enhancement of microstructural and mechanical properties

ISIJ International, Vol. 55 (2015), No. 7

© 2015 ISIJ1439

ISIJ International, Vol. 55 (2015), No. 7, pp. 1439–1447

* Corresponding author: E-mail: [email protected]: http://dx.doi.org/10.2355/isijinternational.55.1439

1. Introduction

Ferritic stainless steel (FSS) with 11–12% Cr developed to feel the gap between high cost austenitic stainless steels (ASS) and the rust prone carbon steels has attracted as low cost utility stainless steels. These modified FSS grades are now commonly used in the coal mining industry for bulk transport of coal and gold, for cane and beet sugar process-ing equipment, road and rail transport, power generation, petrochemical, pulp and paper industries etc. In fact, the use of these steels in the past few years has increased markedly with their successful applications in passenger vehicles, coaches, buses, trucks, freight and passenger wagons.1–3)

The gas metal arc welding (GMAW) is commonly used for fabricating various components of ferritic stainless steel. One of the unique characteristics of GMAW process is the way molten metal is transferred across the arc. Metal trans-fer is controlled by several parameters, including current, voltage, polarity, electrode extension, shielding gas compo-sition, and electrode diameter. Now a day, pulsed GMAW (P-GMAW) process is used to obtain the best of both spray mode of metal transfer (S-mode) and short circuit mode of metal transfer (SC-mode). The pulsed GMAW process works by forming one droplet of molten metal at the end of the electrode per pulse. Then, just the right amount of cur-rent is added to push that one droplet across the arc and into

Enhancement of Microstructural and Mechanical Properties by Pulse Mode of Metal Transfer in Welded Modified Ferritic Stainless Steel

Manidipto MUKHERJEE,1) Anupama DUTTA,1) Prasanta KANJILAL,2) Tapan Kumar PAL1)* and Sunil SISODIA3)

1) Metallurgical and Material Engineering Department, Jadavpur University, Kolkata, 700032 India. 2) National Test House, Saltlake, Kolkata, 700091 India. 3) Salem Steel Plant, Steel Authority of India Ltd., Salem, Tamil Nadu, 636013 India.

(Received on October 3, 2014; accepted on March 13, 2015)

The present study describes the enhancement of microstructural and mechanical properties by pulse mode of metal transfer in welded modified ferritic stainless steel (409 M) sheets (as received) of 4 mm thickness. The welded joints were prepared by varying modes of metal transfer at different heat input conditions (i.e. pulse mode at 0.5 kJ/mm and 0.9 kJ/mm and spray mode at 0.5 kJ/mm), using austenitic filler wire (i.e. 308 L) under Ar + 10% CO2 atmosphere. It has been observed that the pulse mode of metal transfer significantly alters the weld metal composition compare to spray mode which promotes com-paratively stable austenite in the welds and also depicts significant enhancement in grain structure even with the higher heat input condition. Present study clearly shows that pulse mode enhances micro-hard-ness of welded joints and toughness values of weld metals compare to spray mode of metal transfer for a particular heat input.

KEY WORDS: modified ferritic stainless steel; mode of metal transfer; microstructure; mechanical proper-ties.

the puddle. Unlike conventional GMAW, where current is represented by a straight line, pulsed GMAW drops the cur-rent at times when extra power is not needed; therefore the process needs to be cooled off. It is this “cooling off” period that allows pulsed GMAW to weld better on thin materials, control distortion and run at lower wire feed speeds.4) The beneficial effects most often reported in the literature are: 1) the total heat input to the welding decreased; 2) the width of the heat-affected zone is reduced; and 3) the depth/width ratio of weld bead increased by high frequency.5) However, it is not too easy to select the values for these pulse param-eters, since in case of each welding condition (base material, electrode material and diameter, shielding gas type, etc.) there is an optimum parameter combination.6) Several stud-ies reported the beneficial effect of pulse mode, however; there is not much literature available where enhancement by pulse mode had been compared with the other modes of metal transfer. In order to attain maximum benefit of P-GMAW for a given application, such knowledge should be acquired.

The metallurgical changes in microstructure of weld-ments are strongly influenced by the variation in modes of metal transfer. In the ferrite-austenite dissimilar weld metals two types of martensite phases (i.e. ε-martensite and α ′-martensite) can form on cooling which is being a result of a diffusionless phase transformation. The results of X-ray diffraction, magnetic measurements, and transmission electron microscopy lead the researchers to believe that the most probable way of the phase transformation is the γ →


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ε → α ′.7) It was found that α ′-martensite phase is formed almost exclusively at the intersection of twin faults or shear bands (bands of imperfect micro twins), stacking faults, and ε-phase.8) These variations in martensite content lead to the desirable and/or detrimental mechanical properties of welded joints. It was reported previously9) that spray mode induces more amount of metastable grain boundary austenite in the weld metal. This grain boundary austenite ultimately trans-forms to lath martensite due to high heat input and slower cooling associated with this mode. Whereas in short circuit mode, the cooling rate being faster, suppress the formation of grain boundary austenite and thus lath martensite as well. The spray mode of metal transfer (S-mode) produces more amount of grain boundary austenite along with lath martens-ite which inhibits ferrite grain growth in the weld metals and induces higher strength and toughness. Although in case of high temperature heat affected zone (HTHAZ) completely reverse phenomenon are being observed where short circuit mode produces fine grain structures and better metallurgical properties due to low heat input condition. It is also reported that different heat input conditions in particular mode of metal transfer also have profound effects on the formation of martensite and weld metal toughness.10) However, the effect of martensite on mechanical properties is controver-sial and may promote hydrogen induced cracking.11) Hence, it is obvious that control of martensite content is essential to improve the performance of welded ferritic stainless steel in different applications.

In the present study, welding parameters were varied to operate with two different modes of metal transfer i.e. pulse mode and spray mode. Pulse mode (P-GMAW) further var-ied with two heat input conditions i.e. 0.5 kJ/mm and 0.9 kJ/mm and the spray mode was carried out at 0.5 kJ/mm heat input condition using austenitic filler wires i.e. 308 L, under Ar+10% CO2 atmosphere. In the present work minimum heat input of 0.5 kJ/mm was referred from the previous work.9) The aim is to get some insight on the enhancement of microstructural and mechanical properties by pulse mode of metal transfer in welded modified ferritic stainless steel compare to spray mode of metal transfer at the same and high heat input condition.

2. Experimental Procedure

2.1. MaterialThe hot rolled sheets of 4 mm thick 409 M grade modi-

fied ferritic stainless steel were cut into the required dimen-sion (200 mm × 70 mm × 4 mm) and used for the weld joints preparation. Figure 1 shows the details of weld joint preparation and test plate assembly used for the welding process. The chemical compositions of the base metal and austenitic filler metal (308 L of diameter 1.2 mm) are given in Table 1.

2.2. Welding ProcedureThe experiments were conducted using a KEMPPI water

cooled universal MIG/MAG machine (Model: FastMIG Pulse 450) using DC electrode positive (DCEP). The weld-ing conditions and process parameters used to fabricate the joints are given in Table 2. The welding operations were performed using Ar+10%CO2 shielding gas mixture with the gas flow rate of 15 L/min. Square butt joints with a root gap of 1.5 mm were fabricated in GMAW process using the selected welding parameters so the spray and pulsed spray (P-GMAW) modes of metal transfer could be oper-ated. To ascertain the operating mode, oscilloscope (Model: DLM2000; Make: Yokogawa Electronic Co., Japan) and

Table 1. Chemical compositions of base metal and filler metals.

Type Chemical composition (in wt %)

SSP 409 Mbase metal

C% Si% Mn% P% S% Cr% Ni% Mo% Cu% Nb% N%

0.030 0.463 0.79 0.029 0.014 11.10 0.31 0.033 0.026 0.017 0.01

308 L filler metal 0.015 0.53 1.68 0.012 0.03 19.53 9.26 0.117 0.082 0.026 0.053

Table 2. Welding conditions and process parameters.

Sample Specification

Ip(A)

Tp(ms)

Ib(A)

Tb(ms)

F(Hz)

Im(A)

Voltage (V)

WS(mm/min)

HI(kJ/mm) MTM

P1 350 2 110 6 125 170 25.0 400 0.57 Pulse

P2 370 2 130 6 125 190 27.0 300 0.92 Pulse

S1 — — — — — 220 25.0 500 0.59 Spray

Note: Ip= peak current; Tp= peak time; Ib= background current; Tb= background time; F= pulse fre-quency; Im=mean current; WS=welding speed; HI=Heat Input [(I×v×η×60)/(WS×1 000)]; MTM= metal transfer modes.

Fig. 1. Weld preparation and test plate assembly.


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high speed video camera (Model: Phantom V311, Make: Vision Research Inc., USA) were used to record current pulsation, voltage, frequency and drop transfer respectively during welding run as shown in Fig. 2. The spray transfer mode as shown in Fig. 2(a) is identified as almost constant potential drop, whereas Fig. 2(b) reveals pulsed waveform along with one drop per pulse (ODPP) transfer in single pulsation. The mean current of the pulse had been calculated using the following equation.12)

II T I T

T Tm

p p b b

p b

=++

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

Where, Ip, Ib, Tp and Tb are the peak current, background current, peak time and background time, respectively.

Furthermore, all necessary care had been taken to avoid joint distortion by applying proper clamping devices. The soundness of all the welded plates was examined using the radiography testing.

2.3. Metallographic StudyThe specimens of metallography comprising weld metal,

HAZ and base metal were polished using different grades of emery papers followed by final polishing in the disc polishing m/c using diamond compound (2 μm particle size). Microstructural examination was carried out using a light optical microscope (Make: Carl ZEISS India Pvt. Ltd.; Model: Imager. A1m). The XRD measurements were

performed in a Rigaku X-ray diffractometer (made in Japan) using the monochromatic Cu Kα radiation operated at 40 kV tube voltage and 30 kA tube current, with a scan rate of 1° per min, with 2θ ranging from 40° to 100°. The weld metals were also examined under Transmission Elec-tron Microscope (TEM) (Make: Philips Ltd., Netherlands; Model: CM-70). The specimens from weld metal were pre-pared using diamond cutter and fine grades of emery papers up to 0.1 mm followed by final thinning in chemical etching/thinning process up to 10 μm for TEM examination.

2.4. Dilution CalculationThe percentage dilution (pct. DL) was calculated from the

geometrical characteristics of weld joints such as total area of weld deposit (AWD), area of top (ATR) and root (ARR) reinforcement, area of base metal fusion (ABF), and area of root gap (ARG), as schematically shown in Fig. 3. The esti-mation of AWD, ABF and pct. DL was found out as follows:10)

A A A A AWD RG TR BF RR= + + + ................. (2)

A A A A ABF WD RG TR RR= − − − ................. (3)

D AL BF WD% A / %= ×100 ......................... (4)

2.5. Calculation of Cr–Ni Equivalents and Martensite Transformation (MT) Temperatures

Weld metal compositions obtained from dilution were used to calculate the chromium equivalent (Creq) and nickel equivalent (Nieq) values using the following equations.9)

Cr Cr Mo Nbeq = + + ×% % . %0 7 .............................. (5)

Ni Ni C N Cueq = + × + × + ×% % % . %35 20 0 25 ...... (6)

In order to predict the presence of α ′- and ε- martensite, martensite start temperature (Ms) and ε-martensite start tem-perature (Mεs) i:e: martensite transformation (MT) tempere-tures, for different weld metals, were calculated using the following equations:13)

M (K) = As C N3 199 8 1 4 17 9 21 7

6 8 45 0 5

− × + − × − × −× − × −

. ( . ) . Ni . Mn

. Cr . Si 55 9 1 9 1 4

14 4

. Mo . (C . N)

(Mo Cr Mn) . [(Ni Mn)

(Cr Mo Al Si

× − × + ×+ + − × + ×

+ + + ))] /1 2 410− ....................................... (7)

Fig. 2. (a) Potential drop vs. Time curve of spray transfer mode and (b) Typical oscilloscope and high speed video camera record of weld run P1 reveals pulsed waveform with ODPP mode of metal transfer.

Fig. 3. Schematic view of different locations of welded joint con-sidered dilution calclation.


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M K A C Nsε ε( ) . ( . ) . Ni . Mn

. Cr . Si .

= − × + − × − × −× + × −

710 5 1 4 18 5 12 4

8 4 13 4 1 66 22 7 11 6

1 4 3 7

× − × + ×+ × + + − × + ×

+

Mo . Al .

(C . N) (Mo Cr Mn) . [(Ni Mn)

(Cr Mo++ + +Al Si)] /1 2 277 .......................................... (8)

where, A3 is the γ → α phase transformation temperature of pure iron i.e. 910°C, Aε is the fictitious phase transformation temperature of γ → ε, about 390 K.13)

2.6. Grain Size MeasurementGrain size of weld metals were evaluated by linear inter-

cept method (as per ASTM E1382) from the optical micro-graphs under an optical microscope using image analyzing software (AxioVision/AxioCam version 4.6) and average of five readings is reported.

2.7. Ferrite and Martensite MeasurementFerrite percent of each weld metal was evaluated using

Feritscope FMP30 and average of ten readings are reported. The actual α ′-martensite content (Cα ′) had been com-puted from feritscope data F (%Ferrite) using relation Cα ′ (mass%)=1.7×F14) and average of five readings is reported.

2.8. Micro-Hardness TestingMicro-hardness survey was made on flat metallographic

specimen across the joints in Vickers’s microhardness test-ing machine (Make: LECO Co., USA; Model: LM248AT) using 100 gf load at an interval of 500 micron.

2.9. Charpy Impact TestingSub-size charpy impact specimens, due to smaller plate

thickness, were prepared as shown in Fig. 4 to evaluate the impact toughness of weld metal. The V-notch with a notch-tip radius of 0.25 mm was made parallel to the welding direction. Impact test was performed at room temperature using pendulum type impact testing machine as per ASTM E23-07 and average of four readings is reported.

2.10. Scanning Electron MicroscopyThe fractured surfaces of the impact tested specimens

were examined under Scanning Electron Microscope (SEM) (Make: JEOL Ltd., Japan; Model: LSM-6360) to understand the micro-mechanism in fracture.

3. Results and Discussion

3.1. Evolution of Weld Metal MicrostructureThe effect of different modes of metal transfer on geo-

metrical characteristics of welded joints with respect to its total area of weld deposit (AWD), total area of base metal fusion (ABF), and dilution of base metal (DL) are given in

Table 3. It is observed that as the heat input within pulse mode (Table 2) increases base metal dilution also increases. Whereas spray mode of metal transfer has highest DL because the high current situated with this mode increases the current density as well as the heat content of the molten metal which tends to enhance the area of base metal fusion. Weld metal compositions obtained from dilution calcula-tion are given in Table 4. The weld metal compositions were then used to evaluate the chromium equivalent (Creq) and nickel equivalent (Nieq) values9) using the Eqs. (5) and (6) along with Creq/Nieq ratios and the values are given in Table 4. The Creq/Nieq ratios are able to provide important information on stability of austenite (phase transformation) and primary solidification of weld metals. It is revealed from Table 4 that the weld metal P1 has lowest Creq/Nieq ratio followed by P2 and S1. High Creq/Nieq ratio increases the stability of the δ-ferrite by shifting the solidification line away from the triple point (i.e. L+γ+δ zone) into the δ-ferrite region. Whereas, low Creq/Nieq ratio increases the stability of γ-phase in the weld microstructure. Therefore weld metal P1 should have higher stable austenite in the final microstructure compare to other weld metals. Again high Creq/Nieq ratio increases the metastability of austenite phase which ultimately transform into martensite through solid-state phase transformation during cooling. In order to predict the presence of α ′- and ε-martensite, martensite starts temperature (Ms) and ε-martensite start temperature

Fig. 4. Dimension of Sub-size impact specimen.

Table 3. Dilution calculation for different welds.

SI. No. ABF (mm2) AWD (mm2) %DL

P1 8.517 25.27 33.72

P2 17.84 47.76 37.35

S1 16.65 35.46 46.96

Table 4. Composition of different weld metals derived from dilu-tion and Creq, Nieq, Creq/Nieq ratio and Martensite Trans-formation (K) temperature.

Elements P1 P2 S1

C 0.020 0.020 0.022

Si 0.507 0.505 0.498

Mn 1.379 1.347 1.262

P 0.017 0.018 0.019

S 0.024 0.024 0.022

Cr 16.687 16.381 15.572

Ni 6.242 5.917 5.058

Mo 0.088 0.085 0.077

Cu 0.063 0.061 0.055

Nb 0.022 0.022 0.021

N 0.038 0.037 0.032

Creq 16.79 16.482 15.665

Nieq 7.73 7.39 6.49

Creq/Nieq 2.17 2.23 2.41

Ms[K (°C)] 324.37 (51.37) 339.05 (66.06) 378.09 (105.09)

Mεs[K (°C)] 320.97 (47.98) 331.84 (58.85) 360.70 (87.71)


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(Mεs) i.e. martensite transformation (MT) temperetures, for different weld metals were calculated using the Eqs. (7) and (8), respectively. The calculated results of Ms and Mεs for different weld metals are also given in Table 4. Weld met-als, in general, have lower Mεs temperature values compare to Ms temperature and therefore, γ → α ′ transformation will take predominant role over γ → ε transformation.13) However, the Ms and Mεs values for P1 being very close (less than 5°C) to each other compare to weld metal P2 and S1, the possibility of the ε-martensite present in weld metal P1 will be more. Nevertheless, it is well known that γ → α ′ transformation has more thermodynamical stability over γ → ε transformation15,16) thus under any stress such as residual stress during welding, ε-martensite will transform into α′-martensite.8) Therefore, it can be assumed that the weld metals should contain predominantly α′-martensite in their final microstructure. Also, MT temperatures have been correlated with Creq/Nieq ratios of different weld metals as shown in Fig. 5(a). The weld metal compositions in terms of Creq/Nieq ratio have created a variation in the MT tem-peratures among the weld metals and hence the amount of martensite laths. Accordingly, P2 and S1 having higher Creq/Nieq ratio and MT temperatures than weld metal P1 (Table 4) should provide higher amount of martensite laths. To identify different phases in the weld metals X-ray diffraction patterns of three weld metals were taken (Fig. 5(b)) which clearly show the presence of γ-phase, and α′-martensite. However the variation of α′-martensite in the weld met-als are not clearly depicted in XRD analysis. Therefore to clearly understand the variation of α′-martensite content in weld metals, quantitative analysis by the magnetic induc-tion measurement has been carried out and the values are presented in Fig. 5(c). It indicates that, S1 has highest amount of α′-martensite followed by weld metal P2 and P1. Higher Creq/Nieq ratio and MT temperature of S1 should be the primary cause of higher martensite fraction. Whereas stable austenite formation in P1, due to lower Creq/Nieq ratio,

mainly restricted the solid-state phase transformation (γ → α ′) during cooling and decreases the amount of martensite.

Depending on the Creq/Nieq ratio (Table 4), in the pres-ent study, the solidification mode of all weld metals can be categorized in the following way:17,18)

F e L L Cr Nieq eqmod : : / .→ + → → >δ δ γ 1 95

F mode of solidification results in complete formation of ferrite, which may partially transform into austenite during cooling and ultimately leads to the formation of the Widmanstatten structure.19) However, depending upon the metastability of the phases, due to variation in weld compo-sition, solidification and transformation can occur in differ-ent ways, i.e. the precipitation of primary ferrite, plus three phase reaction (ferrite, austenite and liquid) at the terminal solidification stage, and δ → γ continuing below the solidus line.20) Hence, the final microstructure should consist of ver-micular δ as a primary phase within the dendrite arms envel-

Fig. 5. (a) correlation between Creq/Nieq ratio and MT temperature; (b) XRD pattern and (c) α′-martensite content of three weld metals.

Fig. 6. Optical micrographs of weld metal (a) P1; (b) P2 and (c) S1 show various grain morphology along with the presence of GBA, IGA and WA.


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oped by the GBA, IGA and WA.19) The complexity of the weld microstructure arises from the fact that, after a certain amount of primary solidification as ferrite, austenite precipi-tation occurs through a peritectic/eutectic reaction (L+δ → L+δ+γ) and remain into the two phase (δ+γ) solidification region leading ultimately to primary ferrite and different austenite formation in final microstructures. The austenite formed at grain boundaries of welds are metastable depend-ing upon their respective Creq/Nieq ratios and tends to trans-form into lath martensite (denoted as α ′).9) According to Mukherjee and Pal,9) more metastable γ-phase produces higher amount of lath martensite in the welded structure through solid-state phase transformation (γ → α ′). It is worthwhile to mention here that, P2 and S1 weld metals having higher Creq/Nieq ratios and MT temperatures create more metastable γ-phase and should produce higher amount of lath martensite (Fig. 5(c)).

Optical micrographs in Figs. 6 and 7 show that all weld metals primarily consist of three phases i.e. ferrite (δ), austenite (γ) and martensite (α′). However, there are some noticeable variations in between the micrographs of different weld metals. In the micrographs austenite revealed as thin layer of grain boundary austenite (GBA) or Widmanstätten austenite (WA) as shown in Fig. 6(a). GBA can be identified as a bulk single phase structure along the ferrite grain boundaries and sometimes these structures are extended into the matrix as a thin layer which is eventually known as Widmanstätten structure. On the other hand, Figs. 6(b) and 6(c) shows more complex microstructures of weld metals P2 and S1.The area represented by the micrographs (Figs. 6(b) and 6(c)) has undergone a growth of GBA in the interfaces between δ-ferrite grains and a small portion of intragranular austenite (IGA) within the δ-ferrite grains. Furthermore, bright-field TEM micrograph of weld met-als (Fig. 7) clearly reveal the presence of lath martensite (predominantly α ′-martensite) along with the γ phase in the ferrite matrix. From the micrographs it is revealed that the weld metal S1 has higher lath martensite content (density

of darker region) followed by weld metal P2 and P1 which is in sequence with the previous observation in Fig. 5(c). Lath martensite is predominantly formed along the grain boundaries in the weld metals and originated through solid state phase transformation from γ phase during cooling. TEM micrographs also reveal that dislocation pile-ups (Ds) along the grain boundaries in the γ phase are act as a precur-sor to γ → α ′ transformation instead of ε-martensite which may varies with the weld metal composition. Therefore, the presence of ε-martensite in the weld metals is assumed to be negligible and hence, γ → γ +Ds → γ + Ds + α ′ trans-formation will take predominant role over γ → γ + ε → γ + α ′ transformation. From the observation it can be elucidated that after complete solidification and transformation, weld microstructure should contain primary ferrite, austenite and lath martensite at room temperature.

The grain size of weld metals were evaluated from the optical micrographs and the values are shown in Fig. 8. According to Fig. 8, weld metals attribute finer grains com-pared to HTHAZ, irrespective of modes of metal transfer. Figure 8 also reveals that the grain size of weld metals is more or less comparable with the BM. It is interesting to note that finest grains among the weld metals are obtained in weld metal P1. Again, the grain size of weld metal P2 is very close to that of S1. Interestingly, inspite of same heat input weld metal P1 has finer structure compare to weld

Fig. 7. Bright-field TEM micrograph of weld metal (a) P1; (b) P2 and (c) S1 reveal the presence of austenite (γ) and mar-tensite (α′) in the ferrite matrix. Martensite transformation from γ phase is governed by the dislocation pile-ups along the grain boundary.

Fig. 8. Average grain size of base metal, weld metals and HTHAZs.


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Table 5. Average hardness values of weld metal and HAZ.

Sample No. Avg. hardness of weld metal Avg. hardness of HAZ

P1 380.06 294.92

P2 323.54 284.69

S1 359.80 292.83

metal S1. Even with the high heat input (0.9 kJ/mm), associ-ated with weld metal P2, the grain did not grow much and remain comparable with the weld metal S1. This variation in grain structure can be explained by the pulsing mechanism. In pulse mode current pulsing (Table 2) results in periodic variation of arc forces and higher turbulence in weld pool which may reduce the temperature in front of the solidify-ing interface. The amplitude of thermal oscillations in the weld pool enhances with pulse mode compare to spray mode. The temperature fluctuation inherent in pulsed weld-ing also leads to a continuous change in the weld pool size and shape, which favours the growth of new grains at the weld centre. The new grains have been found to be regularly oriented toward the direction of higher thermal gradient which may yield the shift in preferred growth direction21) and generates fine grain structure. However the grain growth in pulse mode is proportional to the heat input. Increase in heat input endures slow cooling and allows grain to grow up to a certain extent in P2 though it remains comparable to S1 due to the grain refining mechanism. On the other hand, the higher amount of martensite formed at grain boundaries from metastable γ phase, due to higher Creq/Nieq ratio and MT temperature (Table 4); mainly control the grain growth of ferrite (δ) and austenite (γ) in weld metal S1 during cooling. Thus, it can be stated that the pulse mode of metal transfer can extensively refine the weld metal grain structure in comparison to spray mode of metal transfer.

3.2. Evolution of HTHAZ MicrostructureThe typical heat affected zone (HAZ) microstructures

of different welded joints as shown in Fig. 9 reveals fer-rite with some martensite. The high temperature HAZ (HTHAZ) adjacent to the fusion line represents base metal heated above the A3 temperature (i.e. 910°C) during the weld thermal cycle, and is characterized by δ-ferrite grains surrounded by grain boundary martensite. Figure 9 also depicts that the amount of martensite present in the micro-structures is inadequate to control the grain coarsening in HTHAZ.9) Therefore, despite the partial solid-state phase transformation from ferrite to austenite on cooling, the HTHAZ is characterized by a coarse grain size. The grain size of HTHAZ (Fig. 8) is directly related with the applied

heat input (Table 2). Increase in heat input decreases the cooling rate and allows more time for HTHAZ grain growth. Therefore lower heat input associated with weld metal P1 produces relatively finer structure compare to P2. However the grain size of P1 and S1 has considerable differences although their heat input values are the same. The higher current density associated with S1 increases the heat content in HTHAZ even with the lower heat input and probably the reason of substantial grain growth.

3.3. Correlation of Mechanical Properties with Micro-structure

3.3.1. Micro-hardnessThe microhardness values are plotted in Fig. 10 and the

average microhardness of different welded zones are given in Table 5. The average hardness of weld metal (~350 Hv) is higher than HAZ (~290 Hv) and BM (~200 Hv). Furthermore, among the weld metals of pulse mode, P1 having comparable martensite content (Fig. 5(c)) and finer grain size (Fig. 8), are expected to provide higher hardness. However, when comparing between pulse mode and spray mode, weld metal P1 and S1 reveal substantial deference in micro-hardness values although both the weld metals having same heat input. This is probably associated with the finer grain size of P1 compare to S1. The substantial differences in micro-hardness values are also prudent between weld metal P2 and S1 although their grain size is comparable. Here dominating role has been played by the formation of higher α ′-martensite at the grain boundaries in weld metal S1 which tends to increase the dislocation density, resulting in higher hardness compare to the weld metal P2 where mar-tensite content is less (Fig. 7). On the other hand, HTHAZ having coarser ferrite grains (Fig. 8) with lesser amount of lath martensite and separately placed dislocations provide lower hardness than weld metals as shown is Fig. 10. Inter-estingly, there are not many differences between HTHAZ hardness values. However P1 having finer grain structure provide slightly higher HAZ hardness where as the S1 and

Fig. 9. Optical micrographs of HTHAZ (a) P1; (b) P2 and (c) S1 show coarse grain structure.

Fig. 10. Micro-hardness of welded joints.


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P2 offers comparable HAZ hardness values.

3.3.2. ToughnessCharpy impact toughness values of all the weld metals are

illustrated in Fig. 11. The impact toughness of base metal is 29.41 J and impact toughness values of weld metal S1, P1 and P2 are 36.23 J, 41.33 J and 33.46 J respectively. It clearly indicates that there is an increase in toughness value of weld metal compared to base metal irrespective of the modes of metal transfer. However, among the weld met-als, P1 exhibits highest impact toughness followed by S1 and P2. Nevertheless, variation in impact toughness values among weld metal S1 and P2 is less.

The weld metals toughness is undoubtedly dependent upon the several factors such as amount of metastable austenite, martensite transformation, carbide precipitation, grain size etc. Gooch and Ginn22) have showed second-ary cleavage to be arrested at martensite colonies through metallographic examination. Due to this effect, total energy absorbed during fracture will increase at higher martensite contents. The grain growth of delta ferrite at high tempera-

Fig. 12. Fractographs of impact specimens (a) P1; (b) P2 and (c) S1.

tures is also be restricted by a higher fraction of martensite on the grain boundaries. Improved toughness should thus being obtained in weld metals containing higher fraction of martensite. In the present study, weld metals, in general, having higher fraction of martensite together with equiva-lent grain structure compare to base metal have attributed higher toughness values (Fig. 11). However, the impact toughness values among weld metals show relatively com-plex behaviour involving several dependence factors. In the present study, weld metal P1 having higher amount of stable austenite together with comparable amount of martensite formation (Figs. 5(c) and 7) along with finest grain struc-ture (Fig. 8), attributed relatively higher toughness. This phenomenon only indicates that the dominant role played by the grain size and the stable austenite. Weld metal P2 attributed comparatively coarse grain structure (Fig. 8) with lower amount of lath martensite (Figs. 5(c) and 7) which ultimately results in comparatively lower toughness by decreasing the absorbed energy during fracture. However in comparison between S1 and P2, higher martensite forma-tion in weld metal S1 hindered secondary cleavages during fracture and thus increases toughness. Although weld metal P2 is having comparable grain structure and comparatively stable austenite increases toughness up to certain extent and make it compatible with the spray mode even with the higher heat input. Therefore, it can be stated that pulse mode of metal transfer can provide better toughness values even with the high heat input condition.

Additionally, SEM fractographs in Fig. 12 show the size and distribution of dimples on the surface of broken charpy impact specimens. By comparing the fractographs in Figs. 12(a) and 12(b), it is observed that the P1 and P2 weld met-als have mainly ductile rupture with very few cleavages, although P1 shows a very fine dimple structure which is a clear indication of enhancement in toughness. Conversely, both weld metal S1 and P2 have extensive distribution of

Fig. 11. Impact toughness values of base metal and weld metals at room temperature.


© 2015 ISIJ1447

cleavage facets along with ductile rupture on the fracture surfaces which undoubtedly support the close proximity of the toughness values as illustrated in Figs. 12(b) and 12(c). The cleavage facets depicts the reduction in toughness up to a certain extent but does not deteriorates so much because of the certain ductile nature of the rupture which is incor-porated by the microstructural constituents.

4. Conclusion

From the present study, following conclusions can be drawn:• Pulse mode of metal transfer reduces the base metal

dilution and significantly alters the weld metal composition. Pulse mode also renders comparatively stable austenite in the weld structure with lower Creq/Nieq ratios and MT tem-peratures compare to spray mode.• Pulse mode significantly enhances the weld metal

grain structure by producing the grain size values analogous to spray mode even with the higher heat input condition. In general, high temperature heat affected zone of all the welded joints are characterized by coarse ferrite grains. However, pulse mode restricts grain coarsening of HTHAZ up to certain extent.• The weld metal is harder than HAZ followed by base

metal. Among the different weld metals, P1 shows relatively higher hardness values compare to P2 and S1. However, HAZ hardness remains comparable with each other.• The Charpy impact toughness of weld metal, in gen-

eral, is better than the base metal. The weld metal P1 exhib-its relatively higher impact toughness values than the other weld metals. However, close proximity in the toughness values of weld metal P2 and S1 clealy depicts enhancement due to pulse mode.• Finally, as a whole it can be concluded that the pulse

mode of metal transfer can produce better microstructural and mechanical properties compare to spray mode of metal

transfer in any heat input condition.

AcknowledgementsThe authors would like to gratefully thank the Council

of Scientific and Industrial Research (CSIR), Pusa, New Delhi, India for funding this research with a senior research fellowship.

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