Sulfide stress corrosion cracking and fatigue crack growth of welded TMCP API 5L X65 pipe-line steel

International Journalof Fatigue

ELSEVIER

International

Journal of Fatigue 23 (2001) 103-113 www.elsevier.com/Iocate/ijfatigue

Sulfide stress corrosion cracking and fatigue crack growth of welded TMCP API 5L X65 pipe-line steel L.W. Tsay a,*, Y.C. Chen a, S.L.I. Chan b a Institute of Materials b Institute

Engineering, National Taiwan Ocean University, Zpei-Ning Road, 202 Keelung, of Materials Science and Engineering, National Taiwan VniversiQ, Taipei, Taiwan,

Received

8 March

2000; received

in revised

form 9 June 2000; accepted

28 August

Taiwan, ROC

ROC

2000

Abstract Slow strain rate tensile test and fatigue crack growth (FCG) test were performed to evaluate the fracture behavior of API 5L X65 steel weldments after hydrogen-charging. Regardless of the testing environments, tensile fracture of all welds was located at the weld metal (WM), which had the lowest hardness as compared to that of the other parts of the weld. When hydrogen-charged in an H,S-saturated solution, all specimens suffered a small drop in tensile strength as compared to those tested in air; however, the loss in ductility was more significant. The susceptibility to hydrogen embrittlement of the welds could be reduced significantly by subjecting the welds to 6OOTY2 h tempering treatment. With the tempering treatment, the number of surface cracks in the WM after hydrogen-charging decreased and the tensile properties improved. Results of FCG test demonstrated that in a higher stress intensity factor range (AK), tempered steel plates with a severely banding structure had the higher fatigue crack growth rates (FCGRs) than those of as-received ones. The WM of an uncharged weld had the lower FCGRs as compared to the parent metal within the same AK range. However, the enhancement of crack growth in the WM was very pronounced after hydrogen-charging. SEM fractographs of tensile and fatigue-fractured surfaces revealed a quasi-cleavage fracture in the embrittled region. 0 2001 Elsevier Science Ltd. All rights reserved. Keywords:

Slow strain rate tensile test; Fatigue crack growth rates; Hydrogen embrittlement; Quasi-cleavage

1. Introduction The resistance of a pipe-line steel to sulfide stress corrosion cracking (SSCC) and hydrogen-induced cracking (HIC) is very important for steel to be used in sour oil/gas applications. Elongated MnS inclusions are the most susceptible sites for HIC initiation [ 1,2] as hydrogen atoms can easily accumulate at the interface between the steel matrix and non-metallic inclusions. The segregation of elements such as P, Mn, C in the steel enhances the formation of hard bands on cooling [3,4]. The existante of the banding structure not only decreases the resistance of the steels to HIC [3,4] but also reduces both the notch impact toughness and the reduction of area in the tensile test [5]. Especially in controlled rolling plates [l], HIC tends to propagate along pearlite bands or low temperature transformation microstructures in the steel.

* Corresponding 24625324.

author. Tel.: +886-2-24622192;

fax: +886-2-

Under the same strength level, microstructural changes from ferrite/pearlite mixtures to tempered martensite in the steel are effective in improving its SSCC resistance [6]. It has also been noted [6] that low carbon bainite (LCB) or LCB/ferrite mixtures can improve the SSCC resistance of the pipe-line steels. In a banded structure, the hydrogen transportation flux is larger along the longitudinal direction relative to the rolling direction (RD) than that along the transverse direction relative to the RD [7,8]. Tau et al. [9] studied the hydrogen-assisted fatigue crack propagation of bainitic and tempered martensitic steel, it was found that the microstructure was a dominant factor in determining the hydrogen-assisted crack growth of the steel. However, no direct relationship between hydrogen permeation behaviour and hydrogen-assisted fatigue crack growth (FCG) has been determined [9]. In general, the heat-affected zone (HAZ) of a weld (especially at the region adjacent to the fusion line) is hardened after welding. The hardened microstructures are usually sensitive to hydrogen embrittlement (HE)

0142-l 123/01/$ - see front matter 0 2001 Elsevier Science Ltd. All rights reserved. PII:SO142-1123(00)00081-5

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Journal of Fatigue 23 (2001) 103-113

Nomenclature BM base metal CGHAZ coarse-grained heat-affected zone CT compact tension FB fusion boundary FCG fatigue crack growth FCGR fatigue crack growth rate FGHAZ fine-grained heat-affected zone HAZ heat-affected zone HE hydrogen embrittlement HIC hydrogen-induced cracking LCB low carbon bainite PWHT postweld heat treatment RD rolling direction SEM scanning electron microscope SMAW shielded metal arc welding SSCC sulfide stress corrosion cracking SSRT slow strain rate tensile TEM transmission electron microscope TMCP thermo-mechanical control process WM weld metal

and stress corrosion cracking. To avoid the formation of such susceptible microstructures, the carbon equivalent or the alloying element additions have to be limited [lo]. It is reported that lowering the carbon content to the range of 0.01-0.05% markedly decreases the hardness of the central segregation zone and reduces the HIC susceptibility of high-strength pipe-line steels [5]. Significant advances in the steel-making process have been achieved by the combination of the controlled rolling with on-line accelerated cooling process, which is generally known as thermo-mechanical control process (TMCP) [ 111. The microstructures of TMCP steels are highly refined by rapid cooling, resulting in a significant improvement on strength and toughness of the steels [ 121. In addition, as the strength of TMCP steels is achieved by refined microstructures, an improved weldability of the steels by reducing the alloying additions can be expected [lo]. Furthermore, the HAZ of TMCP steel welds possesses a higher resistance to cold cracking and a better toughness as compared to that of the conventional steel welds [13,14]. Hydrogen in the steel reduces the grain boundary cohesion strength [15] and decreases the fracture stress [ 16,171. The fatigue crack growth rates (FCGRs) of the steel are accelerated in the presence of hydrogen [9,1820]. In this work, the SSCC and fatigue crack growth (FCG) of a TMCP API 5L X65 steel plate and weldments have been investigated. The resistance to SSCC of pipe-line steels was usually determined by slow strain rate tensile (SSRT) tests [21-231. The FCGRs of the weldment were measured and the results were compared

to those of the parent plate. In addition, the FCGRs in various regions of a weld were determined by aligning the crack growth normal to welding direction. While the welds were usually used in the as-welded condition, the influence of postweld heat treatment (PWHT) on the resistance to HE and FCG of the weld has also been evaluated; by subjecting the welds to tempering at 600°C for 2 h. The fractography associated with the tensile and fatigue-fractured specimen were examined, with special emphasis on the location of fracture and the change of fracture modes after hydrogen-charging.

2. Material and experimental procedures The used material was a 15 mm thick API 5L X65 steel plate produced by the TMCP process. In this process, the extent of banding found in conventional steel plates was effectively reduced by cooling with water through the temperature range of 800-500°C. The chemical composition of the steel by weight percent was O.lOC, 1.49Mn, 0.25Si, 0.25Ni, O.O24Nb, O.O24Cu, O.O5V, O.O14P, 0.001s with the balance being Fe. The shielded metal arc welding (SMAW) process was applied to fill up the joint. The joint geometry is a singlebevel 25” groove, with a root gap of 10 mm. A total of 14 welding passes was needed to complete the weld. A high cellulose type electrode with a diameter of 4.0 mm, which was suitable for the welding of API X60 grade steel, was used in this study. The electrode used satisfied the AWS A5.5 E7010-G standard specification. The

LW. Tsay et al. /International

nominal composition of the deposit by weight percent was O.O9C,057Mn, 0.2OSi, O.O14P, O.OllS, 0.52Mo and the remaning Fe. The welding parameters employed in this study were : welding current=120 A, welding voltage=25 V and welding speed=130 mm/mm. Various specimens with different orientation with respect to the RD sectioned from the welded plate are shown in Fig. 1. Schematic diagrams showing the dimensions of tensile, Charpy impact and FCG specimens employed in this work are shown schematically in Fig. 2. SSRT tests were performed at room temperature either with a strain rate of 5~10~~ s-l in air or 5~10~~ s-l in an H,S-saturated solution. Tensile specimens of steel plate aligned either parallel or transverse to RD are designated as L and T specimens, respectively. For the welded tensile (W) specimen, the weld metal (WM) was located at the centre of the 50 mm gauge length [Fig. 2(a)]. Furthermore, some tensile specimens were subjected to a PWHT at 600°C for 2 h. The testing solution was prepared according to the NACE standard TM-Ol77-86. To ensure the saturation of solution with H,S during the test, continuous purging of H,S bubbles was carried out throughout the test. Tensile testing results represent data from at least three specimens for each test. Charpy impact and FCG tests were also performed at room temperature in air. To measure the impact energy of the weld, the V notch of the impact specimen was located either at 1 mm away from the vertical fusion boundary in the fine-grained HAZ (FGHAZ) labeled HAZ specimen, or at the WM (i.e. WM specimen). The impact values represent results from an average of at least five specimens. For compact tension (CT) specimens of the steel plate, crack growth directions were aligned either parallel or normal to the RD of the steel plate, which were designated as TL and LT specimens, respectively. The effect of various microstructures in the weldment on the FCGRs was investigated by making the crack growth

Fig. 1. Schematic representation tioned from the welded plate.

showing

the testing

specimens

sec-

105

Journal of Fatigue 23 (2001) 103-113

! I

$1

’

i I I

I

IlO! 12

I 1

unit : mm

Fig. 2. Schematic diagrams showing the dimensions of (a) tensile, (b) Charpy impact and (c) compact tension (CT) specimens employed in the test.

direction normal to the welding direction. Such specimens were named as the fatigue-weld (FW) specimens. As shown in Fig. 2(c), the crack would propagate 10 mm long within the WM including a 2 mm precrack, then across the vertical fusion boundary into the HAZ, finally entering the base metal (BM). For the CT specimens (LT, TL and FW specimens) subjected to 6OO”C/2 h tempering treatment, a notation of ‘6’ behind the specified specimens is given, e.g. TL-6 stands for a tempered TL specimen. In case of hydrogen-charged CT specimens, a capital ‘H’ is added to the last symbol of designated specimens, thus, TL-H represents the hydrogencharged TL specimen. The specimen’s profile was cut by an electrodischarge machine with a copper wire electrode, followed by grinding the specimen to the required thickness. A computerized MTS hydraulic servo-controlled testing system was utilized to perform the fatigue crack growth experiment. The loading frequency was 20 Hz with a constant amplitude sinusoidal waveform of the applied load at a stress ratio of 0.1 throughout the test. In order to restrict the crack growth direction along the centreline of the specimen, a 90” V notch of 0.5 mm depth was machined on the opposite surfaces of all CT specimens. In order to study the effect of hydrogen on

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the crack growth, CT specimens were hydrogen-charged in the H&saturated solution for 48 h to ensure that specimens were fully charged with hydrogen. Then, they were washed in a mixed acid to clean off the surface corrodant, followed by rinsing with distilled water. After hydrogen-charging, no detectable defects were found in the specimens, which indicated the steel plate and weld were resistant to HIC. To inhibit the trapped hydrogen from diffusing out of the specimens during fatigue testing, the hydrogen-charged specimens were subjected to an electroless copper coating. The time interval between hydrogen-charging and fatigue testing was maintained at a constant minimum period of time. Microhardness measurements across the fusion boundaries between the deposit and BM were taken using a Vickers microhardness tester. A 3% nital solution was used to reveal the microstructures in various regions of the weld. Fractographic observations of the fractured specimens were made on a Hitachi S4100 scanning electron microscope (SEM). The thin foils examined by a JEOL-100 CX-II transmission electron microscope (TEM) were prepared by a standard jet-polisher using a 5% perchloric acid, 25% glycerol and 70% ethanol solution.

3. Results and discussion 3.1. Microstructure observation Metallographs revealing different microstructures in various regions of the as-welded weldment are given in Fig. 3. The microstructures of WM consisted of grain boundary ferrite, Widmanstatten ferrite, acicular ferrite and small amount of microphases [Fig. 3(a)]. Mainly upper and lower bainite, together with minor amount of Widmanstatten ferrite, were observed in the coarsegrained HAZ (CGHAZ) [Fig. 3(b)]. In the region reheated to a temperature just above AC,, a significantly refined grain size (fine-grained HAZ, FGHAZ) as compared to that of the BM was obtained [Fig. 3(c)]. The FGHAZ as examined by the TEM comprised of finegrained ferrite and pearlite [Fig. 3(d)]. The steel plate used in this study also revealed a mixture of fine ferrite and pearlite [Fig. 3(e)]. It shows that the degree of banding in the steel plate was not as severe as that in conventional steel plates. However, segregation bands reappeared after tempering at 6OO”C/2 h [Fig. 3(f)]. In addition, no significant changes of microstructures in the HAZ and WM between the as-welded and tempered specimens have been found. It has also been observed that the WM contained a large amount of oxide inclusions.

Joumal of Fatigue 23 (2001) 103-113

3.2. Microhardness measurement Microhardness distributions in various regions of the weld are displayed in Fig. 4. The steel plate had a Rockwell hardness of HR, 92.6 (or Hv205) in the asreceived condition. After tempering at 600°C for 2 h, a slight increase in hardness to HR, 94.5 (or Hv215) was obtained. Regardless of the tempering treatment, the most hardened region was located at the CGHAZ [Fig. 4(a)] and the WM was found to be slightly softer than the BM. Fig. 4(b) shows the hardness distribution along the centreline of the deposits from the bottom to top weld sides. Understandably the top weld was harder than the rest of the weld, due to a previous deposit tempered by subsequent welding passes. The associated reheating temperature below AC,, resulted in the decrease in hardness of previous deposit. However, the deviation in hardness was found to be narrow, given a certain fluctuation in hardness. Thus, the tempering effect imposed by subsequent welding thermal cycles on the hardness of the previous deposits seemed to be insignificant. This result could be attributed to the inherent low alloy content and low hardenability of the WM. However, the as-welded deposit of the final welding pass, which was untempered by the welding heat, had a slightly higher hardness than that found in the former deposits, as shown in Fig. 4(b). 3.3. Charpy impact test The room temperature impact energy of L and T specimens sectioned from the steel plate were 206 and 188 J, respectively. This is in accordance with the fact that the refined microstructures of TMCP steel plate have a high impact toughness. The impact energy of the HAZ was also as high as 214 J. However, the impact toughness of the WM was as low as 80 J and could be attributed to the coarse solidified microstructures and the presence of large amount of inclusions in the WM. 3.4. Tensile properties Table 1 lists the tensile properties of various specimens tested in air and in an H&saturated solution. For the steel plates tested in air, tempered specimens (L6, T6) have a slightly higher strength than that of untempered specimens (L, T). This may be due to the reappearance of the banded structure after tempering, which has an anisotropic effect on the tensile properties. Consequently, the tensile properties changed when stressed along the banding direction. Also, L and T specimens had similar tensile properties. The fusion zone was slightly softer than the other regions of the weld, and the tensile fracture of welds was located at the WM when tested in air. Furthermore, a decline in tensile strength was observed for the tempered welds. As a whole, all specimens tested in air revealed a higher

L. W. ?kuy et al. /lnternationnl

Fig. 3. Microstructures and (f) tempered BM.

in various

regions

of the as-welded

Journal

weldment.

strength and excellent ductility. However, as the testing environments changed from air to H,S-saturated solution, a significant degradation in tensile properties was found. A slight decrease in tensile strength and a remarkable drop in ductility occurred for the steel plate tested in the H,S solution, in contrast to those tested in air. Steel plates still had slightly better tensile properties than the welds. It can be noted that tensile properties of tempered welds (W6) were superior to those of the aswelded weldment (W), suggesting that the susceptibility of the welds to HE could be reduced by a PWHT at 6OO”C/2 h. Meanwhile, with or without PWHT, the fracture site of the welds in the H,S-saturated solution always located at the WM. Macrographs showing the fractured tensile specimens tested in an H,S-saturated solutions are shown in Fig. 5. By comparing Fig. 5(a and b), it can be seen that more circumferential (secondary) cracks are present on the surface of T specimens than on the L specimens surface. In addition to these circumferential cracks which were perpendicular to the loading direction, small stepwise cracks aligned parallel to the loading direction of L specimens had been found, as shown in Fig. 5(a). For the W specimens, cracks were mainly initiated and propagated in the WM [Fig. 5(c)]. Furthermore, the

oj’

Fatigue23

(2001)

(a) WM, (b) CGHAZ,

103-113

(c) FGHAZ,

107

(d) TEM micrograph

of (c), (e) BM

CGHAZ adjacent to the fusion boundary (FB) was highly susceptible to HE, as evidenced by the short cracks adjacent to fusion boundary. The amount of cracking was reduced clearly for the welds with PWHT [Fig. 5(d)], noticeable necking was found in this specimen showing again the improved tensile properties for the tempered weld. SEM fractographs of tensile fractured specimens showed ductile dimple fracture for all the specimens tested in air. For the specimens tested in the H,S-saturated solution, quasi-cleavage fracture resulted as shown in Fig. 6. 3.5. Fatigue

crack propagation

Fig. 7 shows the FCGRs (da/dN) vs the stress intensity factor range (AK) for the steel plates and welds tested in air. The LT specimens had similar FCG characteristics as those of the TL specimens, as shown in Fig. 7(a). Thus, the crack growth characteristics of the steel plate were not sensitive to the orientations of the specimens. The extent of banding in these two orientations with respect to the RD have been reduced by the TMCP treatment, thus, it is reasonable to have such results. However, the banded structure reappeared after tempering treatment, as already shown in Fig. 3(f). The FCGRs of

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L. W. Tsay et al. /International

(a) -7.

I

Journal of Fatigue 23 (2001) 103-113

I

P

--nqzzz~~

I I

WIV/:

BM HAZ --.-Tempered -+-As-welded

0'

2

4

6

8

12

10

Location (mm) Top Weld

l

EC 9

200

x Y

7---150 - Bottom Weld

3 2

100

f

--Tempered

50 01 0

-D-As-welded 2

4

6

8

10

12

14

16

Location (mm) Fig. 4. Microhardness distribution in the weld, (a) across boundary and (b) along the centreline of the deposits.

the fusion

tempered TL (i.e., TL-6) specimens were higher than those of untempered ones, especially in the high AK region. Hence, higher FCGRs in the TX-6 specimen could be attributed to the enhanced crack growth along the banded structure. Fig. 7(b) shows the result of crack growth tests for the welds as compared to the steel plates. For the as-welded weld (FW), as the crack propagated across the WM and HAZ into the BM, similar crack growth behavior as that of the TL specimen was found. The results also demonstrated that the FCGRs of WM and HAZ in FW specimens were lower than those of the BM within the same AK range. The WM and HAZ have a significant higher resistance to crack growth than Table 1 Tensile properties

of various

specimens L In air

Materials Properties/conditions UTS (MPa) Elongation (x100%) Reduction in area (x100%) Fracture location a UTS: Ultimate

tensile strength.

636 21 40 _

tested in air and in an H&saturated

In H2S 598 6 3 _

6*: subjected

that of the steel plate (or BM). Similar results were reported for the EH36 steel weldments [24]. The crack growth behavior in various regions of an as-welded weldment is not only affected by the microstructures but also by the presence of different residual stresses in these regions. The distribution of residual stresses in a welded plate has long been known [25]. The longitudinal residual stress, which acts along the welding direction, had a maximum value in the WM and changed from tensile to compressive stresses as it moved away from the fusion zone. Meanwhile, the transverse residual stress varies symmetrically from compressive at the edges to tensile at the centre of the plate. It is known that the presence of compressive residual stress ahead of the crack tip will cause crack closure [26,27], resulting in reduced FCGRs. As the crack propagated across the as-welded WM, the partial relief and redistribution of residual stresses caused the crack face to bend or to rotate [28], hence, an enhanced crack closure occurred. Thus, decreased FCGRs in the WM could be expected. When the crack grew further into the HAZ, the magnitude of longitudinal residual stresses gradually decreased hence its effect on the FCGRs was found to be reduced. Finally, the crack growth characteristics became similar to that of the parent metal as the crack grew in the BM. The fatigue crack growth behaviour of FW specimen was in agreement with the model of Shi et al. [28], who performed a detailed study on the effect of welding residual stresses and the fatigue crack growth of pipeline steel. Tempered welds (FW-6) had higher FCGRs in the WM and HAZ than those in the as-welded weldments (FW). No significant microstructural changes were revealed between the as-welded and tempered welds in different regions of the welds. Such changes of crack growth behavior for tempered welds, thus could be attributed to the relief of residual stresses after tempering. The two crack growth curves of IV-6 and TL-6 specimens intersected at a AK of about 20.6 MPa&. Lower FCGRs in the WM than those in the tempered steel plate were found at the low AK range. However, the situation was reversed for the crack growth in the HAZ of FW6 specimens. The da/dIV values in the HAZ were higher than those in the tempered steel plate at the same AK

solutions*

T In air

In H,S

L6* In air

621 23 40 -

573 5 3 -

646 20 40 _

to 6OO”C/2 h tempering

treatment.

In H2S 624 7 3 _

T6* In air 635 22 40 -

In H2S 602 7 3 -

W In air 575 15 40 WM

In H,S 531 3 4 WM

W6* In air 562 16 40 Wh4

In H2S 547 6 4 WM

109

Fig. 5. Macroscopic photographs showing the different fractured (d) W6 specimens. Arrows indicating examples of circumferential of the L specimen.

Fig. 6. SEM fractographs the H,S-saturated solution (b) W specimens.

of the tensile fractured showing quasi-cleavage

specimens fracture

specimens tested in the H2S-saturated surface cracks. SWC: stepwise crack

tested in of (a) L,

range. As the crack passed through the HAZ, consistent FCGRs were resulted for both FW-6 and TL-6 specimens. The results of fatigue crack growth tests for steel plates and welds after hydrogen-charging are shown in Fig. 8. For the hydrogen-charged TL (i.e. TL-H) specimen, enhanced crack growth only occurred at the low AK range [Fig. 8(a)]. Meanwhile, the effect of hydrogen

solution for the (a) L, (b) T, (c) W and aligned parallel to the loading direction

on accelerating the crack propagation of TL-6-H specimen was more prominent. Such a result may again be attributed to the reappearance of banding structure in the tempered specimen, owing to the banding structure being sensitive to HIC. The hydrogen diffusivity was the highest along the longitudinal direction of a banded structure than that along the other directions [7,8], thus sufficient hydrogen atoms diffused into the strained region ahead of the crack tip and assisted crack propagation. Consequently, the tempered steel plate after hydrogen-charging had higher FCGRs than those of the untempered one within the low AK range. It also demonstrated that the FCGRs of TL-6-H and TL-H specimens were similar to those of TL-6 and TL specimens respectively, as the AK was greater than about 20 MPadm. Such results could be related to the relatively slow transportation of hydrogen atoms to the embrittled region ahead of crack tip during crack growth, when the fatigue crack propagated at a faster rate. Thus, at higher AK values the influence of hydrogen was found to be lesser. For the hydrogen-charged welds, the increase in FCGRs was very significant as compared to those of the uncharged specimens. As shown in Fig. 8(b), the FCGRs in the WM and HAZ of the charged specimens (i.e. FWH) were higher than those of the FW specimens. The results also showed that the degree of enhanced crack growth was more pronounced for the weld than for the steel plate after hydrogen-charging. When the crack propagated through the HAZ into the BM, similar’crack growth properties for the FW and FW-H specimens were obtained, where the influence of hydrogen was lesser owing to the higher crack growth rates relative to hydro-

L. W. Tsay et al. /International

Journal of Fatigue 23 (2001) 103-113

1o-3

(4

t(a)

x

x

LT 4

0

1o-5

.

TL-6-H

.

TL-H

0

TL-6

+

TL

30

20

AK (MPa&)

AK (MPadi)

(b)

+

TL

f

f

4

+

FW

; lo-

40

I

IO

20

30

40

A.K (MPa,h)

Fig. 7. Fatigue crack growth behaviours of (a) steel plates and (b) weldments.

gen diffusion. Significantly accelerated crack growth was found in the WM of FW-6-H specimen. The resistance to FCG in the WM was not improved by tempering treatment. But as the crack growth passed the WM into the HAZ, the FCGRs significantly dropped, suggesting the tempered HAZ with a higher resistance to crack growth than the WM in the FW-6-H specimen. Fig. 9 shows the macroscopic appearances of fatiguefractured surfaces of various specimens. As mentioned earlier, the FW specimen has a higher resistance to crack growth than the TL specimen at the low AK range without hydrogen-charging. Both steel plates and welds revealed a flat and smooth fracture surface, as shown in Fig. 9((a) and (b)) respectively. Furthermore, no obvious differences were found in the macro fatigue-fractured appearance in various regions of the weld between the

AK (MPahi) Fig. 8. Fatigue crack growth behaviours of the hydrogen-charged (a) steel plates and (b) weldments.

as-welded and tempered welds. For the specimens after hydrogen-charging, a change in fatigue fractography was found. Accelerated crack growth of the TL-H and TL6-H specimen was associated with a shinny area ahead of the notch tip [Fig. 9(c)]. This area has been shown to be quasi-cleavage fracture in a later micrograph of Fig. 10(c). The extent of embrittled area was deeper in the central portion of the specimen than in the free surfaces. Such features could be due to a higher constraint developed at the central portion of the specimen, implying the detrimental effect of hydrogen in this region. In an earlier work by Toribio [29,30], it was indicated that the diffusion of hydrogen was governed not only by the concentration gradient but also by the hydrostatic stress field in the material. Both the FW-H and

L. W. T~say et al. / Internutional

Journul of Fatigue 23 (2001) 103-I 13

Ill

precrack

of crack growth

Direction

f Fig. 9.

Macroscopic

appearances

of fatigue-fractured

surfaces

FW-6-H specimens showed similar macro fatigue-fractured features. The coarse columnar grains, which were not refined by the subsequent welding passes, revealed brittle fracture features. It was also found that the HAZ of FW-6-H specimen was less susceptible to HE, and revealed less brittle features (not shown here) as compared to that of the FW-H specimen. SEM fractographs of fatigue-fractured specimens are shown in Fig. 10. Without hydrogen-charging, fatiguefractured appearance of the steel plate revealed a transgranular fatigue fracture with fatigue striations at high AK values [Fig. 10(a)]. Although the WM and HAZ of the FW specimen had a higher resistance to crack growth as compared to the steel plate, fatigue fracture appearance was observed to be similar between them [Fig. 10(b)]. The only difference was the presence of a large amount of oxide inclusions in the WM of the FW specimen. After hydrogen-charging, the fatigue fracture surface of the shinny area in the TL-6-H specimen was mainly quasi-cleavage fracture in the low AK range [Fig. 10(c)]. It was reasonable that the enhanced crack growth of hydrogen-charged steel plates was associated with the more brittle fracture appearance. A remarkable difference was found in fatigue fractography of the FW-H and FW specimens. The portion of fracture surface consisted of a large extent of quasi-cleavage in the WM of FWH specimens [Fig. 10(d)]. In addition, noticeable quasicleavage fracture was also observed in the CGHAZ of FW-H specimens [Fig. 10(e)]. These results imply that

of (a) TL, (b) FW, (c) TL-H and (d) FWH

specimens.

the WM and CGHAZ of the as-welded weldment are highly sensitive to HE, hence, enhanced crack growth in the WM and HAZ of FW-H specimen can be expected. As mentioned previously, the FW-6-H specimen also exhibited high FCGRs in the WM and decreased in FCGRs as the crack grew into the HAZ. The observation on fatigue-fractured surfaces showed similar fracture features in the WM of FW-H and FW-6-H specimens, again confirming that the PWHT had little effect in improving the hydrogen-assisted FCG behaviour of the WM for this pipe-line weld. However, transgranular fatigue fracture was found in the CGHAZ of FW-6-H specimen [Fig. 10(f)]. The improved resistance to crack growth in the CGHAZ of FW-6-H specimen, as shown in [Fig. 8(b)], could be related to the effective reduction of quasi-cleavage fracture in this region.

4. Conclusions The sulfide stress corrosion cracking and fatigue crack growth of welded TMCP API 5L X65 pipe-line steel have been studied in this work. Following conclusions are drawn: 1. Ferrite and pearlite mixed structures were found in the as-received steel plate. Clearly banded structures which would deteriorate the mechanical properties of the steel plate reappeared after tempering at 6OO”C/2

I12

Fig. IO. SEM fractogrdphs CGHAZ of FW-6-H.

L. W. T.sq et (11./International

of fatigue-fractured

specimens

Journal of Fatigue 23 (2001) 103-l 13

(a) TL, (b) WM of FW, (c) TL-6-H.

h. The microstructure of weld metal (WM) consisted of grain boundary ferrite, Widmanstatten ferrite, acicular ferrite and a few microphases. Bainite with a small amount of Widmanstatten ferrite were observed in the coarse-grained heat-affected zone (CGHAZ). The fine-grained heat-affected zone (FGHAZ) comprised of fine-grained ferrite and pearlite. 2. The most hardened region was located at the CGHAZ, and the WM had a slight lower hardness than the base metal (BM). The relatively lower hardness in the WM in essence could be attributed to the inherent low alloy content and low hardenability, resulting in the tensile fracture of welds located at this region. Postweld heat treatment (PWHT) of the weld had only a minor influence on the changes of microstructures and hardness. 3. Tensile properties of steel plates were found to be superior to welds tested either in air or in an H,Ssaturated solution. After tempering at 6OO”C/2 h, the improved tensile properties of the weld tested in an H,S-saturated solution were associated with a reduction in the amount of surface cracks in the WM.

(4 WM of FW-H,

(e) CGHAZ

of N-H

and (f)

4. The WM and HAZ of the as-welded weldment possessed a higher resistance to crack growth than the steel plate. The reappearance of banded structure in the steel plate after tempering treatment resulted in an enhancement of crack growth as compared to the as-received steel plate, as AK was greater than 22.4 MPa.\im. The effect of hydrogen seemed to be diminished as AK increased. This can be attributed to the relatively slow hydrogen diffusion to the crack tip as compared to the fast FCGRs at high AK. In addition, the acceleration of fatigue crack growth rates (FCGRs) was rather significant for the hydrogencharged untempered weld. Meanwhile, the resistance to crack growth in the WM did not improve for the hydrogen-charged tempered weld. However, the tempered HAZ could retard the fatigue crack growth as compared to the WM of a hydrogen-charged tempered weld. 5. Fractographic observations of tensile fractured and fatigue-fractured surfaces by SEM revealed quasicleavage fracture in the embrittled region.

L.W. Tsay et al. /International

Acknowledgements The authors gratefully acknowledge the financial support of the Republic of China National Science Council (Contract No. NSC 87-2216-E-019-008).

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