International Journal of Minerals, Metallurgy and Materials Volume 23, Number 6, June 2016, Page 667 DOI: 10.1007/s12613-016-1279-z
Effect of shot peening on hydrogen embrittlement of high strength steel Xin-feng Li1), Jin Zhang2), Ming-ming Ma3), and Xiao-long Song1) 1) State Key Laboratory for Mechanical Behavior of Materials, Xi’an Jiaotong University, Xi’an 710049, China 2) Department of Geosciences, State University of New York, Stony Brook, New York 11794-2100, USA 3) Department of Heat Treatment and Surface Treatment, AVIC Aircraft Co., Ltd. Xi’an Break Branch, Xingping 713106, China (Received: 5 September 2015; revised: 7 January 2016; accepted: 13 January 2016)
Abstract: The effect of shot peening (SP) on hydrogen embrittlement of high strength steel was investigated by electrochemical hydrogen charging, slow strain rate tensile tests, and hydrogen permeation tests. Microstructure observation, microhardness, and X-ray diffraction residual stress studies were also conducted on the steel. The results show that the shot peening specimens exhibit a higher resistance to hydrogen embrittlement in comparison with the no shot peening (NSP) specimens under the same hydrogen-charging current density. In addition, SP treatment sharply decreases the apparent hydrogen diffusivity and increases the subsurface hydrogen concentration. These findings are attributed to the changes in microstructure and compressive residual stress in the surface layer by SP. Scanning electron microscope fractographs reveal that the fracture surface of the NSP specimen exhibits the intergranular and quasi-cleavage mixed fracture modes, whereas the SP specimen shows only the quasi-cleavage fractures under the same hydrogen charging conditions, implying that the SP treatment delays the onset of intergranular fracture. Keywords: high strength steel; shot peening; hydrogen embrittlement; cracking; fracture; inclusions
1. Introduction Because of excellent strength and toughness, high strength steels have been widely used; however, the occurrence of hydrogen embrittlement (HE) limits their further applications. It has been established that the dependence of HE on strength levels can be divided into two stages. For low and medium strength steels (tensile strength <1000 MPa), the HE is less affected by strength. Meanwhile, when the tensile strength is greater than 1000 MPa, HE increases with strength [1–5]. Therefore, an urgent requirement exists to investigate the interaction between hydrogen and high strength steels. Shot peening technology, a mechanical surface modification technique, can induce compressive residual stresses and lead to the formation of microstructures in the shot peening layer of steels. Thus, the effect of peening treatment on the HE of various metals has been studied. However, the results are inconclusive. For instance, Takakuwa and Soyama [6]
demonstrated that the compressive residual stress introduced by cavitation peening suppressed the invasion of hydrogen for 316L austenite stainless steel; they also found that the crack propagation rate was suppressed by 65% after cavitation peening [7]. Similar results were found in Ref. [8]. However, Marchi et al. [9] revealed that laser peening enhanced the hydrogen-assisted fracture of Alloy 22. Almost no studies have been investigated for the effect of shot peening on HE of high strength steels. Thus, the HE of high-strength steels with or without shot peening was studied by a serious of tests, including electrochemical hydrogen charging, slow strain rate tests (SSRT), and hydrogen permeation tests.
2. Materials and methods Screw-thread high strength steel bars were used, and the chemical composition of the steel is shown Table 1. The steel bars were subjected to heat treatment. The steel billets of 150 mm × 150 mm in size were hot-rolled to 40 mm,
Corresponding author: Xin-feng Li, E-mail:
[email protected]; Xiao-long Song, E-mail:
[email protected] © University of Science and Technology Beijing and Springer-Verlag Berlin Heidelberg 2016
668
Int. J. Miner. Metall. Mater., Vol. 23, No. 6, June 2016
air-cooled to room temperature, and then tempered at 350°C for 5 h. The cylindrical specimens, with a gauge dimension of 5 mm in diameter and 25 mm in length, were machined along the longitudinal direction of the bars and mechanically ground with #800 SiC grit paper. One group of specimens was subjected to the shot peening (SP) process, whereas the other was not (no shot peening, NSP). The applied shot peening parameters were 0.5 MPa pressure and 100% coverage. For the hydrogen permeation tests, the specimens with 0.88 mm thickness were prepared by electrical discharge machining from the steel bars. Subsequently, the surfaces of specimens were mechanically ground to a finish by #800 SiC grit paper. The specimens were then rinsed with deionized water and degreased with alcohol. Table 1. Chemical composition of the high strength steel wt% C
Si
Mn
P
S
Cr
Ni
Cu
Al
Fe
0.27 1.55 2.20 0.011 0.0043 0.97 0.007 0.017 0.008 Balance
The hydrogen permeation instrumentation was composed of the electrolytic cell with two compartments (cathodic and anodic sides), the reference electrode (Hg/HgO/NaOH in 0.1 mol/L NaOH), two auxiliary electrodes (Pt plate), and two potentiostats/galvanostats. A specimen with an exposed surface area of 2.27 cm2 on each side was mounted between the compartments. One side of the specimen acted as the hydrogen entry side. It was galvanostatically polarized at a constant charging current density (20 mA/cm2) in 0.5 mol/L H2SO4 with 0.824 g/L Na4P2O7. The hydrogen exit side of the cell was potentiostatically maintained at a constant potential of 200 mV vs. the reference electrode. The flux of hydrogen through the specimen was measured in terms of the steady-state current density ( I P , A·m2) and then converted to the steady-state hydrogen permeation flux (J, mol·m−2·s1) according to the following equation [10–11]. I P (1) nF The apparent hydrogen diffusivity (Dapp, m2·s1) was determined by the breakthrough method as J
Dapp
L2 6tL
(2)
If the charging surface was in equilibrium, the apparent subsurface concentration (Capp, mol·m3) could be calculated by Capp
J L Dapp
(3)
where n represents the number of electrons transferred; F the Faraday constant, C·mol1; L the specimen thickness, m; and tL the lag time, s, defined as 0.63 times of the steady-state value (t0.63). Hydrogen was introduced to the specimens by electrochemical pre-charging in a solution of 0.5 mol/L NaOH at a cathodic current density ranging from 0 to 10 mA/cm2 for 24 h. To prevent the release of hydrogen from the specimens, cadmium electroplating was conducted on the surface of specimens [1213]. SSRT was performed on an Instron1195 electronic tensile testing machine with a constant strain rate of 0.03 mm/min, equivalent to a normal strain rate of 2 × 105 s1. The HE index (HEI) of relative susceptibility at a given cathodic current density was determined by measuring the reduction of area loss, which can be expressed as H 100% HEI 0 (4)
0
where 0 and H are the area reduction of hydrogen-uncharged and hydrogen-charged specimens, respectively. The tensile fracture surfaces were observed using a Hitachi SU6600 field emission scanning electron microscope (SEM) with an energy dispersive spectrometer (EDS).
3. Results 3.1. Microscopic observations The microstructures of steels are shown as the tempered martensite in Fig. 1(a). Fig. 1(b) shows the spherical inclusions in uniform distribution of the steel. The EDS analysis in Fig. 1(c) shows that the inclusions are primarily composed of O–Al–Si–Ca. Figs. 1(d) and (e) display the transmission electron microscope (TEM) images of the surface layer in NSP and SP specimens, respectively, along the transversal section. It can be seen that the NSP specimen exhibits lath martensite, with lath spacing of approximately 250 nm; meanwhile, a large number of dislocations form at the interior or lath boundaries of martensite in the SP specimen, and the pattern of lath martensite almost disappears. Fig. 2 presents the typical surfaces of NSP and SP steel. Compared with the NSP steel, the surface roughness of the SP steel increases. 3.2. Microhardness measurements A microhardness test was performed by micro Vickers hardness tester with a load of 0.98 N and a dwell time of 10 s. The measurements were performed at each depth three times and the averages were collected for measurement errors.
X.F. Li et al., Effect of shot peening on hydrogen embrittlement of high strength steel
669
Fig. 1. Microstructures of steels and inclusions: (a) optical image of NSP specimens; (b) inclusions in the steel; (c) EDS analysis of inclusions; (d) and (e) TEM images of NSP and SP specimens, respectively, along the transversal section.
Fig. 2. Microgeometrical aspects of the specimen surfaces: (a) NSP specimen; (b) SP specimen.
The hardness distribution on the cross section of specimens is illustrated in Fig. 3. The results indicate that the microhardness is higher on the surface of the SP specimens
and then gradually decreases from the outer surface to the core. The hardness of the base materials is approximately Hv 530 on average. After SP treatment, the surface hardness
670
increases up to Hv 613. The enhancement of microhardness is due to plastic deformation introduced by SP treatment. The effect of SP on HV is notable, reaching a depth of approximately 250 μm.
Int. J. Miner. Metall. Mater., Vol. 23, No. 6, June 2016
strength first decreases and then increases for both the NSP and SP specimens, as shown in Fig. 6(a). However, both the elongation and the reduction of area decrease, especially at high current density, and the loss extent of ductility of the NSP specimens is higher than that of the SP specimens, as shown in Figs. 6(b) and (c). Fig. 6(d) displays the relationship between HEI and current density, indicating that the SP specimens show higher resistance to hydrogen embrittlement in comparison with the NSP specimens. In addition, Table 2 provides the hydrogen permeability data of the NSP and SP specimens. In comparison with the NSP specimen, SP treatment sharply decreases the hydrogen diffusivity from 4.43 × 1011 to 2.04 × 1011 m2·s1 and increases the apparent subsurface concentration from 12.21 to 15.48 mol·m3.
Fig. 3. Hardness variation from the surface to the interior of the NSP and SP specimens.
3.3. Residual stress measurements To determine the macroscopic residual stress of the specimens, X-ray stress analysis was performed on the NSP and SP specimens. The residual stress was determined by the sin2ψ method. In-depth measurement of the residual stress was conducted by chemical corrosion method. A solution of aqua regia, which consisted of hydrochloric acid and fuming nitric acid with a volume ratio of 3:1, was used to thin the specimens. The relationships between the residual stress and depth from the surface are shown in Fig. 4. With increasing depth from the surface, the compressive residual stress decreases for both the NSP and SP specimens. Moreover, it can be seen that the maximum compressive residual stress of the NSP specimen is 225 MPa, whereas it reaches a value of 706 MPa for the SP specimen. Note that the depth of the compressive residual stress layer is about 300 μm for the SP specimen, which is more than three times deeper than that of the NSP specimen.
Fig. 4. Relationship between residual stress and depth for NSP and SP specimens.
3.4. Tensile tests Fig. 5 shows the stress–displacement diagram of the hydrogen-charged or uncharged specimens with NSP and SP treatment at different cathodic current densities. With increasing current density, the displacement of fracture decreases regardless of the NSP and SP specimens. Compared with the NSP specimens at the same cathodic current density, the curve of the SP specimens exhibits a large amount of necking. With increasing current density, the tensile
Fig. 5. Stress–displacement diagram for hydrogen-charged and hydrogen-uncharged NSP and SP treatment specimens at different cathodic current densities.
3.5. Fractograph The SEM fractograph of the hydrogen-uncharged NSP
X.F. Li et al., Effect of shot peening on hydrogen embrittlement of high strength steel
671
Fig. 6. Mechanical properties of the NSP and SP hydrogen-charged specimens at various cathodic current densities: (a) tensile strength; (b) elongation at fracture; (c) reduction of area; (d) HEI. Table 2. Data of experimental hydrogen permeability Treatment
t0.63 / s
J∞ / (mol·m2·s1)
Dapp / (m2·s1)
Capp / (mol·m3)
NSP
2916
6.15 × 107
4.43 × 1011
12.21
6333
7
2.04 × 1011
15.48
SP
3.59 × 10
specimen is shown in Fig. 7. The crack initiates at the center of the specimen (indicated by a white arrow in Fig. 7(a)). High magnification observation in Fig. 7(b) shows the dimples in a range of sizes, revealing a ductile fracture. The propagation region comprises both small dimples and many transgranular secondary cracks, as shown by white arrows in Fig. 7(c)). The shear lip region is composed of finer dimples than those in the crack initiation and propagation regions, as shown in Fig. 7(d). Similar results are also observed in the hydrogen-uncharged SP specimen. Fig. 8 shows the fracture surfaces of the hydrogen-charged NSP specimen at a current density of 1 mA/cm2. The crack nucleation site is located at the inclusion indicated by a black arrow in Fig. 8(a). High magnification observation of the inclusion at the crack origin is presented in Fig. 8(b), implying that the inclusion is not cohesive with the matrix. In addition, the crack initiation site exhibits a mixed fracture mode as quasi-cleavage fractures (QC) and
intergranular fractures (IG) accompanied with the secondary cracks (indicated by white arrows), as shown in Fig. 8(c). High magnification observation of the secondary crack is shown in Fig. 8(d), revealing that the crack growth path is along the prior austenitic grain boundaries. Fig. 9 presents the fracture surfaces of the hydrogen-charged SP specimen at a current density of 1 mA/cm2. High magnification observation of the inclusions is given in Fig. 9(b); it can be observed that the interface between the inclusion and the matrix is incohesive. Unlike the fracture feature of the NSP specimen in Fig. 8, it is composed of quasi-cleavage and dimple fractures, as shown in Fig. 9(c).
4. Discussion 4.1. Effect of shot peening on resistance to hydrogen embrittlement The hydrogen embrittlement phenomenon is related to the hydrogen diffusion. The hydrogen diffusion process is delayed by the presence of traps, such as vacancies, dislocations, interfaces, and microvoids. Addach et al. [14] compared the hydrogen diffusion coefficients of iron which were mechanically polished (hydrogen diffusivity = ~1011 m2·s1) and annealed (hydrogen diffusivity = ~109 m2·s1), and
672
Int. J. Miner. Metall. Mater., Vol. 23, No. 6, June 2016
Fig. 7. SEM fractographs of the hydrogen-uncharged NSP specimen: (a) overview of the fracture surface; (b), (c), and (d) high magnification images of the indicated region in Fig. 7(a).
Fig. 8. SEM fractographs of the hydrogen-charged (1 mA/cm2) NSP specimen: (a) overview of the fracture surface; (b) inclusions at crack origin; (c) and (d) high magnification images of the indicated region in Fig. 8(a).
suggested that mechanical polishing induced a striking decrease in the hydrogen diffusivity because it modified the microstructure and led to strain hardening. Rivera et al. [15] studied hydrogen trapping and diffusion properties in
as-received and cold-rolled state API 5L X60 steel via electrochemical hydrogen permeation tests; they found that the development of the transient permeation was retarded in cold rolled steel in comparison with the as-received steel.
X.F. Li et al., Effect of shot peening on hydrogen embrittlement of high strength steel
This was due to the creation of traps by cold-rolling. Huang et al. [16] revealed that the hydrogen diffusivity decreased with increasing plastic deformation, while elastic deforma-
673
tion had no significant effect; the hydrogen trapping by dislocations was the dominant process in plastic deformation.
Fig. 9. SEM fractographs of the hydrogen-charged (1 mA/cm2) SP specimen: (a) overview of the fracture surface; (b) and (c) high magnification images of the indicated region in Fig. 9(a).
During hydrogen charging, the hydrogen traps of the surface layer are firstly filled with hydrogen, and then hydrogen atoms diffuse into the center of the specimens. The dislocation density of the shot peening layer increases dramatically, as shown in Fig. 1(e). Therefore, the existence of the shot peening layer suppresses the hydrogen invasion into the interior of specimens. On the other hand, the hydrogen diffusion behavior is affected by hydrostatic stress, and then the hydrogen concentrates at the crack tip or elastic-plastic boundary [17–18]. The hydrostatic stress at the surface is associated with residual stress. When tensile residual stress is introduced, it is possible to accelerate the hydrogen invasion. Compressive residual stress reduces hydrostatic stress, reduces lattice spacing, and then suppresses hydrogen invasion. Numerical simulation results [19] revealed that residual stress had a significant impact on the stress field around the crack tip. The hydrogen concentration around the crack tip increased with tensile residual stress but was suppressed under compressive residual stress. Brass et al. [20] suggested that the cold-worked superficial layer caused by shot peening generated a more homogeneous distribution of hydrogen and the compressive stress suppressed crack initiation. Owing to the high hydrogen trapping and compressive
residual stress introduced by SP, the SP specimens exhibit higher resistance to hydrogen embrittlement in comparison with NSP specimens under the same hydrogen-charging conditions. 4.2. Effect of shot peening on hydrogen induced fracture Independent of NSP and SP, the crack origin of the hydrogen-charged specimens is located at the inclusion site, as shown in Figs. 8(a) and 9(a). Inclusions are termed to be irreversible hydrogen traps. After hydrogen charging, a distinct aggregation of hydrogen at the interface between the matrix and the inclusions has been reported [21–22]. On the other hand, the inclusions enriched with O–Al–Si–Ca are incohesive with the matrix, as shown in Figs. 8(b) and 9(b). During the tensile tests, the uncoordinated deformation between the matrix and the inclusions causes a high local stress concentration at the interface. Therefore, hydrogen atoms diffuse toward the interface, driven by the stress gradient; this results in a greater local hydrogen concentration around the stress concentration sites. When the hydrogen concentration at the interface reaches a critical value, a micro-crack forms. Under the same hydrogen charging current density, the
674
Int. J. Miner. Metall. Mater., Vol. 23, No. 6, June 2016
fracture surface of the NSP specimen shows the intergranular and quasi-cleavage fracture modes, as shown in Fig. 8(c), while only quasi-cleavage is observed on the SP specimen fracture surface, as shown in Fig. 9(b). The onset of intergranular fracture depends on the diffusible hydrogen content. Wang et al. [23] indicated that the intergranular fracture occurred when the diffusible hydrogen content was 0.5 × 106 for the strength level of 1050 MPa and less than 0.1 × 106 for the strength level of 1300 MPa. Martin et al. [24] reported that the percentage of intergranular fracture in nickel specimens was related to the amount of hydrogen that accumulated on the grain boundaries. The amount of hydrogen enrichment at the grain boundaries was controlled by either the aging time before testing or the hydrogen concentration in the bulk. After SP treatment, the dislocation density increases and residual compressive stress layers are introduced, as shown in Figs. 1 and 4, respectively. The shot peening layer serves as a barrier to hydrogen invasion by increasing hydrogen trapping and decreasing the hydrogen flux (Table 2). Therefore, the hydrogen concentration in the SP specimen is too low to generate intergranular fractures.
[2]
[3]
[4]
[5]
[6]
[7]
[8]
5. Conclusions (1) Compared with NSP steel, SP treatment enhances the resistance to hydrogen embrittlement of the steel. This is related to the change in the microstructure and the generation of residual stress in the surface layer after SP. (2) In comparison with the NSP steel, SP treatment has sharply decreased the apparent hydrogen diffusivity and increased the subsurface hydrogen concentration. (3) In the hydrogen-uncharged specimen, the crack origin is located at the center of the specimen, whereas the crack originates from inclusions in the hydrogen-charged specimen. (4) SEM fractograph results reveal that the NSP specimen surface exhibits intergranular and quasi-cleavage fractures, whereas the SP specimen shows only quasi-cleavage fractures. This implies that SP treatment delays the onset of intergranular fracture.
[9]
[10]
[11]
[12]
[13]
Acknowledgements
[14]
This work was financially supported by the Doctoral Research Assistant Foundation of Xi′an Jiaotong University.
[15]
References [1]
M.Q. Wang, E. Akiyama, and K. Tsuzaki, Determination of
[16]
[17]
the critical hydrogen concentration for delayed fracture of high strength steel by constant load test and numerical calculation, Corros. Sci., 48(2006), No. 8, p. 2189. Y.S. Ding, L.W. Tsay, M.F. Chiang, and C. Chen, Gaseous hydrogen embrittlement of PH 13-8 Mo steel, J. Nucl. Mater., 385(2009), No. 3, p. 538. T. Michler and J. Naumann, Microstructural aspects upon hydrogen environment embrittlement of various bcc steels, Int. J. Hydrogen Energy, 35(2010), No. 2, p. 821. A. Barnoush and H. Vehoff, Recent developments in the study of hydrogen embrittlement: Hydrogen effect on dislocation nucleation, Acta Mater., 58(2010), No. 16, p. 5274. X.F. Li, J. Zhang, P. Zhang, P. Li, and X.L. Song, Failure analysis of high strength steel bar used in a wind turbine foundation, J. Fail. Anal. Prev., 15(2015), No. 2, p. 295. O. Takakuwa and H. Soyama, Using an indentation test to evaluate the effect of cavitation peening on the invasion of the surface of austenitic stainless steel by hydrogen, Surf. Coat. Technol., 206(2012), No. 18, p. 3747. O. Takakuwa, T. Ohm, M. Nishikawa, Y.A. Toshimitsu, and H. Soyama, Suppression of fatigue crack propagation with hydrogen embrittlement in stainless steel by cavitation peening, Strength Fract. Complex., 7(2011), No. 1, p. 79. O. Takakuwa and H. Soyama, Suppression of hydrogen-assisted fatigue crack growth in austenitic stainless steel by cavitation peening, Int. J. Hydrogen Energy, 37(2012), No. 6, p. 5268. C.S. Marchi, T. Zaleski, S. Lee, N.Y.C. Yang, and B. Stuart, Effect of laser peening on the hydrogen compatibility of corrosion-resistant nickel alloy, Scripta Mater., 58(2008), No. 9, p. 782. C.F. Dong, K. Xiao, Z.Y. Liu, W.J. Yang, and X.G. Li, Hydrogen induced cracking of X80 pipeline steel, Int. J. Miner. Metall. Mater., 17(2010), No. 5, p. 579. C.F. Dong, Z.Y. Liu, X.G. Li, and Y.F. Cheng, Effects of hydrogen-charging on the susceptibility of X100 pipeline steel to hydrogen-induced cracking, Int. J. Hydrogen Energy, 34(2009), No. 24, p. 9879. X.F. Li, Y.F. Wang, P. Zhang, B. Li, X.L. Song, and J. Chen, Effect of pre-strain on hydrogen embrittlement of high strength steels, Mater. Sci. Eng. A, 616(2014), p. 116. X.F. Li, J. Zhang, Y.F. Wang, B. Li, P. Zhang, and X.L. Song, Effect of cathodic hydrogen-charging current density on mechanical properties of prestrained high strength steels, Mater. Sci. Eng. A, 641(2015), p. 45. H. Addach, P. Berçot, M. Rezrazi, and J. Takadoum, Study of the electrochemical permeation of hydrogen in iron, Corros. Sci., 51(2009), No. 2, p. 263. P.C. Rivera, V.P. Ramunni, and P. Bruzzoni, Hydrogen trapping in an API 5L X60 steel, Corros. Sci., 54(2012), p. 106. Y.L. Huang, A. Nakajima, A. Nishikata, and T. Tsuru, Effect of mechanical deformation on permeation of hydrogen in iron, ISIJ Int., 43(2003), No. 4, p. 548. Y. Liang, P. Sofronis, and R.H. Dodds Jr., Interaction of hy-
X.F. Li et al., Effect of shot peening on hydrogen embrittlement of high strength steel drogen with crack-tip plasticity: effects of constraint on void growth, Mater. Sci. Eng. A, 366(2004), No. 2, p. 397. [18] A.H.M. Krom, R.W.J. Koers, and A. Bakker, Hydrogen transport near a blunting crack tip, J. Mech. Phys. Solids, 47(1999), No. 4, p. 971. [19] O. Takakuwa, M. Nishikawa, and H. Soyama, Numerical simulation of the effects of residual stress on the concentration of hydrogen around a crack tip, Surf. Coat. Technol., 206(2012), No. 11-12, p. 2892. [20] A.M. Brass, J. Chene, G. Anteri, J. Ovejero-Garcia, and L. Castex, Role of shot-peening on hydrogen embrittlement of a low-carbon steel and a 304 stainless steel, J. Mater. Sci., 26(1991), No. 16, p. 4517.
675
[21] K. Horikawa, N. Ando, H. Kobayashi, and W. Urushihara, Visualization of hydrogen gas evolution during deformation and fracture in SCM 440 steel with different tempering conditions, Mater. Sci. Eng. A, 534(2012), p. 495. [22] M. Au, High temperature electrochemical charging of hydrogen and its application in hydrogen embrittlement research, Mater. Sci. Eng. A, 454-455(2007), p. 564. [23] M.Q. Wang, E. Akiyama, and K. Tsuzaki, Hydrogen degradation of a boron-bearing steel with 1050 and 1300 MPa strength levels, Scripta Mater., 52(2005), No. 5, p. 403. [24] M.L. Martin, B.P. Somerday, R.O. Ritchie, P. Sofronis, and I.M. Robertson, Hydrogen-induced intergranular failure in nickel revisited, Acta Mater., 60(2012), p. 2739.
