improved resistance towards hydrogen induced … · duplex stainless steels (dss) are widely used...

IMPROVED RESISTANCE TOWARDS HYDROGEN INDUCED STRESS CRACKING (HISC) OF HOT ISOSTATICALLY PPRESSED (HIP) DUPLEX

STAINLESS STEELS UNDER CATHODIC PROTECTION

Gro Østensen Lauvstad SINTEF

Richard Birkelandsvei 3A NO-7465 Trondheim

Norway gro.o.lauvstad@sintef.no

Roy Johnsen Martin Bjurström and Carl-Gustav Hjorth Bård Nyhus

NTNU/SINTEF Metso Powdermet AB SINTEF Richard Birkelandsvei 2B Kontorsvägen 1, P.O. Box 54 Richard Birkelandsvei 3A

NO-7491 Trondheim, SE-735 21 Surahammar NO-7465 Trondheim Norway Sweden Norway

ABSTRACT Hot isostatically pressed (HIP) and forged duplex stainless steels are both candidate materials for subsea components. In the present study, the resistance towards hydrogen induced stress cracking (HISC) of these materials has been compared. Testing was performed by four-point bending of fatigue pre-cracked SENB (single edge notch bend) specimens polarised to -1050 mV vs. SCE in a 3.5 wt% NaCl solution at 4 °C. The specimens were loaded to different crack tip opening displacement (CTOD) values for a maximum duration of 30 days. Three HIP materials, one duplex steel (DSS, UNS S31803) and two super duplex steels (SDSS, UNS S32550 and UNS S32760) were investigated and compared to a forged SDSS material (UNS S32760). For the HIP materials no cracking was detected for specimens loaded at CTOD levels lower than 0.06 mm, while some slight cracking was observed for a CTOD of 0.08 mm. The CTOD threshold value for HISC thus appears to be in the range from 0.06 to 0.08 mm for these materials. The forged SDSS material on the other hand showed significant cracking at CTOD levels above 0.03 mm, while no cracking was observed at CTOD levels below 0.016 mm. The testing thus demonstrates the improved resistance towards HISC for the HIP materials. Key words: duplex stainless steel, hydrogen embrittlement, cathodic protection, microstructure

INTRODUCTION Duplex stainless steels (DSS) are widely used in the Oil and Gas industry due to their excellent mechanical properties as well as their corrosion resistance in chloride/hydrocarbon environments. The susceptibility of DSS to hydrogen induced stress cracking (HISC) when exposed to loading under cathodic protection (CP) is of primary concern1,2. HISC may be regarded as the result of interplay between hydrogen embrittlement and stress/deformation. For duplex materials additional complexity is introduced by the different strength, ductility and microstructure of the ferrite and the austenite phases constituting the material, as well as their different susceptibilities to hydrogen embrittlement3. The susceptibility of an alloy to HISC is strongly influenced by its hydrogen ingress characteristics, which may be described by two parameters, the efficiency of hydrogen entry and the efficiency of hydrogen trapping4. Under CP, the entry efficiency represents the amount of hydrogen entering the alloy relative to the total amount formed cathodically on the alloy surface. Some of the hydrogen entering the alloy is trapped at structural heterogeneities, either reversibly (i.e. the trap can release hydrogen) or irreversibly (i.e. the hydrogen release rate can be considered as negligible). The modeling of the hydrogen transport in DSS must take into account the different mobility of hydrogen in the ferrite (α) and austenite (γ) phases1,5: The ferrite (bcc structure) is characterized by a high diffusion rate and a low solubility of hydrogen, while the austenite (fcc-structure) exhibits a lower diffusion rate and a higher solubility. It is thus commonly assumed that the transport of hydrogen through a DSS occurs mainly through the ferrite phase6. Hydrogen may be assumed to occur in two forms in the duplex matrix, (i) as diffusible hydrogen absorbed in the ferrite phase and (ii) as hydrogen reversibly trapped by microstructural defects and/or at the austenite grains6. Consequently, the microstructure of a DSS influences the susceptibility towards HISC, in terms of (i) the ferrite content, higher ferrite content favoring HISC, (ii) austenite spacing, smaller austenite spacing resulting in a more resistant material, and (iii) ferritic/austenitic grain size, smaller grain size resulting in improved resistance towards HISC1: The following ranking from more susceptible to less susceptible is observed to apply for HISC: forgings > rolled plates > hot isostatically pressed (HIP). The objective of the current work was to investigate further the effect of microstructure on the resistance towards HISC by comparing HIP DSS and super DSS (SDSS) to a forged SDSS material.

EXPERIMENTAL Materials Three HIP materials delivered by Metso Powdermet AB were tested: • SDSS 33086-P3 APM2327 – UNS S32550 (ID P3) • SDSS 32795-70 APM2379 – UNS S32760 (ID 32) • DSS 33092-03 APM2377 – UNS S31803 (ID 33) These materials were compared with a previously tested forged SDSS (UNS S32760). The chemical composition of the materials is given in Table 1. Table 2 summarizes the parameters characterizing the microstructure of the materials, the volume fraction of ferrite (%) and the austenite spacing. The latter parameter was assessed according to the procedure described in Appendix B of DNV RP-F112 7. Figure 1(a) and (b) illustrate the different microstructures of the HIP and the forged materials, respectively.

SENB Testing The testing was performed with single edge notch bend (SENB) samples at different crack tip opening displacements (CTOD), a geometry that has also been used in previous HISC testing8. The dimensions of the test pieces are shown in Figure 2(a). The test environment used was 3.5% NaCl dissolved in distilled water, 4 °C at atmospheric conditions in an open vessel, and a potential of -1050 mV versus SCE applied potentiostatically. The samples were machined from the respective materials and a notch was introduced with spark eroding so that clip gauges could be mounted. In the bottom of the notch a fatigue pre-crack was made on each sample. The target total initial crack depth, a0, was 4 mm, giving an initial a/w ratio of 0.45. The HIP samples were loaded to target CTOD levels of 0.015, 0.035, 0.06 and 0.08 mm. The clip-gauge was used to measure the increase in Crack Mouth Opening Displacement (CMOD) as the loading (deflection) was slowly increased. In order to obtain the correct CTOD levels, a CMOD vs. CTOD relationship from previous CTOD testing was applied (see Figure 2(b)). Post-test Examinations The testing of the HIP materials was terminated after 30 days and the samples were unloaded. Half of the samples were heat tinted and broken up after cooling in liquid nitrogen. The heat tinting was done in order to separate the cracking which had occurred during testing from the residual fracture resulting from final breaking of the samples. The fracture surfaces were inspected in a stereomicroscope as well as by scanning electron microscopy (SEM, Hitachi 3500). The rest of the samples were prepared for metallographic examination of sections taken in the longitudinal specimen direction, i.e. normal to the crack plane. One specimen side edge (S) and one mid-section (M) were examined. The samples where etched in 40% NaOH (20 V, 10-20 sec.) and examined in an MF3 optical microscope. After exposure, samples of the HIP material were analyzed for bulk hydrogen concentration by the melt extraction technique (JUWE H-mat 225). The samples of the forged material were loaded to CTOD-values between 0.014 and 0.101 mm, and were unloaded whenever cracking was observed by visual examination of the specimen side edges. All of the samples were then heat tinted and broken up after cooling in liquid nitrogen.

RESULTS AND DISCUSSION Comparison of SENB results The results of the SENB testing at different CTOD levels for the three HIP materials are summarized in Table 3, Table 4 and Table 5, while the results obtained with forged SDSS (UNS S32760) are summarized in Table 6. The results obtained for the HIP materials can be summarized as follows: • From the inspection of the metallographic sections, no cracking was observed for any of the HIP

materials loaded to CTOD values below 0.08 mm; at CTOD 0.08 mm the maximum crack depth measured was 0.16 mm for the SDSS UNS S32760 (see Figure 3), while the other two materials showed evidence of micro-cracks 0.02 mm long (see Figure 4, SDSS 32550).

• In only one case could a positive identification of additional cracking by analysis of the fracture surface be made, on a sample of SDSS UNS S32550 loaded to a CTOD of 0.08 mm (see Figure 5).

• Generally, an additional area of transgranular cracking of the ferrite phase only was observed ahead of the crack tip.

The forged material on the other hand, showed evident cracking after 21 days at a CTOD level of 0.033 mm. When comparing the results for the HIP and forged materials it should be noted that constant deflection in bending is maintained only as long as the crack does not propagate into the material. Accordingly, no quantitative comparison of crack propagation rates for the materials tested is possible. It should also be noted that when performing constant bend testing of DSS, the effect of cold creep should be taken into account1,2: Cold creep will lead to a redistribution of load at the tip of the fatigue pre-crack, and accordingly the reported CTOD-values should be interpreted as “initial” values, i.e. valid before the creep process has started. On the other hand may the cold creep load redistribution also result in cracking of the surface oxides/scales and may thus facilitate the hydrogen pick-up process and actually promote the HISC process. The load relaxation in a SENB specimen caused by cold creep and crack growth may, however, not be representative for the actual conditions experienced by a subsea component. Even considering these aspects of the SENB testing, the very significant cracking observed in the forged material after 21 days of testing at CTOD 0.03 mm as opposed to the very limited cracking observed in the HIP materials after 30 days testing at CTOD 0.08 mm, demonstrates a marked increase in the resistance towards HISC for the HIP material. Effects of bulk hydrogen concentration The bulk hydrogen concentration of the HIP samples after exposure was measured to 8±3 ppm (SDSS S32550), 10±5 ppm (SDSS S32760) and 8±2 ppm (SDS S31803). For comparison, the forged SDSS S32760 material had acquired a bulk concentration of 11±2 ppm after 28 days of exposure under the same conditions. The difference between the materials in hydrogen concentration thus appears to be negligible. The melt extraction technique does, however, measure the average hydrogen content in the bulk of the sample analyzed, i.e. it does not give information of the local distribution of hydrogen in the austenite and ferrite phases. Also, the contributions from both trapped and diffusible hydrogen are measured by this technique. Effect of microstructure The effect of microstructure on the resistance towards HISC has been attributed to grain boundaries acting as trap sites for hydrogen: In the finer grained materials more hydrogen may be trapped in the grain boundaries and accordingly less diffusible hydrogen is available during the cracking process and the time to failure in a slow strain rate test is increased9. The effective diffusion coefficient for hydrogen, Deff, measured for two DSS materials (SAF2205) was ~30% higher for the more coarse grained material (ferrite unit size 13.9 µm) than for the finer grained material (ferrite unit size 7.6 µm). Due to the low diffusivity of hydrogen in austenite it can, however, be argued that the austenite in itself can act as an irreversible trap5. The amount of austenite in the structure may on the other hand act as a hydrogen reservoir, the hydrogen being delivered by dislocations to the interface with ferrite during straining resulting in induced cracking and consequential loss in ductility10. In the present study the average austenite spacing of the materials tested ranges from typically 14 µm for the HIP materials to 49 µm for the forged material. The results of the SENB testing thus confirm the effect of austenite spacing on the resistance towards HISC.

Results of fracture surface examinations Examination of the fracture surface of sample P3-8 (HIP SDSS 32760 loaded to a target CTOD of 0.08 mm) by SEM revealed a change in the appearance of the fracture surface corresponding to the heat-tinted area indicated in Figure 5, as indicated in Figure 6(a). Figure 6(b) gives a detail from the area next to the fatigue pre-crack, showing that the test induced cracking occurred by transgranular cleavage fracture of the ferrite phase followed by microvoid coalescence of the austenite phase11. Evidence of secondary cracking is also observed. This is consistent with the observations of transgranular cleavage of ferrite grains ahead of the crack tip with intact austenite grains in between, as shown in Figure 3 and Figure 4. Figure 6(c) gives a detail of the residual fracture resulting from the final breaking of the sample. The fracture surface here shows evidence of transgranular fracture of both the ferrite and the austenite phases.

CONCLUSIONS Generally, no cracking was observed for the HIP-specimens loaded to CTOD levels below 0.08 mm. Only some limited cracking was observed at the maximum CTOD of 0.08 mm, with the maximum crack length measured being 0.16 mm (SDSS S32760) after 30 days of testing. The CTOD threshold value for HISC thus appears to be in the range from 0.06 to 0.08 mm for the HIP materials tested. The forged SDSS material (UNS 32760) on the other hand showed marked cracking at CTOD levels above 0.03 mm, the measured extent of cracking after 13 days being approximately 2 mm. The SENB testing, although rather qualitative in nature, demonstrates a marked increase in the resistance towards HISC for HIP DSS and SDSS as compared to a forged SDSS. The average austenite spacing of the HIP materials was typically 14 µm, while for the forged SDSS an austenite spacing of 49 µm was determined. The effect may be attributed to grain boundaries acting as trap sites for hydrogen: In the finer grained materials more hydrogen may be trapped in the grain boundaries and accordingly less diffusible hydrogen is available in the grain interiors during the cracking process. The difference in resistance towards HISC is not reflected in the total bulk concentration of hydrogen measured for the different materials, the HIP and forged materials showing quite similar bulk hydrogen concentration values.

ACKNOWLEDGEMENTS Statoil and Norsk Hydro are gratefully acknowledged for the permission to publish the results obtained on the forged SDSS material.

REFERENCES /1/ T. Cassagne, F. Busschaert, “A review on hydrogen embrittlement of duplex stainless steels”

NACE CORROSION 2005, Paper no. 05098 (NACE, Houston, TX, 2005) /2/ S. Huizinga, B. McLoughlin, I.M. Hannah, S.J. Paterson and B.N.W. Snedden, “Fracture of a

subsea super duplex manifold by HISC and implications for design”, NACE CORROSION 2006, Paper no. 06145 (NACE, Houston, TX, 2006)

/3/ A. A. El-Yazgi, D. Hardie, “Stress corrosion cracking of duplex and super duplex stainless steels

in sour environments”, Corros. Sci., 40 (1998) 909 /4/ B. G. Pound, “Potentiostatic pulse technique to determine the efficiency of hydrogen absorption

by alloys”, J. Appl. Electrochem., 21 (1991) 967 /5/ V. Olden, C. Thaulow, R. Johnsen, Hydrogen diffusion and cracking in supermartensittic, duplex

and super duplex stainless steel – A state of the art review”, Theor. Appl. Fract. Mech., In progress

/6/ E. Owczarek, T. Zakroczymski, “Hydrogen transport in a duplex stainless steel”, Acta Mater., 48

(2000) 3059 /7/ Recommended Practice DNV-RP-F112, “Design of stainless steel subsea equipment exposed

to cathodic protection”, Draft issue, April 2006 /8/ A. Mikkelsen, S. Wästberg, R. Johnsen, B. Nyhus and T. Rogne, “Influence of ambient pressure

on Hydrogen Induced Stress Cracking (HISC) of Duplex Stainless Steels under Cathodic Protection”, NACE CORROSION 2006, Paper no. 06499 (NACE, Houston, TX, 2006)

/9/ S.L. Chou, W.-T. Tsai, “The effect of grain size on the hydrogen-assisted cracking in duplex

stainless steels”, Mat. Sci. Eng., A270 (1999), 219 /10/ A. A. El-Yazgi and D. Hardie, “Effect of heat treatment on susceptibility of duplex stainless

steels to embrittlement by hydrogen”, Mater. Sci. Tech. Ser., 16 (2000) 506 /11/ S. Roychowdhury and V. Kain, “Environmental effects on the fracture toughness of a duplex

stainless steel (UNS-S31803)”, NACE CORROSION 2006, Paper no. 06492 (NACE, Houston, TX, 2006)

TABLE 1

CHEMICAL COMPOSITION OF THE TESTED MATERIALS (%): MATERIALS CERTIFICATE VALUES

MATERIAL C Mn Si P S Cr Ni Mo Cu N W

HIP 22Cr DSS S31803 0.021 1.09 0.74 0.018 0.002 22.5 5.2 2.96 - 0.20 -

HIP 25Cr SDSS S32550 0.026 0.60 0.46 0.019 0.002 26.0 6.2 3.00 1.8 0.29 -

HIP 25Cr SDSS S32760 0.023 0.60 0.40 0.023 0.003 25.6 6.81 3.50 0.66 0.27 0.67

Forged 25Cr SDSS S32760 0.024 1.21 0.26 - 0.002 24.0 7.3 4.0 - - -

TABLE 2 MICTROSTRUCTURAL CHARACTERIZATION OF TESTED MATERIALS:

FERRITE CONTENT AND AUSTENITE SPACING

FERRITE AUSTENITE SPACING MATERIAL

CONTENT NO. OF MEAS.

AVERAGE MAX. MIN.

HIP DSS (UNS S31803) 47% 218 13.8 ± 8.2 µm 50.5 µm 1.4 µm

HIP SDSS (UNS S32550) 43% 201 14.6 ± 9.4 µm 65.3 µm 2.6 µm

HIP SDSS (UNS S32760) 48% 202 12.7 ± 9.4 µm 62.1 µm 1.1 µm

235 47.9 ± 43.7 µm 215.4 µm 3.8 µmForged SDSS (UNS S32760) -

238 50.9 ± 44.1 µm 239.0 µm 6.7 µm

TABLE 3

SENB TESTING OF HIP SDSS UNS S32550 (33086-P3 APM2327): FRACTURE SURFACE ANALYSIS/ METALLOGRAPHIC EXAMINATION (SIDE EDGE (S)/MID-SECTION (M)). ∆a= CRACK

EXTENSION DUE TO TESTING (CONTINUOUS), a0=INITIAL CRACK LENGTH

ID CTOD, mm

FRACTURE SURFACE

MET EXAM

RESULT ∆a, mm a0, mm

P3-4 0.015 X - 4.67

P3-5 0.015 X - 4.38

P3-1 0.035 X - 4.15

S P3-6 0.038 -

M

4.25

S P3-7 0.064 -

M

4.15

S

No cracking

0

P3-3 0.084 -

M Slight cracking 0.02

4.36

P3-8 0.085 X - Cracking 0.08 4.25

TABLE 4 SENB TESTING OF HIP SDSS UNS S32760 (32795-70 APM2379): FRACTURE SURFACE

ANALYSIS/ METALLOGRAPHIC EXAMINATION (SIDE EDGE (S)/MID-SECTION (M)). ∆a= CRACK EXTENSION DUE TO TESTING (CONTINUOUS), a0=INITIAL CRACK LENGTH

ID CTOD,

mm FRACTURE SURFACE

MET EXAM

RESULT ∆a, mm a0, mm

32-1 0.015 X - 4.12

32-2 0.016 X - 4.24

32-4 0.036 X - 4.16

S 32-5 0.032 -

M

3.80

32-8 0.066 X 4.31

S 32-7 0.067 -

M

No cracking

0

4.37

S 0.02 32-3 0.084 -

M

Cracking

0.16

4.28

32-6 0.081 X - No cracking 0 4.24

TABLE 5 SENB TESTING OF HIP DSS UNS S31803 (33092-03 APM2377): FRACTURE SURFACE ANALYSIS/ METALLOGRAPHIC EXAMINATION (SIDE EDGE (S)/MID-SECTION (M)).∆a= CRACK EXTENSION

DUE TO TESTING (CONTINUOUS), a0=INITIAL CRACK LENGTH

ID CTOD, mm

FRACTURE SURFACE

MET EXAM

RESULT ∆a, mm a0, mm

33-2 0.015 X - 4.03

33-7 0.015 X - 4.38

33-3 0.036 X - 4.32

S 33-6 0.035 -

M

4.41

S 33-8 0.060 -

M

4.26

33-1 0.077 X - 4.01

S

No cracking

0

33-4 0.080 -

M Slight cracking 0.02

4.41

TABLE 6

SENB TESTING OF FORGED SDSS UNS S32760: FRACTURE SURFACE ANALYSIS. ∆a= CRACK EXTENSION DUE TO TESTING, a0=INITIAL CRACK LENGTH

ID CTOD, mm

RESULT ∆a, mm a0, mm TEST TIME, DAYS

S-12 0.014* - 4.1 40 S-10 0.016*

No cracking

- 4.1 S-9 0.033 1.6 3.9

S-8 0.032 1.7 3.8

21

S-1 0.044 1.9 4.1

S-5 0.053 2.8 3.7

S-2 0.089 2.1 3.9

S-6 0.101

Cracking

2.1 3.9

13

* Cycling unloading/loading (2 times).

50 µm 50 µm

FIGURE 1 - Microstructure of materials: (a) HIP SDSS 32760 and (b) forged SDSS 32760 (M X200).

(a) (b)

w=9 mm

15 mm120 mm

a0=4 mmw=9 mm

15 mm120 mm

a0=4 mm

FIGURE 2 - SENB sample dimensions and notch/fatigue pre-crack geometry (a), relationship

between measured CMOD and CTOD (b).

100 µm FIGURE 3 - Specimen 32-3(M) SDSS UNS S32760 CTOD=0.08 (Mx500): Crack extends 160 µm

from end of fatigue pre-crack, region of cracked ferrite phase extends 280 µm from end of fatigue pre-crack.

IGURE 4 - Specimen P3-3(M) SDSS UNS S32550 CTOD=0.08 (Mx500): Micro-crack extends

100 µm100 µm F

20 µm from end of fatigue pre-crack, region of cracked ferrite phase extends 60 µm from end of fatigue pre-crack.

2 mm

M x100.5 mm

M x36

Fatigue pre-crack

Spark-eroded notch

Seat for clip gauge

Fatigue pre-crack

Residual fracture

Cracking duringtesting

FIGURE 5 - Fracture surface specimen P3-8 (SDSS UNS S32550) after testing: Heat tinted area

ahead of fatigue pre-crack indicates additional cracking during testing – detail shown in inset.

250 µm250 µm

50 µm50 µm

50 µm

20 µm20 µm

(a)

(b)

(c)

Fatigue pre-crack FIGURE 6 - SEM analysis of fracture surface of specimen P3-8 (SDSS UNS S32550) after testing:

(a) Area corresponding to heat-tinted area in Figure 5. (b) Detail of (a), area next to fatigue pre-crack fracture surface showing mixture of intergranular fracture and cleavage fracture of ferrite phase and microvoid coalescence of austenite phase, as well as of secondary cracking. (c) Detail of (a), residual fracture from final breaking of sample showing transgranular fracture.