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Evaluation of the hydrogen embrittlement susceptibility in DP steel under static and dynamic tensile conditions
International Journal of Impact Engineering 123 (2019) 118–125
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Evaluation of the hydrogen embrittlement susceptibility in DP steel under static and dynamic tensile conditions
T
⁎
T. Depovera, , F. Vercruysseb, A. Elmahdyb, P. Verleysenb, K. Verbekena a b
Department of Materials, Textiles and Chemical Engineering, Ghent University (UGent), Technologiepark 903, B-9052 Zwijnaarde, Belgium Department of Electrical Energy, Metals, Mechanical Constructions & Systems, Ghent University (UGent), Technologiepark 903, B-9052 Zwijnaarde, Belgium
A R T I C LE I N FO
A B S T R A C T
Keywords: Hydrogen embrittlement Static and dynamic tensile testing Hopkinson Strain rate Hydrogen diffusivity DP steel
Hydrogen (H) induced mechanical degradation is studied in DP steel by performing tensile tests under static and dynamic conditions. Tensile specimens were electrochemically H charged and tensile tests were done ex-situ after charging. Different H contents were charged into the samples by modifying the current density. The strain rate is increased from static (1.67*10−2 and 1.67 s−1) to dynamic (450 and 900 s−1) conditions to verify the effect of H diffusivity during the tensile tests on the hydrogen embrittlement (HE) susceptibility. Therefore, a reproducible methodology was established by using a standardized tensile machine for static testing and split Hopkinson bar experiments for dynamic conditions. The HE degree increased with current density due to higher amount of H, as confirmed by melt extraction. The HE% also increased with slower strain rates since H was able to diffuse to a crack tip, hence accelerating failure. Even at the highest strain rate (900 s−1), the material lost about 10% of its ductility due to H present in the sample and not because of H diffusion during the test. This was concluded since H induced brittle failure initiated at the edges of the samples at slow strain rates. Though at a strain rate of 1.67 s−1, fracture initiated in a ductile way from the center similarly as for tests performed without charging. Fractographic visualization of the fracture surfaces revealed an embrittled central line when charged with H, which evolved into a major crack. MnS inclusions were found in this central line accounting for the H induced crack initiation.
1. Introduction In the automotive industry, advanced high strength steels (AHSS), such as dual phase (DP) steels, are widely used due to their outstanding combination of both formability and strength. The use of these AHSS has been stimulated since these materials can both guarantee an increased safety together with the required weight reduction to meet the rigorous CO2 emission regulations. DP steel has a ferrite-martensite microstructure with low yield strength, but a high work hardening exponent at low and moderate deformation levels leading to a high ultimate tensile strength. The typical characteristics of DP steel are determined by its thermo-mechanical processing which incorporates a volume fraction of austenitic islands transforming into martensite. Hence, the corresponding strength level is predominantly dependent on the volume fraction and the hardness of the martensite phase in the DP structure. Additionally, a volumetric expansion of martensite into the ferrite matrix during production results into residual stresses at the interface between martensite and ferrite. Therefore, a high density of mobile dislocations is present in the soft ferrite contributing to excellent
⁎
continuous yielding. This means that there is no yield point phenomenon and no formation of Lüders bands. Although the mechanical properties of AHSS are excellent, they are unfortunately considered to be prone to hydrogen (H) induced mechanical degradation. Moreover, further applicability of AHSS has been reported to be hindered by the H phenomenon [1]. Essentially, in the presence of H, these materials undergo a large ductility loss, mainly referred to as hydrogen embrittlement (HE). Amongst others, DP, transformation induced plasticity (TRIP) and high strength low alloyed (HSLA) steels are widely used grades in the automotive industry and their interaction with H has therefore already been subject of numerous researches [2–7]. The detrimental effect of H on the mechanical properties was specifically considered by Depover et al. [8]. A rather low HE sensitivity of about 8% was found for a HSLA steel which was correlated to the beneficial presence of Ti and Nb based carbo-nitrides. A DP and TRIP steel, however, lost about 54% and 60% of their initial ductility, respectively. Duprez et al. [9] also performed tensile tests on these materials immediately after electrochemical H charging and on samples that were atmospherically discharged for one week. Their main
Corresponding author. E-mail address: [email protected] (T. Depover).
https://doi.org/10.1016/j.ijimpeng.2018.10.002 Received 7 March 2018; Received in revised form 1 October 2018; Accepted 2 October 2018 Available online 06 October 2018 0734-743X/ © 2018 Elsevier Ltd. All rights reserved.
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conclusion was that the H induced ductility loss for DP and TRIP steel was mainly reversible since a large part of the ductility was recovered after discharging the tensile specimen for one week, whereas for the HSLA steel, the difference between immediate and after one week testing was limited. These observations were further correlated with the findings by Pérez Escobar et al. [10], who performed a thermal desorption spectroscopy (TDS) study on these high strength steels. Different times in vacuum, before the TDS measurement started, were applied to evaluate the H effusion. They detected that H present in TRIP and DP steel was mainly diffusible, leaving the material easily when kept in vacuum, whereas the opposite observations were made for HSLA. Additionally, the obtained activation energies for the different traps in these AHSS were in the same range, indicating the complex correlation between a certain microstructural feature and a specific obtained desorption peak. Only the TRIP steel showed a high temperature peak with an Ea of 90 kJ/mol which was proved to come from hydrogen trapped in the retained austenite [11]. Furthermore, a detailed microstructural analysis was performed on this material by Laureys et al. [4, 5] showing that crack initiation occurred in the martensitic islands, mainly along the martensitic lath interfaces, and that crack propagation was mainly stress driven. Nevertheless, the significant contrast in mechanical behavior between the ferrite and martensite phase in DP steels induces certain jeopardies, especially in the presence of H [12]. Sun et al. [13], attributed the relative high sensitivity of DP steels to HE to sub-structural changes, because cracks initiated along the interface and lath boundaries when H charging was applied. However, Gu et al. [14] considered the delayed fracture properties of a H charged 1500 MPa bainite-martensite DP steel and revealed that its crack growth rate is less than that of conventional quenched and tempered high strength steels. They ascribed this to the fine DP microstructure and the lath boundaries which may act as beneficial H traps, slowing down segregation and diffusion of H to the crack tip. Recently, Koyama et al. [15] studied the microcracking behavior in martensite-ferrite DP in the presence of H. They demonstrated that H affected the damage evolution process considerably. Both the H-enhanced decohesion (HEDE) and H-enhanced localized plasticity (HELP) mechanisms were considered to contribute to the damage evolution in DP steels leading to an unpredictable HE failure. The impact of H diffusion was also recently visualized in DP steel by studying the fracture surfaces of an in-situ charged tensile specimen [16]. The calculated H diffusion distance matched perfectly with the observed transition between H induced brittle and ductile fracture. AHSS are often used in body-in-white vehicle structures due to their weight reducing capacity and increased safety guarantee. Crashworthiness is a key parameter of their applicability and many mechanical parameters are strain rate dependent. Therefore, the effect of strain rate on hardening and fracture of AHSS has been studied extensively. Static and dynamic tensile tests on two DP and two TRIP steels performed by Huh et al. [17] showed an increase of the elongation at fracture for one of the TRIP steels, though strain rate independent values for the other steels. Similar results were obtained by Curtze et al. [18] who concluded that over a wide range of strain rates relatively constant values of the fracture elongation are obtained for a TRIP and DP steel. However, the combination of H induced mechanical degradation and dynamic high strain rate conditions has not been given much attention so far. Therefore, the aim of this work is to investigate the effect of H on the impact properties of a specific DP steel grade. This H influence was evaluated by modifying both the H content in the tensile specimens and the applied strain rate going from static to dynamic conditions.
Table 1 Chemical composition of the used materials in wt%. Material/Element
C
Mn
Si
Other
DP
0.07
1.50
0.25
0.4%−0.8% Cr + Mo
2. Experimental procedure 2.1. Material characterization The sensitivity of DP600 steel to H induced mechanical degradation was investigated in this work. Hot and cold rolling was done till a final thickness of 1.2 mm, which was followed by subsequent annealing via industrial annealing parameters required to attain the desired microstructure. The chemical composition is presented in Table 1. The microstructure was characterized using light optical microscopy (LOM) for which the material was grinded, polished and finally etched using a 2% Nital solution. The samples were immediately cleaned using methanol and acetone. The DP600 steel is a ferritic-martensitic dual phase steel with approximately 23.6% of martensite. The corresponding grain sizes were about 7 μm for the ferritic phase and about 2 μm for the martensitic phase, as shown in Fig. 1. Tensile samples were machined with their tensile axis parallel to the rolling direction. A dogbone-shaped specimen geometry was used, as illustrated in Fig. 2. The small size of the samples is required for the dynamic tests [19]. However, in order to exclude effects related to sample geometry, the same sample geometry is used for the static tests as well. 2.2. Determination of total hydrogen content The total H content was determined by melt extraction at 1600 °C. Samples were H charged electrochemically in 1 g/l thiourea 0.5 M H2SO4 electrolyte for two hours, which guaranteed saturation without H induced damage, as shown in [16]. The applied current density was modified to induce different amounts of H into the tensile samples. For absorption from an aqueous solution, a relation is found between H solubility and applied current density during electrochemical charging, i.e. the solubility is proportional to the square root of the current density [20]. Therefore, the current densities increased from 0.8 over 5 to 25 mA/cm2. The time between H charging and H detection was kept constant at 2 min due to the experimental procedure used to perform both the static and dynamic tensile tests, which also required 2 min between H charging and tensile testing. The system to determine the H content includes a pulse furnace in which a pre-weighted sample is
Fig. 1. Light optical microscopic image of DP600 steel. 119
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Table 2 Total amount of hydrogen determined by melt extraction for different applied current densities. H content
Air
0.8 mA/cm2
5 mA/cm2
25 mA/cm2
wppm
0.52 ± 0.08
3.77 ± 0.12
4.67 ± 0.11
4.97 ± 0.13
were taken during the tests to capture fracture initiation and evaluate the fracture mode upon the applied strain rate. 3. Hydrogen induced mechanical degradation 3.1. Hydrogen uptake capacity Fig. 2. Geometry of the tensile specimens (in mm).
heated to 1600 °C. The sample releases its H as gaseous H2 and is dragged along by a nitrogen flow. This mixture (N2–H2) is directed to a thermal conductivity measuring cell. The thermal conductivity of the mixture depends on the H2 concentration since the conductivity of H2 and N2 differs significantly. Hence, the H content is calculated based on the difference in thermal conductivity.
Melt extraction was done to determine the total H content, for the different H charging methods as described above. Different current densities were applied and the amount of H increased with applied current density (cf. Table 2), which was observed elsewhere as well [20]. The tensile specimens were charged with H under identical conditions. Hence, the actual amount of H present in the sample is known when the tensile test started.
2.3. Determination of the hydrogen induced mechanical degradation
3.2. Tensile tests performed under static conditions
The H induced ductility loss was verified by performing tensile tests on uncharged and H charged samples while using the abovementioned H charging parameters. The H impact was quantified by the HE%, as defined in Eq. (1), with εch and εun being the engineering strains at fracture averaged over the sample gauge section of the H charged and uncharged sample, respectively. Hence, the HE degree can vary between 0 and 1, with 0 indicating no ductility drop and insensitivity to HE. When an index of 1 is obtained, the ductility drop is 100% and the HE is maximal.
Tensile tests were performed on uncharged and H charged samples at current densities of 0.8, 5 and 25 mA/cm2. The corresponding stress–strain curves of the tests done at strain rates of 0.0167 and 1.67 s−1 are presented in Figs. 3 and 4, respectively. The degree of HE was determined and summarized for all conditions in Table 3. Generally, the sensitivity to HE increases with higher applied current density, which can be attributed to the increased H content (cf. Table 2) [24]. This corresponds to previously obtained results and confirms the correlation between the HE susceptibility and the H amount present in the sample [25–28]. A slight increase in flow stress was also observed for the H charged tensile tests, which was attributed solid solution strengthening, as H might pin the dislocations. This was also observed in [9, 29]. Furthermore, the HE resistance increased when a higher strain rate was applied. This was correlated to the H diffusion distance which decreases with higher strain rate, as described in [16, 29]. The distance x (cm) over which H can diffuse, can be calculated by taking the square root of the product of the diffusion coefficient D (cm2/s) and the time t (s) of the tensile test, i.e. x = (D x t)1/2 [15, 29]. Therefore, during a test performed at strain rate of 0.0167 s−1, from the onset of deformation till fracture, H was able to diffuse over approximately 30 μm, whereas at 1.67 s−1, H was just able to diffuse over about 3 μm. Towards the end of the tensile tests, a crack is formed. Close to the
ɛ % HE = 100 ⎛1 − ch ⎞ ɛ un ⎠ ⎝ ⎜
⎟
(1)
Initially, the HE sensitivity was studied by static tensile tests performed on both uncharged and H charged specimens using a standardized Instron tensile machine. H charging was done for two hours and tensile tests were executed ex-situ due to experimental requirements. A constant time lag of 2 min between H charging and tensile testing was maintained. The observations were correlated with those obtained by the melt extraction, which was also done 2 min after H charging. In a first series of static tests, the impact of the H content on the HE susceptibility was evaluated by using samples which were H charged at different applied current densities. Secondly, the effect of H diffusivity was verified by considering two different –though still static- strain rates: 0.0167 and 1.67 s−1. Additionally, the strain rate was further increased till dynamic conditions. These high strain rate tests were done on a split Hopkinson tensile bar (SHTB) setup [21, 22]. For this purpose, the specimen was pinned into slots between two long bars, the input and output bar. The impact of a projectile generates a tensile loading wave at the free end of the input bar which propagates along the input bar towards the sample where it interacts with the specimen resulting in a high strain rate tensile loading of the sample. The loading wave is partly reflected back to the input bar and partly transmitted towards the output bar. Strain gauges on the bars measure the strain corresponding with the loading, reflected and transmitted waves. From those waves, the total force and elongation history of the tensile sample can be determined based on the principles of 1-dimensional elastic-wave propagation in bars [23]. The dynamic tests were done at average strain rates of ± 450 and ± 900 s−1. As such, the effect of H can be elucidated at crash impact velocities and the impact of H diffusivity can be analyzed on the fracture mode. High speed camera (HSC) images of the deforming sample
Fig. 3. Engineering stress–strain diagrams of static tensile tests performed at a strain rate of 0.0167 s−1. 120
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diffraction. Nevertheless, a different failure mode was observed when the strain rate was increased to 1.67 s−1 after H charging at 0.8 mA/cm2, as illustrated in Fig. 7. The corresponding fracture initiation occurred clearly in the center of the tensile specimen, similarly to the tensile tests performed on uncharged material (cf. Fig. 5). Basically, the fracture initiation occurred in the center for all tests performed in the uncharged condition at both strain rates, whereas only the H charged samples tested at the higher strain rate of 1.67 s−1 showed fracture initiation in the center. An overview of the fracture initiation point for all tests is summarized in Table 4. A transition of fracture initiation from the edge to the center was observed as a function of strain rate for H charged samples. Therefore, this transition in fracture mode, when the strain rate was increased, can be correlated with a decrease in available H diffusion time and consequently a decreased H diffusion distance x. Zones which are characterized by a stress concentration ahead of a growing crack tip attract H from the neighboring H saturated matrix. The probability for this to occur is increased at the slower strain rate. The fracture initiation transition was observed between the strain rates of 0.0167 s−1 and 1.67 s−1, which is correlated to diffusion distances of about 30 and 3 μm, respectively.
Fig. 4. Engineering stress–strain diagrams of static tensile tests performed at a strain rate of 1.67 s−1. Table 3 Summary of the degrees of H induced mechanical degradation. HE% −1
0.0167s 1.67s−1
0.8 mA/cm²
5 mA/cm²
25 mA/cm²
45 27
58 37
72 45
3.3. Tensile tests performed under dynamic conditions Tensile tests at dynamic conditions were done on uncharged and H charged samples at current densities of 0.8, 5 and 25 mA/cm2. The corresponding stress–strain diagrams of SHTB tests performed at a strain rate of 450 and 900 s−1 are displayed in Fig. 8 and Fig. 9, respectively. The resultant degrees of HE are given in Table 5. Even in these dynamic conditions, a clear impact of H on the mechanical properties was observed, i.e. the ductility dropped with about 10%. In agreement with the static tests, the sensitivity to H induced mechanical degradation increased slightly with higher current densities, confirming the link between the HE% and the H content [25–28]. Moreover, a slightly lower degree of HE was obtained when the strain rate was increased. However, linking this observation to the diffusion distance x, over which H can diffuse during the tensile test, is not appropriate as the estimated H diffusion distances during the entire test were only about 0.17 and 0.23 μm at strain rates of 900 and 450 s−1, respectively. The fracture initiation during the dynamic tensile tests was again captured by HSC and the image sequence, just before fracture, was considered. The specimens tested after H charging at 25 mA/cm2 at a strain rate of 450 s−1 and 900 s−1 are presented in Figs. 10 and 11, respectively. For both, fracture initiation was found to occur in the center of the specimen. This is in agreement with the findings demonstrated above (cf. Table 4). Generally, fracture initiation occurred in the center for all tests performed in the uncharged condition and for the H charged samples tested at the higher strain rate of 1.67 s−1. Since dynamic conditions were applied, failure initiation from the center was expected to occur since a transition of fracture initiation was observed
crack tip, a zone characterized by a stress concentration is generated. H can diffuse to this zone and hence accelerate failure. When the applied strain rate was increased, H showed a smaller potential to assist the failure process since the distance it can diffuse during the fracture process is limited. The fracture initiation in the static tensile tests was studied in detail. The initiation of failure was captured by HSC and a sequence of the images, just before fracture of the tensile test, were further analyzed. The corresponding images of the test performed at a strain rate of 1.67 s−1 on an uncharged sample are presented in Fig. 5. Considerable diffuse necking can be observed. Moreover, the final fracture initiation was located at the center of the specimen, indicating a ductile failure mode affected by the high stress triaxiality in the center of the tensile sample. The failure sequence of the H charged (at current density of 5 mA/cm2) specimen tested at 0.0167 s−1 is shown in Fig. 6. As opposed to the uncharged sample, localized necking occurred and the final fracture initiated at the edge of the sample. This was attributed to the damaging effect of H on the material at the sample edge where lattice H preferentially diffuses to. Therefore, this H is able to embrittle the material locally, leading to failure initiation and a further accelerated fracture propagation from the edge onwards. These observations were also made for in-situ H charged tensile specimens of DP and TRIP steel [4, 5, 16], on which H induced failure initiation and propagation was evaluated by SEM analysis and electron backscatter
Fig. 5. HSC images of tensile test performed in air at strain rate of 1.67 s−1 during fracture. 121
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Fig. 6. HSC images of tensile test performed after H charging (5 mA/cm2) at strain rate of 0.0167 s−1 during fracture.
Fig. 7. HSC images of tensile test performed after H charging (0.8 mA/cm2) at strain rate of 1.67 s−1 during fracture. Table 4 Overview of fracture modes in the static tensile tests. Fracture initiation point
Air
0.8 mA/cm²
5 mA/cm²
25 mA/cm²
0.0167s−1 1.67s−1
Center Center
Edge Center
Edge Center
Edge Center
Fig. 9. Engineering stress–strain diagrams of dynamic SHTB tests performed at a strain rate of 900 s−1. Table 5 Summary of the HE% for the dynamic tensile test conditions.
Fig. 8. Engineering stress–strain diagrams of dynamic SHTB tests performed at a strain rate of 450 s−1.
HE%
0.8 mA/cm²
5 mA/cm²
25 mA/cm²
450s−1 900s−1
9 8
11 9
16 11
properties [10, 25]. However, as shown by melt extraction also performed 6 min after H charging, the major part of H is still present in the material: the H content decreased to 3.1 wppm. The 0.67 wppm, released as compared to the original tests (cf. Table 2), was confirmed to be H present near dislocations as demonstrated in our earlier work [25, 30–34]. The corresponding stress–strain diagram for the referred specimen (*Charged) was added to the stress–strain curves for the samples
as a function of strain rate for H charged samples. To verify the reliability of the obtained results of the dynamic testing and to understand the effect of H on the mechanical properties in more detail, identical SHTB tensile tests were performed 6 min after electrochemical H charging at 0.8 mA/cm2 instead of the applied 2 min for the previous tests. As such, part of the mobile H had been discharged from the tensile specimen and could no longer influence the mechanical 122
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Fig. 10. HSC images of tensile test performed after H charging (25 mA/cm2) at strain rate of 450 s−1 during fracture.
Fig. 11. HSC images of tensile test performed after H charging (25 mA/cm2) at strain rate of 900 s−1 during fracture.
35]. The fracture surface of the uncharged and H charged sample at 25 mA/cm², both tensile tested at 1.67 s−1, is presented in Fig. 13. Comparison between both showed a delamination in the central region for the H charged sample. Furthermore, the fracture surfaces of two dynamic tests at 450 s−1 and H charged at 0.8 and 25 mA/cm2 are presented in Fig. 14. A central crack perpendicular to the fracture surface can again be observed. Its size increased with increasing current density, i.e. it is linked to the higher H content. The brittle zone around the crack exhibited some shear characteristics. The crack width also increased slightly with increasing strain rate. Fig. 15 shows a detailed characterization of the central delamination crack in the fracture surface of the tensile test performed at 900 s−1 after H charging at 5 mA/ cm2. Second phase particles were observed and EDX analysis confirmed it were MnS inclusions, which are known to be specifically harmful in the presence of hydrogen [16, 35]. When the fracture surfaces of the static and dynamic tensile tests are compared, similar conclusions can be drawn: a dimpled, ductile fracture surface with brittle features at the MnS containing central segregation line. However, no extensive delamination crack was present for the static tests, while a clear crack propagating throughout the sample center was detected for the dynamic tests.
Fig. 12. Engineering stress–strain diagrams of dynamic SHTB tests performed at a strain rate of 450 s−1, the intermittent time between H charging and tensile testing was 2 (charged) and 6 min (*charged).
tested after 2 min of intermittent time at strain rate of 450 s−1, as presented in Fig. 12. A similar mechanical response compared to the uncharged specimen was obtained, indicating that the HE effect was negligible although the H content was still about 3.1 wppm. This confirmed the crucial importance of mobile H trapped at dislocations in the HE phenomenon [25–28].
5. Conclusion The H induced mechanical degradation in DP steel was evaluated by performing tensile tests under static and dynamic conditions. The samples were H charged till saturation and tensile tests were subsequently done ex-situ. The applied current density was modified to induce different H amounts into the tensile specimens. The effect of the H diffusivity during the test on the HE sensitivity was estimated by increasing the strain rate from static (1.67*10−2 and 1.67 s−1) to dynamic (450 and 900 s−1) conditions. The HE susceptibility increased with applied current density due to higher amounts of H, as determined
4. Fractographic analysis The fracture surfaces of the tensile specimens were analyzed by SEM and a ductile dimple pattern was generally observed. However, a thin brittle zone was present at the center of the sample due to the typical segregation line for DP steels, as already demonstrated elsewhere [16, 123
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Fig. 13. SEM images of fracture surface of the uncharged (left) and specimen H charged at 25 mA/cm² (right), tensile tested at 1.67 s−1.
Fig. 14. SEM images of fracture surface of specimen H charged at 0.8 mA/cm² (left) and at 25 mA/cm² (right), tensile tested at 450 s−1.
Fig. 15. MnS inclusions in the H induced crack observed in the central segregation line on the fracture surface of the specimen H charged at 5 mA/cm2, tensile tested at 900 s−1, as revealed by EDX analysis (right).
in the tests performed in air. Fractographic analysis revealed a brittle central line when H charged, which evolved into a major crack which size increased with applied strain rate. EDX analysis revealed the presence of MnS inclusions in this segregation line, initiating the H induced cracks.
by melt extraction. Moreover, the HE degree was increased with slower strain rates. This was linked to higher probability of H to diffuse to regions of stress concentration ahead of a crack tip at slower strain rate and as such accelerating failure. Even at the highest strain rate (900 s−1), the material lost about 10% of its ductility, which was related with H present in the samples and not due to H diffusion during testing. This was concluded since H induced failure initiated at the edges of the samples at slow strain rates. However, at a strain rate of 1.67 s−1, fracture initiated in a ductile way from the center similarly as
Acknowledgments TD holds a postdoctoral fellowship via grant no. BOF01P03516. The 124
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authors wish to thank the Special Research Fund (BOF), UGent (BOF15/ BAS/06). The authors also acknowledge the technical staff from the Department Materials, Textiles and Chemical Engineering and the Department of Electrical Energy, Metals, Mechanical Constructions & Systems, UGent, for their help with experiments and/or sample preparation.
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References [1] Loidl M, Kolk O. Hydrogen embrittlement in HSSs limits use in lightweight body. Adv Mater Processes 2011;169:22–5. BMW Group. [2] Ryu JH, Chun YS, Lee CS, Bhadeshia HKDH, Suh DW. Effect of deformation on hydrogen trapping and effusion in TRIP-assisted steel. Acta Mater 2012;60:4085–92. [3] Sun J, Jiang T, Sun Y, Wang Y, Liu Y. A lamellar structured ultrafine grain ferritemartensite DP steel and its resistance to HE. J Alloys Compd 2017;698:390–9. [4] Laureys A, Depover T, Petrov R, Verbeken K. Characterization of hydrogen induced cracking in TRIP-assisted steels. Int J Hydrogen Energy 2015;40(47):16901–12. [5] Laureys A, Depover T, Petrov R, Verbeken K. Microstructural characterization of hydrogen induced cracking in TRIP-assisted steel by EBSD. Mater Charact 2016;112:169–79. [6] Liu Q, Venezuela J, Zhang M, Zhou Q, Atrens A. Hydrogen trapping in some advanced high strength steels. Corros Sci 2016;111:770–85. [7] Ronevich JA, Speer JG, Matlock DK. Hydrogen embrittlement of commercially produced advanced high strength steels. SAE Int J Mater Manuf 2010;3:255–67. [8] Depover T, Pérez Escobar D, Wallaert E, Zermout Z, Verbeken K. Effect of in-situ hydrogen charging on the mechanical properties of advanced high strength steels. Int J Hydrogen Energy 2014;39:4647–56. [9] Duprez L, Verbeken K, Verhaege M. Effect of hydrogen on the mechanical properties of multiphase high strength steels. Proceedings of the International Hydrogen Conference. 2008. [10] Pérez Escobar D, Verbeken K, Duprez L, Verhaege M. Evaluation of hydrogen trapping in high strength steels by thermal desorption spectroscopy. Mat Sci and Eng A 2012;551:50–8. [11] Pérez Escobar D, Depover T, Wallaert E, Duprez L, Verbeken K, 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;60:2593–605. [12] Davies RG. Influence of martensite content on the hydrogen embrittlement of dualphase steels. Scr Mater 1983;17:889–92. [13] Sun S, Gu J, Chen N. The influence of hydrogen on the sub-surface of the martensite and ferrite dual-phase steel. Scr Mater 1989;23:1735–8. [14] Gu JL, Chang KD, Fang HS, Bai BZ. Delayed fracture properties of 1500 MPa bainite/martensite dual-phase high strength steel and its hydrogen traps. ISIJ Int. 2002;42:1560–4. [15] Koyama M, Tasan CC, Akiyama E, Tsuzaki K, Raabe D. Hydrogen-assisted decohesion and localized plasticity in dual-phase steel. Acta Mater 2014;70:174–87.
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International Journal of Impact Engineering journal homepage: www.elsevier.com/locate/ijimpeng
Evaluation of the hydrogen embrittlement susceptibility in DP steel under static and dynamic tensile conditions
T
⁎
T. Depovera, , F. Vercruysseb, A. Elmahdyb, P. Verleysenb, K. Verbekena a b
Department of Materials, Textiles and Chemical Engineering, Ghent University (UGent), Technologiepark 903, B-9052 Zwijnaarde, Belgium Department of Electrical Energy, Metals, Mechanical Constructions & Systems, Ghent University (UGent), Technologiepark 903, B-9052 Zwijnaarde, Belgium
A R T I C LE I N FO
A B S T R A C T
Keywords: Hydrogen embrittlement Static and dynamic tensile testing Hopkinson Strain rate Hydrogen diffusivity DP steel
Hydrogen (H) induced mechanical degradation is studied in DP steel by performing tensile tests under static and dynamic conditions. Tensile specimens were electrochemically H charged and tensile tests were done ex-situ after charging. Different H contents were charged into the samples by modifying the current density. The strain rate is increased from static (1.67*10−2 and 1.67 s−1) to dynamic (450 and 900 s−1) conditions to verify the effect of H diffusivity during the tensile tests on the hydrogen embrittlement (HE) susceptibility. Therefore, a reproducible methodology was established by using a standardized tensile machine for static testing and split Hopkinson bar experiments for dynamic conditions. The HE degree increased with current density due to higher amount of H, as confirmed by melt extraction. The HE% also increased with slower strain rates since H was able to diffuse to a crack tip, hence accelerating failure. Even at the highest strain rate (900 s−1), the material lost about 10% of its ductility due to H present in the sample and not because of H diffusion during the test. This was concluded since H induced brittle failure initiated at the edges of the samples at slow strain rates. Though at a strain rate of 1.67 s−1, fracture initiated in a ductile way from the center similarly as for tests performed without charging. Fractographic visualization of the fracture surfaces revealed an embrittled central line when charged with H, which evolved into a major crack. MnS inclusions were found in this central line accounting for the H induced crack initiation.
1. Introduction In the automotive industry, advanced high strength steels (AHSS), such as dual phase (DP) steels, are widely used due to their outstanding combination of both formability and strength. The use of these AHSS has been stimulated since these materials can both guarantee an increased safety together with the required weight reduction to meet the rigorous CO2 emission regulations. DP steel has a ferrite-martensite microstructure with low yield strength, but a high work hardening exponent at low and moderate deformation levels leading to a high ultimate tensile strength. The typical characteristics of DP steel are determined by its thermo-mechanical processing which incorporates a volume fraction of austenitic islands transforming into martensite. Hence, the corresponding strength level is predominantly dependent on the volume fraction and the hardness of the martensite phase in the DP structure. Additionally, a volumetric expansion of martensite into the ferrite matrix during production results into residual stresses at the interface between martensite and ferrite. Therefore, a high density of mobile dislocations is present in the soft ferrite contributing to excellent
⁎
continuous yielding. This means that there is no yield point phenomenon and no formation of Lüders bands. Although the mechanical properties of AHSS are excellent, they are unfortunately considered to be prone to hydrogen (H) induced mechanical degradation. Moreover, further applicability of AHSS has been reported to be hindered by the H phenomenon [1]. Essentially, in the presence of H, these materials undergo a large ductility loss, mainly referred to as hydrogen embrittlement (HE). Amongst others, DP, transformation induced plasticity (TRIP) and high strength low alloyed (HSLA) steels are widely used grades in the automotive industry and their interaction with H has therefore already been subject of numerous researches [2–7]. The detrimental effect of H on the mechanical properties was specifically considered by Depover et al. [8]. A rather low HE sensitivity of about 8% was found for a HSLA steel which was correlated to the beneficial presence of Ti and Nb based carbo-nitrides. A DP and TRIP steel, however, lost about 54% and 60% of their initial ductility, respectively. Duprez et al. [9] also performed tensile tests on these materials immediately after electrochemical H charging and on samples that were atmospherically discharged for one week. Their main
Corresponding author. E-mail address: [email protected] (T. Depover).
https://doi.org/10.1016/j.ijimpeng.2018.10.002 Received 7 March 2018; Received in revised form 1 October 2018; Accepted 2 October 2018 Available online 06 October 2018 0734-743X/ © 2018 Elsevier Ltd. All rights reserved.
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conclusion was that the H induced ductility loss for DP and TRIP steel was mainly reversible since a large part of the ductility was recovered after discharging the tensile specimen for one week, whereas for the HSLA steel, the difference between immediate and after one week testing was limited. These observations were further correlated with the findings by Pérez Escobar et al. [10], who performed a thermal desorption spectroscopy (TDS) study on these high strength steels. Different times in vacuum, before the TDS measurement started, were applied to evaluate the H effusion. They detected that H present in TRIP and DP steel was mainly diffusible, leaving the material easily when kept in vacuum, whereas the opposite observations were made for HSLA. Additionally, the obtained activation energies for the different traps in these AHSS were in the same range, indicating the complex correlation between a certain microstructural feature and a specific obtained desorption peak. Only the TRIP steel showed a high temperature peak with an Ea of 90 kJ/mol which was proved to come from hydrogen trapped in the retained austenite [11]. Furthermore, a detailed microstructural analysis was performed on this material by Laureys et al. [4, 5] showing that crack initiation occurred in the martensitic islands, mainly along the martensitic lath interfaces, and that crack propagation was mainly stress driven. Nevertheless, the significant contrast in mechanical behavior between the ferrite and martensite phase in DP steels induces certain jeopardies, especially in the presence of H [12]. Sun et al. [13], attributed the relative high sensitivity of DP steels to HE to sub-structural changes, because cracks initiated along the interface and lath boundaries when H charging was applied. However, Gu et al. [14] considered the delayed fracture properties of a H charged 1500 MPa bainite-martensite DP steel and revealed that its crack growth rate is less than that of conventional quenched and tempered high strength steels. They ascribed this to the fine DP microstructure and the lath boundaries which may act as beneficial H traps, slowing down segregation and diffusion of H to the crack tip. Recently, Koyama et al. [15] studied the microcracking behavior in martensite-ferrite DP in the presence of H. They demonstrated that H affected the damage evolution process considerably. Both the H-enhanced decohesion (HEDE) and H-enhanced localized plasticity (HELP) mechanisms were considered to contribute to the damage evolution in DP steels leading to an unpredictable HE failure. The impact of H diffusion was also recently visualized in DP steel by studying the fracture surfaces of an in-situ charged tensile specimen [16]. The calculated H diffusion distance matched perfectly with the observed transition between H induced brittle and ductile fracture. AHSS are often used in body-in-white vehicle structures due to their weight reducing capacity and increased safety guarantee. Crashworthiness is a key parameter of their applicability and many mechanical parameters are strain rate dependent. Therefore, the effect of strain rate on hardening and fracture of AHSS has been studied extensively. Static and dynamic tensile tests on two DP and two TRIP steels performed by Huh et al. [17] showed an increase of the elongation at fracture for one of the TRIP steels, though strain rate independent values for the other steels. Similar results were obtained by Curtze et al. [18] who concluded that over a wide range of strain rates relatively constant values of the fracture elongation are obtained for a TRIP and DP steel. However, the combination of H induced mechanical degradation and dynamic high strain rate conditions has not been given much attention so far. Therefore, the aim of this work is to investigate the effect of H on the impact properties of a specific DP steel grade. This H influence was evaluated by modifying both the H content in the tensile specimens and the applied strain rate going from static to dynamic conditions.
Table 1 Chemical composition of the used materials in wt%. Material/Element
C
Mn
Si
Other
DP
0.07
1.50
0.25
0.4%−0.8% Cr + Mo
2. Experimental procedure 2.1. Material characterization The sensitivity of DP600 steel to H induced mechanical degradation was investigated in this work. Hot and cold rolling was done till a final thickness of 1.2 mm, which was followed by subsequent annealing via industrial annealing parameters required to attain the desired microstructure. The chemical composition is presented in Table 1. The microstructure was characterized using light optical microscopy (LOM) for which the material was grinded, polished and finally etched using a 2% Nital solution. The samples were immediately cleaned using methanol and acetone. The DP600 steel is a ferritic-martensitic dual phase steel with approximately 23.6% of martensite. The corresponding grain sizes were about 7 μm for the ferritic phase and about 2 μm for the martensitic phase, as shown in Fig. 1. Tensile samples were machined with their tensile axis parallel to the rolling direction. A dogbone-shaped specimen geometry was used, as illustrated in Fig. 2. The small size of the samples is required for the dynamic tests [19]. However, in order to exclude effects related to sample geometry, the same sample geometry is used for the static tests as well. 2.2. Determination of total hydrogen content The total H content was determined by melt extraction at 1600 °C. Samples were H charged electrochemically in 1 g/l thiourea 0.5 M H2SO4 electrolyte for two hours, which guaranteed saturation without H induced damage, as shown in [16]. The applied current density was modified to induce different amounts of H into the tensile samples. For absorption from an aqueous solution, a relation is found between H solubility and applied current density during electrochemical charging, i.e. the solubility is proportional to the square root of the current density [20]. Therefore, the current densities increased from 0.8 over 5 to 25 mA/cm2. The time between H charging and H detection was kept constant at 2 min due to the experimental procedure used to perform both the static and dynamic tensile tests, which also required 2 min between H charging and tensile testing. The system to determine the H content includes a pulse furnace in which a pre-weighted sample is
Fig. 1. Light optical microscopic image of DP600 steel. 119
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Table 2 Total amount of hydrogen determined by melt extraction for different applied current densities. H content
Air
0.8 mA/cm2
5 mA/cm2
25 mA/cm2
wppm
0.52 ± 0.08
3.77 ± 0.12
4.67 ± 0.11
4.97 ± 0.13
were taken during the tests to capture fracture initiation and evaluate the fracture mode upon the applied strain rate. 3. Hydrogen induced mechanical degradation 3.1. Hydrogen uptake capacity Fig. 2. Geometry of the tensile specimens (in mm).
heated to 1600 °C. The sample releases its H as gaseous H2 and is dragged along by a nitrogen flow. This mixture (N2–H2) is directed to a thermal conductivity measuring cell. The thermal conductivity of the mixture depends on the H2 concentration since the conductivity of H2 and N2 differs significantly. Hence, the H content is calculated based on the difference in thermal conductivity.
Melt extraction was done to determine the total H content, for the different H charging methods as described above. Different current densities were applied and the amount of H increased with applied current density (cf. Table 2), which was observed elsewhere as well [20]. The tensile specimens were charged with H under identical conditions. Hence, the actual amount of H present in the sample is known when the tensile test started.
2.3. Determination of the hydrogen induced mechanical degradation
3.2. Tensile tests performed under static conditions
The H induced ductility loss was verified by performing tensile tests on uncharged and H charged samples while using the abovementioned H charging parameters. The H impact was quantified by the HE%, as defined in Eq. (1), with εch and εun being the engineering strains at fracture averaged over the sample gauge section of the H charged and uncharged sample, respectively. Hence, the HE degree can vary between 0 and 1, with 0 indicating no ductility drop and insensitivity to HE. When an index of 1 is obtained, the ductility drop is 100% and the HE is maximal.
Tensile tests were performed on uncharged and H charged samples at current densities of 0.8, 5 and 25 mA/cm2. The corresponding stress–strain curves of the tests done at strain rates of 0.0167 and 1.67 s−1 are presented in Figs. 3 and 4, respectively. The degree of HE was determined and summarized for all conditions in Table 3. Generally, the sensitivity to HE increases with higher applied current density, which can be attributed to the increased H content (cf. Table 2) [24]. This corresponds to previously obtained results and confirms the correlation between the HE susceptibility and the H amount present in the sample [25–28]. A slight increase in flow stress was also observed for the H charged tensile tests, which was attributed solid solution strengthening, as H might pin the dislocations. This was also observed in [9, 29]. Furthermore, the HE resistance increased when a higher strain rate was applied. This was correlated to the H diffusion distance which decreases with higher strain rate, as described in [16, 29]. The distance x (cm) over which H can diffuse, can be calculated by taking the square root of the product of the diffusion coefficient D (cm2/s) and the time t (s) of the tensile test, i.e. x = (D x t)1/2 [15, 29]. Therefore, during a test performed at strain rate of 0.0167 s−1, from the onset of deformation till fracture, H was able to diffuse over approximately 30 μm, whereas at 1.67 s−1, H was just able to diffuse over about 3 μm. Towards the end of the tensile tests, a crack is formed. Close to the
ɛ % HE = 100 ⎛1 − ch ⎞ ɛ un ⎠ ⎝ ⎜
⎟
(1)
Initially, the HE sensitivity was studied by static tensile tests performed on both uncharged and H charged specimens using a standardized Instron tensile machine. H charging was done for two hours and tensile tests were executed ex-situ due to experimental requirements. A constant time lag of 2 min between H charging and tensile testing was maintained. The observations were correlated with those obtained by the melt extraction, which was also done 2 min after H charging. In a first series of static tests, the impact of the H content on the HE susceptibility was evaluated by using samples which were H charged at different applied current densities. Secondly, the effect of H diffusivity was verified by considering two different –though still static- strain rates: 0.0167 and 1.67 s−1. Additionally, the strain rate was further increased till dynamic conditions. These high strain rate tests were done on a split Hopkinson tensile bar (SHTB) setup [21, 22]. For this purpose, the specimen was pinned into slots between two long bars, the input and output bar. The impact of a projectile generates a tensile loading wave at the free end of the input bar which propagates along the input bar towards the sample where it interacts with the specimen resulting in a high strain rate tensile loading of the sample. The loading wave is partly reflected back to the input bar and partly transmitted towards the output bar. Strain gauges on the bars measure the strain corresponding with the loading, reflected and transmitted waves. From those waves, the total force and elongation history of the tensile sample can be determined based on the principles of 1-dimensional elastic-wave propagation in bars [23]. The dynamic tests were done at average strain rates of ± 450 and ± 900 s−1. As such, the effect of H can be elucidated at crash impact velocities and the impact of H diffusivity can be analyzed on the fracture mode. High speed camera (HSC) images of the deforming sample
Fig. 3. Engineering stress–strain diagrams of static tensile tests performed at a strain rate of 0.0167 s−1. 120
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diffraction. Nevertheless, a different failure mode was observed when the strain rate was increased to 1.67 s−1 after H charging at 0.8 mA/cm2, as illustrated in Fig. 7. The corresponding fracture initiation occurred clearly in the center of the tensile specimen, similarly to the tensile tests performed on uncharged material (cf. Fig. 5). Basically, the fracture initiation occurred in the center for all tests performed in the uncharged condition at both strain rates, whereas only the H charged samples tested at the higher strain rate of 1.67 s−1 showed fracture initiation in the center. An overview of the fracture initiation point for all tests is summarized in Table 4. A transition of fracture initiation from the edge to the center was observed as a function of strain rate for H charged samples. Therefore, this transition in fracture mode, when the strain rate was increased, can be correlated with a decrease in available H diffusion time and consequently a decreased H diffusion distance x. Zones which are characterized by a stress concentration ahead of a growing crack tip attract H from the neighboring H saturated matrix. The probability for this to occur is increased at the slower strain rate. The fracture initiation transition was observed between the strain rates of 0.0167 s−1 and 1.67 s−1, which is correlated to diffusion distances of about 30 and 3 μm, respectively.
Fig. 4. Engineering stress–strain diagrams of static tensile tests performed at a strain rate of 1.67 s−1. Table 3 Summary of the degrees of H induced mechanical degradation. HE% −1
0.0167s 1.67s−1
0.8 mA/cm²
5 mA/cm²
25 mA/cm²
45 27
58 37
72 45
3.3. Tensile tests performed under dynamic conditions Tensile tests at dynamic conditions were done on uncharged and H charged samples at current densities of 0.8, 5 and 25 mA/cm2. The corresponding stress–strain diagrams of SHTB tests performed at a strain rate of 450 and 900 s−1 are displayed in Fig. 8 and Fig. 9, respectively. The resultant degrees of HE are given in Table 5. Even in these dynamic conditions, a clear impact of H on the mechanical properties was observed, i.e. the ductility dropped with about 10%. In agreement with the static tests, the sensitivity to H induced mechanical degradation increased slightly with higher current densities, confirming the link between the HE% and the H content [25–28]. Moreover, a slightly lower degree of HE was obtained when the strain rate was increased. However, linking this observation to the diffusion distance x, over which H can diffuse during the tensile test, is not appropriate as the estimated H diffusion distances during the entire test were only about 0.17 and 0.23 μm at strain rates of 900 and 450 s−1, respectively. The fracture initiation during the dynamic tensile tests was again captured by HSC and the image sequence, just before fracture, was considered. The specimens tested after H charging at 25 mA/cm2 at a strain rate of 450 s−1 and 900 s−1 are presented in Figs. 10 and 11, respectively. For both, fracture initiation was found to occur in the center of the specimen. This is in agreement with the findings demonstrated above (cf. Table 4). Generally, fracture initiation occurred in the center for all tests performed in the uncharged condition and for the H charged samples tested at the higher strain rate of 1.67 s−1. Since dynamic conditions were applied, failure initiation from the center was expected to occur since a transition of fracture initiation was observed
crack tip, a zone characterized by a stress concentration is generated. H can diffuse to this zone and hence accelerate failure. When the applied strain rate was increased, H showed a smaller potential to assist the failure process since the distance it can diffuse during the fracture process is limited. The fracture initiation in the static tensile tests was studied in detail. The initiation of failure was captured by HSC and a sequence of the images, just before fracture of the tensile test, were further analyzed. The corresponding images of the test performed at a strain rate of 1.67 s−1 on an uncharged sample are presented in Fig. 5. Considerable diffuse necking can be observed. Moreover, the final fracture initiation was located at the center of the specimen, indicating a ductile failure mode affected by the high stress triaxiality in the center of the tensile sample. The failure sequence of the H charged (at current density of 5 mA/cm2) specimen tested at 0.0167 s−1 is shown in Fig. 6. As opposed to the uncharged sample, localized necking occurred and the final fracture initiated at the edge of the sample. This was attributed to the damaging effect of H on the material at the sample edge where lattice H preferentially diffuses to. Therefore, this H is able to embrittle the material locally, leading to failure initiation and a further accelerated fracture propagation from the edge onwards. These observations were also made for in-situ H charged tensile specimens of DP and TRIP steel [4, 5, 16], on which H induced failure initiation and propagation was evaluated by SEM analysis and electron backscatter
Fig. 5. HSC images of tensile test performed in air at strain rate of 1.67 s−1 during fracture. 121
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Fig. 6. HSC images of tensile test performed after H charging (5 mA/cm2) at strain rate of 0.0167 s−1 during fracture.
Fig. 7. HSC images of tensile test performed after H charging (0.8 mA/cm2) at strain rate of 1.67 s−1 during fracture. Table 4 Overview of fracture modes in the static tensile tests. Fracture initiation point
Air
0.8 mA/cm²
5 mA/cm²
25 mA/cm²
0.0167s−1 1.67s−1
Center Center
Edge Center
Edge Center
Edge Center
Fig. 9. Engineering stress–strain diagrams of dynamic SHTB tests performed at a strain rate of 900 s−1. Table 5 Summary of the HE% for the dynamic tensile test conditions.
Fig. 8. Engineering stress–strain diagrams of dynamic SHTB tests performed at a strain rate of 450 s−1.
HE%
0.8 mA/cm²
5 mA/cm²
25 mA/cm²
450s−1 900s−1
9 8
11 9
16 11
properties [10, 25]. However, as shown by melt extraction also performed 6 min after H charging, the major part of H is still present in the material: the H content decreased to 3.1 wppm. The 0.67 wppm, released as compared to the original tests (cf. Table 2), was confirmed to be H present near dislocations as demonstrated in our earlier work [25, 30–34]. The corresponding stress–strain diagram for the referred specimen (*Charged) was added to the stress–strain curves for the samples
as a function of strain rate for H charged samples. To verify the reliability of the obtained results of the dynamic testing and to understand the effect of H on the mechanical properties in more detail, identical SHTB tensile tests were performed 6 min after electrochemical H charging at 0.8 mA/cm2 instead of the applied 2 min for the previous tests. As such, part of the mobile H had been discharged from the tensile specimen and could no longer influence the mechanical 122
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Fig. 10. HSC images of tensile test performed after H charging (25 mA/cm2) at strain rate of 450 s−1 during fracture.
Fig. 11. HSC images of tensile test performed after H charging (25 mA/cm2) at strain rate of 900 s−1 during fracture.
35]. The fracture surface of the uncharged and H charged sample at 25 mA/cm², both tensile tested at 1.67 s−1, is presented in Fig. 13. Comparison between both showed a delamination in the central region for the H charged sample. Furthermore, the fracture surfaces of two dynamic tests at 450 s−1 and H charged at 0.8 and 25 mA/cm2 are presented in Fig. 14. A central crack perpendicular to the fracture surface can again be observed. Its size increased with increasing current density, i.e. it is linked to the higher H content. The brittle zone around the crack exhibited some shear characteristics. The crack width also increased slightly with increasing strain rate. Fig. 15 shows a detailed characterization of the central delamination crack in the fracture surface of the tensile test performed at 900 s−1 after H charging at 5 mA/ cm2. Second phase particles were observed and EDX analysis confirmed it were MnS inclusions, which are known to be specifically harmful in the presence of hydrogen [16, 35]. When the fracture surfaces of the static and dynamic tensile tests are compared, similar conclusions can be drawn: a dimpled, ductile fracture surface with brittle features at the MnS containing central segregation line. However, no extensive delamination crack was present for the static tests, while a clear crack propagating throughout the sample center was detected for the dynamic tests.
Fig. 12. Engineering stress–strain diagrams of dynamic SHTB tests performed at a strain rate of 450 s−1, the intermittent time between H charging and tensile testing was 2 (charged) and 6 min (*charged).
tested after 2 min of intermittent time at strain rate of 450 s−1, as presented in Fig. 12. A similar mechanical response compared to the uncharged specimen was obtained, indicating that the HE effect was negligible although the H content was still about 3.1 wppm. This confirmed the crucial importance of mobile H trapped at dislocations in the HE phenomenon [25–28].
5. Conclusion The H induced mechanical degradation in DP steel was evaluated by performing tensile tests under static and dynamic conditions. The samples were H charged till saturation and tensile tests were subsequently done ex-situ. The applied current density was modified to induce different H amounts into the tensile specimens. The effect of the H diffusivity during the test on the HE sensitivity was estimated by increasing the strain rate from static (1.67*10−2 and 1.67 s−1) to dynamic (450 and 900 s−1) conditions. The HE susceptibility increased with applied current density due to higher amounts of H, as determined
4. Fractographic analysis The fracture surfaces of the tensile specimens were analyzed by SEM and a ductile dimple pattern was generally observed. However, a thin brittle zone was present at the center of the sample due to the typical segregation line for DP steels, as already demonstrated elsewhere [16, 123
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Fig. 13. SEM images of fracture surface of the uncharged (left) and specimen H charged at 25 mA/cm² (right), tensile tested at 1.67 s−1.
Fig. 14. SEM images of fracture surface of specimen H charged at 0.8 mA/cm² (left) and at 25 mA/cm² (right), tensile tested at 450 s−1.
Fig. 15. MnS inclusions in the H induced crack observed in the central segregation line on the fracture surface of the specimen H charged at 5 mA/cm2, tensile tested at 900 s−1, as revealed by EDX analysis (right).
in the tests performed in air. Fractographic analysis revealed a brittle central line when H charged, which evolved into a major crack which size increased with applied strain rate. EDX analysis revealed the presence of MnS inclusions in this segregation line, initiating the H induced cracks.
by melt extraction. Moreover, the HE degree was increased with slower strain rates. This was linked to higher probability of H to diffuse to regions of stress concentration ahead of a crack tip at slower strain rate and as such accelerating failure. Even at the highest strain rate (900 s−1), the material lost about 10% of its ductility, which was related with H present in the samples and not due to H diffusion during testing. This was concluded since H induced failure initiated at the edges of the samples at slow strain rates. However, at a strain rate of 1.67 s−1, fracture initiated in a ductile way from the center similarly as
Acknowledgments TD holds a postdoctoral fellowship via grant no. BOF01P03516. The 124
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authors wish to thank the Special Research Fund (BOF), UGent (BOF15/ BAS/06). The authors also acknowledge the technical staff from the Department Materials, Textiles and Chemical Engineering and the Department of Electrical Energy, Metals, Mechanical Constructions & Systems, UGent, for their help with experiments and/or sample preparation.
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