ORIGINAL ARTICLE Open Access

Effect of filler addition on solidificationbehaviour and hot tensile properties ofGTA-welded tube joints of Super 304Haustenitic stainless steelM. Vinoth Kumar1, V. Balasubramanian1* and A. Gourav Rao2

Abstract

Background: Super 304H austenitic stainless steel tubes containing 2.3 to 3 (wt.%) of Cu are mainly used insuperheaters and reheaters of ultra super critical boilers due their excellent corrosion and oxidation resistance. Cuaddition causes precipitation strengthening effect by fine Cu-rich precipitates which evolve during creep conditionsand results in increased creep strength. The microstructural evolution of stainless steels during welding significantlyaffects the material properties.

Methods: In this work, solidification mode microstructure and hot tensile properties of autogenous and filler-addedSuper 304H gas tungsten arc (GTA) welded tube joints were correlated.

Results: Autogenous welds of Super 304H solidified as austenite with 1.57 % delta ferrite and intercellular(Fe,Cr)23(C,B)6 borocarbides. In filler-added welds, the higher Ni equivalent and addition of carbide stabilizingelements (Nb,Mo) eliminated δ ferrite and segregation of B as borocarbides.

Conclusions: The filler-added welds exhibited superior tensile strength than autogenous welds at both roomand high temperature.

Keywords: Super 304H, Gas tungsten arc welding, Solidification behaviour, Segregation, Hot tensile properties

BackgroundEfficiency of the power plant strongly depends on theoperating steam temperature and pressure. The concernfor reducing the CO2 emission and coal consumptionleads to the efforts to build power plants operating athigher steam parameters (Brozda 2007; Shu Ping et al.2010; David et al. 2013). Austenitic stainless steels areselected for use in sections of superheaters and rehea-ters, which require good oxidation and fire side corro-sion resistance, in addition to the creep strength. Therecent improvements in operating steam temperaturesup to 620 °C were achieved through the development ofcreep-resistant stainless steels, containing Cr, Ni, W, Co,

Cu and N. Future plans of increasing the efficiency ofthe power plants up to 49 % require creep-resistantsteels operating at steam temperatures of 760 °C andpressures of 35 MPa (Viswanathan et al. 2005). Super304H austenitic stainless steel is used in recently installedultra super critical (USC) power units of Japan operatingat steam temperature of about 613 °C (Thomas 2013).Further, Super 304H is listed as a candidate material forsuperheaters and reheaters in the US program for devel-opment of materials for USC with steam temperatures of760 °C (Viswanathan et al. 2005). Super 304H with nom-inal composition of 0.1C-18Cr-9Ni-3Cu-Nb-N derivesexcellent creep strength from the distinct Cu additionwhich precipitates as fine Cu-rich particle under creepconditions, in addition to NbCrN, Nb(C,N) and M23C6

particles (Li et al. 2010).The different modes of solidification mechanism

(A-AF-FA-F) and the methods to predict the volume

* Correspondence: visvabalu@yahoo.com1Center for Materials Joining and Research (CEMAJOR), Department ofManufacturing Engineering, Annamalai University, Annamalai Nagar, TamilNadu 608002, IndiaFull list of author information is available at the end of the article

© 2015 Vinoth Kumar et al. Open Access This article is distributed under the terms of the Creative Commons Attribution 4.0International License (http://creativecommons.org/licenses/by/4.0/), which permits unrestricted use, distribution, andreproduction in any medium, provided you give appropriate credit to the original author(s) and the source, provide a link tothe Creative Commons license, and indicate if changes were made.

Vinoth Kumar et al. International Journal of Mechanicaland Materials Engineering (2015) 10:25 DOI 10.1186/s40712-015-0051-x

percentage of ferrite along with the solidification modesare described in detail elsewhere (Valiente Bermejo2012). A ternary alloy with a composition of Super 304His expected to solidify as austenite under equilibriumconditions (Hunter and Ferry 2002). In non-equilibriumcooling conditions such as welding, the solidificationsequence of the stainless steel depends on the coolingrate and, in highly alloyed steels such as Super 304H,it may result in local compositional variation due tosegregation, which alters the solidified phases at roomtemperature (Darja Steiner et al. 2011).Recently, few investigations (Indrani et al. 2011; Yang

et al. 2006; Li et al. 2010; Bai et al. 2013; Ha and Jung2012) were carried out to understand the ageing behav-iour of Super 304H alloy at high temperatures and eval-uated the effect of varying Cu content on microstructureand mechanical properties of the alloy. Works carriedout on weldability studies and high temperature proper-ties of Super 304H weld joints reported the tendency ofSuper 304H to hot cracking and use of matching Super304H filler metal as an alternate to Ni-based filler metal(Vekeman et al. 2014; Kim et al. 2014). However, theavailable information on solidification behaviour of gastungsten arc (GTA) welds of Super 304H and its effectson high temperature properties are very scant. Hence inthis work, an attempt has been made to compare the so-lidification modes, microstructure and hot tensile prop-erties of autogenous and filler-added GTA-welded tubejoints of Super 304H austenitic stainless steel. Theamount of alloy segregation, composition of segregatedelements in autogenous joints, and the role of alloyingelements added in the filler metal to avoid detrimentalsegregation along with their effects on tensile strength ofthe joints are discussed in this paper.

MethodsSuper 304H austenitic stainless steel tubes of outerdiameter 57.1 mm and wall thickness of 3.5 mm wereused in this investigation. Chemical composition of the

Super 304H tube and filler metal used in this investiga-tion are presented in Table 1. The tensile properties oftube in solution-annealed condition are given in Table 2.For autogenous welds, the joints with square butt prep-aration with no gap was employed, and for filler-addedwelds, the joints with single ‘V’ butt preparation waswelded using GTAW process. Argon was used as theshielding and purging gas for both cases, with flow ratesof 12 and 10 L/min, respectively. The welding parame-ters used in this investigation are shown in Table 3.Photographs of the autogenous and filler-added GTA-

weld joints are shown in Fig. 1a, b, respectively. To iden-tify the mode of weld solidification, microstructural ex-aminations were carried out by obtaining samples fromthe weld joints. The samples were prepared using stand-ard metallographic techniques and etched with Glycere-gia for 5–10 s to reveal the general structure and withboiling Murakami’s reagent to reveal δ ferrite and car-bides. The microstructural examination of the sampleswas carried out using light optical microscope (OM)with image analysing software, field emission scanningelectron microscope (FE-SEM). The compositional vari-ation within the weld regions was analysed using energy-dispersive spectroscopy (EDS) attached with SEM. The δferrite measurement in the welds was carried out usingferritescope. The specimens for transverse tensile testwere extracted from the weld joints using wire-cut elec-tric discharge machining. Figure 1c represents the di-mension of the transverse tensile specimens extractedfrom the weld joints. Instron universal testing machine(UTM) was used to carry out the tensile testing, underconstant crosshead speed mode with a nominal strainrate of 1 × 10−3 S−1. The tensile tests were carried out atfour different test temperatures (room temperature (RT),550 °C, 600 °C and 650 °C). The UTM system wasequipped with a three-zone resistance heating furnacefor high temperature tests and a computer with data ac-quisition system for obtaining digital load-elongationdata. The photograph of hot tensile test specimens after

Table 1 Chemical composition (wt.%) of parent metal (PM) and filler metal (FM)

C Si Mn P S Cr Ni N Cu Nb Mo B Al

PM 0.086 0.23 0.81 0.021 0.0003 18.18 9.06 0.095 3.080 0.045 - 0.0039 0.01

FM 0.1 0.3 3.3 <0.01 <0.01 18.3 15.7 0.16 2.9 0.5 0.7 - -

Table 2 Tensile properties of parent metal in as-receivedcondition

0.2 % yieldstrength (MPa)

Ultimate tensilestrength (MPa)

Elongation in25 mm gaugelength (%)

Annealingtemperature

Parentmaterial

284.2 575.8 71.8 1145 °C

Table 3 Parameters for GTAW welding of Super 304H

Mode of welding Autogenous welding With filler addition

Power source type Constant current Constant current

Current (A) 100 75

Voltage (V) 11 11.3

Welding speed (mm/min) 70 75

Heat input (kJ/mm) 0.943 0.68

Vinoth Kumar et al. International Journal of Mechanical and Materials Engineering (2015) 10:25 Page 2 of 11

the test is shown in Fig. 1d. The hardness was measuredacross the weld centre line using Vickers microhardnesstester with a load of 500 g and dwell time of 15 s. Thefracture surfaces of the tensile specimens were analysedusing SEM to reveal the mode of fracture.

ResultsMicrostructureFigure 2a shows the macrograph of the autogenousGTAW joint and it reveals the various regions of thejoint, consisting of parent metal (PM), heat-affected zone(HAZ) and fusion zone (FZ). Figure 2b shows the micro-structural variation across the joint at low magnification.The micrograph of weld metal (WM) is shown in Fig. 2c,which reveals the presence of austenite (white phase)with δ ferrite (dark phase). Micrograph of the WM atcentre of the joint, shown in Fig. 2d, reveals the presenceof precipitates both within the austenite grains and inthe austenitic grain boundaries. The presence of precipi-tates within the δ ferrite at the grain boundaries was alsorevealed in Fig. 2d. Micrograph of weld fusion lineshown in Fig. 2e reveals the presence of δ ferrite withmixed morphology and preferential orientation towardsthe weld centre.Macrograph of the filler-added GTAW joint is shown

in Fig. 3a and the microstructural variation across thejoint is shown in Fig. 3b. The WM micrograph of thefiller-added weld joint, shown in Fig. 3c, reveals thepresence of fully austenitic equiaxed dendritic structure.

The grain size of the filler-added WM is coarser (80–100 μm) than the PM (30–40 μm). Micrograph of filler-added WM is shown at higher magnification in Fig. 3dand reveals no δ ferrite. However, it consists of finerinterdendritic and grain boundary precipitates than theautogenous weld. The fusion line of filler-added joint isshown in Fig. 3e, which reveals the partially melted zonewith epitaxial growth of austenitic grains towards theweld centre with no evidence of δ ferrite.SEM micrograph of HAZ of the autogenous weld joint

is shown in Fig. 4a, and it reveals coarsening of precipi-tates and austenite grains. Figure 4b shows the SEMmicrograph of autogenous WM, which reveals the pres-ence of finer precipitates within austenitic grains andcoarse precipitates within the δ ferrite. Chemical compos-ition of the WM region determined by EDS is shown inFig. 4c, where the composition of alloying elements inWM matches with the PM. In order to determine the typeof precipitates available within δ ferrite, the sample wasetched with Murakami’s reagent near boiling point.Murakami’s reagent preferentially etches the carbides atthe austenite grain boundaries and δ ferrite. Micrographof WM etched with Murakami’s reagent is shown in Fig. 4dand it shows the intercellular segregation of alloyingelements (grey phase). EDS line scan across the grey phase(L1) for elemental variation in Cr and Ni is presented inFig. 4e. The marginal increase in Cr concentration andreduction in Ni concentration across the grey phasesuggests the presence of δ ferrite. The composition of the

Fig. 1 Photographs of welded tubes and details of tensile specimens. a Autogenously welded tube joint. b Welded tube joint with filler addition.c Dimensions of hot tensile specimen in ‘mm’. d Hot tensile specimens after test

Vinoth Kumar et al. International Journal of Mechanical and Materials Engineering (2015) 10:25 Page 3 of 11

δ ferrite region was measured by EDS analysis on the spot(S1) and found to be rich in Fe, Cr, B and C.SEM micrographs of the filler-added weld joint are

shown in Fig. 5 and the WM micrograph shown in Fig. 5areveals randomly oriented fully austenitic dendritic grainswith no δ ferrite. The elemental spectrum of the WM re-gion marked in Fig. 5a was determined by EDS and pre-sented in Fig. 5b. The WM composition is similar to thecomposition of the filler metal added to the weld. The fu-sion line and HAZ of the filler-added weld joint, shown in

Fig. 5c, reveals the columnar austenitic grains in the weldclose to the fusion line with no δ ferrite and few coarse pre-cipitates within the coarsened austenitic grains of the HAZ.The WM at higher magnification, shown in Fig. 5d, revealsthe presence of interdendritic and grain boundary precipi-tates. The EDS elemental mapping of Nb, Mo and C forthe WM region marked in Fig. 5d is shown in Fig. 5e, f, g,respectively. The precipitates are confirmed to be the richin Nb, Mo and C; segregation of other alloying elements isnot observed from the EDS analysis.

Fig. 2 Optical micrograph of autogenous GTA-welded joint. a Macrograph. b Micrograph across the weld. c Weld metal. d Weld metal at highmagnification. e Fusion line

Vinoth Kumar et al. International Journal of Mechanical and Materials Engineering (2015) 10:25 Page 4 of 11

Ferrite measurementThe amount of δ ferrite in the WM of autogenous andfiller-added weld joints was measured using ferritescopeand presented in Table 4. The average ferrite content ofthe autogenous WM is found to be 1.57 % and no δ fer-rite is found in the WM of filler-added joint. Theamount of δ ferrite determined using the ferritescope forautogenous WM is much lower than volume of darkphase (termed as δ ferrite) in the microstructure. Theamount of dark phase in autogenous weld is found to be8.23 % by image analysis.

Tensile propertiesThe engineering stress-strain curves of autogenousand filler-added weld joints at various test tempera-tures are shown in Fig. 6a, b, respectively, and thetensile properties are presented in Table 5. The RTtensile strength of autogenous weld joint is lesserthan the PM strength (−2 %), with failure located inthe weld centre. In case of filler-added joint, the ten-sile strength exceeded the PM strength (+6 %), withfailure in the parent material. The tensile strength ofboth autogenous and filler-added weld joints

Fig. 3 Optical micrograph of filler-added GTA-welded joint. a Macrograph. b Micrograph across the joint. c Weld metal. d Weld metal at highmagnification. e Fusion line

Vinoth Kumar et al. International Journal of Mechanical and Materials Engineering (2015) 10:25 Page 5 of 11

decreased with increase in test temperature. Thetensile strength decreased by 31 % in autogenousweld joints and by 36 % in filler-added weld jointwith increase in test temperature from RT to 650 °C.The elongation of both the autogenous and fillerweld joints decreased with increase in testtemperature.

HardnessThe mid-thickness microhardness measured across the au-togenous and filler-added weld joint are shown in Fig. 7.The WM of filler-added joint recorded higher hardnessthan the WM of autogenous joint; steep hardness varia-tions were observed within the WM region of both au-togenous and filler-added joints. The lowest hardness value

Fig. 4 SEM micrograph of Autogenous weld joint. a Micrograph (Glyceregia). b Weld metal (Glyceregia). c EDS composition of weld metal. d Weldmetal (Murakami’s). e Line scan across the grey phase (L1). f Spot EDS of grey phase

Vinoth Kumar et al. International Journal of Mechanical and Materials Engineering (2015) 10:25 Page 6 of 11

was recorded in the WM region close to the FL for bothautogenous and filler-added joints. The HAZ region ofboth the autogenous and filler-added joints has not under-gone softening due to the weld thermal cycles and retainedits hardness equivalent to that of the PM (175 HV).

Fracture surfacesThe fracture surface of autogenous joint tensile tested atRT and 550 °C is shown in Fig. 8a, b, respectively. Thereduction in cross-sectional area of the specimen testedat RT is higher (refer to Fig. 8a) than that of the speci-men tested at 550 °C. The high magnification SEM frac-tograph of autogenous joints tested at RT and 550 °C isshown in Fig. 8c, d and invariably consists of dimples.Since the fracture of the autogenous joints was locatedin the WM, the dimple sizes closely match with that ofthe dendrite size. The precipitates segregated along thegrain boundaries (marked by arrow) assisted in crackinitiation.

Fig. 5 SEM micrograph of filler-added weld joint. a Weld metal (Glyceregia). b EDS analysis by weld metal. c Fusion line (Glyceregia). d Weldmetal at high magnification. e Nb elemental map. f Mo elemental map. g C elemental map

Table 4 Percentage of delta ferrite in welds of Super 304H joint

Joint δ ferrite (%) Average δferrite (%)

Standarddeviation (%)

Autogenous 1.5, 1.7, 1.6, 1.7, 1.5,1.7, 1.4, 1.5, 1.6, 1.5

1.57 0.106

Filler added No traces Nil Nil

Vinoth Kumar et al. International Journal of Mechanical and Materials Engineering (2015) 10:25 Page 7 of 11

The SEM fractographs of filler-added joint tested atRT and 550 °C is shown in Fig. 8e, f, respectively, whichreveals a similar behaviour with much lesser reductionin the cross-sectional area of the specimen tested at550 °C than at RT. The fracture in the filler-added joints

was located at the PM, which is evident from the ab-sence of dendritic structure in the high magnificationfractographs shown in Fig. 8g, h. The specimen tested at550 °C reveals coarse and fewer dimples than the speci-men tested at RT (refer to Fig. 8g). Flat featureless re-gions which are much prominent in specimen tested at550 °C evidence the reduction in elongation of the jointat high temperatures. The Nb-rich precipitates (markedby arrow) associated with the voids in the fracture surfaceof RT specimen act as the crack initiation site.

DiscussionThe mode of solidification in stainless steels is stronglydependent on the chemical composition which is usuallyrepresented as Cr and Ni equivalents (Valiente Bermejo2012; Brooks et al. 1983). The WRC-1992 constitutiondiagram Cr/Ni equivalents are commonly used to pre-dict the solidification mode and ferrite content of theweld based on composition of the alloy (Srinivasan et al.2012). The Cr and Ni equivalent of autogenous weldmetal is determined as 18.2 and 14.7 %, respectively.Similarly, the Cr and Ni equivalent for filler-added weldmetal is determined as 19.4 and 23.1 %. The welds of

Fig. 6 Engineering stress-strain curves of weld joints at various testtemperatures. a Autogenous weld joint. b Filler-added weld joint

Table 5 Tensile properties of autogenous and filler-added weld joints of Super 304H

Test temperature(°C)

0.2 % yieldstrength (MPa)

Ultimate tensilestrength (MPa)

Elongation in 25 mmgauge length (%)

Failurelocation

Autogenous joint RT 302.4 ± 5.5 564.3 ± 7.0 38.96 ± 2.0 WM

550 201.6 ± 7.4 420.4 ± 10.0 32.09 ± 2.6 WM

600 173.0 ± 7.4 388.2 ± 7.6 24.97 ± 3.0 WM

650 206.2 ± 5.3 389.0 ± 8.5 34.70 ± 2.5 WM

Filler-added joint RT 349.6 ± 10.2 614.6 ± 11.7 52.3 ± 3.0 PM

550 247.3 ± 10.6 470.2 ± 10.7 39.2 ± 3.0 PM

600 196.3 ± 6.7 413.6 ± 12.5 36.6 ± 2.3 PM

650 240.8 ± 7.6 392.7 ± 13.7 27.3 ± 2.6 PM

Fig. 7 Hardness profile across the weld centre line

Vinoth Kumar et al. International Journal of Mechanical and Materials Engineering (2015) 10:25 Page 8 of 11

both autogenous and filler-added joints are predicted tosolidify as primary austenite (A-mode) with no δ ferritein the WM (as per WRC-1992 diagram). The WM of au-togenous weld consists of 1.57 % δ ferrite (refer toTable 4), which is in contrary to the predicted weldmetal microstructure.

The autogenous weld solidifies by epitaxial growth anda plane front solidification zone near the fusion line(refer to Fig. 2d) and transforms to cellular mode withmicrosegregation to the cellular boundaries (refer toFig. 2b) (Elmer et al. 1989). The vertical section of theFe-Ni-Cr ternary diagram is shown in Fig. 9 with lines

Fig. 8 Fracture surface of autogenous (a–d) and filler-added (e–h) weld joints tested at RT and 550 °C. a, e SEM macrograph, RT. b, f SEM macrograph,550 °C. c, g SEM micrograph, RT. d, h SEM micrograph, 550 °C

Vinoth Kumar et al. International Journal of Mechanical and Materials Engineering (2015) 10:25 Page 9 of 11

representing the compositional range of the Super 304H.The first solidifying structure in the weld depends onthe side in which the eutectic liquidus nominal compos-ition lies and the remaining liquid may solidify as eutec-tic or peritectic based on the level of segregation(Brooks and Thompson 1991).The autogenous weld of Super 304H first solidifies as

δ ferrite near the fusion line due to the peritectic solidi-fication, followed by eutectic single austenite solidifica-tion phase evident from the cellular structure ofaustenite in the weld centre. The austenite continues togrow into the peritectic δ ferrite by solid state trans-formation, with austenite as the more stable phase.Metalographically hard to distinguish peritectic and eu-tectic ferrite was retained in the intercellular austenitic/austenitic and austenitic/δ-ferrite boundaries of the au-togenous weld (refer to Fig. 2b), as the solid state trans-formation of δ ferrite to austenite is not complete due tosegregation or cooling conditions (Brooks et al. 1983).The addition of B to stainless steel enhances the har-denability and creep strength, however the B in solid so-lution is beneficial and precipitated borides aredetrimental. The segregation of Fe, Cr, B and C (referto Fig. 4f ) in the intercellular region of the austenitereveals the presence of (Fe,Cr)23(C,B)6 borocarbides(Hoffman and ASM 1989; Karlsson et al. 1982). Thelimited solubility of B in austenitic stainless steel of 30PPM at 900 °C may cause the excess B to segregatealong the grain boundaries and combine with Cr andFe to form a low melting eutectic with the austenite(Carinci 1994). Such eutectic formation and borocar-bide precipitation within the ferrite stringers alter the

composition of the solidifying liquid ahead of the aus-tenite/liquid interface and hinder the liquid to solidifyas austenite in the intercellular boundaries.In case of filler-added weld, the WM is fully austen-

itic with no δ ferrite as predicted (refer to Fig. 5a)and it is attributed to the increased addition of Ni,Nb, Mo and N to the weld metal. The higher Niequivalent of filler-added WM resulted in A mode ofsolidification with primary austenite dendrites as sin-gle phase (refer to Fig. 5d) (Bonollo et al. 2004). Theprecipitation of (Fe,Cr)23(C,B)6 was to be suppressedby the Mo, Nb and N addition to the WM of filler-added joint. The formation of (Nb,Mo)C clusters or(Nb,Mo) carbonitrides (refer to Fig. 5e, f, g) preventsthe supplement of C to the grain boundaries and re-tards the formation of detrimental carboborides infiller-added joint (Hara et al. 2004).The failure of the autogenous weld joint was located at

the weld centre regardless of the test temperature, re-corded the lowest hardness in the weld joint (refer toFig. 7). In single pass autogenous weld, the weld metalsolidifies as coarse columnar grains by epitaxial growthtowards the weld centre with preferential orientation,which may result in a weak centre line in the weld(Villafuerte and Kerr 1990). The higher tensile strengthof filler-added joint than the autogenous joint at all testtemperatures is attributed to the fine Nb- and Mo-richcarbides. In turn, B retained in the austenitic matrixwithout precipitation of detrimental carboborides in theWM of filler-added joint. The retention of elemental Bin the WM increases the hardenability and therebyresulted in increased strength of the filler-added joint.The reduction in the strength values of weld joints withincrease in test temperature is attributed to the acceler-ated recovery process (Choudhary and Rao Palaparti2012).

Conclusions1. Autogenously welded GTAW joints of Super 304H

austenitic stainless steel resulted in weld metal withaustenite and δ ferrite (1.57 %). Boron tends tosegregate along the intercellular boundaries to form(Fe,Cr)23(C,B)6 borocarbides.

2. In the filler-added GTAW joints, the increased Mo,Nb and N content suppressed the precipitation ofcoarse borocarbides by formation of (Nb,Mo)Cclusters or (Nb,Mo) carbonitrides and retainedelemental B within the matrix to increase thehardenability of the joint.

3. Autogenous GTAW joints exhibited inferiorroom temperature and high temperature tensilestrength compared to filler-added GTAW jointswith increased austenite formers (Ni) and carbideformers (Nb and Mo).

Fig. 9 Vertical section of Fe-Ni-Cr ternary diagram at 70 % Fe content(Brooks et al. 1983)

Vinoth Kumar et al. International Journal of Mechanical and Materials Engineering (2015) 10:25 Page 10 of 11

Competing interestsThe authors declare that they have no competing interests.

Authors’ contributionsAll authors read and approved the final manuscript.

AcknowledgementsThe authors wish to express their sincere thanks to M/s Mailam India Ltd,Pondicherry, India for providing financial assistance to carry out this researchwork through the Mailam India Research (MIR) Fellowship, M/s SalzgitterMannesmann Stainless Tubes Italia Srl of Italy for supplying the Super 304Htubes required to carry out this work and The Director, Naval Materials ResearchLaboratory, Ambernath, Mumbai for providing the facility to carry out hottensile testing.

Author details1Center for Materials Joining and Research (CEMAJOR), Department ofManufacturing Engineering, Annamalai University, Annamalai Nagar, TamilNadu 608002, India. 2Naval Materials Research Laboratory (NMRL),Ambernath, Mumbai 421506, India.

Received: 12 June 2015 Accepted: 9 November 2015

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