Effect of Tempering Pass on HSLA-80 Steel HAZ …...cation of HSLA steels where weight re-duction is required (through wall thick-ness reduction), while ensuring high weldability,

304-s

WE

LD

ING

RE

SE

AR

CH

Introduction

Microalloyed steels of high strengthand low alloy (HSLA) have an excellentcombination of properties (high yieldstrength, toughness, and weldability) dueto their unique characteristics of smallgrain sizes achieved through very low con-tents of alloy elements and thermome-chanical treatment. Thus, Almeida et al.(Ref. 1), who is also supported by otherauthors such as Jing-Hong et al. (Ref. 2),refers to the tendency of increasing appli-cation of HSLA steels where weight re-duction is required (through wall thick-ness reduction), while ensuring highweldability, which is required for fieldwelding. The trend in the development ofthese steels for piping (API 5L X-60, X-80,X-100 steel, according to the classificationof the American Petroleum Institute) is

closely linked to industrial demand for in-creasing the diameters and work pressurein pipelines.

The continuous improvement in theproperties of the HSLA steels has beenachieved by the presence of a very low con-tent of alloying elements, such as Nb, Ti,and V, and a thermomechanically con-trolled treatment during rolling, which(both) contributes to a decrease in grainsize. Other important factors, as reportedby Jing-Hong et al. (Ref. 2), are the for-mation of desired microstructures andprecipitates (acicular ferrite and bainite,which may also appear with retained

martensite/austenite (MA) constituents).On the other hand, since welding

processes are employed in applicationsusing steels from the API 5L class, it is evi-dent that the influence of the thermal cycleson the heat-affected zone (HAZ) mi-crostructures is of vital importance, espe-cially in the region of coarsened grains. Thishas resulted in a large number of papers fo-cused on the study of this area. Das (Ref. 3)suggested that the refinement of the mi-crostructure near the weld interface is veryeffective in improving the HAZ toughnessof microalloyed steels, which is achieved bydelaying the austenite grain growth, withthe transformation to acicular ferrite andformation of precipitates or inclusionswithin the austenite grains. For his part,Moeinifar et al. (Ref. 4) addressed the ef-fect of thermal cycling in X-80 steel, show-ing that the heat input in the submerged arcwelding (SAW) process (leading to differ-ent thermal cycles) has a significant influ-ence on the microstructure and hardness ofthe coarse-grain HAZ. With similar results,Mohandas et al. (Ref. 5) provided a com-parative study of the behavior of HAZ mi-crostructure and properties for three steelsof different compositions with the shieldedmetal arc welding (SMAW), gas tungstenarc welding (GTAW), and gas metal arcwelding (GMAW) processes, relating theprocess energy input to the microstructureand properties of the HAZ.

Zhang et al. (Ref. 6) obtained, in a com-parative study of two HSLA steels with andwithout niobium, based on the physical sim-ulation of thermal cycles in a Gleeble simu-lator, a microstructure with grain boundaryferrite, acicular ferrite, and degeneratepearlite at low cooling rates applied in theno Nb steel specimens, while bainite isformed at fast cooling rates. For the speci-mens of the Nb-bearing steel, the referredauthors obtained granular bainite as a dom-inant microstructure over a wide range ofcooling rates, with martensite appearing atvery fast cooling rates. Although for coolingrates slower than 32°C/s, the initial trans-formation temperature decreases by about20°C in the presence of niobium. Zhang etal. (Ref. 7), in another article, addressed the

Effect of Tempering Pass on HSLA-80 Steel HAZ Microstructures

Continuous cooling transformation diagrams devised from a simulated thermalcycle were applied over a high-temperature, modified Nb-bearing

steel microstructure

BY A. CRUZ-CRESPO, D. BEZERRA DE ARAUJO, AND A. SCOTTI

ABSTRACT

The alloying system and thermal history of the hot rolling process applied to high-strength low-alloy steels (HSLA) leads to a very particular behavior of these materi-als under welding thermal cycles. In this work, microstructures and hardness of a grain-coarsened heat-affected zone (HAZ) were analyzed from API 5L X80 Nbmicroalloyed steel specimens after undergoing simulated thermal cycles to representboth first and tempering passes. The first thermal cycle for each sample reached thepeak temperature of 1350°C, while the second was of 950°C. Using the different cool-ing curves imposed by the simulator, a continuous cooling transformation (CCT) dia-gram was raised for both conditions. The predominant microstructure for the firstthermal cycle was granular bainitic ferrite at low cooling rates, but it changed intobainitic ferrite as cooling rate increases, reaching some presence of martensite at thehighest cooling rates. The microstructure in the second thermal cycle is quasi-polygo-nal ferrite at low cooling rates and bainitic ferrite at the fastest cooling rates. How-ever, the hardness did not exceed 300 HV in any case and the hardness measured waseven lower in the simulated tempering pass specimens. These results indicate that thissteel has high weldability and no special techniques, such as preheating, need to beemployed to prevent cold cracking. However, the study suggests the need for futurework on aging of precipitates in this grain-refined region due to the tempering pass.

A. CRUZ-CRESPO, D. BEZERRA DEARAUJO, and A. SCOTTI([email protected]) are with the Centerfor Research and Development of WeldingProcesses, Federal University of Uberlandia,Brazil.

KEYWORDS

High-Strength Low-Alloy(HSLA)

Nb-Microalloyed SteelThermal CycleSimulationHeat-Affected Zone (HAZ)Tempering Pass

OCTOBER 2013, VOL. 92

Cruz-Crespo Suppment October 2013_Layout 1 9/13/13 4:39 PM Page 304

effect of niobium on the microstructure, onthe coarse-grained HAZ properties, and onthe starting transformation temperatureunder the effect of different energy inputs.Shome (Ref. 8), while also using simulation,established that the peak temperaturemakes the austenitic grain size increase lin-early, influencing the HAZ structure andproperties.

While it is true that the effects of thechemical composition and welding energyon the resulting microstructure of thecoarse region of the microalloyed steelHAZ have been relatively well studied,specialized literature shows little or nodata regarding the effect of the “temper-ing pass” on the coarse region. The mainreason for this is that the coarse-grainedregion is always of concern in studies onweldability of structural steels, as a regionsusceptible to loss of toughness and asbeing a facilitator of hydrogen cracking.The microstructure refined by the “tem-pering pass” technique is usually quitebeneficial for reducing hardenability andincreasing toughness. In addition to notbeing known in detail, the microstructureresulting in the refined region can be ex-pected for steels with higher contents ofNb to present aging of the precipitates.Therefore, the objective of this study wasto evaluate the microstructural formationof a tempering pass on the coarse-grainedregion of the HAZ of Nb microalloyedHSLA-80 steel, through simulating weld-ing thermal cycles from a first pass fol-lowed by a tempering pass.

Materials and Methods

The methodology approach of this workwas the imposition of thermal cycles repre-

senting welding withdifferent energiesthrough simulation onsamples of a microal-loyed steel of highstrength (HSLA), TypeAPI 5L X80, composi-tion being reported inTable 1. The coarse-grained region of theHAZ was initially simu-lated, assuming a peaktemperature of 1350°Cand cooling rates rang-ing from 3° to 130°C/s.Half of the specimensthat underwent thissimulation were heatedagain to 950°C to simu-late the effect of heattreatment that is car-ried out by a secondbead over the region ofthe coarse-grainedHAZ (CGHAZ) of aprecedent bead (tem-pering pass).

As shown in Fig. 1, the base metal has amicrostructure with high prevalence ofpolygonal ferrite. Alongside the polygonalferrite, dark-etching regions are present, re-sulting from transformations of the residual

austenite at relatively low temperatures(probably composed of bainite and MA mi-croconstituent). This microstructure pat-tern coincides with that reported by Jing-Hong et al. (Ref. 2), Cizek et al. (Ref. 9), andBott et al. (Ref. 10), who deal with the char-

305-sWELDING JOURNAL

WE

LD

ING

RE

SE

AR

CH

Fig. 1 — HSLA microstructure of the Nb microalloyed steel (API 5L X-80). Fig. 2 — Cooling and heating curves experienced by the specimens to simu-late the various thermal cycles of the first pass in the region of the coarse grainHAZ (peak temperature = 1350°C) for different diameters of the central partof the specimens: d = 10 mm (ϕ8/5 = 3.4°C/s); d = 9 mm (ϕ8/5 = 3.8°C/s); d= 8 mm (ϕ8/5 = 5.4°C/s); d = 7 mm (ϕ8/5 = 13.9°C/s); d = 6 mm (ϕ8/5 =51°C/s); d = 5 mm (ϕ8/5 = 126.6°C/s).

Table 1 — Chemical Composition of the Nb Microalloyed Steel (API 5L X-80), in % Mass

C Mn Si P S Mo Ni Cr Cu

0.04 1.71 0.25 0.009 0.004 0.006 0.156 0.202 0.214

V Nb Ti Al N B V + Nb + Ti0.003 0.100 0.009 0.018 0.005 0.0001 0.112

Fig. 3 — Cooling and heating curves experienced by the specimens to simu-late the various thermal cycles imposed by the second bead on the region ofthe coarse-grain HAZ of the precedent bead (peak temperature = 950°C) fordifferent diameters of the central part of the specimens: d = 10 mm (ϕ8/5 =3.7°C/s); d = 9 mm (ϕ8/5 = 5.3°C/s); d = 8 mm (ϕ8/5 = 15.5°C/s); d = 7 mm(ϕ8/5 = 52.2°C/s); d = 6 mm (ϕ8/5 = 95.2°C/s); d = 5 mm (ϕ8/5 = 151.5°C/s).


acterization of steels of similar compositionto the present work. Those authors statedthe existence of MA microconstituent asso-ciated with the ferritic grain boundaries,which were revealed only under electronmicroscopy. Other authors, such asAlmeida et al. (Ref. 1), Gorni and Mei (Ref.11), and Shanmugam et al. (Ref. 12), reporthaving found, by means of electron mi-croscopy, some presence of bainite together

with ferrite in steels with compositions rel-atively similar to those shown in Table 1.

These authors associate this fact withthe synergistic effect that presents itselffor particular relations of microalloyingelements. An important group of authors,such as Mishra (Pathak) et al. (Ref. 13)and Moon et al. (Ref. 14), also confirmedthe phases and microconstituents alreadymentioned in the microstructure of steels

of similar composition to that in the pres-ent study, together with the presence ofniobium-rich precipitates associated withthe grain boundaries, which undoubtedlycontribute to the high mechanical resist-ance of these steels.

The simulating of the thermal cyclingwas performed using a simple equipmentdesign (described in Vilarinho and Araujo(Ref. 15)) based on the Joule effect, which

OCTOBER 2013, VOL. 92306-s

WE

LD

ING

RE

SE

AR

CH

Fig. 4 — Corresponding microstructures for the first cycle (A1, B1, C1, D1, E1, F1 ) and second cycle (A2, B2, C2, D2, E2, F2) with different cooling rates (dif-ferent specimen diameters), where αq = quasi-polygonal ferrite, αB = granular bainitic ferrite, αB

o = bainitic ferrite, αW = Widmansta ̈tten ferrite, α´M =martensite [microconstituent nomenclatures based on Krauss and Thompson (Ref. 18)].

E2 — peak = 950° and ϕ8/5 = 95.2 °C/s F1 — peak = 1350° and ϕ8/5 = 126.6 °C/s F2 — peak = 950° and ϕ8/5 = 151.1 °C/s

D1 — peak = 1350° and ϕ8/5 = 13.9 °C/s D2 — peak = 950° and ϕ8/5 = 52.2 °C/s E1 — peak = 1350° and ϕ8/5 = 51.0 °C/s

B2 — peak = 950° and ϕ8/5 = 5.3 °C/s C1 — peak = 1350° and ϕ8/5 = 5.4 °C/s C2 — peak = 950° and ϕ8/5 = 15.5 °C/s

A1 — peak = 1350° and ϕ8/5 =3.4 °C/s A2 — peak = 950° and ϕ8/5 =3.7 °C/s B1 — peak = 1350° and ϕ8/5 = 3.8 °C/s)

Cruz-Crespo Suppment October 2013_Layout 1 9/16/13 9:31 AM Page 306

enables rapid heating of specimens and al-lows for natural cooling of the specimenswith the aid of aluminum heat sinks thatalso function as support and electricalcontact. An electronic controlling devicedisconnects the power source when thepreregulated temperature at the center ofthe test pieces is reached. The thermo-couple placed in the center of the samplealso records the cooling, which due to thesmall dimensions of the sensor (Type K,diameter 0.5 mm) and the characteristicsof continuous natural cooling can detectthe starting and finishing points of themetallurgical transformations.

The specimens were cylindrical, 150mm long, with 10 mm diameter at the ex-tremes. In the central part, along a 10-mmlength, different diameters (5–10 mm)were machined, allowing for variances inthe cooling rate over a wide range. Thesegeometric measures were determined bythe finite element method (Ansys) for thedesired cooling rates. One disadvantage ofthis technique is that in order to obtainvery fast cooling rates, the specimen di-ameter has to be small, and likewise the

material volume from which the thermalcycle is measured. This characteristicmakes the sensitivity of the method lowerto detect activation energy for the startingand finishing of transformations. Cross-sections of the specimens under differentcooling rates were cut in the central regionfor metallographic analysis (2% Nitaletching) and microhardness (10 indenta-tions with a load of 500 g and a load appli-cation time of 10 s).

To determine the starting and finishingpoints of the transformations, a differen-tial analysis has been used, as described byZachrisson (Ref. 16). This analysis con-sists of performing a regression of a partof the experimental cooling curves (thecurve that best fits is the exponential typeTreg = aebt + cedt, where a, b, c, and d arecoefficients obtained by regression analy-sis and where “t” is the time) and extrap-olate it so as to include temperaturesbelow transformation. The difference be-tween the thermal cycle experimentalcurve (T = f(t)) and the regression curvein the region where deviation is perceived(D(T)i = Ti – Tregi) is plotted in relation

to temperature, thereby obtaining thetemperatures of the transformations. Asimilar technique was successfully appliedby Alexandrov and Lippold (Ref. 17), butrather than using determined exponentialequations as reference thermal cycles,they used calculated (analytic formulas ornumerical modeling) ones for generatingthe references.

Results and Discussion

Thermal Cycles of the Coarse-GrainRegion of the HAZ during the First andTempering Passes

The curves of thermal cycles to simu-late the effect of the first bead (peak tem-peratures 1350°C) are shown in Fig. 2.Analogously, the thermal cycles to simu-late the effect of a second weld pass (peaktemperature of 950°C), that is, the over-heating of a region affected by the heat ofthe precedent pass (coarse-grain regionof the HAZ, in this case), are shown inFig. 3. Inflections in the curves shown inFigs. 2 and 3 are observed (more distinct

307-s

WE

LD

ING

RE

SE

AR

CH

WELDING JOURNAL

Table 2 — Resulting Cooling Rates between 800° and 500°C (ϕ8/5) that the Specimens Experienced after Heating up to 1350° and 950°C, toSimulate a First Pass and the Subsequent Effect of a Tempering Pass, Respectively, under Different Simulated Heat Inputs

Specimen center diameters (mm) Cooling rate (°ϕ8/5) from first Cooling rate (°ϕ8/5) from temperingpass CGHAZ simulation pass simulation on the CGHAZ

5 126.6°C/s 151.5°C/s6 51°C/s 92.2°C/s7 13.9°C/s 52.2°C/s8 5.4°C/s 15.5°C/s9 3.8°C/s 5.3°C/s10 3.4°C/s 3.7°C/s

Note: Cooling rates after heating up to 1350°C to simulate a first pass and then heating the same specimen to 950°C to simulate the effect of a subsequent temperingpass.

Fig. 5 — Microhardness after both the first and sec-ond thermal cycles as a function of the cooling ratebetween 800° and 500°C.

Fig. 6 — CCT diagrams resulting from the first ther-mal cycle (simulating the coarse-grain HAZ region):The martensite starting temperature (Ms ≈ 465°C) isa predicted value, calculated from Andrews’s equa-tion Ms(°C) = 539 – 423 C – 30.4 Mn – 17.7 Ni –12.1 Cr –7.5 Mo, employed by Zhang et al. (Ref. 7).

Fig. 7 — CCT diagrams resulting from the secondthermal cycle (simulating the refinement of the coars-ened-grain region of the HAZ-tempering pass): Themartensite starting temperature (Ms ≈ 465°C) is apredicted value, calculated from Andrews’s equationMs(°C) = 539 – 423 C – 30.4 Mn – 17.7 Ni – 12.1Cr – 7.5 Mo, employed by Zhang et al. (Ref. 7).


for slower cooling rates, justified mainlybecause of a larger metal volume in theregion of temperature measurementsand a not high enough dynamic responsefrom the thermocouple — 0.5-mm wires).These are characteristics of the activationenergy for the start and finish of thetransformation in ferritic steels. Thecooling rates (ϕ8/5) in the range of tem-peratures from 800° to 500°C are summa-rized in Table 2 and also reported in thecaptions of the thermal cycle plots.

As seen, the cooling rates becomeproportionately faster for the thermalcycle with peak temperature of 950°Ccompared to the first cycle with peaktemperature of 1350°C. Such behavior istypical of simulations as in the presentwork (less heat to dissipate), but thatshould also be observed in an actualwelding setting.

Influence of the First and SecondThermal Cycles on the Microstructure ofthe Coarse Region HAZ

The difference in behavior of the trans-formations during cooling of the speci-mens with different cooling rates leads tovariations in the microstructure of somewith regard to others. (The microstructureidentification in this work was based on anomenclature for ferritic microcon-stituents taken from the Atlas for BainiticMicrostructures, developed by the Iron andSteel Institute of Japan Bainite Commit-tee and described by Krauss and Thomp-son (Ref. 18).) It is important to mentionthat microconstituent nomenclature forcarbon steel HAZ is still not standardized,but a discussion in this direction or anyproposal of microstructure nomenclatureis out of focus for this paper. Figure 4shows, side by side, the microstructures ofthe specimens under different thermal cy-cles, at 1350°C (simulating the first pass)on the left, and 950°C (simulating the tem-pering pass) on the right. In relation to theoriginal microstructure of steel (Fig. 1),one perceives greater grain sizes in the re-gion heated to 1350°C (Fig. 4, left). Thegrain growth for this type of steel resultingfrom a range of peak temperatures hasbeen studied by Kuziak et al. (Ref. 19) andShome (Ref. 8), for whom up to a temper-ature of about 1100°C the austenite grainpractically does not grow, but then in-creases almost linearly at a high rate. Thisphenomenon could be motivated by thedissolution of the precipitates rich in Nb asa function of the elapsed time at tempera-tures above 1100°C. In contrast, Ivanov etal. (Ref. 20) show a gradual growth of thegrain at peak temperature.

Figure 4A1 and Fig. 4B1 show the mi-crostructure of the coarse-region HAZwith slow cooling (3.4° and 3.8°C/s), char-acterized by the predominant presence of

granular bainitic ferrite (αB). This result isconsistent with those reported by Zhang etal. (Refs. 6, 7), Gorni and Mei (Ref. 11),and Ivanov et al. (Ref. 20), who obtainedCCT diagrams for steels with relativelysimilar compositions to that of the presentwork, where the transformation fromaustenite to bainite is predominant. Andfor the cooling rate of 5.4°C/s, shown onthe microstructure (Fig. 4C1), alongsidethe granular bainitic ferrite (αB) there isthe presence of bainitic ferrite of type αB

o

and Widmanstätten ferrite (αW). For evenfaster cooling rates (13.9° to 51.0°C/s), thepresence of bainitic ferrite of type αB

o iseven more visible (Fig. 4D1 and E1). Inthe case of 4D1, there is still some granu-lar bainitic ferrite αB. Whenever a peaktemperature above the dissolution of theniobium-rich precipitates is reached,which, according to Zhang et al. (Ref. 6)and Moeinifar et al. (Ref. 4), is less than1200°C, this element (Nb), in solid solu-tion for such cooling rates, obstructs thekinetics of ferrite grain formation. Thishappens because solute drags at theaustenite grain boundary, thereby pro-moting the formation of bainitic ferritetype αB

o. In the particular case of the cool-ing rate of 126.6°C/s, Fig. 4F1 shows thepresence of martensite (αM) mixed withbainitic ferrite type αB

o. The presence ofbainitic ferrite (αB

o) with martensite(α´M) in the microstructure at faster cool-ing rates is consistent with that reported byZhang et al. (Ref. 6) and Ivanov et al. (Ref.20), who show a CCT diagram (obtainedfor a steel composition relatively similar tothe present study) with martensitic trans-formation at lower temperatures when thecooling rate exceeds 50°C/s.

According to the manifestation of alarge number of authors, among them Shiand Han (Ref. 21), transformation ofaustenite to bainite takes place during thecooling process experienced by the differ-ent specimens (different cooling rates).This occurs through a migration of carboninto the interior of the austenite grains,thereby stabilizing the austenite and fa-voring martensite transformation at lowertemperatures, which leads to the appear-ance of MA microconstituent. These au-thors show that to the extent that t8/5 in-creases (with decreasing cooling rate) theMA fraction increases.

The microstructures corresponding tothe coarse region HAZ after undergoing asecond thermal cycle with different coolingrates are shown to the right in Fig. 4. Thepeak temperature reached the normaliza-tion temperature (grain refinement) for thesteel under study, considering the high heat-ing rates. In all cases, there is a tendency to-ward the formation of a finer microstruc-ture as compared to the first heat cycle, aswould be expected by the thermal historyimposed (even considering starting the

transformation from coarse grains, thecomparison with Fig. 1 shows the grains willbe as fine as those of the base metal, sug-gesting the action of restriction to graingrowth of Nb carbonitrides).

If one compares the microstructure ofFig. 4A2, corresponding to the slowestcooling rate of the second thermal cycle,with that of Fig. 1, corresponding to thebase metal, a high similarity is noted. Thismicrostructure is characterized by a highpredominance of quasi-polygonal ferrite(αq). In such cases, the most notable dif-ference is that, in the base metal (Fig. 1),the grains have a certain orientation dueto the effect of the controlled thermome-chanical rolling process and the ferritegrain boundaries are smooth and continu-ous, in that for the lowest cooling rate inthe second cycle (Fig. 4A2) the grainspresent irregular grain boundaries.

The presence of quasipolygonal ferrite(αq) is noted in the microstructure afterthe second cycle at low cooling rates —Fig. 4A2, B2, and C2. In the particularcase of 4C2 (cooling rates of 15.5°C/s)alongside the quasipolygonal ferrite (αq)is the presence of granular bainitic ferrite(αB). For faster cooling rates (52.2°C/s),the microstructure is granular bainitic fer-rite (αB) mixed with bainitic ferrite typeαB

o — Fig. 4D2. And for cooling rates of95.2° and 151.2°C/s, the microstructure ispredominantly bainitic ferrite (αB

o) —Fig. 4E2 and 4F2. The fact that no marten-site appears in the microstructure (Fig.4F2), despite a faster cooling rate(151.5°C/s) than in the first cycle(126.6°C/s), is explained, according toShome and Mohanty (Ref. 22), as the re-duction of austenite grain size in the sam-ple under the second thermal cycle, whichreduces the hardenability.

The Influence of the First and SecondThermal Cycles on the Microhardness ofthe Coarse-Region HAZ

The microhardness values obtained forthe HAZ grain-coarsened region, corre-sponding to different cooling rates be-tween 800° and 500°C during the thermalcycles simulating the first bead and thetempering pass, are shown in Fig. 5. It be-comes evident that, in general, the micro-hardness tends to increase with an in-crease in the cooling rate, which isexplained mainly by the modifications un-dergone in the microstructure, which arecovered in detail in the previous section.(The effect of grain sizes on hardness, i.e.,a faster cooling rate would lead to smallergrains, which, in turn, usually presentgreater hardness, cannot be neglected inthis analysis, yet with less significance,considering the fact that the found hard-ness for the tempering pass is lower thanthat of the first pass, although the temper-


WE

LD

ING

RE

SE

AR

CH


ing grain size is smaller.) Such behaviorhas been reported by a large number of au-thors for steels of similar composition tothat used in the present work, includingsuch authors as Zhang et al. (Ref. 6),Gorni and Mei (Ref. 11), Cizek et al. (Ref.9), and Ivanov et al. (Ref. 20), which is ob-viously also justified based on the changesin the microstructure. In the particularcase of Gorni and Mei’s work, one obtainsthe behavior of the Vickers hardness as afunction of cooling rate on a logarithmicscale, showing that a steel of lower carboncontent (0.04%) hardness is practicallyconstant for low values of cooling rate,growing for faster cooling rates, acquiringa parabolic character, similar to thatshown in Fig. 5 of this work.

The microhardness of the base metal(218 ± 5 HV0.5) is slightly superior to thatobtained after the thermal cycles simulat-ing the first pass with slow cooling rate(208 ± 3 HV0.5, 211 ± 3 HV0.5, 206 ± 4HV0.5, 213 ± 3 HV0.5, corresponding to3.4°, 3.8°, 5.4°, and 13.9°C/s, respectively— Fig. 5). This fact is more associated withthe elapsed time at high temperatures,above the dissolution temperature of theNb-rich precipitates, than with the coolingrates between 800° and 500°C. After a suf-ficient elapsed time at high temperaturesfor dissolution to occur, the cooling rate issuch that it prevents the diffusive processand the reprecipitation on the grainboundary, which make a significant hard-ening effect. The loss of the hardening ef-fect of the precipitates is not fully com-pensated for by the hardening effect of thetransformation, which takes place uponcooling due to the thermal cycle.

Interestingly, the particular case ofspecimen microhardness under the cool-ing rate of 5.4°C/s during the first thermalcycle, which although statistically similarto the specimens with slower rates, isslightly lower in terms of its average value.If observed in detail in the microstructureof Fig. 4A1, B1, and C1, one can perceivethat, alongside the granular bainitic ferrite(αB), bainitic ferrite (αB

o), and Wid-manstätten ferrite (αW) microstructuresappeared in the form of coarse plates forthe lowest cooling rate (5.4°C/s), causing asoftening. However, it is appropriate to re-iterate that the softening of the metal isnot significant from the point of view ofthe statistical value variability.

Zhang et al. (Ref. 6) suggest that thehardness of the granular bainitic ferrite(αB) ranges between 210 and 240 HV, sothe microhardness values obtained forspecimens with slower cooling rates (3.4°,3.8°, 5.4°, and 13.9°C/s) in the first heatcycle (Fig. 5) fully correspond to the pres-ence of granular bainitic ferrite (αB) ob-served for these specimens in Fig. 4A1,B1, C1, and D1. For their part, Moeinifaret al. (Ref. 4), who reported the predomi-

nance of granular bainitic ferrite (αB) inthe microstructure of the coarse-grainedregion of the HAZ, obtained microhard-ness values near the maximum range forthis already mentioned microconstituent(around 240 HV). The difference betweenthe microhardness values for granularbainitic ferrite (αB) obtained by these au-thors and those obtained in the presentwork is related to the higher carbon con-tent in the steel for the aforementionedwork.

For the highest cooling rates (51.0°C/sand 126.6°C/s) during the first thermalcycle, the growth of the microhardness(Fig. 5) is undoubtedly linked to the highpresence of bainitic ferrite (αB

o) (Fig.4E1), bainitic ferrite (αB

o), and martensite(α´M) — Fig. 4F1. Such an increase in themicrohardness together with an increasein the cooling rate could also be related toincreasing the volume fraction of the MAmicroconstituent. As such, Moeinifar etal. (Ref. 4) concluded that increasing thecooling rate increases the fraction of theMA microconstituent, which, in this case,plays a governing role on the hardnessprogress in the region of the coarse grainin the HAZ.

Zhang et al. (Ref. 6) state that marten-site (α´M) presents microhardness valuesof 320 to 340 HV, which were not reachedin this work, even for the highest coolingrate (126.6°C/s) during the first thermalcycle — Fig. 5. Ivanov et al. (Ref. 20) de-clared the presence of martensite for highcooling rates, obtaining a microhardnessvalue of around 325 HV at 50°C/s, whenthe peak temperature is 1350°C. Such ahigh value in microhardness obtained bythe above authors, in relation to this work,is primarily related to the higher carboncontent (0.06%) compared to that re-ported in Table 1.

From Fig. 5, the second thermal cycle(simulated tempering pass) also leads toan increased microhardness as the coolingrate is faster, consistent with the modifi-cations undergone in the microstructure(Fig. 4A2–F2). Such an increase in the mi-crohardness could also be related to thepossible increase in the volume fraction ofthe MA microconstituent. When compar-ing microhardness as a function of coolingrate between conditions after the firstcycle and after the second cycle, it is clearthat there is regularity in behavior be-tween the two conditions, but lower valuesfor the second thermal cycle (the differ-ence becomes more evident as the coolingrate increases). The parabolic characterobtained for the microhardness, both afterthe first thermal cycle and the second, iscompletely coincident with results ob-tained by Gorni and Mei (Ref. 11). Thedecrease in microhardness is a general in-dication that the second thermal cycle hasa positive effect on the coarse-grained

HAZ region. This confirms the reductionof hardenability due to the finer granular-ity achieved in the recrystallization zone ofthe coarse HAZ under the second cycle.

Obtaining Continuous CoolingTransformation (CCT) Curves for the Firstand Second Thermal Cycles of theCoarse-Region HAZ

After thermal cycling of Figs. 2 and 3,the continuous cooling transformation di-agrams referring to the first and the sec-ond thermal cycles were obtained, as seenin Figs. 6 and 7. In Fig. 6, it can be seenthat the changes during the first thermalcycle occur at temperatures lower thanthose reported by several authors for APIX80 steels, amongst them Zhang and Far-rar (Ref. 23), Zhao et al. (Ref. 24), Stal-heim et al. (Ref. 25), Cizek et al. (Ref. 26),Jing-Hong et al. (Ref. 2), and Liu et al.(Ref. 27). However, these authors haveworked with diagrams related to thermaltreatment processes, from which it followsthat the heating is slower and with longerelapsed time in the austenitic region. Thisdrop in temperature of phase transforma-tion in the case of thermal cycles of theHAZ has been reported by Zhang et al.(Ref. 6). Zhang et al. also showed that asone increases the cooling rate there is afurther decrease in the transformationtemperature, due to lower elapsed timesabove 900°C, limiting the diffusiveprocesses of phase transformation. Ac-cording to Zhang et al. (Ref. 6), and Gorniand Mei (Ref. 11), as the cooling rate be-tween 800° and 500°C grows there is alsoan improvement in the stability of theaustenite, which lowers the transforma-tion temperature.

In total agreement with the CCT dia-gram of Fig. 6, the curves of Fig. 2 suggestthat diffusional transformations (austen-ite to bainite) take place for any coolingrate (manifested by sensitive variation oflatent heat, causing deflection in therecorded thermal cycle). As observed inFig. 2, deflections were not observed onlyin the curve of the fastest cooling rate.Latent heat of austenite to martensitetransformation is not well divulged incurrent literature. Radaj (Ref. 28, p. 292)cites other sources to say that the latentheat of the austenite-pearlite transfor-mation is 92 J/g and the latent heat of theaustenite-martensite transformation is83 J/g for a 1.2% C. Even considering thehigher content of C and no data providedfor austenite-bainite transformation, thedifferences in the figure are not remark-able to eliminate the deflection. In addi-tion, Alexandrov and Lippold (Ref. 17)detected this transformation measuringin-situ weld metal continuous coolingtransformation. Thus, the main reasonfor not detecting the latent heat of anaustenite to martensite transformation in


WE

LD

ING

RE

SE

AR

CH


the cooling curves might be the limitationof the simulation (low metal volume inthe region of temperature measurementsand not too fast dynamic response fromthe thermocouple).

Upon analysis of microstructure re-sulting from transformations (Fig.4A1–F1), the presence of bainite is de-clared in all specimens (coarse HAZ dur-ing the first thermal cycle). This result isalso coincident with those reported byZhang et al. (Ref. 6), Gorni and Mei(Ref. 11), and Ivanov et al. (Ref. 20), whoobtained CCT diagrams for steels of rel-atively similar composition to the pres-ent work, where the transformation fromaustenite to bainite was predominant.The presence of martensite (α´M), (Fig.4F1) evidenced by hardness close to 300HV and by the martensite starting tem-perature line in Fig. 6, may be possible iftogether with bainitic ferrite (αB

o) in themicrostructure, the curve deflection isstill present. This is consistent with thatreported by Gorni and Mei (Ref. 11) andZhang et al. (Ref. 6) for steels similar tothe one in this work. From these authors,the martensitic transformation defined isbased either on the microstructure andhardness, in the case of Zhang et al. (Ref.6), or established in calculated (pre-dicted) CCT diagrams, as in the case ofGorni and Mei (Ref. 11) at a constanttemperature. For their part, Ivanov et al.(Ref. 20) also obtained a CCT diagramthat shows the martensitic transforma-tion temperature for high cooling rates,without declaring how they define thistemperature. Zhang et al. (Ref. 6) andIvanov et al. (Ref. 20) show the marten-sitic transformation as being possiblewhen the cooling rate exceeds 50°C/s,which is in line with the present work. Ifone compares the CCT diagram repre-senting the coarse-grain HAZ as for thefirst pass (Fig. 6) with that obtained ac-cording to the second thermal temperingpass cycle (Fig. 7), one can see that in thislast mentioned there also appears atransformation of austenite to ferritecurve at low cooling rates, which obvi-ously occurs with a higher level of diffu-sion. Data from Gorni and Mei (Ref. 11)coincide with this result, which states theferrite transformation at low coolingrates, below 0.5°C/s. The difference incooling rates at which this transforma-tion occurred in the referred work and inthis present work may be linked to a highretention in the austenite region, whichundoubtedly increases the stability ofthis phase. Also, Ivanov et al. (Ref. 20)agree with the possible occurrence offerrite transformation, while reporting aCCT diagram where this transformationis reflected even at cooling rates above50°C/s, which could be associated prima-rily with higher carbon content (0.06%)

in relation to that of the present work(0.04%).

Bainitic transformation for the secondthermal cycle (tempering pass) (Fig. 7) oc-curs at higher temperatures than in thefirst — Fig. 6. The foregoing is also relatedto the elapsed time in the austenitic re-gion. In the first cycle, there is more timespent in the austenitic region than in thesecond due to a higher peak temperature(1350°C for the first cycle and 950°C forthe second) leading to increased solubilityand slow transformation temperaturesduring cooling. The presence of marten-site in the tempering pass would not bepossible within the cooling range under in-vestigation (Fig. 7), which agrees with themicrostructures obtained — Fig. 4.

It is important to point out that An-drews’s equation for MS temperature pre-diction used in Figs. 6 and 7 was not de-veloped for steels with such a low-Ccontent. However, Capdevila (Ref. 29)points out that although the relationshipbetween the martensite start transforma-tion temperature and steel compositionhas been investigated previously by sev-eral researchers (for instance, Grange andStewart, 1946; Payson and Savage, 1944;and Kung and Rayment, 1978), it was thestudy by Andrews (1965) that has provedto be the most reliable, because it consid-ered the largest number of samples. Sour-maila and Garcia-Mateo (Ref. 30) show acomparison between their proposed neu-tral network model results and the predic-tions from Andrews and concluded thatthe neural network model performs atleast equally as well as the thermodynamicapproach. Applying Sourmaila and Gar-cia-Mateo’s model for the steel of thiswork, a very similar MS temperature wasfound, i.e., 460°C.

General Discussion

Considering the different behaviorsbetween metallurgical characteristics (mi-crostructure and hardness) when appliedto a first cycle peak temperature of 1350°Cand a second cycle on the coarse-grainedzone formed (peak temperature of900°C), this can be explained by the theoryand supported by results from other au-thors, demonstrating the validity of the ap-plied simulation. The possibility of work-ing with natural cooling in the region ofstudy and application of the technique ofdifferential analysis made it possible todraw up CCT diagrams for different ther-mal experiences undergone by the steel,similar to what happens in welding.

The CCT diagrams show that for themicrostructure of the coarse-grainedHAZ zone, either the primary formation(simulating the first pass) or recrystalliza-tion (simulating a second pass or temper-ing pass) are mainly granular bainitic fer-

rite (αB) or bainitic ferrite (αBo). There is

a perceived improvement in properties ofthe coarse HAZ region when tempered byreducing the grain size and decrease inhardness (the hardenability is lower),showing that the steel in question also hashigh weldability. From the standpoint ofhardness, which according to Ivanov et al.(Ref. 20) has a direct relationship with themechanical properties for HSLA steels,only a very high cooling rate, greater than100°C/s, puts the union at risk due to hy-drogen-assisted cold cracking initiated inthe HAZ. That is, the material could bewelded by a very large range of processesand parameters, and no preheating proce-dure is needed.

Continuous cooling transformationdiagrams could be obtained for other re-gions of the HAZ (including HAZ in thebead metal recrystallized by precedentpasses), facilitating the programming ofsequences of passes to optimize the jointproperties. On the other hand, althoughit was not verified in the work, carboni-tride precipitates or the forming of mi-croconstituent MA could happen, espe-cially in HAZ regions in which thetemperature was below the dissolutiontemperatures for Nb-bearing precipi-tates. These microconstituents can dete-riorate properties of the HAZ, especiallyconcerning toughness. This means that,despite the high weldability demon-strated by this study for this steel, furtherinvestigation on all regions of the HAZ isadvised and possible with the same ex-perimental approach.

Conclusions

1. The system used to evaluate the effectof thermal cycling on the microstructureformation of the HAZ of a HSLA steel(simulator and method of differential analy-sis of cooling curve) was effective and prac-tical, sensitive enough to produce CCT dia-grams of different regions of the HAZ.

2. For the CCT diagrams of the originalHAZ (the first heat cycle) of the steel understudy, the microstructure is predominantlygranular bainitic ferrite (αB) at low coolingrates, but is transformed into bainitic ferrite(αB

o), to the extent that increases in thecooling rate made it finer. Even for veryhigh cooling rates, the microstructure is pre-dominantly bainitic ferrite (αB

o) with thepresence of martensite. In all cases, the mi-crohardness is less than 300 HV, qualifyingthe steel studied as of good weldability.

3. From the point of view of the basic mi-crostructure, the CCT diagram shows thatthe recrystallized region (simulating thetemperature for a second pass) further im-proved the weldability of the material understudy, by refining the grain while reducingthe hardness, even though not significantlyaltering the type of microstructure.


WE

LD

ING

RE

SE

AR

CH


Acknowledgments

The authors thank engineers LuciaBasilio P. de Souza and Ivy Jorge Francofor their support in performing the ther-mal simulation of the specimens. Theyalso are pleased to thank the Brazilianagency CAPES, which, by means of proj-ect No. 059/2009, financially supportedthe experimental work and enabled inter-action between the researchers involved.

References

1. Almeida, D. M., Maciel, T. M., andBracarence, A. Q. 2004. ZTA fracture compar-ative study of bainitic and pearlitic API steelafter implant test. Soldagem & Inspeção 9(4):198–204 (ISSN 0104-9224 printed/ISSN 1980-6973 online) (in Portuguese).

2. Jing-Hong, Y., Qing-You, L., Dong-Bai,S., and Xiang-Yang, L. 2010. Microstructureand transformation characteristics of acicularferrite in high niobium-bearing microalloyedsteel. Journal of Iron and Steel Research, Inter-national 17(6): 53–59.

3. Das, A. K. 2010.The present and the fu-ture of line pipe steels for petroleum industry.Materials and Manufacturing Processes 25(1-3):14–19 (DOI: 10.1080/10426910903202427).

4. Moeinifar, S., Kokabi, A. H., and MadaahHosseini, H. R. 2011. Role of tandem sub-merged arc welding thermal cycles on proper-ties of the heat-affected zone in X80 microal-loyed pipe line steel. Journal of MaterialsProcessing Technology 211(3): 368–375(doi:10.1016/j.jmatprotec.2010.10.011).

5. Mohandas, T., Madhusudan, R. G., andKumar, B. S. 1999. Heat-affected zone soften-ing in high-strength low-alloy steels. Journal ofMaterials Processing Technology 88(1-3): 284–294 (doi:10.1016/S0924-0136(98)00404-X).

6. Zhang, Y. Q., Zhang, H. Q., Liu, W. M.,and Hou, H. 2009(a). Effects of Nb on mi-crostructure and continuous cooling transfor-mation of coarse grain heat-affected zone in610 MPa class high-strength low-alloy struc-tural steels. Materials Science and Engineering A499, p. 182–186.

7. Zhang, Y. Q., Zhang, H. Q., Li, J. F., andLiu, W. M. 2009(b). Effect of heat input on mi-crostructure and toughness of coarse grainheat-affected zone in Nb microalloyed HSLAsteels. Journal of Iron and Steel Research, Inter-national 16(5), p. 73–80 (doi:10.1016/S1006-706X(10)60014-3).

8. Shome, M. 2007. Effect of heat-input onaustenite grain size in the heat-affected zone ofHSLA-100 steel. Materials Science and Engi-neering A, p. 454–460.

9. Cizek, P., Wynne, B. P., Davies, C. H. J.,Muddle, B. C., and Hodgson, P. D. 2002. Effectof composition and austenite deformation on the

transformation characteristics of low-carbon andultralow-carbon microalloyed steels. Metallurgi-cal and Materials Transactions A 33(5): 1331–1349(doi: 10.1007/s11661-002-0059-8).

10. Bott, I. S., De Souza, L. F. G., Teixeira,J. C. G., and Rios, P. R. 2005. High-strengthsteel development for pipelines: a Brazilianperspective. Metallurgical and Materials Trans-actions A Vol. 36A, p. 443–454(ISSN 1073-5623 Printed/ISSN 1543-1940 On-line) (10.1007/s11661-005-0315-9).

11. Gorni, A. A., and Mei, P. R. 2004. Austen-ite transformation and age hardening of HSLA-80 and ULCB steels. Journal of Materials Process-ing Technology, Vol. 155–156, p. 1513–1518(doi:10.1016/j.jmatprotec.2004.04.245).

12. Shanmugam, S., Ramisetti, N. K., Misra,R. D. K., Hartmann, J., and Jansto, S. G. 2008.Microstructure and high strength–toughnesscombination of a new 700 MPa Nb-microal-loyed pipeline steel. Materials Science and Engi-neering A, 478 (1-2): 26–37 (doi:10.1016/j.msea.2007.06.003).

13. Mishra (Pathak), S. K., Ranganathan, S.,Das, S. K., and Das, S. 1998. Investigation onprecipitation characteristics in a high strengthlow alloy (HSLA) steel. Scripta Materialia39(2): 253–259.

14. Moon, J., Kim, S., Lee, J., Hwang, B., Lee,Ch. G., and Lee, Ch. 2008. Effect of Cu and B ad-dition on tempering behavior in the weldCGHAZ of high strength low alloy plate steel.Materials Science and Engineering A 497 (1-2):153–159 (doi: 10.1016/j.msea.2008.06. 031).

15. Vilarinho, L. O., and Araujo, D. B. 2010.Low-cost methodology for obtaining CCT weld-ing diagrams. Soldagem & Insp. 15(3): 184–190(doi:10.1590/S0104-92242010000300004) (inPortuguese)

16. Zachrisson, J. 2006. In-Situ Detectionand Characterization of Phase Transformationsin Weld Metals. Master’s thesis. Master of Sci-ence Programme, Engineering Physics. LuleaUniversity of Technology, p. 44.

17. Alexandrov, B. T., and Lippold, J. C. 2005.In-Situ Weld Metal Continuous Cooling Transfor-mation Diagrams. IIW Doc. IX-2162-05.

18. Krauss, G., and Thompson, S. W. 1995.Ferritic carbon microstructures in steels. ISIJInternational. Vol. 35 (8), pp. 937–945. Iron andSteel Institute of Japan Bainite Committee,1992. Atlas for Bainitic Microstructures, Vol. 1,ISIJ, Tokyo, pp. 4, 5.

19. Kuziak, R., Botd, T., and Yi-Wen, Ch.1995. Microstructure control of ferrite-pearlite high strength low alloy steels utilizingmicroalloying additions. Journal of MaterialsProcessing Technology, Vol. 53, p. 255–262(http://dx.doi.org/10.1016/0924-0136(95)01983-L).

20. Ivanov, A. Yu., Sulyagin, R. V., Orlov, V.V., and Kruglova, A. A. 2011. Structure andproperties of weld joints in X80, X90, and K70pipe steel. Steel in Translation 41(7): 611–616

(doi: 10.3103/S0967091211070047).21. Shi, Y., and Han, Zh. 2008. Effect of

weld thermal cycle on microstructure and frac-ture toughness of simulated heat-affected zonefor a 800 MPa grade high strength low alloysteel. Journal of Materials Processing Technology207 (1-3): 30–39 (doi:10.1016/j.jmatpro-tec.2007.12.049).

22. Shome, M., and Mohanty, O. N. 2006.Continuous cooling transformation diagramsapplicable to the heat-affected zone of HSLA-80 and HSLA-100 steels. Metallurgical and Ma-terials Transactions A 37(7): 2159–2169 (doi:10.1007/BF02586136).

23. Zhang, Zh., and Farrar, R. A. 1995. Anstlas of Continuous Cooling Transformation(CCT) Diagrams Applicable to Low Alloy WeldMetals. The Institute of Materials, p. 95.

24. Zhao, M. C., Yang, K., Xiao, F. R., andShan, Y. Y. 2003. Continuous cooling transfor-mation of underformed and deformed low car-bon pipeline steels. Materials Science and Engi-neering A 355(1-2): 126–136 (doi:10.1016/S0921-5093(03)00074-1).

25. Stalheim, D. G., Barnes, R., Keith, E.,and McCutcheon, B. 2007. Alloy designs forhigh strength oil and gas transmission linepipesteel. International Symposium on MicroalloyedSteel for the Oil and Gas Industry. Eds. W. J.Fazackerley, P. Bordignon, K. Hulka, and F. Si-ciliano. TMS (The Minerals, Metals & Materi-als Society), p. 73–108.

26. Cizek, P., Wynne, B. P., Hodgson, P., andMuddle, B. C. 2005. Effect of simulated ther-momechanical processing on the transforma-tion characteristics and microstructure of anX80 pipeline steel. The 1st International Confer-ence on Super-High Strength Steels. Associ-azione Italiana di Metallurgia, Italy (ISBN8885298567).

27. Liu, Ch., Wang, D. Z., Liu, Y. X., Zhu,Q. H., and Zhau, Yu. 1997. Composition designof a new type low-alloy high-strength steel. Ma-terials & Design 18(2): 53–59 (doi:10.1016/S0261-3069(97)00029-0).

28. Radaj, D. 2003. Welding Residual Stressesand Distortion — Calculation and Measurement,English Edition, Vol. 2, Düsseldorf, DVS-Verlag.

29. Capdevila, C. 2012. Chapter 15: Neuralnetworks modeling of phase transformations insteels. Phase Transformations in Steels, Vol. 2:Diffusionless transformations, high strengthsteels, modelling and advanced analytical tech-niques. Edited by Elena Pereloma and David V.Edmonds, Woodhead Publishing Ltd., pp.464–500, ISBN 978-1-84569-971-0 (print)/ISBN 978-0-85709-611-1 (online).

30. Sourmaila, T., and Garcia-Mateo, C.2005. Critical assessment of models for predict-ing the Ms temperature of steels. Computa-tional Materials Science 34(4): 323–334,http://dx.doi.org/10.1016/j.com-matsci.2005.01.002.


WE

LD

ING

RE

SE

AR

CH


Documents

Effect of Tempering Pass on HSLA-80 Steel HAZ …...cation of HSLA steels where weight re-duction is required (through wall thick-ness reduction), while ensuring high weldability,