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CORROSION SCIENCE SECTION

CORROSIONVol. 62, No. 5 375

Submitted for publication December 2004; in revised form,August 2005.

Corresponding author. E-mail: [email protected].* UNAM, Facultad de Qumica Circuito Interior, C.U., Edificio B,

C.P. 04510, Mxico, D.F. Mexico.** Instituto Mexicano del Petrleo, Programa de Investigacin y

Desarrollo de Ductos, Eje Central Lzaro Crdenas #152, SanBartolo Atepehuacan, 07730, Mxico, D.F. Mexico.

*** UAEM-CIICAP, Av. Universidad 1001, Col. Chamilpa, 6225-Cuernavaca, Mor., Mexico.

Sulfide Stress Cracking Susceptibilityof Welded X-60 and X-65 Pipeline Steels

C. Natividad,* M. Salazar,** A. Contreras,** A. Albiter,** R. Prez,** and J.G. Gonzalez-Rodriguez,***

ABSTRACT

The susceptibility of API X-60 and X-65 longitudinal weld

beads to sulfide stress cracking (SSC) has been evaluated

using slow strain rate tests (SSRT) in the NACE solution

saturated with hydrogen sulfide (H2S). The tests were supple-

mented by potentiodynamic polarization curves and hydrogen

permeation measurements. The weld beads were produced

using the submerged arc welding (SAW) process. Three differ-

ent temperatures were used: room temperature (25C), 37C,

and 50C. The corrosion rate, taken as the corrosion current

density, Icorr, the amount of hydrogen uptake for the weld-

ments, C0, and the SSC susceptibility increased with an in-

crease in the temperature from 25C to 50C. Although anodic

dissolution seems to play an important role in the cracking

mechanism, the most likely mechanism for the cracking sus-

ceptibility of X-60 and X-65 weldments in H2S solutions seems

to be hydrogen embrittlement.

KEY WORDS: hydrogen embrittlement, sulfide stress corrosion

cracking, X-60 and X-65 steel weldments

INTRODUCTION

The oil industry contains a wide variety of corrosive

environments. Mexican crude oil and gas commonly

contain entrained water, carbon dioxide (CO2), and

hydrogen sulfide (H2S). The transport of these types

of products always induces failures in the pipeline

systems, and not less frequently in the weld beads.

The welding industry has recognized that weld-

induced stresses play an important role in certain

localized corrosion phenomena. Each year, tens of

million of dollars are expended to replace or repairpipes and vessels that suffer excessive localized metal

loss, stress corrosion cracking (SCC), or hydrogen

embrittlement (HE). When sulfide is present, this type

of brittle failure is known as sulfide stress cracking

(SSC), and it has been established as a particular case

of HE.

Weld metal corrosion is normally attributed to

the differences in composition and to differences in

electrochemical potentials between the parent metal,

heat-affected zone (HAZ), and weld metal. A lower

electrochemical potential of the weld bead is com-

monly related to the composition, microstructure,

and distribution of inclusions.1

The suitable sour service materials are listed in

NACE MR0175,2 whereas the TM01773 lists Method

A as one of the tests to be performed to determine

the inclusion of a material in MR0175. Many high-

strength, low-alloy steels are precluded from under-

going this test, especially in the as-welded condition,

due to either high parent material hardness or the

formation of high-localized stressed weld regions in

the weld HAZ. Specifications for welded sample test-

ing are not addressed in the NACE standard primarily

because this test is geared toward wrought samples.

Welded samples differ from homogeneous samplesbecause of their local variations in microstructure and

0010-9312/06/000071/$5.00+$0.50/0 2006, NACE International

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376 CORROSIONMAY 2006

compositions. The development of multi-phase micro-

structures is important for the attainment of certain

mechanical properties, but it can be detrimental for

resistance to SSC. Carbon-rich phases such as pearl-

ite, bainite, or martensite can be particularly suscep-

tible to this mode of HE.4-7

The goal of this study was to evaluate the sus-

ceptibility to SCC in H2S-containing environments of

weld beads obtained from API X-60 and X-65 pipeline

steels produced by the submerged arc welding (SAW)

techniques using the slow strain rate testing (SSRT)

technique, although the Method A test in TM0177 isnot an SSRT test but a constant load tensile test.

EXPERIMENTAL PROCEDURES

In this work, longitudinal weld beads of API X-60

and X-65 pipelines steels were analyzed. These pipes

are typically used in the Mexican pipeline systems.

They have an outside diameter of 42 in. (1,066 mm)

(X-60) and 24 in. (609.6 mm) (X-65), with a wall

thickness of 0.5 in. (12.7 mm) (X-60) and 0.562 in.

(14.27 mm) (X-65). The chemical composition of these

steels is shown in Table 1. Table 2 shows the weld-

ing parameters used to join the pipeline steels. A filler

wire electrode with a chemical composition consisting

of 0.08% C, 1.1% Mn, and 0.6% Si was used.

Cylindrical tensile specimens with a 25.00-mm

gauge length and a 2.50-mm gauge diameter were

machined from an unused pipeline perpendicularto the rolling direction, as shown in Figure 1, which

shows a schematic illustration of the orientation and

location of the sample with respect to the longitudi-

nal seam weld. Before testing, the specimens were

abraded longitudinally with 600-grade emery paper,

degreased, and masked, with the exception of the

gauge length. The masking agent used was an inert

resin and it has been observed that it does not induce

crevice-type corrosion at the end of the test. Speci-

mens were subjected to conventional, monotonic,

SSRT testing in air, as an inert environment, and the

standard NACE solution (5% sodium chloride [NaCl],0.5% acetic acid [CH3COOH], and saturated with hy-

drogen sulfide [H2S]) at a strain rate of 1.00 106 s1

at room temperature, (25C), 37C, and 50C. All the

tests were performed at the open-circuit potential.

According to the authors experience, stress/strain

curves give some conflicting results; instead of this,

TABLE 1

Chemical Compositions of X-60 and X-65 Steels (wt%)

Steel C Mn Si P S Al Nb Cu Cr Ni Mo V Ti

X-60 0.025 1.57 0.14 0.012 0.002 0.044 0.097 0.31 0.29 0.17 0.03 0.002 0.014X-65 0.04 1.48 0.25 0.012 0.002 0.041 0.047 0.09 0.02 0.5 0.069 0.017

TABLE 2

Welding Parameters

Parameter Value

Current 320 AVoltage 20 VTravel speed 3.6 mm/sArgon flow rate 22 L/minPreheating temperature 250CHeat input 17.11 KJ/cmStick out 12 mm

(a) (b)

FIGURE 1. Schematic illustration of the orientation and location of the sample with respect to the longitudinal seam weld.

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the loss in ductility was assessed in terms of the per-

centage reduction in area (%RA) using:

%

RA

A A

A

i f

i

= 100

(1)

where Aiand A

fare the initial and final areas, respec-

tively. A susceptibility index to SCC (ISCC) was calcu-

lated as follows:

I

RA RA

RASCC

AIR NACE

AIR

=% %

% (2)

where %RAAIRand %RANACE are the percentage reduc-

tion in area values in air and in the H2S-saturated

NACE solution, respectively, previously deaerated with

nitrogen gas. The fracture surfaces were then exam-

ined using scanning electron microscopy (SEM). Po-

tentiodynamic polarization curves were performed at

a sweep rate of 1.0 mV/s using a fully automated po-tentiostat controlled with a desktop computer. These

tests were done in duplicate.

Hydrogen permeation tests were carried out using

the two-component Devanathan-Stachurski8 cell. The

specimen, which was fabricated from the same pipe

sample as the tensile specimens, was mounted be-

tween the two compartments, giving an effective area

of 3.14 cm2 exposed to the H2S-containing solution,

under open-circuit conditions to generate hydrogen.

The hydrogen collection compartment contained an

electrolyte of 0.5 M sodium hydroxide (NaOH) solution

purged with nitrogen (N2) gas. The Pd-plated specimensurface exposed to this solution was potentiostatically

passivated at a constant potential of 300 mV vs. a

saturated calomel electrode (SCE), which was used as

reference electrode in the polarization curves also. In

both cases, a graphite rod was used as an auxiliary

electrode. Potentiodynamic polarization curves were

performed once the free corrosion potential value

(Ecorr) was stable, i.e., it did not change more than

2 mV/min. The scanning started at 500 mV, with

respect to Ecorr, and finished at 300 mV more positive

than Ecorr, at a scanning rate of 1 mV/s. Corrosion

current values (Icorr) were calculated using Tafel ex-

trapolation. When the passive current reached aconstant value, then the H2S-containing solution

was poured into the other compartment to start the

hydrogen permeation experiments. Polarization curves

were repeated at least three times when there was

no more than 10 mV to 15 mV in the Ecorrvalues and

the difference in current densities was within an error

of 5%. Temperatures were kept constant by using a

heating tape.

The hydrogen coefficient diffusion (D) was calcu-

lated using:

D L

tlag=

2

6 (3)

where L is the specimen thickness (0.7 mm) and

tlag, the lag time, is the time elapsed when 63% of

the steady-state permeation current, Jss, has been

reached. The number of Pd-coated hydrogen atoms at

the entrance surface, C0, was calculated using:

JDC

Lss =

0

(4)

RESULTS AND DISCUSSION

Figure 2 shows micrographs of the X-60 and

X-65 steels weldments. These figures clearly show the

different microstructures found in a weldment, which

itself consists mainly of polygonal and coarse acicular

ferrite. This microstructure optimizes the strength

and the toughness of the weld beads.9-12 The results of

the hardness measurements obtained from the differ-ent zones of the weld bead are presented in Figure 3.

(a)

(b)

FIGURE 2. Micrographs obtained by optical microscopy of the

weldments: (a) X-65 and (b) X-60 steel.

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The hardness was measured at the middle of the weld

bead (Figures 3[a] and [c]), known as the reheated

zone. It was also measured at the middle of the upper

weld bead. Remarkable differences can be observed

between both measurements. A considerable decrease

in hardness can be observed in the reheated zone.

The hardness of ferritic-perlitic steels increases by

increasing the perlite content. The obtained hardness

values are within the recommended limits for avoiding

cracking and fracture in the weld bead.

Figures 4 and 5 show the effect of temperature

on the polarization curves for X-60 and X-65 pipeline

steels, respectively. As expected in these solutions,

there is no passive region in any of the cases, only

active dissolution. For both steels, the Ecorrdecreases

as the temperature is increased, with the most noble

value at 25C and 600 mV, and the most active at

50C and around 800 mV. The Ecorrvalue for the

X-60 steel at 37C was 650 mV whereas for the X-65

steel, it was 700 mV. For the X-60 steel, the anodic

current density and the Icorrvalue increased as thetemperature increased, but not for the X-65 steel.

Figure 6 shows two typical hydrogen permeation

current transients obtained at 25C for both steels.

It can be seen that the hydrogen permeation current

is higher in the X-60-type steel than in the X-65 one

by almost five times, and the steady state is reached

faster in the former than in the later. This figure

shows some current-time transients, especially for the

X-65 steel, which are consistent with diffusional pro-

cesses. Figure 7 shows the effect of the temperature

on the hydrogen uptake for both steels. It is clear that

the amount of hydrogen uptake increases with tem-

perature for both steels, being always higher in the

X-60 than in the X-65-type steel. All these results are

consistent with those found in the literature,12 as will

be seen later.

The susceptibility toward SSC was measured with

the ISCC index, and the results are plotted in Figure 8.

Values close to the unit mean that the steel is highly

susceptible to SSC, whereas values close to zero mean

that the steel is immune to SSC. Thus, Figure 8

clearly shows that, in all cases, regardless of the tem-perature, both steels are highly susceptible to SCC,

(c) (d)

(a) (b)

FIGURE 3. Hardness measurements in weld beads: (a) and (b) X-60 steel and (c) and (d) X-65 steel.

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and the effect of the temperature is negligible, al-

though the tendency is that this susceptibility in-

creases with increasing temperature. X-60 pipeline

steel was more susceptible toward SCC than the X-65

steel, although this difference seems to be negligible.

An analysis of the fracture surfaces made by SEM

revealed corroded surfaces in all the specimens, as

shown in Figure 9, which show the fracture surface

of the X-65 steel strained at 37C. However, even in

conditions where the corrosion rate was the highest,

i.e., at 50C, some quasi-cleavage features, arrowed,

similar to those produced by HE, were found on the

fracture surface of the X-60 steel (Figure 10). A

surface-initiated crack propagated through a brittle

fracture perpendicular to the applied stress until it

reached a critical size where ductile rupture thenoccurred, usually at 45 to the tensile axis. This gen-

FIGURE 5. Effect of temperature on the polarization curves in the

saturated NACE solution with H2S for X-65 steel weldments.

FIGURE 4. Effect of temperature on the polarization curves in the

saturated NACE solution with H2S for X-60 steel weldments.

FIGURE 6. Typical time variation of hydrogen permeation current

in the saturated NACE solution with H2S for X-60 and X-65 steel

weldments at 25C.

FIGURE 7. Effect of temperature on the C0values for X-60 and X-65

steel weldments.

eral fracture morphology was exhibited regardless

of whether the fracture initiated in the base metal,

fusion zone, or HAZ. Post-fracture exposure of the

fracture surface, in most cases, caused the formation

of a corrosion product, making the determination of

the fracture mode at times difficult or impossible in

the SEM, even when the specimens were cleaned

according to ASTM G 1.13 A chemical analysis of the

inclusions found in the fracture surfaces of both

steels was performed by energy-dispersion of x-rays

and are given in Figure 11. The chemical composition

of the inclusions found in the base metal were differ-

ent for each steel, as can be seen from Figure 11,

although they were similar in shape (globular) and

size (7 m to 10 m). These inclusions are composed

mainly of Fe, Mn, S, Ca, Si, Al, C, O, which can form(Fe, Mn)S, CaO, MnS-Al2O3 mixtures, Ca-Al-Si-Mg-O

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globular inclusions, or Fe-Mn-S-O inclusions. These

differences, mainly in chemical composition, can

explain the differences observed in both hydrogen

permeation current density (Figure 6) and SCC sus-

ceptibility (Figure 8).

Metallographic cross sections of X-60 steel frac-

tured at 25C and 50C are shown in Figures 12

and 13, respectively. It seems that the cracks were

predominantly transgranular in nature, but this

could have been better shown with an etched speci-

men. Just as indicated by the polarization curve,

that the corrosion rate increased as the temperature

increased, the amount of corrosion products inside

the cracks is more pronounced at 50C than at 25C,

and the crack length was longer at 50C than at 25C,

although this might be random. In any case, thesepictures show the important role of anodic dissolu-

tion in the cracking process. For X-65 steel, no cracks

were observed, only pits, just as shown in Figure 14,

possibly because the cracks were polished away. It

should be noted that the term pitting refers to the

localized breakdown of a thin, protective passive film.

Such films do not form on carbon steels at low po-

tentials in acidic sour environments, so the observed

attack might be better described as localized corrosion

that occurs at a break in the somewhat protective iron

sulfide film.

The requirements for SSC based on the HE

mechanism include a susceptible microstructure, athreshold level of hydrogen to induce cracking, and an

applied or residual stress.14 First, it can be assumed

that the failed round tensile bars used in the H2S

study expose every weld microstructure directly to

the test solution. In this study, environmental factors

that enhanced hydrogen uptake by the welds also en-

hanced the SSC susceptibility. Thus, the tensile tests

showed that these welds are highly susceptible to

SCC (Figure 8), since in air the failure was completely

ductile. In the testing solution the fracture mode was

very brittle, with a very small percentage reduction in

area values (%RA) (Figures 9 and 10), and there was a

large number of cracks induced by the solution (Fig-

ures 11 through 13). Second, permeation studies

have consistently shown that the H2S solution to pro-

duce hydrogen flux transients peaks at early times,

and these measured fluxes are directly proportional

to surface hydrogen concentrations (Figure 5). Since

the SSC susceptibility, ISCC, increases with tempera-

ture in the same fashion as the hydrogen concentra-

tion, Co, does, presumably HE would be most likely

when the concentration of hydrogen is at a maximum.

However, the corrosion current values observed in

the polarization curves, Figures 4 and 5, were almost

unaffected by temperature in the same fashion as theISCC and Co values were. Additionally, the anodic Tafel

FIGURE 8. Effect of temperature on the ISCC values for X-60 and

X-65 steel weldments.

FIGURE 9. SEM micrographs of X-65 steel weldment fractured by

SSR tests in saturated NACE solution with H2S at 37C.

FIGURE 10. SEM micrographs of X-60 steel weldment fractured by

SSR tests in saturated NACE solution with H2S at 50C.

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slopes on the polarization curves were not affected by

the temperature, which may indicate that the HE is

the main mechanism explaining why there might be

failures observed in this study, and the anodic disso-

lution component plays a secondary role, i.e., provid-

ing electrons for the reduction of hydrogen. Asahi, et.

al.,12 for instance, using the four-point bend tests and

hydrogen permeation measurements, also found that

the cracking susceptibility and the hydrogen uptake

of a steel increased as the H2S concentration and the

temperature increased, and with a decrease in the

pH value, concluding that the SSC mechanism was

HE. Griffiths and Turnbull,14 performing SSRT ex-

periments and hydrogen permeation measurements,

also found that the cracking susceptibility increased

with the amount of hydrogen uptake, and the latter

increased with a decrease in the pH, increasing the

H2S concentration and the cathodic current charging,

concluding the presence of an HE mechanism, and

finding a threshold total hydrogen content for crack-ing under SSRT conditions between 100 ppm and

250 ppm for the alloy tested. Thus, it seems that the

most likely mechanism in the SSC susceptibility ofX-60 and X-65 weldments is HE, but anodic dissolu-

FIGURE 14. Cross section of X-65 steel weldment fractured at

50C.

(a) (b)

FIGURE 11. Chemical composition of inclusions in X-60 and X-65 steels in fractured specimen by tensile tests.


50C.


25C.

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tion plays a secondary role. However, more research is

necessary to clarify this point.

CONCLUSIONS

The effects of temperature on the corrosion rate,

hydrogen uptake, and SSC susceptibility of X-60 and

X-65 weldments has been investigated. The most im-

portant results are as follows:

The corrosion rate, taken as Icorr, for X-60 weld-

ment, increased with an increase in temperature from

25C to 50C, but not for X-65.

The amount of hydrogen uptake for the weldments

increased with an increase in temperature from 25C

to 50C.

The SSC susceptibility also increased with an in-

crease in temperature for both weldments from 25C

to 50C.

The most likely mechanism for the cracking sus-

ceptibility of X-60 and X-65 weldments in H2S solu-tions seems to be HE, and anodic dissolution seems

to play a secondary role in the cracking mechanism,

but more research is necessary.

ACKNOWLEDGMENTS

The authors acknowledge DGEP-UNAM and

CONACYT, from Mexico, for their financial support.

REFERENCES

1. J.L. Dawson, J.W. Palmer, P.J. Moreland, G.E. Dicken, Ad-vances in Corrosion Control and Materials in Oil and Gas Pro-duction, European Federation of Corrosion Publications, no. 6,P.S. Jackman, L.M. Smith, eds. (Great Britain: IOM Communica-tions Ltd., 1999), p. 155-169.

2. NACE MR0175, Sulfide Stress Cracking Resistance MetallicMaterials for Oilfield Equipment (Houston, TX: NACE Interna-tional, 2005), p. 1-40.

3. NACE TM0177, Laboratory Testing of Metals for Resistance toSulfide Stress Cracking and Stress Corrosion Cracking in H2SEnvironments (Houston, TX: NACE, 2005).

4. J.C. Charbonnier, H. Margot-Marette, F. Moussy, D. Bridoux, C.Perdrix, Rev. Mtal. 1 (1988): p. 91-105.

5. A. Gingell, X. Garat, Advances in Corrosion Control and Materi-als in Oil and Gas Production, European Federation of CorrosionPublications, no. 6, eds. P.S. Jackman, L.M. Smith (GreatBritain: IOM Communications Ltd., 1999), p. 127-134.

6. G. Echaniz, C. Morales, T. Perez, Advances in Corrosion Controland Materials in Oil and Gas Production, European Federationof Corrosion Publications, no. 6, eds. P.S. Jackman, L.M. Smith(Great Britain: IOM Communications Ltd., 1999), p. 135-140.

7. H.-Y. Liou, R.-I. Hsieh, W.-T. Tsai, Corros. Sci. 44 (2002): p.2,841-2,856.

8. M.A.V. Devanathan, Z. Stachurski, Proc. R. Soc. London A 270(1962): p. 90.

9. A.R. Bhatti, M.E. Saggese, D.N. Hawkins, J.A. Whiteman, M.S.Golding, Welding Journal 63 (1984): p. 224-230.

10. R.E. Dolby, Factors Controlling Weld Toughness-The PresentPosition, Pt. II-Weld Metals, The Welding Institute MembersReport 14/1976/M (1976).

11. P.R. Kirkwood, Met. Constr. 5 (1978): p. 260-264.12. H. Asahi, M. Ueno, T. Yonezawa, Corrosion 50, 7 (1994): p. 537.13. ASTM G 1, Standard Practice for Preparing, Cleaning, and

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14. A.J. Griffiths, A. Turnbull, Corrosion 53, 9 (1997): p. 700.

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