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Hydrogen embrittlement behaviors of additive manufactured maraging steel investigated by in situ high-energy X-ray diffraction
Materials Science & Engineering A 766 (2019) 138341
Contents lists available at ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Hydrogen embrittlement behaviors of additive manufactured maraging steel investigated by in situ high-energy X-ray diffraction
T
Shilei Lia, Mingming Liua, Yang Renb, Yandong Wanga,∗ a
Beijing Advanced Innovation Center for Materials Genome Engineering, State Key Laboratory for Advanced Metals and Materials, University of Science and Technology Beijing, Beijing, 100083, China b X-ray Science Division, Advanced Photon Source, Argonne National Laboratory, Argonne, Illinois, 60439, USA
A R T I C LE I N FO
A B S T R A C T
Keywords: Additive manufacturing Hydrogen embrittlement High-energy X-ray diffraction Martensitic transformation
The microstructures and mechanical properties of additively manufactured (AM) maraging steel were investigated before and after hydrogen charging. The hydrogen-charged specimen suffers severe hydrogen embrittlement and its strength and plasticity decline obviously. The mechanism of ductile-to-brittle transition of AM maraging steel after hydrogen charging was studied by in situ high energy X-ray diffraction and electron backscatter diffraction techniques. Martensitic transformation occurs in the hydrogen-charged specimen under low stress. The premature phase transformation along with hydrogen-induced microcracks resulting in the brittle fracture of the hydrogen-charged specimen.
1. Introduction Additive manufacturing (AM) is fast gaining acceptance as a way to manufacture structural components with high structural integrity at a low cost [1–3]. Several AM alloys with mechanical properties comparable to alloys conventionally manufactured have been produced over the past years by reducing the residual porosity, insufficient layer bonding and hot cracking [4–7]. Selective laser melting (SLM), a typical metal AM technology, builds up three dimensional (3D) parts layer-bylayer out of powder particles. Maraging steels have an excellent combination of ultra-high strength with good toughness and ductility and therefore are widely applied in the aircraft, aerospace and tooling industries [8–11]. Maraging steels, consist of the martensitic matrix and various precipitates, are typically manufactured via a series of complex processes: casting, homogenizing, forging, quenching, machining, hot rolling, and a final aging treatment [9,10]. The as-produced maraging steels by SLM usually achieve significantly higher strength than conventional processing methods without heat treatment, because of their quenched martensitic matrix with fine cellular microstructures induced by the high cooling rate during SLM [3,12]. The microstructure and mechanical properties of maraging steels manufactured by SLM, as well as the effects of heat treatment have been investigated in the previous literature [13–15]. Conventional maraging steels are susceptible to hydrogen
∗
embrittlement and their mechanical properties suffer a great loss [16,17]. In this work, the hydrogen embrittlement behaviors of AM maraging 300 steel were investigated. The synchrotron-based high-energy X-ray diffraction (HE-XRD) technique was used to reveal the deformation and damage mechanisms of both the hydrogen-free and hydrogen-charged specimens. 2. Materials and experimental procedures The AM grade 300 maraging steel used in this study was fabricated by a SLM facility EOS M290 (EOS GmbH, Krailling, Germany) equipped with a 400 W Yb-fiber laser. Spherical gas atomized powders with particle sizes ranging from 20 to 45 μm were selected and the following parameters were used during the building of AM 316 L: laser power (350 W), beam size (100 μm), layer thickness (40 μm), scanning speed (1100 mm/s) and line spacing (150 μm). A pillar with a geometry of 10 mm (length) × 10 mm (width) × 50 mm (height) was built and no post-built heat treatment was applied. The microstructure of the as-fabricated specimen perpendicular to the building direction (BD) was observed using an optical microscope (OM) and a ZEISS Supra 55 scanning electron microscope (SEM) equipped with an electron backscatter diffraction (EBSD) detector. Hydrogen was introduced into the specimens by electrochemical charging with current density 50 mA/cm2 at room temperature in a 3 wt% NaCl +0.3 wt% NH4SCN solution for 24 h. The concentration of
Corresponding author. E-mail address: [email protected] (Y. Wang).
https://doi.org/10.1016/j.msea.2019.138341 Received 28 May 2019; Received in revised form 13 August 2019; Accepted 27 August 2019 Available online 28 August 2019 0921-5093/ © 2019 Elsevier B.V. All rights reserved.
Materials Science & Engineering A 766 (2019) 138341
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hydrogen was 5 wt ppm, measured using the pulse heating inert gas fusion infrared absorption method. Both the hydrogen-free and hydrogen-charged AM steels were cut perpendicular to the BD into flat dog-bone shaped tensile specimens, with a gauge length of 10 mm, width 3 mm and thickness 0.5 mm. The in situ tensile experiments were performed at beamline 11-ID-C of the Advanced Photon Source (APS) at Argonne National Laboratory. A monochromatic X-ray beam (500 × 500 μm) with an energy of 106 keV (with a wavelength of 0.11730 Å) was used to investigate the evolution of lattice strain and diffraction intensity during tensile loading. The fracture surface morphologies of both hydrogen-free and hydrogencharged specimens were observed in SEM. The deformed microstructures near the fractures of the tensile specimens were investigated. During the in situ tensile testing, a two-dimensional (2D) detector was used to collect the diffraction patterns at different strain levels. The specimen was strained incrementally at room temperature until fracture. The accurate distance between the specimen and the detector was calibrated by Ceria (CeO2) powders. The loading direction (LD) and transverse direction (TD) correspond to azimuth angles ϕ = 90° and ϕ = 180°, respectively. The 2D diffraction patterns were caked and integrated along specified azimuth angles over a range of ± 5° to obtain one-dimensional HE-XRD diffraction spectrums. A Pseudo-Voigt function was used to determine the position and the integrated intensity of the diffraction peaks. The lattice strain evolution in the retained austenite (RA) and martensite in the studied steel can be obtained by calculating the lattice spacing change during deformation using the function εhkl= (dhkl-d0)/d0, where εhkl is the lattice strain of a specific {hkl} plane, and dhkl and d0 are the interplanar spacings with and without stress.
Fig. 2. Engineering stress-strain curves of the hydrogen-free and hydrogencharged AM grade 300 maraging steel.
specimens before deformation. Different diffraction peaks with various interplanar spacings are identified as those belonging to the bodycentered cubic (BCC) structured martensite (M), as well as the retained austenite (A) with a face-centered cubic (FCC) structure. Comparisons among the diffraction peaks for the hydrogen-free and hydrogencharged AM maraging steel specimens suggest that no phase transformation occurs during hydrogen charging. 3.2. Macroscopic tensile behavior
3. Results and discussion
Fig. 2 shows the engineering stress-strain curves of the hydrogenfree and hydrogen-charged specimens tested at room temperature. Compared with the hydrogen-free specimen, the hydrogen-charged specimen has obvious lower tensile strength and elongation after fracture. The hydrogen-charged specimen fails without obvious plastic deformation and its ultimate tensile strength is lower than the yield strength of the hydrogen-free specimen. The macroscopic and microscopic morphologies of the fracture surfaces of the hydrogen-free and hydrogen-charged specimens, as
3.1. Microstructures Fig. 1a and b shows the OM microstructure and inverse pole figure (IPF) of the as-fabricated AM grade 300 maraging steel. Similar to a conventional grade 300 maraging steel [18], the AM steel also has a martensitic matrix with uneven coarse plate martensites. Fig. 1c shows the HE-XRD profiles of the hydrogen-free and hydrogen-charged
Fig. 1. OM (a) and IPF (b) micrographs of the studied AM grade 300 maraging steel. (c) HE-XRD profiles of the hydrogen-free and hydrogen-charged specimens before deformation. 2
Materials Science & Engineering A 766 (2019) 138341
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Fig. 3. The macroscopic and microscopic morphologies of the fracture surfaces of the hydrogen-free and hydrogen-charged specimens showing the obvious hydrogen embrittlement.
hydrogen-free and hydrogen-charged specimens with the applied stress along the LD are displayed in Fig. 5. In the elastic deformation stage, lattice strains of all reflections increase approximately linearly with increasing the applied stress. The slopes of the lattice strain vs. stress curves of reflections for martensite and RA in the hydrogen-charged specimen are almost the same as those in the hydrogen-free specimen. The evolution of diffraction intensity with the applied stress reflects the deformation and phase transformation information of phases in the material. For the FCC-structured RA and BCC-structured martensite, their intensities of A111/A200 and M110 reflections increase during uniaxial tensile deformation because of the development of fiber texture. While the martensitic transformation from RA to the newly formed martensite (NFM) will cause the decline of intensity for RA. As the small volume of RA in the studied material, this transformation may not significantly increase the intensity of martensite. From the evolution of normalized integral intensity, as shown in Fig. 6, it is easy to identify the critical stress for martensitic transformation in the hydrogen-free and hydrogen-charged specimens. The RA in the hydrogen-free specimen begins to transform to NFM when the applied stress reaches 760 MPa. This critical stress of martensitic transformation is lower than the macro yielding of the hydrogen-free specimen because the RA is easy to deform and has a lower micro yield stress. The presence of hydrogen can decrease the stacking fault energy of austenite and
shown in Fig. 3, confirm the transition from ductile to brittle in the AM maraging steel induced by hydrogen charging. Non-uniform dimples, as the typical feature of ductile fracture, were observed on the fracture surface of the hydrogen-free specimen. SEM fractography of the hydrogen-charged specimen reveals a brittle fracture mode which presents a mixture of dimples and cleavage planes. In addition, a large number of secondary cracks were observed on the fracture surface of the hydrogen-charged specimen. 3.3. Micromechanical behavior during in situ HE-XRD In situ HE-XRD was used to investigate the deformation behaviors of both hydrogen-free and hydrogen-charged specimens during tensile loading. Fig. 4 shows the 1D HE-XRD profiles of the specimens before and after hydrogen charging along the LD under different loadings. The peak positions along the LD continually shift to the higher value of dspacing during the tensile process. The diffraction intensity of reflections belong to RA along the LD in the hydrogen-free specimen decreases with increasing applied stress during tensile loading, and the peak finally merges with the background. However, the diffraction peaks of RA along the LD in the hydrogen-charged specimen remains almost unchanged during deformation until failure, as shown in Fig. 4b. The evolution of lattice strain of martensite and RA in both
Fig. 4. The 1D HE-XRD profiles of the hydrogen-free (a) and hydrogen-charged (b) specimens along the LD under different loadings (bottom to top represents 0 MPa to fracture). 3
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Fig. 5. Lattice strain evolution of the martensite and RA in the hydrogen-free (a) and hydrogen-charged (b) specimens with the applied stress along the LD.
promote the transformation of RA [19,20]. Some RA grains in the hydrogen-charged specimen may begin to plastic deform locally and transform to NFM under a very low applied stress [21]. In the present study, the hydrogen-charged specimen has much lower critical stress of 300 MPa for martensitic transformation compared with the hydrogenfree specimen.
3.4. Cracking in the hydrogen-charged specimen The proportion of RA in the hydrogen-free specimen significantly decreases after deformation, while the proportion for the deformed hydrogen-charged specimen is similar to the undeformed condition, as shown in Fig. 4. This confirms that the premature martensitic transformation in the hydrogen-free specimen only occurs in some local areas. Fig. 7 shows the evolution of normalized full width at half maximum (FWHM) for martensite and RA along the LD of the hydrogen-free and hydrogen-charged specimens. The normalized FWHM for A111 reflection of the hydrogen-charged sample has an obvious decrease at 464 MPa, which is higher than the critical stress for the martensitic transformation. The formation of NFM introduces localized stress fields and hydrogen atoms rapidly diffuse under localized stress gradients. High concentration of hydrogen may decrease the cohesive force between the atoms in the crystal and cause localized plastic deformation or cleavage [17], leading to the cracking along the {111}⊥LD planes in the RA. Longitudinal section morphologies near the fractures, cut along the LD from the fractured tensile specimens, are investigated in OM.
Fig. 7. Normalized FWHM for martensite and RA along the LD of the hydrogenfree and hydrogen-charged specimens.
Secondary cracks were observed in the hydrogen-charged specimen near the tensile fracture, as shown in Fig. 8b. In contrast, no crack was observed on the longitudinal section near the fracture of the hydrogenfree specimen (Fig. 8a). The IPF in Fig. 9 shows the orientation distribution vertical to the LD of the cracking region near the fracture of the hydrogen-charged specimen. The RA in the hydrogen-charged fractured specimen was not
Fig. 6. Normalized integral intensity evolution of the martensite and RA in the hydrogen-free (a) and hydrogen-charged (b) specimens with the applied stress along the LD. 4
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Fig. 8. OM micrographs of the longitudinal section near the fractures of the hydrogen-free (c) and hydrogen-charged (d) specimens.
treatments by combining annealing, solution treatment and aging is expected to improve the hydrogen embrittlement resistance of AM maraging steel. 4. Conclusions The deformation and damage mechanism of AM maraging steel before and after hydrogen charging was investigated using the in situ HE-XRD technique. Results reveal that the hydrogen-charged AM specimen suffers a loss in both plasticity and strength. The observation of fracture morphologies also confirms the ductile-to-brittle transition after hydrogen charging. The presence of hydrogen causes the locally premature martensitic transformation and cracking along the {111}⊥LD planes in the RA. Acknowledgements Fig. 9. The IPF of the cracking region near the fracture of the hydrogen-charged specimen.
This work was supported by the National Key Research and Development Program of China (No. 2017YFA0403804), the National Natural Science Foundation of China (NSFC) (Nos. 51601013 and 51527801) and the State Key Laboratory for Advanced Metals and Materials (Grant No. 2018Z-25). Use of the Advanced Photon Source is supported by the U. S. Department of Energy, Office of Science, Office of Basic Energy Sciences, under Contract No. DE-AC02-06CH11357.
observed in the EBSD result, while its existence is confirmed by the HEXRD results in Fig. 4b. It is considered that the RA in the EBSD specimen surface transfers into martensite due to the mechanical polishing. As RA and the martensite matrix in maraging steel usually follows the NW orientation relationship {111}A //{110}M [18]. The IPF of the cracked region in Fig. 9 indicates that the microcracks near the fracture of the hydrogen-charged specimen are along the {110}⊥LD planes of martensite. Thus, these microcracks initiate and propagate along the {111}⊥LD planes of RA, which is consistent with the HE-XRD results in Fig. 7. These microcracks release the localized stress fields in RA and thus inhibit the further phase transformation from RA to NFM. Under the applied stress, these microcracks coalesce and propagate quickly, resulting in the brittle fracture of the hydrogen-charged specimen. The content of RA will change and intermetallic particles may precipitate in the AM maraging steel during the post heat treatment, and the microstructure evolution significantly affects the mechanical properties [15]. The investigation on the hydrogen embrittlement behaviors of the conventional 300 grade maraging steel after different heat treatments indicated that the ranking of hydrogen embrittlement susceptibility in the decreasing order was peak-aged, over-aged and solution annealed condition [17]. It was reported that the elongation of conventional grade 300 maraging steel decreased from 8% to 0.3% under hydrogen charging [18]. In the present study, the AM maraging steel has a decline in elongation after hydrogen charging from 14.3% to 2.5%. The local residual stress, the content of RA and precipitation of intermetallic particles are believed to affect the hydrogen embrittlement behaviors of the AM maraging steel. Optimizing the heat
Appendix A. Supplementary data Supplementary data to this article can be found online at https:// doi.org/10.1016/j.msea.2019.138341. References [1] D. Herzog, V. Seyda, E. Wycisk, C. Emmelmann, Additive manufacturing of metals, Acta Mater. 117 (2016) 371–392. [2] J. Suryawanshi, K.G. Prashanth, S. Scudino, J. Eckert, O. Prakash, U. Ramamurty, Simultaneous enhancements of strength and toughness in an Al-12Si alloy synthesized using selective laser melting, Acta Mater. 115 (2016) 285–294. [3] C. Tan, K. Zhou, W. Ma, P. Zhang, M. Liu, T. Kuang, Microstructural evolution, nanoprecipitation behavior and mechanical properties of selective laser melted high-performance grade 300 maraging steel, Mater. Des. 134 (2017) 23–34. [4] B.E. Carroll, T.A. Palmer, A.M. Beese, Anisotropic tensile behavior of Ti–6Al–4V components fabricated with directed energy deposition additive manufacturing, Acta Mater. 87 (2015) 309–320. [5] K.G. Prashanth, S. Scudino, J. Eckert, Defining the tensile properties of Al-12Si parts produced by selective laser melting, Acta Mater. 126 (2017) 25–35. [6] E. Cakmak, M.M. Kirka, T.R. Watkins, R.C. Cooper, K. An, H. Choo, W. Wu, R.R. Dehoff, S.S. Babu, Microstructural and micromechanical characterization of IN718 theta shaped specimens built with electron beam melting, Acta Mater. 108 (2016) 161–175. [7] Z.G. Zhu, Q.B. Nguyen, F.L. Ng, X.H. An, X.Z. Liao, P.K. Liaw, S.M.L. Nai, J. Wei, Hierarchical microstructure and strengthening mechanisms of a CoCrFeNiMn high
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[8]
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entropy alloy additively manufactured by selective laser melting, Scr. Mater. 154 (2018) 20–24. L. Sun, T.H. Simm, T.L. Martin, S. McAdam, D.R. Galvin, K.M. Perkins, P.A.J. Bagot, M.P. Moody, S.W. Ooi, P. Hill, M.J. Rawson, H.K.D.H. Bhadeshia, A novel ultra-high strength maraging steel with balanced ductility and creep resistance achieved by nanoscale β-NiAl and Laves phase precipitates, Acta Mater. 149 (2018) 285–301. Y. Wang, J. Sun, T. Jiang, Y. Sun, S. Guo, Y. Liu, A low-alloy high-carbon martensite steel with 2.6 GPa tensile strength and good ductility, Acta Mater. 158 (2018) 247–256. W. Xu, P.E.J. Rivera-Díaz-del-Castillo, W. Yan, K. Yang, D. San Martín, L.A.I. Kestens, S. van der Zwaag, A new ultrahigh-strength stainless steel strengthened by various coexisting nanoprecipitates, Acta Mater. 58 (2010) 4067–4075. H. Leitner, R. Schnitzer, M. Schober, S. Zinner, Precipitate modification in PH13-8 Mo type maraging steel, Acta Mater. 59 (2011) 5012–5022. E.A. Jägle, P. Choi, J. Van Humbeeck, D. Raabe, Precipitation and austenite reversion behavior of a maraging steel produced by selective laser melting, J. Mater. Res. 29 (2014) 2072–2079. J. Suryawanshi, K.G. Prashanth, U. Ramamurty, Tensile, fracture, and fatigue crack growth properties of a 3D printed maraging steel through selective laser melting, J. Alloy. Comp. 725 (2017) 355–364. Y. Bai, Y. Yang, D. Wang, M. Zhang, Influence mechanism of parameters process
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and mechanical properties evolution mechanism of maraging steel 300 by selective laser melting, Mater. Sci. Eng. A 703 (2017) 116–123. J. Mutua, S. Nakata, T. Onda, Z. Chen, Optimization of selective laser melting parameters and influence of post heat treatment on microstructure and mechanical properties of maraging steel, Mater. Des. 139 (2018) 486–497. K.G. Reddy, S. Arumugam, T.S. Lakshmanan, Hydrogen embrittlement of maragingsteel, J. Mater. Sci. 27 (1992) 5159–5162. L.P.M. Santos, M. Béreš, I.N. Bastos, S.S.M. Tavares, H.F.G. Abreu, M.J. Gomes Da Silva, Hydrogen embrittlement of ultra high strength 300 grade maraging steel, Corros. Sci. 101 (2015) 12–18. M. Béreš, L. Wu, L.P.M. Santos, M. Masoumi, F.A.M. Da Rocha Filho, C.C. Da Silva, H.F.G. de Abreu, M.J. Gomes Da Silva, Role of lattice strain and texture in hydrogen embrittlement of 18Ni (300) maraging steel, Int. J. Hydrogen Energy 42 (2017) 14786–14793. A.E. Pontini, J.D. Hermida, X-Ray diffraction measurement of the stacking fault energy reduction induced by hydrogen in an AISI 304 steel, Scr. Mater. 37 (1997) 1831–1837. I.M. Robertson, The effect of hydrogen on dislocation dynamics, Eng. Fract. Mech. 68 (2001) 671–692. Y.H. Fan, B. Zhang, H.L. Yi, G.S. Hao, Y.Y. Sun, J.Q. Wang, E.H. Han, W. Ke, The role of reversed austenite in hydrogen embrittlement fracture of S41500 martensitic stainless steel, Acta Mater. 139 (2017) 188–195.
Contents lists available at ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Hydrogen embrittlement behaviors of additive manufactured maraging steel investigated by in situ high-energy X-ray diffraction
T
Shilei Lia, Mingming Liua, Yang Renb, Yandong Wanga,∗ a
Beijing Advanced Innovation Center for Materials Genome Engineering, State Key Laboratory for Advanced Metals and Materials, University of Science and Technology Beijing, Beijing, 100083, China b X-ray Science Division, Advanced Photon Source, Argonne National Laboratory, Argonne, Illinois, 60439, USA
A R T I C LE I N FO
A B S T R A C T
Keywords: Additive manufacturing Hydrogen embrittlement High-energy X-ray diffraction Martensitic transformation
The microstructures and mechanical properties of additively manufactured (AM) maraging steel were investigated before and after hydrogen charging. The hydrogen-charged specimen suffers severe hydrogen embrittlement and its strength and plasticity decline obviously. The mechanism of ductile-to-brittle transition of AM maraging steel after hydrogen charging was studied by in situ high energy X-ray diffraction and electron backscatter diffraction techniques. Martensitic transformation occurs in the hydrogen-charged specimen under low stress. The premature phase transformation along with hydrogen-induced microcracks resulting in the brittle fracture of the hydrogen-charged specimen.
1. Introduction Additive manufacturing (AM) is fast gaining acceptance as a way to manufacture structural components with high structural integrity at a low cost [1–3]. Several AM alloys with mechanical properties comparable to alloys conventionally manufactured have been produced over the past years by reducing the residual porosity, insufficient layer bonding and hot cracking [4–7]. Selective laser melting (SLM), a typical metal AM technology, builds up three dimensional (3D) parts layer-bylayer out of powder particles. Maraging steels have an excellent combination of ultra-high strength with good toughness and ductility and therefore are widely applied in the aircraft, aerospace and tooling industries [8–11]. Maraging steels, consist of the martensitic matrix and various precipitates, are typically manufactured via a series of complex processes: casting, homogenizing, forging, quenching, machining, hot rolling, and a final aging treatment [9,10]. The as-produced maraging steels by SLM usually achieve significantly higher strength than conventional processing methods without heat treatment, because of their quenched martensitic matrix with fine cellular microstructures induced by the high cooling rate during SLM [3,12]. The microstructure and mechanical properties of maraging steels manufactured by SLM, as well as the effects of heat treatment have been investigated in the previous literature [13–15]. Conventional maraging steels are susceptible to hydrogen
∗
embrittlement and their mechanical properties suffer a great loss [16,17]. In this work, the hydrogen embrittlement behaviors of AM maraging 300 steel were investigated. The synchrotron-based high-energy X-ray diffraction (HE-XRD) technique was used to reveal the deformation and damage mechanisms of both the hydrogen-free and hydrogen-charged specimens. 2. Materials and experimental procedures The AM grade 300 maraging steel used in this study was fabricated by a SLM facility EOS M290 (EOS GmbH, Krailling, Germany) equipped with a 400 W Yb-fiber laser. Spherical gas atomized powders with particle sizes ranging from 20 to 45 μm were selected and the following parameters were used during the building of AM 316 L: laser power (350 W), beam size (100 μm), layer thickness (40 μm), scanning speed (1100 mm/s) and line spacing (150 μm). A pillar with a geometry of 10 mm (length) × 10 mm (width) × 50 mm (height) was built and no post-built heat treatment was applied. The microstructure of the as-fabricated specimen perpendicular to the building direction (BD) was observed using an optical microscope (OM) and a ZEISS Supra 55 scanning electron microscope (SEM) equipped with an electron backscatter diffraction (EBSD) detector. Hydrogen was introduced into the specimens by electrochemical charging with current density 50 mA/cm2 at room temperature in a 3 wt% NaCl +0.3 wt% NH4SCN solution for 24 h. The concentration of
Corresponding author. E-mail address: [email protected] (Y. Wang).
https://doi.org/10.1016/j.msea.2019.138341 Received 28 May 2019; Received in revised form 13 August 2019; Accepted 27 August 2019 Available online 28 August 2019 0921-5093/ © 2019 Elsevier B.V. All rights reserved.
Materials Science & Engineering A 766 (2019) 138341
S. Li, et al.
hydrogen was 5 wt ppm, measured using the pulse heating inert gas fusion infrared absorption method. Both the hydrogen-free and hydrogen-charged AM steels were cut perpendicular to the BD into flat dog-bone shaped tensile specimens, with a gauge length of 10 mm, width 3 mm and thickness 0.5 mm. The in situ tensile experiments were performed at beamline 11-ID-C of the Advanced Photon Source (APS) at Argonne National Laboratory. A monochromatic X-ray beam (500 × 500 μm) with an energy of 106 keV (with a wavelength of 0.11730 Å) was used to investigate the evolution of lattice strain and diffraction intensity during tensile loading. The fracture surface morphologies of both hydrogen-free and hydrogencharged specimens were observed in SEM. The deformed microstructures near the fractures of the tensile specimens were investigated. During the in situ tensile testing, a two-dimensional (2D) detector was used to collect the diffraction patterns at different strain levels. The specimen was strained incrementally at room temperature until fracture. The accurate distance between the specimen and the detector was calibrated by Ceria (CeO2) powders. The loading direction (LD) and transverse direction (TD) correspond to azimuth angles ϕ = 90° and ϕ = 180°, respectively. The 2D diffraction patterns were caked and integrated along specified azimuth angles over a range of ± 5° to obtain one-dimensional HE-XRD diffraction spectrums. A Pseudo-Voigt function was used to determine the position and the integrated intensity of the diffraction peaks. The lattice strain evolution in the retained austenite (RA) and martensite in the studied steel can be obtained by calculating the lattice spacing change during deformation using the function εhkl= (dhkl-d0)/d0, where εhkl is the lattice strain of a specific {hkl} plane, and dhkl and d0 are the interplanar spacings with and without stress.
Fig. 2. Engineering stress-strain curves of the hydrogen-free and hydrogencharged AM grade 300 maraging steel.
specimens before deformation. Different diffraction peaks with various interplanar spacings are identified as those belonging to the bodycentered cubic (BCC) structured martensite (M), as well as the retained austenite (A) with a face-centered cubic (FCC) structure. Comparisons among the diffraction peaks for the hydrogen-free and hydrogencharged AM maraging steel specimens suggest that no phase transformation occurs during hydrogen charging. 3.2. Macroscopic tensile behavior
3. Results and discussion
Fig. 2 shows the engineering stress-strain curves of the hydrogenfree and hydrogen-charged specimens tested at room temperature. Compared with the hydrogen-free specimen, the hydrogen-charged specimen has obvious lower tensile strength and elongation after fracture. The hydrogen-charged specimen fails without obvious plastic deformation and its ultimate tensile strength is lower than the yield strength of the hydrogen-free specimen. The macroscopic and microscopic morphologies of the fracture surfaces of the hydrogen-free and hydrogen-charged specimens, as
3.1. Microstructures Fig. 1a and b shows the OM microstructure and inverse pole figure (IPF) of the as-fabricated AM grade 300 maraging steel. Similar to a conventional grade 300 maraging steel [18], the AM steel also has a martensitic matrix with uneven coarse plate martensites. Fig. 1c shows the HE-XRD profiles of the hydrogen-free and hydrogen-charged
Fig. 1. OM (a) and IPF (b) micrographs of the studied AM grade 300 maraging steel. (c) HE-XRD profiles of the hydrogen-free and hydrogen-charged specimens before deformation. 2
Materials Science & Engineering A 766 (2019) 138341
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Fig. 3. The macroscopic and microscopic morphologies of the fracture surfaces of the hydrogen-free and hydrogen-charged specimens showing the obvious hydrogen embrittlement.
hydrogen-free and hydrogen-charged specimens with the applied stress along the LD are displayed in Fig. 5. In the elastic deformation stage, lattice strains of all reflections increase approximately linearly with increasing the applied stress. The slopes of the lattice strain vs. stress curves of reflections for martensite and RA in the hydrogen-charged specimen are almost the same as those in the hydrogen-free specimen. The evolution of diffraction intensity with the applied stress reflects the deformation and phase transformation information of phases in the material. For the FCC-structured RA and BCC-structured martensite, their intensities of A111/A200 and M110 reflections increase during uniaxial tensile deformation because of the development of fiber texture. While the martensitic transformation from RA to the newly formed martensite (NFM) will cause the decline of intensity for RA. As the small volume of RA in the studied material, this transformation may not significantly increase the intensity of martensite. From the evolution of normalized integral intensity, as shown in Fig. 6, it is easy to identify the critical stress for martensitic transformation in the hydrogen-free and hydrogen-charged specimens. The RA in the hydrogen-free specimen begins to transform to NFM when the applied stress reaches 760 MPa. This critical stress of martensitic transformation is lower than the macro yielding of the hydrogen-free specimen because the RA is easy to deform and has a lower micro yield stress. The presence of hydrogen can decrease the stacking fault energy of austenite and
shown in Fig. 3, confirm the transition from ductile to brittle in the AM maraging steel induced by hydrogen charging. Non-uniform dimples, as the typical feature of ductile fracture, were observed on the fracture surface of the hydrogen-free specimen. SEM fractography of the hydrogen-charged specimen reveals a brittle fracture mode which presents a mixture of dimples and cleavage planes. In addition, a large number of secondary cracks were observed on the fracture surface of the hydrogen-charged specimen. 3.3. Micromechanical behavior during in situ HE-XRD In situ HE-XRD was used to investigate the deformation behaviors of both hydrogen-free and hydrogen-charged specimens during tensile loading. Fig. 4 shows the 1D HE-XRD profiles of the specimens before and after hydrogen charging along the LD under different loadings. The peak positions along the LD continually shift to the higher value of dspacing during the tensile process. The diffraction intensity of reflections belong to RA along the LD in the hydrogen-free specimen decreases with increasing applied stress during tensile loading, and the peak finally merges with the background. However, the diffraction peaks of RA along the LD in the hydrogen-charged specimen remains almost unchanged during deformation until failure, as shown in Fig. 4b. The evolution of lattice strain of martensite and RA in both
Fig. 4. The 1D HE-XRD profiles of the hydrogen-free (a) and hydrogen-charged (b) specimens along the LD under different loadings (bottom to top represents 0 MPa to fracture). 3
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Fig. 5. Lattice strain evolution of the martensite and RA in the hydrogen-free (a) and hydrogen-charged (b) specimens with the applied stress along the LD.
promote the transformation of RA [19,20]. Some RA grains in the hydrogen-charged specimen may begin to plastic deform locally and transform to NFM under a very low applied stress [21]. In the present study, the hydrogen-charged specimen has much lower critical stress of 300 MPa for martensitic transformation compared with the hydrogenfree specimen.
3.4. Cracking in the hydrogen-charged specimen The proportion of RA in the hydrogen-free specimen significantly decreases after deformation, while the proportion for the deformed hydrogen-charged specimen is similar to the undeformed condition, as shown in Fig. 4. This confirms that the premature martensitic transformation in the hydrogen-free specimen only occurs in some local areas. Fig. 7 shows the evolution of normalized full width at half maximum (FWHM) for martensite and RA along the LD of the hydrogen-free and hydrogen-charged specimens. The normalized FWHM for A111 reflection of the hydrogen-charged sample has an obvious decrease at 464 MPa, which is higher than the critical stress for the martensitic transformation. The formation of NFM introduces localized stress fields and hydrogen atoms rapidly diffuse under localized stress gradients. High concentration of hydrogen may decrease the cohesive force between the atoms in the crystal and cause localized plastic deformation or cleavage [17], leading to the cracking along the {111}⊥LD planes in the RA. Longitudinal section morphologies near the fractures, cut along the LD from the fractured tensile specimens, are investigated in OM.
Fig. 7. Normalized FWHM for martensite and RA along the LD of the hydrogenfree and hydrogen-charged specimens.
Secondary cracks were observed in the hydrogen-charged specimen near the tensile fracture, as shown in Fig. 8b. In contrast, no crack was observed on the longitudinal section near the fracture of the hydrogenfree specimen (Fig. 8a). The IPF in Fig. 9 shows the orientation distribution vertical to the LD of the cracking region near the fracture of the hydrogen-charged specimen. The RA in the hydrogen-charged fractured specimen was not
Fig. 6. Normalized integral intensity evolution of the martensite and RA in the hydrogen-free (a) and hydrogen-charged (b) specimens with the applied stress along the LD. 4
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Fig. 8. OM micrographs of the longitudinal section near the fractures of the hydrogen-free (c) and hydrogen-charged (d) specimens.
treatments by combining annealing, solution treatment and aging is expected to improve the hydrogen embrittlement resistance of AM maraging steel. 4. Conclusions The deformation and damage mechanism of AM maraging steel before and after hydrogen charging was investigated using the in situ HE-XRD technique. Results reveal that the hydrogen-charged AM specimen suffers a loss in both plasticity and strength. The observation of fracture morphologies also confirms the ductile-to-brittle transition after hydrogen charging. The presence of hydrogen causes the locally premature martensitic transformation and cracking along the {111}⊥LD planes in the RA. Acknowledgements Fig. 9. The IPF of the cracking region near the fracture of the hydrogen-charged specimen.
This work was supported by the National Key Research and Development Program of China (No. 2017YFA0403804), the National Natural Science Foundation of China (NSFC) (Nos. 51601013 and 51527801) and the State Key Laboratory for Advanced Metals and Materials (Grant No. 2018Z-25). Use of the Advanced Photon Source is supported by the U. S. Department of Energy, Office of Science, Office of Basic Energy Sciences, under Contract No. DE-AC02-06CH11357.
observed in the EBSD result, while its existence is confirmed by the HEXRD results in Fig. 4b. It is considered that the RA in the EBSD specimen surface transfers into martensite due to the mechanical polishing. As RA and the martensite matrix in maraging steel usually follows the NW orientation relationship {111}A //{110}M [18]. The IPF of the cracked region in Fig. 9 indicates that the microcracks near the fracture of the hydrogen-charged specimen are along the {110}⊥LD planes of martensite. Thus, these microcracks initiate and propagate along the {111}⊥LD planes of RA, which is consistent with the HE-XRD results in Fig. 7. These microcracks release the localized stress fields in RA and thus inhibit the further phase transformation from RA to NFM. Under the applied stress, these microcracks coalesce and propagate quickly, resulting in the brittle fracture of the hydrogen-charged specimen. The content of RA will change and intermetallic particles may precipitate in the AM maraging steel during the post heat treatment, and the microstructure evolution significantly affects the mechanical properties [15]. The investigation on the hydrogen embrittlement behaviors of the conventional 300 grade maraging steel after different heat treatments indicated that the ranking of hydrogen embrittlement susceptibility in the decreasing order was peak-aged, over-aged and solution annealed condition [17]. It was reported that the elongation of conventional grade 300 maraging steel decreased from 8% to 0.3% under hydrogen charging [18]. In the present study, the AM maraging steel has a decline in elongation after hydrogen charging from 14.3% to 2.5%. The local residual stress, the content of RA and precipitation of intermetallic particles are believed to affect the hydrogen embrittlement behaviors of the AM maraging steel. Optimizing the heat
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