58 (2007) 721–729

Materials Characterization

Effect of artificial aging on the microstructure of weldment onAPI 5L X-52 steel pipe

B. Vargas-Arista, J.M. Hallen, A. Albiter ⁎

Departamento de Ingeniería Metalúrgica, IPN-ESIQIE, Laboratorios Pesados de Metalurgia, UPALM Zacatenco, 07738 México DF, Mexico

Received 16 June 2006; received in revised form 11 October 2006; accepted 16 November 2006

Abstract

The effects of artificial aging on the microstructure in the weldment of an API 5L X-52 steel pipe were studied. Aging wasperformed at 250 °C over a period of 1000 h and values were recorded at every 100 h intervals. Transmission electron microscopyobservations showed precipitation strengthening from nearly circular Nb–C containing nanoparticles for the base metal and heataffected zone, and cementite for the weld metal. The largest amount of precipitation in the weldment zone was obtained at 500 h,due to peak-aging, which showed the highest particle density. The weld metal was more susceptible to aging, exhibiting the highestincrease in precipitation at 500 h, followed by the heat affected zone. After 500 h, the deterioration in the microstructure wascaused by the coarsening of particles due to over-aging. The base metal showed the larger increment in particle size after 900 h ofaging accompanied by a bigger decrease in fine particles than in the weld metal.© 2006 Elsevier Inc. All rights reserved.

Keywords: Microstructure; Precipitation; Artificial aging; API 5L X-52 steel pipe; Weldment

1. Introduction

API 5L steel pipe is used for natural gas, hydrocarbonsand crude oil transport in the petroleum industry. The pipeis manufactured using microalloyed steel with elementssuch as Nb, Ti, V and a low carbon content (b0.1 wt.%),from which its strength, toughness and weldability [1,2]are obtained. A thermomechanically controlled rollingprocess is employed for the control of grain refinementand subsequent precipitation strengthening, resulting in

⁎ Corresponding author. Instituto Mexicano del Petróleo, Eje CentralLázaro Cárdenas No. 152 Col. San Bartolo, Atepehuacan, México DF07730, Mexico.

E-mail addresses: aalbiter@imp.mx (A. Albiter),bvarista26@yahoo.com.mx (B. Vargas-Arista), j_hallen@yahoo.com(J.M. Hallen).

1044-5803/$ - see front matter © 2006 Elsevier Inc. All rights reserved.doi:10.1016/j.matchar.2006.11.004

five classes of precipitates: (1) very large (2 to 10 μm insize) cuboidal TiN particles; (2) large (0.1 to 1.0 μm) TiNparticles; (3) medium size (30 to 50 nm) irregular shaped(Nb, Ti) carbo-nitrides; (4) fine (b20 nm) rounded Nbcarbo-nitrides; (5) very fine (b5 nm) dispersed precipi-tates [2–4]. In the rolling process, the steel continuouslydeforms and recrystallizes, producing a fine austenitegrain size promoted by TiN particles from casting, whichprevents austenite grain growth by pinning the grainboundaries. During the finishing stage at 700 to 900 °C,austenite recrystallization is delayed by the strain inducedprecipitation of (Nb, Ti) carbo-nitrides pinning austenitegrain boundaries; a resultant structure is obtained ofaustenite grains elongated along the rolling direction. Atthe end of rolling, ferrite grains nucleate at the prioraustenite grain boundaries growing across the grains; thisleads to a very fine acicular ferritic grain structure.

Table 1Chemical composition (wt.%) of the base and weld metals

Materialcondition

Element (wt.%)

C Mn Si Ti V Nb P S Al Ni Ceq

Base metal 0.080 1.06 0.26 0.003 0.054 0.041 0.019 0.003 0.039 0.019 0.274Weld metal 0.046 1.14 0.27 0.020 0.030 0.050 0.021 0.010 – – 0.242

Table 2Mechanical properties of the base and weld metals of API 5L X-52steel pipe

Materialcondition

Hardnessvickers HV

σYS 0.2%MPa

UTSMPa

Elongation(%)

Base metal 188.3 419 514 30Weld metal 217.1 518 589 22

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Interphase precipitation can accompany the austenite-to-ferrite transformation with Nb (C, N) precipitates of20 nm size nucleating at phase boundaries during thegrowth of ferrite. Deformation may be continued after theaustenite has transformed to ferrite, which leads toextensive dislocation networks. During coiling of thesteel below 600 °C, extremely small Nb (C, N)precipitates (b5 nm) nucleate on the dislocation networkswithin the ferrite grains [3,4].

As with most alloys, the steel pipe is also subjected toa natural aging phenomenon during long periods of timeat room temperature, which is likely to be accelerated bythe service conditions within the temperature range from25 to 70 °C, producing changes in the microstructureand mechanical properties after long-term service [6,7].

There are only a few investigations reported regardingthe aging in ferrous alloys, particularly in the weldmentof API 5L steel pipes that have aimed at understandingthe aging, such as the deterioration or degradation of themicrostructure. Some studies are related to the weldmetal of multi-layer steel aged at 550 °C for a time up to30,000 h; the fine grain microstructure in the heataffected zone shows structural changes such as theagglomeration and coarsening of precipitates [8]. Asecond case is the weld metal of 308 stainless steelproduced by submerged arc welding aged at 500 °C for24 h, resulting in chromium carbide precipitation [9]. Athird case related to aging is the API 5L steel pipe aged at250 °C exhibiting an increment in ferrite grain size as afunction of aging time [10]. A further study on a titaniumalloy aged in the range of 300 to 500 °C for 1 to 2778 hshowing an increment in ω and α precipitates in thematrix [11]. Additionally, the 50.7% Ni–Ti alloy aged at500 °C from 1 to 100 h, resulted in Ni4Ti3 precipitation,whichmodified the mechanism of nucleation and growthof precipitates [12].

There are studies on carbon steel, indicating thatnatural aging can be accelerated by heat treatment, in thetemperature range from 100 to 350 °C [6,7]. Additionally,Homma et al. [13] reported that low strength steel issusceptible to aging at 250 °C for 1 h that is equivalent to2 years of natural aging. Most of the line pipe used for thetransport of hydrocarbons is manufactured with this typeof steel, having a service life of about 20 years. Therefore,

it was decided to accelerate the natural aging by heating at250 °C to study the long-term changes of natural aging, onthe microstructure of the weldment zone.

The aim of this study is to characterize themicrostructural changes in the weldment zone, as theresult of the accelerated artificial aging at 250 °C fordifferent times. The contribution of this study is toprovide experimental information that will be useful inunderstanding the microstructural deterioration ofcommercial line pipe, with more than 20 years ofservice, without considering the work pressure andchemical composition of the fluid transported.

2. Experimental procedures

2.1. Material

Weldment specimens were obtained from an API 5LX-52 steel pipe, which was 9.5 mm (0.375 in) thick and914 mm (36 in) in diameter. The pipe was longitudinallywelded by a double-pass from a submerged arc weldingprocess, according to the API 1104 welding procedurespecification [14]. The chemical composition and themechanical properties for the untreated base and weldmetals are given in Tables 1 and 2. These results were inaccordance to API 5L [15]. The dimensions of thetransverse specimens of weldment were 10 mm(0.394 in) width and 50 mm (2 in) length. The specimenswere artificially aged in a Carbolite furnace at 250±5 °C[8,11] for 1000 h and the values were monitored every100 h in order to cover a complete aging process.

2.2. Microstructure

In order to characterize the microstructure of the basemetal, the secondary recrystallization zone of the heat

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affected zone and the weld metal, the aged specimenswere prepared by metallographic techniques and etchedusing Nital 2 vol.% solution in order to reveal themicrostructure. The microstructural features were ob-served using a JEOL 6300 scanning electron microsco-py operated at 20 kV. The morphological and chemicalcharacterizations were achieved by means of a JEOL-2000FX-II transmission electron microscopy operated at200 kV, which was equipped with an energy dispersiveX-ray spectrometer. Thin foils of 3 mm (0.118 in)diameter were prepared by grinding and jet-polishing,using a 30 vol.% HNO3+70 vol.% CH3OH solution at−55 °C.

3. Experimental results and discussion

3.1. Micro-structural analysis of the untreated weld-ment using scanning electron microscope (SEM)

SEM images of the longitudinal weldment on API 5LX-52 steel pipe are shown in Fig. 1. SEM observationsshowed that the weldment zone generated by doublesubmerged arc welding (DSAW) was formed of threedifferent microstructures (Fig. 1a). The as-rolledmicrostructure of the base metal (Fig. 1b) was mainlyformed by a ferritic matrix (dark areas), and pearliticcolonies (white areas) with a percentage of each

Fig. 1. (a) Untreated weldment by the SAW process, (b) the base metal as-rol

constituent of about 66 and 34, respectively. Bothphases showed the rolling direction related to thethermomechanically controlled rolling process [3,4].The measured average ASTM grain size number of theuntreated condition was 9.36 ASTM, according toASTM E 112. This averaged ASTM grain size numberfulfils the Mexican specification for this type of steel,indicated in the NRF-001-PEMEX [16].

The heat affected zone generated by the weldingthermal cycle (Fig. 1c) showed a complex recrystallizedmicrostructure located near the fusion line, formed bycoarse-grained ferrite (dark areas), acicular ferrite (whiteareas), and small pearlitic colonies at grain boundaries(white points) [8,13,17]. The weld metal (Fig. 1d)exhibited a solidification microstructure resulting fromthe welding thermal cycle. This was mainly composedof large columnar ferrite grains preferentially oriented inthe direction of heat flow (dark areas), a mix of phases ofpearlite and acicular ferrite between columnar grains(white areas), and spherical non-metallic inclusions of1 μm diameter (white points) [9,18,19].

3.2. Transmission electron microscopic (TEM) analysisof the untreated weldment

Fig. 2a and b show typical TEM bright field images ofthe base metal structure. Fig. 2a clearly shows the

led, (c) the HAZ as recrystallized, and (d) the weld metal as solidified.

Fig. 2. TEM bright field (BF) images obtained from the base metal untreated: (a) cementite at a grain boundary, (b) discontinuous pearlite, (c) Ti–Ccontaining particle, and (d) its corresponding EDS spectrum showing Ti and C peaks.

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presence of cementite located in ferrite grain boundaries(dark lines), and discontinuous pearlite with discontinu-ous carbide and precipitation of particles (Fig. 2b) due tothe pearlitic transformation [1,3,4]. Additionally, largenearly circular Ti–C containing particles of 220 nmaverage size were located within ferrite grains, having ahigh solubility product in the austenite phase of2.5×10−4 wt.% at 900 °C [1]. Therefore, particles actedas a refinement agent of austenite grain due to a straininduced and strengthening precipitation generated by thethermomechanically controlled rolling process [1,3,4], asobserved on the TEM image and its corresponding EDSspectrum in Fig. 2c and d, respectively.

The heat affected zone (HAZ) structure of the second-ary recrystallized zone produced by the welding thermalcycle shows coarse ferrite grains (Fig. 3a), and plateshape cementite located in a ferrite grain boundary (darklines), which is clearly seen in the TEM image (Fig. 3b)[2,6,17]. Also, it was observed that there are circularshaped Nb–C containing particles from 20 to 40 nm size

within ferrite grains, and a low solubility in the ferritephase of the order of 10−6 wt.% at 700 °C, as indicated inthe TEM image and its corresponding EDS spectrum inFig. 3c and d, respectively. The particles were obtainedby a partial dissolution and re-precipitation processeswith continuous cooling from 1350 and 1200 °C(welding thermal cycle). Therefore, the precipitationstrengthening produced by the thermomechanicallycontrolled rolling process was modified [1,8,20].

Fig. 4a and b show TEM observations of the weldmetal structure obtained from the welding thermal cycleby the DSAW process, exhibiting large columnar ferritegrains, and cementite with preferential nucleationlocated in a ferrite grain boundary (dark lines). TheTEM image in Fig. 4c shows large nearly circular Fe–Ccontaining particles with an average size of 200 nm. TheEDS analysis of the particles assumed the formation ofcementite (Fe and C peaks shown in Fig. 4d) and theseparticles were located within ferrite grains and acted as arefinement agent of grains, which is associated with the

Fig. 3. TEM BF images obtained from the HAZ untreated: (a) plate ferrite, (b) cementite at a grain boundary, (c) Nb–C containing particle, and (d) itsEDS spectrum showing Nb and C peaks.

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continuous cooling from 1350 °C and through theaustenite–ferrite temperature range during the weldingthermal cycle [6,18,19].

3.3. TEM structural analysis of the weldment aged at250 °C

Three significant effects on the microstructure wereobtained from the weldment aged at 250 °C. The effectswere an increment in ferritic grain size in the base metal,precipitation strengthening from nanoparticles [1,5,11],and their coarsening process [1,8,12]. The ASTM grainsize number was reported in our previous paper [21].The effect of aging on the microstructure of the basemetal, secondary recrystallization of the HAZ, and theweld metal was the nanoparticles precipitation process.Fine nearly circular Nb–C containing particles wereprecipitated (dark points) on the dislocation networkswithin ferrite grains of the base metal, as observed in theTEM image (Fig. 5a). The EDS spectrum of these

particles clearly displayed the characteristic Nb and Cpeaks (Fig. 5b), which is identified as the Nb2Ccompound reported in a previous study [21]. Similarresults of the particles (dark points) on the dislocationswithin ferrite grains in the HAZ and weld metal werefound and indicated in Figs. 6a and 7a.

The main difference between the particles of the baseand weld metals was their chemical composition. In theFe–C containing particles for the weld metal, Fe and Cpeaks were detected in the EDS spectrum (Fig. 7b),which is identified as the Fe3C cementite compound[21]. The base metal and the HAZ showed the samechemical composition of Nb–C containing particles, asindicated in the EDS spectra (Figs. 6b and 5b). Themaximum particle precipitation was obtained for theweld metal, the HAZ, and the base metal after 500 h,which showed an increase in fine particles, due toprecipitation strengthening described by the Ashby–Orowan process [1–3]. The particles showed heteroge-neous nucleation on dislocation networks, giving

Fig. 4. TEM BF images obtained from the weld metal untreated: (a) columnar grain, (b) cementite at a grain boundary, (c) Fe–C containing particle,and (d) its EDS spectrum showing Fe and C peaks.

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stronger interactions, with more intersections andobstacles for dislocation motion, thus, impartingadditional strength to the final microstructure [1–3].

TEM image analysis shows the apparent nanoparticledensity as a function of the aging time, in Fig. 8 this isgiven for weldments aged at 250 °C. An evidentincrease was observed with the advance of aging timefor the weld metal, the HAZ, and the base metal,achieving a maximum density at 500 h as a consequenceof the peak-aging condition, after this time a decrease indensity clearly appeared.

The base metal exhibited a larger particle densitythan that of the weld metal and the HAZ as a function ofthe precipitation strengthening by the thermomechani-cally controlled rolling process [1]. The HAZ showed alower particle density, which is attributed to the partialdissolution and re-precipitation of particles produced bythe welding thermal cycle [20]. However, the weldmetal exhibited the largest increase in particle density(47 vol.%) than that of the HAZ (42 vol.%) and the base

metal (22 vol.%) peak-aged compared to that of theuntreated condition. Therefore, the weld metal showedmore susceptibility to aging, promoting a largerprecipitation. The base metal showed a minor increasein particle density due to the reduced extent ofprecipitation.

The over-aging condition was obtained after 500 h ofaging, producing an effect on the weldment microstruc-ture, which generated a decrease in particle density andparticle coarsening by the Ostwald ripening process. Aselective growth of the larger particles at the expense ofthe other smaller ones was the dominant process. Thelatter diminished continuously until they disappearedcompletely. The new medium size particles had weakinteractions with dislocation networks, and representedweaker obstacles for plastic deformation and generated adecrease in strength of the microstructure [1,8,11,12].From TEM images analysis, the particle coarseningmeasured as average particle size is shown in Fig. 9.Coarsening of the particle size in the base metal, in the

Fig. 5. TEM BF image obtained from the base metal peak-aged for500 h: (a) fine Nb–C containing nanoparticles on dislocations, and (b)its corresponding EDS spectrum.

Fig. 6. TEM BF image obtained from the HAZ peak-aged for 500 h:(a) fine Nb–C containing nanoparticles on dislocations, and (b) itsEDS spectrum.

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HAZ, and in the weld metal was not observed during theearly stages of aging. Only the precipitation processprevailed. After 500 h, the coarsening of particles wasdetected due to over-aging. The base metal showed alower average particle size than that of the HAZ and theweld metal produced by the thermomechanicallycontrolled rolling process [1]. The HAZ exhibited alarger average particle size than that of the weld and thebase metals due to the heating effect by the weldingthermal cycle [17,20].

Therefore, another difference among the three mi-crostructures was the particle size. The HAZ showed thelargest average particle size in the range of 15 to 30 nm,followed by the weld metal (10 to 30 nm) and the basemetal (10 to 25 nm) after 900 h aging, as can be observedin the TEM images in Fig 9. However, the base metalachieved the largest increase in average particle size(75 vol.%) more than that of the weld metal (33 vol.%),

and the HAZ (29 vol.%) after 900 h aging compared tothat of the untreated condition. The base metal was moreaffected by the over-aging due to the fine particlesprecipitated by the thermomechanically controlled roll-ing process contributing to the coarsening of the largerparticles [1–3].

In summary, the weld metal was more susceptible toartificial aging at 250 °C than theHAZ and basemetal dueto the larger amount of nanoparticle precipitation and theincrement in particle density, producing higher precipita-tion strengthening in the peak-aging condition. The agingevolution of the weld metal, the HAZ and the base metalwas obtained frommicrostructural differences, such as thedifference in particle precipitation and subsequentcoarsening among the alloys. The main effect of artificialaging was the particle precipitation at short times, andtheir coarsening at long times. The aging results suggestedthat during the natural aging at a service temperatureranging from 25 to 70 °C of the commercial line pipe, thecarbide particles are precipitated suggesting an impro-vement in the strength, and after long-service times,

Fig. 7. TEM BF image obtained from the weld metal peak-aged for500 h: (a) fine Fe–C containing nanoparticles on dislocations, and (b)its EDS spectrum.

Fig. 8. Particle density vs. aging time for the we

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coarsening of carbide particle begins, and an increase inthe ferritic grain size producing a deterioration in themicrostructure. We believe that our results are importantto the pipeline community dealing with the microstruc-tural behavior, which shows significant changes producedby natural aging.

4. Conclusions

The weld metal, the HAZ, and the base metal weresusceptible to artificial aging at 250 °C. Aging affectedtheir microstructure with significant changes, such as anincrease in grain size in the base metal, precipitationstrengthening from nanoparticles and coarsening ofnanoparticles.

The effect of artificial aging was related to theprecipitation strengthening for the peak-aging conditionat 500 h, and to a particle coarsening process, above500 h.

The weld metal was more susceptible to acceleratedaging due to a larger increase in particle density andlower particle coarsening.

The evolution of artificial aging strongly dependedon microstructural features and precipitation processes,such as the morphological type of the ferrite grains, thenanoparticle size, the chemical composition of theprecipitates and the particle density.

Acknowledgements

The authors are grateful for the experimental andfinancial support received from Instituto Mexicano del

ldment aged at 250 °C for different times.

Fig. 9. Average particle size vs. aging time for the weldment aged at 250 °C for different times.

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Petroleo, Program of Technology Development PEP-IPN and CONACYT.

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