Corrosion Science 50 (2008) 1865–1871

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Corrosion Science

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Effect of microstructure on the hydrogen trapping efficiency and hydrogeninduced cracking of linepipe steel

Gyu Tae Park a,*, Sung Ung Koh b, Hwan Gyo Jung b, Kyoo Young Kim a

a Department of Materials Science and Engineering, Pohang University of Science and Technology (POSTECH), San 31 Hyoja-Dong, Nam-Ku, Pohang 790-784, Republic of Koreab Technical Research Center, POSCO, Geidong-Dong, Nam-Ku, Pohang 790-784, Republic of Korea

a r t i c l e i n f o

Article history:Received 23 October 2007Accepted 29 March 2008Available online 1 May 2008

Keywords:A. Linepipe steelB. Hydrogen permeationC. Hydrogen induced crackingC. Hydrogen diffusionC. Hydrogen trapping efficiency

0010-938X/$ - see front matter � 2008 Elsevier Ltd. Adoi:10.1016/j.corsci.2008.03.007

* Corresponding author. Tel.: +82 54 279 2830; faxE-mail address: pentium@postech.ac.kr (G.T. Park)

a b s t r a c t

The hydrogen trapping efficiency in different microstructures is compared, and the critical hydrogen fluxfor hydrogen induced cracking (HIC) is determined for API X65 grade linepipe steel. By controlling thestart cooling temperature (SCT) and the finish cooling temperature (FCT) in thermomechanically con-trolled process (TMCP), three different kinds of microstructure such as ferrite/degenerated pearlite (F/DP), ferrite/acicular ferrite (F/AF), and ferrite/bainite (F/B) are obtained. A modified ISO17081(2004) stan-dard method is used to evaluate the hydrogen trapping by measuring the permeability (JssL) and theapparent diffusivity (Dapp). Microstructures affecting both hydrogen trapping and hydrogen diffusionare found to be DP, AF, BF and martensite/austenite (M/A) constituents. The hydrogen trapping efficiencyis increased in the order of DP, BF and AF, with AF being the most efficient. HIC is initiated at the local M/Aconcentrated region when the steel has such microstructures as F/AF or F/B. Although the trapping effi-ciency of bainite is lower than that of AF, bainite is more sensitive microstructure to HIC than to AF.

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1. Introduction

For transportation of oil and gas which contain hydrogen sulfide(H2S), the application of linepipe steel is limited due to mechanicaland corrosion problems induced by hydrogen. Especially, the hydro-gen induced cracking (HIC) and sulfide stress cracking (SSC) are themajor problems of linepipe steels induced by hydrogen during ser-vice in the sour environment [1]. When the linepipe steel is exposedto a sour environment, hydrogen atom is produced by surface corro-sion of the steel. In the presence of hydrogen sulfide, recombinationreaction of hydrogen atoms to the molecular hydrogen is retarded,consequently allowing, hydrogen atoms to diffuse into the steel.Diffused hydrogen atoms are trapped at sensitive metallurgical de-fects such as at the interface between non-metallic inclusion andsteel matrix. Cracking can occurs if the critical amount of hydrogennecessary for crack initiation is accumulated [2].

Although many investigators have studied on the HIC and SSCphenomena of the high strength linepipe steels, a great deal ofattention has been focused on the effect of metallurgical parameterssuch as microstructure. However, many of those investigators haddifficulty explaining the combined effect of microstructure andhydrogen on the cracking nature of HIC and SSC, exclusively [3,4].Few researchers have reported that acicular ferrite structure is mostdesirable microstructure for the high strength linepipe steels to as-sure high resistance to SSC, however, the role of AF has not been yet

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properly explained with respect to hydrogen diffusion [5]. In med-ium and high strength steels, lower bainite can generally offer high-er strength than acicular ferrite and show lower susceptibility tohydrogen embrittlement as compared to quenched and temperedmartensite. Nonetheless, the comparison between acicular ferriteand bainite in terms of hydrogen diffusion has not been exclusivelyexamined [6]. Therefore, it is necessary to study the relationship be-tween various microstructures and hydrogen diffusion for a clearunderstanding of cracking problems induced by hydrogen.

The aim of the present investigation is to understand the effect ofmicrostructure on hydrogen diffusion in high strength linepipesteels. Using different types of heat treatment during TMCP, steelswith three different microstructures such as F/DP, F/AF, and F/B wereprepared, and standard hydrogen permeation measurements werecarried out using a modified ISO17081(2004) method [7]. Precisionanalytical microscopy was carried out to characterize the microstruc-ture; furthermore, HIC sensitivity was evaluated in reference to NACETM0284-96 method and the crack initiation site was examined after-ward by using an electron microscopy [8]. The correlation betweenthe microstructural variations and the hydrogen permeability, appar-ent diffusivity, and solubility has been described in detail.

2. Experimental method

2.1. Test material preparation and microstructure analysis

Table 1 lists the chemical composition of the test steel used inthis study. This steel is equivalent to the API X65 grade steel. All

Fig. 1. A schematic graph representing thermomechanically controlled process(TMCP) conditions applied to make tested steels.

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the steel samples were obtained from the same slab. By using thesame slab for all the test steel samples, the level of inclusions in thesteel can be fixed in same proximity since the inclusions signifi-cantly affect the hydrogen trapping property. The steel slab was di-vided into four equal parts and each part was reheated up to1180 �C and hot-rolled to 20 mm thick and finally cooled at a rateof 12 �C/s by the accelerated cooling process. In the acceleratedcooling process, various microstructures were obtained by control-ling the SCT and FCT with three different levels while other coolingparameters were kept constant. Fig. 1 shows the TMCP conditionsand the notation of specimens used in this research.

For microstructural observations, all specimens were ground upto 2000 grit paper and polished with 1 lm diamond suspension.They were degreased with acetone and etched with nital solution(a mixture of 5% nitric acid and ethanol). The microstructure wasexamined with optical microscope (OM) and scanning electronmicroscope (SEM) at t/4 position since HIC often occurs at this po-sition, where t is the thickness of steel plate. The area fractions ofphases were quantitatively measured by an image analyzer. Espe-cially, the area fraction of M/A constituents was determined aftertint etching. To characterize the kind of phases, the hardness val-ues, measurable by an ultra-micro Vicker’s hardness tester, werealso utilized. To measure the distribution of non-metallic inclu-sions with different size, all specimens were polished up to0.25 lm and OM micrographs of non-etched clean steel surfacewere observed at X500 magnification and then, all micrographsof 60 fields were analyzed by image analysis program to measurethe area, total number and size of inclusions.

2.2. Hydrogen permeation test

The modified ISO17081(2004) standard method was used tomeasure the hydrogen permeation flux through a thin steel mem-brane. The steel membrane had the surface area of 30 � 30 mm2

and the thickness of 1 mm. The steel membrane was prepared fromthe portion at t/4 position parallel to the rolling direction. As pre-viously mentioned, this is because HIC actually occurs at t/4 posi-tion where the microstructure is homogeneous. Both sides of thesteel membrane were ground up to 2000 grit paper, and the detec-tion side of steel membrane was electroplated with palladium (Pd)in order to eliminate flux-limiting surface impedances and to en-sure the reliability of the hydrogen oxidation current. After clean-ing the specimen surface with ethanol, the steel membrane waslocated and sealed between two electrochemical cells. An area of10.5 � 10.5 mm2 was exposed to each electrolyte inside the cells.Although the morphology of steel membrane and the ratio of ra-dius to thickness were different with those recommended by thestandard method, our additional experiments showed that thedimension of the specimen in this research was sufficient enoughto ensure the volume controlled diffusion. The deaerated 0.1 NNaOH solution and the deaerated NACE TM0284-96A solution (amixture of 5.0 wt% NaCl and 0.5 wt% glacial acetic acid dissolvedin distilled water) were selected as test solutions in the hydrogenoxidation cell and the hydrogen charging cell, respectively. Firstly,the prepared NaOH solution was poured into the hydrogen oxida-tion cell. Then to ensure the continuous deaeration of the NaOHsolution during the hydrogen permeation test, N2 gas was bubbledthrough the solution for 30 min at a flow rate of 10 ml min�1 perliter of solution. The potential of the steel membrane on the hydro-gen oxidation cell was maintained at a potential of 250 mVSCE, at

Table 1Chemical composition of tested steel (wt%)

C Mn Si P S Al

0.05 1.25 0.2 0.005 0.002 0.1

which the dominant oxidation reaction was oxidation of hydrogenwhich diffused through the steel membrane. Once the backgroundcurrent was stabilized, the prepared NACE solution was pouredinto the hydrogen charging cell. During the hydrogen permeationtest, H2S gas was continuously bubbled through the NACE solutionto maintain the positive pressure. The steel membrane in thehydrogen charging cell was cathodically polarized with appliedcurrent density of 500 lA cm�2 at room temperature. Then, thehydrogen oxidation current was continuously measured for onehour with data logging every three seconds and analyzed in termsof the hydrogen diffusion and the hydrogen trapping.

2.3. Data analysis

The hydrogen flux (mol H m�2 s�1) through the steel specimenwas measured in terms of steady-state current density, iss, and con-verted into hydrogen flux according to the following equation [7].

JSS ¼iSS

nF: ð1Þ

The permeability (mol H m�1 s�1) is defined by

JSSL ¼ iSSLnF

; ð2Þ

where n is the number of electrons transferred, F the Faraday‘s con-stant, L the specimen thickness, and Jss is the steady-state hydrogenflux. The apparent diffusivity (Dapp) was determined by the break-through method

Dapp ¼L2

2p2tb; ð3Þ

where tb is the breakthrough time when a diffused hydrogen wasfirstly oxidized on a Pd coated layer. If the surface hydrogen is inthermodynamic equilibrium with the sub-surface hydrogen, theapparent hydrogen solubility (mol H m�3) is defined by

Capp ¼JSSLDapp

: ð4Þ

2.4. HIC resistance and fractography analysis

To evaluate the HIC resistance of the steels, the HIC test wasperformed in reference to the NACE TM0284-96A. Then the values

Nb + Ti + V Cu + Ni Ca Cr Mo

�0.1 �0.4 0.002 �0.1 �0.1

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of crack area ratio (CAR) and the size, position, and length of crackswere measured using an ultrasonic detector. Based on this result, afresh fracture surface was revealed by a mechanical processing,and the fractrography was investigated using SEM and energy dis-persive spectroscopy (EDS).

3. Results and discussion

3.1. Microstructure

The effect of material factors on the hydrogen diffusion dependson two factors. One is the concentration of atomic hydrogen on thesteel surface which forms due to the environmental factors such asthe pH of solution, the partial pressure of H2S gas, i.e., potentialamount of hydrogen which may diffuse through the steel. Theother is the steel microstructure consisted of the primary and sec-ondary phases including non-metallic inclusions and precipitateswhich affect both the hydrogen trapping and diffusion in the steel.Therefore, in order to understand the HIC phenomena, it is neces-sary to study the role of the microstructure on the hydrogen diffu-sion behavior in the steel.

Fig. 2 shows typical microstructures examined at t/4 position ofsteel, and they are F/DP, F/AF, and F/B. As summarized in Table 2,Steels A1 and A3 had F/DP microstructure, Steel A2 had F/AF, andSteel A4 had F/BF. A detailed SEM observation made after nitaletching or tint etching of the specimens revealed the characteris-

Fig. 2. Optical microstructures for the tested steel after 5% nital etching with: (a), (c

tics of the second phase in these steels. It is concluded that the typ-ical second phases were DP, B, and M/A constituents as shown inFig. 3. The structure of DP phase, which was referred as pseudo-pearlite without the banding pattern, was different from that ofpearlite evolved by normalizing and slow cooling treatment sincethe cooling rate of DP was higher than that of typical pearlite,and thus the carbon diffusion during cooling was not sufficientto form the lamellar structure of cementites [9]. The AF observedin this research had irregular grain boundary and there were anumber of small-sized sub-grain boundaries in the grain. This isdifferent from AF formed during welding, which had randomly-ori-ented grain boundary and showed a needle-like shape [10]. Bainiteof A4 had a number of inter-lath cementites and its micro-hard-ness value was about Hv280 with a denting load of 10gf. For allsteels, M/A constituents were observed and their average sizewas below 5 lm.

The microstructural observation showed that all steels manu-factured by TMCP consisted of fine elongated ferrite as primaryphase with various kinds of second phases which were determinedby the cooling condition such as SCT and FCT. In low FCT, the sec-ond phase of steel cooled from single phase region of stable austen-ite was AF, while that of steel cooled from two-phase region ofstable ferrite/austenite was bainite having a number of discretecementite particles. This is due to the carbon enrichment in aus-tenite with different SCT. The carbon enriched austenite in fer-rite/austenite two-phase region can easily form a hard bainite

) ferrite/degenerated pearlite, (b) ferrite/acicular ferrite, and (d) ferrite/bainite.

Fig. 3. SEM photographs showing the typical second phases and average ultra-micro Vicker‘s hardness values: (a) degenerated pearlite, (b) martensite/austenite(M/A) constituents, and (c) bainite.

Fig. 4. Total number of non-metallic inclusions with the size.

Fig. 5. Effect of thickness of steel membrane on the apparent diffusivity.

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structure which increases the hardness due to the large amount offinite cementite in the bainite. At a constant SCT condition, theamount of M/A constituents increased with the decrease in FCTindicating that FCT affects the fraction of M/A constituents whichhave a low transformation temperature.

Fig. 4 illustrates the distribution of non-metallic inclusions withthe size. The distribution of non-metallic inclusions with differentsize were analyzed by image analysis program on the micrographs(X500 magnifications) of 60 fields. The inclusions in the steel wereanalyzed as complex oxides of Al–Ca–Si–O type and the mean sizeof inclusion was 1.5 lm while the mean area fraction being 0.49%.The Ca-treatment performed during steel making process resultedin oxide inclusions with a small size of globular shape. Elongatedmanganese sulfides were never observed. It was assumed thatthe all information of inclusions such as the size, distribution,

and area fraction in steels were quite similar since all steel speci-mens were made of the same slabs produced from identical steelmaking process before TMCP. Therefore, the effect of inclusionson the hydrogen diffusion could be assumed to be similar for allspecimens and was therefore ignored in this research. In the testedsteels, the expected precipitates were Ti,Nb(C,N) though the addi-tional analysis about precipitates was not performed. The degree ofprecipitation for all specimens was assumed to be almost similarsince the precipitates were already formed at the same reheatingtemperature of 1150 �C for all steels; thus, TMCP did not affectsthe transformation of precipitates in the current research. There-fore, the effect of precipitates on the hydrogen diffusion also wasignored similar to the case of inclusion.

3.2. Hydrogen permeation test

There are several experimental methods to evaluate the hydro-gen diffusion behavior in steels. Among them, an electrochemicalmethod developed by Devanathan and Starchurski is easy to per-form in a laboratory while some experimental cautions are re-quired to obtain reliable and reproducible results [7,11]. Effortshave been put in to standardize the experimental procedures; thus,the ISO17081(2004), measurement method of determining hydro-gen permeation and hydrogen uptake and transport in metals byan electrochemical technique, was accepted as the standard meth-

Fig. 7. Change in potential of the steel membrane in the hydrogen charging cellwith time during the permeation test.

Fig. 8. Permeation transients for tested steels.

G.T. Park et al. / Corrosion Science 50 (2008) 1865–1871 1869

od. In this research, some modification of experimental proceduresin reference to the standard method was performed to simulate thepractical environment of linepipe steels. Although the morphologyand size of the steel membrane was insufficient as per the recom-mended one due to the specimen preparation, our preliminary re-search revealed that the dimension of the steel membrane used inthis research was sufficient to ensure the volume controlled diffu-sion and to ignore the contribution of surface reaction in the diffu-sion kinetics. In our additional research, the apparent diffusivityincreased according to the increase in specimen thickness whenthe specimen thickness was below 1 mm as shown in Fig. 5. Whenthe specimen thickness exceeded 1 mm, however, there was aslight change of apparent diffusivity. It means that the volume con-trolled diffusion is ensured when specimen thickness is over 1mm.Aqueous solution containing H2S gas is highly corrosive, and it cancorrode the surface of steel membrane in the hydrogen chargingcell and form sulfide scale which may impede the hydrogen injec-tion to steel matrix. Therefore, the cathodic current density of500 lA cm�2, which stabilized the behavior of permeation curveand did not induce internal crack during the permeation test,was applied to the steel membrane in the hydrogen charging cell.Fig. 6 shows the effect of applied cathodic current density on thestability of hydrogen permeation curve. The applied current den-sity of 2 mA cm�2 or 1 mA cm�2 triggered the hydrogen perme-ation curve to become unstable because of the formation ofsulfide scale hindering the hydrogen injection or the severe evolu-tion of hydrogen bubbles on the steel surface in the hydrogencharging part. On the other hands, a relatively stable hydrogen per-meation curve was obtained and some corrosion has occurred onthe steel surface in the hydrogen charging cell during the test.Fig. 7 illustrates the potential of the steel surface in the hydrogencharging cell when the applied current density was 500 lA cm�2.The potential was maintained at about �690 mVSCE during the testand the pH value of the hydrogen charging solution changed from2.95 to 2.98. In these potential and pH conditions, the Pourbaixdiagram illustrates that the iron is immune to corrosion [12].Although a little corrosion on the steel surface was observed afterthe test, we concluded that this could be ignored and the modifiedexperimental conditions were sufficient to obtain reliable resultson the hydrogen diffusion.

The hydrogen permeation test provides various information onthe hydrogen diffusion through the steel membrane. A permeationcurve imparts important factors of the hydrogen diffusion andtrapping such as apparent diffusivity (Dapp), permeability (JssL),

Fig. 6. Effect of applied cathodic current density on the hydrogen permeationtransients in NACE solution at room temperature.

hydrogen solubility in steel (Capp), and amount of irreversiblytrapped hydrogen. The Dapp represents the apparent lattice diffu-sivity dissolved and reversibly trapped hydrogen and the Capp cor-responds to the hydrogen in lattice and reversible traps [7]. It isoften believed that the decrease in Dapp, JssL and the increase in Capp

mean that more hydrogen can be trapped in steel. In this research,the effect of irreversible trapping was not a concern since it was as-sumed that distribution of the expected irreversible trapping sitesof non-metallic inclusions and precipitates in tested steels is quitesimilar.

For the steels having various microstructures, the hydrogen per-meation curve was determined in the NACE solution containingH2S gas at the applied cathodic current density of 500 lA cm�2

as shown in Fig. 8. The calculated diffusion parameters of the Dapp,the JssL, and the Capp are listed in Table 2. As shown from the resultsof Steel A1 and Steel A2 in Table 2, AF and M/A constituents withlow phase transformation temperature can be easily formed by adecrease in FCT when SCT is given in a single phase region of stableaustenite. The steel with AF microstructure has smaller Dapp, JssLand larger Capp than that with F/DP microstructure. The true diffu-sivity of hydrogen in a-iron is in the order of 10�9 m2 s�1 [13].Since for Steel A1 and Steel A2, Dapp is one order of magnitudelower, the presence of reversible traps can be inferred, and theAF and M/A constituents are expected to be the reversible trap.

Table 2Hydrogen permeation data for tested steels

Specimen Microstructure Fraction of DP/AF/B M/A fraction Dapp (�10�10 m2 s�1) JssL (�10�9 mol H m�1 s�1) Capp (�mol H m�3

Al F/DP 3.75 1.28 9.27 13.3 14.33A2 F/AF 8.12 5.73 4.05 8.47 20.91A3 F/DP 3.93 0.88 9.38 12.9 13.79A4 F/B 9.38 4.45 4.44 12 27.13

F, ferrite; DP, degenerated pearlite; AF, acicular ferrite; B, bainite; Dapp, apparent diffusivity; JssL, permeability; Capp, apparent solubility.

Fig. 9. Crack area ratio (CAR) for tested steels by ultrasonic analysis after NACETM0294 method. CAR value of the Steel A1 and A3 is zero.

1870 G.T. Park et al. / Corrosion Science 50 (2008) 1865–1871

The previous study on AF in our laboratory has shown that the AFin HSLA steels has randomly-oriented grain boundaries and highdislocation density due to tangled dislocation in it [14]. Hydrogencan be trapped reversibly at the tangled dislocation, and it can re-sult in decrease in the Dapp, and JssL and increase in Capp [5]. At thesame FCT, decrease in SCT from the single phase region of stableaustenite to the two-phase region of ferrite/austenite changedthe dominant second phase from AF to bainite. Steel A4 with F/Bmicrostructure had higher Dapp and JssL than Steel A2 with F/AFmicrostructure although the tendency of Capp did not have goodagreement due to the effect of complex microstructure. For deter-mining the reliable Capp in bainite structure, the additional exper-iment was performed in reference to JIS Z3113 called theglycerin method [15]. We could obtain the amount of diffusiblehydrogen by immersing the standard HIC specimens charged withhydrogen in glycerin in which the temperature was about 45 �Cand by measuring the volume of escaping hydrogen gas from thespecimens. As a result, the diffusible hydrogen of steel with F/AFwas much larger than that of steel with F/B when the area fractionof all kinds of phases except for ferrite was similar. This means thatthe reversible hydrogen trapping efficiency of AF is higher thanthat of bainite and more hydrogen can be trapped in the AF.

The amount of trapped hydrogen can be affected by fraction ofspecific microstructure acting as a hydrogen trapping site.Although the fraction of DP and bainite can be controlled to bethe same by controlling the TMCP parameters, the amount of allother important phases can be controlled accordingly. The M/Aconstituents critically affect HIC; however, more M/A constituentsare formed in bainite steel. Thus the fraction of all the secondphases like DP, bainite and M/A constituents can not be controlledat once. Therefore, it is difficult to compare the hydrogen trappingbehavior between F/DP and F/B microstructures due to the effect ofM/A constituents. It is reported that the bainitic structure has low-er permeability than the pearlitic structure [16]. For pearlitic struc-ture, the dominant trap sites of diffused hydrogen are located atthe interface between ferrite and cementite in lamellar pearliteor the pearlite boundary, while the interfaces of a number of fineinter-lath cementites in a bainitic structure act as an inhibitor forhydrogen diffusion [17,18]. The retained austenite (RA) itself inM/A constituents does not trap hydrogen significantly, but theinterfaces between the RA and martensitic layer may be the possi-ble trapping sites [19]. Interface area of cementites in bainiticstructure will be larger than that in pearlite and it can increasethe reversible hydrogen trapping efficiency of bainitic structure.

3.3. HIC resistance and fractography analysis

For the specimens which had undergone a standard HIC test,the CAR was measured using an ultrasonic detection, and the resultis shown in Fig. 9. No internal cracks were observed in Steel A1 andSteel A3 with F/P microstructure, while HIC occurred in Steel A2with F/AF and Steel A4 with F/B. All the information about cracksuch as the size, width, and position was measured and the frac-ture surface obtained by designed mechanical processing was ob-served by SEM as shown in Fig. 10. The detailed observation

illustrated that an initiation site of internal crack was the locallyagglomerated M/A constituents. Non-metallic inclusions, usuallyreported as the typical initiation sites of cracks, were not detectedin both Steel A2 and Steel A4. This implies that the HIC initiationsite is more closely related to the local enrichment of M/A agglom-erates rather than the average distribution of M/A constituents. M/A constituents in F/B microstructure are likely to agglomerate dur-ing accelerated cooling since the carbon enriched prior-austenitein ferrite/austenite two-phase region may result in the agglomera-tion of hard phases. As the microstructure changes from F/AF to F/B, CAR value increases for all steels and HIC is strongly affected bymicrostructure. This can be explained by high toughness of the AFimpeding the crack propagation. Therefore, it can be emphasizedthat HIC can be effectively reduced when the microstructure iscontrolled from F/B into AF and also the change in microstructurefrom F/B into the AF can meet the requirement on high strengthand toughness along with high sour resistance [5,20–22].

Diffusible hydrogen embrittles the steel microstructure and en-hances the crack propagation. It has been reported that the conceptof threshold hydrogen concentration is important to understandthe HIC phenomena in terms of both hydrogen and metallurgicalfactors and a larger hydrogen entry flux to the steel leads to a fasterbuild-up in the local concentration of hydrogen to some thresholdlevel necessary to initiate cracking [23]. In spite of the high effi-ciency of AF on hydrogen trapping, which is a potentially danger-ous site on crack occurrence, the CAR of steel with F/AF waslower than that of steel with F/B due to high toughness of the AFand high threshold hydrogen concentration for inducing HIC. Fora requirement of high strength linepipe steel with low HIC suscep-tibility, the microstructure having a low transformation tempera-ture such as the AF or bainite is inevitable. Therefore, the changein microstructure from F/B into the AF is desirable and it is neces-sary to prevent the local agglomeration of hard M/A constituents inorder to decrease the HIC susceptibility.

Fig. 10. Hydrogen induced cracking (HIC) initiation sites for Steel A4: (a) macro-scopic view of fracture surface and (b), detailed SEM image (position 1 in (a)).

G.T. Park et al. / Corrosion Science 50 (2008) 1865–1871 1871

4. Conclusions

1. The key microstructures of the high strength linepipe steel,affecting both hydrogen trapping and hydrogen diffusion, werethe degenerated pearlite (DP), acicular ferrite (AF), bainite (B)and M/A constituents.

2. The ability of microstructure to trap hydrogen was explained interms of the apparent diffusivity (Dapp), permeability (JssL), andsolubility of hydrogen in steel (Capp). The lower the values ofDapp and JssL are, the more the hydrogen trapping occurs inthe steel.

3. The results of this research suggested that the AF and B act asthe reversible trapping site and the hydrogen trapping effi-ciency was increased in the order of DP, B and AF, with AF beingthe most efficient.

4. Initiation site of internal crack was proved to be not non-metal-lic inclusions as often reported as the typical initiation sites ofcracks, but the locally agglomerated M/A constituents for thesteel with F/AF or F/B microstructure since M/A constituentsare easily embrittled by hydrogen.

5. As the microstructure changed from F/AF to F/B, the HIC suscep-tibility increased, and this can be explained by high toughnessof the AF impeding the crack propagation. Therefore, the changein microstructure from F/B into the AF was desirable, and it wasnecessary to prevent the local agglomeration of hard M/A con-stituents in order to decrease the HIC susceptibility.

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