Hybrid laser arc welding of thick high-strength pipeline steels of grade X120 with adapted heat input

Journal of Materials Processing Tech. 275 (2020) 116358

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Hybrid laser arc welding of thick high-strength pipeline steels of grade X120 with adapted heat input

T

Ömer Üstündağa, , Sergej Gooka, Andrey Gumenyuka,b, Michael Rethmeiera,b,c ⁎

a

Fraunhofer Institute for Production Systems and Design Technology, Berlin, Germany Bundesanstalt für Materialforschung und -prüfung, Berlin, Germany c Institute of Machine Tools and Factory Management, Technische Universität Berlin, Berlin, Germany b

ARTICLE INFO

ABSTRACT

Associate Editor: C.H. Caceres and S.J. Na

The influence of heat input and welding speed on the microstructure and mechanical properties of single-pass hybrid laser arc welded 20 mm thick plates of high-strength pipeline steel X120 were presented. The heat input was varied in the range of 1.4 kJ mm−1 to 2.9 kJ mm−1, while the welding speed was changed between 0.5 m min−1 and 1.5 m min−1. A novel technique of bath support based on external oscillating electromagnetic field was used to compensate the hydrostatic pressure at low welding velocities. A major advantage of this technology is, that the welding speed and thus the cooling time t8/5 can be variated in a wide parameter window without issues regarding the weld root quality. The recommended welding thermal cycles for the pipeline steel X120 can be met by that way. All tested Charpy-V specimens meet the requirements of API 5 L regarding the impact energy. For higher heat inputs the average impact energy was 144 ± 37 J at a testing temperature of −40 °C. High heat input above 1.6 kJ mm−1 leads to softening in the weld metal and heat-affected-zone resulting in loss of strength. The minimum tensile strength of 915 MPa could be achieved at heat inputs between 1.4 kJ mm−1 and 1.6 kJ mm−1.

Keywords: High-strength low-alloy steel Hybrid laser-arc welding Mechanical-technological properties Microstructure Toughness Pipeline steel of grade X120

1. Introduction Ultra-high-strength steel grades allow higher operating pressure in natural gas transport lines without increasing the tube wall thickness. This in turn brings many economic benefits, such as lower material consumption and a reduction in transport and manufacturing costs. Witek (2015) reports that the use of higher strength steel for onshore pipelines results in significantly lower manpower, lower logistical costs and lower extend of welding costs. The demand for higher strength steels continues to rise steadily, particularly with the exploration proceeding into deep sea and arctic regions. The operation of pipelines for the efficient and safe transport of gaseous fuels under harsh environmental conditions requires excellent impact toughness and tensile strength of the materials used. Villalobos et al. (2018) report in their review study of current development of micro-alloyed steels, that the pipeline steel grade X120 produced by thermo-mechanical controlled processing (TMCP) followed by accelerated cooling (ACC) combines excellent ductility, high strength and good impact toughness at low temperatures. For the high yield strength of 120 ksi which corresponds to 830 MPa, the concept of microalloying also plays a decisive role. The base alloying elements such as Cu, Ni and Mo as well as micro-alloying elements such as V, Nb, Ti, B and Mn are used to impart the desired properties to the steel. Okaguchi et al. (2003) worked on the

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development of X120 and found that the micro-alloying elements suppress grain coarsening to ensure a very fine grain size of approx. 1 μm–2 μm which leads to an optimum balance of strength and toughness. Ishikawa et al. (2006) developed a cooling concept that leads to the formation of a high ductility by producing a biphasic microstructure of ferrite and bainite when rolling the X120 sheet. An alloying of B causes a positive effect to the HAZ toughness while producing a lower bainite phase. The sensitivity to hot cracks can be eliminated by addition of Mn due to setting of S as observed by Lippold (2014). Beidokhti et al. (2009) found that a low percentage alloying addition of Ti promotes a nucleation of acicular ferrite (AF), which affects an increase of weld toughness. It has to be mentioned that the manufacturing procedures of this perspective steel grade are constantly being investigated and improved. Kong et al. (2015) showed in their work that the addition of the elements Ti and B into steels increase yield and tensile strength level. The strength can also be increased gradually with decreasing final accelerated cooling temperature during rolling. Asahi et al. (2004) reported that the large diameter tubes of X120 can be manufactured for mass production under proper forming conditions and certain tool design criteria. In general, pipeline steels of grade X120 have a reduced carbon content of less than 0.1 wt % and can be characterized as having good weldability from a metallurgical point of view.

Corresponding author at: Pascalstraße 8-9, 10587, Berlin. E-mail address: [email protected] (Ö. Üstündağ).

https://doi.org/10.1016/j.jmatprotec.2019.116358 Received 26 February 2019; Received in revised form 30 July 2019; Accepted 4 August 2019 Available online 12 August 2019 0924-0136/ © 2019 Elsevier B.V. All rights reserved.

Journal of Materials Processing Tech. 275 (2020) 116358

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Nevertheless, there is also a high demand for future works to develop suitable welding consumables especially for grades like X120. A literature review shows that several studies have been carried out in this regard. Usually, it is expected that the filler wire has overmatching strength compared to the base material. One of the latest developments of matching welding wire for SAW of X120 is described in patent by Zhang et al. (2017). The developed welding wire ensures a tensile strength of approx. 960 MPa. However, the modern steels X120 can have a tensile strength of over 1000 MPa. By simply addition of higher alloy content, the problem cannot be solved since it can lead to lower welded joint toughness or cold cracking. In some circumstances, an undermatched filler wire is preferable and can lead to acceptable results regarding to mechanical properties of welded joints. Gook et al. (2014) showed that the minimum tensile strength and impact energy according to API 5 L could be achieved even with an undermatching strength of metal-cored wire by HLAW of 20 mm thick plates of pipeline steel X120. As it stated by Liu and Bhole (2013), the development of a suitable welding system is another fundamental challenge in welding X120. Arcbased welding processes such as manual shielded metal arc welding (SMAW) and mechanized gas metal arc welding (GMAW) can be used to weld thick X120 sheets without significant problems, as it described by Gräf et al. (2003). These processes have a small melt pool, which is preferable for position welding, and are mostly used as field girth welding methods for pipe laying. A major disadvantage of these processes is the limitation of the penetration depth and the low productivity. A multi-wire SAW welding is required to fill the large-volume grooves of longitudinal seams in pipeline production. This high productivity process results in high heat input, resulting in a subsequent wide softening zone near the HAZ and a deterioration in toughness due to grain growth. To take this aspect into account, the X120 steel contains some amount of V for its precipitation hardening effect, as shown by Hillenbrand et al. (2004). In addition to coarse grain, larger amounts of martensite-retained austenite (MA) islands can significantly reduce HAZ toughness, as reported by Huda et al. (2016). The steel manufacturer Voestalpine (2012) recommends the optimal temperature-time cycles for the high strength TMCP steels. The characteristic cooling time from 800 °C to 500 °C (t8/5-time) should be in the range of about 3 s to 15 s. An empirical estimate shows that a multi-wire SAW process with an energy per unit length of approx. 9.5 kJ mm−1 is required to fill the 20 mm thick V-groove. Depending on the total wall thickness of the welding part, t8/5-times of 90 s and more are to be expected. It is therefore to be expected with a high degree of softening of the base material. The production experience available today is not sufficient to be able to assess the softening occurring in the base material in adjacent the weld. In each case, the excessive heat input has to be controlled and minimized. Lan et al. (2016) demonstrated that the acceptable impact toughness can be obtained first at a reduced heat input of 1.43 kJ mm− 1 . In another study by Wang et al. (2012) the same approach was indicated. In this work 12.7 mm thick plates X120 were double side welded with the SAW process and a compressed air was blown into the joint to induce a fast cooling process. The absorbed impact energy was increased when using this method. For thicker plates and from the perspective of production safety is impossible to reduce the heat input. An alternative welding process that provides high weld productivity with low heat input is laser beam welding (LBW). It is characterized by deep penetration, high welding speed, narrow weld zone and HAZ. On the other hand, LBW causes, due to the high cooling rate, an increase in hardness and deteriorates toughness of the weld metal owing to formation of martensite. Efimenko et al. (2010) recommended that for cold cracking susceptibility the cooling rate should not exceed 80 °C s−1. Cooper et al. (2005) conducted welding trials with preheating on API 5 L-X80 weld joints and concluded that attempts to reduce the cooling rate by preheating up to 100 °C showed no significant differences regarding to weld metal toughness and can be neglected. As a compromise, the hybrid laser arc welding (HLAW) was developed in the end of the 1970s and enables by coupling of LBW and GMAW in the

same weld pool a reduced cooling rate comparing to LBW and a lower heat input, an increase of the productivity and a low consumption of filler wire in comparison to GMAW or SAW (Eboo et al., 1978). Several studies show that the HLAW was applied successfully for welding of 9.3 mm thick HSLA steels (Cao et al., 2011), 12.5 mm with different joints (Atabaki et al., 2014) or double-side welding technique (Chen et al., 2013). Rethmeier et al. (2009) reported that 32 mm thick materials could be welded using multi-layer technique combining HLAW and GMAW processes. The application of single-pass HLAW for welding of thick sections is still limited due to certain technological aspects. One of the limiting factors is the formation of gravity drop-outs, which usually occurs when welding thick plates in flat position and reduced welding velocity. In this case, the hydrostatic pressure exceeds the Laplace pressure, which is dependent on the surface tension. An increase of the weld root width at reduced welding speeds leads to a reduction of the surface tension, why a stable process can be realized only for high enough welding speed especially in flat position, which leads to high cooling rates. To overcome the problem of sagging and to control the heat input due to the adapted welding speed, a novel technique of weld pool support based on generating Lorentz forces in the weld pool, due to oscillating magnetic field and induced eddy currents is used in this work. This technique works contactless, which is a significant advantage over usual bath supports. The oscillating magnetic field B is perpendicular to the welding direction and is produced by an AC magnet. The electric density j is parallel to the welding direction. The resulting Lorentz force FL=B x j is directed upwards and counteracting the hydrostatic pressure ph and the arc pressure parc. The scheme of the electromagnetic weld pool support system is shown in Fig. 1. The effectiveness of this melt pool control system has already been demonstrated for laser beam welding of up to 30 mm thick AlMg3 plates by Avilov et al. (2012). It could shown by Üstündağ et al. (2018a) that the electromagnetic weld pool support system can also be applied for HLAW on ferromagnetic steels. The most important assumption here is that the skin layer depth δ is less than the plate thickness, so that the stability of the arc is not influenced by the oscillating magnetic field. Üstündag et al. (2018b) provided a study on tolerances regarding geometrical nonuniformity of the weld edges. A good gap bridgeability of up to 1 mm as well as misalignment tolerances of up to 2 mm at HLAW of 25 mm thick steel sheets with electromagnetic weld pool support system could be proved. The proposed technology is intended to be introduced for the manufacturing of gas transition pipelines as well as containers and vessels in apparatus construction. In such applications, the assurance of the weld quality, in particular, the control of the lack of root fusion and full penetration, is of primary importance. As described by Dupriez and Truckenbrodt (2016), recent experiences with modern methods of process monitoring in laser material processing such as optical

Fig. 1. Scheme of electromagnetic weld pool support system according to Avilov et al. (2012). 2

Journal of Materials Processing Tech. 275 (2020) 116358

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coherence tomography (OCT) enable highly accurate automatic seam tracking and offer the possibility of non-destructive inline process monitoring and quality assurance during laser welding. The aim of this study is to investigate the influence of heat input and welding speed on the mechanical properties such as tensile strength and toughness of HLAW-welded thick steel sheets X120 while using electromagnetic weld pool support and maintaining quality requirements for the welds.

from one side in butt joint. The edge preparation was a single V groove with a broad root face of 14 mm and an opening angle of 45°. A metalcored wire Megafil 742M (ISO 187276: T 69 6 Mn2NiCrMo M M 1 H5) with a diameter of 1.2 mm was used for the HLAW experiments. Table 1 and Table 2 show the chemical composition and mechanical properties of the materials used. The chemical composition of the base material and filler wire was identified by plasma spectrometric analysis at the BAM. The carbon equivalent (CEPcm), which is an indicator for weldability and provides details about hardenability, crack susceptibility, change in microstructure and mechanical properties after welding, amounts to 0.19. This indicates a good weldability. The CEPcm was calculated according to the API specification 5L (2004). For steels with a carbon content less than 0.12 wt % the CEPcm is estimated as follows:

2. Experimental setup 2.1. Experimental equipment The high power fibre laser IPG YLR-20000 with a maximum output power of 20 kW, an emission wave-length of 1070 nm and a beam parameter product of 11 mm x mrad served as laser beam source. The laser radiation was transmitted through an optical fiber with a core diameter of 200 μm. A laser processing head BIMO HP from HIGHYAG with a focal length of 350 mm providing a spot focus diameter of 0.56 mm was used. A microprocessor-controlled welding machine Qineo Pulse 600 with a maximum current of 600 A was applied as a power source for the arc. The laser optics and GMAW torch were mounted on the robot arm, where the laser axis was positioned 90° to the weld specimen surface and the GMA torch was tilted 25° relative to the laser axis. The experiments were carried out with an arc leading position and a distance of 4 mm between the two heat sources. The laser beam was underfocused 11 mm below the top surface of the weld specimen. The wire stick-out was kept at 18 mm, see Fig. 2(a). The AC magnet was positioned 2 mm below the workpiece, where the distance between the two magnet poles was 25 mm. The laser optics and GMAW torch were positioned exactly in the middle between the magnet poles. The workpiece was moved by an x-y-positioning table. The processing head and the magnet remained in a fixed position during the welding. The experimental setup can be seen in Fig. 2(b).

CEPcm = C +

Si Mn Cu Ni Cr Mo V + + + + + + + 5B 30 20 20 60 20 15 10

(1)

The welds were performed under the gas shielding. The gas mixture consisted of 82% Ar and 18% CO2. The gas flow rate was 20 l min−1. 2.3. Variables and analysis A series of experiments was designed with five levels of welding speed and the resulting heat input by adjusting the power of the laser and arc. Table 3 summarizes the main welding parameters. The arc current and voltage are averages according to the arc power source. The total line energy input for HLAW (Qtotal in kJ mm−1) is calculated as the sum of the laser beam energy input (QL) and the arc energy input (QArc). The equation is also used by Bunaziv et al. (2018) to calculate the heat input. The efficiency factor for GMAW is taken in account with 0.8.

Qtotal = QL + Q Arc =

60 P 60 I U + x 0.8 vw 1000 vw

(2)

where P is the laser power (kW), I is the arc current (A), U is the arc voltage (V) and vw is the welding velocity (mm min−1). During the experiments the AC magnet was operated at an oscillating frequency of 1.2 kHz and a magnet power of 1.7 kW to 2.1 kW. The magnetic field strength for an ideal compensation of the root excess weld metal was between 118 mT and 128 mT. The influence of the magnetic field strength was investigated by Üstündağ et al. (2018a). It could be shown, that higher strength of the magnetic field will cause an

2.2. Welding materials The welding experiments of the present study were conducted on 20 mm thick plates of X120 grade HSLA steel. The welds were performed

Fig. 2. HLAW with electromagnetic weld pool support system: (a) process arrangement and (b) experimental setup. 3

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Table 1 Chemical composition of materials used, shown in wt %. Material/Element

C

Si

Mn

P

S

Al

Cr

Ni

base material X120 welding wire Megafil 742 (T 69 6 Mn2NiCrMo M M 1 H5)

0.05 0.05 Mo 0.22 0.45

0.32 0.09 Cu 0.03 –

1.58 1.84 V 0.038 –

0.009 0.022 Nb 0.03 –

< 0.001 0.010 Ti 0.01 –

0.02 – B 0.001 –

0.39 0.52 N 0.003 –

0.04 1.99 Fe bal. bal.

base material X120 welding wire Megafil 742 (T 69 6 Mn2NiCrMo M M 1 H5)

Table 2 Mechanical properties of materials used. Material

Rp0.2 in MPa

Rm in MPa

A in %

AV in J

base material X120 welding wire Megafil 742* (T 69 6 Mn2NiCrMo M M 1 H5)

816 > 690

1020 770 940

14.7 17

270 (−40 °C) 75 (−40 °C)

* Mechanical properties of the pure weld metal. Table 3 Main welding parameters for single-pass HLAW of 20 mm thick X120 steels. Sample Arc current Arc voltage Welding in A in V velocity in cm min−1

Laser power in kW

Heat input in kJ mm−1

1 2 3 4 5

15.3 15.3 16.5 17 17.5

2.9 2.2 1.9 1.6 1.4

330 430 500 530 530

34 37 38 40 40

50 75 100 125 150

Fig. 3. Schematic representation of the hardness measurement lines.

16.5 kW laser power with the electromagnetic weld pool support technique. The root side of this sample is also visible. It can be observed that the root is ideally compensated by the generated electromagnetic pressure over the entire seam length, why the welded joint satisfied the requirements related to quality B according to ISO 12932. The metallographic evaluation of the cross-sectional weld shape formation lets conclude that the full penetrated welds are formed to a wine-cup shape which is typical for HLAW. The most important welding parameters that influence weld seam quality and weld shape are welding speed, laser power, arc power and HLAW process configuration such as leading mode of the GMAW arc. The shape of the thick walled HLAW weld can be categorized in two parts. The upper part is an arc dominated zone with a wide fusion area and HAZ, while the lower part is laser dominated and is characterized by narrow and nearly parallel seam flanks. Fig. 6 shows exemplary cross sections of single-layer welded 20 mm thick plates of pipeline steel X120 with different welding velocities and resulting heat inputs. The 20 mm thick HLAW welds could be performed without sagging, gravity drop-out and inadmissible root reinforcement with help of the electromagnetic weld pool support system. The hydrostatic pressure could be effectively compensated by the upward directed Lorentz force for all welded joints, produced with different welding velocities variated in a wide range between 0.5 m min−1 and 1.5 m min−1. The higher wire feed speed was needed to fill the groove, especially at high welding velocities 1.25 m min−1 and 1.5 m min−1. For a qualitative analysis of the root over the entire seam length and visualization that the weld pool support system works safely, the root excess weld metal was measured with a laser profile scanner. It is evident, that there is a slight overcompensation of the root through the electromagnetic weld pool support system. According to results of Üstündağ et al. (2018b), the acting magnetic force can be optimized so that the overcompensating effect can be resolved. This point was not further investigated within the framework of this study, since the objective was to examine the mechanical properties of the welds. In this context the tensile and impact test specimens were partly extracted in the mid-thickness. The laser profile scan of the weld root side is exemplarily shown in Fig. 7. Shallower temperature gradients in the HAZ resulting from high heat input such as at a welding speed of 0.5 m min−1 lead to formation of wider coarse-grained heat affected zone (CGHAZ). The grain size

overcompensation of the root due to higher electromagnetic pressure. A high oscillating frequency was necessary to protect the electric arc on the top side from oscillating magnetic field and current. At this frequency the skin layer depth, where the eddy currents have an effect, was 16.8 mm, which is less than the material thickness of 20 mm. Otherwise, the electric arc will be deflected by the currents (Üstündağ et al., 2018a). After welding, transverse sections in the mid-length position were cut for metallographic inspection. The cut surfaces were prepared by polishing and etching using 2% nital solution. The Vickers hardness measurement was carried out with Leitz microhardness tester Miniload 2 and data logger Leitz RZD-DO. The tests were performed under a load of 4.9 N and a dwell time of 15 s. The Vickers hardness testing machine was calibrated according to DIN EN ISO 6507-3 with a maximum deviation of the HV0.5 measurements of ± 3%. The measurements were performed on three levels over the material thickness: upper part (2 mm below the surface), middle and root part (2 mm above the bottom). A schematic representation of the hardness measuring lines is shown in Fig. 3. Charpy V-notched specimens were extracted 4 mm below the top surface of the welded specimens and then machined according to DIN EN ISO 148-1 with standard 55 x 10 x 10 mm3 dimensions. The notch was placed in the WM center. All the Charpy tests were carried out at -40 °C. Furthermore, round tensile test specimens according to DIN 50125 form B with a diameter of 12 mm were extracted from the middle of the plate thickness. For each sample five to ten impact tests and three tensile tests were performed. The specimen-taking plans of Charpy V-notched specimens and tensile test specimens are shown in Fig. 4(a) and (b) respectively. 3. Results and discussion The visual test of the welded joint did not reveal imperfections, such as cracks, porosity, incomplete fusion or lack of penetration. A welded joint is shown in Fig. 5 exemplary. The specimen was welded with a heat input of 1.9 kJ mm−1 at a welding speed of 1 m min−1 using 4

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Fig. 4. Specimen-taking plan: (a) Charpy V-notched specimens; (b) tensile specimens and dimension.

within the HAZ gradually varied from the fused zone (FZ) to the finegrained heat affected zone (FGHAZ) due to different temperatures prevailing in these areas during the welding. Therefore, the HAZ can be divided in CGHAZ, FGHAZ and inter-critical heat affected zone (ICHAZ). The CGHAZ surrounds the weld metal. A wide CGHAZ can reduce the hot crack sensitivity because of a larger grain boundary area over which stress relaxation can be accommodated. Loss of strength is to be expected due to the grain coarsening. It is also recognizable that the HAZ becomes narrower with decreasing heat input. At the lowest welding velocity, the HAZ in the mid-thickness has a width of approx. 5.4 mm. The width of the HAZ is reduced with increasing welding speed to 0.8 mm. The width of the FZ is also decreased with increasing welding speed. The whole width of the welding zone including FZ and HAZ should be minimized. For analyses of the weld shape and the different zones in the HAZ, the width of this areas was measured as schematic can be seen in Fig. 8. Fig. 9 clearly shows the change of the geometrical shape of the FZ in different layers of the material thickness and the width of the HAZ in the mid-thickness. It is apparent that the width of the FZ and HAZ rises with increasing heat input. The width of the upper part is larger because of the V groove, which has to be filled by welding wire. Due to the high wire feeding rate related to the welding velocity at a high heat input of 2.9 kJ mm−1 the width of the upper part of the weld rising faster in comparison to the root part, where the arc process has no significant influence.

During the welding experiments the temperature was measured via a pyrometer Metis MQ22 with a measuring range from 350 °C to 1300 °C on the top side in a horizontal distance of 1.5 mm to the laser beam to determine the cooling conditions. The characteristic cooling time from 800 °C to 500 °C (t8/5-time), grows with increasing heat input. For the lowest welding speed of 0.5 m min−1 the t8/5-time time is approx. 13 s, which corresponds to the cooling rate of 23 °C s−1. A relatively short t8/5-time of 2 s could be measured with decreasing heat input at higher welding velocity of 1.5 m min−1, see Fig. 9. The microstructure of the welded samples is primary determined by the local thermal cycles, the cooling conditions and the peak temperature that is achieved during the welding process. It is also of great practical interest whether the high heat input leads to the formation of a soft microstructure, whereby the deterioration of the strength can be caused. Microscopic images were made, to investigate the impact of welding parameters such as welding velocity and heat input on the HAZ size, the microstructure and grain size distribution in the HAZ. Fig. 10 shows the microstructural development of HLAW welds X120 performed with different heat inputs. The metallographic examination was carried out in the wide area in adjacent of fusion line (FL) including HAZ and BM. The average grain size was defined on the metallographic micrographs by software and in accordance with ASTM Standard E 112. The correlation between the heat input and the width of the HAZ can be recognized.

Fig. 5. The face and the root of the hybrid laser arc welded joint with electromagnetic weld pool support technique using 16.5 kW laser power at a welding speed of 1 m min−1. 5

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Fig. 6. Cross sections of single-layer HLAW on 20 mm thick plates X120 with different heat inputs. Fig. 7. Laser profile scan of the root over the total seam length.

A high heat input of 2.9 kJ mm−1 leads to a wide HAZ, where grain coarsening could be observed. The average grain size in the CGHAZ and FGHAZ is 18.9 μm and 3.3 μm respectively. The grain size in the FGHAZ varies in a range between 2.4 μm and 4.7 μm depending on the heat input. Since the FGHAZ, ICHAZ and sub-critical HAZ (SCHAZ) are not exposed to high peak temperatures due to increasing distance to the FZ and to the heat sources, these zones consist mainly of small grains and for these areas of the weld joint a high impact toughness can be expected. The grain sizes in the CGHAZ fall to 10.6 μm even for the lowest heat input of 1.4 kJ mm−1 investigated. It is to be expected that the tensile strength will increase with decreasing heat input due to finer grains. Fig. 11 shows the grain size distribution in the CGHAZ and FGHAZ for the different heat inputs. Both the CGHAZ and the FCHAZ have the bainitic or bainitic-martensitic microstructure, while the BM mainly consists of a lower bainite due to niobium and boron additions in the manufacturing process. The comparison of the WM microstructure depending on the heat input is shown in Fig. 12. For all samples, the microstructure dominated by bainite of different morphologies. The granular bainitic (GB) microstructure, which is characteristic for the WM at the lowest welding speed of 0.5 mm min−1 and corresponding t8/5-time of 13 s is shown in Fig. 12(a). The GB grains are coarse and stochastically distributed. This observation correlates with results of Bruce and Boring (2006), who showed in their work that GB forms at lower cooling rates. It can be also seen that the diameter of the prior austenite grains decreases with increasing welding speed. For higher welding velocities of 0.75 m min-1 and 1 m min-1 (Fig. 12(b)–(c)) with a cooling time (t8/5) of 7 s and 5 s correspondingly, the WM is predominately composed of GB with some

Fig. 9. Weld shape geometry depending on heat input and t8/5-time.

amount of lath bainite (LB). The prior austenite grain boundaries are clearly visible. Fine bainite laths grow from the grain border to the grains middle and within the prior austenite grains. Laths having the same orientation form packets and blocks, and the laths with different orientations split the prior austenite grain into different grains. As is discussed

Fig. 8. Schematic drawing for the analyses of the weld shape geometry. 6

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Fig. 10. Microstructural development of HLAW welds X120 in adjacent of FL including HAZ and BM at different heat inputs.

In addition to the thermal cycle during welding, the resulting microstructure of the WM is also adjusted by the metallurgy of the welding wire. Thus, 0.45% Mo is added to the welding wire as a trace element, whereby the strength of the weld metal and the impact resistance at low temperatures is improved. As it shown by Villalobos et al. (2018) the addition of a certain amount of Mo reduces the phase transition temperature during the post-weld cooling process. This refines weld metal structure while expanding the formation temperature range of bainite. Fig. 13 shows the average values of the microhardness HV0.5 of WM (Fig. 13(a)) and HAZ (Fig. 13(b)) for different heat inputs. The increasing deterioration of the hardness with increasing heat input can be clearly seen. The largest softening effect can be observed for the weld, which is performed at a welding velocity of 0.5 m min−1 and thus highest heat input. The average hardness values in WM and HAZ fall accordingly to 240 HV and 290 HV in this case. Such a degradation in hardness compared to the hardness of the base metal of 310 HV can lead to a significant deterioration of the weld strength. It is also remarkable that the process heat is distributed homogeneously through the thickness of the weld at a low welding speed of 0.5 m min-1. The indication for this effect is that the hardness values for the upper part, the middle part and the root part of the WM are at the level of 240 ± 20 HV, causing about the same degree of softening. At higher welding speeds, differences in local hardness levels can be observed for different parts of the welds. The reason for this effect is an uneven distribution of the process heat during the welding in the upper part and root part of the weld. For the welding speeds of 1 m min-1 and above, it is characteristic that the hardness of the root part is higher due to the lower heat input and the associated higher cooling rate than in the upper part. In other words, the supplied energy of the arc has no significant thermal effect on the root part of the seam for the weld thickness of 20 mm. In general, the results show that a high heat input in the range of 1.9 kJ mm-1 to 2.9 kJ mm-1 leads to a noticeable softening in the HAZ and WM. Only at higher welding speeds from 1.25 m min-1 and a heat input of 1.6 kJ mm-1, the hardness of the WM reaches that of the base material, especially in the root part. This suggests that

Fig. 11. Grain sizes in the HAZ of HLAW welds X120 depending on the heat input.

by Zhou et al. (2018), the boundaries of laths can effectively inhibit the crack propagation and improve the toughness, therefore the mix microstructure of both GB and LB is advantageous for the toughness. With further increasing of welding velocity and reducing of heat input to 1.6 kJ mm-1 and 1.4 kJ mm-1, the microstructure transforms to bainiticmartensitic microstructure consisted of LB and M/A constituents which is a mixture of martensite and retained austenite. M/A constituents mainly distributed on the prior austenite grain boundaries and lath boundaries. The WM toughness can be positively influenced by small amount and small size of the M/A constituents. However, as it shown by Shin et al. (2009) the M/A constituents are hard phases of the WM microstructure and a large number of M/A constituents can easily initiate cracking due to high stress level at M/A matrix interfaces. As a result, it can be expected that the fracture probability of M/A constituents increases with increasing number and size of M/A constituents. 7

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Fig. 12. WM microstructure of HLAW welds X120 depending on the heat input.

no strength losses are to be expected in these welds. The maximum allowed hardness is 35 HRC according to the standards of the American Petroleum Institute API 5 L which corresponds to 332 HV. Samples welded with the highest welding speed of 1.5 m min−1 show peak WM hardness of 364 HV in the root part of the weld, which exceeds the permitted level. Nevertheless, the average hardness in the laser zone of 326 ± 31 HV meets the requirements of the standards. When analyzing the hardness measurements, it can be seen that the highest WM hardness is present in the root of the weld (Fig. 13(a)), whereas for the HAZ, the maximum hardness is at the weld seam top (Fig. 13(b)). However, a relatively large scattering of measured hardness values in the HAZ has to be taken into account when interpreting these results. During the hardness test, the surrounding material is deformed by the imprints, which changes the material properties. In order to prevent a wrong interpretation of the determined hardness, the distance between the imprints during the Vickers hardness test was chosen to be 0.5 mm. On the other hand, the HAZ of laser hybrid welds is relatively narrow, consists of heterogeneity regions, and it is therefore technically difficult to reproducibly place the hardness imprints in

the HAZ regions having the same hardness properties. Fig. 14 shows exemplarily the fluctuations of the hardness values in HAZ measured symmetrically on the right and left side of the weld performed with a heat input of 2.2 kJ mm−1 at a welding speed of 0.75 m min−1. The results speak for the fact that the hardness values of the HAZ are on average slightly lower than those of the base material for different heat inputs, but no reliable correlation can be demonstrated between the degree of material softening in the HAZ and the heat input. The results of the tensile test show that all specimens are broken in the WM or at the transition to the HAZ. The welded samples with lower heat input are broken near the HAZ because of the grain coarsening at the border to the WM. The minimum tensile strength of 915 MPa according to API 5 L is reached only for higher welding speeds higher than 1.25 m min−1. Due to the reduced heat input and increasing hardness in the weld metal it was expected that these samples achieve the requirements regarding the tensile strength. The reduced heat input leads also to a lower grain coarsening which impact the strength of the seam positively. The results of the tensile testing according to DIN EN ISO 6892-1 for various welding velocities are shown in Fig. 15. The joint

Fig. 13. Results of the Vickers hardness for the HLAW welds performed with different welding velocities: a) microhardness of WM; b) microhardness of HAZ. 8

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Fig. 14. Individual Vickers microhardness values of HAZ for the HLAW weld performed with a heat input of 2.2 kJ mm−1 at a welding speed 0.75 m min−1.

Fig. 15. Results of the tensile test and joint efficiency for the HLAW welds performed with different welding velocities.

Fig. 17. Stress strain curve for investigated samples with different welding velocities.

efficiency, the tensile strength of the used materials and the minimum tensile strength for welded joints of 915 MPa according to API 5 L and EN ISO 3183 can be seen in Fig. 15 as well. Pandey et al. (2019a, 2019b) show that joint efficiency is a meaningful indicator of weld quality for various welding conditions when joining high strength steels for high-temperature applications. In the present work, this coefficient was calculated according to the following

relationship:

Weld joint efficiency =

tensile strength of welded joint * 100% tensile strength of base metall

(3)

From the results it can be deduced that the efficiency of the welded joint is in direct correlation with the heat input. The lowest joint efficiency of 75% characterizes the soft WM caused by high heat input of

Fig. 16. Fracture location of the transverse tensile specimens welded by HLAW; (a) heat input 2.9 kJ mm−1 and (b) heat input 1.4 kJ mm−1. 9

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Fig. 18. SEM images of a tensile test specimen welded using 16.5 kW laser power and 19.3 kW arc power at a welding speed of 1 m min−1.

2.9 kJ mm−1. Nevertheless, the minimum WM strength of 770 MPa guaranteed by the welding wire (see Table 2) could be achieved. The maximum joint efficiency of 93% could be calculated for the heat input of 1.4 kJ mm−1. Reason for this increased joint efficiency is a shorter residence time of the melt in the high temperature range and the formation of a much narrower softening zone. The results in terms of weld metal hardness and tensile strength are consistent with the regression equations of Turichin et al. (2018) with an accuracy of up to 15%. Fig. 16 shows exemplarily two tensile test samples, which were produced by the HLAW process with the highest investigated heat input and the lowest heat input, with fracture location in the WM and on the border to the HAZ. All samples show a cup and cone fracture with a necking in the WM.

Table 4 Results of notched-bar impact test (at −40 °C) of hybrid laser welded samples with different heat inputs. vw in m min−1

Energy per unit length in kJ mm−1

Mean absorbed energy (KV) at −40 °C testing temperature in J

Standard Deviation in J

0.5 0.75 1 1.25 1.5

2.9 2.2 1.9 1.6 1.4

144 149 75 117 107

11 28 14 35 35

Fig. 19. Fracture surfaces of tested Charpy impact test specimens for different heat inputs. 10

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Fig. 20. Fracture path images of notched-bar impact test results.

The remaining samples tested showed a ductile fracture behavior (honeycombed surfaces). The Charpy impact test was conducted in accordance with EN ISO 148-1 at a test temperature of -40 °C. All tested samples exceed the required minimum impact toughness of 40 J (for 0 °C) according to API 5 L. The results of the Charpy impact test are shown in Table 4. The evaluation of the fracture surface morphology revealed that there is a mixed ductile-brittle fracture for all samples. The fracture surfaces of tested samples once per parameter set are shown in Fig. 19. It could be recognized a tendency that an increase of the heat input leads to an increase of the impact energy. This effect is supported by the fact that with an increased t8/5 time, the resulting WM structure contains less amount of the hard-martensitic phase that promotes the embrittlement. The mixed WM microstructure of GB and LB, which is characteristically for the heat inputs 2.9 kJ mm−1 and 2.2 kJ mm-1, is beneficial for the toughness. It has to be mentioned that one of the tested samples has a rather low value of absorbed impact energy (Fig. 19(c)). The fracture surface of the sample which was welded with a welding speed of 1 m min- 1 and resulting heat input of 1.9 kJ mm−1 shows a relatively big area of the brittle fracture zone of 52%. A solidification crack is also visible on the fracture surface, which is a cause for the reduction of toughness. The area of the brittle fracture zone was measured planimetrically and the percentage data refers to a total area of 80 mm2 for an ISO sample 10 mm × 10 mm with a 2 mm deep Vnotch. The fracture surfaces of the remaining samples show a ductile brittle mixed fracture, thus have a plastic-elastic character of the fracture. The ratio of brittle to ductile fracture zone for the remaining samples are in a range of 13% to 30%. The V-notch Charpy test of laser or hybrid laser welds can sometimes have a fracture path deviation (FPD) into a softer base metal (Ohata et al. (2015)) and the following fraction can therefore be classified as ductile fracture. The problem of the occurrence of FPD is associated with the difficulties to correctly evaluate the toughness values

By using a undermatched filler wire with respect to tensile strength, the tensile specimens taken from welds performed with high heat input break in the WM and do not meet the tensile strength requirements. A wide weld seam with a high heat input means a high dilution of the WM with the BM, which indicates that the tensile strength of the weld joint is determined by the undermatched filler wire. For lower heat inputs and resulting thin weld seams, the microstructure is fine grained and the required minimum tensile strength of 915 MPa is achieved. According to the standards API 5 L and EN ISO 3183, a fracture in the WM is acceptable if the requested minimum tensile strength of 915 MPa is reached. The mechanical properties given by the filler wire manufacturer, see Table 2, refer to the tensile strength of the pure weld metal at typically higher heat input and without regard to the dilution effect. The stress strain diagram for tested tensile samples are shown in Fig. 17. It is recognizable, that the tensile strength rises with increasing welding velocity and decreasing heat input. One of the samples has a characteristic course (blue continuous line) and drop very rapidly by increase of the strain. Thereupon, the fracture surface of this sample was examined with the scanning electron microscopy (SEM), which is shown in Fig. 18. The study by Pandey et al. (2017) on the analysis of the tensile fracture behavior of high-strength steel Grade 91 shows that the fracture morphology strongly depends on the microstructure and the presence of secondary phase particles. For the specimen shown in Fig. 18 the fractured surface mainly indicates the presence of transgranular ductile dimples in the big area around the hot crack surface. The small amount of cleavage facets could be also found in vicinity of hot crack region. SEM examinations, conducted on the crack surface (Fig. 18(b)) revealed that the crack type exhibits a dendritic structure, which is a characteristic typical of solidification cracks. The solidification crack surface seen on the SEM images is rounded and smooth. Furthermore, holes on the examined surface can be indicated, what is a sign for an insufficient liquid flow in the interdendritic areas (Quiroz et al. (2010)). 11

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of the weld metal. The fracture path images of notched-bar impact test results are shown in Fig. 20. It can be seen, that the fracture path propagates mainly through the weld metal without PDF into a base metal.

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4. Conclusions Experimental investigations on single-pass HLAW of 20 mm thick pipeline steel of grade X120 were made with five different heat inputs. The experiments were carried out with a welding velocity in a range of 0.5 m min−1 and 1.5 m min−1. The gravity drop-outs were prevented by using a novel technique of bath support based on external oscillating magnetic field. The main findings of the work are listed as follows:

• HLAW of 20 mm thick plates of grade X120 could be welded in • • •

single-pass without root excess weld metal, where the AC magnet was operated with an AC power between 1.7 kW and 2.1 kW at an oscillating frequency of 1.2 kHz; as opposed to conventional HLAW parameters with resulting short cooling times t8/5 of approx. 1 s, the welding thermal cycle can be shifted towards lower cooling rates by using an electromagnetic weld pool support; in contrast to the specifications provided by the steel manufacturers, satisfactory results regarding the tensile strength and Charpy toughness could be achieved with a cooling time of 2 s to 3 s; higher cooling times t8/5 lead to softening, grain coarsening and loss of strength due to the higher heat input.

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