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Testing for hydrogen—a-106 steel compatibility
Int. 1. Hydrogen Energy, Vol. 5, pp. 597--608 Pergamon Press Ltd. 1980, Printed in Great Britain IntcrnatfonM .Maodation for Hydrogen Energy
0360-3199/80/1101--0597 $02.00/0
TESTING F O R HYDROGEN--A-106 STEEL COMPATIBILITY* J. MURALI, T. A. ADLER, T. S. SUDARSrtAN,M. R. LOUTHAN,JR. and R. P. McNrrr Virginia Polytechnic Institute and State University, Blacksburg, VA 24061, U.S.A.
(Received .for publication 16 April 1980) Abstract--Hydrogen compatibility of A-106 pipeline steel was systematically investigated for various stress states, strain rates and stress ranges using different specimen geometries and loading configurations. Test samples were either exposed to hydrogen during testing or had been precharged by exposure to gaseous hydrogen at pressure to 14 MPa. Ductility reduction, changes in fracture mode, decreases in fracture energy and reductions in low cycle "ratcheting" (fatigue) life were observed although no delayed failures were seen. These adverse effects were associated with the severity of the stress state to which the material was subjected and with geometrical constraints such as machining marks, scratches and notches. INTRODUCTION
THE DEVASTATINGeffects of hydrogen on the mechanical properties of high strength steels have been documented by a number of researchers [1-5]. In particular, phenomena such as catastrophic delayed failure or loss of ductility are aspects of hydrogen "embrittlement" that most engineers consider when designing hydrogen systems. Because most hydrogen embrittlement has been observed in high strength steel, the natural recourse is to use low strength steels for hydrogen systems. This tendency, however, is thwarted by recent results which demonstrate changes in mechanical response of low-carbon, mild steels exposed to hydrogen [6-8]. These changes include surface cracking, alterations in macroscopic and microscopic fracture modes, change in fatigue life and loss of ductility. Such laboratory results are in apparent conflict with numerous "successful" applications of mild steels in the transmission and storage of hydrogen, e.g. gas storage bottles. Furthermore, published results of materials embrittled by hydrogen are often in disagreement. These contradictions may be reconciled by an examination of the stress state, stress level, strain rates and cycling conditions employed in service and by comparing service results with a wide range of laboratory tests. The stress state is important in two different aspects of the degradation process: (1) aiding the accumulation of hydrogen at the failure initiation site, and (2) promoting the crack growth process [9-11]. It is not clear that the stress state most effective in process (1) is similarly favorable for process (2). The accumulation of hydrogen at a possible failure site may be accomplished by: (a) strain-aided diffusion where a localized hydrostatic tensile stress enhances the transport and capacity of hydrogen [12], and/or (b) dislocation transport of hydrogen [9]. It is clear that the stress state required to achieve strain aided diffusion (triaxial stresses) is different from that required for dislocation transport where shear stresses are operative. Also strain aided diffusion can occur throughout the elastic-plastic range while dislocation transport would be primarily associated with the post yield behaviour. In material compatibility tests in the laboratory an investigator may be required to employ high strain rates which allow sufficient times for dislocation transport but not for strain aided diffusion. Conversely, low strain rate or static load tests may emphasize strain aided diffusion. Thus the results of any single laboratory test may not duplicate service conditions. The importance of stress level in the embrittling process is quite evident if one considers the fact that a threshold stress level exists in most cases for hydrogen affected cracking
[13]. The role of hydrogen in the "embrittling" processes, subsequent to actual accumulation at fracture initiation site(s) is, apparently, strongly dependent on the stress state [9-12, 14-16] with models invoking either plasticity (shear stress dependence) or decohesion (hydrostatic tensile stress dependence) and experimental evidence can be cited to support either rationale [17, 18]. *Work supported by DOE Contract No. E(40-1)5255. 597
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In view of the above it is crucial that qualification of a material for hydrogen service be established via a test program involving a variety of tests that model actual operating environments, particularly in regards to stress states, strain rates and stress range. The following describes the results of a program to assess hydrogen "embrittlement" in the ferritic-pearlitic pipe steel, A-106 grade B. M A T E R I A L S AND P R O C E D U R E S The material used in this investigation was ASTM type A-106 grade B pipe steel having the nominal composition of 0.25% C, 0.27-0.93% Mn, 0.10% Si. This composition is typical of the plain carbon steels often used for pipeline and pressure vessel applications. Tensile, notched tensile, compact tension and tubular samples were machined from a single heat of as-received, seamless pipe having the banded microstructure shown in Fig. 1. The
FIG. 1. Microstructure of A-106-B steel showing banded ferrite and regions of pearlite.
axis of the tensile specimen was parallel to the hot drawn direction in the pipe. Disk rupture tests were made on samples prepared by cold rolling and annealing sections machined from the plate stock. Hydrogen exposures and the tests in high pressure hydrogen were made in the equipment described in [18]. All tests were conducted at room temperature. RESULTS Tensile tests
Tensile deformation of smooth bar samples in hydrogen gas at pressures between 3 and 15 MPa was accompanied by yield point suppression, surface cracking, losses in fracture strength and a decrease in the strain-hardening exponent. These effects were strain rate dependent and increased
TESTING FOR HYDROGEN--A-106 STEEL COMPATIBILITY
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L
FIG. 2. Hydrogen-inducedsurfacecrackinginA-106-B steeltensilesampletestedin 15MPa hydrogen.
as the initial strain rate decreased. Surface finish and hydrogen gas purity were found to be important test variables. Rough surfaces increased the tendency for embrittlement while polished surfaces and minor additions of oxygen to the test gas improved hydrogen compatibility. Both of these effects are consistent with other studies [16] and help explain the large scatter in the literature estimates of apparent hydrogen compatibility. The hydrogen induced surface cracking in the smooth bar samples was often associated with the machining marks and other surface defects (Fig. 2). In an attempt to maximize the role of surface defects on the failure process a series of "ratcheted" tensile tests were made (Fig. 3). "Ratcheting" was begun after the sample had yielded and strain cycling to about 0.7% plastic strain per cycle was continued to failure. This quasi fatigue test was used in hope that the unloading would keep sharp any cracks which formed, thereby maximizing the effect of surface finish. Samples with 32 rms, 320 grit, 600 grit and electropolished surfaces were tested at initial strain rates of approx 10 -4 s -~ in hydrogen or oxygen gas. No effect of environment on the cycles to failure was observed
•7•YCYCLED TO
/ ~' /
/ / / / PLASTICSTRAIN 1234 PER CYCLE
/ DISPLACEMENT FIG. 3. "Ratcheting" in a typical test on A-106-B smooth bar sample.
aH = stress in hoop direction. at = stress in the axial direction.
0.4064
OH = aL
and
OH = 2aL
Cycled between
oH=or.
and
OH = 20L
Cycled between
oH - aL
0.4699
0.508
= 2oL
oH
0.4064
Fatigue tests
oH = 2oL
0.508
Static tests
= 2oL
0.4953 on
Stress state
Control sample
Test type
Minimum wall thickness (mm)
Hydrogen
Hydrogen
Hydrogen
Hydrogen
Hydrogen
Oxygen
Gas
7
4
8
7
11
11
Pressure (MPa)
590
525
560
710
730
Hoop
295
390
525
280
355
365
Longitudinal
Stress (MPa)
TABLE 1. Results of internal pressurization tests on A-106 steel cylinders
Held in biaxial stress state at 4 MPa for 572 h after 4030 cycles at the same loadings. Failed by overload.
Failed by tensile overloading after 700 cyles.
Burst by overpressure after being exposed to hydrogen gas at 4 MPa for 3400 h. Burst by overpressure after being e ~ d to hydrogen gas at 3.5 MPa for 650 h. Exposed to 7 MPa hydrogen for 7 h 20 rain and failed when the pressure was increased slightly ( - 0.2 MPa).
Burst by overpressure.
Remarks
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on samples with either the 600 grit or the electropolished surfaces. However, samples with surfaces having the 32 rms and the 320 grit finishes showed hydrogen induced reductions in cycle lifetime (Table 1). MetaUographic observations showed that fewer surface cracks were present in samples tested by ratcheting as compared to similar samples pulled to failure in equivalent hydrogen environments. In both the tensile and ratcheted samples cracks close to the failure were long and blunt while cracks remote from the failure were short and sharp. The blunt cracks were primarily in the necked regions indicating that blunting occurs as a result of necking and overload while the sharpness of the other cracks indicates that during uniform deformation the hydrogen induced cracks remained sharp and thus more susceptible to hydrogen induced crack growth. This conclusion was supported by ratchet tests of notched samples. The notch root radius of 0.025 cm concentrated the strain during testing and reduced the fatigue or ratchet life of these samples. Thus 20-50% fewer cycles were required to cause failures in hydrogen environments than in equivalent oxygen environments. The fracture mode in hydrogen in both the smooth bar and notched samples was quasi-cleavage while microvoid coalescence was observed in oxygen.
Tubular samples Cylindrical pressure vessel samples (Fig. 4a) machined from the tube stock used for the tensile samples have been exposed to high-pressure hydrogen gas for as long as 3000 h without delayed failure. The stress state in the wall of these vessels is controlled by the application of an axial load in combination with the internal pressure; a cylindrical stress state (OL = 1/2Oc) is obtained when the axial load is zero and a spherical stress state (aL = ac) is obtained by the application of the appropriate axial load. Samples have been exposed under "cylindrical" loading, "spherical" loading and under cyclic loading where the stress state was cycled from cylindrical to spherical. In all cases
TABLE2.
Cycles
to failure in ratchet test
Sample type
Surface finish
Number of cycles
Test environment
Smooth bar
320 grit 320 grit 320 grit 32 rms 600 grit 7 in. 7 in.
135 115 104 103 119 30 12
5.2 MPa 02 5.2 MPa H2 10.4 MPa H2 10.4 MPa H2 10.4 MPa H2 5.2 MPa 02 5.2 MPa H2
Notched bar
the maximum principal exposure stress exceeded the yield strength of the material. No delayed failures have been observed in any of these pressure vessels. Furthermore in all but one case the failures which occurred by overpressurization were accompanied by large scaleplastic deformation. However, in some cases the exposure apparently reduced the burst stress of the samples (Table 2). A cylinder burst by overpressurization with hydrogen after long-time hydrogen exposure in a cylindrical stress state is shown in Fig. 4b. Metallographic examination of the failed samples showed that the macroscopic and microscopic appearances of the failure surfaces of the oxygen and hydrogen exposed samples were noticeably different. Surface cracking, losses in microductility and fracture mode changes were observed in the hydrogen exposed samples. However, the surface cracks were typically very shallow (Fig. 5a) and accompanied by regions of brittle fracture (Fig. 5b) which were only several hundredths of a centimeter deep. The primary fracture mode on all samples was microvoid coalescence. The tubes used in this study contained no intentional flaws and, in fact, considerable care was
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41a FIO. 4(a) Thin walled pressure vessel of A-106-B steel subjected to internal pressure and axial loading. (b) Pressure vessel of A-106-B steel burst by internal hydrogen after soaking in 600 psi (4 MPa) hydrogen for 3400 h. Burst pressure--1650 psi (11 MPa). Wall thickness--O.02 in. (0.518 ram).
TESTING FOR HYDROGEN--A-106 STEEL COMPATIBILITY
Flo. 5(a). SEM fractograph of sample in Fig. 4(b) showing surface cracking on inner, hydrogen-exposed surface. (b) SEM fractograph of the fracture surface of sample in Fig. 4(b) showing quasi-cleavage on the side exposed to hydrogen.
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taken to provide smooth flaw-free surfaces. The combination of the lack of "delayed" failure in the tubular samples and the large effects of hydrogen on the tensile properties of A-106 steel show that tensile test data alone are poor predictors of long-time hydrogen compatibility. However, the results of the ratcheted tensile tests indicate that large strain amplitude fatigue may be a problem in hydrogen environments. Large strains during nominal loadings may also occur at the tip of major flaws and thus provide the necessary conditions for in service hydrogen embrittlement. The one tubular sample which developed a leak without being intentionally overpressurized provides some support for this conclusion. This tubular pressure vessel was filled with hydrogen gas to 7 MPa, or approx 70% of its rupture pressure in hydrogen. The sample was then loaded axially until the vessel walls were in a "spherical" state of stress. The elastic and plastic strains which accompanied axial loading increased the internal volume of the vessel causing the internal pressure to drop and the stress state to become more cylindrical. The sample remained under these loads for approx 7 h before the pressure was increased to obtain the spherical stress state. This slight increase in pressure to achieve spherical state (0.2 MPa) caused failure by cracking through the sample wall. The crack was very short and the failure simply resulted in a leak in the sample. This type of failure by the development of through thickness cracks without unstable crack propagation is in contrast to the catastrophic failures produced by overpressurization. Furthermore, fractographic examination of the failure showed that quasi-cleavage was the dominate fracture mode throughout the crack surface. This fracture mode is identical to that which causes the surface cracking in tensile samples deformed in hydrogen environments and indicates that hydrogeninduced stable growth lead to failure in this case. Disk rupture tests under cyclic conditions provided further support for this conclusion.
FIG. 6. Preflawed disk samples. The pencil points to the sample failed in hydrogen by leaking. The other deformed sample was tested in oxygen and failed by bursting catastrophically.
D i s k rupture tests
The susceptibility of A-106 steel to hydrogen-induced failure under cyclic loading conditions was further evaluated by stressing samples under biaxial loading to 90% of the rupture pressure. Static loading to this level caused large-scale yielding in the hydrogen gas but did not cause rupture. Cyclic hydrogen pressurization to 90% of the rupture pressure caused large reductions
TESTING FOR HYDROGEN--A-106 STEEL COMPATIBILITY
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in the life of the disks as failure occurred in less than 30 cycles (one cycle being pressurization to 90% of the rupture pressure and depressurization to one atm hydrogen) and similar cyclic exposures in oxygen gas did not cause failures after 100 cycles. These failures were similar to the pressure vessel failure, in that the crack growth in hydrogen was apparently stable as the sample leaked rather than burst. These data thus indicate that the lack of crack growth under sustained loading does not predict the suitability of a material for service under cyclic loading and that cyclic loadings in hydrogen can cause rapid losses in mechanical integrity. Confirmation of such significant hydrogen effects on crack growth were obtained in studies with compact tension samples. Overpressurization to burst in H2 required only approx 85% the pressure as for burst samples in 02. Further the hydrogen test samples leaked while most of those in O2 burst violently (Fig. 6).
Compact tension samples Fatigue precracked compact tension samples were tested in either 14 MPa hydrogen or oxygen at room temperature. The hydrogen environment promoted crack growth [Figs. 7(a) and 7(b)], caused the fracture mode to change from microvoid coalescence to cleavage [Figs. 7(c) and 7(d)] and lowered the apparent Jc for unstable crack growth (Table 3). The J methodology used was TABLE3. Jtc for unstable crack growth in compact tension sample of A-106 grade B steel Environment
J (maximumload)
13.8 MPa (2000psi) 02
MJ in. •lb 413~ (2360in-q-ff~.)
13.8MPa (2000psi) H2
119~_~2 (680 ~in"lb)
similar to the one described in [20]. The samples were tested to the same crack-opening displacement by constant crosshead motion. The crosshead displacement at maximum load was less in hydrogen than in oxygen and the crack length at the final crack opening was greater in hydrogen than in oxygen. The crack fronts in the H2 sample were relatively straight whereas those in 02 were quite curved with maximum growth in the center. These results confirm the previous conclusion that hydrogen environments promote crack growth. The observation that large amounts of stable crack growth accompanied the deformation in both hydrogen and oxygen supports the disk rupture and pressure vessel observations that hydrogen effects on "stable" crack growth under cyclic or increasing loads are the major cause for concern in the use of A- 106 steels in hydrogen environments. CONCLUSIONS The compatibility tests described in this paper demonstrate that environmental hydrogen may significantly lower the mechanical properties of A-106 steel. However no delayed failures were produced and adverse effects were noted only when the yield strength was exceeded. Thus, it is likely that shear stress plays an important role in the environmental degradation process. The following points should be emphasized: (1) the large-scale surface cracking observed in uniaxial tests of A-106 steel in hydrogen gas was macroscopically absent when biaxial tests were conducted using thin walled tubular pressure vessels; (2) hydrogen exposures reduced the low-cycle fatigue life of A-106 steel in both biaxial and uniaxial tests; (3) hydrogen exposures caused changes in fracture mode and lowered the toughness (Jc) of A106 steel. These results, particularly the lack of delayed failure coupled with the effects of stress state demonstrate the importance of a wide range of testing for materials compatibility in hydrogen.
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Fi6. 7(a). Fracture surface of the compact tension sample tested in 13,8 MPa oxygen. The center, darker area was formed during the test in 02 and the crack grew from the bottom to top. (b) Fracture surface of the compact tension sample tested in 13.8 MPa hydrogen. The center area was formed during the test in hydrogen. The samples in Figs. 7(a) and 7(b) were tested to the sample displacement. (c) SEM fractograph of compact tension sample tested in 13.8 MPa 02. (d) SEM fractograph of a compact tension sampled tested in 13.8 MPa H2.
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REFERENCES 1. I. M. BERNSTEIN,Mater. Sci. Engng., 6, 1 (1970). 2. A. R. TROIAr~O, Trans. Am. Soc. Metals 52, 54 (1960). 3. J. P. b-~DELLE,ASTM STP-543, ASTM, Phil., PA., p. 221 (1974). 4. M. R. LOUTHA~, JR., Hydrogen in Metals, ASM, Metals Park, Ohio, p. 53 (1974). 5. A. W. THOMPSON,Environmental Degradation of Engineering Materials, VPI, Blacksburg, VA, p. 3 (1977). 6. T. S. StrDARSHAr~,M. R. LoLrrnAN and R. P. McNrrr, Scr. metal. 12, 799 (1978). 7. R. GA~ER and I. M. BEI~STEIN, Environmental Degradation of Engineering Materials, VPI, Blacksburg, VA, p. 963 (1977). 8. T. D. LEE, T. GOLDENBERGand J. P. HIRTER, Advances in Research on the Strength and Fracture of Materials, Pergamon Press, p. 243 (1977). 9. M. R. LOUTHAN,JR., G. R. CASKEY,JR., J. A. DONOVANand D. E. RAWL, JR., Mater. Sci. Engng. 10, 357 (1972). 10. J. K. TIEN, ANTHONYW. THoMPsoN, I. M. BERNSTEINand REBECCAJ. RICHARDS,Metal Trans A 7A, 821 (1976). 11. C. D. BEACHEM,Metal Trans. 3, 289 (1972). 12. R. A. OmAr~, Ber. Bunsenges, Phys. Chem. 76, 848 (1973). 13. C. F. BARTHand E. A. S'mXt~ERWALD,Metal Trans. 1, 3451 (1970). 14. P. DOIG and G. T. J o s s , Metal Trans. 8A, 1993 (1977). 15. N. J. PETCrI, Phil. Mag. 1,331 (1956). 16. H. P. VAN LEEUWEN, Effect of Hydrogen on Behavior of Materials, AIME, New York, p. 480 (1976). 17. M. R. LOUTHAN,JR. and R. P. McNrrr, Effect of Hydrogen on Behavior of Materials, AIME, New York, p. 496 (1976). 18. A. W. THOMPSONand I. M. BERNSTEIN,Rockwell International Report SC-PP-75-63, Thousand Oaks, CA 91360 (1976). 19. M. R. LOUTHAN,JR., R. P. McNrrT and N. SRIDHAR,Environmental Degradation of Engineering Materials, VPI, Blacksburg, VA, p. 745 (1977). 20. J. R. RACE,P. C. PARIS and J. G. MERELE, ASTM STP-536, pp. 231-245 (1973).
0360-3199/80/1101--0597 $02.00/0
TESTING F O R HYDROGEN--A-106 STEEL COMPATIBILITY* J. MURALI, T. A. ADLER, T. S. SUDARSrtAN,M. R. LOUTHAN,JR. and R. P. McNrrr Virginia Polytechnic Institute and State University, Blacksburg, VA 24061, U.S.A.
(Received .for publication 16 April 1980) Abstract--Hydrogen compatibility of A-106 pipeline steel was systematically investigated for various stress states, strain rates and stress ranges using different specimen geometries and loading configurations. Test samples were either exposed to hydrogen during testing or had been precharged by exposure to gaseous hydrogen at pressure to 14 MPa. Ductility reduction, changes in fracture mode, decreases in fracture energy and reductions in low cycle "ratcheting" (fatigue) life were observed although no delayed failures were seen. These adverse effects were associated with the severity of the stress state to which the material was subjected and with geometrical constraints such as machining marks, scratches and notches. INTRODUCTION
THE DEVASTATINGeffects of hydrogen on the mechanical properties of high strength steels have been documented by a number of researchers [1-5]. In particular, phenomena such as catastrophic delayed failure or loss of ductility are aspects of hydrogen "embrittlement" that most engineers consider when designing hydrogen systems. Because most hydrogen embrittlement has been observed in high strength steel, the natural recourse is to use low strength steels for hydrogen systems. This tendency, however, is thwarted by recent results which demonstrate changes in mechanical response of low-carbon, mild steels exposed to hydrogen [6-8]. These changes include surface cracking, alterations in macroscopic and microscopic fracture modes, change in fatigue life and loss of ductility. Such laboratory results are in apparent conflict with numerous "successful" applications of mild steels in the transmission and storage of hydrogen, e.g. gas storage bottles. Furthermore, published results of materials embrittled by hydrogen are often in disagreement. These contradictions may be reconciled by an examination of the stress state, stress level, strain rates and cycling conditions employed in service and by comparing service results with a wide range of laboratory tests. The stress state is important in two different aspects of the degradation process: (1) aiding the accumulation of hydrogen at the failure initiation site, and (2) promoting the crack growth process [9-11]. It is not clear that the stress state most effective in process (1) is similarly favorable for process (2). The accumulation of hydrogen at a possible failure site may be accomplished by: (a) strain-aided diffusion where a localized hydrostatic tensile stress enhances the transport and capacity of hydrogen [12], and/or (b) dislocation transport of hydrogen [9]. It is clear that the stress state required to achieve strain aided diffusion (triaxial stresses) is different from that required for dislocation transport where shear stresses are operative. Also strain aided diffusion can occur throughout the elastic-plastic range while dislocation transport would be primarily associated with the post yield behaviour. In material compatibility tests in the laboratory an investigator may be required to employ high strain rates which allow sufficient times for dislocation transport but not for strain aided diffusion. Conversely, low strain rate or static load tests may emphasize strain aided diffusion. Thus the results of any single laboratory test may not duplicate service conditions. The importance of stress level in the embrittling process is quite evident if one considers the fact that a threshold stress level exists in most cases for hydrogen affected cracking
[13]. The role of hydrogen in the "embrittling" processes, subsequent to actual accumulation at fracture initiation site(s) is, apparently, strongly dependent on the stress state [9-12, 14-16] with models invoking either plasticity (shear stress dependence) or decohesion (hydrostatic tensile stress dependence) and experimental evidence can be cited to support either rationale [17, 18]. *Work supported by DOE Contract No. E(40-1)5255. 597
598
J. MURALI et al.
In view of the above it is crucial that qualification of a material for hydrogen service be established via a test program involving a variety of tests that model actual operating environments, particularly in regards to stress states, strain rates and stress range. The following describes the results of a program to assess hydrogen "embrittlement" in the ferritic-pearlitic pipe steel, A-106 grade B. M A T E R I A L S AND P R O C E D U R E S The material used in this investigation was ASTM type A-106 grade B pipe steel having the nominal composition of 0.25% C, 0.27-0.93% Mn, 0.10% Si. This composition is typical of the plain carbon steels often used for pipeline and pressure vessel applications. Tensile, notched tensile, compact tension and tubular samples were machined from a single heat of as-received, seamless pipe having the banded microstructure shown in Fig. 1. The
FIG. 1. Microstructure of A-106-B steel showing banded ferrite and regions of pearlite.
axis of the tensile specimen was parallel to the hot drawn direction in the pipe. Disk rupture tests were made on samples prepared by cold rolling and annealing sections machined from the plate stock. Hydrogen exposures and the tests in high pressure hydrogen were made in the equipment described in [18]. All tests were conducted at room temperature. RESULTS Tensile tests
Tensile deformation of smooth bar samples in hydrogen gas at pressures between 3 and 15 MPa was accompanied by yield point suppression, surface cracking, losses in fracture strength and a decrease in the strain-hardening exponent. These effects were strain rate dependent and increased
TESTING FOR HYDROGEN--A-106 STEEL COMPATIBILITY
599
L
FIG. 2. Hydrogen-inducedsurfacecrackinginA-106-B steeltensilesampletestedin 15MPa hydrogen.
as the initial strain rate decreased. Surface finish and hydrogen gas purity were found to be important test variables. Rough surfaces increased the tendency for embrittlement while polished surfaces and minor additions of oxygen to the test gas improved hydrogen compatibility. Both of these effects are consistent with other studies [16] and help explain the large scatter in the literature estimates of apparent hydrogen compatibility. The hydrogen induced surface cracking in the smooth bar samples was often associated with the machining marks and other surface defects (Fig. 2). In an attempt to maximize the role of surface defects on the failure process a series of "ratcheted" tensile tests were made (Fig. 3). "Ratcheting" was begun after the sample had yielded and strain cycling to about 0.7% plastic strain per cycle was continued to failure. This quasi fatigue test was used in hope that the unloading would keep sharp any cracks which formed, thereby maximizing the effect of surface finish. Samples with 32 rms, 320 grit, 600 grit and electropolished surfaces were tested at initial strain rates of approx 10 -4 s -~ in hydrogen or oxygen gas. No effect of environment on the cycles to failure was observed
•7•YCYCLED TO
/ ~' /
/ / / / PLASTICSTRAIN 1234 PER CYCLE
/ DISPLACEMENT FIG. 3. "Ratcheting" in a typical test on A-106-B smooth bar sample.
aH = stress in hoop direction. at = stress in the axial direction.
0.4064
OH = aL
and
OH = 2aL
Cycled between
oH=or.
and
OH = 20L
Cycled between
oH - aL
0.4699
0.508
= 2oL
oH
0.4064
Fatigue tests
oH = 2oL
0.508
Static tests
= 2oL
0.4953 on
Stress state
Control sample
Test type
Minimum wall thickness (mm)
Hydrogen
Hydrogen
Hydrogen
Hydrogen
Hydrogen
Oxygen
Gas
7
4
8
7
11
11
Pressure (MPa)
590
525
560
710
730
Hoop
295
390
525
280
355
365
Longitudinal
Stress (MPa)
TABLE 1. Results of internal pressurization tests on A-106 steel cylinders
Held in biaxial stress state at 4 MPa for 572 h after 4030 cycles at the same loadings. Failed by overload.
Failed by tensile overloading after 700 cyles.
Burst by overpressure after being exposed to hydrogen gas at 4 MPa for 3400 h. Burst by overpressure after being e ~ d to hydrogen gas at 3.5 MPa for 650 h. Exposed to 7 MPa hydrogen for 7 h 20 rain and failed when the pressure was increased slightly ( - 0.2 MPa).
Burst by overpressure.
Remarks
TESTING FOR I-IYDROGEN--A-106 STEEL COMPATIBILITY
601
on samples with either the 600 grit or the electropolished surfaces. However, samples with surfaces having the 32 rms and the 320 grit finishes showed hydrogen induced reductions in cycle lifetime (Table 1). MetaUographic observations showed that fewer surface cracks were present in samples tested by ratcheting as compared to similar samples pulled to failure in equivalent hydrogen environments. In both the tensile and ratcheted samples cracks close to the failure were long and blunt while cracks remote from the failure were short and sharp. The blunt cracks were primarily in the necked regions indicating that blunting occurs as a result of necking and overload while the sharpness of the other cracks indicates that during uniform deformation the hydrogen induced cracks remained sharp and thus more susceptible to hydrogen induced crack growth. This conclusion was supported by ratchet tests of notched samples. The notch root radius of 0.025 cm concentrated the strain during testing and reduced the fatigue or ratchet life of these samples. Thus 20-50% fewer cycles were required to cause failures in hydrogen environments than in equivalent oxygen environments. The fracture mode in hydrogen in both the smooth bar and notched samples was quasi-cleavage while microvoid coalescence was observed in oxygen.
Tubular samples Cylindrical pressure vessel samples (Fig. 4a) machined from the tube stock used for the tensile samples have been exposed to high-pressure hydrogen gas for as long as 3000 h without delayed failure. The stress state in the wall of these vessels is controlled by the application of an axial load in combination with the internal pressure; a cylindrical stress state (OL = 1/2Oc) is obtained when the axial load is zero and a spherical stress state (aL = ac) is obtained by the application of the appropriate axial load. Samples have been exposed under "cylindrical" loading, "spherical" loading and under cyclic loading where the stress state was cycled from cylindrical to spherical. In all cases
TABLE2.
Cycles
to failure in ratchet test
Sample type
Surface finish
Number of cycles
Test environment
Smooth bar
320 grit 320 grit 320 grit 32 rms 600 grit 7 in. 7 in.
135 115 104 103 119 30 12
5.2 MPa 02 5.2 MPa H2 10.4 MPa H2 10.4 MPa H2 10.4 MPa H2 5.2 MPa 02 5.2 MPa H2
Notched bar
the maximum principal exposure stress exceeded the yield strength of the material. No delayed failures have been observed in any of these pressure vessels. Furthermore in all but one case the failures which occurred by overpressurization were accompanied by large scaleplastic deformation. However, in some cases the exposure apparently reduced the burst stress of the samples (Table 2). A cylinder burst by overpressurization with hydrogen after long-time hydrogen exposure in a cylindrical stress state is shown in Fig. 4b. Metallographic examination of the failed samples showed that the macroscopic and microscopic appearances of the failure surfaces of the oxygen and hydrogen exposed samples were noticeably different. Surface cracking, losses in microductility and fracture mode changes were observed in the hydrogen exposed samples. However, the surface cracks were typically very shallow (Fig. 5a) and accompanied by regions of brittle fracture (Fig. 5b) which were only several hundredths of a centimeter deep. The primary fracture mode on all samples was microvoid coalescence. The tubes used in this study contained no intentional flaws and, in fact, considerable care was
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41a FIO. 4(a) Thin walled pressure vessel of A-106-B steel subjected to internal pressure and axial loading. (b) Pressure vessel of A-106-B steel burst by internal hydrogen after soaking in 600 psi (4 MPa) hydrogen for 3400 h. Burst pressure--1650 psi (11 MPa). Wall thickness--O.02 in. (0.518 ram).
TESTING FOR HYDROGEN--A-106 STEEL COMPATIBILITY
Flo. 5(a). SEM fractograph of sample in Fig. 4(b) showing surface cracking on inner, hydrogen-exposed surface. (b) SEM fractograph of the fracture surface of sample in Fig. 4(b) showing quasi-cleavage on the side exposed to hydrogen.
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taken to provide smooth flaw-free surfaces. The combination of the lack of "delayed" failure in the tubular samples and the large effects of hydrogen on the tensile properties of A-106 steel show that tensile test data alone are poor predictors of long-time hydrogen compatibility. However, the results of the ratcheted tensile tests indicate that large strain amplitude fatigue may be a problem in hydrogen environments. Large strains during nominal loadings may also occur at the tip of major flaws and thus provide the necessary conditions for in service hydrogen embrittlement. The one tubular sample which developed a leak without being intentionally overpressurized provides some support for this conclusion. This tubular pressure vessel was filled with hydrogen gas to 7 MPa, or approx 70% of its rupture pressure in hydrogen. The sample was then loaded axially until the vessel walls were in a "spherical" state of stress. The elastic and plastic strains which accompanied axial loading increased the internal volume of the vessel causing the internal pressure to drop and the stress state to become more cylindrical. The sample remained under these loads for approx 7 h before the pressure was increased to obtain the spherical stress state. This slight increase in pressure to achieve spherical state (0.2 MPa) caused failure by cracking through the sample wall. The crack was very short and the failure simply resulted in a leak in the sample. This type of failure by the development of through thickness cracks without unstable crack propagation is in contrast to the catastrophic failures produced by overpressurization. Furthermore, fractographic examination of the failure showed that quasi-cleavage was the dominate fracture mode throughout the crack surface. This fracture mode is identical to that which causes the surface cracking in tensile samples deformed in hydrogen environments and indicates that hydrogeninduced stable growth lead to failure in this case. Disk rupture tests under cyclic conditions provided further support for this conclusion.
FIG. 6. Preflawed disk samples. The pencil points to the sample failed in hydrogen by leaking. The other deformed sample was tested in oxygen and failed by bursting catastrophically.
D i s k rupture tests
The susceptibility of A-106 steel to hydrogen-induced failure under cyclic loading conditions was further evaluated by stressing samples under biaxial loading to 90% of the rupture pressure. Static loading to this level caused large-scale yielding in the hydrogen gas but did not cause rupture. Cyclic hydrogen pressurization to 90% of the rupture pressure caused large reductions
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in the life of the disks as failure occurred in less than 30 cycles (one cycle being pressurization to 90% of the rupture pressure and depressurization to one atm hydrogen) and similar cyclic exposures in oxygen gas did not cause failures after 100 cycles. These failures were similar to the pressure vessel failure, in that the crack growth in hydrogen was apparently stable as the sample leaked rather than burst. These data thus indicate that the lack of crack growth under sustained loading does not predict the suitability of a material for service under cyclic loading and that cyclic loadings in hydrogen can cause rapid losses in mechanical integrity. Confirmation of such significant hydrogen effects on crack growth were obtained in studies with compact tension samples. Overpressurization to burst in H2 required only approx 85% the pressure as for burst samples in 02. Further the hydrogen test samples leaked while most of those in O2 burst violently (Fig. 6).
Compact tension samples Fatigue precracked compact tension samples were tested in either 14 MPa hydrogen or oxygen at room temperature. The hydrogen environment promoted crack growth [Figs. 7(a) and 7(b)], caused the fracture mode to change from microvoid coalescence to cleavage [Figs. 7(c) and 7(d)] and lowered the apparent Jc for unstable crack growth (Table 3). The J methodology used was TABLE3. Jtc for unstable crack growth in compact tension sample of A-106 grade B steel Environment
J (maximumload)
13.8 MPa (2000psi) 02
MJ in. •lb 413~ (2360in-q-ff~.)
13.8MPa (2000psi) H2
119~_~2 (680 ~in"lb)
similar to the one described in [20]. The samples were tested to the same crack-opening displacement by constant crosshead motion. The crosshead displacement at maximum load was less in hydrogen than in oxygen and the crack length at the final crack opening was greater in hydrogen than in oxygen. The crack fronts in the H2 sample were relatively straight whereas those in 02 were quite curved with maximum growth in the center. These results confirm the previous conclusion that hydrogen environments promote crack growth. The observation that large amounts of stable crack growth accompanied the deformation in both hydrogen and oxygen supports the disk rupture and pressure vessel observations that hydrogen effects on "stable" crack growth under cyclic or increasing loads are the major cause for concern in the use of A- 106 steels in hydrogen environments. CONCLUSIONS The compatibility tests described in this paper demonstrate that environmental hydrogen may significantly lower the mechanical properties of A-106 steel. However no delayed failures were produced and adverse effects were noted only when the yield strength was exceeded. Thus, it is likely that shear stress plays an important role in the environmental degradation process. The following points should be emphasized: (1) the large-scale surface cracking observed in uniaxial tests of A-106 steel in hydrogen gas was macroscopically absent when biaxial tests were conducted using thin walled tubular pressure vessels; (2) hydrogen exposures reduced the low-cycle fatigue life of A-106 steel in both biaxial and uniaxial tests; (3) hydrogen exposures caused changes in fracture mode and lowered the toughness (Jc) of A106 steel. These results, particularly the lack of delayed failure coupled with the effects of stress state demonstrate the importance of a wide range of testing for materials compatibility in hydrogen.
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Fi6. 7(a). Fracture surface of the compact tension sample tested in 13,8 MPa oxygen. The center, darker area was formed during the test in 02 and the crack grew from the bottom to top. (b) Fracture surface of the compact tension sample tested in 13.8 MPa hydrogen. The center area was formed during the test in hydrogen. The samples in Figs. 7(a) and 7(b) were tested to the sample displacement. (c) SEM fractograph of compact tension sample tested in 13.8 MPa 02. (d) SEM fractograph of a compact tension sampled tested in 13.8 MPa H2.
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