Physical simulation-based analysis of multipass welding in S500 shipbuilding steel

doi:10.21203/rs.3.rs-4983691/v1

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Physical simulation-based analysis of multipass welding in S500 shipbuilding steel

https://doi.org/10.21203/rs.3.rs-4983691/v1

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Novel generations of shipbuilding steels have outstanding toughness due to the improved steel producing processes. Their microstructure mainly consists of ferrite and bainite, whilst the presence of acicular ferrite has a role in high impact energy of the welded joint. This research aims to analyze the effect of multipass welding on weld characteristics of S500 shipbuilding steel. A Gleeble 3500 simulator machine is used to produce the welding thermal cycles by the Rykalin-3D model on 70 x 10 x 10 mm samples manufactured in transversal direction from a submerged arc welded joint of 16 mm plate. For the simulation of the coarse-grained zone forming in the weld metal (CGHAZ-W) 1350°C and for the intercritical zone (ICHAZ-W) 815°C were selected, while the combination of these peak temperatures was applied for the intercritically reheated coarse-grained zone (ICCGHAZ-W). The examined t_8/5 interval was 5…30 s. The weld properties were examined by microstructural examination, hardness test, and instrumented Charpy V-notch impact toughness test. The impact energy values of subzones were below the unaffected weld metal. Longer cooling time resulted in lower impact energy in ICHAZ-W. However, this tendency was not observed in CGHAZ-W. ICHAZ-W and ICCGHAZ-W resulted the lowest impact toughness, which was indicated by the large instable crack propagation.

S500 shipbuilding steel

physical simulation

welding

instrumented impact test

microstructural characterization

Modern low-carbon thermomechanically treated steels are popularly used in shipbuilding and marine structures (oil rigs, foundations for wind turbines), where the components of the product must withstand extreme environmental load and the resulting increased mechanical stress [1]. These structures are often subjected to dynamic loading at negative temperatures. Therefore, the steel producers strive to provide these steels with excellent toughness properties, against the regulations imposed quite strict requirements. Depending on the steel grades offshore steels are produced by normalizing (N), thermomechanically controlled process (TMCP), quenching and tempering (Q + T). For the TMCP steels in the medium strength range (R_eH=400…500 MPa) are characterized by low carbon content and low carbon equivalent (CE) with the mixed microstructure of ferrite and bainite. Offshore steels with the highest toughness are typically manufactured by TMCP that leads to an extremely fine-grained microstructure. The presence of acicular ferrite (AF) forming across the small precipitates and inclusions has a determining role in achieving the demanded toughness [2][3]. In welded joints ferrite plates growing in many radial directions on small-sized non-metallic inclusions, as crystallization nuclei, can contribute to the improvement of toughness properties due to their heterogeneous structure [2]. However, the presence of M-A constituents can reduce the impact energy especially in the heat-affected zone [4]. Due to the favorable chemical composition, these steels have good weldability which is verified by its location in the first zone of Graville diagram [5]. Cold cracking is not expected, however, they may initiate in larger cross sections from the existing hot cracks at the end crater [6]. Deterioration of toughness in the HAZ for example is a known problem with offshore steels [3][7]. Furthermore, it is a challenge to guarantee that the impact energy level in the dendritic, mostly multi-pass weld, is the same as that in the fine-grained and micro-alloyed steel plate, which is rolled through specific rolling processes. The applied standards [8] have certain acceptance criteria for the toughness properties of the HAZ and the weld, which are typically required to be tested with Charpy V-notch (CVN) impact tests and with fracture toughness tests, such as crack-tip opening displacement (CTOD). Nowadays, numerous research deal with the role of AF in the heat-affected zone (HAZ), however, there is less focus on the weld microstructure and mechanical properties.

Tervo et al. [3] investigated low-temperature toughness properties of 500 MPa offshore steels and their simulated coarse-grained heat-affected zones (CGHAZ). It was found that both impact and fracture toughness of an S500 offshore steel with coarse titanium-based nitrides were lower and the scatter higher than in the steel with a higher quantity of calcium-based inclusions and without the coarse nitrides. In [7] three experimental steels were studied in order to find optimal conditions for the AF formation in the CGHAZ. One of the steels was Al-deoxidized, while the other two were Ti-deoxidized. The focus was to distinguish whether the deoxidation practice affected the AF formation in the simulated CGHAZ. It was found that AF formed in the simulated CGHAZ of one of the Ti-deoxidized steels and its fraction increased with increasing cooling time. In this steel, the inclusions consisted mainly of small (1–4 µm) TiOx-MnS, and the tendency for prior austenite grain coarsening was the highest. In another [9] study, the effects of varying fractions of AF (0–49 vol.%) were assessed in the simulated, unaltered CGHAZ of three experimental steels. Two steels were deoxidized using Ti and one using Al. The characterization was carried out by using electron microscopy, energy-dispersive X-ray spectrometry, electron backscatter diffraction and X-ray diffraction. The fraction of AF varied with the heat input and cooling time applied in the Gleeble thermomechanical simulator. AF was present in one of the Ti-deoxidized steels with all the applied cooling times, and its fraction increased with increasing cooling time. However, in other materials, only a small fraction (13–22%) of AF was present and only when the longest cooling time was applied. The impact toughness of the simulated specimens was evaluated using instrumented Charpy V-notch testing. Contrary to the assumption, the highest impact toughness was obtained in the conventional Al-deoxidized steel with little or no AF in the microstructure, while the variants with the highest fraction of AF had the lowest impact toughness. It was concluded that the decreased impact toughness in the CGHAZ of steels with AF was due to the coarser microstructural and inclusion features, and also the fraction of AF may not have been great enough to improve the CGHAZ toughness of the steels investigated. The effects of titanium content on the weld microstructure, mechanical properties, and inclusion characteristics were investigated by Seo et al [10] in the as-deposited bainitic gas metal arc weld metals having a nearly constant level of oxygen content. It was found that titanium addition enhanced the formation of AF with the maximum proportion being obtained at ∼0.07 wt.% Ti. The resultant change in weld microstructure with titanium content was well reflected in the Charpy impact energy showing the lowest ductile-brittle transition temperature at 0.07 wt.% Ti.

The effect of multipass welding on the weld metal has been studied previously e.g. by Kang et al. [11]. They utilized physical simulation to study the microstructures and mechanical properties in reheated zones in two different weld metals. They concluded that the toughness reduction in reheated zones of low hardenability weld metal was attributed to the increase of the grain boundary ferrite in the microstructure, whereas in the high hardenability weld metal the reason for the reduced toughness in the reheated zones was the coalesced bainite.

The goal of present research is to investigate the properties of multipass welded joint of a 500 MPa grade shipbuilding steel by physical simulation in the industrially relevant t_8/5 cooling time range of arc welding processes. The experimental program is aimed at the production and analysis of physically simulated HAZ subzones (coarse-grained, intercritical, intercritically reheated coarse-grained) occurring in the weld, which can be considered critical zones in terms of toughness, by microstructural characterization, hardness test and instrumented Charpy V-notch impact toughness test.

1.1. Base and filler materials

The chemical composition of the 500 MPa grade (S500ML) shipbuilding steel, classified by EN 10025-4 [12], the selected S3Ni1Mo0.2 (ESAB OK 13.24) according to EN ISO 14171 [13] filler metal for submerged arc welding (SAW) and the weld are summarized in Table 1. In case of the base material which is under development the publication of the exact chemical composition was not allowed. The copper coated wire electrode is alloyed by nickel and molybdenum with a diameter of 4 mm which is qualified by CTOD (Crack Tip Opening Displacement) tests. The recommended flux for the given filler material is ESAB OK Flux 10.62, which is a high basicity, neutral, bonded flux intended primarily for multipass butt welding of carbon and low alloy steel plates. It produces weld metal that is very clean metallurgically and exhibits exceptional impact toughness at low temperatures. This filler material and flux combination is directly developed for fine grained steels in offshore industry and shipbuilding.

Table 1

Chemical compositions of the base, filler material and the weld
Material	C	Si	Mn	P	S	Cr	Mo	Ni	Nb	V	Ti	Cu	Al	B	CEV	CET
Base material	≤ 0.14	≤ 0.6	≤ 1.7	≤ 0.02	≤ 0.01	Cr + Mo ≤ 0.65		≤ 2.0	Nb + V + Ti ≤ 0.26			≤ 0.55	≥ 0.02	N/A	N/A	N/A
Filler material	0.067	0.184	1.33	0.015	0.0078	0.053	0.187	0.784	0.003	0.003	0.003	0.058	0.014	0.002	0.49	0.27
Weld*	0.053	0.366	1.46	0.011	0.0058	0.061	0.201	0.893	0.024	0.006	0.007	0.202	0.022	0.0002	0.42	0.25
* Optical emission spectrometer (OES) analysis

The guaranteed mechanical properties of the base and filler material are summarized in Table 2.

Table 2

Mechanical properties of the base and filler material
Material	R_eH, MPa	R_m, MPa	A₅, %	CVN, J
Base material	500	580–760	15	20 J [-40°C]
Filler material	530	620	N/A	N/A

1.2. Experimental circumstances

In contrary to the conventional application [14][15] of Gleeble physical simulators for the HAZ tests, the goal of present experiments was the simulation-based analysis of multipass welding heat cycles on weld properties. For this purpose, welded joints were prepared from 16 mm thick plates by submerged arc welding (SAW). The reason of SAW is to produce a relatively homogenous weld in one layer with this technology, which is sufficient in volume for the production of physical simulation specimens. According to Fig. 1 symmetric one side joint was prepared with Y-groove, the edge gap was 3 mm, the depth of root face was 4 mm, the angle of bevel was 40°. In terms of the applied welding parameters the welding current was 751 A, the welding voltage was 30.8 V, the wire feeding speed was 1.75 m/min, the welding speed was 27.9 cm/min, the welding heat input was 5 kJ/mm. The physical simulation specimens were machined transversely from the welded joint in such a way that the seam falls in the middle part of the specimens. The specimen size was 70\(\:\times\:\)10\(\:\times\:\)10 mm. The physical simulation experiments were performed on Gleeble 3500 thermophysical simulator. By the equipment the desired subzones of welded joint can be precisely and homogeneously produced in a sufficient volume for the further material tests.

The physical simulation experiments were aimed at producing the most unfavorable microstructure occurring in the multilayer weld in terms of the toughness properties. Therefore, 1350°C was selected for the simulation of coarse-grained heat-affected zone forming in the weld (CGHAZ-W) and 815°C peak temperature was used for the intercritical heat-affected zone (ICHAZ-W). In case of CGHAZ the motivation was to produce largest possible grains in the heat-affected zone. In the case of the ICHAZ, when choosing the temperature, we assumed that the estimated value of the A_c1 temperature for steels with a low carbon equivalent is 765°C, and for the intercritical simulation, based on practical experience, it is worth choosing A_c1+50 = 815°C. The combination of the two thermal cycles were applied for the simulation of the local intercritically reheated coarse-grained heat-affected zone (ICCGHAZ), which is often the most brittle part of welded joint. Figure 2 illustrates as an example how, in the case of a three-layer weld construction, the different subzones are formed in the weld as a result of the heat effect of each row of weld seams. In the ICCGHAZ located in the first welding pass the grains coarsen due to the heat cycle of the second pass, then the third weld pass intercritically reheat this area with heating between A_c1 and A_c3.

The examined t_8/5 interval was 5…30 s considering the characteristic welding heat input of arc welding processes. Rykalin-3D [16] model was used for the production of heat cycles which was set by the thermophysical properties of the investigated 500 MPa carbon steel as follows: density ρ = 7.7 g/cm³, specific heat at constant pressure c_p = 0.46 J/g·°C and thermal conductivity λ = 0.50 J/cm·s·°C. Figures 3 and 4 show the CGHAZ-W and ICHAZ-W cycles produced using the Rykalin model. In case of CGHAZ-W the heating rate was 1200°C/s and the holding time at peak temperature was 0.1 s, whilst in case of ICHAZ the heating rate was 200°C/s and holding time at peak temperature was 0.5 s. In both thermal cycles the cooling according to Rykalin-3D model finished at 250°C, after which the specimens cooled down freely to room temperature in the copper grips. In case of the simulation of ICCGHAZ the second cycle started when the first cycle decreased to 150°C. For the temperature control and measurement K-type thermocouple was used during the simulations. Thermocouples were welded to the surface with 1 mm distance from the weld centerline in order to avoid the effect of possible segregations.

The properties of the simulated HAZ subzones were analyzed by optical and scanning electron microscopy, macroscopic and hardness test, and instrumented Charpy V-notch impact test.

3.1 Microstructural characterization

Specimens were mechanically cut at the thermocouples. Then the surface was grinded and polished before etching with 2% Nital (alcoholic solution containing 2% HNO₃). The original weld microstructure consisting of acicular ferrite nucleated from the oxide inclusions is still visible after the thermal cycles (Fig. 5). However, as a result, the acicular ferrite laths have become finer. This is shown especially in CGHAZ-W with the t_8/5 cooling time of 5 s. In ICHAZ-W and ICCGHAZ-W samples it is seen that carbon has diffused at the prior austenite grain boundaries forming carbon-rich constituents. This tendency is expected to increase along with the increasing cooling time, although not clearly visible in the images. Grain boundary ferrite with ferrite side plates can be observed at the prior austenite grain boundaries in all cases. This is due to the initial allotriomorphic ferrite formation, and subsequent change of the transformation mechanism to promote the lath-like ferrite side plate formation.

More detailed observation into the weld metal microstructures was carried out using field emission scanning electron microscopy (FESEM, Zeiss Sigma) for the Nital-etched specimens. Figure 6 shows the as-deposited weld microstructure, whereas the HAZ-Ws are shown in Fig. 7. The applied acceleration voltage was 5 kV, and the working distance was in the range of 4–6 mm. As seen in Fig. 6a, the microstructure of the original weld metal consists mainly of acicular ferrite, although grain boundary ferrite is also present. The higher magnification image in Fig. 6b shows the acicular ferrite having nucleated from small (< 1 µm) non-metallic inclusions. The HAZ-W microstructures (Fig. 7) are presented with the corresponding hardness values of each sample to clarify the change in the microstructure. Proper analysis of the hardness results will follow in the Chap. 3.2. Generally, it can be observed that the acicular ferritic microstructure mostly remains in the samples after thermal cycles of HAZ-W simulations. However, the CGHAZ-W simulation with t_8/5 = 5 s evidently hardens the microstructure by refining the acicular ferrite laths. Though, with long cooling time i.e. t_8/5 = 30 s, the acicular ferrite laths either coarsen or the fraction of acicular ferrite decreases, which results as decreased hardness.

3.2 Macroscopic and hardness test

A Reicherter UH250 universal macro hardness tester was used for the hardness measurement. Evaluation was based on the EN ISO 15614-1 [17] welding qualification standard which allows HV_max = 380 HV10 hardness for the non-heat-treated welded joints of the thermomechanically controlled processed steels belonging to the 2nd group of CEN ISO/TR 15608 [18]. The measured transversal hardness distribution before the Gleeble simulation together with the macroscopic image is illustrated in Fig. 8.

Thanks to the selected matching filler material the weld has nearly the same hardness as the base material, whilst slight (10%) softening can be observed in the heat-affected zone.

Five hardness measurements were taken on the cross sections at the thermocouples of the Gleeble specimens. The average hardness results with the standard deviations are shown in Fig. 9. Compared to the 233 ± 1.5 HV10 base material and the 233 ± 2.2 HV10 original weld slight hardening can be observed in the investigated HAZ subzones, however the measured 265 HV10 peak hardness is safely under the permitted 380 HV10 value thanks for the mostly acicular ferrite microstructure. The medium (t_8/5 = 15 s) cooling time resulted equal hardness with the base material and the weld. In case of the examined HAZ zones a decreasing trend in hardness can be observed as the cooling time increases.

3.3 Instrumented Charpy V-notch impact toughness test

The Charpy V-notch (CVN) impact toughness tests have been performed on PSD 300 instrumented impact tester. During the evaluation it was considered that the required impact energy of the S500ML base material is 20 J at -40°C according to EN 10025-4 [12]. Based on the diagram in Fig. 10, thanks to the acicular ferrite microstructure, both the original weld and the examined HAZ subzones safely exceeded this relatively low requirement for impact energy. Compared to ferrite-pearlite and martensitic structural steels the HAZ subzones of S500ML are significantly more ductile. In contrary to the preliminary assumptions the CGHAZ-W results the highest toughness among the investigated zones. In case of ICHAZ and ICCGHAZ the impact energy decreased by the increasing cooling time, which tendency was not observed in CGHAZ. Based on the fracture surface analysis the original weld and the CGHAZ-W were also ductile, whilst ICHAZ and ICCGHAZ showed ductile-brittle characteristics. The ratio of brittle part was the highest in ICHAZ.

Figure 11 shows the macroscopic images of the fracture surfaces of the tested CVN specimens, whereas a closer look at the fracture surface of the original weld and ICCGHAZ-W with t_8/5 = 30 s is provided in Fig. 12. Original weld showed mainly ductile fracture behavior, and this is evident in the fractographies in Fig. 12a and b. However, regions of brittle facture are present also in the original weld, which could be attributed to the presence of grain boundary ferrite among the mostly acicular ferritic microstructure. As seen in Fig. 12b, a secondary crack was spotted in one of these brittle regions. Similar observations were made by Kang et al. [11] who concluded that the grain boundary ferrite regions in the otherwise acicular ferritic microstructure are the weakest locations in the low hardenability weld metal with similar composition to the present study. Fractographies in Fig. 12c and d are taken from the ICCGHAZ-W sample with t_8/5 = 30 s. This particular sample was one of the most brittle ones among the tested samples showing the absorbed impact energy of only 120.1 J. Therefore, the fracture surface was full of brittle fracture regions with only a slight ductile fracture regions appearing between them.

The relationship between expansion and Charpy impact energy is illustrated in Fig. 13. The expansion varied between 1.5 and 2.5 mm indicating favorable toughness properties.

Comparing to the conventional Charpy V-notch impact toughness test where the whole energy absorbed during the fracture is determined, the instrumented impact testing can provide more detailed information about the fracture process and the ductile/brittle behavior of the material. Using strain gauge measurement technology, the load-time diagram can be determined, and the characteristic points of the fracture process can be identified (the start of the plastic strain, maximal force, start of the unstable crack propagation, end of the unstable crack propagation) [19]. In the registered diagram, the unstable crack propagation stage is correlated with the amount of brittle fracture on the fracture surface. From the load-time diagram, the force-displacement diagram can be calculated. Assuming that the crack initiation occurs at the maximal force, the registered diagram can be divided into two parts according to the maximal force. Until the maximum force, the area under the curve is considered as the absorbed energy for crack initiation (W_i), while the remaining area is for crack propagation (W_p). As the ratio of the absorbed energy for the crack initiation increases, the toughness of the examined material reduces [19]. The detailed method of the instrumented Charpy V-notch pendulum impact test is described in EN ISO 14556 [20]. During the instrumented Charpy V-notch impact tests force-displacement diagrams were determined for the unaffected weld metal and the simulated HAZ subzones. Figure 14 shows the force-displacement diagram of the original weld metal, and the determined diagrams for the simulated HAZ subzones for the t_8/5=15 s. In case of the original weld and the CGHAZ-W the diagrams show typical ductile behavior, while ICHAZ-W and ICCGHAZ-W have ductile-brittle characteristics due to the observed instable crack propagation. The measured and calculated data of the Charpy V-notch impact toughness test is summarized in Table 3.

Table 3

Measured and calculated data of instrumented Charpy V-notch impact test
Zone	T_max, °C	t_8/5, s	F_max, kN	e, mm	CVN, J	W_i, J	W_ic, %	W_p, J	W_pc, %	Fracture mode
BM	-	-	22.7	2.23	183.4	55.5	30.6	126.7	69.4	mostly ductile
CGHAZ-W	1350	5	26.5	1.93	164.1	56.4	34.4	107.6	65.6	mostly ductile
		15	24.8	2.02	183.1	62.0	33.9	121.0	66.1	mostly ductile
		30	23.3	2.14	178.3	66.9	37.5	113.4	62.5	little instable crack propagation
ICHAZ-W	815	5	24.2	1.80	153.8	59.3	38.9	94.6	61.1	little instable crack propagation
		15	23.9	1.83	146.8	67.1	46.2	79.7	53.8	ductile-brittle
		30	23.1	1.65	124.7	54.2	37.7	72.9	57.3	ductile-brittle
ICCGHAZ-W	1350 815	5	24.9	1.83	153.4	61.0	40.2	92.5	59.8	little instable crack propagation
		15	24.0	1.83	144.5	62.4	43.0	82.1	57.0	ductile-brittle
		30	23.9	1.73	131.5	80.6	61.0	50.6	39.0	brittle-ductile

On the basis of the material tests carried out, it can be concluded that thanks to the acicular ferrite weld microstructure, even the HAZ subzones considered critical in terms of toughness are acceptable, they show a toughness significantly exceeding the standard requirement (20 J at -40°C), and also the degree of hardening and softening can be considered minimal. However, despite of the relatively high impact energy values instable crack propagation occurred in ICHAZ-W and ICCGHAZ-W indicating ductile-brittle behavior. Increasing the t_8/5 cooling time also reduces the hardness of the HAZ subzones formed in the weld and the toughness of the ICHAZ and ICCGHAZ, therefore, overall, it is advisable to weld these steels with a low welding heat input.

Physical simulation experiments were carried out on S500ML shipbuilding steel for the analysis of the effect of multipass welding on weld properties. The main findings of the research work are summarized as follows:

With the help of physical simulation, it was possible to produce HAZ subzones critical to toughness in the multi-layer weld seam of the examined steel, which were as follows: coarse-grained heat-affected zone (CGHAZ-W), intercritical heat-affected zone (ICHAZ-W), intercritical coarse-grained zone (ICCGHAZ-W).

Significant hardening and softening were not observed in the examined CGHAZ-W, ICHAZ-W and ICCGHAZ-W. The t_8/5 = 15 s resulted in nearly equal hardness with the original weld and base material.

Mainly acicular ferrite was observed in simulated HAZ subzones formed in the weld, although also grain boundary ferrite and ferrite side plates were present. The ferrite laths had become finer and carbon-rich constituents had formed at the prior austenite grain boundaries due to the thermal cycles.

The toughness of the original weld metal and the simulated HAZ subzones formed in the weld significantly exceeded the base material requirement with more than 100 J impact energy at -40°C.

The measured impact energy of the examined HAZ subzones formed in the weld was under the original weld toughness in all cases. The longer t_8/5 cooling time resulted lower impact energy in ICHAZ-W, although this tendency was not observed in CGHAZ-W. Overall, the ICHAZ-W showed the lowest toughness among all subzones, which was indicated by the relatively large instable crack propagation in the impact energy diagrams.

Grain boundary ferrite and ferrite side plate regions were suggested to have been the weakest regions in otherwise mainly acicular ferritic microstructures.

The authors declare no competing interests.

Funding

The financial support of the János Bolyai Research Scholarship of the Hungarian Academy of Sciences (Grant number: Bo/00643/22/6) and Business Finland, in project Fossil-Free Steel Applications: Phase II (Grant number: Dnro 5562/31/2023) is gratefully acknowledged.

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Physical simulation-based analysis of multipass welding in S500 shipbuilding steel

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Abstract

Figures

1. Introduction

2. Materials and methods

1.1. Base and filler materials

1.2. Experimental circumstances

3. Materials tests and discussion

3.1 Microstructural characterization

3.2 Macroscopic and hardness test

3.3 Instrumented Charpy V-notch impact toughness test

4. Conclusions

Declarations

Funding

References

Status:

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