Welding of Different Metals Across Ferritic Substrates: A Thermal Modeling

Steels with increased creep strength, such as 9 wt% chromium Grade 91 and 2.25 wt% chromium Grade 23, have been made possible by the ever-increasing efficiency of today's fossil fuel power plants. These materials' extraordinary creep capabilities result from the interplay of several various types of strengthening processes, including dispersion hardening, solution hardening, and subgrain boundary hardening. When micro-alloying elements like vanadium, niobium, carbon, and nitrogen are added to a material, they generate carbides and carbonitrides that improve the bainitic microstructure [1]. However, the strength of Grade 91 materials comes from their precipitation-strengthened martensitic microstructure. Dissimilar materials must be transitioned over short distances in today's power plants to satisfy a wide range of service conditions. To back up best practise fabrication, estimate remaining life, and back up future inspection decisions, knowing how dissimilar metal welds (DMW) behave under creep conditions is essential. At high temperatures in service, a phenomena known as uphill-diffusion of carbon from the lower alloyed chromium steel to the one with greater chromium content can be observed [2–5]. This is caused by the rapid transition between materials with extremely different chemical compositions. Because of the differing chemical activities of carbon in the two materials employed in the DMW, carbon will migrate from one to the other, permanently altering the microstructure close to the fusion line. Carbon enrichment on the more highly alloyed side causes carbides to precipitate near the fusion line in the carbon enriched zone (CEZ), which depletes the surrounding matrix of carbide forming elements like chromium and consequently changes the properties of this area relative to those of the base metal matrix. In contrast, carbon is depleted from the microstructure on the less alloyed side. In order to provide the carbon migration with atoms, a carbon-depleted zone (CDZ) is created where all the carbides are dissolved. However, at higher temperatures, the microstructure tends to recrystallize to a weaker ferritic crystal structure, as the carbides in the lower alloyed material no longer provide sufficient support for the martensitic/bainitic microstructure. These structural changes are localised to an area less than 1 mm across the fusion line, and they lead to a change in mechanical characteristics. As a result, strain localization can become more pronounced, making the component more vulnerable to damage under real-world conditions [6, 7].

Computing simulations were carried out using MatCalc Version 5.62 [8]. To learn about the precipitation habits of the materials, equilibrium phase fraction calculations were performed. Carbon activity calculations were also performed since carbon activ- ity variations in the material constitute the driving factor for carbon migration in DMWs [9]. In addition, a 1-dimensional finite ele- ment mesh with 100 elements was used to perform the diffusion calculations. Please note that the mc fe 2.058.tdb and mc fe 2.010.ddb MatCalc calculation. This paper's calculations are based on those of a different DMW [10]. Metallographic analysis of creep-tested cross-weld samples of T91 welded with T23 weld metal confirmed the accuracy of the calculations. For example, the most noticeable microstructure changes occur at the fusion line region between the T23 weld metal and the T91 base material when it is exposed to creep. There were Grade 23 weld metal electrodes and a Grade 91 base material, and their alloying compositions are summarised in Table 1. Using three different Grade 23 type weld electrodes ("E9015-G") of varying electrode diameters, the welds were created by hand using manual shielded metal arc welding (SMAW) (2.5 mm, 3.2 mm, 4 mm). Sequence diagrams of the weld layers are depicted in Fig. 1. The thickness of the plate was 48 mm, the preheat temperature was 150 C, and the interpass temperature was 315 C. Table 2 provides a summary of the observed welding parameters; the "Pass"-column numbers correlate to the pass numbers shown in the layer sequence diagrams in Figure 1. Samples were PWHT at 745°C for 2 hours after welding. After being subjected to welding and PWHT, the samples were machined so that creep testing could take place. Sample geometries for creep testing at 600 and 625 degrees Celsius and 80 MPa are shown in Figure 2.

It is feasible to analyse equilibrium conditions to explain pre-cipitation behaviour during cooling of the base metals with only the chemical composition of the applied material. To do thermodynamic calculations, MatCalc employs the CALPHAD approach, which permits the user to add flexible databases for various alloying systems. Because of this, we can make predictions about material behaviour in settings when no experimental data exists. A large number of simulations can be computed if the chemical composition of the system is known. Precipitation patterns in the materials were estimated by simulating the equilibrium phase

A wide range of precipitates was considered in the analyses. After reviewing the relevant literature [2, 5, 11–16], we learn that M23C6 carbides and MX carbonitrides are the most abundant precipitate species in the tested substances. Different studies have discovered both M6C and M7C3 carbides in low- and high-chromium-alloyed steels. Figure 3 shows a clear link between M7C3 and M23C6 precipitation. The simulation results indicate that M7C3 precipitation occurs during cooling after austenite transforms into ferrite. At temperatures of about 510 °C, M7C3 precipitates are being dissolved in favour of M23C6 carbides. The simulation findings demonstrate an earlier commencement of M23C6 precipitation at 780 °C if M7C3 carbides are not included in the calculations (Fig. 4). This illustration should make it clear that accurate material behaviour simulation requires knowledge of precipitates present in the materials. The following carbon migration calculations are based on the carbon activity calculations performed on Grade 91 base metal and Grade 23 weld metal, respectively. Carbon activity findings for the sample materials are displayed in Figure 5. Compared to Grade 91 base material, Grade 23 weld metal electrodes have much lower carbon activity by more than an order of magnitude. What's fascinating is that despite the fact that the weld metal electrodes share a virtually identical chemical makeup, they have noticeably different carbon activity. Considering the disparities in the weld metal electrodes' chemical makeup, it's intriguing to speculate on the feasibility of bridging the carbon activity gap by adjusting the electrodes' chemical makeups within the bounds of the materials' stipulated elements (see Table 3) [17, 18]. A reduction in the driving force for carbon migration will occur if the gap in carbon activity can be closed or at least substantially narrowed by altering the concentration of the alloying elements. There needs to be an initial assessment of how various factors affect car- bon activity. One way to achieve this is to change the content of just one element inside its defined borders while leaving the rest of the components unchanged. Carbon activity versus ele- ment content plots reveal how different elements affect carbon activity in a given substance (compare Fig. 6). of Grade 23 material. The similar approach can be used to increase the carbon activity of Grade 91 material. What we get after the procedure is depicted in Fig. 7. The discrepancy between Grade 91 and Grade 23 materials' carbon activities cannot be eliminated entirely. At 550 °C, for instance, T23/T91 has a ratio of 51.25 (large gap) while T23v/T91v (materials with a tailored composition that is aimed to reduce the carbon activity difference) has a ratio of 3.33 (small gap). This reduction in the carbon activity ratio by a factor of 15 should significantly lessen the driving force for the carbon migration.

To verify the results achieved via simulations, microstruc- tural investigations in the area of the fusion line with a light-optical microscope and SEM examination were carried out. Directed carbon migration from the lower alloyed weld metal to the higher alloyed base material caused negative microstructural changes as a result of the applied stress and the elevated temperature. As previously indicated, the carbon depletion leads to the breakdown of microstructure-stabilizing carbides in the Grade 23 weld metal which results in undesirable recrystallized ferrite in the carbon-depleted zone (CDZ) (CDZ). Furthermore, photos obtained using a scanning electron microscope support these results .This photo was captured by the SEM's secondary electron (SE) detector. On the T91 side, little finely scattered carbides may be recognised, while on the T23 side, very coarse in the lower alloyed weld metal, the carbon depletion weakens the microstructure and makes it softer. Based on EDS examination, the coarse white particles found on the Grade 23 weld metal side are Laves phase, which is rich in tungsten. Weld metal microstructure has only undergone partial recrystallization, as shown by microstructural investigation, hence ferrite islands are only seen at the fusing point. Reasons for this include the use of micro-alloying to increase strength, and the minimal amount of time required for testing (1000 to 10,000 h). Table 4 is a summary of the information gathered about ferritic recrystallization. It is interesting to note that the diffusion and, ultimately, the recrystallization behaviour of the weld metal appear to be affected by the state of the Grade 91 base material. In contrast to the "B" samples, which were made using Grade 91 material that had been renormalized and tempered before welding, the "A" samples were made with untreated ex-service material.

After creep exposure, the renormalized and tempered samples clearly exhibit higher recrystallization than the untreated samples, despite the fact that the testing time for both groups was shorter. The normalising process is responsible for this improvement by decreasing the microstructure's dislocation density. It is possible to expedite the recrystallization process by increasing the temperature of the sample during the creep exposure so that more dislo- cations are formed and accumulated.

The influence of PWHT and creep exposure on the welds was studied by measuring the joints' hardness. Renormalized and tempered Grade 91 base material appears to be slightly tougher than the ex- service ("A") sample, and this trend continues in the as-welded condition. There is not much of a distinction in the weld metal. Weld metal hardness decreases after PWHT, by 113 HV10 for "A" samples and 128 HV10 for "B" samples. There was a greater reaction to PWHT in the renormalized and temperature-treated base metal specimen, with a fall in hardness of 38 HV10, compared to the "A"- sample, which only showed a decrease in hardness of 22 HV10. Micro- hardness line scanning (HV0.05) across the fusion line of sample 7A-2 yielded the image in Fig. 11. There is a clear decrease in hardness on the weld metal side, and its breadth of around 100 m matches well with the width of the recrystallized microstructure. Finely distributed carbides precipitated across the fusing line as a result of carbon migration, creating a zone of increased hardness known as the carbon enriched zone (CEZ). The findings of the microstructural investigation and the outcomes of the computational simulations are corroborated by this line scan.

The purpose of this research was to illustrate the potential of using thermodynamic simulation software to predict the behaviour of creep-exposed dissimilar metal welds. Precipitation sequences during equilibrium cooling can be predicted by basing thermodynamic calculations on the chemical composition. These findings are useful inputs for more involved calculations like 1-dimensional diffusion simulations where the simulation parameters are set to mimic creep experiments or service conditions. Under conditions resembling those of real-world creep exposure tests, carbon migration calculations revealed the magnitude of carbon depletion in the lower alloyed material and carbon enrichment in the higher alloyed material. The extent of the difference between high- and low-chromium steels was calculated using carbon activity calculations. In addition, a concrete illustration of the usefulness of element activity calculations in the context of alloying optimization was provided. The unfavourable microstructural changes that these welded joints undergo under service conditions were revealed by microstructural studies of creep-exposed T91 to T23 DMWs. The dissolution of carbides in the carbon-depleted zone resulted in a partial recrystallization of the weld metal microstructure. In addition, on the side of the weld that contains the more highly alloyed Grade 91 base metal, finely distributed carbides were discovered to correlate to a dramatic increase in hardness right next to the fusion line. Good agreement was found between the computational simulations and the creep experiments after an analysis and comparison with the data accumulated with 1-dimensional diffusion simulations. It was also demonstrated that familiarity with the materials and their precise microstructural features, specifically their inherent precipitate phases, is necessary for improved results from the computational simulations.

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Research Scholar, Sunrise University, Alwar, Rajasthan