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Archives of Civil and Mechanical Engineering (2020) 20:82 https://doi.org/10.1007/s43452-020-00087-1

ORIGINAL ARTICLE

Weld joint properties of nitrogen‑alloyed austenitic stainless steel using multi‑pass GMA welding

Naveen Kumar1  · Navneet Arora1 · S. K. Goel2

Received: 6 March 2020 / Revised: 5 June 2020 / Accepted: 22 June 2020 / Published online: 30 June 2020 © Wroclaw University of Science and Technology 2020

AbstractThe welding of nitrogen (0.29 wt%)-alloyed austenitic steel (grade 23-8-N) was performed with gas metal arc welding process. Solution treatment was performed at 950 °C and 1150 °C on base metal prior to weld. Base metal after second treatment has maximum ultimate tensile strength of 942 MPa and impact toughness 66 J. The microstructures of different zones of the weld joint were characterized using an optical microscope and field scanning electron microscope (FESEM). The microhardness, tensile and impact toughness tests of the weldments were conducted along with weld ferrite evaluation. ER2209 duplex filler metal used for welding has lower C and N content which changed the weld solidification mode. Weld has microstructure containing austenite + ferrite. Being a strong austenite former, nitrogen caused minimum ferrite near weld–HAZ interface while maximum ferrite content was observed at weld centre. Weld metal has minimum while base metal has maximum microhardness. UTS (892 MPa) and impact strength (96 J) of weld made on 1150 °C solution-treated base metal were maximum as compared to other weld joints.

Keywords 23-8-N nitronic steel · ER 2209 duplex filler · Solution treatment · Characterization · Carbide · Hardness · Ferrite

1 Introduction

Alloying with nitrogen improves the corrosion resistance and mechanical properties of austenitic stainless steels (ASS) and duplex steels. There are several grades of steels available with varying nitrogen content. Any steel is called high-nitrogen steel (HNS) if it contains more than 0.4 wt% of N [1]. Nickel content of steel used for producing medical equipment causes allergy and other side effects to the human body, so it is necessary to abolish the nickel content of the equipment. Nitrogen is gaining attention as a replacement of nickel in steels due to its lower price and easy availability. Moreover, being a strong austenite stabiliser, it improves the mechanical strength without affecting the ductility and impact toughness properties. Further, nitrogen addition to ASS is found to decrease the stacking fault energy (SFE)

and improve the work-hardening capability of these mate-rials. In the last few years, nitrogen-containing ASS are getting much attention due to their superior properties [2]. Sometimes these steels are called ‘Nitronic steels’ depending upon the nitrogen content (≤ 0.4 wt%). The optimum content of nitrogen in Cr–Mn austenitic stainless steels improves mechanical and tribological properties [3].

In recent studies, new grades of nitronic steels (21-4-N and 23-8-N) have been developed. The existing applications of these steels are in internal combustion engine valves due to their high tensile strength and impact toughness after solution annealing. In addition, the nitronic steels also have excellent slurry erosion resistance which make them suitable for hydro-turbine underwater parts [4–6]. Generally, the weldability of ASS is found to be good, but nitrogen as alloying element affects mechanical and metallurgical properties of weld and base metal. During the joining of nitrogen-alloyed steels, it is also found that nitrogen is partially lost from the weld pool [7–9]. Loss of nitrogen in welding can be avoided by provid-ing sufficient nitrogen content to the weld pool which can be achieved using nitrogen-containing filler metal or enriching the shielding gas with nitrogen [1, 10]. Pure nitrogen shielding gas has increased the nitrogen content of weld metal up to 1.25%

* Naveen Kumar nkumar3@me.iitr.ac.in

1 Mechanical and Industrial Engineering Department, Indian Institute of Technology, Roorkee, Roorkee, India

2 Star Wire (India) Ltd, Faridabad, India

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and supressed ferrite formation [11]. There is no matching filler metal available for nitronic steels; therefore, the selec-tion/development of suitable filler metal to successfully weld these steels is still a challenge. There is not much literature available for welding nitrogen containing austenitic stainless steels. Some of the researchers have tried to weld nitronic steels with a variety of electrodes but still, no specific filler or shielding gas composition has been standardised. In the fric-tion stir welding (FSW) of nitrogen containing steels hardness and tensile strength were improved but carbides deteriorated its corrosion resistance [12, 13]. Earlier studies have revealed that massive carbides in nitronic steels are injurious to impact toughness and erosion resistance of these steels [5, 14, 15]. Attempts to control formation of carbides in weld metal are still not reported. Thus, this study is aimed to limit the forma-tion of carbides in weldment using a low-carbon filler metal.

In the present study, the solution treatment of 23-8-N nitronic steel was done at two different temperatures (950 ºC and 1150 ºC) with an objective to get superior mechanical and metallurgical properties. ER2209 filler wire was used to avoid weld solidification cracking. Weld joints made in each case were characterized for microhardness, tensile strength and impact strength along with change in ferrite content. The effect of carbides on hardness, tensile and impact strength has been studied in detail. In addition, the weld ferrite content was measured with optical microscope, Ferritescope, and compared with the ferrite as predicted by modified Schaeffler diagram using Cr/Ni equivalent ratio.

2 Materials and methods

2.1 Base and filler material

Base metal used in the present study was nitrogen-alloyed steel grade 23-8-N, an austenitic steel (~ 0.29 wt% N). The material for investigation was supplied by M/s Star Wire India Ltd., Ballabhgarh, Haryana, India. The chemical composition of the as-received (AR) base metal was analysed by optical emission spectrometer (Model: Metalvision 1008i). In the present study, welding was performed using ER 2209 duplex stainless steel filler wire which has a composition nearly similar to base metal for major elements (Cr, Mn, Ni, etc.) except for carbon and nitrogen. Details of the chemical compositions of the base, filler and weld metals are given in Table 1.

2.2 Solution annealing

The base metal in as-received (AR) condition was solution treated at two different temperatures, i.e. 950 °C (ST-1) and 1150 °C (ST-2) for 120 min and then water quenched. Car-bide clusters were observed in the austenitic matrix without heat treatment due to high-carbon content. The carbides are visible as bright areas, whereas austenite appears dark in the micrograph. Solution annealing resulted in the dissolution of carbide networks and also mitigated its re-precipitation along with larger average grain size.

2.3 Welding method

Welding of the base metal was done in all three conditions, i.e. AR, ST-1 and ST-2. The samples for welding were pre-pared in the dimension of 100 mm × 50 mm × 13 mm single-V butt joint with 60° groove angle having a root height 2 mm and gap 1.5 mm. Gas Metal Arc Welding (GMAW) process was used for welding with ER2209 duplex steel filler wire (1.2 mm diameter). Table 2 shows welding parameters used in the present study. Pure Argon (16 L/min) gas was used for shielding during the welding. Welding was accomplished in four passes, including the root pass. Weld tacking and fixtures were used to minimize the weld distortion.

2.4 Characterization of weldments

The samples were prepared for metallographic study as per standard procedure. Etching of base metal samples was done with aqua-regia (3HCl + HNO3) and weld metal was electrolytically etched in 20% NaOH solution to reveal the microstructure. Microstructural study of the base and weld metal was carried out using an optical microscope, FESEM

Table 1 Chemical analysis of the base metal, filler wire and weld metal (wt%)

Metal C N Cr Ni Mn S P Mo Si Fe

23-8-N base 0.32 0.29 23 7.5 2 0.03 0.04 0.12 – BalER2209 filler 0.03 0.15 22.5 8.5 1.5 0.002 0.02 3 0.60 BalWeld centre 0.08 0.22 22.8 8.47 1.73 0.004 0.03 2.37 0.45 Bal

Table 2 Welding parameters for root and filler passes

Process Pass no. Current (Amp)

Voltage (V)

Filler feed (m/min)

Welding speed (cm/min)

GTAW Root 140 24 4.3 32GMAW 1 200 32 7 27GMAW 2 191 28 5.8 27GMAW 3 182 26 5.5 25

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equipped with X-ray spectroscopy (EDS) and X-ray diffrac-tion (XRD).

Microhardness measurement was carried out in transverse directions at an interval of 0.5 mm using 300 g load and 10 s dwell time. ASTM E8/E8M standard was followed for preparing tensile test specimens. Charpy impact toughness test samples were machined as per ASTM 23-12c standard considering the notch location at two different locations (i.e. weld centre and HAZ) as shown in Fig. 1. Three samples for each condition (AR, ST-1 and ST-2) of base metal and weld metal (Weld-1, Weld-2 and Weld-2) were tested for tensile and impact energy tests for the accuracy of the results.

3 Results and discussion

3.1 Microscopic observations

Microstructures of base metal in AR and solution-treated (ST-1 and ST-2) conditions are shown in Fig. 2. In this steel, nickel is partially replaced with nitrogen without consid-erable effect on microstructure and mechanical properties. Figure 2a, b shows the clusters of carbide in AR base metal due to its high carbon content. These carbides are randomly distributed alongside the grain boundaries and metal matrix. Solution annealing at two different temperatures (ST-1 and ST-2) resulted in maximum dissolution of these carbides and also suppressed the re-precipitation of carbide clusters [16]

as shown in Fig. 2c–f. The change in the carbide content due to solution treatment was measured using Image-J software and was found to be 26%, 18.58% and 9.70% for AR, ST-1 and ST-2 conditions, respectively. The area fraction (%) was measured for carbide and austenite phases. The average of 3 different optical microscopic images at 50X magnifica-tion was taken to minimize the error. Figure 2e, f shows the residual carbides and maximum grain coarsening after solution annealing. Grain size measured for AR, ST-1 and ST-2 base metals was 17.25 μm, 20.83 μm and 27.37 μm, respectively. Generally, M7C3 and M23C6 (where M can be Cr, Mn, Fe) type of carbides were observed [5, 17, 18].

Figure 3 shows FESEM micrograph of AR base metal with EDS analysis. Spectra 1 and 2 in Fig. 3a show the varia-tion in composition at austenite and carbide regions, respec-tively. Figure 3b shows the EDS line scanning for the AR base metal throughout the different regions. The elemental concentration varies with different zones and peaks at Cr graph, and valleys in Fe graph show the massive carbide for-mation. Valleys in the graph of austenite forming elements (N and Ni) also support the carbide formation.

In the present study, the solidification mode was theoreti-cally predicted by modified Schaeffler diagram. Cr and Ni equivalents were calculated using the following equations as per modified Schaeffler diagram [19]:

(1)Creq = Cr + 1.5Mo + 1.5W + 0.48Si + 2.3V + 1.75Nb + 2.5Al,

Fig. 1 Schematic of weld joint preparation and samples extraction scheme

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The Creq and Nieq were calculated employing the above equations and considering the chemical composition from Table 1. The base metal having Creq 23.18 and Nieq 22.56 has a fully austenitic structure with Creq/Nieq ratio of 1.03. Figure 4a shows the coordinates of base, filler and weld metal on modified Schaeffler diagram (MSD) and Fig. 4b shows Fe–Cr–Ni pseudo-binary diagram indicating the metal solidification modes according to Creq/Nieq ratio.

(2)Nieq = Ni + Co + 0.1Mn + 0.01Mn2 + 18N + 30C. The Creq/Nieq ratio for base and filler metal was 1.03 and 2.18, respectively. As per MSD, the Creq/Nieq ratio of weld metal valued 1.69. The solidification mode changes from fully austenite (A) to austenite–ferrite (AF) between Creq /Nieq ratio 1.25–2.18. Further increment in Creq/Nieq ratio from 1.48 to 1.95 leads to change of solidification mode from AF to FA (ferrite–austenite) [20]. Weld joints Weld-1, Weld-2 and Weld-3 were made on AR, ST-1 and ST-2 conditions of base metal. Figure 5a–c shows the opti-cal micrographs of weld joints at the middle pass of weld

Fig. 2 Optical micrographs of base metal: a and b AR, c and d ST-1, and e and f ST-2 conditions

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centre locations. Chemical composition and solidification mode of weld metal depends on the welding conditions along with base and filler metal compositions [21]. C and N content was minimum at weld centre and gradually increased towards weld–HAZ interface. The possible rea-son for changing chemical composition from weld centre to weld–HAZ interface may be the compositional differ-ence between base and filler metals and elemental dilu-tion due to multi-pass welding. The microstructure of base metal and filler metal were fully austenitic and duplex,

respectively. Therefore, as per MSD, weld microstructure must lie between these two [22].

In agreement with the above, the weld microstructures contained dual microstructure, i.e. ferrite (δ) and austenite (γ). In the micrographs, ferrite can be seen as a dark phase, while austenite appears as a bright phase. The ferrite micro-structure has two different morphologies, vermicular (δV) and lathy (δL) depending upon the cooling rate and chemi-cal composition (Fig. 5b, c). High cooling rate resulted in more δL as compared to δV [23]. Also, it was observed that

Fig. 3 SEM EDS analysis of as-received (AR) base metal

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ferrite content of weld metal was not uniform and dependent on solidification mode and solid-state transformation from ferrite to austenite. It increased from weld–HAZ interface to weld centre and from first pass to last (top) pass of weld as last weld pass has more exposure to environment. The filler metal ER2209 has low C and N content than the base metal, but fusion zone nearby the interface enriched with C and N due to dilution and migration from base metal to weld.

Figure 6 shows the XRD patterns of base and weld metal in different conditions. The carbide peaks were vanished after solution treatment in the base metal. The austenite peak (110) has higher intensity in weld metal near weld–HAZ interface whereas the same peak at weld centre region has comparatively low intensity. The overall intensity of aus-tenite (γ) peaks was very high than the ferrite (δ) peaks at different 2θ angles.

Figure 7a–c shows the weld–HAZ interface of different weld joints. Carbide dissolution and grain coarsening were observed in HAZ region due to heating effect of the weld thermal cycle in multi-pass welding. Fine carbides were still present in the HAZ region which caused higher hard-ness as compared to weld [24]. AF (austenite–ferrite) mode of solidification was observed near weld–HAZ interface region. Variation in ferrite morphology and its content from weld–HAZ interface to weld side is seen in Fig. 7b. Solution treatment prior to weld also contributed to grain coarsening of HAZ observed in W-3 weld joint as shown in Fig. 7c.

Ferrite content of the weld metal was measured using image analysing software (Image-J), Ferritescope and modi-fied Schaeffler diagram, and the same is reported in Table 3. To ensure the reproducibility of results, an average of three readings was taken in case of Image-J and Ferritescope. The optical micrographs from weld centre locations were taken at 500X magnification for image analysing and ferrite

prediction. The maximum ferrite content was predicted by Ferritescope for each weld. Modified Schaeffler diagram pre-dicted 5–10 FN for weld centre. The ferrite content in all the above methods was found different due to various assump-tions made in MSD and errors associated with Ferritescope and non-equilibrium welding conditions.

Image of weld joint and FESEM/EDS analysis at the weld–HAZ interface is shown in Fig. 8. The weld–HAZ images have been shown at low and high magnifications. The EDS line scanning and mapping at the weld–HAZ inter-face shows the variations in chemical composition for HAZ and weld metal. Base and filler metal have compositional differences in which HAZ has higher C, N, and Cr content than weld metal near interface but dilution and diffusion resulted in slight migration of elements from high- to low-concentration region throughout the interface.

3.2 Mechanical tests

3.2.1 Microhardness measurement

The microhardness profiles of weld joints Weld-1, Weld-2 and Weld-3 in the transverse direction are shown in Fig. 9. Average hardness values of base metal in AR, ST-1 and ST-2 conditions were 347 HV, 340 HV and 321 HV, respectively. The decrease in hardness values was observed as a result of solution annealing, which caused the dissolution of carbides and nitrides along with grain coarsening. ST-2 treatment of base metal resulted in minimum hardness with maximum grain coarsening. The maximum variation of hardness was measured in AR base metal due to its heterogeneous micro-structure and carbide clusters. The peak hardness indent on carbide location is shown for curve Weld-1 with optical micrograph in Fig. 9.

Fig. 4 Locations of base, filler and weld metal on a modified Schaeffler diagram [19], b Fe–Cr–Ni pseudo-binary diagram [20]

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Solution annealing leads to alleviation of carbides den-sity and promoting homogeneity of the microstructure. C and N present at the interstitial locations resulted in solid solution strengthening. The dissolution of carbides released the C and Cr which resulted in higher C con-tent in localized zones with higher hardness. Hardness in the HAZ zone was minimum for curve Weld-3 shown in Fig. 9 due to additional effect of weld thermal cycle in multipass welding on ST-2 base metal which already has lower carbides and coarser grains. Weld zones were

observed to have a comparatively lower hardness than base and HAZ. Average hardness values of Weld-1, Weld-2 and Weld-3 joints were 308 HV, 312 HV and 304 HV respec-tively. Ratio of austenite–ferrite content and heat input in the subsequent passes affect hardness in the fusion zone. From SEM/EDS and OES analysis, it was confirmed that weld centre has minimum content of C and N. Both these elements are austenitic stabilizer, and further contribute towards enhancing the hardness of the resulted austenite phase.

Fig. 5 Optical micrographs of weld in different conditions a weld-1, b weld-2, and c weld-3

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3.2.2 Tensile testing

The tensile test results are shown in Table 4. Ultimate tensile strength (UTS) of the AR, ST-1 and ST-2 base metal was

810 MPa, 912 MPa and 942 MPa, respectively. UTS and ductility (% EL) of base metal were found to be higher, and yield strength (YS) has shown a decreasing trend with an increase in solution treatment temperature. The higher con-centration of carbides in AR base metal induces the region of high-stress concentrations at carbide/austenite interface which is the region of maximum stress concentration and has higher possibility of de-cohesion. UTS of weld joints Weld-1, Weld-2 and Weld-3 were 860 MPa, 898 MPa and 879 MPa, respectively. Stress–strain plot for base and weld are shown in Fig. 10.

C and N play a significant role in UTS and ductility of nitronic steel. Carbon released from the dissolution of car-bides, enhances the C content of the matrix resulting in higher tensile strength [25]. In addition, reduced volume of carbide clusters decreased the carbide/austenite interface area which was considered as the weakest zone. Nitrogen increases the concentration of free electrons, and hence promotes metallic character of atomic interactions, whereas carbon enhances localization of electrons. Further, higher N content of this steel resulted in strain hardening of the samples during tensile testing and led to higher tensile strength. Nitrogen atoms interact with dislocations in the

Fig. 6 XRD patterns of base and weld metal in different conditions

Fig. 7 Optical micrographs of weld interface in conditions: a Weld-1, b Weld-2, and c Weld-3

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austenite matrix and increase the internal friction and lower the stacking fault energy (SFE). This reduction in SFE fur-ther restricts the dislocation movement. In ASS, nitrogen increases the coefficient ‘k’ of the Hall–Petch equation, which resulted in higher strength of the material [26]. Ten-sile strength of the weld samples was lower than base metal in each condition and the same has been shown in Table 4. Change in the ferrite morphology due to cooling rate and

Table 3 Ferrite measurement at weld centre

Weld condition Measured Ferri-tescope (%)

PredictedMicroscopy (area %) Modified Schaef-

fler diagram (FN)

Weld-1 20.7 23 5–10Weld-2 17.5 20 5–10Weld-3 16 18 5–10

Fig. 8 SEM/EDS analysis at weld–HAZ interface. a Lower magnification and b higher magnification image of weld joint

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compositional thrust resulted in different UTS and ductility of weld joints. The presence of nitrogen can affect the weld cooling rate and results in diverse weld microstructures for various passes of weld [17].

Figure 11a–c shows the tensile fractured surfaces of base metal in AR, ST-1 and ST-2 conditions. The presence of cleavage facets with comparatively flat surface in AR condi-tion shows the brittle type of fracture. Increase in dimples on the surface of ST-2 resulted in mixed mode of fracture. Weld metal shows comparatively ductile fracture with the presence of dimples as compared to the base metal.

3.2.3 Impact toughness testing

Table 5 shows the Charpy impact toughness test results for base metal, weld metal and HAZ. The base metal exhibited toughness values of 21 J, 28 J and 66 J in the AR, ST-1 and ST-2 conditions, respectively. Impact toughness is observed to be primarily governed by the presence of carbides in the base matrix. The carbide/austenite interface (de-cohesion region) has high-stress concentration and contributes to crack initiation. Further, carbides are brittle in nature with sharp edges and, therefore, if a notch is made in this region, the fracture takes place in brittle manner resulting in low-ering of impact toughness in the region of high carbide concentration [17]. Very low impact toughness in AR base metal is due to presence of massive carbides. Improvement

in toughness in ST-1 and ST-2 condition is due to dissolu-tion of carbides as a result of solution annealing followed by water quenching. The base metal in ST-2 condition showed a high impact value of 66 J owing to effecting dissolution of carbides in the matrix. The concentration of dissolved C and N in the base metal also increases as a result of solution treatment. This leads to solid solution strengthening which also contributes towards improvement in impact toughness.

The welded samples with notch at weld centre and HAZ zone shows higher impact toughness than the base metal. The impact toughness of Weld-1, Weld-2 and Weld-3 joints was 90 J, 94 J and 96 J, respectively. The minimum hardness at weld centre of Weld-3 supported the maximum toughness. Ferrite morphology and ferrite/austenite ratio also contrib-uted to impact toughness of weld [27]. Increasing acicular ferrite in the weld microstructure contributes to the higher impact strength due to its interlocking nature and fine grain size which increases resistance to crack propagation.

The impact toughness of the HAZ was also analysed by making a notch in the HAZ zone. The impact-fractured specimens with notch in HAZ region are shown in Fig. 12f. It was found that impact values for HAZ were improved with solution treatments and was maximum (52 J) for HAZ in weld-3. The impact strength of HAZ region (notch made in HAZ) was dependent on the direction of crack propaga-tion. In case when the crack propagated forms notch at HAZ

Fig. 9 Microhardness profile of weld joints

Table 4 Tensile test results Metal/condition YS (MPa) UTS (MPa) Elongation (%) Fracture location

Base AR 556 810 ± 22 Sample centre Sample centreBase ST-1 478 912 ± 18 Sample centre Sample centreBase ST-2 370 942 ± 9 Sample centre Sample centreWeld-1 559 860 ± 15 19.56 WeldWeld-2 537 898 ± 13 15.88 WeldWeld-3 435 879 ± 10 21.40 Weld–HAZ interface

Fig. 10 Stress–strain plots for tensile test results of base and weld metal

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to base side, it shows lower impact strength whereas the crack propagated within HAZ or towards weld side resulted in higher impact toughness. The presence of fine residual

carbides in the HAZ region also plays a key role on the crack propagation and impact energy as compared to base metal.

Figure 12 shows the FE-SEM micrographs of impact frac-ture surfaces. In AR condition (Fig. 12a), fracture surface reveals the presence of cleavage facets that confirms the brit-tle fracture of the specimen and it was also justified with the impact results (21 J). After solution annealing, impact toughness was found to be increased which was justified with the presence of dimples at fracture surface. For ST-1, mixed mode of the fracture is observed and fracture surface shows the presence of ductile area of fine shallow dimples as well as cleavage facets (Fig. 12b). In ST-2 condition, a ductile mode of fracture is observed and fracture surface is characterized with fine and coarse dimples, micro-voids with fewer amount of the brittle area. The fractographs of Charpy test specimens for weld metal and HAZ are represented in Fig. 12d, e. The Charpy test specimens of weld fusion zone exhibited fully ductile fracture and had mesh of very fine

Fig. 11 FESEM fractographs of tensile test specimens a AR, b ST-1, c ST-2 and d weld metal at centre

Table 5 Results of Charpy impact toughness test

Metal/condition Notch location Impact energy (J)

Fracture propa-gation

Base AR Centre 21 ± 5 –Base ST-1 Centre 28 ± 4 –Base  ST-2 Centre 66 ± 2.5 –Weld-1 Weld centre 90 ± 3 Within weld

HAZ 32 ± 4.5 Towards weldWeld-2 Weld centre 94 ± 3 Within weld

HAZ 38 ± 4 Towards baseWeld-3 Weld Centre 96 ± 2.5 Within weld

HAZ 52 ± 3 Towards base

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dimples as shown in Fig. 12d. Fracture surface in HAZ zone shows the mixed mode (Fig. 12e) with the presence of both cleavage facets and dimples.

4 Conclusions

The present study compares the mechanical and metallurgi-cal properties of 23-8-N austenitic steel welded with duplex steel filler wire. The conclusions are

• Solution treatment of the nitronic steel (grade 23-8-N) at 950 °C and 1150 °C dissolved the carbide clusters, and area fraction (%) was decreased from 26% for AR to 18.57% and 9.70% for ST-1 and ST-2 base metal, respec-tively.

• Ultimate tensile strength of base metal was improved from 810 MPa for AR to 912 MPa 942 MPa for ST-1 and ST-2 conditions. The impact toughness was also found 21 J, 28 J, and 66 J for AR, ST-1 and ST-2 conditions, respectively.

• The dilution of carbon and nitrogen at the weld–HAZ interface enriched the nearby weld region and resulted

in austenite–ferrite mode of weld solidification. Ferrite content was found maximum at weld centre and it was reduced towards weld–HAZ interface.

• The average hardness of base metal in each condition (AR, ST-1 and ST-2) was higher than the HAZ and weld metal.

• Ferrite content and morphology affected the tensile strength and ductility of the weld metal. Maximum impact toughness of weld metal was found to be 96 J which was higher than each condition of the base metal.

• The maximum impact toughness in HAZ region was 52 J which was less than the weld metal. Location of notch tip has considerable effect on crack formation.

Acknowledgements The authors gratefully acknowledge Prof. Desh Bandhu Goel, Indian Institute of Technology, Roorkee, for giving their technical inputs. The authors are also thankful to Star Wire (India) Ltd. for providing material to perform the study.

Compliance with ethical standards

Conflict of interest There is no conflict of interest.

Ethical statement I ensure that the work described has been carried out in accordance with Publishing Ethics.

Fig. 12 Fractured surfaces of base metal: a AR, b ST-1, c ST-2, d weld metal, e HAZ and f fractured samples

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