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Hydrogen embrittlement behavior of palladium modified PH 13-8 Mo stainless steel as a function of age hardening
Pergamon
ScriptaMetaUurgicaet Materialia,Vol. 31, No. 2, pp. 125-130,1994 1994ElsevierScienceLtd Printedin the USA.All rights reserved 0956-716X/94$6.00+ 00
HYDROGEN EMBRITrLEMENT BEHAVIOR OF PALLADIUM MODIFIED PH 13-8 MO STAINLESS STEEL AS A FUNCTION OF AGE HARDENING
J.R. Sallly Department of MaterialsScienceand Engineering Universityof Virgi',i~ Charlottesville,VA 22903-2442
MJ. Cieslak and J.A. Van Den Avyle MetallurgyDepartment Saudia NationalLaboratories Albuquerque, blewMexleo,87185 (Received March 22, 1994) Introduction The hydrogen embrittlement (liE) susceptibility of precipitation age hardened stainless steels (17-4 PH, PH 15-5, PH 13-8 Me, and others) is well established in the literature [1-5]. Susceptibility is a strong function of strength and hence lower aging temperatures produce alloys which are more prone to HE [5]. Recently its was shown that uniformly distributed PdAI precipitates improve the HE resistance of a palladium (Pd) modified version of PH 13-8 Me stainless steel at a yield strength of 1250 MPa [6]. Intergranular cracking on prior austenite grain boundaries is minimized and transgranular cracking is observed. Trapping analysis of hydrogen permeation data supports the notion that PdAI particles act as beneficial trap sites [6]. However, the effect of PdAI on the HE resistance of higher strength tempers has not been established. Secondly, the influence of dual phase microstructures typical of east or as-welded PH 13-8 Me on the PdA1 effect has not been investigated. The concept of alloying high strength ferrous base alloys with noble metals to improve HE resistance is not unique to the work described above. Small Pd additions also reduce HE susceptibility in quenched and tempered AISI 4130 steel [7-9]. Klscc and threshold stress in sustained load tests are significantly improved. Transgranular tearing and ductile microvoid formation are observed instead of intergranular cracking with 1 wt.% Pd at yield strengths of both 745 MPa [7,8] and 1170 MPa [9]. Interracial segregation of Pd to interphase interfaces such as MnS inclusions and lath boundaries, it is argued, repels hydrogen from these otherwise sitmificant trapping sites [10]. Pd segregation at MnS inclusions was found to dimlni~h tritium trapping at the sulfide/matrix interface [11]. The implication is that Pd segregation alters the distribution of trapped hydrogen, which in turn affects resistance to HE. In the present study, we examine the influence of Pd on the HE susceptibility of PH 13-8 Me stainless steel (SS) over a range of strengths levels and also for a dual phase as-welded microstructure. Specifically, we compare the HE behavior of conventional wrought PH 13-8 Me SS to wrought PH 13-8 SS alloyed with 0.4, 0.75 and 1.0 wt. % palladium and similarly alloyed, autogenous electron beam welded PH 13-8 Me SS. Wrought alloys were examined after aging to yield strengths of nominally 1000, 1140 and 1440 MPa in accordance with the Hll00, H1050 and H925 tempers, respectively. The as-welded material was similar in hardness to the wrought Hll00 condition.
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Pd additions were made to 3 heats of alloy PH 13.8 Mo (0.05% C, 0.10% Mn, 0.01% P, 0.008% S, 0.10% Si, 1325% Cr, 8.50% Hi, 135% A1, 2.50% Mo, Bal Fe). The Pd levels chosen correspond to two levels (0.75 and 1.0 wt. %) previously observed to improve hydrogen cracking resistance and one level (0.39%) which had no effect [7,8]. Complete chemical compositions and processing details are reported elsewhere [6,12]. Solutionized samples (930°C) were aged to the Hll00, H1050, and H925 conditions for 4 h producing the mechanical properties indicated in Table 1. Table I. Mechanical Properties of PH 13.8 Mo SS in Laboratory Air at High and Low Compositional Limits with Respect to Pd Concentration. PH temper and wt. pct. Pd
Yield Strength 0va'a)
Tensile Strength
H 925 - 0% Pd
1415
H 925 - 1.02% Pd
Elongation
Hardness (R~ scale)
1522
14.7
46
1373
1468
15.1
47
H 1050 - 0% Pd
1248
1283
15.8
42
H 1050 - 1.02% Pd
1174
1236
18.9
40
H 1100 - 0% Pd
1031
1129
20.1
36
0,~Pa)
H 1100 - L02% Pd 1128 21.1 36 979 Iote~ 1) reported results are the average from three smooth bar tensile samples. 2) hardness of welds, for which tensile properties were not available, ranged from Rockwell C 37 to 38 Constaat extension rate tensile (CERT) tests were conducted with circumferentiaily notched round bar specimens with a semicircular polished notch of radius equal to 0.07 cm and a notch diameter of 0.497 cm. The constraint level was 1.9 (0.33=uniaxial tension, 2.5=sharp notch). The weld fusion zone of electron beam welds was centered at the root of the notch as confirmed metallographically. Specimens were strained to failure at lxl0 ~ on/second crosshead displacement rate in deaerated 0.6 M NaC1 solution. Specimens were subjected to continual cathodic polarization at -1.1 Volt vs. a Saturated Calomel Electrode {SCE} for 30 hours prior to commencement and also during CERT testing. These conditions have been observed to cause si~nlflcant H E in AISI 4340 steel at a xlmilar yield strength [I3]. Bridgman's analysis of naturally necking bars was employed to aid in the interpretation of notched tensile results [14]. A Hitachi Model S-500 microscope was used for fi'actography. Results PH 13-8 Mo is a fully martensitic precipitation hardening stainless steel which transforms to martensite at relatively low temperature (M, = 130°C, M t = 20°C). PH 13-8 Mo is anstenitic above 8000C. Precipitation of the intermetalfic phase, ~-NiAl, occurs within the martensitic matrix in the temperature range of 425-620°C and is responsible for precipitation strengthening. Examlnation of as-cast (vacuum-arc-remelted) PH 13-8 Mo showed that (a) Pd lowers the M~ temperature, Co) Pd is completely soluble in austenite and therefore is present in solid solution in the martensitic phase of the duplex as-cast structure, and (c) 1 ~m size PdAl precipitates are formed in the ferrite phase of the duplex as-cast structure [12]. Similarly, Pd is also expected to partition between PdAl parfides in the ~ ferrite phase and substitutional solid solution in the martensite of dual phase weld zone material (Figure 1). In the single phase, martensitic wrought product, submicrometer sized PdAl precipitates are observed after age hardening heat treatments [6]. These particles are uniformly distributed throughout the
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martensite and are not uniquely associated with either prior austenite grain boundaries or marteusitic lath boundaries [6]. PdAI formation causes a slight decrease in post-aged yield and tensile strengths because of its competition with Ni for the consumption of AI. A corresponding slight increase in ductility is observed with Pd, particularly at the 1% Pd level, due to this effect (Table 1). Grain and lath boundaries were analyzed for segregation of Pd [6]. Segregation of Pd to austenitic grain boundaries during solutionlzing is not anticipated as the microstructure is wholly austenltic and Pd is completely soluble. For the aged material, a Pd concentration at the grain boundary statistically indistinguishable from the matrix was found [6]. Susceptibility of wrought PH 13-8 Mo SS as a function of strength and Pd concentration The influence of Pd concentration on the mechanical properties of wrought H 13-8 Mo at three strength levels is shown in Figure 2. Here, HE susceptibility is expressed as the ratio of the maximum load attained in the CERT test to the maximum load obtained from CERT testing in air following the procedure of Hancock [15]. Similar trends exist in the case of susceptibility expressed as the ratios of times to failure; these plots are not shown. Ductile microvoid formation and growth dominated fracture is observed after tests in air at all strength levels. With charging, the highest strength conventional alloy (H925) exhibits the greatest susceptible to HE as expected [5]. A combination of intergranular separation on prior austenite grain boundaries and brittle transgranular fracture is observed for the hydrogen charged conventional alloys at all strength levels (Figure 3). In all cases, this brittle outer zone extends radially from the root of the notch and transitions gradually to microvoid coalescence characteristic of ductile overload in the center of the notched sample even though the ma,fimum longitudinal applied stress occurs at the center [14]. Concerning Pd additions, the Hll00 temper exhibits the lowest HE susceptibility over the entire range of Pd concentrations while the H925 temper exhibits the greatest. However, the Pd effect is still operative at highest strength temper although 0.39% Pd has only marginal benefit. Intergranular fracture is completely absent for Pd containing alloys at all strength levels (Figure 3). Instead, hydrogen assisted cracking is transgranular with an appearance that is consistent with inter-lath and tram-lath tearing and; possibly, cleavage. HE cracking also extends radially from the root of the notch and transitions to microvoid nucleation and growth dominated fracture in the center of the sample. Since the longest time to failure in these tests was 60-70 hours, hydrogen charging to obtain a uniform hydrogen concentration across the entire cross-section of the notched diameter was not obtained during the testing period even after the 30 h pre-charge. This conclusion is reached based on the low hydrogen diffusion coefficients for PH 13-8 Mo SS and calculated times required to obtain a uniform hydrogen concentration [6]. To verify that the Pd effect is still operative after long charging times to insure uniform hydrogen concentration at the highly stressed centerline of the notch an additional specimen (0.75% Pd at H1050) was cathodicaUy charged in 0.1 molar NaOH. This specimen was continually charged at -1130mA/cm 2, producing a hydrogen fugacity far exceeding the -1.1 V vs. SCE condition. Failure was not observed even after over 200 days at a constant load greater than the maximum load producing failure in conventional H1050 alloys by CERT testing. Suscentibilitv of welded PH 13-8 Mo as a function of Pd concentration Figure 2 also illustrates the effect of hydrogen on as-welded PH 13-8 Mo SS as a function of Pd. The weld is not as resistant to hydrogen as the wrought Hll00 alloy, which is of comparable hardness (Table 1). The welded dual phase microstructure exhibits greater susceptibility than the wrought marteusitic material at a similar strength level. Metallographic examination of longitudinal sections from fracture surfaces as well as SEM fractography indicate that cracking is transgranular across the ~ ferrite and marteusite phases and is not uniquely associated with the ferrite/marteusite interface. Discussion and Conclusions The addition of Pd to aged PH 13-8 Mo SS results in (a) formation of PdA1 precipitates, (b) decreased bulk hydrogen diffusion coefficients compared to the conventional alloy, (c) changes in hydrogen assisted fracture modes, and (d) improved HE resistance with increasing Pd concentration [6]. Three possible explanations are:
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(i) inhibited hydrogen ingress due to Pd induced catalysis of H z recombinative desorption, (//) Pd segregation to interfaces and subsequent repulsion of hydrogen, and (///) formation of a strong reversible trap (FdAI particles) that diminishes hydrogen segregation at other competing trap sites and retards occlusion within the stress field of the notch. Explanation (i) is dismissed because neither differences in exchange current densities, Tafel slopes, nor hydrogen production rates were observed with Pd [6]. Consequently, Pd is not expected to inhibit hydrogen ingress. In regard to mechanism (//), sitmificant segregation of Pd to metallurgical interfaces was not detected [6]. Moreover, Shirley and Hall predict that Pd only weakly repels hydrogen in iron [17]. Mechanism (///) is most consistent with the present study and is supported by trapping analysis of permeation data [6]. To summarize this view, permanently reduced apparent hydrogen diffusion coe~cients are obtained in aged Pd modified alloys, but not in solutionized Pd containing versions with Pd is solid solution. This supports the notion that PdA1 particles create a high density of traps that absorb hydrogen both irreversibly and reversibly. If such traps were irreversible and easily saturated, the steady state diffusion data produced from a series of permeation rise and decay transients would be the same for both the aged Pd modified alloy and the conventional alloy after the initial charging step saturated such irreversible traps [18]. This was not observed [6]. Operating on the assertion that PdAI particles are a strong reversible trap, it is worth reviewing the possible benefits. Recall that the criteria for optimum utilization of beneficial traps include (a) small in size so as not to serve themselves as fracture initiation sites, (b) high density and uniformly distributed, (c) revers~le and non-saturable [20-22]. An explanation involving beneficial trapping by PdA1 particles can account for all of the results of the present study. HE resistance improves with increasing Pd concentration and is not completely absent even at 0.38% Pd. This can be interpreted in connection with the benefits of an increased density of PdA] particles. Pd was more beneficial at lower yield strengths. This observation is consistent with the notion that the critical lattice hydrogen concentration required to trigger HE is inversely proportional to alloy strength [23]. Therefore, mobile hydrogen that reaches the fracture process zone is more "potent" with increasing strength [23]. Assnmlng a finite extent to which PdAI particles can trap hydrogen and retard its repartitionlng to the stress field of the notch, higher strength alloys would be more adversely affected by remaining mobile hydrogen. Finally, the PdA1 effect is diminished for the dual phase weld. This is attributed to the non-uniform distribution of PdA1 particles which are present in the 6 ferrite but not the martensite phase of the as-cooled weld metal (Figure 1). Other factors such as intrinsic susceptibility may affect the behavior of the dual phase weld metal as well.
This work was supported by the U.S. Department of Energy under contract number DE-AC04-76DP00789. The authors wish to acknowledge F. Bovard, G. Young, Jr. and D. Enos for various experimental contributions. We would also like to thank the staff of the Sandia National Laboratories Melting and Solidification Facility for preparing the alloys. References 1. C.T. Fujii, ASTM STP 610, H. Craig, ed., ASTM, Philadelphia, PA, 213, (1976). 2. A.W. Thompsoi~ MetaU. Trans., 4, 2819, (1973). 3. G.T. Murray, MetaH. Trans. A, 12(A), 2138, (1981). 4. G.T. Murray, H.H. Honegger, T. Mousel, Corrosion J., 40, 146, (1984). 5. P. Munn~ B. Andersson, Corrosion J., 46, 286, (1990). 6. J.R. Scully, J.A. Van Den Avyle, M.J. Cieslak, A.D. Romig, C.R. Hills, MetaH. Trans. A, 22A, 2429, (1991). 7. B. E. Wilde, C.D. Kim, and J.C. Turn Jr., Corrosion, 38, 515, (1982). 8. J.B. l~rnsden, B.E. Wilde, and P.J. Stocker, Scripta Met., 17, 971, (1983). 9. B.E. Wilde, I. Chattoraj, T.A. Mozhi, Scripta Met., 21, 1369, (1987). 10. M.IC Miller, S.S. Brenner, M.G. Burke, Met Trans A, 18A, 519, (1987). 11. T. D. I.e, and B.E. Wilde, In Current Solutions to Hydrogen Problems in Steels, Eds. C.G. Interrante, G.M. Pressouyre, ASM, Metals Park, Ohio, 413 (1982). 12. M.J. Cieslak, j,a~. Van Den Avyle, C.R. Hills, R.E. Semarge, Met Trans A, 19A, 3063 (1988).
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HYDROGEN EMBRITTLEMENT
13. J.R. Seully, PJ. Moran, Corrosion, 44, 130, (1988). 14. P.W. Bridgman, Studies in Large Plastic Flow, McGraw-Hill, Ine, 9 (1952)• 15. J.W. Hancock, A.C. Mackenzie, J. Mech. Phys. Solids, 24, 147 (1976). 16. H.W. Picketing, M. Zamanzadeh, In Hydrogen Effects in Metals, Met. Soe. of AIME, I.M. Bernstein, A.W. Thompson, eds., 143, (1981). 17. A.I. Shirley, C.K. Hall, Seipta Met., 17, 1003, (1983). 18. G.M. Pressouryre, F.M. Faure., ASTM STP 962, L. Raymond, Ed., ASTM, philadelphia, 353 (1980). 19. R.A. Oriani, Acta Met., 18, 147, (1970). 20. G.M. Pressouyre and I.M. Bernstein, Acta Met., 27, 89, (1979). 21. G.M. Pressouyre, Met Tram A, 14A, 2189 (1983). 22. G.M. Pressouyre, Met Tram A, 10A, 1571 (1979). 23. K.N. Akhurst, TJ. Baker, Met Tram A, 12A, 1059, (1981). .
.
.
.
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!
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~
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o I
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20
POSITION
9 30
•
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40
(MICROMETERS)
Figure 1. (a) X-ray map showing Pd distribution (Pd Le) in the martensite and 6 ferrlte phases of an as-cast PH 13-8 Mo mi~rostrueture with 1% Pd simulating a as-cooled weld, and Co) microprobe analysis across ferrite/martemite interface [12]. ;
1
0
I 1031
~
i
0
MPa
~
~
~
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~
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1248 weld
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¢
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.
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- 1031
MPa
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MPe
[.
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[ ----)<-
0.2
I
- 1415 - Weld
I
I
0.55
0.825
0 0.275
Weight
%
MPa
1,1
Pd
Figure 2. Relationship between ratio of maximum load during slow straining with concurrent hydrogen charging to that obtained in air as a function of Pd concentration. Results are shown for H925, H1050, Hll00 tempers and as-cooled welds.
130
HYDROGENEMBRJTFLEMENT
0% Pd
Vol. 31, No. 2
I% Pd
A
B
C
Figure 3. Fractography from CERT at -1.1V~ for: (a) H l l 0 0 at 0% (left) and 1% Pd (right), (b) H1050 at 0% (left) and 1% Pd (right), and (c) H925 at 0% 0eft) and 1% I'd (right). Transgranular fracture was observed for weld metal. Microvoid fracture was observed in air CERT for all tempers and weld metal. Micron markers are indicated on each micrograph.
ScriptaMetaUurgicaet Materialia,Vol. 31, No. 2, pp. 125-130,1994 1994ElsevierScienceLtd Printedin the USA.All rights reserved 0956-716X/94$6.00+ 00
HYDROGEN EMBRITrLEMENT BEHAVIOR OF PALLADIUM MODIFIED PH 13-8 MO STAINLESS STEEL AS A FUNCTION OF AGE HARDENING
J.R. Sallly Department of MaterialsScienceand Engineering Universityof Virgi',i~ Charlottesville,VA 22903-2442
MJ. Cieslak and J.A. Van Den Avyle MetallurgyDepartment Saudia NationalLaboratories Albuquerque, blewMexleo,87185 (Received March 22, 1994) Introduction The hydrogen embrittlement (liE) susceptibility of precipitation age hardened stainless steels (17-4 PH, PH 15-5, PH 13-8 Me, and others) is well established in the literature [1-5]. Susceptibility is a strong function of strength and hence lower aging temperatures produce alloys which are more prone to HE [5]. Recently its was shown that uniformly distributed PdAI precipitates improve the HE resistance of a palladium (Pd) modified version of PH 13-8 Me stainless steel at a yield strength of 1250 MPa [6]. Intergranular cracking on prior austenite grain boundaries is minimized and transgranular cracking is observed. Trapping analysis of hydrogen permeation data supports the notion that PdAI particles act as beneficial trap sites [6]. However, the effect of PdAI on the HE resistance of higher strength tempers has not been established. Secondly, the influence of dual phase microstructures typical of east or as-welded PH 13-8 Me on the PdA1 effect has not been investigated. The concept of alloying high strength ferrous base alloys with noble metals to improve HE resistance is not unique to the work described above. Small Pd additions also reduce HE susceptibility in quenched and tempered AISI 4130 steel [7-9]. Klscc and threshold stress in sustained load tests are significantly improved. Transgranular tearing and ductile microvoid formation are observed instead of intergranular cracking with 1 wt.% Pd at yield strengths of both 745 MPa [7,8] and 1170 MPa [9]. Interracial segregation of Pd to interphase interfaces such as MnS inclusions and lath boundaries, it is argued, repels hydrogen from these otherwise sitmificant trapping sites [10]. Pd segregation at MnS inclusions was found to dimlni~h tritium trapping at the sulfide/matrix interface [11]. The implication is that Pd segregation alters the distribution of trapped hydrogen, which in turn affects resistance to HE. In the present study, we examine the influence of Pd on the HE susceptibility of PH 13-8 Me stainless steel (SS) over a range of strengths levels and also for a dual phase as-welded microstructure. Specifically, we compare the HE behavior of conventional wrought PH 13-8 Me SS to wrought PH 13-8 SS alloyed with 0.4, 0.75 and 1.0 wt. % palladium and similarly alloyed, autogenous electron beam welded PH 13-8 Me SS. Wrought alloys were examined after aging to yield strengths of nominally 1000, 1140 and 1440 MPa in accordance with the Hll00, H1050 and H925 tempers, respectively. The as-welded material was similar in hardness to the wrought Hll00 condition.
125
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Pd additions were made to 3 heats of alloy PH 13.8 Mo (0.05% C, 0.10% Mn, 0.01% P, 0.008% S, 0.10% Si, 1325% Cr, 8.50% Hi, 135% A1, 2.50% Mo, Bal Fe). The Pd levels chosen correspond to two levels (0.75 and 1.0 wt. %) previously observed to improve hydrogen cracking resistance and one level (0.39%) which had no effect [7,8]. Complete chemical compositions and processing details are reported elsewhere [6,12]. Solutionized samples (930°C) were aged to the Hll00, H1050, and H925 conditions for 4 h producing the mechanical properties indicated in Table 1. Table I. Mechanical Properties of PH 13.8 Mo SS in Laboratory Air at High and Low Compositional Limits with Respect to Pd Concentration. PH temper and wt. pct. Pd
Yield Strength 0va'a)
Tensile Strength
H 925 - 0% Pd
1415
H 925 - 1.02% Pd
Elongation
Hardness (R~ scale)
1522
14.7
46
1373
1468
15.1
47
H 1050 - 0% Pd
1248
1283
15.8
42
H 1050 - 1.02% Pd
1174
1236
18.9
40
H 1100 - 0% Pd
1031
1129
20.1
36
0,~Pa)
H 1100 - L02% Pd 1128 21.1 36 979 Iote~ 1) reported results are the average from three smooth bar tensile samples. 2) hardness of welds, for which tensile properties were not available, ranged from Rockwell C 37 to 38 Constaat extension rate tensile (CERT) tests were conducted with circumferentiaily notched round bar specimens with a semicircular polished notch of radius equal to 0.07 cm and a notch diameter of 0.497 cm. The constraint level was 1.9 (0.33=uniaxial tension, 2.5=sharp notch). The weld fusion zone of electron beam welds was centered at the root of the notch as confirmed metallographically. Specimens were strained to failure at lxl0 ~ on/second crosshead displacement rate in deaerated 0.6 M NaC1 solution. Specimens were subjected to continual cathodic polarization at -1.1 Volt vs. a Saturated Calomel Electrode {SCE} for 30 hours prior to commencement and also during CERT testing. These conditions have been observed to cause si~nlflcant H E in AISI 4340 steel at a xlmilar yield strength [I3]. Bridgman's analysis of naturally necking bars was employed to aid in the interpretation of notched tensile results [14]. A Hitachi Model S-500 microscope was used for fi'actography. Results PH 13-8 Mo is a fully martensitic precipitation hardening stainless steel which transforms to martensite at relatively low temperature (M, = 130°C, M t = 20°C). PH 13-8 Mo is anstenitic above 8000C. Precipitation of the intermetalfic phase, ~-NiAl, occurs within the martensitic matrix in the temperature range of 425-620°C and is responsible for precipitation strengthening. Examlnation of as-cast (vacuum-arc-remelted) PH 13-8 Mo showed that (a) Pd lowers the M~ temperature, Co) Pd is completely soluble in austenite and therefore is present in solid solution in the martensitic phase of the duplex as-cast structure, and (c) 1 ~m size PdAl precipitates are formed in the ferrite phase of the duplex as-cast structure [12]. Similarly, Pd is also expected to partition between PdAl parfides in the ~ ferrite phase and substitutional solid solution in the martensite of dual phase weld zone material (Figure 1). In the single phase, martensitic wrought product, submicrometer sized PdAl precipitates are observed after age hardening heat treatments [6]. These particles are uniformly distributed throughout the
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HYDROGEN EMBRITTLEMENT
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martensite and are not uniquely associated with either prior austenite grain boundaries or marteusitic lath boundaries [6]. PdAI formation causes a slight decrease in post-aged yield and tensile strengths because of its competition with Ni for the consumption of AI. A corresponding slight increase in ductility is observed with Pd, particularly at the 1% Pd level, due to this effect (Table 1). Grain and lath boundaries were analyzed for segregation of Pd [6]. Segregation of Pd to austenitic grain boundaries during solutionlzing is not anticipated as the microstructure is wholly austenltic and Pd is completely soluble. For the aged material, a Pd concentration at the grain boundary statistically indistinguishable from the matrix was found [6]. Susceptibility of wrought PH 13-8 Mo SS as a function of strength and Pd concentration The influence of Pd concentration on the mechanical properties of wrought H 13-8 Mo at three strength levels is shown in Figure 2. Here, HE susceptibility is expressed as the ratio of the maximum load attained in the CERT test to the maximum load obtained from CERT testing in air following the procedure of Hancock [15]. Similar trends exist in the case of susceptibility expressed as the ratios of times to failure; these plots are not shown. Ductile microvoid formation and growth dominated fracture is observed after tests in air at all strength levels. With charging, the highest strength conventional alloy (H925) exhibits the greatest susceptible to HE as expected [5]. A combination of intergranular separation on prior austenite grain boundaries and brittle transgranular fracture is observed for the hydrogen charged conventional alloys at all strength levels (Figure 3). In all cases, this brittle outer zone extends radially from the root of the notch and transitions gradually to microvoid coalescence characteristic of ductile overload in the center of the notched sample even though the ma,fimum longitudinal applied stress occurs at the center [14]. Concerning Pd additions, the Hll00 temper exhibits the lowest HE susceptibility over the entire range of Pd concentrations while the H925 temper exhibits the greatest. However, the Pd effect is still operative at highest strength temper although 0.39% Pd has only marginal benefit. Intergranular fracture is completely absent for Pd containing alloys at all strength levels (Figure 3). Instead, hydrogen assisted cracking is transgranular with an appearance that is consistent with inter-lath and tram-lath tearing and; possibly, cleavage. HE cracking also extends radially from the root of the notch and transitions to microvoid nucleation and growth dominated fracture in the center of the sample. Since the longest time to failure in these tests was 60-70 hours, hydrogen charging to obtain a uniform hydrogen concentration across the entire cross-section of the notched diameter was not obtained during the testing period even after the 30 h pre-charge. This conclusion is reached based on the low hydrogen diffusion coefficients for PH 13-8 Mo SS and calculated times required to obtain a uniform hydrogen concentration [6]. To verify that the Pd effect is still operative after long charging times to insure uniform hydrogen concentration at the highly stressed centerline of the notch an additional specimen (0.75% Pd at H1050) was cathodicaUy charged in 0.1 molar NaOH. This specimen was continually charged at -1130mA/cm 2, producing a hydrogen fugacity far exceeding the -1.1 V vs. SCE condition. Failure was not observed even after over 200 days at a constant load greater than the maximum load producing failure in conventional H1050 alloys by CERT testing. Suscentibilitv of welded PH 13-8 Mo as a function of Pd concentration Figure 2 also illustrates the effect of hydrogen on as-welded PH 13-8 Mo SS as a function of Pd. The weld is not as resistant to hydrogen as the wrought Hll00 alloy, which is of comparable hardness (Table 1). The welded dual phase microstructure exhibits greater susceptibility than the wrought marteusitic material at a similar strength level. Metallographic examination of longitudinal sections from fracture surfaces as well as SEM fractography indicate that cracking is transgranular across the ~ ferrite and marteusite phases and is not uniquely associated with the ferrite/marteusite interface. Discussion and Conclusions The addition of Pd to aged PH 13-8 Mo SS results in (a) formation of PdA1 precipitates, (b) decreased bulk hydrogen diffusion coefficients compared to the conventional alloy, (c) changes in hydrogen assisted fracture modes, and (d) improved HE resistance with increasing Pd concentration [6]. Three possible explanations are:
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(i) inhibited hydrogen ingress due to Pd induced catalysis of H z recombinative desorption, (//) Pd segregation to interfaces and subsequent repulsion of hydrogen, and (///) formation of a strong reversible trap (FdAI particles) that diminishes hydrogen segregation at other competing trap sites and retards occlusion within the stress field of the notch. Explanation (i) is dismissed because neither differences in exchange current densities, Tafel slopes, nor hydrogen production rates were observed with Pd [6]. Consequently, Pd is not expected to inhibit hydrogen ingress. In regard to mechanism (//), sitmificant segregation of Pd to metallurgical interfaces was not detected [6]. Moreover, Shirley and Hall predict that Pd only weakly repels hydrogen in iron [17]. Mechanism (///) is most consistent with the present study and is supported by trapping analysis of permeation data [6]. To summarize this view, permanently reduced apparent hydrogen diffusion coe~cients are obtained in aged Pd modified alloys, but not in solutionized Pd containing versions with Pd is solid solution. This supports the notion that PdA1 particles create a high density of traps that absorb hydrogen both irreversibly and reversibly. If such traps were irreversible and easily saturated, the steady state diffusion data produced from a series of permeation rise and decay transients would be the same for both the aged Pd modified alloy and the conventional alloy after the initial charging step saturated such irreversible traps [18]. This was not observed [6]. Operating on the assertion that PdAI particles are a strong reversible trap, it is worth reviewing the possible benefits. Recall that the criteria for optimum utilization of beneficial traps include (a) small in size so as not to serve themselves as fracture initiation sites, (b) high density and uniformly distributed, (c) revers~le and non-saturable [20-22]. An explanation involving beneficial trapping by PdA1 particles can account for all of the results of the present study. HE resistance improves with increasing Pd concentration and is not completely absent even at 0.38% Pd. This can be interpreted in connection with the benefits of an increased density of PdA] particles. Pd was more beneficial at lower yield strengths. This observation is consistent with the notion that the critical lattice hydrogen concentration required to trigger HE is inversely proportional to alloy strength [23]. Therefore, mobile hydrogen that reaches the fracture process zone is more "potent" with increasing strength [23]. Assnmlng a finite extent to which PdAI particles can trap hydrogen and retard its repartitionlng to the stress field of the notch, higher strength alloys would be more adversely affected by remaining mobile hydrogen. Finally, the PdA1 effect is diminished for the dual phase weld. This is attributed to the non-uniform distribution of PdA1 particles which are present in the 6 ferrite but not the martensite phase of the as-cooled weld metal (Figure 1). Other factors such as intrinsic susceptibility may affect the behavior of the dual phase weld metal as well.
This work was supported by the U.S. Department of Energy under contract number DE-AC04-76DP00789. The authors wish to acknowledge F. Bovard, G. Young, Jr. and D. Enos for various experimental contributions. We would also like to thank the staff of the Sandia National Laboratories Melting and Solidification Facility for preparing the alloys. References 1. C.T. Fujii, ASTM STP 610, H. Craig, ed., ASTM, Philadelphia, PA, 213, (1976). 2. A.W. Thompsoi~ MetaU. Trans., 4, 2819, (1973). 3. G.T. Murray, MetaH. Trans. A, 12(A), 2138, (1981). 4. G.T. Murray, H.H. Honegger, T. Mousel, Corrosion J., 40, 146, (1984). 5. P. Munn~ B. Andersson, Corrosion J., 46, 286, (1990). 6. J.R. Scully, J.A. Van Den Avyle, M.J. Cieslak, A.D. Romig, C.R. Hills, MetaH. Trans. A, 22A, 2429, (1991). 7. B. E. Wilde, C.D. Kim, and J.C. Turn Jr., Corrosion, 38, 515, (1982). 8. J.B. l~rnsden, B.E. Wilde, and P.J. Stocker, Scripta Met., 17, 971, (1983). 9. B.E. Wilde, I. Chattoraj, T.A. Mozhi, Scripta Met., 21, 1369, (1987). 10. M.IC Miller, S.S. Brenner, M.G. Burke, Met Trans A, 18A, 519, (1987). 11. T. D. I.e, and B.E. Wilde, In Current Solutions to Hydrogen Problems in Steels, Eds. C.G. Interrante, G.M. Pressouyre, ASM, Metals Park, Ohio, 413 (1982). 12. M.J. Cieslak, j,a~. Van Den Avyle, C.R. Hills, R.E. Semarge, Met Trans A, 19A, 3063 (1988).
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13. J.R. Seully, PJ. Moran, Corrosion, 44, 130, (1988). 14. P.W. Bridgman, Studies in Large Plastic Flow, McGraw-Hill, Ine, 9 (1952)• 15. J.W. Hancock, A.C. Mackenzie, J. Mech. Phys. Solids, 24, 147 (1976). 16. H.W. Picketing, M. Zamanzadeh, In Hydrogen Effects in Metals, Met. Soe. of AIME, I.M. Bernstein, A.W. Thompson, eds., 143, (1981). 17. A.I. Shirley, C.K. Hall, Seipta Met., 17, 1003, (1983). 18. G.M. Pressouryre, F.M. Faure., ASTM STP 962, L. Raymond, Ed., ASTM, philadelphia, 353 (1980). 19. R.A. Oriani, Acta Met., 18, 147, (1970). 20. G.M. Pressouyre and I.M. Bernstein, Acta Met., 27, 89, (1979). 21. G.M. Pressouyre, Met Tram A, 14A, 2189 (1983). 22. G.M. Pressouyre, Met Tram A, 10A, 1571 (1979). 23. K.N. Akhurst, TJ. Baker, Met Tram A, 12A, 1059, (1981). .
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130
HYDROGENEMBRJTFLEMENT
0% Pd
Vol. 31, No. 2
I% Pd
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B
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Figure 3. Fractography from CERT at -1.1V~ for: (a) H l l 0 0 at 0% (left) and 1% Pd (right), (b) H1050 at 0% (left) and 1% Pd (right), and (c) H925 at 0% 0eft) and 1% I'd (right). Transgranular fracture was observed for weld metal. Microvoid fracture was observed in air CERT for all tempers and weld metal. Micron markers are indicated on each micrograph.