Weldability and Microstructural Aspects of Shielded Metal ...

1. Introduction

The HY (High yield) series of steels are traditionallyused for construction of naval ships and submarines.1)

These are essentially quenched and tempered low to medi-um carbon low alloy steels which derive their strengthmainly from carbon. But carbon being deleterious for weld-ing,2) over past two decades there have been attempts tosubstitute these HY series of steels by developing a new se-ries of low-carbon low-alloy copper-bearing steels with im-proved weldability maintaining similar strength and tough-ness.1,3) Lower carbon in these steels helps in achieving bet-ter weldability, whereas, copper enhances strength by pre-cipitation hardening.4)

ASTM A710 grade steel (0.07% C, 0.5% Mn, 0.4% Si,0.75% Cr, 0.9% Ni, 0.2% Mo, 1.15% Cu and 0.02% Nb)was first to be developed in these low-carbon copper-bear-ing steels in late seventies for use in offshore structures.5–10)

Based on its chemistry, HSLA-80 steel was developed byUnited States Navy in early eighties as a substitute for HY-80 grade of steel.1,11–14) High strength (Minimum yieldstrength of 552 MPa), good low temperature impact tough-ness (81 J at �85°C) and good weldability of HSLA-80steel made it a candidate for construction of naval ships andsubmarines. It can also be used for making components ofmany engineering bodies such as heavy duty tracks,bridges, earth moving equipment and off-shore struc-tures.15,16) After successful development of HSLA-80 steel,

United States Navy initiated a programme on developmentof HSLA-100 steel with improved strength (Minimum yieldstrength of 690 MPa) keeping good impact toughness (81 Jat �85°C) and weldability. This was necessary for makingnaval and engineering structures which encounter complexdynamic loading and require strength higher than that ofHSLA-80 steel. This steel has a nominal composition of0.06% C, 1% Mn, 0.02% P, 0.006% S, 0.4% Si, 1.6% Cu,3.5% Ni, 0.6% Cr, 0.6% Mo, 0.03% V, 0.04% Nb.1) Fewtrial heats were made in this regard in late eighties and earlynineties.1,17) Since then there have been efforts to obtain acritical understanding of the structure–property correlationsof this steel. Although, in the past several studies weremade in that direction,18–22) very few attempts were made tounderstand the weldability aspects of this steel.

Weldability is considered to be critical for any newly de-veloped steel for its successful commercial use. The basicreason for development of HSLA-100 steel to replace HY-100 is that of its improved weldability. Although, both thesegrade of steels have the same carbon equivalent (�0.81),the HSLA-100 steel is expected to have better weldabilityowing to its lower carbon content as indicated in Gravillediagram2) shown in Fig. 1.

Again for any welding process, the welding consumablesplay an important role. In the past Smith et al.23) as well asWallace et al.24) attempted to weld HSLA-100 steel with themilitary specified consumables meant for HY-100 steel.Smith et al. had worked with 25 mm thick plates and adopt-

ISIJ International, Vol. 42 (2002), No. 3, pp. 290–298

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Weldability and Microstructural Aspects of Shielded Metal ArcWelded HSLA-100 Steel Plates

Sanjay Kumar DHUA, Debasis MUKERJEE and Darbha Subrahmanya SARMA1)

Research and Development Centre for Iron and Steel, Steel Authority of India Ltd., Ranchi-834 002 India.E-mail: [email protected] 1) Department of Metallurgical Engineering, Banaras Hindu University, Varanasi-221005India. E-mail: [email protected]

(Received on August 16, 2001; accepted in final form on December 19, 2001 )

HSLA-100 steel with 14 mm thickness in quenched and tempered condition was shielded metal arc weld-ed (SMAW) with 2 kJ/mm heat input using basic flux coated filler rods without any pre or post welding heattreatments. The steel was found to be welded satisfactorily in this condition without developing any defect.Optical microscopy studies revealed typical cast dendritic structure in the weld metal and coarse bainite ingrain-coarsened area of the heat-affected zone (HAZ). Transmission electron microscopy (TEM) study con-firmed incidence of mixed structure of martensite laths and bainite in weld metal, while, it was mainly ofbainite laths in HAZ with evidence of martensite–austenite (M–A) constituent and massive ferrite. The yieldstrength (YS), ultimate tensile strength (UTS) and Charpy V-notch (CVN) impact energy of the weld metal(YS-695 MPa, UTS-842 MPa and CVN-105 J at �50°C) and HAZ (YS-790 MPa, UTS-891 MPa and CVN-130 Jat �50°C) were found satisfactory although HAZ properties were inferior to the base metal properties. Thehardening of HAZ was not very significant in this steel under the present welding condition.

KEY WORDS: steels; weldability; heat affected zone; transmission electron microscopy; martensite; impacttoughness.

ed the process of submerged arc welding (SAW) with a pre-heat temperature of 93°C and achieved the desired weld-ment properties. Wallace et al.24) as a part of the certifica-tion program, welded HSLA-100 steels of various thicknessranging from 9.5 to 51 mm adopting different techniqueslike gas metal arc welding (GMAW), shielded metal arcwelding (SMAW) and submerged arc welding (SAW) andthey used preheat temperatures ranging from 15 to 52°Cdepending on the plate thickness. They achieved desiredweld properties upto 19 mm thick plates with the existingconsumables, but failed to achieve properties for higherthickness plates with low preheat temperature (15°C). Morerecently, Park et al.25) worked on the effects of post weldingheat treatments on the toughness of the weld HAZ in a Cu-containing HSLA-100 steel. They used a thermomechanicalsimulator to simulate the weld HAZ and studied the effectof various thermal cycles on Cu precipitation and its effecton the HAZ properties.

In the current study, an attempt is made to weld 14 mmthick plates of HSLA-100 steel with an existing militaryspecified welding consumable by the process of shieldedmetal arc welding (SMAW) without any pre or post weldingtreatments. The objectives are to find out whether the steelis weldable in this condition without developing any cracksor fissures and also to study in detail the microstructure andmechanical properties of the steel in its weld and the heataffected zones.

2. Experimental Procedure

The steel used in the experiments was supplied by thePhoenix Steel Corporation, USA in the form of austenizedand quenched plates of 51 mm thickness from a 150 t trialproduction heat. The chemical analysis of the steel (basemetal) was conducted by optical emission spectroscopy, andthe composition is provided in Table 1. For welding experi-ments, plates of 200�100�14 (T) mm size were cut fromthe top/bottom portion of the 51 mm as received plate. Theplates were austenitised at 950°C for 40 min and quenchedin water followed by tempering at 650°C for 1 h. All the

aforesaid heat-treatments were carried out in an air-mufflefurnace. Subsequently single ‘V’ groove angles (30 degree)were cut in the heat-treated plates with 4 mm root faces fora total 60 degree included angle between two plates. A typi-cal weld groove design is shown in Fig. 2.

The shielded metal arc welding experiments were con-ducted using a M/S Advani Oerlikon make Triodyn-K320model DC-arc welding machine. The welding electrodesused for this study were of Atom Arc 12018-M2 grade(M/S ESAB make as per US military specification MIL-12018M2). The electrodes were basic flux coated and3.15 mm in diameter. The composition of the filler rod asprovided in the test certificate is given in Table 1.

The welding experiments were carried out manually in 5passes with welding voltage and current of 25 V and 130 Arespectively. The average welding speed was 97 mm/min.Based on the above parameters, the welding heat input wascalculated to be 2.01 kJ/mm as per Eq. (1).26)

............................(1)

Where: H�heat input (kJ/mm), E�arc voltage (V), I�welding current (A), S�arc travel speed (mm/min).

The composition of the weldmetal was analyzed in opti-cal emission spectrometer taking sparks in the weld deposit.The same is provided in Table 1. For metallographic char-

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Table 1. Chemical composition of welding filler rod, weldmetal and base metal.

Fig. 1. Graville’s diagram showing weldability of steel as a func-tion of carbon content and carbon equivalent.2)

Fig. 2. Schematic diagrams showing typical weld groove designfor shielded metal arc welding (SMAW) of HSLA-100steel plates: (a) and (b) location and orientation of tensileand Charpy specimens in HAZ and in weld metal respec-tively.

acterization of base metal, HAZ and weld metal, 10�20�14 (L�W�T) mm specimens were cut. The transversethrough-thickness (T–T) sections of these specimens werestudied. The specimens were mechanically polished andetched with a 2% nital solution for observation of the mi-crostructure using a Neophot-30 optical microscope. Thinfoils were prepared for transmission electron microscopy(TEM) in a twin jet electropolisher in a solution of 95%glacial acetic acid and 5% perchloric acid at 10°C and 25 V.For making thin foils special care was taken and polishedthin sections (T–T) were etched by 2% nital to delineateclearly the weld metal and HAZ zones before punching 3mm discs from the respective locations for final electropol-ishing. The foils were examined in a JEOL 4000 EX micro-scope at 200 KV accelerating voltage. Energy DispersiveSpectrometry (EDS) of few thin foils were conducted usinga Link’s EDS system.

For mechanical property evaluation, standard tensilesamples were taken along the weld direction from differentzones as indicated in Figs. 2(a) and 2(b). Shouldered tensilespecimens of 25 mm gauge length and 6.25 mm gauge di-ameter (as per ASTM A 370) were used in the present in-vestigation. Tensile testing was carried out with a Instron1195 universal testing machine at a cross head speed of1 mm/min. Standard Charpy samples of 10�10�55 mm di-mensions with 2 mm V-notch were prepared according toASTM E 23. Samples were cut in the transverse throughthickness (T–T) direction, keeping the position of the notchin the respective HAZ and weld metal locations as illustrat-ed in Figs. 2(a) and 2(b). The samples were light etched by2% nital solution to position the notch exactly in the HAZand weld metal locations. The tests were conducted with aTinius–Olsen model Charpy impact testing machine at vari-ous temperatures viz. �20°C and �50°C. The Charpy im-pact toughness of the welding electrodes for these steels isusually specified24) at �50°C, and that was the reason be-hind conducting the Charpy testing up to �50°C. For bothtensile and impact testing three specimens were tested foreach condition, and average values were reported. Thehardness was measured in a Vicker’s hardness tester using a30 kg load and an average hardness of 10 indented fields fora particular sample was reported. The error in hardnessmeasurement was �5 Vicker’s hardness number (VHN),while the error in case of yield strength (YS) and ultimatetensile strength (UTS) were �10 MPa. The error in mea-surement of elongation (EL) percent and reduction in area(RA) percent were �1%, whereas, it was �5 Charpy V-notch (CVN) energies in case of Charpy impact toughness.

Fractography studies of the tested impact specimens werecarried out in a JEOL make JSM-840A model scanningelectron microscope (SEM). X-ray diffractometric studiesof the base metal (water quenched and tempered at 650°C)were conducted to detect the presence of retained austenitewith a Siemen’s X-ray diffractometer, using a Mo Ka tar-get.

3. Results

3.1. Visual Examination

The visual examination of the welded joints, after ma-chining and surface grinding has indicated that no surface


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(a)

(b)

(c)

(d)

(e)

Fig. 3. Optical microstructures of the SMAW plate: (a) weld fu-sion line, left hand side of which shows weld metal,whereas, right hand side indicates HAZ, (b) a typical castdendritic structure of the weld metal, (c) weld metalstructure at higher magnification, (d) grain coarsenedarea of the HAZ depicting big prior austenite grains adja-cent to the fusion line and coarse bainite laths with fewmassive ferrite grains, and (e) base metal depicting tem-pered martensite.

or underbead cracks or fissures were present in the weld-ment. Thus, it can be observed that quenched and temperedHSLA-100 steel plates can be welded at 2 kJ/mm heat inputby the process of shielded metal arc welding without anyrequirement of pre or post welding heat treatments.

3.2. Optical Microscopy

Typical microstructure of the weldment near the fusionline of the SMAW plate is shown in Fig. 3(a) in which thefusion line is clearly seen. The microstructure of the weldmetal is shown at 100� and 500� magnifications in Figs.3(b) and 3(c) respectively. Both showed typical dendriticpattern found in a cast structure. It was however difficult toidentify phases present in the weld metal structure at thismagnification. The microstructural phases seems to consistof martensitie and/or bainitie.

The HAZ microstructure adjacent to the fusion line (Fig.3(d)) however appeared to be an aligned structure consist-ing of parallel bainitic ferrite laths separated by elongatedsecond phase particles. A few large angular ferrite grainscould also be observed. The prior austenite grain size ofHAZ near fusion line was of the order of 30–75 mm in di-ameter. The prior austenite grain size of the HAZ observedby Smith et al.23) in a submerged arc welded HSLA-100steel at 2 kJ/mm heat input was in the order of 25–60 mm.

The microstructure of the base metal about 15 mm awayfrom the fusion line was that of tempered martensite asshown in Fig. 3(e). The average prior austenite grain size ofthe base metal was measured and found to be 15 mm whichwas significantly less than that of grain coarsened HAZ.

3.3. Transmission Electron Microscopy

The transmission electron microscopy of the weld metalrevealed a mixed microstructure of martensite laths andbainitic ferrite. The martensite lath structure is depicted in

Figs. 4(a) and 4(b). Figures 4(c) and 4(d) illustrate massiveferrite with high dislocation density coexisting withmartensite–austenite (M–A) constituent which appeared asdark grain boundary phase typically found in a low carbonlow alloy steel weldment. Figure 4(e) exhibits a Ti basedoxide inclusion observed in the weld metal structure, theEDS analysis of which is shown in Fig. 5.

The TEM microstructures of HAZ are shown in Figs.6(a)–6(g). Figure 6(a) depicts coarse bainite laths, whereas,Fig. 6(b) shows a highly dislocated massive ferrite grainalong with M–A constituent (appeared as dark grain bound-ary phase) with a evidence of twinning in it. Typical M–A


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(a) (b) (c)

(d) (e)

Fig. 4. TEMs of the weld metal: (a) and (b) BF micrographs depicting martensite laths, (c) and (d) massive ferrite grainscoexisting with M–A (dark boundary phase) constituents, (e) BF showing big globular inclusions.

Fig. 5. EDS spectrum of the big globular inclusion observed inthe weld metal.

constituent found in HAZ structure is illustrated in Fig.6(c). More M–A constituents located at lath boundaries areexhibited in Fig. 6(d). The Cu and Nb(C, N) precipitatescould be observed at few places in the HAZ and a typicalBF micrograph depicting those precipitates is shown in Fig.6(e). A SAD and its schematic confirming the precipitatetypes are exhibited in Figs. 6(f) and 6(g) respectively.

The water quenched and 650°C tempered microstructureof the base metal has been shown in Fig. 7. The structureexhibits martensite laths.21) At lath boundaries the marten-site freshly formed from the intercritical austenite couldalso be seen as a dark phase (Fig. 7(a)). The diffractionanalyses of the lath boundary phase revealed that it was es-sentially a bcc phase, could be either ferrite or marten-site.21,22) Precipitates of Cu and Nb(C, N) were found withinthe laths as well as at the lath boundaries. The Cu precipi-tates were bigger and rod shaped (50 nm long), the evidenceof which are shown in BF and DF images (Figs. 7(b) and7(c)). Some of the fine and spherical particles observed inthe above BF and DF images were of Nb(C, N). TheNb(C, N) precipitates were inherited from the austenitestructure, and were not expected to grow on tempering. The

SAD pattern and its sketch taken from the same field areexhibited in Figs. 7(d) and 7(e), which confirmed presenceof both Cu and Nb(C, N) rings. The DF image shown inFig. 7(c) was essentially corresponding to (111) Cu ringand expected to have some area from the closely located(111) Nb(C, N) ring.

3.4. YS, UTS and VHN

The YS, UTS and VHN of the base metal, HAZ andweld metal of the MAW plates are shown in Table 2. TheYS of the base metal, HAZ and weld metal were 833, 790and 695 MPa respectively. The UTS of them were 938, 891and 842 MPa respectively. Both YS and UTS were maxi-mum in the base metal and lowest in the weld metal. Therewas a little deterioration of the HAZ properties comparedto base metal, the YS was reduced by 43 MPa and UTS wasreduced by 47 MPa. In case of VHN, however, peak value(320 VHN) was found in the HAZ and lowest value (277VHN) was found in the weld metal. The base metal hard-ness was 289 VHN.

3.5. Elongation Percent and Reduction in Area Percent

The %El of the base metal, HAZ and weld metal were re-


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(a) (b) (c)

(d) (e)

(f) (g)

Fig. 6. TEMs of the HAZ: (a) BF depicting coarse bainite laths, (b) a massive ferrite grain along with M–A constituent(dark boundary phase) with a evidence of twinning in it, (c) and (d) BF micrographs showing M–A constituents,(e) BF showing precipitates, (f ) SAD, and (g) schematic of SAD identifying Cu and Nb(C, N) precipitates.

spectively 23, 18 and 26% (Table 2). The maximum %Elwas achieved in the weld metal and the minimum was inthe HAZ. The %El of the HAZ was little less compared tothe weld metal and base metal owing to the coarse grainsize and presence of hard phase (M–A) in it. The % RA ofthe weld metal, HAZ and base metal varied within a closerange (70–73%). The %RA of the base metal and HAZwere found to be same of 73%, whereas, it was only slight-ly less at 70% in case of weld metal (Table 2).

3.6. Charpy Impact Toughness

The Charpy V-notch (CVN) impact energies of the basemetal, HAZ and weld metal at �20°C testing temperaturewas 190, 149 and 127 J (Table 2). The HAZ toughness wasless by 41 J from the base metal. The weld metal toughnesswas the lowest and 63 J less than that of the base metal. At�50°C testing temperature the base metal, HAZ and weldmetal toughness was found to be 179, 130, 105 J respective-ly.

Fractographic studies of the impact tested specimen werecarried out through SEM. Figures 8(a)–8(c) show typicalfractographs of the broken Charpy specimen tested at�20°C respectively from weld metal, HAZ, and base metal

locations. All three specimen exhibited dimples in the frac-ture surfaces confirming ductile mode of failures.

4. Discussion

4.1. Microstructure

The transmission electron microstructures of the weldmetal reveal low carbon lath martensite (Figs. 4(a) and4(b)), bainitic ferrite (Fig. 4(c)) along with massive ferritesubgrains (Fig. 4(d)). The higher Mn and Ni content in theweld metal composition (Table 1) and faster cooling rateowing to lower welding heat input (2.01 kJ/mm) has result-ed in the formation of bainitic or massive ferrite with a highdislocation density along with martensite. Similar observa-tions were made by Deb et al.27) in their study with GMAWweldments in HY-100 steel welded with Linde 120 (MIL-120S-1) electrode having similar composition like that usedin the present study. They observed typical microstructureof low carbon lath martensite and bainite phases in theirGMAW weld deposit with welding heat input 2.2 kJ/mm.

Large and globular (600–800 nm) non-metallic inclu-sions could be observed in the weld metal (Fig. 4(e)). TheEDS analysis (Fig. 5) confirmed this to be of Ti based in-clusion and most likely it was a complex combination ofTi2O3 and (Mn, Si) O. The origin of Ti could be the basicflux cover of the electrodes used in the welding. The chemi-cal analysis of the flux used in the submerged arc weldingof HSLA-100 and HSLA-80 steels reported by Dixon etal.28) and Sen et al.29) respectively showed the occurrence ofTi in the basic flux. Sen et al.29) also reported the existenceof big globular Ti based inclusion in the weldmetal of asubmerged arc welded HSLA-80 steel plate. According tothem, as well as other previous workers,30,31) these inclu-


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Table 2. Mechanical properties of the shielded metal arcwelded HSLA-100 steel weldment.

(a) (b) (c)

(d) (e)

Fig. 7. TEMs of the base metal: (a) partially recovered matrix with freshly formed martensite appearing as dark phase atlath boundaries, (b) BF and (c) DF showing elongated and rod shaped Cu precipitates with a few fine and spheri-cal Nb(C, N) precipitates, (d) SAD, and (e) schematic of SAD, identifying Cu and Nb(C, N) precipitates.

sions assist in the nucleation of acicular ferrite in the weldstructure which is desirable from toughness point of view.According to Losz et al.,30) these titanium oxides are ther-mally stable and do not lose their ability to promote acicu-lar ferrite formation even at locations in the welding fusionline. The acicular ferrite is considered to be very effectivein enhancing weld toughness.

The HAZ microstructure consists of bainite laths asshown in Fig. 6(a), separated by interlath M–A constituent.M–A constituent is seen in all the BF micrographs depictedin Figs. 6(a)–6(e). Figure 6(b) shows clear evidence oftwinning within a M–A island and large ferrite subgrains inthe HAZ. The air cooling of the HAZ from high austenitis-ing temperature and relatively short cooling time was ex-pected to produce coarse bainite structure. Krishnadev etal.32) expressed similar opinion while studying the simulat-ed HAZ properties of copper-precipitation strengthenedHSLA steels with HSLA-80 and HSLA-100 type composi-tions. Losz et al.,30) while studying the HAZ microstruc-tures in HSLA steel weldments observed that the micro-

structure of low carbon steels (�0.10% C) were largelybainitic. According to them, a combination of large austen-ite grain size and a limited number of nuclei sites give riseto coarse transformation structures. In the present case too,bainite was of lath morphology instead of granular typenormally found in an air-cooled HSLA-100 steel. Accordingto Losz et al.30) the bainite grows as packets of dislocatedparallel laths separated by low angle boundaries and veryoften the boundaries are decorated with coarse carbon richsecond phase particles. This corroborates the present obser-vation of M–A constituent at lath boundaries. The HAZmicrostructure observed by Deb et al.27) in a GMAW HY-100 weldment was of mixed type, consisting the lathmartensite and upper bainite. This could be due to the high-er carbon content of HY-100 steel compared to HSLA-100.

The HAZ microstructure also showed precipitates. Al-though, most of the Cu and Nb(C, N) precipitates were like-ly to go into the solution when the peak HAZ temperaturewas reached in a weld thermal cycle, some of these parti-cles were expected to reprecipitate during its cooling cycleowing to high Cu (�1.77 wt%) content of this steel.Moreover, the effect of tempering due to subsequent weld-ing passes in a multipass welding must have also partiallyassisted in reprecipitation of Cu and Nb(C, N) particles inthe HAZ.

The TEM microstructure of the base metal revealed tem-pered martensite laths with elongated and rod shaped Cuprecipitates. A dark phase, which was established to bemartensite or acicular ferrite by electron diffraction ap-peared along the lath boundaries. Mujahid et al.20) has alsoobserved this lath boundary dark phase. According to themthe new austenite formed on tempering above 665°C trans-formed to martensite on subsequent air cooling to roomtemperature, leaving behind substantial amounts of retainedaustenite which appeared as a dark phase at lath bound-aries. Our investigation, however, revealed that this lathboundary dark phase (Fig. 7(a)) was martensite, whichcould be freshly formed from intercritical austenite on cool-ing.21,22) This is also supported by the fact that the presentstudy did not reveal any austenite peak in X-ray diffraction.

4.2. Relation between the Microstructure and theStrength

The YS (695 MPa) of the weld metal was the lowestcompared to the HAZ and base metal owing to the presenceof untempered martensite and bainitic ferrite and also dueto absence of Cu and Nb(C, N) precipitates. The UTS (842MPa) and VHN (277) are however reasonably good becauseof the higher dislocation density in ferrite and martensite ofthe as cast structure. The YS, UTS and VHN of the weldmetal are very close to the values (YS- 725 MPa, UTS-889MPa, VHN-282) obtained by Smith et al.23) in their sub-merged arc weld deposit in a HSLA-100 steel weldment.The values obtained in the present investigation are stillabove the minimum specified values for HSLA-100 steel,viz. YS-690 MPa.

The presence of large amount of ferrite and M–A con-stituent in the HAZ must have resulted in inferior YS com-pared to base metal. However, YS of the HAZ is still betterthan weld metal because of the benefit of precipitates in theformer. Highly dislocated bainite laths and ferrite in the


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(a)

(b)

(c)

Fig. 8. SEM fractographs of the Charpy specimen tested at�20°C revealing dimples: (a) weld metal, (b) HAZ, and(c) base metal.

HAZ enabled it to achieve high UTS (891 MPa). The hard-ness of the HAZ was expected to be little higher than basemetal and weld metal owing to the incidence of largeamount of M–A constituent in it. According to Shiga etal.,33) the adverse effect of M–A constituent in HAZ on itstoughness is enhanced owing to the segregation of C, Nduring g to a transformation. The segregation of these in-terstitials is brought about by thermal cycles of subsequentwelding passes in a multipass welding. The hardening ofthe HAZ is a known phenomenon and reported by manyearlier workers in HSLA and HY steels.16,23,27,28,34) As com-pared to HY-100 steel, the hardening of HAZ was howevernot very significant in the present study. Deb et al.27) report-ed average HAZ hardening to the tune of 165 VHN inGMAW HY-100 steel, welded with 2.2 kJ/mm heat input,whereas, in the present study it was found to be only 30VHN for HSLA-100 steel welded with 2.01 kJ/mm heatinput. The higher carbon content of HY-100 steel could beresponsible for its higher HAZ hardness. Czyryca34) also re-ported insignificant HAZ hardening (25–30 VHN) inHSLA-100 steel weldments. Philips et al.16) in their workwith GMAW HSLA-80 steel found peak hardening of HAZat 1 kJ/mm heat input and HAZ hardness was 300 VHN,whereas it was 420–450 VHN for HY-80 steel under similarcondition.

4.3. Influence of Microstructure on the Impact Tough-ness

The CVN energies of the weld metal, HAZ and basemetal at �20°C were reasonably good, although the weldmetal had the lowest toughness as compared to HAZ andbase metal. The cast columnar structure of the weld metalcan be responsible for its comparatively lower CVN energy.The HAZ CVN is less with respect to base metal owing toits coarse grain size and presence of large grain boundaryM–A constituents.

The CVN energy of the weld metal at �50°C (105 J)meets the minimum stipulated requirement of 61 J at�51°C as per US military specification, MIL-E-22200/10Ameant for MIL-12018-M2 type electrodes,24) which wasused in the present study.

The CVN energies of the weld metal at �20°C and�50°C obtained in the present study are also superior tothe values (93 J at �20°C and 61 J at �50°C) reported bySmith et al.23) of a SA weld deposit of HSLA-100 steel.They were also found to be superior to the values (90 J at�20°C and 70 J at �50°C) reported by Deb et al.27) for theweld deposit of a GMAW weldment of HY-100 steel.

The fractographic studies of the fractured specimen test-ed at �20°C revealed coarser dimples in the HAZ com-pared to weld metal and base metal locations, which was inaccordance with their CVN energies found at that tempera-ture.

The present welding experiments are restricted to 2.01kJ/mm following the US military requirement of maintain-ing low energy inputs in welding of HY-100 grade steel toobtain optimum HAZ properties.23) Although welding athigher heat input is preferred in industry for saving weldingtime, it however results in considerable variation of weld-metal as well as HAZ microstructures when heat input israised. Owing to slower cooling rate at larger heat input,

both weld metal and HAZ microstructure changes andcoarsening of prior austenite grain size and precipitation ofundesirable hard phases takes place leading to the deterio-ration of strength and toughness.16,24,28,32,33,35,36) The effectis however more pronounced in case of HY-series (HY-80/HY-100) of steels compared to Cu-strengthened HSLA-80/100 steels owing to the higher carbon content of the for-mer.23)

5. Conclusions

(1) The HSLA-100 steel plates of 14 mm thickness canbe welded by shielded metal arc welding at 2 kJ/mm heatinput without any pre or post welding heat treatments.

(2) The weldmetal microstructure was comprised oflath martensite and bainitic ferrite. The HAZ structure wascoarse bainitic laths with interlath M–A constituents andmassive ferrite.

(3) Both weld metal and HAZ properties were abovethe stipulated requirements of HSLA-100 steel thoughproperties were slightly less than the base metal.

(4) The CVN of the weld metal at �20 and �50°Ctesting temperatures were 127 and 105 J which are abovethe minimum stipulated values for the specific kind of elec-trode used in this study. The HAZ toughness was good andfound to be 149 and 130 J at �20 and �50°C respectively.

Acknowledgements

The authors acknowledge the financial assistance provid-ed through an INDO-US project coordinated by NavalResearch Laboratory, USA and National MetallurgicalLaboratory, India. Authors are also grateful to the manage-ment of R&D Centre, Steel Authority of India Limited,Ranchi for providing permission and the necessary experi-mental facilities for this work.

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