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Hydrogen compatibility of HT-9 martensitic stainless steel
877
HYDROGEN COMPATIBILITYOF HT-9 MARTENSITIC STAINLESS STEEL* J. PI.Hyzak and R. E. Stoltz Sandia National Laboratories Livermore, California 94550 The hydrogen compatibility of the alloy HT-9 is addressed. The results of tensile tests performed on specimens both in a gaseous hydrogen environment and after cathodic charging are described. The findings show that there is only a small effect of hydrogen (at the levels tested) on the ductility of quenched-and-temperedXT-9, However, larger reductions in ductility are observed for the as-quenched specimens. It is also shown that increasing the hydrogen concentrationand the strength level generally increases the likelihood of surface crack initiationwhich is an indication of a decrease in damage tolerance. 1.
INTRODUCTION
The proposed use of a ferritic/martensiticsteel for first wall and blanket structures raises a number of materials compatibilityquestions, including possible hydrogen embrittlement. Hydrogen can be produced in the material from bath direct injection of deuterium and tritium as well as from (n,p) reactions. The presence of interstitialhydrogen in levels as low as l-2 weight ppm is known to severely affect the fracture behavior of martensitic steels[l,21 Embrittlement'increases with yield strength 'and generally decreases with temperatureabove ambient. In this work, HT-9 tensile specimens exposed to both external (gaseous environment)and internal (cathodicallycharged) hydrogen have been tested. The effects of temperature,hydrogen pressure and alloy strength level on the tensile embrittlementof HT-9 are reported herein. 2.
EFFECT OF EXTERNAL HYDROGEN EXPOSURES ON TENSILE PROPERTIES OF HT-9
2.2 Test Results The results of all tensile tests are presented in Table I; the values are the average of two tests. In no instance did the presence of hydrogen significantly affect the yield or ultimate strength of the alloy. The effect of the hydrogen was manifest in the ductility as measured by reduction-in-area(RA) at fracture. At 298K, 0.10 MPa hydrogen reduced the tensile ductility from 68 to 58%. Although this loss in ductility is small, it does suggest that a hydrogen-inducedchange in fracture behavior has occurred. This effect was verified by comparison of the fracture surfaces of samples tested in air and in 0.10 MPa hydrogen, shown in Figures 1 and 2. The fracture process in air is entirely one of ductile void initiation, growth and coalesence. The overall fracture surface is typical of cup-cone type fracture, Figure 1, including a flat central region with radial cracks extending to an outer shear lip. All the air-fracture surfaces exhibit a ductile void fracture morphology. Comparison of the void size with the microstructure indicates that void initiation occurs at the tempered carbides,
2.1 Experimental Details To determine the effect of external hydrogen on HT-9, previously uncharged tensile specimens were tested to failure in a gaseous hydrogen environment. All tests were performed on samples taken from a pipe section of HT-9 with composition: Chemistry of HT-9, Wt.% C Si Mn Cr Ni MOW ------_--.22 .38 .52 11.3 .50 .85 .50
V .27
P S .019 ,006
Prior to testing, the specimens were austenitized at 1313K/30 min. and air cooled, followed by tempering at 1023K/60 min., with a final hardness of R,Zl. The resulting microstructure was fully martensitic, with a distribution of tempered carbides. In the hydrogen experiments samples were equilibrated at pressure and temperature for 30 min. before testing. *This work supported by U.S. Department of Energy, DOE, under contract number DE-AC04-76DP 00789.
0022-3115/81/0000-0000/%02.75 0 1981North-Holland
Figure 1. Cup-cone fracture of HT-9 tested in air at 298K.
878
J M. Hyzak,
R.E. Stoltz
/ Hydrogen
compatibility
of HT-9
martensitic
SS
air and 0.10 MPa hydrogen. No hydrogen-induced interlath fracture was observed; instead, ductile dimple fracture occurred across the 473673K temperature range of test.
Figure
2.
HT-9 tested in 0.10 MPa hydrogen at 298K with A and B indicating regions of hydrogen induced fracture.
The large radial fissures are regions where the ductile cracks join, and the side surfaces of these fissures exhibit elongated dimples similar In hydrogen, to those in the shear lip region. Figure 2, the overall features of the fracture surfaces are similar to those in air. However, there are regions of surface crack initiation, areas A and B, which exhibit an alternate fracture mode. The fracture is dominated by shallow The secondary cracks and a "woody" appearance. size of the cracks suggest that fracture occurs at or between the blocky martensite laths. At elevated temperatures, hydrogen effects are suppressed and the ductilities are comparable in
Mechanical
In order to study the effect of hydrogen pressure at ambient temperature, a series of tests was performed in the range from 6.9-34.5 MPa hydrogen. As indicated in Table I, a progressive but slight decrease in ductility occurs up to 20.7 MPa, with a sharp decrease manifest at 34.5 MPa hydrogen pressure. In the range 6.920.7 MPa hydrogen, the fracture behavior is similar to that at 0.10 MPa. Small regions of martensite interlath fracture are observed at the outer part of the fracture surface. As the test pressure is increased, the area percentage of this surface crack initiation mode increases. At 34.5 MPa, Figure 3, the hydrogen affected regions are much larger (A and B), comprising approximately 50% of the fracture area. The sample exhibits two fracture regions perpendicular to the tensile axis connected by a ductile shear region. Little overall necking of the sample was observed. When seen at higher magnifications, the flat regions show martensite interlath fracture with secondary cracks similar to that in 0.10 MPa hydrogen. In some regions of the sample, the martensite was oriented nearly perpendicular to the tensile axis, so that the interlath fracture exhibits a stepped appearance. 3.
EFFECT OF INTERNAL HYDROGEN PROPERTIES OF HT-9
ON THE TENSILE
To better
represent the hydrogen distribution in a first wall structure, specimens were internal-
Table I Properties of HT-9 in Air and Hydrogen Ultimate Strength MPa (ksi)
Ductility (Reduction %
Temperature
Atmosphere
Yield Strength MPa (ksi)
25OC
Air
592 (85.8)
799 (115.8)
67
25OC
0.10 MPa H2
565 (81.9)
837 (121.3)
58
2oooc
Air
533 (77.3)
719 (104.3)
64
2oooc
0.10 MPa H2
545 (79.0)
733 (106.2)
64
3oooc
Air
513 (74.4)
695 (100.7)
65
3oooc
0.10 MPa H2
524 (76.0)
689 (99.9)
67
4oo"c
Air
465 (67.4)
628 (91.0)
62
4oooc
0.10 MPa H2
496 (71.9)
637 (92.3)
63
250C
6.9 MPa Hp.
715 (103.6)
800 (116.0)
65
25'C
13.8 MPa H2
707 (102.5)
788 (114.2)
56
25OC
20.7 MPa H2
717 (104.2)
792 (114.8)
53
25'C
34.5 MPa H2
724 (105.1)
793 (115.0)
26
in Area)
819
J.M. Hyzak, R.E. Stoltz 1 Hydrogen compatibility of HT-9 martensitic SS
ly charged with hydrogen prior to testing. Hydrogen was introduced by cathodic charging, and specimens were immediately copper-plated to retard outgassing. The charging solution was 4 percent sulfuric acid containing 5 mg of sodium The current densities employed were arsenate. O.OO~A/C~* (O.O*A/in*) and 0.006 A/cm* (O.O4A/ in*) for 8 and 150 minutes (Table II). Specimens were aged at room temperature for 24 hours before testing in order to evenly distribute the hydrogen. Preliminary data indicate that even the least severe charging condition (.003A/ cm*, 8 minutes) would result in a hydrogen concentration greater than 5 wt. ppm (275 appm). Material was tested in both the as-quenched and quenched-and-tempered conditions. Since HT-9 will be subject to irradiation hardening, the as-quenched microstructurewas tested as a Preliminary means of assessing the effect of hardening on hydrogen compatibility. All specimens were austenitized as above and air cooled; half the specimens were tested as is and the remainder tempered as above. 3.1 Experimental
Figure 3.
Results
surrounded by a uniform shear lip. The fracture mode in the central section was dimpled rupture. Examination of metallographic cross-sections of the fracture surfaces revealed that fracture occurred in the tensile (rolling) direction along martensite lath boundaries as the result of nucleation and growth of voids at tempered carbides and stringers. The as-quenched microstructure had an uncharged yield strength substantially higher than that of the quench and tempered condition, 980 MPa versus 650 MPa. There was also a much larger effect of hydrogen charging on the tensile properties in the
The results of all the tensile tests are given in Table II. For the quenched-and-tempered microstructure, there was no significant effect of internal hydrogen on the ductility. However, there did appear to be a small but consistent increase in both yield and ultimate strength as a result of hydrogen charging. The fracture process for both the charged and uncharged specimens was the same, classical cup-cone failure. The fracture surfaces were rough with many large secondary cracks in the interior
Mechanical
Properties
Cond.**
HT-9 specimen tested in 34.5 MPa hydrogen at 298K with A and B indicating regions of hydrogen induced fracture at the surface.
Table II of HT-9 Subject
to Cathodic
Yield Strength MPa (ksi)
Ultimate Strength MPa (ksi)
Charging
Reduction in Area (X)
Total Elongation (X)
Spec #*
Charging
Q&T 18
Uncharged
650 (94.3)
792 (114.9)
58.8
16.9
Q&T 16
0.003 A/cm*-8 min.
656 (95.2)
823 (119.4)
62.1
17.8
Q&T 20
0.003 A/cm*-150
min. 676 (98.0)
816 (118.4)
59.6
16.5
Q&T 23
0.006 A/cm*-150
min. 673 (97.6)
836 (121.3)
60.0
17.2
Q 17
Uncharged
980 (142.2)
1676 (243.1)
28.0
8.8
Q 15
0.003 A/cm*-8 min.
943 (136.7)
1616 (234.4)
27.1
8.0
Q 19
0.003 A/cm*-150
min.
1261 (182.9)
1564 (226.9)"""
6.7
1.7
Q 26
0.003 A/cm*-150
min.
1225 (177.6)
1731 (251.0)*X*
7.2
3.7
Q 24
0.006 A/cm*-150
min.
554 (80.3)***
6.0
.l
*Q&T-Quenched-and-Tempered Q-As-Quenched **Current ***Fracture
Microstructure
Microstructure
Density Stress
(amps/cm*)
and Charging
Time
J. M. Hyzak, R.E. Stoltz
880
f H,,drogen
quenched material compared to the quenched-andtempered condition. Charging at O.O03A/cm2 for 2.5 hours increased the yield strength by 260 MPa and reduced the ductility by 78%, from 28% to 7.2%. Charging for a comparable time at twice that current density caused further embrittlement; the specimen failed at 554 MPa (80.3 ksi), before general yielding occurred. The fracture appearance of the as-quenched specimens differed in both the charged and uncharged conditions from that observed for the quenched-and-tempered specimens. The uncharged specimen with the as-quenched microstructure had a cup-cone type failure with a blocky topography on the fracture surface and some secondary cracking. The fracture surface exhibited regions of both dimpled rupture and a more brittle faceted fracture. Examination of cross-sections of the fracture surfaces suggested that fracture was transgranular proceeding mainly along untempered martensitic lath boundaries. Fracture of the highly charged samples (Q19, 024, Q26) was much more flat and brittle. Specimen 919 failed in a cup-cone mode, but the fracture surface was dominated by faceted fracSpeciture along untempered lath boundaries. men Q24 which was the most severely charged exhibited surface crack initiation and no shear lip. The fracture initiated at surface flaws associated with the machining process. 4.
computibilit~~
oftIT-
murtetrsitic
SS
bide/matrix interfaces, which are incoherent. Thus raising the temperature decreases the amount of hydrogen associated with the lath interfaces compared to that retained at the carbides. In this way ductile fracture at the carbides would be enhanced. Raising the total Pressure at 298K does not affect the partitioning of hydrogen, but rather increases the amount at the carbides and lath interfaces proportionally. Secondly, the mixed mode of fracture exhibited in the 298K/0.10 MPa test suggests a competition between the two fracture processes. Ductile void initiation is the result of imcompatibility of flow between the hard carbides and softer At elevated temperature this incompatimatrix. bility is increased due to the further softening of the matrix. Thus, increasing temperature may not suppress the hydrogen-induced martensite fracture but rather enhance the ductile void Future tests at elevated temperatures process. and higher pressures are planned and will be useful in verifying this hypothesis. At hydrogen concentrations above those calculated for first wall structures[+l, neither internal nor external hydrogen have a large effect on the tensile ductility of quenched-and-tempered HT-9. The most severe cathodic charge ($20 wt. PPm H2) resulted in no ductility loss while testing in 0.10 MPa hydrogen (~5 wt. ppm 112) lowered the ductility by only 13%.
DISCUSSION
The results of the tests with external hydrogen indicate that-increasing the hydrogen pressure has a far more detrimental effect on the compatibility of HT-9 than do increases in the It is interesting to compare test temperature. the theoretical increase in solubility at 0.10 The lattice solubility of hyMPa to 34.5 MFa. drogen in Fe[3lis given by: C = 3.7p'exp[-6500/RT];
cm3H2(STP)/cm3Fe
An increase in temperature from 298 to 673K results in an approximately 500 fold increase in solubility while an increase from 0.10 MPa to 34.5 MPa results in an 18 fold increase. On this basis, increasing temperature should increase the effect of hydrogen far more than inA number of factors creasing the pressure. suggest the opposite is true and support the data of Table I. First, the above equation At 298K and 0.10 is for lattice solubility. MPa H2, a concentration of 0.04 atomic ppm is predicted, while a value of l-5 wt. ppm is obThe solubility is thus dominated by served. hydrogen trapped at interfaces and defects in The amount of hydrogen trapped the material. is related to the trap depth and concentration. As the temperature is raised, less hydrogen is retained in the weaker traps, until finally One may speculate all the hydrogen is mobile. that the martensite lath interfaces are weaker traps (since they are coherent) than the car-
The higher strength, as-quenched microstructure (980 MPa yield strength) was more severely The intermediate charge embrittled, however. (0.003 A/cm2, 2.5 hours) resulted in a 78% loss in ductility which was accompanied by a 24% increase in yield strength. At twice that current density, the as-quenched specimen failed at 554 MPa, well below the anticipated yield stress. This increased effect of hydrogen compared to that observed for the quenched-and-tempered specimens may be the result of the quench hardened microstructure having ret i ed austenite at the martensite lath boundaries a59. The austenite phase is mechanically unstable, and it most likely transforms to a twinned martensite structure during tensile straining. The resulting high carbon martensite is particularly brittle resulting in interlath fracture [61. Even in the uncharged condition, the HT-9 as-quenched tensile specimens failed by a combination of interlath fracture and dimpled rupture, However when hydrogen charged, the fracture mode was entirely interlath fracture indicating that hydrogen further reduces the effective strength at that boundary. The as-quenched specimens were used to simulate the yield strength increases observed following irradiation. Besides the presence of retained austenite at the lath boundaries, there are other important differences between a quench hardened microstructure and a quenched-and-tem-
J. M. Hyzak, R. E. Stoltz / Hydrogen
pered microstructure that has been irradiation hardened. Irradiationwould introduce a large number of hydrogen traps along with the increase in strength, In this case, the traps may reduce the susceptibilityof the irradiated alloy to hydrogen embrittlementby binding the hydrogen to innocuous sites, This is the same type of process that is thought to give the quenched-andtempered microstructure its relative immunity to hydrogen embrittlement. However, hydrogen testing of irradiated material is critical to understanding this process, and should be pursued. Finally the results for both the quenchhardened samples and the quenched-and-tempered specimens tested at the higher hydrogen pressures show that increasing the hydrogen concentration can change the fracture morphology and increase the likelihood of surface crack initiation. Uncharged specimens as well as those with a small hydrogen concentration generally fail by cup-cone fracture. Cracking occurs at the centerline of the specimen generally by void nucleation, growth, and coalescence before surface crack initiation occurs. Higher hydrogen levels decrease the damage tolerance of the material such that surface cracks nucleate at lower strains. As a result, subsequent loading after crack initiation has occurred accelerates hydrogen assisted crack growth, 5.
CONCLUSIONS
1.
Both internal and external hydrogen have only a small effect on the tensile ductility of quenched-and-temperedHT-9 at concentration levels of 5-10 wt. ppm, well above those calculated for first wall structures.
2.
Increasing the temperature to 473-673K eliminates any effect of hydrogen in the 510 wt. ppm range on thetensile properties of the quenched-and-temperedspecimens.
3.
Increasing the external hydrogen test pressure to the range of 34.5 MPa, severely reduces the ductility of quenched-and-tempered HT-9 at 298K. The fracture process shifts from ductile void growth and coalescence to a predominately interlath fracture mode.
4.
As-quenched HT-9, which has a significantly higher yield strength than the quenched-andtempered microstructure,is more severely affected by internal hydrogen concentrations in the lo-20 wt. ppm range. Cathodic charging to these levels reduced the ductility 78% and changed the fracture mode to brittle faceted fracture along untempered martensite lath boundaries.
5.
Increasing the hydrogen concentrationand the alloy strength level generally increases the likelihood of surface crack initiation over the cup-cone type fracture found in uncharged samples. Surface cracking, which represents a decrease in defect tolerance, enhances embrittlementby interrupting the centerline cracking process of void nucleation and growth,
compatibility
of HT-9 martemitic
SS
881
REFERENCES 1. Williams, D. P. and Nelson, H. G., Met. Trans. 1 (1970) 63-68. 2. Bernstein, I. M., Matl. Sci. Engr. 6 (1970) 1-19. 3.
Gonzalez, 0. D. (quoted by Oriani, R. A.), Proceedings of Conference on Fundamental Aspects of Stress Corrosion Cracking, NACE (1969), 43.
4.
Stoltz, R. E., Baskes, M. I., and Look, G. W., "Calculationsof Hydrogen Isotope Loading in HT-9 First Wall Structures",ADIP Quarterly Progress Report, September 30, 1980, DOE~ER-0045~4,151-159.
5.
Lechtenberg, T. A., "The Effect of Heat Treatment and Thermal Exposure on the Microstructure and Fracture Properties of a 12CrlMo-0.3V Martensitic Stainless Steel (HT-9)", this conference proceedings.
6. Horn, R. M., and Ritchie, R. O., Met. Trans. 9A (1978) 1039-1053.
HYDROGEN COMPATIBILITYOF HT-9 MARTENSITIC STAINLESS STEEL* J. PI.Hyzak and R. E. Stoltz Sandia National Laboratories Livermore, California 94550 The hydrogen compatibility of the alloy HT-9 is addressed. The results of tensile tests performed on specimens both in a gaseous hydrogen environment and after cathodic charging are described. The findings show that there is only a small effect of hydrogen (at the levels tested) on the ductility of quenched-and-temperedXT-9, However, larger reductions in ductility are observed for the as-quenched specimens. It is also shown that increasing the hydrogen concentrationand the strength level generally increases the likelihood of surface crack initiationwhich is an indication of a decrease in damage tolerance. 1.
INTRODUCTION
The proposed use of a ferritic/martensiticsteel for first wall and blanket structures raises a number of materials compatibilityquestions, including possible hydrogen embrittlement. Hydrogen can be produced in the material from bath direct injection of deuterium and tritium as well as from (n,p) reactions. The presence of interstitialhydrogen in levels as low as l-2 weight ppm is known to severely affect the fracture behavior of martensitic steels[l,21 Embrittlement'increases with yield strength 'and generally decreases with temperatureabove ambient. In this work, HT-9 tensile specimens exposed to both external (gaseous environment)and internal (cathodicallycharged) hydrogen have been tested. The effects of temperature,hydrogen pressure and alloy strength level on the tensile embrittlementof HT-9 are reported herein. 2.
EFFECT OF EXTERNAL HYDROGEN EXPOSURES ON TENSILE PROPERTIES OF HT-9
2.2 Test Results The results of all tensile tests are presented in Table I; the values are the average of two tests. In no instance did the presence of hydrogen significantly affect the yield or ultimate strength of the alloy. The effect of the hydrogen was manifest in the ductility as measured by reduction-in-area(RA) at fracture. At 298K, 0.10 MPa hydrogen reduced the tensile ductility from 68 to 58%. Although this loss in ductility is small, it does suggest that a hydrogen-inducedchange in fracture behavior has occurred. This effect was verified by comparison of the fracture surfaces of samples tested in air and in 0.10 MPa hydrogen, shown in Figures 1 and 2. The fracture process in air is entirely one of ductile void initiation, growth and coalesence. The overall fracture surface is typical of cup-cone type fracture, Figure 1, including a flat central region with radial cracks extending to an outer shear lip. All the air-fracture surfaces exhibit a ductile void fracture morphology. Comparison of the void size with the microstructure indicates that void initiation occurs at the tempered carbides,
2.1 Experimental Details To determine the effect of external hydrogen on HT-9, previously uncharged tensile specimens were tested to failure in a gaseous hydrogen environment. All tests were performed on samples taken from a pipe section of HT-9 with composition: Chemistry of HT-9, Wt.% C Si Mn Cr Ni MOW ------_--.22 .38 .52 11.3 .50 .85 .50
V .27
P S .019 ,006
Prior to testing, the specimens were austenitized at 1313K/30 min. and air cooled, followed by tempering at 1023K/60 min., with a final hardness of R,Zl. The resulting microstructure was fully martensitic, with a distribution of tempered carbides. In the hydrogen experiments samples were equilibrated at pressure and temperature for 30 min. before testing. *This work supported by U.S. Department of Energy, DOE, under contract number DE-AC04-76DP 00789.
0022-3115/81/0000-0000/%02.75 0 1981North-Holland
Figure 1. Cup-cone fracture of HT-9 tested in air at 298K.
878
J M. Hyzak,
R.E. Stoltz
/ Hydrogen
compatibility
of HT-9
martensitic
SS
air and 0.10 MPa hydrogen. No hydrogen-induced interlath fracture was observed; instead, ductile dimple fracture occurred across the 473673K temperature range of test.
Figure
2.
HT-9 tested in 0.10 MPa hydrogen at 298K with A and B indicating regions of hydrogen induced fracture.
The large radial fissures are regions where the ductile cracks join, and the side surfaces of these fissures exhibit elongated dimples similar In hydrogen, to those in the shear lip region. Figure 2, the overall features of the fracture surfaces are similar to those in air. However, there are regions of surface crack initiation, areas A and B, which exhibit an alternate fracture mode. The fracture is dominated by shallow The secondary cracks and a "woody" appearance. size of the cracks suggest that fracture occurs at or between the blocky martensite laths. At elevated temperatures, hydrogen effects are suppressed and the ductilities are comparable in
Mechanical
In order to study the effect of hydrogen pressure at ambient temperature, a series of tests was performed in the range from 6.9-34.5 MPa hydrogen. As indicated in Table I, a progressive but slight decrease in ductility occurs up to 20.7 MPa, with a sharp decrease manifest at 34.5 MPa hydrogen pressure. In the range 6.920.7 MPa hydrogen, the fracture behavior is similar to that at 0.10 MPa. Small regions of martensite interlath fracture are observed at the outer part of the fracture surface. As the test pressure is increased, the area percentage of this surface crack initiation mode increases. At 34.5 MPa, Figure 3, the hydrogen affected regions are much larger (A and B), comprising approximately 50% of the fracture area. The sample exhibits two fracture regions perpendicular to the tensile axis connected by a ductile shear region. Little overall necking of the sample was observed. When seen at higher magnifications, the flat regions show martensite interlath fracture with secondary cracks similar to that in 0.10 MPa hydrogen. In some regions of the sample, the martensite was oriented nearly perpendicular to the tensile axis, so that the interlath fracture exhibits a stepped appearance. 3.
EFFECT OF INTERNAL HYDROGEN PROPERTIES OF HT-9
ON THE TENSILE
To better
represent the hydrogen distribution in a first wall structure, specimens were internal-
Table I Properties of HT-9 in Air and Hydrogen Ultimate Strength MPa (ksi)
Ductility (Reduction %
Temperature
Atmosphere
Yield Strength MPa (ksi)
25OC
Air
592 (85.8)
799 (115.8)
67
25OC
0.10 MPa H2
565 (81.9)
837 (121.3)
58
2oooc
Air
533 (77.3)
719 (104.3)
64
2oooc
0.10 MPa H2
545 (79.0)
733 (106.2)
64
3oooc
Air
513 (74.4)
695 (100.7)
65
3oooc
0.10 MPa H2
524 (76.0)
689 (99.9)
67
4oo"c
Air
465 (67.4)
628 (91.0)
62
4oooc
0.10 MPa H2
496 (71.9)
637 (92.3)
63
250C
6.9 MPa Hp.
715 (103.6)
800 (116.0)
65
25'C
13.8 MPa H2
707 (102.5)
788 (114.2)
56
25OC
20.7 MPa H2
717 (104.2)
792 (114.8)
53
25'C
34.5 MPa H2
724 (105.1)
793 (115.0)
26
in Area)
819
J.M. Hyzak, R.E. Stoltz 1 Hydrogen compatibility of HT-9 martensitic SS
ly charged with hydrogen prior to testing. Hydrogen was introduced by cathodic charging, and specimens were immediately copper-plated to retard outgassing. The charging solution was 4 percent sulfuric acid containing 5 mg of sodium The current densities employed were arsenate. O.OO~A/C~* (O.O*A/in*) and 0.006 A/cm* (O.O4A/ in*) for 8 and 150 minutes (Table II). Specimens were aged at room temperature for 24 hours before testing in order to evenly distribute the hydrogen. Preliminary data indicate that even the least severe charging condition (.003A/ cm*, 8 minutes) would result in a hydrogen concentration greater than 5 wt. ppm (275 appm). Material was tested in both the as-quenched and quenched-and-tempered conditions. Since HT-9 will be subject to irradiation hardening, the as-quenched microstructurewas tested as a Preliminary means of assessing the effect of hardening on hydrogen compatibility. All specimens were austenitized as above and air cooled; half the specimens were tested as is and the remainder tempered as above. 3.1 Experimental
Figure 3.
Results
surrounded by a uniform shear lip. The fracture mode in the central section was dimpled rupture. Examination of metallographic cross-sections of the fracture surfaces revealed that fracture occurred in the tensile (rolling) direction along martensite lath boundaries as the result of nucleation and growth of voids at tempered carbides and stringers. The as-quenched microstructure had an uncharged yield strength substantially higher than that of the quench and tempered condition, 980 MPa versus 650 MPa. There was also a much larger effect of hydrogen charging on the tensile properties in the
The results of all the tensile tests are given in Table II. For the quenched-and-tempered microstructure, there was no significant effect of internal hydrogen on the ductility. However, there did appear to be a small but consistent increase in both yield and ultimate strength as a result of hydrogen charging. The fracture process for both the charged and uncharged specimens was the same, classical cup-cone failure. The fracture surfaces were rough with many large secondary cracks in the interior
Mechanical
Properties
Cond.**
HT-9 specimen tested in 34.5 MPa hydrogen at 298K with A and B indicating regions of hydrogen induced fracture at the surface.
Table II of HT-9 Subject
to Cathodic
Yield Strength MPa (ksi)
Ultimate Strength MPa (ksi)
Charging
Reduction in Area (X)
Total Elongation (X)
Spec #*
Charging
Q&T 18
Uncharged
650 (94.3)
792 (114.9)
58.8
16.9
Q&T 16
0.003 A/cm*-8 min.
656 (95.2)
823 (119.4)
62.1
17.8
Q&T 20
0.003 A/cm*-150
min. 676 (98.0)
816 (118.4)
59.6
16.5
Q&T 23
0.006 A/cm*-150
min. 673 (97.6)
836 (121.3)
60.0
17.2
Q 17
Uncharged
980 (142.2)
1676 (243.1)
28.0
8.8
Q 15
0.003 A/cm*-8 min.
943 (136.7)
1616 (234.4)
27.1
8.0
Q 19
0.003 A/cm*-150
min.
1261 (182.9)
1564 (226.9)"""
6.7
1.7
Q 26
0.003 A/cm*-150
min.
1225 (177.6)
1731 (251.0)*X*
7.2
3.7
Q 24
0.006 A/cm*-150
min.
554 (80.3)***
6.0
.l
*Q&T-Quenched-and-Tempered Q-As-Quenched **Current ***Fracture
Microstructure
Microstructure
Density Stress
(amps/cm*)
and Charging
Time
J. M. Hyzak, R.E. Stoltz
880
f H,,drogen
quenched material compared to the quenched-andtempered condition. Charging at O.O03A/cm2 for 2.5 hours increased the yield strength by 260 MPa and reduced the ductility by 78%, from 28% to 7.2%. Charging for a comparable time at twice that current density caused further embrittlement; the specimen failed at 554 MPa (80.3 ksi), before general yielding occurred. The fracture appearance of the as-quenched specimens differed in both the charged and uncharged conditions from that observed for the quenched-and-tempered specimens. The uncharged specimen with the as-quenched microstructure had a cup-cone type failure with a blocky topography on the fracture surface and some secondary cracking. The fracture surface exhibited regions of both dimpled rupture and a more brittle faceted fracture. Examination of cross-sections of the fracture surfaces suggested that fracture was transgranular proceeding mainly along untempered martensitic lath boundaries. Fracture of the highly charged samples (Q19, 024, Q26) was much more flat and brittle. Specimen 919 failed in a cup-cone mode, but the fracture surface was dominated by faceted fracSpeciture along untempered lath boundaries. men Q24 which was the most severely charged exhibited surface crack initiation and no shear lip. The fracture initiated at surface flaws associated with the machining process. 4.
computibilit~~
oftIT-
murtetrsitic
SS
bide/matrix interfaces, which are incoherent. Thus raising the temperature decreases the amount of hydrogen associated with the lath interfaces compared to that retained at the carbides. In this way ductile fracture at the carbides would be enhanced. Raising the total Pressure at 298K does not affect the partitioning of hydrogen, but rather increases the amount at the carbides and lath interfaces proportionally. Secondly, the mixed mode of fracture exhibited in the 298K/0.10 MPa test suggests a competition between the two fracture processes. Ductile void initiation is the result of imcompatibility of flow between the hard carbides and softer At elevated temperature this incompatimatrix. bility is increased due to the further softening of the matrix. Thus, increasing temperature may not suppress the hydrogen-induced martensite fracture but rather enhance the ductile void Future tests at elevated temperatures process. and higher pressures are planned and will be useful in verifying this hypothesis. At hydrogen concentrations above those calculated for first wall structures[+l, neither internal nor external hydrogen have a large effect on the tensile ductility of quenched-and-tempered HT-9. The most severe cathodic charge ($20 wt. PPm H2) resulted in no ductility loss while testing in 0.10 MPa hydrogen (~5 wt. ppm 112) lowered the ductility by only 13%.
DISCUSSION
The results of the tests with external hydrogen indicate that-increasing the hydrogen pressure has a far more detrimental effect on the compatibility of HT-9 than do increases in the It is interesting to compare test temperature. the theoretical increase in solubility at 0.10 The lattice solubility of hyMPa to 34.5 MFa. drogen in Fe[3lis given by: C = 3.7p'exp[-6500/RT];
cm3H2(STP)/cm3Fe
An increase in temperature from 298 to 673K results in an approximately 500 fold increase in solubility while an increase from 0.10 MPa to 34.5 MPa results in an 18 fold increase. On this basis, increasing temperature should increase the effect of hydrogen far more than inA number of factors creasing the pressure. suggest the opposite is true and support the data of Table I. First, the above equation At 298K and 0.10 is for lattice solubility. MPa H2, a concentration of 0.04 atomic ppm is predicted, while a value of l-5 wt. ppm is obThe solubility is thus dominated by served. hydrogen trapped at interfaces and defects in The amount of hydrogen trapped the material. is related to the trap depth and concentration. As the temperature is raised, less hydrogen is retained in the weaker traps, until finally One may speculate all the hydrogen is mobile. that the martensite lath interfaces are weaker traps (since they are coherent) than the car-
The higher strength, as-quenched microstructure (980 MPa yield strength) was more severely The intermediate charge embrittled, however. (0.003 A/cm2, 2.5 hours) resulted in a 78% loss in ductility which was accompanied by a 24% increase in yield strength. At twice that current density, the as-quenched specimen failed at 554 MPa, well below the anticipated yield stress. This increased effect of hydrogen compared to that observed for the quenched-and-tempered specimens may be the result of the quench hardened microstructure having ret i ed austenite at the martensite lath boundaries a59. The austenite phase is mechanically unstable, and it most likely transforms to a twinned martensite structure during tensile straining. The resulting high carbon martensite is particularly brittle resulting in interlath fracture [61. Even in the uncharged condition, the HT-9 as-quenched tensile specimens failed by a combination of interlath fracture and dimpled rupture, However when hydrogen charged, the fracture mode was entirely interlath fracture indicating that hydrogen further reduces the effective strength at that boundary. The as-quenched specimens were used to simulate the yield strength increases observed following irradiation. Besides the presence of retained austenite at the lath boundaries, there are other important differences between a quench hardened microstructure and a quenched-and-tem-
J. M. Hyzak, R. E. Stoltz / Hydrogen
pered microstructure that has been irradiation hardened. Irradiationwould introduce a large number of hydrogen traps along with the increase in strength, In this case, the traps may reduce the susceptibilityof the irradiated alloy to hydrogen embrittlementby binding the hydrogen to innocuous sites, This is the same type of process that is thought to give the quenched-andtempered microstructure its relative immunity to hydrogen embrittlement. However, hydrogen testing of irradiated material is critical to understanding this process, and should be pursued. Finally the results for both the quenchhardened samples and the quenched-and-tempered specimens tested at the higher hydrogen pressures show that increasing the hydrogen concentration can change the fracture morphology and increase the likelihood of surface crack initiation. Uncharged specimens as well as those with a small hydrogen concentration generally fail by cup-cone fracture. Cracking occurs at the centerline of the specimen generally by void nucleation, growth, and coalescence before surface crack initiation occurs. Higher hydrogen levels decrease the damage tolerance of the material such that surface cracks nucleate at lower strains. As a result, subsequent loading after crack initiation has occurred accelerates hydrogen assisted crack growth, 5.
CONCLUSIONS
1.
Both internal and external hydrogen have only a small effect on the tensile ductility of quenched-and-temperedHT-9 at concentration levels of 5-10 wt. ppm, well above those calculated for first wall structures.
2.
Increasing the temperature to 473-673K eliminates any effect of hydrogen in the 510 wt. ppm range on thetensile properties of the quenched-and-temperedspecimens.
3.
Increasing the external hydrogen test pressure to the range of 34.5 MPa, severely reduces the ductility of quenched-and-tempered HT-9 at 298K. The fracture process shifts from ductile void growth and coalescence to a predominately interlath fracture mode.
4.
As-quenched HT-9, which has a significantly higher yield strength than the quenched-andtempered microstructure,is more severely affected by internal hydrogen concentrations in the lo-20 wt. ppm range. Cathodic charging to these levels reduced the ductility 78% and changed the fracture mode to brittle faceted fracture along untempered martensite lath boundaries.
5.
Increasing the hydrogen concentrationand the alloy strength level generally increases the likelihood of surface crack initiation over the cup-cone type fracture found in uncharged samples. Surface cracking, which represents a decrease in defect tolerance, enhances embrittlementby interrupting the centerline cracking process of void nucleation and growth,
compatibility
of HT-9 martemitic
SS
881
REFERENCES 1. Williams, D. P. and Nelson, H. G., Met. Trans. 1 (1970) 63-68. 2. Bernstein, I. M., Matl. Sci. Engr. 6 (1970) 1-19. 3.
Gonzalez, 0. D. (quoted by Oriani, R. A.), Proceedings of Conference on Fundamental Aspects of Stress Corrosion Cracking, NACE (1969), 43.
4.
Stoltz, R. E., Baskes, M. I., and Look, G. W., "Calculationsof Hydrogen Isotope Loading in HT-9 First Wall Structures",ADIP Quarterly Progress Report, September 30, 1980, DOE~ER-0045~4,151-159.
5.
Lechtenberg, T. A., "The Effect of Heat Treatment and Thermal Exposure on the Microstructure and Fracture Properties of a 12CrlMo-0.3V Martensitic Stainless Steel (HT-9)", this conference proceedings.
6. Horn, R. M., and Ritchie, R. O., Met. Trans. 9A (1978) 1039-1053.