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A lamellar structured ultrafine grain ferrite-martensite dual-phase steel and its resistance to hydrogen embrittlement
Accepted Manuscript A lamellar structured ultrafine grain ferrite-martensite dual-phase steel and its resistance to hydrogen embrittlement Junjie Sun, Tao Jiang, Yu Sun, Yingjun Wang, Yongning Liu PII:
S0925-8388(16)34153-6
DOI:
10.1016/j.jallcom.2016.12.224
Reference:
JALCOM 40146
To appear in:
Journal of Alloys and Compounds
Received Date: 17 November 2016 Revised Date:
15 December 2016
Accepted Date: 17 December 2016
Please cite this article as: J. Sun, T. Jiang, Y. Sun, Y. Wang, Y. Liu, A lamellar structured ultrafine grain ferrite-martensite dual-phase steel and its resistance to hydrogen embrittlement, Journal of Alloys and Compounds (2017), doi: 10.1016/j.jallcom.2016.12.224. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
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ACCEPTED MANUSCRIPT A lamellar structured ultrafine grain ferrite-martensite dual-phase steel and its resistance to hydrogen embrittlement Authors: Junjie Sun, Tao Jiang, Yu Sun, Yingjun Wang, Yongning Liu* Affiliation: State Key Laboratory for Mechanical Behavior of Materials, School of Materials Science and Engineering, Xi’an Jiaotong University, Xi’an 710049, China;
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Corresponding author: Yongning Liu ([email protected]), Tel.:+86 029 82664602; Fax:+86 029 82663453.
Abstract: A lamellar structured ultrafine grain dual-phase (UFG DP) steel was prepared by intercritical annealing and subsequent warm rolling of a low-carbon martensite steel. The ultrafine structure is composed of alternate ferrite and martensite
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strips (i.e., lamellar structure) parallel to the rolling direction, and the strips are composted of very fine grains in an average size of 0.96µm. Hydrogen embrittlement
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(HE) of the UFG DP steel was investigated by a slow strain rate tensile (SSRT) of hydrogen charged specimens. Compared with uniform structured steel obtained in normal quenching and tempering (QT) technology, the UFG DP steel exhibits markedly high resistance to HE at a tensile strength level of 1300MPa. A fracture model based on plastic zone and stress distribution was proposed to explain the fracture process. The lamellar structure leaves weaker ferrite/martensite interface
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parallel to the longitudinal direction that leads to delamination, which relaxes stress concentration and makes crack deflection, resulting in the higher HE resistance of the lamellar structured UFG DP steel. Moreover, the UFG structure increases fracture strength that also increases HE resistance.
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Keywords: ultrafine grain dual-phase steel, hydrogen embrittlement, delamination fracture, ferrite, martensite 1. Introduction
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Recently, ultrafine grained (UFG) steels with mean grain size of less than 1µm have
been studied aggressively[1-12]. Both higher strength and a lower ductile-brittle transition temperature (DBTT) were obtained in this kind of materials[13-15]. However, as strain-hardening ability was decreased seriously when grains were refined to UFG level, the uniform elongation in tensile tests was considerably deteriorated[5, 14, 16, 17]. Such loss of ductility severely restricts their potential applications, therefore many efforts have been made to improve the ductility of UFG steels and there are mainly three possible ways to solve the problem: (1) fabrication of a bimodal grain size distribution, in which the fine grains enhance strength and the coarse grains provide plasticity[18, 19]; (2) introduction of nano-scale carbides into 1
ACCEPTED MANUSCRIPT microstructures, in which the nano-scale carbides enhance dislocation tangle during plastic deformation, which increases strain hardenability and ductility [10, 20]; (3) introduction of martensite phase into the microstructure[3, 4, 15, 21-23], in which the martensite phase transformation induces more dislocations into ferrite adjacent to the martensite. The dislocation–dislocation interaction and dislocation pileup at
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ferrite/martensite (F/M) interface produce long-range elastic back stresses which contribute to rapid strain hardening, thus high strength is obtained in UFG dual-phase (DP) steels without much sacrifice of ductility compared with coarse grained materials. Moreover, the UFG DP steels exhibit continuous yielding, which is beneficial to surface quality for cold forming parts. Therefore, the UFG DP steels
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have been drawing great attention from the researchers and industries.
However, the susceptibility of hydrogen embrittlement (HE) increases with
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increasing strength in steels and leads to a significant loss of ductility and strength especially when the strength is larger than 1200MPa. Worse, HE takes place in DP steels even at much lower strength level at about 690MPa or less[24, 25]. The HE susceptibility is thought to be associated with the high contrast in mechanical properties between martensite and ferrite in DP steels. Hard martensite and local strained martensite-ferrite interface are much more susceptible to HE, and cracks are hydrogen
enhanced
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more easily initiated in the high strength martensite or at the F/M interface by decohesion
(HEDE)
mechanism[24-27].
Moreover
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crack-arresting capability of the hydrogen-bearing ferrite is significantly reduced by hydrogen-enhanced localized plasticity (HELP) mechanism[26]. However, it has been
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reported that ferrite strength and martensite toughness were enhanced by grain refinement to UFG level, which would lead to less severe stress/strain partitioning and better interface cohesion[3]. At present, the effect of ultrafine grain refinement on HE
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of DP steels has not been clear.
Grain refinement effects on the hydrogen embrittlement are found to be conflicting
results in literatures[28-32]. Difference may be caused by different microstructures (grain boundary structure, crystal structure, retained austenite, undissolved carbide, etc.) that exist in fine grained steels. However, few studies have been carried out for UFG steels with a grain size of 1
2µm or less, in which some conflicting results
were shown[33-37]. In a Al bearing high-Mn twin induced plasticity (TWIP) steel[33], HE was successfully suppressed when grains were refined to UFG level. Nie et.al.[34] reported that superior HE resistance of an ultrafine elongated grain steel with composition of 0.6C-2Si-1Cr (wt.%) was obtained and the hydrogen-induced cracks 2
ACCEPTED MANUSCRIPT occurred in association with the ultrafine elongated grain. Kimura et.al.[36] reported that an UFG ferrite steel (with 0.6 wt.% O) exhibited markedly high resistance to hydrogen embrittlement at the tensile strength of 1300MPa, but it has not been clarified if the grain refinement was the reason in increasing the resistance to HE because of hydrogen trapping effect related to nanometer-size oxide particles.
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However, Haruna et.al.[35] reported that HE became more serious with grain refinement of an UFG interstitial free steel produced by accumulative roll-bonding. As for the martensitic steel[37], it was reported that the HE resistance was enhanced in a 0.4C-1Cr-0.5Mo-0.3V-0.04Nb (wt.%) steel with tensile strength from 1300MPa to 1600MPa by grain refining from 20 to 4µm. But, intergranular fracture was not
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suppressed when the grain was refined to a few micrometers. There are no directly relevant researches on the grain refinement effect on HE susceptibility for DP steels
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till now.
In this study, a new way to produce lamellar structured UFG ferrite-martensite DP steels was developed for a low carbon steel with 0.17 wt.% C. HE susceptibility of the UFG DP steel was evaluated by a slow strain rate tensile test (SSRT) and compared with a conventional quenched and tempered (QT) steel. The mechanisms of HE in this lamellar structured UFG DP steel was also discussed. 2.1 Materials
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2. Experimental procedure
A low-carbon steel with composition (in wt.%) of 0.17C, 1.96Mn, 1.52Si, 0.96Cr, 0.17Mo, 0.0096N, 0.012P and 0.0032S was produced by vacuum induction melting
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and casted into an ingot with diameter of 100mm. The composition was designed in order to ensure austenite stability in rolling below Ac1 temperature as shown in Fig. 1 and martensite can be formed after rolling and cooling in air. The processing schedule
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to obtain UFG DP steel is outlined in Fig. 1. The ingot was homogenized at 1200 °C for 2h, and then was rolled into a square bar with a sectional area of 40×40 mm2 (with total area reduction, ε, about 80%) in the temperature range of 1100
950 °C.
The hot-rolled steels were cooled in air to room temperature (RT) and lath martensite microstructure was obtained. Then the martensite bars were reheated to a temperature between Ac1 and Ac3 and held for 1h. The Ac1 and Ac3 temperatures are 745.4°C and 820.2°C measured by differential scanning calorimeter (DSC). The intercritical annealing temperature was set at 800°C for the present study. Then the intercritical annealed specimens were cooled down to 650
700°C and warm rolled with an
accumulative area reduction (ε) about 50% (equivalent strain of 0.58). The final 3
ACCEPTED MANUSCRIPT sectional area of the steel plate after warm rolling was about 40×20 mm2. Finally the warm rolled bars were air cooled to RT. Commercial AISI4142 steel (with a composition of 0.42C, 0.61Mn, 0.33Si, 0.98Cr, 0.21Mo) was used as the QT steel for comparison of the HE property at a same strength level. AISI4142 steel is an important high-strength engineering structural
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steel that has wide applications, such as shaft parts, high-strength bolts and so on, but it often suffers from HE. The QT steel bar was austenitized at 850 °C for 30 minutes, oil quenched and tempered at 550 °C for 1h, and then water cooled. 2.2 Microstructure characterization
Samples for optical microscope (OM), scanning electron microscope (SEM) and
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Electron Back-Scattered Diffraction (EBSD) were prepared by standard mechanical grinding and polishing procedures. Samples parallel to rolling direction were sliced
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from the UFG DP steel plate for the EBSD measurements. And the sample surfaces should be extremely clean, smooth and no deformation in order to obtain high-quality Kikuchi patterns. Therefore, they were finely polished on a vibration polishing machine for 3h. For TEM observation, 0.3mm thick slice was cut and mechanically ground to 50µm in thickness, then the foil was prepared by a twin-jet polishing using a mixture of 15% perchloric acid and 85% methanol at an applied potential of 40 V at
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-40 °C. The martensite volume fraction was determined on the basis of 20 SEM micrographs with a magnification of 3000 for the UFG DP steel, and the volume fraction statistic was conducted on Image Proplus software. The grain size was analyzed by EBSD measurements and metallographic analysis software (JX-2000).
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2.3 Electrochemical hydrogen permeation Electrochemical permeation technique originally developed by Devabathan and Stachurski[38] was employed to study the hydrogen permeation behavior of the
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lamellar structured UFG DP and QT steels. The hydrogen permeation apparatus was composed of an electrolytic cell with two compartments (cathodic and anodic sides), a reference electrode (Hg/HgO/NaOH 0.1M NaOH), two auxiliary electrodes (Pt plate), an electrochemical workstation (CHI660D) and a galvanostat. The hydrogen permeation samples with a dimension of about 35×35×2mm3 were cut from the processed steels. Subsequently, all specimens were mechanically ground with SiC grinding paper up to 2000 grit and rinsed with acetone. The specimen with an exposed surface area of 2.27 cm2 on each side was clamped between the two compartments. One side of the specimen acted as hydrogen entry side. It was polarized at a constant charging current density (40mA/cm2) in 0.5M H2SO4 with 1g/L thiourea. Prior to 4
ACCEPTED MANUSCRIPT hydrogen charging, the specimens were depleted of residual hydrogen until the current lowered to 1µA. The hydrogen exit side of the cell was maintained at a constant potential of 300 mV versus reference electrode. Time-lag method was used to calculate the hydrogen diffusivity (D)[38]: D=L2/6tL
(1)
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Where L was the specimen thickness, tL was the lag time which corresponded to the point on the permeation curve at which it=0.63i∞, and i∞ was the steady-state permeation current density.
2.4 Electrochemical hydrogen charging and mechanical tests
Smooth cylindrical specimens (Φ5mm×25mm), which were machined from the
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processed steels with axial direction in the rolling direction (RD), were used for the SSRT. Hydrogen charging was performed for 12h in a 0.5M H2SO4 with 1g/L thiourea
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at cathodic current density (ichar) of 0.43A/m2 and 0.6 A/m2 for each kind of steel. The thiourea was added in the solution as a hydrogen recombination poison. In order to ensure that only the gauge section of the specimen was charged with hydrogen, the surface of other part was covered with insulated glue. After hydrogen charging, the specimens were electroplated with a cadmium coating, and then immediately subjected to SSRT. And the SSRTs were conducted on Instron1195 tensile test machine at a crosshead speed of 0.03mm/min, corresponding to a strain rate of
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2×10-5/s. Meanwhile the specimens without hydrogen charging were also tested as a reference. The index of relative susceptibilities to HE of the steels was determined by measuring the relative elongation loss, which can be expressed as:
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HEI (%) =
ε0 − ε H × 100% ε0
(2)
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Where ε0 is the total elongation of the hydrogen-free specimen and εH is the total elongation of hydrogen-charged specimen. After tensile tests, the fracture surfaces of tensile specimens were observed by SEM.
3. Results
3.1 Microstructures
After intercritical annealing and subsequent warm rolling of the lath martensite steel, a lamellar structured UFG DP structure is obtained as shown in Fig. 2. The horizontal and vertical arrows represent normal and rolling direction (ND and RD) respectively. The microstructure is in lamellar structure and it is hard to distinguish martensite and ferrite in OM as shown in Fig. 2(a). The details of the microstructure can be seen in SEM photo (Fig. 2(b)), martensite is in light grey and ferrite matrix is 5
ACCEPTED MANUSCRIPT in dark grey. Martensite is in lath morphology and all the laths are parallel to RD, and the volume fraction of martensite is about 40.6%. A small amount of precipitated nano-scale carbides can be seen in the steel. TEM microstructure (Fig. 2(c)) reveals that the ferrite laths are composed of UFGs with average grain size about 1µm and dislocation density is high in the grains. And the lamellar martensite is composed of
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ultrafine martensite lath/plate with dislocation substructure as shown in the enlarged image in Fig. 2(d). The ultrafine martensite lath/plate is formed after warm rolling below Ac1 temperature.
EBSD was carried out for the lamellar structured UFG DP steel as shown in Fig. 3. Significantly refined grains are revealed in the image-quality (IQ) map (Fig. 3(a)). In
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the IQ map, martensite is characterized by a lower IQ value because of the larger lattice distortions, thus appearing dark gray, whereas the ferrite appears brighter. The
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grain size distribution determined by the EBSD analysis, using a critical misorientation angle of 10 degree to define grains, is illustrated in Fig. 3(b), and the average grain size is measured to be 0.96µm that is similar with the value determined by TEM observation.
Fig. 4 shows the OM and SEM microstructures for the QT sample, which shows a typical tempered martensitic structure with fine carbides segregated at lath boundaries.
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The average prior-austenite grain size is measured to be 18.2µm using the grain boundary revealing and analyzing method reported in literature[39, 40]. 3.2 Hydrogen permeation tests
Fig. 5 shows the hydrogen permeation transient currents for QT and lamellar
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structured UFG DP steels. The time-lag method is used to calculate the hydrogen diffusion coefficient, and the measured hydrogen permeability data are summarized in Table 1. The UFG DP steel demonstrates much lower hydrogen diffusivity compared
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with that of QT steel.
3.3 HE susceptibility
Fig. 6 shows the stress-strain curves of the hydrogen-charged QT and lamellar
structured UFG DP steels with varied hydrogen charging current density. Total elongation decreases with increasing hydrogen charging current density for both steels, but there is no significant change in the work-hardening behavior. By contrast, the lamellar structured UFG DP steel displays much lower HE susceptibility than that of QT steel. The yield and tensile strengths of the UFG DP and QT steels do not change significantly when they are pre-hydrogen charged at a current density of 0.43A/m2, but the total elongation to fracture deteriorates. With the hydrogen current density 6
ACCEPTED MANUSCRIPT increases to 0.6A/m2, the hydrogen content in the specimen increases, both the tensile strength and the elongation deteriorate severely for the QT steel, while the UFG DP steel still maintains a decent level of strength and elongation. Table 2 summarizes the mechanical properties of the two kinds of steels with varied hydrogen charging current density.
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3.4 Fractography Fig. 7 displays the fracture appearances of hydrogen charged and hydrogen free UFG DP steels after the SSRT. For the hydrogen charged specimen, both specimens charged at current density 0.43A/m2 and 0.6A/m2 display a similar fracture mode and only the fracture appearance of the specimen charged at 0.6 A/m2 is shown as a
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representative. It can be seen clearly that the fractures contain two distinct areas for both the hydrogen free (Fig. 7(a)) and hydrogen charged (Fig. 7(e)) samples:
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peripheral shear lip region (indicated by the black arrows) and center ductile and delamination fracture region (indicated by the blue arrows). The fractographs of the shear lip region consists of dimples as shown in the enlarged figures (Fig. 7(c and g)) of the black box marked areas, but the dimples become shallower for the hydrogen charged specimen. For the center region, fractures show a layered topography, in which cracks are branched in the longitudinal direction (
200µm that is much wider than
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specimen, but the width of the layer is about 50
RD) of the tensile
ferrite or martensite lath width (0.5
2µm) as shown in Fig. 2. And the center region
contains some smooth areas and humps or rough regions as indicated by the orange arrows and red bocks respectively as shown in Fig. 7(b and f). The smooth region is
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composed of dimple as shown in Fig. 7(b and f), and the hump region is also composed of mainly dimples as shown in Fig. 7(d and g). Fig. 7(d) shows the dimples and a small amount of micro-delamination cracks (as indicated by the purple arrows).
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In contrast, the micro-delamination is more striking and cracks (as indicated by the yellow dotted arrows) between two layers become larger and more clear for the hydrogen charged specimen as shown in Fig. 7(f). Brittle fracture characteristics become apparent and more micro-delamination cracks (as indicated by the purple arrows) appear on the delamination fracture surface as shown in Fig. 7(h). The increased micro-delamination cracks and brittle fracture demonstrate that the martensite becomes more brittle and the interface cohesion between martensite and ferrite becomes weaker with introducing hydrogen. Therefore, it can be considered that the martensite becomes more brittle and the delamination becomes more pronounced with hydrogen charging, which indicates a hydrogen-assisted 7
ACCEPTED MANUSCRIPT delamination in the UFG DP steels. In the QT steel, fracture surfaces of all the samples with varied hydrogen charging current density display two regions, the peripheral region with shear lip and the center crack initiation and propagation region, as shown in Fig. 8(a and c). The shear lip is a ductile fracture region which consists of dimples. For the crack initiation and
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propagation region, it displays different characteristics with increasing hydrogen charging current density. The fracture mode is fully ductile for the uncharged specimen (Fig. 8(b)), while the fractographs become brittle intergranular for the hydrogen charged specimens as shown in Fig. 8(d) and the intergranular cracks propagate along the prior-austenite grain boundaries. Such hydrogen induced 4. Discussion
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4.1 Microstructure effect on hydrogen diffusion
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intergranular fracture has been commonly observed in conventional QT steels[41-43].
Hydrogen diffusion plays an important role on hydrogen induced crack. The hydrogen traps are related to microstructural features such as dislocations, interfaces, vacancies, impurity atoms, micro voids or any other lattice defect[44],
which
remarkably delay hydrogen diffusion in materials. In the UFG DP steel, martensite transformation introduces more dislocation in the ferrite grains; moreover, the UFG
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DP microstructure increases both F/M interfaces and grain boundaries, which lead to hydrogen traps increase. As for the QT steel, carbides precipitation increases the traps, but the high-temperature tempering process makes dislocation density decreased seriously, which results in the QT steel has a higher hydrogen diffusion coefficient.
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The time (t) required to attain the equilibrium hydrogen concentration in the studied steels can be roughly estimated from Eq. (3) [45], t=0.334R2/D
(3)
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Where R is the radius of the gauge section of the tensile specimen, and R=2.5×10-3m. Taking D equals 5.2×10-11 m2/s and 1.3×10-10 m2/s for the UFG DP and QT steels respectively as shown in Table 1, calculated t is about 4.01×104s and 1.6×104s, respectively, i.e., about 11h and 4.5h. In this experiment, hydrogen charging time is 12h. Therefore, hydrogen concentration in lattice sites is in balance. 4.2 Delamination fracture mechanism As shown in the tensile specimen fracture surface of the UFG DP steel in Fig. 7, distinct delamination can be observed, and the delamination fracture becomes more significant with hydrogen charging. In addition, the layer thickness does not match the martensite ferrite or martensite lath width and the center region contains some plastic 8
ACCEPTED MANUSCRIPT shear areas (i.e., smooth areas). This phenomenon can be interpreted by the fracture mechanics theory on plastic zone at crack tip as illustrated in Fig. 9. The loading direction (parallel to RD) and main direction of crack propagation (perpendicular to RD) are defined as y and x axes respectively, stress and plastic zone width (r0) ahead
σy =σx =
r0
(1 − 2ν ) = 2π
KI 2πr 2
K IC σs
(4) 2
(5)
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of the crack tip can be expressed by Eq. (4) and (5) respectively[46]:
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Where r is the distance from the crack tip on x axis, KI is stress field intensity factor, ν is Poission’s ratio, KIC is fracture toughness and σs is the material yield stress. The measured transverse yield strength (about 485.2MPa) is lower than that of the
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longitudinal direction (about 797.2MPa), delamination cracks (i.e., crack deflection) easily initiate at the F/M interface ahead of the crack tip as illustrated in Fig. 9(a). Similar delamination fracture has been reported by Kimura et al in ultrafine elongated grain steel[34, 47-49], and the delamination was thought to be associated with the texture caused strength differences between the longitudinal and transverse directions. The formation of layered fracture is illustrated in Fig. 9(b). After several
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micro-delamiantions, plastic shear deformation takes place along the maximum shear stress direction. When the plastic shear deformation goes through the plastic zone, shear deformation will stop and the crack turns to the main direction of the crack (x axis). Then the second delamination layer begins to form. The above process is
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repeated until the material breaks and forms the layered topography. Common value of the fracture toughness of DP steel is about 50
100MPa·m1/2[50], and the yield
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strength of the experimental UFG DP steel is approximately 800MPa. Substituting these data into Eq. (4), the calculated r0 is on the order of 0.1mm, thus, the shear deformation will cross many ferrite and martensite laths, and the result is in conformity with the value of layer width shown in Fig. 7. With hydrogen charging, cohesion strength between the crystal planes or F/M interfaces is reduced that will promote delamination fracture. This is consistent with the experimental result that delamination increases with increasing hydrogen charging current density. 4.3 Improvement of HE resistance and hydrogen induced cracking mechanism of the UFG DP steel The significant improvement of the HE resistance was demonstrated in the lamellar-structured UFG DP steel when compared with the QT samples at the same 9
ACCEPTED MANUSCRIPT strength level of 1300MPa. The enhanced intrinsic fracture resistance to HE in the lamellar structured UFG DP steel is thought to be associated with the oriented F/M interface (i.e., lamellar structure), ultrafine grains and dual phase structure. It has been demonstrated that key factors affecting the HE of high strength steels are dissolved hydrogen content, microstructure and stress intensity factor[51].
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Hydrogen was reported to exist on prior austenite grain boundaries preferentially in martensite[52, 53] that caused significant reduction in grain boundary cohesive strength. Therefore, hydrogen-induced intergranular fracture preferentially occurs at the prior austenite grain boundaries in the conventional QT steels[34, 41-43] and the mechanical properties of the hydrogen charged QT samples is seriously deteriorated
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by HEDE mechanism. As for the DP steels, hydrogen-induced martensite cracking has similarities with that in martensite steels[26]. And the strength difference between
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martensite and ferrite introduces strain/stress partition[54, 55] which promotes hydrogen concentration at the F/M interface. Hereby, the propensity for cracking is enhanced in that local as a result of significant reduction in lattice cohesive strength caused by locally high hydrogen concentrations, i.e., cracking at the F/M interface by HEDE. The weakened F/M interface has a similar effect on crack formation like that of oxide films. The oxide films give rise to matrix interface weakening and are seen to
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constitute cracks within aluminum alloys[56,57]. The lamellar structure leaves weaker F/M interfaces parallel to the longitudinal direction (//RD) of the bar, and the hydrogen concentration at the F/M interface promotes cracks propagate along the longitudinal direction more easily rather than along the transverse direction. This
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analysis is consistent with the experimental result that the delamination fracture becomes more significant with hydrogen charging. In fact, the HE resistance of the lamellar structured UFG DP sample is mainly dependent on crack propagation
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resistance in transverse direction rather than along the longitudinal direction. Therefore, the oriented lamellar structure is thought to be the dominating reason why the UFG DP steel has much higher resistance than the QT steel. As has been demonstrated by Bai et al [33] that grain refinement suppressed HE in
high-Mn TWIP steel, because grain refinement increased grain boundaries that reduced hydrogen concentration at grain boundaries. In this experiment, the UFG DP structure introduces much more grain boundaries and F/M interfaces that can reduce local accumulation of hydrogen at grain boundaries and interfaces, which may also contribute to the suppression of HE. In addition, the oriented F/M interfaces result in a large number of micro-delaminations with increasing tensile strain, which can relax 10
ACCEPTED MANUSCRIPT stress concentration and deflect crack propagation direction that can reduce HE in the UFG DP steel. Moreover, the critical value of the local stress for transgranular or intergranular fracture (σfc) and the effective grain size (d) are roughly comparable on the same scale according to the relation: σfc=Kd-1/2 [34, 49, 58]. When compared to the QT steel with an average prior-austenite grain size of 18.2µm, the UFG DP steels with an average
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grain size of 0.96µm should have much higher σfc. Therefore, the grain refinement also increases HE resistance of the lamellar structured UFG DP steel. 5. Conclusion
In this study, lamellar structured UFG DP steel was produced by intercritical
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annealing and subsequent warm rolling of a martensite starting structure in a low carbon steel. HE property of the lamellar structured UFG DP steel was investigated
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using SSRT, and was compared with that of the conventional QT samples at a similar strength level of 1300MPa. The following conclusions can be drawn. 1. Lamellar structured UFG DP steel is formed by intercritical annealing and subsequent warm rolling of a martensite starting structure. The ultrafine structure is composed of alternate ferrite and martensite strips (i.e., lamellar structure) parallel to the rolling direction, and the strips are composted of very fine grains in
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an average size of 0.98µm.
2. For SSRT tests, the lamellar structured UFG DP steel exhibits markedly high resistance to HE compared with conventional QT steel at a tensile strength level of 1300MPa. In contrast to the intergranular fracture in the QT samples, the
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hydrogen-induced fracture exhibits a ductile mode and micro-delamination in the UFG DP steel.
3. The lamellar structure in UFG DP steel leaves weaker F/M interfaces parallel to
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the longitudinal direction, and the strength difference between martensite and ferrite promotes hydrogen concentration at the F/M interfaces resulting in reduction in the lattice cohesive strength by HEDE mechanism. Both aspects promote cracks initiation and propagation along the F/M interface, i.e., hydrogen-induced cracks are more easily propagating along the longitudinal direction rather than along the transverse direction, which promotes delamination fracture. And a proposed fracture model based on plastic zone and stress distribution explains the delamination well.
4. The enhanced resistance to HE in the lamellar structured UFG DP steel is thought to be associated with the lamellar structure and UFG DP structure. The lamellar 11
ACCEPTED MANUSCRIPT structure promotes delamination along the longitudinal direction which relaxes stress concentration and deflects crack propagation direction, leading to the higher HE resistance of the UFG DP steel. Moreover, the UFG structure increases fracture strength that also increases HE resistance. Acknowledgements 51271137 and 50871082) is gratefully acknowledged. Reference:
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Financial support from the National Natural Science Foundation of China (No.
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ACCEPTED MANUSCRIPT [15] M. Calcagnotto, D. Ponge, D. Raabe, Effect of grain refinement to 1μm on strength and toughness of dual-phase steels, Mater. Sci. Eng., A, 527 (2010) 7832-7840. [16] K.T. Park, Y.S. Kim, J.G. Lee, H.S. Dong, Thermal stability and mechanical properties of ultrafine grained low carbon steel, Mater. Sci. Eng., A, 293 (2000) 165-172. [17] N. Tsuji, Y. Minamino, Y. Ito, Y. Saito, Strength and ductility of ultrafine grained aluminum and iron produced by ARB and annealing, Scripta Mater., 47 (2002) 893-899. distributions in low carbon steels, Scripta Mater., 57 (2007) 1065-1068.
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[30] I.J. Park, S.M. Lee, H.H. Jeon, Y.K. Lee, The advantage of grain refinement in the hydrogen embrittlement of Fe-18Mn-0.6C twinning-induced plasticity steel, Corros. Sci., 93 (2015) 63–69. [31] R. Padmanabhan, W.E. Wood, Hydrogen induced cracking in a low alloy steel, Metall. Mater. Trans. A, 14 (1983) 2347-2356.
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ACCEPTED MANUSCRIPT [36] Y. Kimura, Y. Sakai, T. Hara, A. Belyakov, K. Tsuzaki, Hydrogen induced delayed fracture of ultrafine grained 0.6% O steel with dispersed oxide particles, Scripta Mater., 49 (2003) 1111-1116. [37] W. Hui, H. Dong, Y. Weng, J. Shi, Y. Nie, Z. Chu, Y. Chen, Delayed Fracture Behavior of Ultrafine Grained High Strength Steel, Acta Metall. Sin., 40 (2004) 561-568. [38] M.A.V. Devanathan, Z. Stachurski, The Adsorption and Diffusion of Electrolytic Hydrogen in Palladium, Proc. R. Soc. A, 270 (1962) 90-102. transformation in hypereutectoid steel, J. Mater. Res., DOI (2012).
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[39] F.L. Lian, H.J.Liu, J.J. Sun, X.J. Sun, S.W. Guo, Y.N. Liu, L.X. Du, Ultrafine grain effect on pearlitic [40] J.J. Sun, F.L. Lian, H.J. Liu, T. Jiang, S.W. Guo, L.X. Du, Y.N. Liu, Microstructure of warm rolling and pearlitic transformation of ultrafine-grained GCr15 steel,. Mater. Charact., 95 (2014) 291-298.
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fracture of high strength steel by constant load test and numerical calculation, Corros. Sci., 48 (2006) 2189-2202.
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2007.
[47] Y. Kimura, T. Inoue, F. Yin, O. Sitdikov, K. Tsuzaki, Toughening of a 1500MPa class steel through formation of an ultrafine fibrous grain structure, Scripta Mater., 57 (2007) 465-468. [48] Y. Kimura, T. Inoue, K. Tsuzaki, Tempforming in medium-carbon low-alloy steel, J. Alloys Compd., 1 (2013) S538-S542.
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[49] Y. Kimura, T. Inoue, E. Akiyama, K. Tsuzaki, Hydrogen Embrittlement in Ultrafine Elongated Grain Structure Steel Processed by Warm Tempforming, International Hydrogen Conference (IHC 2012), ASME Press, 2014.
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[50] ASM International Handbook Committee, Properties and Selection: Irons, Steels, and
High-performance Alloys, tenth ed., The Materials Information Company, USA, 1995. [51] M. Koyama, E. Akiyama, K. Tsuzaki, D. Raabe, Hydrogen-assisted failure in a twinning-induced plasticity steel studied under in situ hydrogen charging by electron channeling contrast imaging, Acta Mater., 61 (2013) 4607-4618. [52] K. Takai, Jun, rsquo, i. Seki, Y. Homma, Observation of Trapping Sites of Hydrogen and Deuterium in High-Strength Steels by Using Secondary Ion Mass Spectrometry, Mater. Trans., JIM, 36 (1995) 1134-1139. [53] A. Nagao, C.D. Smith, M. Dadfarnia, P. Sofronis, I.M. Robertson, The role of hydrogen in hydrogen embrittlement fracture of lath martensitic steel, Acta Mater., 60 (2012) 5182-5189. [54] S. Morooka, O. Umezawa, S. Harjo, K. Hasegawa, Y. Toji, Quantitative Analysis of Tensile Deformation Behavior by In-Situ Neutron Diffraction for Ferrite-Martensite Type Dual-Phase Steels, Tetsu to Hagane, 98 (2012) 311-319. 14
ACCEPTED MANUSCRIPT [55] J. Kadkhodapour, S. Schmauder, D. Raabe, S. Ziaei-Rad, U. Weber, M. Calcagnotto, Experimental and numerical study on geometrically necessary dislocations and non-homogeneous mechanical properties of the ferrite phase in dual phase steels, Acta Mater., 59 (2011) 4387-4394. [56] J. Campbell, M. Tiryakioğlu, Modelling Microstructure and Properties: The Contributions of Grain Size, DAS and Bifilms, Mater. Sci. Forum, 519-521 (2006) 1453-1460. [57] N.R. Green, J. Campbell, Statistical distributions of fracture strengths of cast Al-7Si-Mg alloy, Mater. Sci. Eng., A, 173 (1993) 261-166.
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[58] J.W. Morris, Z. Guo, C.R. Krenn, K. Y.-H., The Limits of Strength and Toughness in Steel, ISIJ Int., 41
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(2001) 599-611.
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ACCEPTED MANUSCRIPT Table 1 Measured hydrogen permeability data for QT and lamellar structured UFG DP steels. Thickness (mm)
tL (s)
D (m2/s)
QT steel
1.82
4251
1.3×10-10
UFG DP steel
1.93
11941
5.2×10-11
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Materials
ACCEPTED MANUSCRIPT Table 2 Summary of the tensile properties of the hydrogen-charged QT and lamellar structured UFG DP steels with varied hydrogen charging current density. ichar
Yield strength
2
QT steel
HEI
(A/m )
(MPa)
(MPa)
(%)
(%)
0
1228.8
1345.8
7.3
0
0.43
1223.6
1335.0
2.65
63.7
0.6
—
1081.9
0
797.2
1375.1
0.43
798.1
1369.8
0.6
793.5
1354.0
0
100
11.68
0
8.65
25.9
5.3
54.6
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UFG DP steel
Tensile strength Elongation
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Materials
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Fig. 1 Processing routes to produce UFG DP steel. AC: air cooling; RT: room
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temperature; ε: area reduction.
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Fig. 2 Microstructure of the lamellar structured UFG DP steel: (a) optical microstructure; (b) SEM microstructure; (c) TEM microstructure; (d) substructure of
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the martensite.
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Fig. 3 Typical imaging maps obtained from the FE-SEM/EBSD for the UFG DP steel:
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(a) image quality (IQ) map, (b) grain size distribution.
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Fig. 4 Microstructure of the QT steel: (a) optical microstructure; (b) SEM
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microstructure.
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Fig. 5 Hydrogen permeation transient current curves for QT and lamellar structured
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UFG DP steels.
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Fig. 6 Tensile properties of the hydrogen-charged QT and lamellar structured UFG DP steels with varied hydrogen charging current density: (a) lamellar structured UFG DP steel; (b) QT steel; (c) HEI of the QT and lamellar structured UFG DP steels.
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Fig. 7 Fractographs of the lamellar structured UFG DP steels after the SSRT with varied hydrogen charging current density: (a), (b), (c), (d) hydrogen free specimen; (e), (f), (g), (h) 0.6A/m2. The black, blue, and yellow arrows show the shear lip, delamination region and delamination, respectively.
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Fig. 8 Fractographs of the QT steels after the SSRT with and without hydrogen charging: (a) over view of the fracture surface and (b) crack initiation and propagation region of the uncharged sample, (c) over view of the fracture surface and (d) crack
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current density.
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initiation and propagation region of the hydrogen charged specimen with 0.6A/m2
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Fig. 9 Schematic of the delamination fracture mechanism: (a) plastic zone and stress
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distribution ahead of crack tip; (b) schematic of layered fracture formation.
ACCEPTED MANUSCRIPT Lamellar structured and ultrafine grained dual-phase (UFG DP) steel is produced. The UFG DP steel exhibits high resistance to hydrogen embrittlement (HE).
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Delamination along lamellar structure contributes to high resistance to HE. Grain refinement enhances hydrogen embrittlement property of DP steel.
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A model based on fracture mechanics explains the delamination well.
S0925-8388(16)34153-6
DOI:
10.1016/j.jallcom.2016.12.224
Reference:
JALCOM 40146
To appear in:
Journal of Alloys and Compounds
Received Date: 17 November 2016 Revised Date:
15 December 2016
Accepted Date: 17 December 2016
Please cite this article as: J. Sun, T. Jiang, Y. Sun, Y. Wang, Y. Liu, A lamellar structured ultrafine grain ferrite-martensite dual-phase steel and its resistance to hydrogen embrittlement, Journal of Alloys and Compounds (2017), doi: 10.1016/j.jallcom.2016.12.224. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
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ACCEPTED MANUSCRIPT A lamellar structured ultrafine grain ferrite-martensite dual-phase steel and its resistance to hydrogen embrittlement Authors: Junjie Sun, Tao Jiang, Yu Sun, Yingjun Wang, Yongning Liu* Affiliation: State Key Laboratory for Mechanical Behavior of Materials, School of Materials Science and Engineering, Xi’an Jiaotong University, Xi’an 710049, China;
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Corresponding author: Yongning Liu ([email protected]), Tel.:+86 029 82664602; Fax:+86 029 82663453.
Abstract: A lamellar structured ultrafine grain dual-phase (UFG DP) steel was prepared by intercritical annealing and subsequent warm rolling of a low-carbon martensite steel. The ultrafine structure is composed of alternate ferrite and martensite
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strips (i.e., lamellar structure) parallel to the rolling direction, and the strips are composted of very fine grains in an average size of 0.96µm. Hydrogen embrittlement
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(HE) of the UFG DP steel was investigated by a slow strain rate tensile (SSRT) of hydrogen charged specimens. Compared with uniform structured steel obtained in normal quenching and tempering (QT) technology, the UFG DP steel exhibits markedly high resistance to HE at a tensile strength level of 1300MPa. A fracture model based on plastic zone and stress distribution was proposed to explain the fracture process. The lamellar structure leaves weaker ferrite/martensite interface
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parallel to the longitudinal direction that leads to delamination, which relaxes stress concentration and makes crack deflection, resulting in the higher HE resistance of the lamellar structured UFG DP steel. Moreover, the UFG structure increases fracture strength that also increases HE resistance.
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Keywords: ultrafine grain dual-phase steel, hydrogen embrittlement, delamination fracture, ferrite, martensite 1. Introduction
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Recently, ultrafine grained (UFG) steels with mean grain size of less than 1µm have
been studied aggressively[1-12]. Both higher strength and a lower ductile-brittle transition temperature (DBTT) were obtained in this kind of materials[13-15]. However, as strain-hardening ability was decreased seriously when grains were refined to UFG level, the uniform elongation in tensile tests was considerably deteriorated[5, 14, 16, 17]. Such loss of ductility severely restricts their potential applications, therefore many efforts have been made to improve the ductility of UFG steels and there are mainly three possible ways to solve the problem: (1) fabrication of a bimodal grain size distribution, in which the fine grains enhance strength and the coarse grains provide plasticity[18, 19]; (2) introduction of nano-scale carbides into 1
ACCEPTED MANUSCRIPT microstructures, in which the nano-scale carbides enhance dislocation tangle during plastic deformation, which increases strain hardenability and ductility [10, 20]; (3) introduction of martensite phase into the microstructure[3, 4, 15, 21-23], in which the martensite phase transformation induces more dislocations into ferrite adjacent to the martensite. The dislocation–dislocation interaction and dislocation pileup at
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ferrite/martensite (F/M) interface produce long-range elastic back stresses which contribute to rapid strain hardening, thus high strength is obtained in UFG dual-phase (DP) steels without much sacrifice of ductility compared with coarse grained materials. Moreover, the UFG DP steels exhibit continuous yielding, which is beneficial to surface quality for cold forming parts. Therefore, the UFG DP steels
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have been drawing great attention from the researchers and industries.
However, the susceptibility of hydrogen embrittlement (HE) increases with
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increasing strength in steels and leads to a significant loss of ductility and strength especially when the strength is larger than 1200MPa. Worse, HE takes place in DP steels even at much lower strength level at about 690MPa or less[24, 25]. The HE susceptibility is thought to be associated with the high contrast in mechanical properties between martensite and ferrite in DP steels. Hard martensite and local strained martensite-ferrite interface are much more susceptible to HE, and cracks are hydrogen
enhanced
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more easily initiated in the high strength martensite or at the F/M interface by decohesion
(HEDE)
mechanism[24-27].
Moreover
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crack-arresting capability of the hydrogen-bearing ferrite is significantly reduced by hydrogen-enhanced localized plasticity (HELP) mechanism[26]. However, it has been
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reported that ferrite strength and martensite toughness were enhanced by grain refinement to UFG level, which would lead to less severe stress/strain partitioning and better interface cohesion[3]. At present, the effect of ultrafine grain refinement on HE
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of DP steels has not been clear.
Grain refinement effects on the hydrogen embrittlement are found to be conflicting
results in literatures[28-32]. Difference may be caused by different microstructures (grain boundary structure, crystal structure, retained austenite, undissolved carbide, etc.) that exist in fine grained steels. However, few studies have been carried out for UFG steels with a grain size of 1
2µm or less, in which some conflicting results
were shown[33-37]. In a Al bearing high-Mn twin induced plasticity (TWIP) steel[33], HE was successfully suppressed when grains were refined to UFG level. Nie et.al.[34] reported that superior HE resistance of an ultrafine elongated grain steel with composition of 0.6C-2Si-1Cr (wt.%) was obtained and the hydrogen-induced cracks 2
ACCEPTED MANUSCRIPT occurred in association with the ultrafine elongated grain. Kimura et.al.[36] reported that an UFG ferrite steel (with 0.6 wt.% O) exhibited markedly high resistance to hydrogen embrittlement at the tensile strength of 1300MPa, but it has not been clarified if the grain refinement was the reason in increasing the resistance to HE because of hydrogen trapping effect related to nanometer-size oxide particles.
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However, Haruna et.al.[35] reported that HE became more serious with grain refinement of an UFG interstitial free steel produced by accumulative roll-bonding. As for the martensitic steel[37], it was reported that the HE resistance was enhanced in a 0.4C-1Cr-0.5Mo-0.3V-0.04Nb (wt.%) steel with tensile strength from 1300MPa to 1600MPa by grain refining from 20 to 4µm. But, intergranular fracture was not
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suppressed when the grain was refined to a few micrometers. There are no directly relevant researches on the grain refinement effect on HE susceptibility for DP steels
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till now.
In this study, a new way to produce lamellar structured UFG ferrite-martensite DP steels was developed for a low carbon steel with 0.17 wt.% C. HE susceptibility of the UFG DP steel was evaluated by a slow strain rate tensile test (SSRT) and compared with a conventional quenched and tempered (QT) steel. The mechanisms of HE in this lamellar structured UFG DP steel was also discussed. 2.1 Materials
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2. Experimental procedure
A low-carbon steel with composition (in wt.%) of 0.17C, 1.96Mn, 1.52Si, 0.96Cr, 0.17Mo, 0.0096N, 0.012P and 0.0032S was produced by vacuum induction melting
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and casted into an ingot with diameter of 100mm. The composition was designed in order to ensure austenite stability in rolling below Ac1 temperature as shown in Fig. 1 and martensite can be formed after rolling and cooling in air. The processing schedule
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to obtain UFG DP steel is outlined in Fig. 1. The ingot was homogenized at 1200 °C for 2h, and then was rolled into a square bar with a sectional area of 40×40 mm2 (with total area reduction, ε, about 80%) in the temperature range of 1100
950 °C.
The hot-rolled steels were cooled in air to room temperature (RT) and lath martensite microstructure was obtained. Then the martensite bars were reheated to a temperature between Ac1 and Ac3 and held for 1h. The Ac1 and Ac3 temperatures are 745.4°C and 820.2°C measured by differential scanning calorimeter (DSC). The intercritical annealing temperature was set at 800°C for the present study. Then the intercritical annealed specimens were cooled down to 650
700°C and warm rolled with an
accumulative area reduction (ε) about 50% (equivalent strain of 0.58). The final 3
ACCEPTED MANUSCRIPT sectional area of the steel plate after warm rolling was about 40×20 mm2. Finally the warm rolled bars were air cooled to RT. Commercial AISI4142 steel (with a composition of 0.42C, 0.61Mn, 0.33Si, 0.98Cr, 0.21Mo) was used as the QT steel for comparison of the HE property at a same strength level. AISI4142 steel is an important high-strength engineering structural
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steel that has wide applications, such as shaft parts, high-strength bolts and so on, but it often suffers from HE. The QT steel bar was austenitized at 850 °C for 30 minutes, oil quenched and tempered at 550 °C for 1h, and then water cooled. 2.2 Microstructure characterization
Samples for optical microscope (OM), scanning electron microscope (SEM) and
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Electron Back-Scattered Diffraction (EBSD) were prepared by standard mechanical grinding and polishing procedures. Samples parallel to rolling direction were sliced
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from the UFG DP steel plate for the EBSD measurements. And the sample surfaces should be extremely clean, smooth and no deformation in order to obtain high-quality Kikuchi patterns. Therefore, they were finely polished on a vibration polishing machine for 3h. For TEM observation, 0.3mm thick slice was cut and mechanically ground to 50µm in thickness, then the foil was prepared by a twin-jet polishing using a mixture of 15% perchloric acid and 85% methanol at an applied potential of 40 V at
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-40 °C. The martensite volume fraction was determined on the basis of 20 SEM micrographs with a magnification of 3000 for the UFG DP steel, and the volume fraction statistic was conducted on Image Proplus software. The grain size was analyzed by EBSD measurements and metallographic analysis software (JX-2000).
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2.3 Electrochemical hydrogen permeation Electrochemical permeation technique originally developed by Devabathan and Stachurski[38] was employed to study the hydrogen permeation behavior of the
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lamellar structured UFG DP and QT steels. The hydrogen permeation apparatus was composed of an electrolytic cell with two compartments (cathodic and anodic sides), a reference electrode (Hg/HgO/NaOH 0.1M NaOH), two auxiliary electrodes (Pt plate), an electrochemical workstation (CHI660D) and a galvanostat. The hydrogen permeation samples with a dimension of about 35×35×2mm3 were cut from the processed steels. Subsequently, all specimens were mechanically ground with SiC grinding paper up to 2000 grit and rinsed with acetone. The specimen with an exposed surface area of 2.27 cm2 on each side was clamped between the two compartments. One side of the specimen acted as hydrogen entry side. It was polarized at a constant charging current density (40mA/cm2) in 0.5M H2SO4 with 1g/L thiourea. Prior to 4
ACCEPTED MANUSCRIPT hydrogen charging, the specimens were depleted of residual hydrogen until the current lowered to 1µA. The hydrogen exit side of the cell was maintained at a constant potential of 300 mV versus reference electrode. Time-lag method was used to calculate the hydrogen diffusivity (D)[38]: D=L2/6tL
(1)
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Where L was the specimen thickness, tL was the lag time which corresponded to the point on the permeation curve at which it=0.63i∞, and i∞ was the steady-state permeation current density.
2.4 Electrochemical hydrogen charging and mechanical tests
Smooth cylindrical specimens (Φ5mm×25mm), which were machined from the
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processed steels with axial direction in the rolling direction (RD), were used for the SSRT. Hydrogen charging was performed for 12h in a 0.5M H2SO4 with 1g/L thiourea
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at cathodic current density (ichar) of 0.43A/m2 and 0.6 A/m2 for each kind of steel. The thiourea was added in the solution as a hydrogen recombination poison. In order to ensure that only the gauge section of the specimen was charged with hydrogen, the surface of other part was covered with insulated glue. After hydrogen charging, the specimens were electroplated with a cadmium coating, and then immediately subjected to SSRT. And the SSRTs were conducted on Instron1195 tensile test machine at a crosshead speed of 0.03mm/min, corresponding to a strain rate of
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2×10-5/s. Meanwhile the specimens without hydrogen charging were also tested as a reference. The index of relative susceptibilities to HE of the steels was determined by measuring the relative elongation loss, which can be expressed as:
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HEI (%) =
ε0 − ε H × 100% ε0
(2)
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Where ε0 is the total elongation of the hydrogen-free specimen and εH is the total elongation of hydrogen-charged specimen. After tensile tests, the fracture surfaces of tensile specimens were observed by SEM.
3. Results
3.1 Microstructures
After intercritical annealing and subsequent warm rolling of the lath martensite steel, a lamellar structured UFG DP structure is obtained as shown in Fig. 2. The horizontal and vertical arrows represent normal and rolling direction (ND and RD) respectively. The microstructure is in lamellar structure and it is hard to distinguish martensite and ferrite in OM as shown in Fig. 2(a). The details of the microstructure can be seen in SEM photo (Fig. 2(b)), martensite is in light grey and ferrite matrix is 5
ACCEPTED MANUSCRIPT in dark grey. Martensite is in lath morphology and all the laths are parallel to RD, and the volume fraction of martensite is about 40.6%. A small amount of precipitated nano-scale carbides can be seen in the steel. TEM microstructure (Fig. 2(c)) reveals that the ferrite laths are composed of UFGs with average grain size about 1µm and dislocation density is high in the grains. And the lamellar martensite is composed of
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ultrafine martensite lath/plate with dislocation substructure as shown in the enlarged image in Fig. 2(d). The ultrafine martensite lath/plate is formed after warm rolling below Ac1 temperature.
EBSD was carried out for the lamellar structured UFG DP steel as shown in Fig. 3. Significantly refined grains are revealed in the image-quality (IQ) map (Fig. 3(a)). In
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the IQ map, martensite is characterized by a lower IQ value because of the larger lattice distortions, thus appearing dark gray, whereas the ferrite appears brighter. The
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grain size distribution determined by the EBSD analysis, using a critical misorientation angle of 10 degree to define grains, is illustrated in Fig. 3(b), and the average grain size is measured to be 0.96µm that is similar with the value determined by TEM observation.
Fig. 4 shows the OM and SEM microstructures for the QT sample, which shows a typical tempered martensitic structure with fine carbides segregated at lath boundaries.
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The average prior-austenite grain size is measured to be 18.2µm using the grain boundary revealing and analyzing method reported in literature[39, 40]. 3.2 Hydrogen permeation tests
Fig. 5 shows the hydrogen permeation transient currents for QT and lamellar
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structured UFG DP steels. The time-lag method is used to calculate the hydrogen diffusion coefficient, and the measured hydrogen permeability data are summarized in Table 1. The UFG DP steel demonstrates much lower hydrogen diffusivity compared
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with that of QT steel.
3.3 HE susceptibility
Fig. 6 shows the stress-strain curves of the hydrogen-charged QT and lamellar
structured UFG DP steels with varied hydrogen charging current density. Total elongation decreases with increasing hydrogen charging current density for both steels, but there is no significant change in the work-hardening behavior. By contrast, the lamellar structured UFG DP steel displays much lower HE susceptibility than that of QT steel. The yield and tensile strengths of the UFG DP and QT steels do not change significantly when they are pre-hydrogen charged at a current density of 0.43A/m2, but the total elongation to fracture deteriorates. With the hydrogen current density 6
ACCEPTED MANUSCRIPT increases to 0.6A/m2, the hydrogen content in the specimen increases, both the tensile strength and the elongation deteriorate severely for the QT steel, while the UFG DP steel still maintains a decent level of strength and elongation. Table 2 summarizes the mechanical properties of the two kinds of steels with varied hydrogen charging current density.
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3.4 Fractography Fig. 7 displays the fracture appearances of hydrogen charged and hydrogen free UFG DP steels after the SSRT. For the hydrogen charged specimen, both specimens charged at current density 0.43A/m2 and 0.6A/m2 display a similar fracture mode and only the fracture appearance of the specimen charged at 0.6 A/m2 is shown as a
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representative. It can be seen clearly that the fractures contain two distinct areas for both the hydrogen free (Fig. 7(a)) and hydrogen charged (Fig. 7(e)) samples:
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peripheral shear lip region (indicated by the black arrows) and center ductile and delamination fracture region (indicated by the blue arrows). The fractographs of the shear lip region consists of dimples as shown in the enlarged figures (Fig. 7(c and g)) of the black box marked areas, but the dimples become shallower for the hydrogen charged specimen. For the center region, fractures show a layered topography, in which cracks are branched in the longitudinal direction (
200µm that is much wider than
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specimen, but the width of the layer is about 50
RD) of the tensile
ferrite or martensite lath width (0.5
2µm) as shown in Fig. 2. And the center region
contains some smooth areas and humps or rough regions as indicated by the orange arrows and red bocks respectively as shown in Fig. 7(b and f). The smooth region is
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composed of dimple as shown in Fig. 7(b and f), and the hump region is also composed of mainly dimples as shown in Fig. 7(d and g). Fig. 7(d) shows the dimples and a small amount of micro-delamination cracks (as indicated by the purple arrows).
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In contrast, the micro-delamination is more striking and cracks (as indicated by the yellow dotted arrows) between two layers become larger and more clear for the hydrogen charged specimen as shown in Fig. 7(f). Brittle fracture characteristics become apparent and more micro-delamination cracks (as indicated by the purple arrows) appear on the delamination fracture surface as shown in Fig. 7(h). The increased micro-delamination cracks and brittle fracture demonstrate that the martensite becomes more brittle and the interface cohesion between martensite and ferrite becomes weaker with introducing hydrogen. Therefore, it can be considered that the martensite becomes more brittle and the delamination becomes more pronounced with hydrogen charging, which indicates a hydrogen-assisted 7
ACCEPTED MANUSCRIPT delamination in the UFG DP steels. In the QT steel, fracture surfaces of all the samples with varied hydrogen charging current density display two regions, the peripheral region with shear lip and the center crack initiation and propagation region, as shown in Fig. 8(a and c). The shear lip is a ductile fracture region which consists of dimples. For the crack initiation and
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propagation region, it displays different characteristics with increasing hydrogen charging current density. The fracture mode is fully ductile for the uncharged specimen (Fig. 8(b)), while the fractographs become brittle intergranular for the hydrogen charged specimens as shown in Fig. 8(d) and the intergranular cracks propagate along the prior-austenite grain boundaries. Such hydrogen induced 4. Discussion
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4.1 Microstructure effect on hydrogen diffusion
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intergranular fracture has been commonly observed in conventional QT steels[41-43].
Hydrogen diffusion plays an important role on hydrogen induced crack. The hydrogen traps are related to microstructural features such as dislocations, interfaces, vacancies, impurity atoms, micro voids or any other lattice defect[44],
which
remarkably delay hydrogen diffusion in materials. In the UFG DP steel, martensite transformation introduces more dislocation in the ferrite grains; moreover, the UFG
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DP microstructure increases both F/M interfaces and grain boundaries, which lead to hydrogen traps increase. As for the QT steel, carbides precipitation increases the traps, but the high-temperature tempering process makes dislocation density decreased seriously, which results in the QT steel has a higher hydrogen diffusion coefficient.
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The time (t) required to attain the equilibrium hydrogen concentration in the studied steels can be roughly estimated from Eq. (3) [45], t=0.334R2/D
(3)
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Where R is the radius of the gauge section of the tensile specimen, and R=2.5×10-3m. Taking D equals 5.2×10-11 m2/s and 1.3×10-10 m2/s for the UFG DP and QT steels respectively as shown in Table 1, calculated t is about 4.01×104s and 1.6×104s, respectively, i.e., about 11h and 4.5h. In this experiment, hydrogen charging time is 12h. Therefore, hydrogen concentration in lattice sites is in balance. 4.2 Delamination fracture mechanism As shown in the tensile specimen fracture surface of the UFG DP steel in Fig. 7, distinct delamination can be observed, and the delamination fracture becomes more significant with hydrogen charging. In addition, the layer thickness does not match the martensite ferrite or martensite lath width and the center region contains some plastic 8
ACCEPTED MANUSCRIPT shear areas (i.e., smooth areas). This phenomenon can be interpreted by the fracture mechanics theory on plastic zone at crack tip as illustrated in Fig. 9. The loading direction (parallel to RD) and main direction of crack propagation (perpendicular to RD) are defined as y and x axes respectively, stress and plastic zone width (r0) ahead
σy =σx =
r0
(1 − 2ν ) = 2π
KI 2πr 2
K IC σs
(4) 2
(5)
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of the crack tip can be expressed by Eq. (4) and (5) respectively[46]:
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Where r is the distance from the crack tip on x axis, KI is stress field intensity factor, ν is Poission’s ratio, KIC is fracture toughness and σs is the material yield stress. The measured transverse yield strength (about 485.2MPa) is lower than that of the
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longitudinal direction (about 797.2MPa), delamination cracks (i.e., crack deflection) easily initiate at the F/M interface ahead of the crack tip as illustrated in Fig. 9(a). Similar delamination fracture has been reported by Kimura et al in ultrafine elongated grain steel[34, 47-49], and the delamination was thought to be associated with the texture caused strength differences between the longitudinal and transverse directions. The formation of layered fracture is illustrated in Fig. 9(b). After several
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micro-delamiantions, plastic shear deformation takes place along the maximum shear stress direction. When the plastic shear deformation goes through the plastic zone, shear deformation will stop and the crack turns to the main direction of the crack (x axis). Then the second delamination layer begins to form. The above process is
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repeated until the material breaks and forms the layered topography. Common value of the fracture toughness of DP steel is about 50
100MPa·m1/2[50], and the yield
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strength of the experimental UFG DP steel is approximately 800MPa. Substituting these data into Eq. (4), the calculated r0 is on the order of 0.1mm, thus, the shear deformation will cross many ferrite and martensite laths, and the result is in conformity with the value of layer width shown in Fig. 7. With hydrogen charging, cohesion strength between the crystal planes or F/M interfaces is reduced that will promote delamination fracture. This is consistent with the experimental result that delamination increases with increasing hydrogen charging current density. 4.3 Improvement of HE resistance and hydrogen induced cracking mechanism of the UFG DP steel The significant improvement of the HE resistance was demonstrated in the lamellar-structured UFG DP steel when compared with the QT samples at the same 9
ACCEPTED MANUSCRIPT strength level of 1300MPa. The enhanced intrinsic fracture resistance to HE in the lamellar structured UFG DP steel is thought to be associated with the oriented F/M interface (i.e., lamellar structure), ultrafine grains and dual phase structure. It has been demonstrated that key factors affecting the HE of high strength steels are dissolved hydrogen content, microstructure and stress intensity factor[51].
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Hydrogen was reported to exist on prior austenite grain boundaries preferentially in martensite[52, 53] that caused significant reduction in grain boundary cohesive strength. Therefore, hydrogen-induced intergranular fracture preferentially occurs at the prior austenite grain boundaries in the conventional QT steels[34, 41-43] and the mechanical properties of the hydrogen charged QT samples is seriously deteriorated
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by HEDE mechanism. As for the DP steels, hydrogen-induced martensite cracking has similarities with that in martensite steels[26]. And the strength difference between
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martensite and ferrite introduces strain/stress partition[54, 55] which promotes hydrogen concentration at the F/M interface. Hereby, the propensity for cracking is enhanced in that local as a result of significant reduction in lattice cohesive strength caused by locally high hydrogen concentrations, i.e., cracking at the F/M interface by HEDE. The weakened F/M interface has a similar effect on crack formation like that of oxide films. The oxide films give rise to matrix interface weakening and are seen to
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constitute cracks within aluminum alloys[56,57]. The lamellar structure leaves weaker F/M interfaces parallel to the longitudinal direction (//RD) of the bar, and the hydrogen concentration at the F/M interface promotes cracks propagate along the longitudinal direction more easily rather than along the transverse direction. This
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analysis is consistent with the experimental result that the delamination fracture becomes more significant with hydrogen charging. In fact, the HE resistance of the lamellar structured UFG DP sample is mainly dependent on crack propagation
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resistance in transverse direction rather than along the longitudinal direction. Therefore, the oriented lamellar structure is thought to be the dominating reason why the UFG DP steel has much higher resistance than the QT steel. As has been demonstrated by Bai et al [33] that grain refinement suppressed HE in
high-Mn TWIP steel, because grain refinement increased grain boundaries that reduced hydrogen concentration at grain boundaries. In this experiment, the UFG DP structure introduces much more grain boundaries and F/M interfaces that can reduce local accumulation of hydrogen at grain boundaries and interfaces, which may also contribute to the suppression of HE. In addition, the oriented F/M interfaces result in a large number of micro-delaminations with increasing tensile strain, which can relax 10
ACCEPTED MANUSCRIPT stress concentration and deflect crack propagation direction that can reduce HE in the UFG DP steel. Moreover, the critical value of the local stress for transgranular or intergranular fracture (σfc) and the effective grain size (d) are roughly comparable on the same scale according to the relation: σfc=Kd-1/2 [34, 49, 58]. When compared to the QT steel with an average prior-austenite grain size of 18.2µm, the UFG DP steels with an average
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grain size of 0.96µm should have much higher σfc. Therefore, the grain refinement also increases HE resistance of the lamellar structured UFG DP steel. 5. Conclusion
In this study, lamellar structured UFG DP steel was produced by intercritical
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annealing and subsequent warm rolling of a martensite starting structure in a low carbon steel. HE property of the lamellar structured UFG DP steel was investigated
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using SSRT, and was compared with that of the conventional QT samples at a similar strength level of 1300MPa. The following conclusions can be drawn. 1. Lamellar structured UFG DP steel is formed by intercritical annealing and subsequent warm rolling of a martensite starting structure. The ultrafine structure is composed of alternate ferrite and martensite strips (i.e., lamellar structure) parallel to the rolling direction, and the strips are composted of very fine grains in
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an average size of 0.98µm.
2. For SSRT tests, the lamellar structured UFG DP steel exhibits markedly high resistance to HE compared with conventional QT steel at a tensile strength level of 1300MPa. In contrast to the intergranular fracture in the QT samples, the
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hydrogen-induced fracture exhibits a ductile mode and micro-delamination in the UFG DP steel.
3. The lamellar structure in UFG DP steel leaves weaker F/M interfaces parallel to
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the longitudinal direction, and the strength difference between martensite and ferrite promotes hydrogen concentration at the F/M interfaces resulting in reduction in the lattice cohesive strength by HEDE mechanism. Both aspects promote cracks initiation and propagation along the F/M interface, i.e., hydrogen-induced cracks are more easily propagating along the longitudinal direction rather than along the transverse direction, which promotes delamination fracture. And a proposed fracture model based on plastic zone and stress distribution explains the delamination well.
4. The enhanced resistance to HE in the lamellar structured UFG DP steel is thought to be associated with the lamellar structure and UFG DP structure. The lamellar 11
ACCEPTED MANUSCRIPT structure promotes delamination along the longitudinal direction which relaxes stress concentration and deflects crack propagation direction, leading to the higher HE resistance of the UFG DP steel. Moreover, the UFG structure increases fracture strength that also increases HE resistance. Acknowledgements 51271137 and 50871082) is gratefully acknowledged. Reference:
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Financial support from the National Natural Science Foundation of China (No.
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ACCEPTED MANUSCRIPT Table 1 Measured hydrogen permeability data for QT and lamellar structured UFG DP steels. Thickness (mm)
tL (s)
D (m2/s)
QT steel
1.82
4251
1.3×10-10
UFG DP steel
1.93
11941
5.2×10-11
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ACCEPTED MANUSCRIPT Table 2 Summary of the tensile properties of the hydrogen-charged QT and lamellar structured UFG DP steels with varied hydrogen charging current density. ichar
Yield strength
2
QT steel
HEI
(A/m )
(MPa)
(MPa)
(%)
(%)
0
1228.8
1345.8
7.3
0
0.43
1223.6
1335.0
2.65
63.7
0.6
—
1081.9
0
797.2
1375.1
0.43
798.1
1369.8
0.6
793.5
1354.0
0
100
11.68
0
8.65
25.9
5.3
54.6
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UFG DP steel
Tensile strength Elongation
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Materials
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Fig. 1 Processing routes to produce UFG DP steel. AC: air cooling; RT: room
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temperature; ε: area reduction.
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Fig. 2 Microstructure of the lamellar structured UFG DP steel: (a) optical microstructure; (b) SEM microstructure; (c) TEM microstructure; (d) substructure of
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the martensite.
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Fig. 3 Typical imaging maps obtained from the FE-SEM/EBSD for the UFG DP steel:
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(a) image quality (IQ) map, (b) grain size distribution.
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Fig. 4 Microstructure of the QT steel: (a) optical microstructure; (b) SEM
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microstructure.
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Fig. 5 Hydrogen permeation transient current curves for QT and lamellar structured
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UFG DP steels.
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Fig. 6 Tensile properties of the hydrogen-charged QT and lamellar structured UFG DP steels with varied hydrogen charging current density: (a) lamellar structured UFG DP steel; (b) QT steel; (c) HEI of the QT and lamellar structured UFG DP steels.
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Fig. 7 Fractographs of the lamellar structured UFG DP steels after the SSRT with varied hydrogen charging current density: (a), (b), (c), (d) hydrogen free specimen; (e), (f), (g), (h) 0.6A/m2. The black, blue, and yellow arrows show the shear lip, delamination region and delamination, respectively.
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Fig. 8 Fractographs of the QT steels after the SSRT with and without hydrogen charging: (a) over view of the fracture surface and (b) crack initiation and propagation region of the uncharged sample, (c) over view of the fracture surface and (d) crack
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current density.
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initiation and propagation region of the hydrogen charged specimen with 0.6A/m2
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Fig. 9 Schematic of the delamination fracture mechanism: (a) plastic zone and stress
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distribution ahead of crack tip; (b) schematic of layered fracture formation.
ACCEPTED MANUSCRIPT Lamellar structured and ultrafine grained dual-phase (UFG DP) steel is produced. The UFG DP steel exhibits high resistance to hydrogen embrittlement (HE).
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Delamination along lamellar structure contributes to high resistance to HE. Grain refinement enhances hydrogen embrittlement property of DP steel.
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A model based on fracture mechanics explains the delamination well.