Microstructures and mechanical properties in two X80 weld metals produced using similar heat input

Accepted Manuscript Title: Microstructures and Mechanical Properties in Two X80 Weld Metals Author: A.R.H. Midawi E.B.F. Santos N. Huda A.K. Sinha R. Lazor A.P. Gerlich PII: DOI: Reference:

S0924-0136(15)30070-4 http://dx.doi.org/doi:10.1016/j.jmatprotec.2015.07.019 PROTEC 14497

To appear in:

Journal of Materials Processing Technology

Received date: Revised date: Accepted date:

23-2-2015 20-7-2015 23-7-2015

Please cite this article as: Midawi, A.R.H., Santos, E.B.F., Huda, N., Sinha, A.K., Lazor, R., Gerlich, A.P., Microstructures and Mechanical Properties in Two X80 Weld Metals.Journal of Materials Processing Technology http://dx.doi.org/10.1016/j.jmatprotec.2015.07.019 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

Microstructures and Mechanical Properties in Two X80 Weld Metals Produced Using Similar Heat Input A.R.H. Midawia,1, E.B.F. Santosb , N. Hudaa, A. K. Sinhaa, R. Lazorc, A.P. Gerlicha a

Centre for Advanced Materials Joining (CAMJ), Department of Mechanical and Mechatronic Engineering, University of Waterloo, 200 University Avenue West, Waterloo, ON, Canada, N2L 3G1 b Laboratory of Materials Characterization (LCAM), Faculty of Mechanical Engineering (FEM), Federal University of Pará (UFPA), Rua Augusto Corrêa, 01 - Guamá, Belém-PA, CEP 66075110. c Materials Engineering Department, TransCanada PipeLines, Calgary, AB, Canada. 1

Corresponding author: Phone.: +1-226-700-5152, Email address: [email protected] (A.R.H. Midawi)

Abstract Two welding consumables suitable for joining X80 linepipe steel are compared in terms of weld metal microstructures, hardness, impact toughness, and tensile properties. The chemical compositions of the consumables were similar, where one of the consumables had a rich wire chemistry containing higher C, Ni, Ti alloying additions compared to a lean one. Beads on plate welding were performed using the gas metal arc welding (GMAW) process set to achieve the same heat input of 0.66 kJ/mm. The results revealed that for both wires the weld metal microstructure was mainly consist of acicular ferrite. The consumable with richer chemistry (C, Ni and Ti) exhibited higher strength and hardness due to its finer final weld metal microstructure; however, Charpy impact tests results revealed the lean chemistry wire had higher toughness at low temperature.

Since both weld metals exhibited similar acicular ferrite

structures, the lower toughness of the richer chemistry weld was attributed to the presence of titanium inclusions which may provide crack initiation sites.

Key words: GMAW consumables, X80 linepipe, tensile strength, impact toughness.

Highlights

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Two welding consumables, of different chemistry, suitable to join X80 pipeline steel were studied.

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The variation in chemical composition between the wires led to differences in the microstructure and mechanical properties.

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The weld metal with higher Ti and Ni content presented finer microstructure, higher strength and lower toughness.

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The Ti inclusions helped to form a finer acicular ferrite microstructure, which generally is beneficial for toughness.

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The large number and size of Ti inclusions acted as crack initiation sites, having detrimental effect on toughness.

1. Introduction

Steel pipelines are ubiquitous in the transport of oil and gas from the extraction sites to refineries. Due to the increasing demand for resources, there is a need for increased capacity, either by operating at higher service pressures, using larger pipe diameters or both. From the materials viewpoint, increasing the pipeline pressures implies in employing materials of higher strength. Currently, the majority of linepipe steel with minimum specified yield strengths above 550 MPa (ie: API X80 grade or higher) is based on high strength low alloy (HSLA) steels, which

consist of a fine grained ferrite or so-called bainitic ferrite that are produced using accelerated cooling and thermo-mechanically controlled processing (TMCP). One of the main requirements for linepipe joints is to achieve weld metal with equal or higher strength than the base material to avoid strain localization or fracture at the weld under loading. However, sufficient toughness is also required, which is typically verified using Charpy impact tests. A common solution is to design a weld metal with an acicular ferrite (AF) which provides a balance between strength and toughness (Babu, 2004). This has stimulated extensive research on the mechanisms of AF formation in weld metals, and determine which factors control its formation (Gretoft et al., 1986). A key factor in AF formation is the welding wire consumable composition, both in terms of the isolated effect of each element, and also the combined effect of the overall composition (Bhadeshia and Svensson, 1993). As described in a review on the formation of acicular ferrite in carbon-manganese weld deposits (Farrar and Harrison, 1987), the following elements have been reported to affect the formation of acicular ferrite: C, Mn, Si, Ni, Ti, Al, Mo and Nb. More recently, it was also reported (Avazkonandeh-Gharavol et al., 2009) that copper contributes to AF formation when employing manual metal arc welding. These elements will combine with oxygen present in the weld metal (Olson and Liu, 1986), which may be controlled by the shielding gas (Terashima and Bhadeshia, 2006) and/or the weld metal composition in (Harrison and Farrar, 1981). The reaction of the oxygen influences the formation of AF by either promoting or suppressing the formation of non-metallic inclusions, such as oxides. Some oxides act as nucleation sites for AF (Ilman et al., 2014), and so an increase in the oxygen content favour the formation of AF (Harrison and Farrar, 1981). For example, it was reported that

increasing oxygen content to 300 ppm changed the weld metal from Widmanstätten side plates an AF microstructure (Ilman et al., 2014). The formation of AF is also promoted by coarse austenite grains with a large number of inclusions of diameters larger than 0.2 µm (Olson and Liu, 1986). The details of AF formation have now been well described as a variation of the bainitic structure (Bhadeshia, 2001), and shown to be a particular variant which depends on intra-granular formation (Babu, 2004). As such, it is necessary to achieve a sufficient prior austenite grain size and number density of non-metallic inclusions of a favourable chemistry, especially those based on titanium oxides (Thewlis, 2004). However, it was also noted that if the amount of non-metallic inclusions reach certain levels depending on the oxygen content, there is a detrimental effect on toughness as crack initiation sites overcome the benefits of achieving a fine AF structure (Terashima and Bhadeshia, 2006) The main strengthening contribution in most HSLA steels is based on grain size refinement (Poorhaydari-Anaraki, 2005). This strengthening mechanism for microalloyed steels was systematically investigated by (Lu et al., 2012) and found to be well described by Hall-Petch strengthening behaviour (Hall, 1951), according to Eq. (1): (1) where σgs is the grain size contribution to yield strength in MPa; k’y is the strengthening coefficient in MPa mm1/2, and (m.l.i) is the mean linear intercept in mm. The value of k’y varies between 17.5-19.2 MPa mm1/2 for high strength low alloy steels as reported by (PoorhaydariAnaraki, 2005). Although titanium is known to be one of the principal elements promoting AF nucleation (Snyder and Pense, 1982); it has been reported in the literature that the optimum value of

titanium content in weld metal should be between 0.02 and 0.05 %wt (Beidokhti et al., 2009); with optimum manganese values around 1.27-wt% (Farrar, 1987). These values tend to promote nucleation of the AF structure with good toughness and tensile strength, however the combinations of elements also play a complex role on properties (Bose-Filho et al., 2007). Specific ratios between the content of Ti and Mn (Snyder and Pense, 1982), Ni and Mn (Farrar, 1987), Mn and Mo amounts (Biermann et al., 1988) also contribute to high toughness, while maintaining high strength. Based on existing literature, there is an optimum weld metal composition to provide high strength and toughness. The main goal of the present investigation is to compare microstructure and mechanical properties (hardness, tensile, and toughness) of two weld metals made using wires of similar chemical composition but mainly varying the Ni and Ti content, while the same heat input is used during welding. These two elements play a crucial role in determining the lower temperature toughness and AF nucleation respectively. Both wires investigated are commercial consumables for joining X80 pipeline steels and the subtle differences in microstructures are used to explain how mechanical performance varies between the weld metals. 2. Experimental Procedure

2.1 Material and Welding Procedure

Bead on plate welds were performed using gas metal arc welding (GMAW), with a robotic welder equipped connected to a Lincoln Power Wave 455M power supply, and a mixture of 85%Ar- 15%CO2 shielding gas. The base material used was a high strength steel ASTM A516-G70 and thickness of 12.5 mm (0.5′′). Single and multipass beads on plate were

performed. The preheating temperature was 120ºC for all welds, and for the multipass beads on plate welds the interpass temperature was 120ºC as well. The multipass welds were made with the purpose of examine toughness properties for both consumables. Two different welding consumables were used to perform the welds, designated Wire A and Wire B. The first, Wire A, has a higher content of Nickel (Ni) and Titanium (Ti). Table 1 indicates the nominal chemical composition of the base metal and Table 2 summarizes the chemical compositions of the weld beads produced by each wire consumable.

* The weight concentration of oxygen on the weld metal is being expressed in ppmw, parts per million by weight.

Both welding wires had a nominal wire diameter of 0.9 mm, and the welding parameters were set to achieve the same heat input for each. However, measurements of wires diameter showed that Wire A had a slightly higher actual wire diameter, 0.92, compared to 0.88 mm of wire B, although this is still within the specification for a 0.9 mm nominal diameter. This led to a 15% lower resistance for Wire A, which demanded slightly higher current when using the same wire feed speed. Table 3 shows the applied parameters for each wire, where the wire feed speed and travel speed were kept constant and the voltage was varied to provide the same heat input. The power supply was set to perform the welds using the standard constant voltage welding mode.

2.2 Metallographic Preparation and Microscopic Analysis

For microstructural characterization, the specimens for microscopic examination were cut from the middle of the cross-section of the weld bead, followed by grinding and polishing. The samples were etched with 2% Nital solution. To investigate differences in grain size produced when using the two different consumables, three areas near the center of the weld bead were selected for each microstructure and the average grain size was calculated, using for this the linear intercept method, the same procedure used by (Lu et al., 2012). The microstructural characterization was made using optical microscopy technique. For the examination of the fracture surface, scanning electron microscopy (SEM) was applied.

2.3 Mechanical Testing

2.3.1 Hardness Test

The Vickers micro hardness measurements were made in the cross section of the weld using an instrumented indentation machine microhardness tester (Nanovea, Model M1). For the single and multipass weld beads the indentations were performed along a straight line crossing the weld from top to bottom, reaching the heat affected zone of the base metal. The spacing between the indentations were 0.25 mm in the weld metal and 0.20 mm in the heat affected zone. The load used was 300g and a dwell time of 15s, and the average hardness values for the weld metal was calculated based on at least 10 repeat tests.

2.3.2 Tensile Test

In order to examine the tensile properties of both weld metals, tensile test specimens were machined from single pass bead-on-plate welds as shown in Fig. 1(a). The tensile test specimens had a gauge length of L0 = 12.7 mm with a cross-sectional area of 2 x 4 mm. The testing

procedure was performed according to ASTM E8M, using a Tinius-Olsen HK10T tensile test frame, with strain rate of 1mm/min.

2.3.3 Charpy Impact Toughness Test

In order to produce sufficient weld metal to compare Charpy impact toughness properties, multi-pass welds were produced using the same welding parameters presented in Table 3. For the Charpy impact toughness test V-notch Charpy impact toughness test specimens were extracted from the multi-pass weld metal. A sub-sized 2/3 Charpy impact specimen was used in this study (55×10×6.67) mm, the notch was 2 mm through 10 mm depth as shown in Fig. 1(b), with impact tests conducted at temperatures of 0, -20, and -40oC, with 3 specimens for each temperature. The Charpy impact machine used was a Satec Systems (Model SI-1K3) of 400 J capacity. Figure 1: Schematic diagram 1(a) tensile test coupons out of weld metal and Figure 1(b) Impact Charpy toughness samples.

3. Results and Discussion

3.1 Macrostructure

Fig. 2 compares the bead shape of the single bead produced for both consumables. The bead produced using both consumables presented similar reinforcement, however the beads produced by using wire A were narrower and deeper when compared to the ones produced using wire B. The higher dilution produced when using wire A is consistent with the higher current values observed and presented in Table 3. The columnar macroscopic structure in the weld metal can be observed for all conditions.

Figure 2: Macrographs for different welding condition, indicating dilution rates and bead dimensions (a) Wire A-19V, 0.66 kJ/mm, 30% Dilution; (b) Wire B-22V-0.67 kJ/mm-21.8% Dilution. 3.2. Microstructure

For both wires, the microstructure predominantly contains an acicular ferrite (AF) interlocking structure with a small fraction of grain-boundary ferrite (GF) and polygonal ferrite (PF), which may also be considered quasi-polygonal ferrite as others describe (Thewlis, 2004). A close observation on the weld microstructure revealed that the weld metal for Wire B exhibited a coarser microstructure than that produced when using wire A. The grain size measurement results confirmed that the weld metal produced using wire A, which has a higher Titanium content, has an average grain size of 1.86 µm, while a value of 2.36 µm was observed for the Wire B weld metal. This is in accordance with what was reported by (Harrison and Farrar, 1981) when investigating the influence of reducing the oxygen content of two submerged-arc high-strength low-alloy (HSLA) steel weld metals, where they suggest that titanium inclusions mainly promote the nucleation of acicular ferrite in the interior of prior austenite grains. This is supported by the refined microstructure observed when using Wire A, since titanium. Figure 3: Optical and SEM micrographs of single bead weld metal microstructures produced using both welding wires with a heat input of 0.66 kJ/mm. (a) Wire A, 19V, 0.66kJ/mm, Top of weld; (b) Wire B, 22V, 0.67kJ/mm, Top of weld; (c) Wire A, 19V, 0.66kJ/mm, Middle of weld; (d) Wire B, 22V, 0.67kJ/mm, Middle of weld; (e) SEM from Wire A, 19V, 0.66kJ/mm, Middle of weld; (f) SEM from Wire B, 22V, 0.67kJ/mm, Middle of weld. As shown in Figure 4, the microstructure in the multipass reheated zones of weld metals produced using both wires contained a slightly higher fraction of GF or PF ferrite morphologies

compared to single pass welding, agreeing to what was reported by (Harrison and Farrar, 1981); however, AF remains the dominant constituent for both weld metals. The slightly higher AF content observed when using Wire A can also be attributed to the combination of the higher titanium content and higher oxygen content which could promote the formation of titaniumbased oxides that help to nucleate fine AF.

Figure 4: Microstructures of reheated zones in multipass welds for the two selected conditions: (a) Wire A and (b) Wire B. 3.4. Hardness behaviour Microhardness profiles for both wires are compared in Fig. 5. As shown in Fig. 5, for the weld metal of wire A the microhardness was around 300.4±7.5HV versus 280.4±4.9HV for wire B when a same heat input (~ 0.66 kJ/mm) was applied in a single pass weld bead. It can be observed that the Wire A produces consistently higher hardness weld metal compared to Wire B. The higher hardness can be attributed to a finer grain size observed in the Wire A weld metal. In addition, a higher Ti content with combination with higher carbon, nitrogen, and oxygen in the weld metal of Wire A may also promote formation of TiN, TiC, or Ti(C,N) precipitates (Beidokhti et al., 2014), which would provide further strengthening. The presence of titaniumrich particles in the weld metal of Wire A was confirmed by EDX analysis as shown in Fig. 12. The hardness value at the heat affected zone right adjacent to the fusion line, is an average value of five measurements 100 µm away from fusion line. The HAZ hardness immediately adjacent to the fusion line exhibits a higher hardness compared to the deposited weld metal using both consumables, indicating similar values and trend which may be expected when identical heat input is used. Figure 5: Hardness profile through thickness for as deposited condition.

Fig. 6 presents the hardness profile for the multipass welding beads, reheated condition, and again the weld deposits performed using the Wire A presented higher hardness of 290.5±13.9 HV throughout the thickness. However the weld metal deposited using the Wire B, presented lower values around 251.9±20.3 HV for the first passes deposited, bottom. This difference in hardness profile can be attributed to its coarser grain size and higher volume fraction of PF in the weld zone as noted previously in Figure 4. In addition, the hardenability of weld metal derived from Wire A is higher, based on the carbon equivalent indicated in Table 3, which is 0.41 for wire A compared to 0.35 for wire B. Figure 6: Hardness profile through thickness for multipass samples.

3.5. Tensile tests

Fig. 7 presents the stress-strain curve for both wires in the as deposited condition for single pass weld beads and Table 4 summarizes the tensile properties of these wires. According to (American Petroleum Institute, 2000) the yield strength for both weld deposits overmatch the X80 pipeline steel minimum requirement of 550MPa. An important parameter for pipeline design, considering the strain based approach, is the relationship between yield strength and tensile strength (YS/UTS), and Table 4 also shows that for the wires studied herein are in the range of 0.8 to 0.9, which is acceptable values for pipeline applications according to reported by (Shinohara et al., 2007) and (Mohr, 2003). Another important parameter is the uniform elongation, which for the present weld metals were 5.8% for Wire A compared to 3.8% for the Wire B weld metal.

Figure 7: Stress-strain curves for the as-deposited weld metals produced using similar heat input.

When using identical heat input, the strength for the weld metal of Wire A is higher than that of Wire B by 10%, and consequently differences in the microstructure resulting from the chemical composition (and resulting difference in phase transformations during cooling) can account for the different behaviour under tensile test. As showed in Table 2 the elements that vary are mainly Ni, Mo, Ti, and C. The influence of grain size refinement on the yield strength for both wire was quantified by the Hall-Petch equation, Eq. (1), and its contribution is summarized in Table 5. The values for the constant k’y varied between 17.5 and 19.2 MPa mm1/2 as reported by (Poorhaydari-Anaraki, 2005). From the results in Table 5 it can be shown that most of the yield strength of the weld metals are contributed by grain size refinement, which is greater for the finer grained weld metal deposited using wire A. The percentage contribution found in the present work is in agreement to what was found by (Poorhaydari-Anaraki, 2005) and (Lu et al., 2012), who reported 66% and 70% respectively.

3.6. Toughness Results The toughness results are presented in Fig. 8 and Table 6, where it can be noted that for the temperature range investigated for the present work, wire B is tougher than wire A. Table 6 presents the difference in toughness values at the temperatures investigated. As the test temperature decreases, the differences in absorbed energy for each weld metal decreases.

Figure 8: Charpy impact toughness results for wire A and wire B weld metal at temperature range of 0 to -40°C.

The V-notch Charpy tests results, Fig. 8, indicate that the weld metal of Wire B exhibits higher upper shelf impact toughness than that of Wire A, and based on the values observed to testing temperatures of -40ºC, it is expected that the ductile to brittle transition temperature is below this value. The higher upper shelf values produced using Wire B are in agreement with the lower hardness results observed in this study, Fig.6. From Table 2, it can be seen that Wire A has 0.795-wt% nickel compared to 0.030-wt% nickel in Wire B. It was reported by (Taylor and Evans, 1983) that the addition of Ni can flatten the Charpy impact transition curve and increase lower shelf toughness, at the same time that the upper shelf toughness will be decreased for weld metal deposits as the Ni content increases. This behaviour was observed in Wire A, and justifies due to high Ni content compared to Wire B. Moreover, according to (Thuvander, 2003) increasing the summation of Mo and Ni has the effect of harden the weld metal and decrease impact toughness. Since the combination of Ni and Mo for both weld metals are rather different (with values of Ni = 0.795wt%, and Mo = 0.292wt% for Wire A, versus Ni = 0.030wt% and Mo = 0.345wt % for Wire B), it appears the reduction in Ni in Wire B has helped to avoid a detriment to toughness. One can also note from Fig. 8 and Table 6 that as the test temperature decreases, the decrease in absorbed energy is less drastic for wire A. This can be explained by the presence of Ni in wire A, since Ni has been shown to increases the cohesive strength of the ferrite lattice, resulting in higher impact toughness values at low temperatures, and decreasing the transition temperature (Leslie, 1972). The nickel content present in Wire A accounts for more gradual decline in impact toughness for Wire A.

3.7. Fractography and Analysis of Inclusions

Further SEM examination of the fracture surfaces was conducted close to the notch area and at the middle area of the fracture surfaces, as shown in Fig. 9, and both wires presented ductile fracture surfaces characterized by the present of dimples. The fracture features close to the notch area exhibits an elongate dimple structure, for both wires, however the strain seems to be higher in the case of Wire B, Fig. 9(c). When one compares these features, along with the decrease in the hardness and lower carbon equivalent value for Wire B, it can explain the higher impact toughness values for this weld metal. The SEM fractography analysis revealed the presence of inclusions, which appear to be oxides, at the center the dimples, Figs. 10 and 11. Most of the inclusions observed in the fracture surface for both wires have size less than 1 µm, which is an ideal size to nucleate AF according to (Prokić-Cvetković et al., 2006), and thus explains the AF structure shown in Fig. 4. The number of inclusions visible on the fracture surface was measured based on an average of five images, each of them having an area of 100 µm2. The number of particles present on the Wire A weld metal was higher, being 3.1 ± 1.8 particles, when compared to 2 ± 0.6 particles in the weld metal of Wire B. Fig. 9. SEM images of Charpy fracture surfaces (at -20°C) for Wire A and B weld metals produced using 0.66 and 0.67 kJ/mm respectively: 11(a) Wire A near the notch; 11(b) Wire A towards middle; 11(c) Wire B near the notch; 11(d) Wire B towards middle. In order to investigate the nature of the inclusions, EDX analysis was performed. The results indicated that the inclusions present on Wire A were mainly titanium and oxygen based, Fig. 10, while in Wire B they were silicon and manganese rich inclusions, as shown in Fig. 11. The Silicon, manganese based inclusions observed on the fracture surface of Wire B, shown in Fig. 11(b), are most likely correspond to (Mn,Si)O inclusions, which are commonly observed as a nucleation point for AF in steels free of titanium (Beidokhti et al., 2009), and this would explain microstructures observed in Figs. 3(b), 3(d), 3(f) and 4(b).

Figure 11: SEM images of the Charpy fracture surfaces showing inclusions at the center of dimples (black arrows) for the weld metal deposited using Wire B; (b) EDX spectrum showing that the inclusions are mainly formed of silicon (Si), manganese (Mn) and oxygen (O).

The titanium inclusions in the wire A weld metal, which are likely to be TiO/TiO2, would promote more refined microstructure by both suppressing the coarsening of austenite grains, agreeing to the stated by (Harrison and Farrar, 1981) for the reheated condition weld metal, and providing more nucleation points for acicular ferrite, which presents good toughness properties, corroborating with the reposted by (Babu, 2004), (Sung et al., 2014). Beidokhti investigated the influence of titanium and manganese on properties of HSLA weld metal using the SAW process (Beidokhti et al., 2009) even reported an optimum value of titanium content to be between (0.02– 0.05 %wt), which provides a good combination of strength and toughness. Since a higher oxygen content was present in the weld made using Wire A (180 versus 110 ppm), this provided a higher density of inclusions and microvoid nucleation sites on the fracture surface. This is also consistent prior work (Abson and Pargeter, 1986) showing a higher frequency of fine inclusions will provide sites for micro-void coalescence during ductile fracture, which may decrease upper shelf fracture toughness.

4. Conclusion

The microstructure and mechanical properties (hardness, tensile strength and impact toughness) for two weld consumables, of different chemical composition, suitable to join X80 pipeline steel was studied. Joints were made using GMAW with the same heat input using both wires, and it was found that:

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Both as-deposited weld metals exhibited an AF microstructure, however the higher titanium, nickel and oxygen content in Wire A provided a finer microstructure.

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For both the as deposited and reheated weld metal, the hardness of the weld metal deposited using Wire A was higher, and attributed to the grain refinement in accordance with the Hall-Petch equation.

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The weld metal produced using Wire B, which was leaner in Ti, Ni and O content, which resulted in higher toughness due to a lower density of oxide inclusions on the fibrous fracture surface observed by SEM and EDX analyses.

These results were observed even when the same heat input was applied for the two weld metals, and the slight differences in chemical compositions resulted in the differences in microstructure and mechanical properties.

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Table 1: Nominal chemical composition of the base metal (A516-Gr70), in wt%. Table 2: Chemical compositions of deposited weld metals, in wt%. Table 3: Processing parameters for each wire and heat input values. Table 4: Tensile properties of weld metals investigated using a heat input of 0.66 kJ/mm. Table 5: Grain size effect on the yield strength for each consumable. Table 6. Charpy impact properties tested at temperature of 0, -20 and -40°C for both weld metal. Figure Caption

Figure 1: Schematic diagram (a) tensile test coupons out of weld metal and Figure (b) impact Charpy toughness samples. Figure 2: Macrographs for different welding condition, indicating dilution rates and bead dimensions (a) Wire A-19V, 0.66 kJ/mm, 30% Dilution and (b) Wire B-22V-0.67 kJ/mm-21.8% Dilution.

Figure 3: Optical and SEM micrographs of single bead weld metal microstructures produced using both welding wires with a heat input of 0.66 kJ/mm. (a) Wire A, 19V, 0.66kJ/mm, Top of weld; (b) Wire B, 22V, 0.67kJ/mm, Top of weld; (c) Wire A, 19V, 0.66kJ/mm, Middle of weld; (d) Wire B, 22V, 0.67kJ/mm, Middle of weld; (e) SEM from Wire A, 19V, 0.66kJ/mm, Middle of weld; (f) SEM from Wire B, 22V, 0.67kJ/mm, Middle of weld. Figure 4: Microstructures of reheated zones in multipass welds for the two selected conditions: (a) Wire A and (b) Wire B. Figure 5: Hardness profile through thickness for as deposited condition. Figure 6: Hardness profile through thickness for multipass samples. Figure 7: Stress-strain curves for the as-deposited weld metals produced using similar heat input. Figure 8: Charpy impact toughness results for wire A and wire B weld metal at temperature range of 0 to -40°C. Fig. 9. SEM images of Charpy fracture surfaces (at -20°C) for Wire A and B weld metals produced using 0.66 and 0.67 kJ/mm respectively: 11(a) Wire A near the notch; 11(b) Wire A towards middle; 11(c) Wire B near the notch; 11(d) Wire B towards middle. Figure 10: (a) SEM images of the Charpy fracture surfaces showing inclusions at the center of dimples for the weld metal deposited using Wire A. (b) EDX spectrum showing that the inclusions are mainly formed of titanium (Ti) and oxygen (O). Figure 11: SEM images of the Charpy fracture surfaces showing inclusions at the center of dimples for the weld metal deposited using Wire B. (b) EDX spectrum showing that the inclusions are mainly formed of silicon (Si), manganese (Mn) and oxygen (O).

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