Effects of Thermal Aging and Absorbed Hydrogen on Fatigue Crack Growth Behavior of Welded C-Mn Steels M. A. Khattak 1 and Mohd N Tamin 2 Abstract— The relatively high sulfur content in C-Mn steels and the hydrogen-rich operating environment renders high susceptibility of the steels to hydrogen-induced cracking, particularly in the welded steel plate connections. The objective of this study is to examine effects of thermal aging and absorbed hydrogen on fatigue crack growth behavior of the heat affected zone (HAZ) in welded Type A516 Gr 70 steels. Results showed that the fatigue crack growth rate behavior of HAZ displays a threshold growth stage followed by a power-law growth region until final fracture of the specimen. The threshold ΔK th for HAZ (13.2 MPa√m) is lower than that for the base metal (15.5 MPa√m). Both thermally aged and hydrogen-absorbed specimens showed a lower threshold ΔK th of 11.4 and 12.2 MPa√m, respectively. In the high ΔK region near final fracture of the specimen, the fracture mode is transgranular with the formation of dimples and micro-voids. Keywords—Fatigue crack growth rate, heat affected zone, thermal aging, absorbed hydrogen, crack growth mechanism. I. INTRODUCTION HEMICAL reactor vessels and pipelines are commonly constructed using welded C-Mn (A516) steels and stainless steel liners. The operating temperature typically ranges from -29 to 427 °C. Thermal aging of these steels under prolonged exposure to intermediate service temperatures lead to deleterious effects such as embrittlement, loss of toughness and creep rupture of the steel [1]. Previous research showed that the ductile-to-brittle transition temperature (DBTT) increases with an increase in thermal aging temperature [2]. These conditions could result in failure of pressure vessels and pressure piping related accidents that are often fatal and involved loss of capital investment [3]-[4]. As most of the problems in heavy-section piping occur at welded joints therefore damage assessment techniques should focus on these regions rather than the base metal alone [5-6]. M. A. Khattak 1 and Mohd N Tamin 2 , Faculty of Mechanical Engineering, Universiti Teknologi Malaysia, 81310 UTM Skudai, Johor, MALAYSIA, Email: 1 [email protected], 2 [email protected]In this fusion zone the resulting microstructures is are very complex and comprised of as many as seven distinct regions [7]. During deposition of single pass welds in C-Mn and low- carbon micro alloyed steels, four microstructural regions in the HAZ can be identified as Coarse-grained HAZ (CGHAZ), Fine-grained HAZ (FGHAZ), Inter critical HAZ (ICHAZ) and Subcritical HAZ (SCHAZ) [8]. High thermal gradient and fast cooling rates across the HAZ render the zone susceptible to materials defects including hard inclusions, blisters by trapped gas and microcracks. In previous studies, highest number of microcracks was found in HAZ of the welded joint [6, 8]. These microcracks are likely to join to form a longer structural crack and propagate under applied fatigue loading. In such case, fatigue crack growth behavior of the welded joint is of prime concern. In addition, HAZ of welded 2.2Cr-1Mo steel joint exhibits inferior rupture strength compared to the base plate [9]. The formation of carbides further deteriorates the corrosion resistance of HAZ when compared to WM and BM [10]. Exposure to hydrogen-rich environment and thermal aging accelerate the deleterious effects on welded joint including the initiation and propagation of brittle cracks. Fatigue crack growth rates of numerous structural steels were enhanced in hydrogen environment [15]-[17]. In high strength steels, microstructure degradation by hydrogen embrittlement is a serious problem [13]-[14]. Cracking associated with a sour environment is classified as hydrogen-induced cracking (HIC)[13]. Failure in this case is similar to environmental cracking [14]. A number of factors such as pH, volume of hydrogen diffused, volume-fraction, and the shape of inclusions present, and the surrounding microstructure influence the process. Stress in this instance is not as critical compared with SCC and SSC [15]. Hydrogen reduces cleavage strength, adhesion strength and the resistance to shear slip; therefore, hydrogen assists intergranular cracking [16]. In controlled laboratory experiment hydrogen can be introduced in the specimen by gaseous hydrogen charging in an oxygen-free copper chamber with tantalum hydride at temperature in excess of 700 o C [18]. Others use high-pressure (20-35 MPa) hydrogen autoclave at 350 o C [19]. In the electrochemical method the metal specimen is soaked in dilute solution of sulphuric acid with a corrosion inhibitor. In this paper, effects of thermal aging and absorbed hydrogen on fatigue crack growth behavior of welded C-Mn C 2nd International Conference on Mechanical, Electronics and Mechatronics Engineering (ICMEME'2013) June 17-18, 2013 London (UK) 120
Transcript
Effects of Thermal Aging and Absorbed Hydrogen on Fatigue Crack Growth Behavior of
Welded C-Mn Steels
M. A. Khattak1 and Mohd N Tamin2
Abstract— The relatively high sulfur content in C-Mn steels and
the hydrogen-rich operating environment renders high susceptibility of the steels to hydrogen-induced cracking, particularly in the welded steel plate connections. The objective of this study is to examine effects of thermal aging and absorbed hydrogen on fatigue crack growth behavior of the heat affected zone (HAZ) in welded Type A516 Gr 70 steels. Results showed that the fatigue crack growth rate behavior of HAZ displays a threshold growth stage followed by a power-law growth region until final fracture of the specimen. The threshold ΔK th for HAZ (13.2 MPa√m) is lower than that for the base metal (15.5 MPa√m). Both thermally aged and hydrogen-absorbed specimens showed a lower threshold ΔK th of 11.4 and 12.2 MPa√m, respectively. In the high ΔK region near final fracture of the specimen, the fracture mode is transgranular with the formation of dimples and micro-voids.
I. INTRODUCTION HEMICAL reactor vessels and pipelines are commonly constructed using welded C-Mn (A516) steels and stainless steel liners. The operating temperature typically
ranges from -29 to 427 °C. Thermal aging of these steels under prolonged exposure to intermediate service temperatures lead to deleterious effects such as embrittlement, loss of toughness and creep rupture of the steel [1]. Previous research showed that the ductile-to-brittle transition temperature (DBTT) increases with an increase in thermal aging temperature [2]. These conditions could result in failure of pressure vessels and pressure piping related accidents that are often fatal and involved loss of capital investment [3]-[4]. As most of the problems in heavy-section piping occur at welded joints therefore damage assessment techniques should focus on these regions rather than the base metal alone [5-6].
M. A. Khattak1 and Mohd N Tamin2, Faculty of Mechanical Engineering, Universiti Teknologi Malaysia, 81310 UTM Skudai, Johor, MALAYSIA, Email: [email protected], [email protected]
In this fusion zone the resulting microstructures is are very complex and comprised of as many as seven distinct regions [7]. During deposition of single pass welds in C-Mn and low-carbon micro alloyed steels, four microstructural regions in the HAZ can be identified as Coarse-grained HAZ (CGHAZ), Fine-grained HAZ (FGHAZ), Inter critical HAZ (ICHAZ) and Subcritical HAZ (SCHAZ) [8].
High thermal gradient and fast cooling rates across the HAZ render the zone susceptible to materials defects including hard inclusions, blisters by trapped gas and microcracks. In previous studies, highest number of microcracks was found in HAZ of the welded joint [6, 8]. These microcracks are likely to join to form a longer structural crack and propagate under applied fatigue loading. In such case, fatigue crack growth behavior of the welded joint is of prime concern. In addition, HAZ of welded 2.2Cr-1Mo steel joint exhibits inferior rupture strength compared to the base plate [9]. The formation of carbides further deteriorates the corrosion resistance of HAZ when compared to WM and BM [10].
Exposure to hydrogen-rich environment and thermal aging accelerate the deleterious effects on welded joint including the initiation and propagation of brittle cracks. Fatigue crack growth rates of numerous structural steels were enhanced in hydrogen environment [15]-[17]. In high strength steels, microstructure degradation by hydrogen embrittlement is a serious problem [13]-[14].
Cracking associated with a sour environment is classified as hydrogen-induced cracking (HIC)[13]. Failure in this case is similar to environmental cracking [14]. A number of factors such as pH, volume of hydrogen diffused, volume-fraction, and the shape of inclusions present, and the surrounding microstructure influence the process. Stress in this instance is not as critical compared with SCC and SSC [15]. Hydrogen reduces cleavage strength, adhesion strength and the resistance to shear slip; therefore, hydrogen assists intergranular cracking [16].
In controlled laboratory experiment hydrogen can be introduced in the specimen by gaseous hydrogen charging in an oxygen-free copper chamber with tantalum hydride at temperature in excess of 700 oC [18]. Others use high-pressure (20-35 MPa) hydrogen autoclave at 350 oC [19]. In the electrochemical method the metal specimen is soaked in dilute solution of sulphuric acid with a corrosion inhibitor.
In this paper, effects of thermal aging and absorbed hydrogen on fatigue crack growth behavior of welded C-Mn
C
2nd International Conference on Mechanical, Electronics and Mechatronics Engineering (ICMEME'2013) June 17-18, 2013 London (UK)
steel plates are examined. The measured crack growth behavior is described in terms of crack growth rate da/dN, stress intensity factor range ∆K and threshold level ∆KTH. Fracture mechanisms due to thermal aging and absorbed hydrogen in the HAZ are identified and described. The applied stresses simulate a typical loading experienced by a cylindrical tube and a pre-existing defect in the form of a long crack is assumed to exist in HAZ which is subjected to an opening mode of cracking. The test is performed at ambient temperature.
II. MATERIALS AND EXPERIMENTAL PROCEDURES The material employed in this study is an ASTM A516-
Grade 70 steel. Cross section of the welded joint showing the different
regions of weld metal, HAZ and the base metal is shown in Figure 1(a). Figure 1(b) shows microstructures of the various phases of the welded A516 steel in the as-received condition. The base metal displays an aggregate of well-defined ferrite-pearlite bands. The HAZ is characterized by a network of coarser grain boundary ferrite with colonies of fine-grain acicular ferrite. Faster cooling of the weld metal region results in fine-grained Widmanstatten ferrite microstructure.
The variations of chemical composition and/or microstructure in the weld metal as well as in the heat-affected zone (HAZ) are crucial to the failure of the welded components in service. The microstructural dependence of the corrosion behavior of steel weldment has also been reported in literature [17].
The chemical composition of the BM, HAZ and WM is shown (see Table II), the balance being Fe. Tensile and yield strength of welded samples in the as-received condition is 480 and 360 MPa, respectively.
Extensive research has been carried out to investigate the microstructural and compositional dependences of mechanical properties [18]-[24] and environmentally assisted cracking resistance [25]-[32] of steel weldments. Only limited information is available in the literature concerning the corrosion behavior of steel weldments. Researchers indicated that Mn addition has a detrimental effect on HAZ corrosion of carbon steels in 0.5 N NaCl solution [33]. They also showed that Nb additions cause a significant increase in the HAZ corrosion in the lower (0.06 wt%) carbon steels, but have little effect on the higher (0.12 wt%) carbon steels. They also reported that the corrosion rate of underwater weld which had a high Mo content was lower than that of the base metal [34].
TABLE I SUMMARY OF WELDING PARAMETERS AND THERMAL CYCLE
Compact tension (CT) specimens were designed such that fatigue crack will propagate along the HAZ. Thermal aging of CT specimens was performed at 420 oC for durations of 500, 800 and 1200 hours. Hydrogen charging of polished and pre-cracked C-Mn CT specimens was carried out in an electrolytic
cell with a stainless steel rod as anode. This facilitates hydrogen attack at the crack tip. The electrolyte consists of dilute sulphuric acid solution (0.5M H2SO4 solution) containing 1.97 gm of arsenic trioxide (As2O3) as corrosion inhibitor. A constant current density of 20 mA/cm2 was maintained throughout the 9-hour charging process.
Fatigue pre-cracking procedure is employed to establish and initial crack, a, in the HAZ region of centre-cracked specimen [20].
TABLE II CHEMICAL COMPOSITION OF A516
GRADE 70 STEEL (WT. %) (wt. %) BM HAZ WM
C 0.263 0.236 0.295 Mn 1.12 1.10 1.45 P 0.014 0.011 0.035 Si 0.46 0.44 1.00 Cu 1.41 0.22 1.28 Ni 0.02 0.02 0.03 V 0.005 0.004 0.016 Mo 0.012 0.011 0.008 S 0.004 0.003 0.013 Cr 0.018 0.018 0.029 Sn 0.04 0.04 0.14
Microstructure study was performed on optical micrographs while scanning electron microscope (SEM) was employed for fractographic analysis.
(a)
BM HAZ WM
(b) Fig. 1(a) Cross-section of the fusion zone of the welded joint, (b) Microstructures of base metal (BM), HAZ and
weld region (WM) of as-received sample. Mag. 10X.
2nd International Conference on Mechanical, Electronics and Mechatronics Engineering (ICMEME'2013) June 17-18, 2013 London (UK)
121
III. RESULTS AND DISCUSSION A. Effects of absorbed hydrogen and aging on
microstructure changes
Hardness measurements across the fusion zone of the welded C-Mn steel joint were made as shown in Figure 3. A widely used hardness value below which it is generally agreed that hydrogen cracking is not expected to occur is 350 HV [35]-[36].
The absorbed hydrogen in the alloy degraded mechanical properties of the material as reflected in lower relative hardness of the phases, as shown in Figure 3(a). Average hardness measure for the hydrogen-embrittled base metal is 29 pct. lower at 106.2 than the as-received plates. Greater hardness reduction of 37 pct. from their respective value in the as-received condition is experienced by HAZ weld metal due to the absorbed hydrogen.
Effects of prolonged thermal exposure of welded C-Mn steel joint on hardness measures are depicted in Figure 3 (b). The 500-hr aging time duration was observed to significantly lower the hardness measures of HAZ by 20 pct. from the as-received condition. However, the hardness measures of the base metal are unaffected. However, hardness measures of the base metal decreases to 146 after thermal exposure of 800 up to 1200 hours.
Fig. 3 Hardness profile across the fusion zone of welded A516 steel
for as-received, hydrogen charged and aged samples.
TABLE III SUMMARY OF HARDNESS VALUES ACROSS THE WELD
Fatigue crack growth behavior of thermally aged and
hydrogen charged A516 welded steel is compared in Fig. 4 b. in terms of normalized crack lengths, a versus loading cycles, N. It is noted that HAZ-AR sample endured the longest life at 375900 cycles while HAZ-9 hour’s hydrogen charged specimen the shortest at 311400 cycles.
Fig. 4 a. Summary of Hardness values across the weld versus aging time
B. Fatigue Crack Growth Behavior
Fig. 4 b. Fatigue crack growth behavior of A516 steel
The threshold ∆K th ranges from 11.4 MPa√m for HAZ-800
hrs aged sample to the highest magnitude of 15.32 MPa√m for base metal plate. In this region, da/dN ranges from 10-5 to 10-2 mm/cycle while the crack-tip driving force, ∆K varies from 13 to 39 MPa√m. The magnitude of ∆K at the final fracture is about 39.51 MPa√m which approximate the fracture toughness of the A516 steel.
Fig. 5 Fatigue crack growth rates of A516 steel
Fig. 5 compares the FCGR behavior of A516 steel. The
variation of ∆K th with thermal aging and absorbed hydrogen investigated is listed (see table IV).
2nd International Conference on Mechanical, Electronics and Mechatronics Engineering (ICMEME'2013) June 17-18, 2013 London (UK)
122
TABLE IV
VARIATION OF THRESHOLD STRESS INTENSITY FACTOR RANGE, ∆KTH WITH THERMALLY AGED AND ABSORBED
Fractrographic features of the weld corresponding to the fatigue crack growth process at Paris region are shown in Figure 7. The rough morphology is a result of uneven crack front propagating through the multi-pass weld.
Fig. 6 Paris region of fatigue crack growth of A516 steel
The threshold ∆K th ranges from 11.4 MPa√m for HAZ-800
hrs aged sample to the highest magnitude of 15.32 MPa√m for base metal plate. In this region, da/dN ranges from 10-5 to 10-2 mm/cycle while the crack-tip driving force, ∆K varies from 13 to 39 MPa√m. The magnitude of ∆K at the final fracture is about 39.51 MPa√m which approximate the fracture toughness of the A516 steel.
IV. CONCLUSIONS
• The base metal displays an aggregate of well-defined ferrite-pearlite bands.
• The low crack growth rate region is characterized by crack growth bridging process.
• Fracture features during crack growth bridging process of the HAZ is primarily transgranular while during final fracture, the mode is intergranular.
Fig. 7 Fracture surface morphology of A516 steel at Paris region of fatigue crack growth. (2500 X)
As Received
800 Hrs Aged
9-Hrs Hydrogen
Secondary Cracks
2nd International Conference on Mechanical, Electronics and Mechatronics Engineering (ICMEME'2013) June 17-18, 2013 London (UK)
123
REFERENCES
[1] J. Spence, D.H. Nash, “Milestones in pressure vessel technology,” International Journal of Pressure Vessels and Piping 81 (2004)89–118
[2] S-H. Song, H.Zhuang, J.Wu, L.-Q.Weng, Z.-X.Yuan,T.-H.Xi, “Dependence of ductile-to-brittle transition temperature on phosphorous grain boundary segregation for a 2.25Cr1Mo steel”, Material Science and Engineering A 486 (2008) 433-438
[3] Tesman, A.B, “Materials of construction for process plants - I,” Chemical engineering, Reproduced by courtesy of Japan steel works (1973).
[4] N.V.Challenger, R.Phaal and S.J.Garwood. “Fracture mechanics assessment of industrial pressure vessel failures,” (1995).
[5] Zeljko Domazet, Lovre Krstulovic-Opara, Mladen Stupalo, “Fatigue cracks and failures in cement industry, shipbuilding and power plant facilites”, Engineering failure analysis 12 (2005) 819-833
[6] R.K.Pandey, “Analysis of cracking in LPG Horton spherical vessel”, Engineering failure analysis. (2004)
[7] R. Viswanathan, “Damage Mechanisms and Life Assesment of High-Temperature Components”, ASM International, (1989)
[8] Matt Johnson et al., “State of the Art Report on Welding and Inspection”, Federal Emergency Management Agency, FEMA-355B/September 2000.
[9] Easterling, K., “Introduction to the Physical Metallurgy of Welding, Butterworth- Heinemann”, 1992.
[10] Burdekin, F.M., and Suman, A., “Further Thoughts on the Relationship Between Toughness”, Workmanship and Design for Earthquake Resistance Structures, IIW DOC X- 1431-98, XV-998-98, and XVG-43-98, 1998.
[11] Harrison, P.L., and Hart, P.H.M., “Influences of Steel Composition and Welding Procedure on the HAZ Toughness of Thick Section Structural Steels,” Proc. Int. Conf. Metallurgy, Welding, and Qualification of Microalloyed (HSLA) Steel Weldments, AWS, pp. 626-658, American Welding Society, 1990.
[12] Johnson, M.Q., “Preliminary Evaluation of Heat Affected Zone Toughness in Structural Shapes Used in the Construction of Seismic Moment Frames”, SAC Joint Venture Report SAC/BD-00/13, 2000.
[13] R. A. Farrar, L. G. Taylor And E. M. Harrison Metals Technology 6,380 (1979).
[14] G. M. Evans, Welding Journal61, 125s (1982). [15] C, B. Dallam, S. Liu And L. Olson, Welding Journal 64, 14s
(1985). [16] O. Grong And D. K. Matlock, Int. Metalrev. 31, 27 (1986). [17] M. Es-Soun1, P. A. Beaven And G. M. Evans, Welding Journal 70,
805s (1991). [18] E. Surian, J, Troi~I'[, R. Herrer And L. A. De Weda, Welding
Journal 70, 133s (1991). [19] D. P. Fairchild, N. V. Bangaru, J. Y. Koo, P. L. Harrison And A.
Ozekcin, Welding Journal 70, 321s (1991). [20] C.T. Fujll, Metall. Trans. A 12A, 1099 (1981). [21] D. Tromans, S. Ramakrlshina And E. B. Hawsolt, Corrosion 42, 63
(19861. [22] H. I. Mchenrv, P. T. Purtscher And T. R. Shives, Corros. Sci. 27,
1041 (1987). [23] Y. Kobayahsl, K. Um , T. Hyodo And T. Taira, Corros. Sci. 27,
1117 (1987). [24] H. Pircher And G. Sussek, Corros. Sci. 27, 1183 (1987). [25] V. B. Lawson, C. Duncan And R. S. Treseder, Mater. Perf. 27, 10
(19881. [26] M. I. Onsoien, O. M. Akselsen, 0. Grong And P. E. Kvaale,
Welding Journal69, 45s (1990). [27] S. A. Gedeon And T. W. Eagar, Welding Journal 69, 213s (1990). [28] G, C. Savva, G. C. Weatherly And K. T. Aust, Corrosion 45,243
(1989). [29] T. M. Majid, S. A. R. Sultan, S. D. Mahdi And S. Jasim, Corrosion
46, 37 (1990). [30] API Standard 1104 – Twentieth Edition, "Welding of Pipelines and
Related Facilities," American Petroleum Institute, Washington, DC, November 2005.
[31] CSA Z662-07, “Oil and Gas Pipeline Systems,” CSA International, June 2007.
2nd International Conference on Mechanical, Electronics and Mechatronics Engineering (ICMEME'2013) June 17-18, 2013 London (UK)
![Fatigue performance evaluation of modified asphalt binder ...mahmoudameri.com/Articles/Fatigue performance evaluation.pdf · to represent the short term aging of asphalt binder [21,23].](https://static.documents.pub/doc/80x56/5eade3a5fac1863a1f405679/fatigue-performance-evaluation-of-modified-asphalt-binder-performance-evaluationpdf.jpg)
