempering (stage 3) treatment was considered in the present study. The specimens were heated in an electric furnace with a temperature controlling accuracy of ±1 C. The heating rates were controlled at 180 C/h (stages 1 and 2) and 120 C/h (stage 3), respectively.
Hydrogen charging and tensile tests Immersion method was used for hydrogen charging and the thermal desorption spectroscopy (TDS) tests were conducted to measure the hydrogen concentration. The cylindrical specimens with dimensions of ∅5 15 mm were charged with hydrogen by immersing them into a 20 mass% ammonium thiocyanate (NH4SCN) solution (pH ¼ 4.8) for 48 h. Then the hydrogen content was immediately measured by TDS. The TDS measurement mainly consists of the following three
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steps: 1) grinding the surface with 1000 grit abrasive paper and washing of the specimens in ultrasonic bath for 3 min, 2) preprocessing in vacuum (105 Pa) for 10e15 min, and 3) heating the specimen to 700 C with a heating rate of 100 C/h and measurement of the hydrogen concentration. The tensile tests were conducted immediately after hydrogen charging. The thickness of specimen was 2 mm with a gage length of 25 mm. The displacement rate was 0.65 mm/ min, corresponding to a strain rate of 3.6 104 s1. Each tensile test lasted for less than 10 min, thus it was assumed that the majority of hydrogen remained in the steel during the tensile tests.
Morphology observation Carbides were extracted by the gold replica technique. Specimens with dimensions of 5 10 10 mm3 were mounted using Bakelite molding powder, ground with abrasive paper and polished with 1 mm diamond paste. The specimens were deeply etched with 5% nital for 2 min, and coated with a gold film for 30 s at a current density of 20 mA in a vacuum sputtering machine (MSP-300B). The gold film was divided into several smaller squares (1 2 mm2), etched with a 5% nital, collected with copper grids, and washed in alcohol and distilled water. Then the gold films with extracted carbides were examined by scanning electron microscopy (SEM, Tescan, MAIA3 LMHT) to obtain the typical morphologies. Highresolution SEM images in ten randomly selected regions were obtained to confirm the typical morphologies of carbides in non-treated and QT-treated specimens, respectively. Transmission electron microscpy (TEM, JEM-2100) was utilized to characterize the high-resolution morphologies of carbides and the corresponding structures, according to the following procedures. The thin foils with dimensions of 10 10 0.3 mm3 were cut from the steel plate, ground down to 70 mm (thickness) with silicon carbide abrasive papers and twin-jet electropolished in an electrolyte, comprising 5% of perchloric acid and 95% of ethanol, with the polishing current controlled at 30 mA. The small-sized carbides in steel were characterized by using APT, which was performed using a CAMEA 4000HR local electrode atom probe. The samples with dimensions of 0.5 0.5 15 mm3 were cut from the as-received steel plate and electro-polished to needle-shape. Then the APT measurement was conducted according to the standard procedures [29]. The electron back scattering diffraction (EBSD) analysis was performed with a beam step size of 0.1 mm on ionpolished surface. For the fractographic observation, the specimens with 15 mm height were cut from the fractured specimens, cleaned with acetone and dried in the air. The morphologies were observed by SEM (Tescan, MAIA3 LMHT).
Results and discussion Microstructures The 2.25Cr-1Mo-0.25 V steel is carbide-rich steel, and the evolution of nano-sized carbides during high temperature
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aging is crucial for understanding the mechanical properties at macroscale. The variation of microstructures in as-received (AR) and QT-treated specimens is shown in Fig. 1. The predominance of strip-shaped carbides and their clusters could be confirmed from Fig. 1a and c for AR specimens. The structure of strip-shaped carbides was determined as M3C (cementite), as shown in Fig. 1e. For QT specimens, no trace of cementite could be found while the predominance of granular carbides was evidenced, as shown in Fig. 1b and d. Most of the granular carbides were determined as M7C3 carbides, as shown in Fig. 1f. Pilling and Ridley [30] reported that the granular carbides could be M6C, M23C6 or M7C3 carbides, for 2.25Cr-1Mo steel. These three types of carbides are more stable during high temperature tempering than M3C, M2C and MC. For example, R.G. Baker [31] studied the carbide evolution of 2.25Cr-1Mo steel during tempering (temperature ranging from 400 to 750 C, for up to 1000 h) and proposed the precipitation sequence of carbides as: M3C / (M2C þ M3C) / M23C6 / M6C, or M3C / (M2C þ M3C) / M7C3 / M6C. Moreover, it has also been reported that the transformation process into M23C6 and M6C carbides can be accelerated when the tempering temperature is above 600 C, due to the higher diffusion rate of Mo and Cr elements. Jiang et al. [26,32] identified the presence of MC/M7C3/M23C6 carbides and reported the absence of M3C after tempering (700 C for 128 h) treatment of 2.25Cr-1Mo-0.25 V steel, by X-ray diffraction (XRD) and selected area electron diffraction (SAED) analysis. The granular carbides were identified as M23C6. Pilling and Ridley [30] investigated the tempering of 2.25Cr-1Mo steel with very low carbon level (0.018e0.09 wt %), and identified the carbides as M6C/M23C6 after tempering at 700 C for 12 h, by electron probe microanalysis (EMPA) analysis. The decomposition of cementite and precipitation of M7C3 carbides after QT treatment observed in the present study are consistent with the results reported in literature. This implies that no trace of cementite could be found in the steel when exposed to the in-service environment because most of them have already decomposed during the fabrication processes of the pressure vessel. The presence of fine V-rich carbides is the main cause of excellent mechanical properties of 2.25Cr-1Mo-0.25 V steel due to the addition of vanadium element, compared to traditional 2.25Cr-1Mo steel. Therefore, the variation of carbides with small sizes prior to and after QT treatment was also investigated. For the AR specimens, the small-sized carbides are uniformly distributed in the steel, as shown in Fig. 2a. These carbides were further characterized by APT technique, which is capable of revealing the three-dimensional structure of nano-sized precipitations at the atomic scale. Fig. 2b shows the element distribution mapping for V using 1% isoconcentration surfaces, which clearly sketches the shapes of carbides. The cross-section profiles in Fig. 2c show the chemical compositions of typical carbide presented in Fig. 2b, indicating the presence of VC carbide. It is seen that some of the VC carbides are exhibiting lamella-like shapes, which is invisible in SEM or TEM images because they can only present the cross-section morphologies in two dimensions. Furthermore, granular VCs are also observed in TEM images and APT
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Fig. 1 e (a, b) SEM images, (c, d) TEM images and (e, f) SAED patterns for the carbides in AR and QT specimens.
map, which was also reported for forged 2.25Cr-1Mo-0.25 V steels [32]. For the QT specimens, no such uniformly distributed carbides could be found in TEM images. Several groups of APT tests were performed but none of them indicated any trace of element segregation. It is believed that many of the fine VCs were decomposed during QT treatment. Fig. 3 shows the recrystallization maps, kernel average maps (KAM), phase maps, and inverse pole figure (IPF) maps for AR and QT specimens, respectively. Clearly, the recrystallization fraction (Fig. 3a and b) increases and dislocation density (Fig. 3c and d) decreases, which can lead to the decrease of residual stress in steel. The phase maps shown in Fig. 3e and f demonstrate the fully bainitic structure of the steel and no trace of retained M/A islands could be found, which was however reported in the forgings of 2.25Cr-1Mo0.25 V steel [32], in contrast with steel plate used in the present study. Furthermore, no significant difference of grain size and orientation distribution could be observed for the AR and QT specimens, as shown in Fig. 3g and h. Therefore, the effect of grain boundaries could also be ruled out.
Hydrogen content TDS tests were conducted for the AR specimens, as shown in Fig. 4a. A considerable amount of hydrogen was detected for the AR specimens, the content of which was determined as 0.94 ppm. Three local peaks were identified from the TDS curve, with the peak temperature as 206.9 C, 244.5 C and 362.6 C respectively. These peaks correspond to different
hydrogen traps in steel, and peaks at higher temperature correspond to deeper traps with higher binding energy. Furthermore, based on the binding energy between hydrogen atoms and the local trap sites, hydrogen traps can be classified into two categories: reversible and irreversible traps. All the hydrogen states shown in Fig. 4a can be classified as irreversible hydrogen because the as-received steel had already been aged in air for more than two years. Therefore, all the hydrogen states with peak temperature above 206.9 C can be classified as irreversible hydrogen for the following TDS curves, which were obtained using the identical testing parameters with that in Fig. 4a. Therefore, the AR specimens should be dehydrogenated before the study of HE. A group of dehydrogenation tests was conducted to determine the appropriate parameters to completely degas the hydrogen present in AR specimens, as shown in Fig. 4b (DeH for dehydrogenated specimens). By comparing the results between DeH-200 C and DeH-300 C, it is seen that hydrogen in steel tends to escape faster when it is aged at higher temperature. The irreversible hydrogen in steel could be fully degassed by aging at 350 C for 60 min, as indicated by the green line in Fig. 4b. Notably, considerable background noise could be detected especially in the high temperature range when the hydrogen content in specimen was very low. A custom-designed TDS instrument can be used to eliminate the background signal [33]. Therefore, all the AR specimens were dehydrogenated at 350 C for 60 min, both in TDS tests and tensile tests. Fig. 5a shows desorption curves after hydrogen charging for QT specimens and dehydrogenated AR specimens. The total hydrogen
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Many studies showed that the local peaks in TDS curves corresponding to the hydrogen trapped at dislocations were typically at low temperature with a peak temperature below 100 C [35e37]. Note that the peak temperature for a particular hydrogen trap tends to shift positively with the increase in the heating speed in TDS tests. A heating speed of 100 C/h was utilized in the present study, which is lower than those used in the studies on the trapping behavior of dislocations. Thus all the hydrogen trapped at dislocations should be desorbed below 100 C, which is however not detected in any of the TDS tests in present study. It is likely that most of the reversible hydrogen desorbed below 100 C had already escaped during the pre-processing procedure of specimens in high vacuum environment during the TDS tests. This phenomenon has also been reported elsewhere for some other materials [17]. Therefore, the hydrogen trapping at dislocations could be ruled out for the local peaks shown in Figs. 4 and 5. Although it is extremely difficult to build one-to-one correlation between the local peaks in Fig. 5 and the local microstructures in steel, it is assumed here that the peak 2 hydrogen is attributed to those trapped at the small-sized carbides such as VCs, while peak 1 hydrogen corresponds to the hydrogen at the other carbides with larger sizes, considering that small-sized precipitations tend to have stronger trapping ability of hydrogen atoms [17,19,38]. Whether the irreversible trapping sites indeed correspond to VC carbides still needs more solid evidence, because the desorption of hydrogen from VC carbides in V-bearing high strength steel has been reported [39].
Tensile tests
Fig. 2 e (a) High-resolution TEM images of fine carbides with needle-like and granular shapes, (b) element map for V with 1% iso-surfaces obtained through APT technique and (c) the corresponding cross-section profile of element distribution, in AR specimens.
content for DeH-HC (dehydrogenated and recharged with hydrogen) and QT-HC (quenched-tempered and hydrogencharged) specimens was 2.37 and 1.04 ppm respectively, with two local peaks identified, as shown in Fig. 5b and c. The peak 2 hydrogen in these two curves, with peak temperature being 214.6 and 212.9 C, can be confirmed as irreversible hydrogen, based on the evidence drawn from Fig. 4a. The peak 1 hydrogen, with peak temperature being 127.0 and 129.4 C, can be classified as reversible hydrogen. The results clearly indicate significant decease in the amount of both reversible and irreversible hydrogen, which is probably due to the decreased hydrogen diffusion coefficient after heat treatment as previously demonstrated by electrochemical permeation tests [34].
The results of tensile tests are presented in Fig. 6. It is found that the QT treatment results in significantly decreased strength and increased ductility. For the AR specimens, no obvious difference was observed after heat treatment at 350 C, which indicates that residual stress should have little effect on mechanical properties. This also demonstrates the feasibility of dehydrogenation treatment at 350 C, which should not have significant influence on the microstructures and mechanical properties. The small deviation between the above mentioned two curves is possibly due to the different locations of specimens which were randomly cut from the steel plate without considering the effect of plate thickness. For the DeH specimens, obvious ductility loss could be observed after hydrogen charging. The susceptibility to HE was decreased to a negligible level after QT treatment, as indicated by the dashed lines in Fig. 6. The HE resistance was apparently improved by QT treatment. This agrees well with the SEM fractography result. The SEM fractography images for AR specimens are shown in Fig. 7aed. In the absence of hydrogen, dimples and secondary cracks could be found on the fractured surface, indicating ductile fracture. However, the fractured surface for the hydrogen-charged specimens featured quasi-cleavages and decreased number of dimples. The transition from ductile fracture to brittle-ductile fracture after hydrogen charging was verified. For the QT specimens, the fractographies are all featured as ductile fracture whether charging or not, as shown in Fig. 7e and f.
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Fig. 3 e (a, b) Recrystallization, (c, d) KAM, (e, f) phase maps, and (g, h) IPF maps for AR and QT specimens.
Hydrogen embrittlement mechanism The variations of microstructures, hydrogen concentration and mechanical properties were confirmed prior to and after QT treatment. It would be necessary to provide an explanation for the improved HE resistance after QT treatment. First, significant decrease of strength after QT treatment was observed. The influence of strength on HE should be clarified, which is strongly affected by the microstructures of material. For example, the susceptibility to HE tends to increase with increasing strength for pipeline steels [40]. However, the opposite trends were reported for the grain refinement strengthened steels, in which the decrease of grain size results in high strength but the susceptibility to HE is decreased due to the reduced diffusive hydrogen per grain boundary area [41]. This contradiction also agrees well with the significant scatter that exhibits in the studies of HE for different grades of steels, as reviewed by Barnoush [42]. The variety of influencing factors leads to the scatter of data and one cannot just conclude whether the strength level is beneficial or detrimental to HE without taking all the influencing factors into account. The effect of strength on HE for high strength steels should be discussed in combination with the strengthening mechanisms of steel. Among all these factors that may determine the brittle fracture, the main factor should be always related to the principal strengthening mechanism for the involved high strength steel. For the AR and QT-treated 2.25Cr-1Mo0.25 V steel in this paper, the effect precipitation, dislocation density and grain size should be considered when discussing the effect of strength.
For the AR specimens, the numerous small-sized carbides in steel result in high strength due to the impediment to dislocation movement by precipitations, and concurrently high resistance to HE, due to the hydrogen trapping at carbides. After QT treatment, the impediment effect of dislocation movement by precipitations was weakened due to the coarsening of carbides, which results in the decrease of strength. At the same time, the HE resistance was improved due to the decrease of hydrogen concentration including both reversible and irreversible hydrogen, which was mainly attributed to the decomposition, precipitation and coarsening of carbides. It is believed that M3C acts as reversible trapping sites, while VC as irreversible trapping sites, although not rigorously proved by experiments. The decreased amount of reversible hydrogen due to the decomposition of M3C is mainly responsible for the improved HE resistance. This finding agrees well with the results reported in literature [38], which indicates that the increased number of reversible trapping sites by carbides can result in detrimental effect on HE. The decreased amount of hydrogen at VC (peak 2 hydrogen) is believed not to be responsible for the variation of HE sensitivity, as it is wellacknowledged that irreversible hydrogen has little effect on HE resistance [16,43]. Hydrogen charging method is also an important influencing factor when discussing the HE mechanism. The ex-situ hydrogen charging method was used in the present study, which enabled to rule out the effect of charging condition, considering that significantly different tensile time could be encountered for different specimens depending on the HE susceptibility for in-situ hydrogen charging. On the other
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Fig. 4 e TDS results for (a) AR specimens without hydrogen charging, and (b) AR specimens after different dehydrogenation treatments (DeH specimens).
hand, the ex-situ hydrogen charging also rules out the influence of hydrogen refilling, and this is crucial when discussing the effect of reversible and irreversible hydrogen. In in-situ hydrogen charging, the reversible hydrogen traps are immediately refilled with hydrogen. The amount of diffusive hydrogen in steel between AR and QT specimens changes during tensile tests and is not simply affected by the original microstructures any more, which makes the analysis of HE mechanism more complicated. However, the usage of ex-situ charging can rule out this influence. It is believed that QT can decrease the amount of hydrogen content including both reversible and irreversible hydrogen, as demonstrated by TDS tests. The improved HE resistance can be explained as follows. It is known that the presence of hydrogen can enhance the local plasticity by increasing the dislocation mobility, which is known as the hydrogen enhanced local plasticity (HELP) mechanism [44]. The increased dislocation mobility by hydrogen can be explained from the aspect of hydrogen
Fig. 5 e (a) TDS results after hydrogen charging for DeH specimens and QT specimens, with the total hydrogen concentration shown in the inset figure. The fitted results of these two curves are shown in (b) and (c) respectively.
shielding effect. The aggregated hydrogen at the stress fields of dislocations modifies the stress field to which it is attached, and reduces the interaction energy of the dislocation and the elastic obstacles allowing dislocation motion at lower stress level in certain directions [45]. The premise of this mechanism
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Fig. 6 e Stress-strain curves for AR, DeH, and QT specimens with and without hydrogen charging.
is that local hydrogen is present in sufficient concentration [46]. For the QT specimens, the amount of diffusive hydrogen was decreased, which weakens the enhancement effect on the dislocation mobility. No obvious ductility loss was observed because the local hydrogen concentration does not reach a critical level so that the dislocation mobility can be affected. Additionally, the applied stress is also a key factor for the HE mechanism and the critical hydrogen concentration that is required for the dislocation initiation is probably dependent
on the local stress concentration. Significant local stress level is expected for the AR and QT specimens due to the difference of strength. Whether this stress difference has significant influence on the critical hydrogen concentration still needs further investigation. Examining the HE resistance at different hydrogen concentrations for both AR and QT specimens would be a viable route. The susceptibility to HE is also strongly affected by strain rate. It is possible that apparent ductility loss could be observed for the QT-HC specimens at lower strain rate, although the strain rate in the present work is relatively low (3.6 104 s1). The hydrogen release during tensile tests also has influence on the final fracture, because the time to fracture for QT specimen is larger than that for AR specimen. The largest time difference for the tensile tests is estimated as 4.6 min. Compared with the influence of hydrogen concentration, we believe this is not the primary influencing factor. This influence can be ruled out by electroplating the specimens with a cadmium coating to prevent hydrogen release [47], which may be investigated in the future work. It is also known that the susceptibility to HE tends to increase at higher hydrogen concentrations [48e50]. For example, Zhu et al. [51], also reported a beneficial effect of tempering on HE for a quenchingepartitioningetempering (QP-T) treated steel, but the improvement of HE was observed only at very low hydrogen concentration and the effect was eliminated after further extending the charging duration. Therefore, exploring the HE resistance of QT specimens at higher hydrogen concentration would be an interesting topic. All the tests in this study were conducted at room temperature, thus the discovered phenomenon and the
Fig. 7 e SEM fractographies for (a) AR-NC (as-received, non-charged) and (b) AR-HC (as-received, hydrogen-charged) specimens, with the enlarged figures shown in (c) and (d) respectively; SEM fractographies for (e) QT-NC (quenchedtempered, non-charged) and (f) QT-HC (quenched-tempered, hydrogen-charged) specimens.
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corresponding HE mechanism may provide guidance for the hydrogen induced cracking during fabrication procedures or startups and shutdowns of reactors, which take place at low temperature. However, the difference between the hydrogen trapping behavior of carbides at room temperature and elevated temperature should be noted. For example, those carbides which are recognized as irreversible trapping sites at room temperature may become reversible at elevated temperature. According to the TDS results for the dehydrogenated specimens in Fig. 4, the hydrogen which is irreversible at room temperature can be completely degassed at 350 C. This implies that all the hydrogen atoms in steel are probably diffusible and can contribute to the macroscale ductility loss, in a high temperature hydrogen environment. Additionally, the interaction mechanism of hydrogen with dislocations may also become different at elevated temperature. The hydrogen atoms may impede the motion of the dislocations by acting as solid solution strengthening agent because they do not remain bound to dislocations at elevated temperature [45].
Conclusions The effect of quenching-tempering (QT) treatment on the hydrogen embrittlement (HE) of 2.25Cr-1Mo-0.25 V steel was investigated through SEM, TEM, APT, TDS, and tensile tests. The following conclusions were obtained: (1) The QT treatment resulted in the decrease of strength, significant increase of ductility and apparently enhanced HE resistance. (2) Two hydrogen states were identified in the hydrogen desorption profiles in TDS tests, with peak 1 hydrogen desorbed at low temperature (~127 C), and peak 2 hydrogen at high temperature (~215 C). It was demonstrated that peak 2 hydrogen is irreversible hydrogen by comparing its peak temperature with that detected in the AR specimens which had been aged in air for 2 years. The hydrogen content of both reversible and irreversible hydrogen decreased after QT treatment. (3) QT treatment resulted in the decomposition of M3C/VC and precipitation of M7C3. The VCs are more probably irreversible traps, corresponding to peak 2 in TDS curves, while the other carbides are reversible traps, corresponding to peak 1. The decrease of reversible hydrogen is believed to be responsible for the improved HE resistance.
Acknowledgements This work is supported by the National Basic Research Program of China (973 Program, Grant No.2015CB057602), China Postdoctoral Science Foundation (Grant No. BX20180245, 2018M643637) and Fundamental Research Funds for the Central Universities (Grant No. XZY012019024).
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[20] Depover T, Verbeken K. The effect of TiC on the hydrogen induced ductility loss and trapping behavior of Fe-C-Ti alloys. Corros Sci 2016;112:308e26. [21] Depover T, Verbeken K. Thermal desorption spectroscopy study of the hydrogen trapping ability of W based precipitates in a Q&T matrix. Int J Hydrogen Energy 2018;43:5760e9. [22] Depover T, Verbeken K. Hydrogen trapping and hydrogen induced mechanical degradation in lab cast Fe-C-Cr alloys. Mater Sci Eng A 2016;669:134e49. [23] Lee J, Lee T, Kwon YJ, Mun D-J, Yoo J-Y, Lee CS. Effects of vanadium carbides on hydrogen embrittlement of tempered martensitic steel. Met Mater Int 2016;22:364e72. [24] Peddle BE, Pickles CA. Carbide development in the heat affected zone of tempered and post-weld heat treated 2.25Cr-1Mo steel weldments. Can Metall Q 2001;40:105e25. [25] Yongtao Z, Haibo H, Lede M, Hanqian Z, Jinfu L. Quantitative carbide analysis using the Rietveld method for 2.25Cre1Moe0.25V steel. Mater Char 2009;60:953e6. [26] Jiang Z, Wang P, Li D, Li Y. The evolutions of microstructure and mechanical properties of 2.25Cr-1Mo-0.25V steel with different initial microstructures during tempering. Mater Sci Eng A 2017;699:165e75. [27] Tezuka H, Sakurai T. A trigger of Type IV damage and a new heat treatment procedure to suppress it. Microstructural investigations of long-term ex-service CreMo steel pipe elbows. Int J Press Vessel Pip 2005;82:165e74. [28] Wang Y, Cheng G, Qin M, Li Q, Zhang Z, Chen K, et al. Effect of high temperature deformation on the microstructure, mechanical properties and hydrogen embrittlement of 2.25Cre1Mo-0.25 V steel. Int J Hydrogen Energy 2017;42:24549e59. [29] Miller MK, Kenik EA. Atom probe tomography: a technique for nanoscale characterization. Microsc Microanal 2004;10:336e41. [30] Pilling J, Ridley N. Tempering of 2.25 pct Cr-1 pct Mo low carbon steels. Metall Trans A 1982;13:557e63. [31] B RG. The tempering of 2.25Cr-1Mo after quenching and normalizing. J Iron Steel Inst 1959;192:257e68. [32] Jiang Z, Wang P, Li D, Li Y. Influence of the decomposition behavior of retained austenite during tempering on the mechanical properties of 2.25Cr-1Mo-0.25 V steel. Mater Sci Eng A 2019;742:540e52. [33] Merzlikin SV, Borodin S, Vogel D, Rohwerder M. Ultra high vacuum high precision low background setup with temperature control for thermal desorption mass spectroscopy (TDA-MS) of hydrogen in metals. Talanta 2015;136:108e13. [34] Song Y, Han Z, Chai M, Yang B, Liu Y, Cheng G, et al. Effect of cementite on the hydrogen diffusion/trap characteristics of 2.25Cr-1Mo-0.25V steel with and without annealing. Materials 2018;11. rez Escobar D, Depover T, Duprez L, Verbeken K, [35] Pe Verhaege M. Combined thermal desorption spectroscopy, differential scanning calorimetry, scanning electron microscopy and X-ray diffraction study of hydrogen trapping in cold deformed TRIP steel. Acta Mater 2012;60:2593e605.
[36] Depover T, Verbeken K. The detrimental effect of hydrogen at dislocations on the hydrogen embrittlement susceptibility of Fe-C-X alloys: an experimental proof of the HELP mechanism. Int J Hydrogen Energy 2018;43:3050e61. rez Escobar D, Depover T, Wallaert E, Duprez L, [37] Pe Verhaege M, Verbeken K. Thermal desorption spectroscopy study of the interaction between hydrogen and different microstructural constituents in lab cast FeeC alloys. Corros Sci 2012;65:199e208. [38] Depover T, Verbeken K. Evaluation of the effect of V4C3 precipitates on the hydrogen induced mechanical degradation in Fe-C-V alloys. Mater Sci Eng A 2016;675:299e313. [39] Hagihara Y, Shobu T, Hisamori N, Suzuki H, Takai K-i, Hirai K. Delayed fracture using CSRT and hydrogen trapping characteristics of V-bearing high-strength steel. ISIJ Int 2012;52:298e306. [40] Hardie D, Charles EA, Lopez AH. Hydrogen embrittlement of high strength pipeline steels. Corros Sci 2006;48:4378e85. [41] Park C, Kang N, Liu S. Effect of grain size on the resistance to hydrogen embrittlement of API 2W Grade 60 steels using in situ slow-strain-rate testing. Corros Sci 2017;128:33e41. [42] Barnoush A, Vehoff H. Recent developments in the study of hydrogen embrittlement: hydrogen effect on dislocation nucleation. Acta Mater 2010;58:5274e85. [43] Takai K, Watanuki R. Hydrogen in trapping states innocuous to environmental degradation of high-strength steels. ISIJ Int 2003;43:520e6. [44] Birnbaum HK, Sofronis P. Hydrogen-enhanced localized plasticityda mechanism for hydrogen-related fracture. Mater Sci Eng A 1994;176:191e202. [45] Robertson IM, Sofronis P, Nagao A, Martin ML, Wang S, Gross DW, et al. Hydrogen embrittlement understood. Metall Mater Trans A 2015;46:2323e41. [46] Martin ML, Dadfarnia M, Nagao A, Wang S, Sofronis P. Enumeration of the hydrogen-enhanced localized plasticity mechanism for hydrogen embrittlement in structural materials. Acta Mater 2019;165:734e50. [47] Akiyama E, Wang M, Li S, Zhang Z, Kimura Y, Uno N, et al. Studies of evaluation of hydrogen embrittlement property of high-strength steels with consideration of the effect of atmospheric corrosion. Metall Mater Trans A 2013;44:1290e300. [48] Zakroczymski T, Glowacka A, Swiatnicki W. Effect of hydrogen concentration on the embrittlement of a duplex stainless steel. Corros Sci 2005;47:1403e14. [49] Koyama M, Akiyama E, Tsuzaki K. Effect of hydrogen content on the embrittlement in a FeeMneC twinning-induced plasticity steel. Corros Sci 2012;59:277e81. [50] Dong CF, Liu ZY, Li XG, Cheng YF. Effects of hydrogencharging on the susceptibility of X100 pipeline steel to hydrogen-induced cracking. Int J Hydrogen Energy 2009;34:9879e84. [51] Zhu X, Li W, Hsu TY, Zhou S, Wang L, Jin X. Improved resistance to hydrogen embrittlement in a high-strength steel by quenchingepartitioningetempering treatment. Scr Mater 2015;97:21e4.
Available online at www.sciencedirect.com
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Improved hydrogen embrittlement resistance after quenchingetempering treatment for a Cr-Mo-V high strength steel Yafei Wang*, Songyan Hu, Yun Li, Guangxu Cheng School of Chemical Engineering and Technology, Xi’an Jiaotong University, Xi’an 710049, China
highlights Decomposition, precipitation and coarsening of carbides confirmed after QT treatment. Beneficial effect of QT-treatment on HE resistance. Decreased reversible hydrogen traps responsible for improved HE resistance.
article info
abstract
Article history:
The effect of quenching-tempering (QT) treatment on the hydrogen embrittlement (HE)
Received 14 August 2019
resistance of a reactor pressure vessel steel was studied. Decomposition of M3C/VC car-
Received in revised form
bides and precipitation of M7C3 carbides were confirmed by transmission electron micro-
17 September 2019
scopy and atom probe tomography observations. Tensile tests showed that HE sensitivity
Accepted 19 September 2019
decreased to a negligible level after QT treatment. The improvement of HE resistance was
Available online 11 October 2019
mainly attributed to the decreased number of M3C carbides which act as the reversible trapping sites for hydrogen. This was supported by the decreased concentration of
Keywords:
reversible hydrogen as measured by thermal desorption spectroscopy. The amount of
Steel
irreversible hydrogen (probably trapped at VC carbides) also decreased, which is however
Heat treatment
not considered responsible for the HE improvement.
Carbides
© 2019 Hydrogen Energy Publications LLC. Published by Elsevier Ltd. All rights reserved.
Hydrogen embrittlement
Introduction Quenching-tempering (QT) treatment is an important heat treatment procedure during the fabrication of hydrogenation reactors, which operate at elevated temperature and high hydrogen partial pressure. Moreover, high temperature hydrogen attack (HTHA) is the main failure mode for these equipment and components that work in a high temperature hydrogen environment. Significant research efforts have
been devoted in the past decades on the study of HTHA mechanism and prevention strategies. The Nelson curves have been extensively used in the industry to determine the operation limit of steels for hydrogen service at elevated temperatures and pressures, according to the American Petroleum Institute (API) Recommended Practice 941 [1]. When the temperature and pressure are controlled below the operation limit, the HTHA is not considered a problem, although the exact values of operation limit for some steels have been revised from time to time based on the experiences
* Corresponding author. E-mail address: [email protected] (Y. Wang). https://doi.org/10.1016/j.ijhydene.2019.09.142 0360-3199/© 2019 Hydrogen Energy Publications LLC. Published by Elsevier Ltd. All rights reserved.
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from some cases of accidental failures (below the operation limit) due to HTHA [2]. It is well-recognized that the carbides in steel can reduce the amount of carbon that is available to react with hydrogen to form methane, which is the key pathway of HTHA. Therefore, the carbide-rich 2.25Cr-1Mo-0.25 V steel has been widely accepted as one of the optimal materials for many components working under high temperature hydrogen environment. According to the Nelson curves, the temperature limit for 2.25Cr-1Mo-0.25 V steel is 510 C when the hydrogen partial pressure is higher than 12.5 MPa, and the temperature limit increases linearly with the further decreasing of hydrogen pressure. The latest version of Nelson curve for 2.25Cr-1Mo-0.25 V steel is based on the data obtained from 10000 þ hours of laboratory tests. Assuming that this curve is accurate enough, the HTHA is not a major concern for hydrogenation reactors made of this material, considering that many reactors operate even far below the temperature limit. In this way, some other failure modes such as hydrogen embrittlement (HE) or hydrogen induced cracking (HIC) may become the important factors that need to be considered for the safe operation of reactors. On the other hand, some cases of failures have been reported during the fabrication of reactors, mostly during welding or heat treatments [3e5]. In these cases, the adsorbed hydrogen, either from the high temperature hydrogen gas in service or from the water molecules in atmosphere during fabrication processes, can play a critical role. The interaction of hydrogen with local microstructures may become the controlling factor that determines the final failures. Therefore, the HE of 2.25Cr-1Mo and 2.25Cr1Mo-0.25 V steel has been reported in some studies [6e9]. However, the hydrogen enhanced delayed fracture is influenced by many factors that take place in different scales, which make the failure mechanism complicated. Thus it is usually difficult to determine the underlying controlling factors by using a single method or methods in similar scales, and subtle change in microstructures can result in significantly different macroscale behaviors [10]. The 2.25Cr-1Mo-0.25 V steel is a carbide-rich high strength low alloy steel, and the interaction of hydrogen with carbides is crucial to understand the hydrogen enhanced fracture because the carbides have been widely acknowledged as typical hydrogen trapping sites. The hydrogen trapping at carbides has been directly evidenced by atom probe tomography (APT) observation [11e13]. Notably, the trapped hydrogen at carbides are typically but not necessarily “irreversible hydrogen”, while hydrogen atoms trapped at dislocations and grain boundaries are typically “reversible hydrogen” [14]. Many literature studies showed that the HE is mainly attributed to the reversible hydrogen rather than irreversible hydrogen [15e17]. Thus the addition of carbides, which leads to the decrease of the fraction of diffusive hydrogen, has been considered as an effective method to mitigate HE [18]. However, hydrogen trapping ability of carbides is affected by the structures, sizes and distributions [19]. Both detrimental and beneficial effect of carbides on HE has been reported [20e23]. For the 2.25Cr-1Mo and 2.25Cr-1Mo-0.25 V steel, up to six types of carbides have been reported including MC, M2C, M7C3, M3C, M23C6 and M6C [24e27]. Additionally,
these carbides are sensitive to high temperature exposure, during which complicated precipitation, decomposition and coarsening processes may occur. The 2.25Cr-1Mo-0.25 V steel heavy pressure vessels undergo multiple heat treatments, including preheating before hot forming, post-weld heat treatment, quenching and tempering, during their fabrication process. However, the effect of heat treatment on the carbide evolution and HE resistance has not been well understood. In the present study, the effect of quenching-tempering in a head-fabrication heat treatment procedure on the HE resistance of 2.25Cr-1Mo-0.25 V steel was investigated. The variations of microstructures, mechanical properties and HE prior to and after heat treatment were studied. The correlations among the carbide precipitation, hydrogen trapping and macroscale ductility loss were clarified. The present study may contribute to the better understanding of hydrogen trapping behavior of carbides and its role in HE mechanism, which may help explain some cases of failures during the complicated fabrication procedures and starts and shutdowns of reactors, which take place at low temperature.
Experimental Materials and heat treatment The material used in this study was the 2.25Cr-1Mo-0.25 V steel plate, with chemical compositionas follows: (wt.%): C 0.15, Cr 2.30, Mo 0.98, V 0.30, Mn 0.54, Si 0.10, P 0.009, S 0.01, Cu 0.02, Ni 0.05, Al 0.05 and Fe balance. The steel plate with a thickness of 98 mm was manufactured by the ArcelorMittal Company and provided by Lanzhou LS Heavy Equipment Co., Ltd. The heavy pressure vessels, for example hydrogenation reactors, mainly consist of two parts: a shell body and two heads. The fabrication of heads includes four steps: preheating for high temperature stamping (940 C for 100 min) þ hot stamping þ quenching (930 C for 150 min, water cooling) þ tempering (720 C for 200 min, air cooling). Previous study showed that the effect of high temperature plastic deformation on HE was mainly attributed to the heat treatment effect instead of plastic deformation [28]. Therefore, the hot stamping procedure was ignored and the combination of preheating (stage 1) þ quenching (stage 2) þ tempering (stage 3) treatment was considered in the present study. The specimens were heated in an electric furnace with a temperature controlling accuracy of ±1 C. The heating rates were controlled at 180 C/h (stages 1 and 2) and 120 C/h (stage 3), respectively.
Hydrogen charging and tensile tests Immersion method was used for hydrogen charging and the thermal desorption spectroscopy (TDS) tests were conducted to measure the hydrogen concentration. The cylindrical specimens with dimensions of ∅5 15 mm were charged with hydrogen by immersing them into a 20 mass% ammonium thiocyanate (NH4SCN) solution (pH ¼ 4.8) for 48 h. Then the hydrogen content was immediately measured by TDS. The TDS measurement mainly consists of the following three
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steps: 1) grinding the surface with 1000 grit abrasive paper and washing of the specimens in ultrasonic bath for 3 min, 2) preprocessing in vacuum (105 Pa) for 10e15 min, and 3) heating the specimen to 700 C with a heating rate of 100 C/h and measurement of the hydrogen concentration. The tensile tests were conducted immediately after hydrogen charging. The thickness of specimen was 2 mm with a gage length of 25 mm. The displacement rate was 0.65 mm/ min, corresponding to a strain rate of 3.6 104 s1. Each tensile test lasted for less than 10 min, thus it was assumed that the majority of hydrogen remained in the steel during the tensile tests.
Morphology observation Carbides were extracted by the gold replica technique. Specimens with dimensions of 5 10 10 mm3 were mounted using Bakelite molding powder, ground with abrasive paper and polished with 1 mm diamond paste. The specimens were deeply etched with 5% nital for 2 min, and coated with a gold film for 30 s at a current density of 20 mA in a vacuum sputtering machine (MSP-300B). The gold film was divided into several smaller squares (1 2 mm2), etched with a 5% nital, collected with copper grids, and washed in alcohol and distilled water. Then the gold films with extracted carbides were examined by scanning electron microscopy (SEM, Tescan, MAIA3 LMHT) to obtain the typical morphologies. Highresolution SEM images in ten randomly selected regions were obtained to confirm the typical morphologies of carbides in non-treated and QT-treated specimens, respectively. Transmission electron microscpy (TEM, JEM-2100) was utilized to characterize the high-resolution morphologies of carbides and the corresponding structures, according to the following procedures. The thin foils with dimensions of 10 10 0.3 mm3 were cut from the steel plate, ground down to 70 mm (thickness) with silicon carbide abrasive papers and twin-jet electropolished in an electrolyte, comprising 5% of perchloric acid and 95% of ethanol, with the polishing current controlled at 30 mA. The small-sized carbides in steel were characterized by using APT, which was performed using a CAMEA 4000HR local electrode atom probe. The samples with dimensions of 0.5 0.5 15 mm3 were cut from the as-received steel plate and electro-polished to needle-shape. Then the APT measurement was conducted according to the standard procedures [29]. The electron back scattering diffraction (EBSD) analysis was performed with a beam step size of 0.1 mm on ionpolished surface. For the fractographic observation, the specimens with 15 mm height were cut from the fractured specimens, cleaned with acetone and dried in the air. The morphologies were observed by SEM (Tescan, MAIA3 LMHT).
Results and discussion Microstructures The 2.25Cr-1Mo-0.25 V steel is carbide-rich steel, and the evolution of nano-sized carbides during high temperature
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aging is crucial for understanding the mechanical properties at macroscale. The variation of microstructures in as-received (AR) and QT-treated specimens is shown in Fig. 1. The predominance of strip-shaped carbides and their clusters could be confirmed from Fig. 1a and c for AR specimens. The structure of strip-shaped carbides was determined as M3C (cementite), as shown in Fig. 1e. For QT specimens, no trace of cementite could be found while the predominance of granular carbides was evidenced, as shown in Fig. 1b and d. Most of the granular carbides were determined as M7C3 carbides, as shown in Fig. 1f. Pilling and Ridley [30] reported that the granular carbides could be M6C, M23C6 or M7C3 carbides, for 2.25Cr-1Mo steel. These three types of carbides are more stable during high temperature tempering than M3C, M2C and MC. For example, R.G. Baker [31] studied the carbide evolution of 2.25Cr-1Mo steel during tempering (temperature ranging from 400 to 750 C, for up to 1000 h) and proposed the precipitation sequence of carbides as: M3C / (M2C þ M3C) / M23C6 / M6C, or M3C / (M2C þ M3C) / M7C3 / M6C. Moreover, it has also been reported that the transformation process into M23C6 and M6C carbides can be accelerated when the tempering temperature is above 600 C, due to the higher diffusion rate of Mo and Cr elements. Jiang et al. [26,32] identified the presence of MC/M7C3/M23C6 carbides and reported the absence of M3C after tempering (700 C for 128 h) treatment of 2.25Cr-1Mo-0.25 V steel, by X-ray diffraction (XRD) and selected area electron diffraction (SAED) analysis. The granular carbides were identified as M23C6. Pilling and Ridley [30] investigated the tempering of 2.25Cr-1Mo steel with very low carbon level (0.018e0.09 wt %), and identified the carbides as M6C/M23C6 after tempering at 700 C for 12 h, by electron probe microanalysis (EMPA) analysis. The decomposition of cementite and precipitation of M7C3 carbides after QT treatment observed in the present study are consistent with the results reported in literature. This implies that no trace of cementite could be found in the steel when exposed to the in-service environment because most of them have already decomposed during the fabrication processes of the pressure vessel. The presence of fine V-rich carbides is the main cause of excellent mechanical properties of 2.25Cr-1Mo-0.25 V steel due to the addition of vanadium element, compared to traditional 2.25Cr-1Mo steel. Therefore, the variation of carbides with small sizes prior to and after QT treatment was also investigated. For the AR specimens, the small-sized carbides are uniformly distributed in the steel, as shown in Fig. 2a. These carbides were further characterized by APT technique, which is capable of revealing the three-dimensional structure of nano-sized precipitations at the atomic scale. Fig. 2b shows the element distribution mapping for V using 1% isoconcentration surfaces, which clearly sketches the shapes of carbides. The cross-section profiles in Fig. 2c show the chemical compositions of typical carbide presented in Fig. 2b, indicating the presence of VC carbide. It is seen that some of the VC carbides are exhibiting lamella-like shapes, which is invisible in SEM or TEM images because they can only present the cross-section morphologies in two dimensions. Furthermore, granular VCs are also observed in TEM images and APT
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Fig. 1 e (a, b) SEM images, (c, d) TEM images and (e, f) SAED patterns for the carbides in AR and QT specimens.
map, which was also reported for forged 2.25Cr-1Mo-0.25 V steels [32]. For the QT specimens, no such uniformly distributed carbides could be found in TEM images. Several groups of APT tests were performed but none of them indicated any trace of element segregation. It is believed that many of the fine VCs were decomposed during QT treatment. Fig. 3 shows the recrystallization maps, kernel average maps (KAM), phase maps, and inverse pole figure (IPF) maps for AR and QT specimens, respectively. Clearly, the recrystallization fraction (Fig. 3a and b) increases and dislocation density (Fig. 3c and d) decreases, which can lead to the decrease of residual stress in steel. The phase maps shown in Fig. 3e and f demonstrate the fully bainitic structure of the steel and no trace of retained M/A islands could be found, which was however reported in the forgings of 2.25Cr-1Mo0.25 V steel [32], in contrast with steel plate used in the present study. Furthermore, no significant difference of grain size and orientation distribution could be observed for the AR and QT specimens, as shown in Fig. 3g and h. Therefore, the effect of grain boundaries could also be ruled out.
Hydrogen content TDS tests were conducted for the AR specimens, as shown in Fig. 4a. A considerable amount of hydrogen was detected for the AR specimens, the content of which was determined as 0.94 ppm. Three local peaks were identified from the TDS curve, with the peak temperature as 206.9 C, 244.5 C and 362.6 C respectively. These peaks correspond to different
hydrogen traps in steel, and peaks at higher temperature correspond to deeper traps with higher binding energy. Furthermore, based on the binding energy between hydrogen atoms and the local trap sites, hydrogen traps can be classified into two categories: reversible and irreversible traps. All the hydrogen states shown in Fig. 4a can be classified as irreversible hydrogen because the as-received steel had already been aged in air for more than two years. Therefore, all the hydrogen states with peak temperature above 206.9 C can be classified as irreversible hydrogen for the following TDS curves, which were obtained using the identical testing parameters with that in Fig. 4a. Therefore, the AR specimens should be dehydrogenated before the study of HE. A group of dehydrogenation tests was conducted to determine the appropriate parameters to completely degas the hydrogen present in AR specimens, as shown in Fig. 4b (DeH for dehydrogenated specimens). By comparing the results between DeH-200 C and DeH-300 C, it is seen that hydrogen in steel tends to escape faster when it is aged at higher temperature. The irreversible hydrogen in steel could be fully degassed by aging at 350 C for 60 min, as indicated by the green line in Fig. 4b. Notably, considerable background noise could be detected especially in the high temperature range when the hydrogen content in specimen was very low. A custom-designed TDS instrument can be used to eliminate the background signal [33]. Therefore, all the AR specimens were dehydrogenated at 350 C for 60 min, both in TDS tests and tensile tests. Fig. 5a shows desorption curves after hydrogen charging for QT specimens and dehydrogenated AR specimens. The total hydrogen
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Many studies showed that the local peaks in TDS curves corresponding to the hydrogen trapped at dislocations were typically at low temperature with a peak temperature below 100 C [35e37]. Note that the peak temperature for a particular hydrogen trap tends to shift positively with the increase in the heating speed in TDS tests. A heating speed of 100 C/h was utilized in the present study, which is lower than those used in the studies on the trapping behavior of dislocations. Thus all the hydrogen trapped at dislocations should be desorbed below 100 C, which is however not detected in any of the TDS tests in present study. It is likely that most of the reversible hydrogen desorbed below 100 C had already escaped during the pre-processing procedure of specimens in high vacuum environment during the TDS tests. This phenomenon has also been reported elsewhere for some other materials [17]. Therefore, the hydrogen trapping at dislocations could be ruled out for the local peaks shown in Figs. 4 and 5. Although it is extremely difficult to build one-to-one correlation between the local peaks in Fig. 5 and the local microstructures in steel, it is assumed here that the peak 2 hydrogen is attributed to those trapped at the small-sized carbides such as VCs, while peak 1 hydrogen corresponds to the hydrogen at the other carbides with larger sizes, considering that small-sized precipitations tend to have stronger trapping ability of hydrogen atoms [17,19,38]. Whether the irreversible trapping sites indeed correspond to VC carbides still needs more solid evidence, because the desorption of hydrogen from VC carbides in V-bearing high strength steel has been reported [39].
Tensile tests
Fig. 2 e (a) High-resolution TEM images of fine carbides with needle-like and granular shapes, (b) element map for V with 1% iso-surfaces obtained through APT technique and (c) the corresponding cross-section profile of element distribution, in AR specimens.
content for DeH-HC (dehydrogenated and recharged with hydrogen) and QT-HC (quenched-tempered and hydrogencharged) specimens was 2.37 and 1.04 ppm respectively, with two local peaks identified, as shown in Fig. 5b and c. The peak 2 hydrogen in these two curves, with peak temperature being 214.6 and 212.9 C, can be confirmed as irreversible hydrogen, based on the evidence drawn from Fig. 4a. The peak 1 hydrogen, with peak temperature being 127.0 and 129.4 C, can be classified as reversible hydrogen. The results clearly indicate significant decease in the amount of both reversible and irreversible hydrogen, which is probably due to the decreased hydrogen diffusion coefficient after heat treatment as previously demonstrated by electrochemical permeation tests [34].
The results of tensile tests are presented in Fig. 6. It is found that the QT treatment results in significantly decreased strength and increased ductility. For the AR specimens, no obvious difference was observed after heat treatment at 350 C, which indicates that residual stress should have little effect on mechanical properties. This also demonstrates the feasibility of dehydrogenation treatment at 350 C, which should not have significant influence on the microstructures and mechanical properties. The small deviation between the above mentioned two curves is possibly due to the different locations of specimens which were randomly cut from the steel plate without considering the effect of plate thickness. For the DeH specimens, obvious ductility loss could be observed after hydrogen charging. The susceptibility to HE was decreased to a negligible level after QT treatment, as indicated by the dashed lines in Fig. 6. The HE resistance was apparently improved by QT treatment. This agrees well with the SEM fractography result. The SEM fractography images for AR specimens are shown in Fig. 7aed. In the absence of hydrogen, dimples and secondary cracks could be found on the fractured surface, indicating ductile fracture. However, the fractured surface for the hydrogen-charged specimens featured quasi-cleavages and decreased number of dimples. The transition from ductile fracture to brittle-ductile fracture after hydrogen charging was verified. For the QT specimens, the fractographies are all featured as ductile fracture whether charging or not, as shown in Fig. 7e and f.
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Fig. 3 e (a, b) Recrystallization, (c, d) KAM, (e, f) phase maps, and (g, h) IPF maps for AR and QT specimens.
Hydrogen embrittlement mechanism The variations of microstructures, hydrogen concentration and mechanical properties were confirmed prior to and after QT treatment. It would be necessary to provide an explanation for the improved HE resistance after QT treatment. First, significant decrease of strength after QT treatment was observed. The influence of strength on HE should be clarified, which is strongly affected by the microstructures of material. For example, the susceptibility to HE tends to increase with increasing strength for pipeline steels [40]. However, the opposite trends were reported for the grain refinement strengthened steels, in which the decrease of grain size results in high strength but the susceptibility to HE is decreased due to the reduced diffusive hydrogen per grain boundary area [41]. This contradiction also agrees well with the significant scatter that exhibits in the studies of HE for different grades of steels, as reviewed by Barnoush [42]. The variety of influencing factors leads to the scatter of data and one cannot just conclude whether the strength level is beneficial or detrimental to HE without taking all the influencing factors into account. The effect of strength on HE for high strength steels should be discussed in combination with the strengthening mechanisms of steel. Among all these factors that may determine the brittle fracture, the main factor should be always related to the principal strengthening mechanism for the involved high strength steel. For the AR and QT-treated 2.25Cr-1Mo0.25 V steel in this paper, the effect precipitation, dislocation density and grain size should be considered when discussing the effect of strength.
For the AR specimens, the numerous small-sized carbides in steel result in high strength due to the impediment to dislocation movement by precipitations, and concurrently high resistance to HE, due to the hydrogen trapping at carbides. After QT treatment, the impediment effect of dislocation movement by precipitations was weakened due to the coarsening of carbides, which results in the decrease of strength. At the same time, the HE resistance was improved due to the decrease of hydrogen concentration including both reversible and irreversible hydrogen, which was mainly attributed to the decomposition, precipitation and coarsening of carbides. It is believed that M3C acts as reversible trapping sites, while VC as irreversible trapping sites, although not rigorously proved by experiments. The decreased amount of reversible hydrogen due to the decomposition of M3C is mainly responsible for the improved HE resistance. This finding agrees well with the results reported in literature [38], which indicates that the increased number of reversible trapping sites by carbides can result in detrimental effect on HE. The decreased amount of hydrogen at VC (peak 2 hydrogen) is believed not to be responsible for the variation of HE sensitivity, as it is wellacknowledged that irreversible hydrogen has little effect on HE resistance [16,43]. Hydrogen charging method is also an important influencing factor when discussing the HE mechanism. The ex-situ hydrogen charging method was used in the present study, which enabled to rule out the effect of charging condition, considering that significantly different tensile time could be encountered for different specimens depending on the HE susceptibility for in-situ hydrogen charging. On the other
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Fig. 4 e TDS results for (a) AR specimens without hydrogen charging, and (b) AR specimens after different dehydrogenation treatments (DeH specimens).
hand, the ex-situ hydrogen charging also rules out the influence of hydrogen refilling, and this is crucial when discussing the effect of reversible and irreversible hydrogen. In in-situ hydrogen charging, the reversible hydrogen traps are immediately refilled with hydrogen. The amount of diffusive hydrogen in steel between AR and QT specimens changes during tensile tests and is not simply affected by the original microstructures any more, which makes the analysis of HE mechanism more complicated. However, the usage of ex-situ charging can rule out this influence. It is believed that QT can decrease the amount of hydrogen content including both reversible and irreversible hydrogen, as demonstrated by TDS tests. The improved HE resistance can be explained as follows. It is known that the presence of hydrogen can enhance the local plasticity by increasing the dislocation mobility, which is known as the hydrogen enhanced local plasticity (HELP) mechanism [44]. The increased dislocation mobility by hydrogen can be explained from the aspect of hydrogen
Fig. 5 e (a) TDS results after hydrogen charging for DeH specimens and QT specimens, with the total hydrogen concentration shown in the inset figure. The fitted results of these two curves are shown in (b) and (c) respectively.
shielding effect. The aggregated hydrogen at the stress fields of dislocations modifies the stress field to which it is attached, and reduces the interaction energy of the dislocation and the elastic obstacles allowing dislocation motion at lower stress level in certain directions [45]. The premise of this mechanism
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Fig. 6 e Stress-strain curves for AR, DeH, and QT specimens with and without hydrogen charging.
is that local hydrogen is present in sufficient concentration [46]. For the QT specimens, the amount of diffusive hydrogen was decreased, which weakens the enhancement effect on the dislocation mobility. No obvious ductility loss was observed because the local hydrogen concentration does not reach a critical level so that the dislocation mobility can be affected. Additionally, the applied stress is also a key factor for the HE mechanism and the critical hydrogen concentration that is required for the dislocation initiation is probably dependent
on the local stress concentration. Significant local stress level is expected for the AR and QT specimens due to the difference of strength. Whether this stress difference has significant influence on the critical hydrogen concentration still needs further investigation. Examining the HE resistance at different hydrogen concentrations for both AR and QT specimens would be a viable route. The susceptibility to HE is also strongly affected by strain rate. It is possible that apparent ductility loss could be observed for the QT-HC specimens at lower strain rate, although the strain rate in the present work is relatively low (3.6 104 s1). The hydrogen release during tensile tests also has influence on the final fracture, because the time to fracture for QT specimen is larger than that for AR specimen. The largest time difference for the tensile tests is estimated as 4.6 min. Compared with the influence of hydrogen concentration, we believe this is not the primary influencing factor. This influence can be ruled out by electroplating the specimens with a cadmium coating to prevent hydrogen release [47], which may be investigated in the future work. It is also known that the susceptibility to HE tends to increase at higher hydrogen concentrations [48e50]. For example, Zhu et al. [51], also reported a beneficial effect of tempering on HE for a quenchingepartitioningetempering (QP-T) treated steel, but the improvement of HE was observed only at very low hydrogen concentration and the effect was eliminated after further extending the charging duration. Therefore, exploring the HE resistance of QT specimens at higher hydrogen concentration would be an interesting topic. All the tests in this study were conducted at room temperature, thus the discovered phenomenon and the
Fig. 7 e SEM fractographies for (a) AR-NC (as-received, non-charged) and (b) AR-HC (as-received, hydrogen-charged) specimens, with the enlarged figures shown in (c) and (d) respectively; SEM fractographies for (e) QT-NC (quenchedtempered, non-charged) and (f) QT-HC (quenched-tempered, hydrogen-charged) specimens.
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corresponding HE mechanism may provide guidance for the hydrogen induced cracking during fabrication procedures or startups and shutdowns of reactors, which take place at low temperature. However, the difference between the hydrogen trapping behavior of carbides at room temperature and elevated temperature should be noted. For example, those carbides which are recognized as irreversible trapping sites at room temperature may become reversible at elevated temperature. According to the TDS results for the dehydrogenated specimens in Fig. 4, the hydrogen which is irreversible at room temperature can be completely degassed at 350 C. This implies that all the hydrogen atoms in steel are probably diffusible and can contribute to the macroscale ductility loss, in a high temperature hydrogen environment. Additionally, the interaction mechanism of hydrogen with dislocations may also become different at elevated temperature. The hydrogen atoms may impede the motion of the dislocations by acting as solid solution strengthening agent because they do not remain bound to dislocations at elevated temperature [45].
Conclusions The effect of quenching-tempering (QT) treatment on the hydrogen embrittlement (HE) of 2.25Cr-1Mo-0.25 V steel was investigated through SEM, TEM, APT, TDS, and tensile tests. The following conclusions were obtained: (1) The QT treatment resulted in the decrease of strength, significant increase of ductility and apparently enhanced HE resistance. (2) Two hydrogen states were identified in the hydrogen desorption profiles in TDS tests, with peak 1 hydrogen desorbed at low temperature (~127 C), and peak 2 hydrogen at high temperature (~215 C). It was demonstrated that peak 2 hydrogen is irreversible hydrogen by comparing its peak temperature with that detected in the AR specimens which had been aged in air for 2 years. The hydrogen content of both reversible and irreversible hydrogen decreased after QT treatment. (3) QT treatment resulted in the decomposition of M3C/VC and precipitation of M7C3. The VCs are more probably irreversible traps, corresponding to peak 2 in TDS curves, while the other carbides are reversible traps, corresponding to peak 1. The decrease of reversible hydrogen is believed to be responsible for the improved HE resistance.
Acknowledgements This work is supported by the National Basic Research Program of China (973 Program, Grant No.2015CB057602), China Postdoctoral Science Foundation (Grant No. BX20180245, 2018M643637) and Fundamental Research Funds for the Central Universities (Grant No. XZY012019024).
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