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Weldability of cast CoCrFeMnNi high-entropy alloys using various filler metals for cryogenic applications
Journal Pre-proof Weldability of cast CoCrFeMnNi high-entropy alloys using various filler metals for cryogenic applications Hyunbin Nam, Sangwon Park, Nokeun Park, Youngsang Na, Hyoungseop Kim, SunJoon Yoo, Young-Hoon Moon, Namhyun Kang PII:
S0925-8388(19)34524-4
DOI:
https://doi.org/10.1016/j.jallcom.2019.153278
Reference:
JALCOM 153278
To appear in:
Journal of Alloys and Compounds
Received Date: 28 August 2019 Revised Date:
2 December 2019
Accepted Date: 3 December 2019
Please cite this article as: H. Nam, S. Park, N. Park, Y. Na, H. Kim, S.-J. Yoo, Y.-H. Moon, N. Kang, Weldability of cast CoCrFeMnNi high-entropy alloys using various filler metals for cryogenic applications, Journal of Alloys and Compounds (2020), doi: https://doi.org/10.1016/j.jallcom.2019.153278. This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. 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. © 2019 Published by Elsevier B.V.
HEA
STS 308L
2 mm
FCC single phase
2 mm
&
FCC single phase
The HEA welds with STS 308L fillers had no BCC (δ-ferrite) → due to dilution of FCC stabilising elements provided from the BM. Significant deformation twins and dislocation formed at 77 K
Cryogenic tensile properties of the welds were superior to those at room temperature → due to formation of deformation twins and dislocations at cryogenic temp.
10 µm
298 K
Commercial STS 308L fillers can be used to CoCrFeMnNi welding for cryogenic applications.
10 µm
77 K
Weldability of cast CoCrFeMnNi high-entropy alloys using various filler metals for cryogenic applications
Hyunbin Nama, Sangwon Parka, Nokeun Parkb, Youngsang Nac, Hyoungseop Kimd, Sun-Joon Yooe, Young-Hoon Moonf, and Namhyun Kanga,*
a
Department of Materials Science and Engineering, Pusan National University, Busan 46241, Korea
b
School of Materials Science and Engineering, Yeungnam University, Gyeongsan 38541, Korea
c
Titanium Department, Korea Institute of Materials Science, Gyeongnam 51508, Korea
d
Department of Materials Science and Engineering, Pohang University of Science and Technology, Pohang 37673, Korea
e
I-Dasan LINC+ Educational Development Institute, Dankook University, Yongin 16890, Korea
f
School of Mechanical Engineering, Pusan National University, Busan 46241, Korea
*Corresponding author: Namhyun Kang ([email protected], +82-51-510-3027)
Abstract This study investigates the effect of stainless steel (STS) 308L and high-entropy alloy (HEA) filler metals on gas tungsten arc (GTA) welding. The weldability of the cast CoCrFeMnNi HEAs was determined based on the microstructural and mechanical properties of the welds. The cast HEA exhibited larger dendrite packets than the weld metals (WMs). The hardness in the WM was superior compared with that in the base metal (BM). The WM using STS 308L exhibited a fully face-centred cubic (FCC) structure with no indication of δ-ferrite and lower hardness than that using HEA filler. The GTA welds using both fillers showed tensile
properties comparable to the cast BM at 298 K, and the tensile fracture of the transverse welds occurred near the cast BM. The cryogenic tensile properties in the GTA welds were superior compared with the room-temperature property due to the significant formation of deformation twins and high dislocation density at 77 K. This was probably due to the decrease in the stacking fault energy at the cryogenic temperature compared with that at the room temperature. Therefore, it is possible to apply commercial STS 308L filler metal for the CoCrFeMnNi HEA in cryogenic applications.
Keywords: metals and alloys, liquid-solid reactions, crystal structure, microstructure, metallography, mechanical properties
1. Introduction Owing to recent the demand for liquefied natural gas (LNG), the need for cryogenic equipment for LNG storage and transport has rapidly increased [1,2]. Furthermore, for the mining of LNG in the polar regions, high-strength and low-temperature toughness of structural steel is required [3,4]. For cryogenic applications, austenitic stainless and 9%-Ni steels are commonly used [5–7]. However, the weld metal (WM) of austenitic stainless and 9%-Ni steels is susceptible to hot cracking and nickel-based consumables need to be employed [8–10]. Specifically, the welding properties of these steels deteriorate significantly at cryogenic temperatures. Therefore, recently, high-entropy alloys (HEAs) with excellent cryogenic properties have received significant attention [11–13]. For the application of HEAs as structural material, weldability assessments are crucial [14]. However, until now, studies on the weldability assessment of HEAs have been limited to the
range of low heat input welding [15–24]. Arc welding of high heat input has mainly been used in various industrial fields where HEAs can be applied. Recently, studies on the weldability assessment of HEAs applying gas tungsten arc (GTA) welding have been reported [25,26]. For the welding of HEAs using GTA, the filler metals need to be properly selected, and commercial filler metals are needed. This study investigates the GTA weldability using the commercial stainless steel (STS) 308L and CoCrFeMnNi filler metals on cast HEAs.
2. Materials and experimental procedures Ingots with equiatomic HEA compositions (Co0.2Cr0.2Fe0.2Mn0.2Ni0.2) were produced by vacuum induction melting. Cast HEA plates were prepared by homogenising the ingot at 1100 °C for 24 h, followed by air cooling. The parts of the ingot that were unaffected by shrinkage defect were sliced to produce cast HEA plates with a thickness of 3 mm. The plates of the base metal (BM) had a dimension of 55 mm (W) × 100 mm (L) × 3 mm (T). The cast HEA plates were cleaned with acetone before welding. A schematic of the V-groove for the GTA welding is shown in Fig. 1. The cast HEA BM had a groove angle of 30° for feeding the filler metals, and the root face length and root gap were 0.5 and 1 mm, respectively. And the filler metals were introduced to fill space of V-groove. Two types of filler metals were used with the HEA and STS 308L wires. The HEA filler metal had the same composition as the BM (Co0.2Cr0.2Fe0.2Mn0.2Ni0.2) and the STS 308L filler metal had the nominal composition of 20 wt% Cr, 10 wt% Ni, 1.6 wt% Mn, 0.4 wt% Si, 0.11 wt% (Mo + C), and balanced Fe. Single-pass GTA welding was applied using the following conditions: welding current of 90 A, welding velocity of 12.5 cm/min, wire diameter of 2.1 Φ, and Ar shielding gas. The GTA welding conditions were chosen such that a full penetrating weld was achieved.
Fig. 1. Schematics of specimens for GTA welding: (a) cast BM and (b) V-groove.
The cross-sectional macro- and microstructure of the BM transverse welds were observed using scanning electron microscopy (SEM) with backscattered electron (BSE) analysis. The crystal structures of the BM and WM were characterised by X-ray diffraction (XRD) at a scan speed of 2° min−1 in a range of 20–90°, with a voltage of 40 kV and current 30 mA using Cu Kα radiation. In addition, electron backscattered diffraction (EBSD) was employed to observe the microstructural properties, such as phase contents, grain size, and dendrite growth behaviour, in the BM and WM. To quantitatively analyse the alloy composition in each region of the transverse welds, electron probe microanalysis (EPMA) was performed. Vickers hardness measurements in the welds were carried out using a load of 500 gf (4.903 N) with 10-s loading cycles. After flattening the upper/lower beads of the welds, tensile tests were performed on a sub-size based on ASTM E8. The tensile tests were carried out at room (298 K) and cryogenic (77 K) temperatures at a strain rate of 8.3 × 10− 4 s− 1.
3. Results and discussion 3.1 GTA weldability of HEAs using various filler metals Fig. 2 shows the shape of the weld pool of the GTA welds produced by HEA and STS 308L
filler metals. Good weldability with no macro-defects such as internal pores and cracks were observed for all welds. Full penetration welds were obtained regardless of the filler metals. The upper and lower beads were nearly flat in the welds produced by various filler metals. Furthermore, the fusion lines (blue-dotted lines) of the welds with HEA filler metal were difficult to observe, while those in the welds with STS 308L filler metal was clearly visible.
Fig. 2. Shapes of the weld pools of the GTA welds produced by various filler metals: (a) HEA and (b) STS 308L. Red- and blue-dotted lines indicate the location of the quantitative EPMA analysis and fusion lines, respectively.
3.2 Microstructural behaviour of GTA welds using various filler metals Fig. 3 shows the crystal structure obtained by XRD for the BM and GTA welds. Regardless of the filler metals used with HEA and STS 308L, the observed face-centred cubic (FCC) solid solution phase with diffraction peaks (2θ = 43.4°, 51.6°, and 74.7°) was the same as that of the BM. The XRD patterns of the STS 304 weld produced with the STS 308L filler metal exhibit FCC and body-centred cubic (BCC) diffraction peaks. Normally, the welds of austenitic STS contain a small amount of δ-ferrite to prevent solidification cracking during welding [27–29]. However, the cast HEA weld produced using the STS 308L filler metal exhibited only FCC and no BCC diffraction peaks.
Fig. 3. XRD patterns of cast HEA BM and GTA welds using HEA and STS 308L filler metals.
The GTA welds using filler metal involved dilution to produce the WM. To confirm the elemental distribution across the weld produced using various filler metals (HEA and STS 308L), quantitative EPMA was performed for the BM across the weld centreline following the red-dotted lines shown in Fig. 2. Fig. 4 shows the line analysis of the chemical compositions in each weld. The various components, such as Co, Cr, Fe, Mn and Ni, exhibited constant composition in the BM, heat-affected zone (HAZ), and WM (Fig. 4(a)). This was due to the same composition of the BM and HEA filler metal, and it could be associated with the undefined fusion line in the welds produced using HEA filler metal (Fig. 2(a)). Furthermore, it provided reasonable evidence on the negligible vaporisation of Mn with a low boiling point during the GTA welding. The application of STS 308L filler metal on the HEA induced dilution, and the composition of the BM and WM with 308L filler metal were significantly different, as shown in Fig. 4(b). The Fe component in the WM was larger than that in the BM, and the Ni, Mn, and Co components in the WM were lower than those in the BM. However, the Cr composition was mostly constant between the BM and WM. It should be noted that, although not shown in Fig. 4, small amount of Si, C, Mo, and Nb components diluted from the 308L filler metal were detected in the WM. Due to the significantly different composition of the HEA BM and
STS 308L filler metal, the fusion line was clearly visible in the welds (Fig. 2(b)). Furthermore, the δ-ferrites of the BCC peaks were not observed in the weld using 308L filler metal (Fig. 3(a)). This was because the austenite-stabilising elements, such as Ni and Mn, of the HEA BM were diluted to the STS WM, therefore decreasing the Fe content and increasing the Ni and Mn contents in the welds using STS 308L filler metal (Fig. 4(b)). By applying the composition measured in the weld centreline (Fig. 4(b)), the Ni and Cr equivalents can be calculated as follows: Creq = (Cr) + 2(Si) + 1.5(Mo) + 5(V) + 5.5(Al) + 1.75(Nb) + 1.5(Ti) + 0.756(W) = 19.1, Nieq = (Ni) + (Co) + 0.5(Mn) + 0.3(Cu) + 25(N) +30(C) = 39.5.
(Eq. 1) (Eq. 2)
As the WM using the STS 308L filler metal exhibited a high Nieq, it was confirmed from the Schaeffler diagram that the STS WM consisted of austenite single phase, based on the relationship between the Creq and Nieq [30–32]. In addition, Fe-rich HEAs with compositions comparable to that of STS WM were reported to have a FCC single phase [33–35]. Therefore, the HEA weld produced using STS 308L filler metal exhibited no δ-ferrite, as shown in Fig. 3(b).
Fig. 4. Elemental behaviour of the GTA welds produced using various filler metals: (a) HEA and (b) STS 308L.
Fig. 5 shows the microstructural behaviour near the area of the HAZ and WM centreline for the welds produced using HEA and STS 308L filler metals. The cast BM exhibited coarse grains of 1.0 ± 0.2 mm (Fig. 5(a)) and the HAZ mostly exhibited the same grain size with that of the BM. Columnar grains growing epitaxially from the fusion line to the WM centreline were observed for the welds produced using the HEA filler metal (Fig. 5(b)). The epitaxial growth of the columnar grains can also be associated with the constant composition (Fig. 4(a)) and undefined fusion line of the weld produced using the HEA filler metal (Fig. 2(a)). However, the clearly visible weld–fusion line produced using the 308L filler metal can be associated with the incomplete epitaxial growth of the columnar grains (Fig. 5(c)). This result is in correlation with the significantly different composition of the weld produced using the 308L filler metal (Figs. 2(b), 4(b), and 5(c)). Towards the centreline, the columnar grains near the fusion line were transformed to equiaxed grains. The equiaxed grains near the centreline of each WM are typical microstructures arising for high heat input welding such as GTA welding [25,26]. The dendrite packet sizes near the WM centreline were nearly the same, regardless of the HEA (270 µm) and STS 308L (285 µm) filler metals.
Fig. 5. Microstructural behaviour of (a) cast BM and near the fusion line and WM centreline using various filler metals: (b and d) HEA and (c and e) STS 308L.
Fig. 6(a) shows the elemental distribution map near the fusion line in the welds using HEA filler metal. The compositional micro-segregations in the WM showed the same pattern as those in the BM [25]; the dendritic microstructure consisting of Co-, Cr-, and Fe-rich dendrite cores and Mn- and Ni-rich interdendritic regions. The elemental distributions near the fusion line using the 308L filler metal are shown in Fig. 6(b). The main elements, such as Co, Cr, Fe, Mn, and Ni, show a different degree of segregation for the WMs using HEA and 308L filler metals. Specifically, the compositional segregations of Cr between the dendrite cores and the interdendritic regions in the STS WM disappeared and those of Co diminished in the WM using 308L filler metal. The STS 308L filler metal contains a larger amount of Fe (~65 wt.%) and Cr (~ 20 wt.%) than the cast HEA, and the main components of the STS 308L filler are diluted in the cast HEA. The high cooling rates and the associated higher undercooling during welding can result in reduced diffusion distances and micro-segregation. However, the origin of the different micro-segregations of Cr and Co in the WMs using the STS 308L filler has not been fully understood in this work and further studies to understand this phenomenon are in progress.
Fig. 6. Elemental distribution map near the fusion line in the WM using various filler metals: (a) HEA and (b) STS 308L.
3.3 Mechanical properties of GTA welds using various filler metals 3.3.1 Hardness behaviour of the WMs at room temperature
Fig. 7 shows the hardness distribution in the transverse welds produced using HEA and STS 308L filler metals. The average hardness of the cast BM was approximately 132 ± 1 Hv0.5. All WMs showed higher average hardness than the cast BM. This significant difference in the hardness of the BM and WMs was primarily due to the grain size of the BM, which was approximately forty times larger than the dendrite packet size of each WM. Furthermore, the average hardness in the WM produced using STS 308L filler metal (150 ± 1 Hv0.5) was lower than that using HEA filler metal (165 ± 1 Hv0.5). These results can be partly due to the variation of the grain size and dendrite packet size in each WM using HEA (270 µm) and STS 308L (285 µm). However, the main effect is due to the significant difference in the constituent components of each WM with HEA and STS 308L filler metal. As shown in Figs. 4 and 6, the HEA and STS WMs exhibited a different distribution and degree of segregation of the main components. The Co, Ni, and Mn components in the WM using the HEA filler were larger than that in the WM using the STS 308L filler metal. Amount and number of constituent components should be associated to the inclusion and the lattice distortion of the WMs [36,37]. Therefore, the average hardness of the HEA weld having a large amount and type of constituent components indicated a large hardness than that of the STS 308L weld.
Fig. 7. Hardness distribution of the GTA welds produced using HEA and STS 308L filler metals.
3.3.2 Tensile properties of welds using various filler metals tested at various temperatures
Figs. 8(a) and 8(b) show typical tensile stress–strain curves and fracture positions of cast BM and GTA welds produced using HEA and STS 308L filler metals. The tensile tests were conducted at different temperatures: 298 and 77 K. The yield strength (YS), tensile strength (TS), and elongation of the cast BM increased by approximately 52%, 44%, and 24%, respectively, as the temperature was decreased from 298 to 77 K. Regardless of the filler metals, the cryogenic tensile properties of all welds were superior compared with their tensile properties at 298 K. The YS and TS of all welds increased by ~100% and the elongation increased by ~20% when the temperature was decreased from 298 to 77 K. Regardless of the testing temperature and filler metals used, the YS and TS of each weld were mostly the same and they were larger than those of the cast BM. To clearly distinguish the tensile fracture position, the red lines are indicated to show the fusion lines of the welds (Fig. 8b). Among the welds, tensile fractures occurred only in the cast BM (red-dotted boxes). Among the welds, tensile fractures occurred only in the cast BM. Therefore, the cast HEA welds produced using HEA and STS 308L filler metals exhibited superior tensile properties compared with those of the cast BM. Furthermore, the STS 308L filler metal exhibited reasonable tensile property for the GTA welding of the cast CoCrFeMnNi HEA.
Fig. 8. Tensile properties of GTA welds with HEA and STS 308L filler metals tested at 298 and 77 K: (a) tensile stress–strain curves and (b) fracture positions in transverse welds.
Fig. 9 shows the microstructure near the crack tip of the tensile fracture position in GTA welds produced using HEA filler metal at 298 and 77 K. Among the welds, tensile fractures occurred only in the cast BM (red-dotted boxes). As the tensile fractures occurred in the cast BM with a coarse grain size, a microstructure analysis near the crack tip (red-dotted boxes in Fig. 8b) was performed in the cast BM of the transverse welds. Following the tensile testing, significant deformation twins and dislocation density were observed at 298 K. Deformation twins can be observed as the coincidence site lattice (CSL) boundary of Σ3 (red lines in the magnified figure in Fig. 9(a)). The fraction of the deformation twins and dislocation density tested at 77 K was higher than that at 298 K. Therefore, the cryogenic properties of all welds were superior compared with their properties at room temperature. This was due to the predominant formation of deformation twins and high dislocation densities at cryogenic temperature [19,38–40].
Fig. 9. Microstructures near the crack tip of the tensile fracture position tested at different temperatures: (a) 298 K and (b) 77 K.
As shown in Fig. 9, among the welds fractures occurred only in the cast BM (Fig. 8(b)). Fig. 10 shows the fracture morphology of the tensile specimens tested at 298 and 77 K. Typical ductile fractures appeared on the fracture surface. The tensile specimen tested at 298 K exhibited large dimples and the non-metallic inclusions in the dimples. The non-metallic inclusions were previously investigated as CrMn oxides [18,19]. However, the fine and coarse dimples coexisted on the fracture surface at 77 K. All welds exhibited significantly smaller dimples at 77 K than those observed at 298 K. As shown in Fig. 9, the deformation twins and tangled dislocations exist in the microstructure near the fracture at 77 K. Because these twins terminate the crack propagation, fine and coarse dimples coexist on the fracture surface at 77 K [19,35].
Fig. 10. Tensile fractographs of welds tested at different temperatures: (a) 298 K and (b) 77 K
4. Conclusions In this study, GTA weldability with HEA and STS 308L filler metal was investigated for cast CoCrFeMnNi HEAs. Based on the study of the microstructures and mechanical properties of the welds produced by GTA welding at 90 A and 12.5 cm/min at room and cryogenic temperatures, the following conclusions can be drawn: 1) No macro-defects, such as internal pores or cracks, were observed for all GTA welds. 2) Despite the application of various filler metals, an FCC solid solution phase was observed in all GTA welds. In the STS 308L WM, the FCC stabilising elements, such as Ni, Co and Mn of the cast HEA BM was diluted into the STS WM and the BCC phase (δ-ferrite) in the STS WM was disappeared. Furthermore, the columnar grains were observed growing epitaxially from the fusion line to the WM in the HEA welds, and the columnar grains exhibited incomplete epitaxial growth from the fusion line in the STS 308L weld. 3) The hardness distribution of the BM and WM in each weld exhibited a sharp variation, because the dendrite packet of the WM was forty times smaller than the grain size of the cast BM. In addition, the average hardness in the STS WM was lower than that of the HEA WM.
4) Regardless of temperature of tensile test (77 and 298 K), all GTA welds and cast BM showed mostly the same tensile properties, and tensile fracture occurred only at the cast BM among all welds. In addition, the cryogenic tensile properties of all welds were superior compared with their tensile properties at 298 K. This is due to the predominant formation of deformation twins and high dislocation densities at 77 K. Therefore, it is possible to apply the commercial STS 308L filler metal as welding material for CoCrFeMnNi HEA.
Acknowledgment This work was supported by the Future Material Discovery Project of the National Research Foundation of Korea (NRF) funded by the Ministry of Science, ICT and Future Planning (MSIP) of Korea [NRF-2016M3D1A1023534], and by the National Research Foundation of Korea funded by Ministry of Science and ICT [GCRC-SOP 2011-0030013].
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Figures and Tables
Fig. 1. Schematics of specimens for GTA welding: (a) cast BM and (b) V-groove. Fig. 2. Shapes of the weld pools of the GTA welds produced by various filler metals: (a) HEA and (b) STS 308L. Red- and blue-dotted lines indicate the location of the quantitative EPMA analysis and fusion lines, respectively. Fig. 3. XRD patterns of cast HEA BM and GTA welds using HEA and STS 308L filler metals. Fig. 4. Elemental behaviour of the GTA welds produced using various filler metals: (a) HEA and (b) STS 308L. Fig. 5. Microstructural behaviour of (a) cast BM and near the fusion line and WM centreline using various filler metals: (b and d) HEA and (c and e) STS 308L. Fig. 6. Elemental distribution map near the fusion line in the WM using various filler metals: (a) HEA and (b) STS 308L. Fig. 7 Hardness distribution of the GTA welds produced using HEA and STS 308L filler metals Fig. 8. Tensile properties of GTA welds with HEA and STS 308L filler metals tested at 298 K and 77 K: (a) tensile stress-strain curves and (b) fracture positions in the transverse welds. Fig. 9. Microstructures near the crack tip of the tensile fracture position tested at different temperatures: (a) 298 K and (b) 77 K. Fig. 10. Tensile fractographs of welds tested at different temperatures: (a) 298 K and (b) 77 K.
Highlights ▪ GTA welds using the HEA and STS 308L fillers had good weldability ▪ STS 308L welds had FCC phase as same as the phase of CoCrFeMnNi HEA ▪ Tensile properties of the GTA welds at 77 K were superior to those at 298 K ▪ Deformation twins and high dislocation densities were significantly formed at 77 K ▪ STS 308L filler produced the reasonable tensile property for welding of cast HEA.
Author Contributions Section Hyunbin Nam : Conceptualization, Writing - Original Draft, Data Curation, Sangwon Park : Investigation, Visualization Nokeun Park : Investigation, Validation Youngsang Na : Resources, Methodology Hyoungseop Kim : Resources Sun-Joon Yoo : Formal analysis Young-Hoon Moon : Formal analysis Namhyun Kang : Writing - Review & Editing, Supervision
Declaration of interests ☒ The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. ☒The authors declare the following financial interests/personal relationships which may be considered as potential competing interests:
No potential conflict of interest was reported by the authors.
S0925-8388(19)34524-4
DOI:
https://doi.org/10.1016/j.jallcom.2019.153278
Reference:
JALCOM 153278
To appear in:
Journal of Alloys and Compounds
Received Date: 28 August 2019 Revised Date:
2 December 2019
Accepted Date: 3 December 2019
Please cite this article as: H. Nam, S. Park, N. Park, Y. Na, H. Kim, S.-J. Yoo, Y.-H. Moon, N. Kang, Weldability of cast CoCrFeMnNi high-entropy alloys using various filler metals for cryogenic applications, Journal of Alloys and Compounds (2020), doi: https://doi.org/10.1016/j.jallcom.2019.153278. This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. 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. © 2019 Published by Elsevier B.V.
HEA
STS 308L
2 mm
FCC single phase
2 mm
&
FCC single phase
The HEA welds with STS 308L fillers had no BCC (δ-ferrite) → due to dilution of FCC stabilising elements provided from the BM. Significant deformation twins and dislocation formed at 77 K
Cryogenic tensile properties of the welds were superior to those at room temperature → due to formation of deformation twins and dislocations at cryogenic temp.
10 µm
298 K
Commercial STS 308L fillers can be used to CoCrFeMnNi welding for cryogenic applications.
10 µm
77 K
Weldability of cast CoCrFeMnNi high-entropy alloys using various filler metals for cryogenic applications
Hyunbin Nama, Sangwon Parka, Nokeun Parkb, Youngsang Nac, Hyoungseop Kimd, Sun-Joon Yooe, Young-Hoon Moonf, and Namhyun Kanga,*
a
Department of Materials Science and Engineering, Pusan National University, Busan 46241, Korea
b
School of Materials Science and Engineering, Yeungnam University, Gyeongsan 38541, Korea
c
Titanium Department, Korea Institute of Materials Science, Gyeongnam 51508, Korea
d
Department of Materials Science and Engineering, Pohang University of Science and Technology, Pohang 37673, Korea
e
I-Dasan LINC+ Educational Development Institute, Dankook University, Yongin 16890, Korea
f
School of Mechanical Engineering, Pusan National University, Busan 46241, Korea
*Corresponding author: Namhyun Kang ([email protected], +82-51-510-3027)
Abstract This study investigates the effect of stainless steel (STS) 308L and high-entropy alloy (HEA) filler metals on gas tungsten arc (GTA) welding. The weldability of the cast CoCrFeMnNi HEAs was determined based on the microstructural and mechanical properties of the welds. The cast HEA exhibited larger dendrite packets than the weld metals (WMs). The hardness in the WM was superior compared with that in the base metal (BM). The WM using STS 308L exhibited a fully face-centred cubic (FCC) structure with no indication of δ-ferrite and lower hardness than that using HEA filler. The GTA welds using both fillers showed tensile
properties comparable to the cast BM at 298 K, and the tensile fracture of the transverse welds occurred near the cast BM. The cryogenic tensile properties in the GTA welds were superior compared with the room-temperature property due to the significant formation of deformation twins and high dislocation density at 77 K. This was probably due to the decrease in the stacking fault energy at the cryogenic temperature compared with that at the room temperature. Therefore, it is possible to apply commercial STS 308L filler metal for the CoCrFeMnNi HEA in cryogenic applications.
Keywords: metals and alloys, liquid-solid reactions, crystal structure, microstructure, metallography, mechanical properties
1. Introduction Owing to recent the demand for liquefied natural gas (LNG), the need for cryogenic equipment for LNG storage and transport has rapidly increased [1,2]. Furthermore, for the mining of LNG in the polar regions, high-strength and low-temperature toughness of structural steel is required [3,4]. For cryogenic applications, austenitic stainless and 9%-Ni steels are commonly used [5–7]. However, the weld metal (WM) of austenitic stainless and 9%-Ni steels is susceptible to hot cracking and nickel-based consumables need to be employed [8–10]. Specifically, the welding properties of these steels deteriorate significantly at cryogenic temperatures. Therefore, recently, high-entropy alloys (HEAs) with excellent cryogenic properties have received significant attention [11–13]. For the application of HEAs as structural material, weldability assessments are crucial [14]. However, until now, studies on the weldability assessment of HEAs have been limited to the
range of low heat input welding [15–24]. Arc welding of high heat input has mainly been used in various industrial fields where HEAs can be applied. Recently, studies on the weldability assessment of HEAs applying gas tungsten arc (GTA) welding have been reported [25,26]. For the welding of HEAs using GTA, the filler metals need to be properly selected, and commercial filler metals are needed. This study investigates the GTA weldability using the commercial stainless steel (STS) 308L and CoCrFeMnNi filler metals on cast HEAs.
2. Materials and experimental procedures Ingots with equiatomic HEA compositions (Co0.2Cr0.2Fe0.2Mn0.2Ni0.2) were produced by vacuum induction melting. Cast HEA plates were prepared by homogenising the ingot at 1100 °C for 24 h, followed by air cooling. The parts of the ingot that were unaffected by shrinkage defect were sliced to produce cast HEA plates with a thickness of 3 mm. The plates of the base metal (BM) had a dimension of 55 mm (W) × 100 mm (L) × 3 mm (T). The cast HEA plates were cleaned with acetone before welding. A schematic of the V-groove for the GTA welding is shown in Fig. 1. The cast HEA BM had a groove angle of 30° for feeding the filler metals, and the root face length and root gap were 0.5 and 1 mm, respectively. And the filler metals were introduced to fill space of V-groove. Two types of filler metals were used with the HEA and STS 308L wires. The HEA filler metal had the same composition as the BM (Co0.2Cr0.2Fe0.2Mn0.2Ni0.2) and the STS 308L filler metal had the nominal composition of 20 wt% Cr, 10 wt% Ni, 1.6 wt% Mn, 0.4 wt% Si, 0.11 wt% (Mo + C), and balanced Fe. Single-pass GTA welding was applied using the following conditions: welding current of 90 A, welding velocity of 12.5 cm/min, wire diameter of 2.1 Φ, and Ar shielding gas. The GTA welding conditions were chosen such that a full penetrating weld was achieved.
Fig. 1. Schematics of specimens for GTA welding: (a) cast BM and (b) V-groove.
The cross-sectional macro- and microstructure of the BM transverse welds were observed using scanning electron microscopy (SEM) with backscattered electron (BSE) analysis. The crystal structures of the BM and WM were characterised by X-ray diffraction (XRD) at a scan speed of 2° min−1 in a range of 20–90°, with a voltage of 40 kV and current 30 mA using Cu Kα radiation. In addition, electron backscattered diffraction (EBSD) was employed to observe the microstructural properties, such as phase contents, grain size, and dendrite growth behaviour, in the BM and WM. To quantitatively analyse the alloy composition in each region of the transverse welds, electron probe microanalysis (EPMA) was performed. Vickers hardness measurements in the welds were carried out using a load of 500 gf (4.903 N) with 10-s loading cycles. After flattening the upper/lower beads of the welds, tensile tests were performed on a sub-size based on ASTM E8. The tensile tests were carried out at room (298 K) and cryogenic (77 K) temperatures at a strain rate of 8.3 × 10− 4 s− 1.
3. Results and discussion 3.1 GTA weldability of HEAs using various filler metals Fig. 2 shows the shape of the weld pool of the GTA welds produced by HEA and STS 308L
filler metals. Good weldability with no macro-defects such as internal pores and cracks were observed for all welds. Full penetration welds were obtained regardless of the filler metals. The upper and lower beads were nearly flat in the welds produced by various filler metals. Furthermore, the fusion lines (blue-dotted lines) of the welds with HEA filler metal were difficult to observe, while those in the welds with STS 308L filler metal was clearly visible.
Fig. 2. Shapes of the weld pools of the GTA welds produced by various filler metals: (a) HEA and (b) STS 308L. Red- and blue-dotted lines indicate the location of the quantitative EPMA analysis and fusion lines, respectively.
3.2 Microstructural behaviour of GTA welds using various filler metals Fig. 3 shows the crystal structure obtained by XRD for the BM and GTA welds. Regardless of the filler metals used with HEA and STS 308L, the observed face-centred cubic (FCC) solid solution phase with diffraction peaks (2θ = 43.4°, 51.6°, and 74.7°) was the same as that of the BM. The XRD patterns of the STS 304 weld produced with the STS 308L filler metal exhibit FCC and body-centred cubic (BCC) diffraction peaks. Normally, the welds of austenitic STS contain a small amount of δ-ferrite to prevent solidification cracking during welding [27–29]. However, the cast HEA weld produced using the STS 308L filler metal exhibited only FCC and no BCC diffraction peaks.
Fig. 3. XRD patterns of cast HEA BM and GTA welds using HEA and STS 308L filler metals.
The GTA welds using filler metal involved dilution to produce the WM. To confirm the elemental distribution across the weld produced using various filler metals (HEA and STS 308L), quantitative EPMA was performed for the BM across the weld centreline following the red-dotted lines shown in Fig. 2. Fig. 4 shows the line analysis of the chemical compositions in each weld. The various components, such as Co, Cr, Fe, Mn and Ni, exhibited constant composition in the BM, heat-affected zone (HAZ), and WM (Fig. 4(a)). This was due to the same composition of the BM and HEA filler metal, and it could be associated with the undefined fusion line in the welds produced using HEA filler metal (Fig. 2(a)). Furthermore, it provided reasonable evidence on the negligible vaporisation of Mn with a low boiling point during the GTA welding. The application of STS 308L filler metal on the HEA induced dilution, and the composition of the BM and WM with 308L filler metal were significantly different, as shown in Fig. 4(b). The Fe component in the WM was larger than that in the BM, and the Ni, Mn, and Co components in the WM were lower than those in the BM. However, the Cr composition was mostly constant between the BM and WM. It should be noted that, although not shown in Fig. 4, small amount of Si, C, Mo, and Nb components diluted from the 308L filler metal were detected in the WM. Due to the significantly different composition of the HEA BM and
STS 308L filler metal, the fusion line was clearly visible in the welds (Fig. 2(b)). Furthermore, the δ-ferrites of the BCC peaks were not observed in the weld using 308L filler metal (Fig. 3(a)). This was because the austenite-stabilising elements, such as Ni and Mn, of the HEA BM were diluted to the STS WM, therefore decreasing the Fe content and increasing the Ni and Mn contents in the welds using STS 308L filler metal (Fig. 4(b)). By applying the composition measured in the weld centreline (Fig. 4(b)), the Ni and Cr equivalents can be calculated as follows: Creq = (Cr) + 2(Si) + 1.5(Mo) + 5(V) + 5.5(Al) + 1.75(Nb) + 1.5(Ti) + 0.756(W) = 19.1, Nieq = (Ni) + (Co) + 0.5(Mn) + 0.3(Cu) + 25(N) +30(C) = 39.5.
(Eq. 1) (Eq. 2)
As the WM using the STS 308L filler metal exhibited a high Nieq, it was confirmed from the Schaeffler diagram that the STS WM consisted of austenite single phase, based on the relationship between the Creq and Nieq [30–32]. In addition, Fe-rich HEAs with compositions comparable to that of STS WM were reported to have a FCC single phase [33–35]. Therefore, the HEA weld produced using STS 308L filler metal exhibited no δ-ferrite, as shown in Fig. 3(b).
Fig. 4. Elemental behaviour of the GTA welds produced using various filler metals: (a) HEA and (b) STS 308L.
Fig. 5 shows the microstructural behaviour near the area of the HAZ and WM centreline for the welds produced using HEA and STS 308L filler metals. The cast BM exhibited coarse grains of 1.0 ± 0.2 mm (Fig. 5(a)) and the HAZ mostly exhibited the same grain size with that of the BM. Columnar grains growing epitaxially from the fusion line to the WM centreline were observed for the welds produced using the HEA filler metal (Fig. 5(b)). The epitaxial growth of the columnar grains can also be associated with the constant composition (Fig. 4(a)) and undefined fusion line of the weld produced using the HEA filler metal (Fig. 2(a)). However, the clearly visible weld–fusion line produced using the 308L filler metal can be associated with the incomplete epitaxial growth of the columnar grains (Fig. 5(c)). This result is in correlation with the significantly different composition of the weld produced using the 308L filler metal (Figs. 2(b), 4(b), and 5(c)). Towards the centreline, the columnar grains near the fusion line were transformed to equiaxed grains. The equiaxed grains near the centreline of each WM are typical microstructures arising for high heat input welding such as GTA welding [25,26]. The dendrite packet sizes near the WM centreline were nearly the same, regardless of the HEA (270 µm) and STS 308L (285 µm) filler metals.
Fig. 5. Microstructural behaviour of (a) cast BM and near the fusion line and WM centreline using various filler metals: (b and d) HEA and (c and e) STS 308L.
Fig. 6(a) shows the elemental distribution map near the fusion line in the welds using HEA filler metal. The compositional micro-segregations in the WM showed the same pattern as those in the BM [25]; the dendritic microstructure consisting of Co-, Cr-, and Fe-rich dendrite cores and Mn- and Ni-rich interdendritic regions. The elemental distributions near the fusion line using the 308L filler metal are shown in Fig. 6(b). The main elements, such as Co, Cr, Fe, Mn, and Ni, show a different degree of segregation for the WMs using HEA and 308L filler metals. Specifically, the compositional segregations of Cr between the dendrite cores and the interdendritic regions in the STS WM disappeared and those of Co diminished in the WM using 308L filler metal. The STS 308L filler metal contains a larger amount of Fe (~65 wt.%) and Cr (~ 20 wt.%) than the cast HEA, and the main components of the STS 308L filler are diluted in the cast HEA. The high cooling rates and the associated higher undercooling during welding can result in reduced diffusion distances and micro-segregation. However, the origin of the different micro-segregations of Cr and Co in the WMs using the STS 308L filler has not been fully understood in this work and further studies to understand this phenomenon are in progress.
Fig. 6. Elemental distribution map near the fusion line in the WM using various filler metals: (a) HEA and (b) STS 308L.
3.3 Mechanical properties of GTA welds using various filler metals 3.3.1 Hardness behaviour of the WMs at room temperature
Fig. 7 shows the hardness distribution in the transverse welds produced using HEA and STS 308L filler metals. The average hardness of the cast BM was approximately 132 ± 1 Hv0.5. All WMs showed higher average hardness than the cast BM. This significant difference in the hardness of the BM and WMs was primarily due to the grain size of the BM, which was approximately forty times larger than the dendrite packet size of each WM. Furthermore, the average hardness in the WM produced using STS 308L filler metal (150 ± 1 Hv0.5) was lower than that using HEA filler metal (165 ± 1 Hv0.5). These results can be partly due to the variation of the grain size and dendrite packet size in each WM using HEA (270 µm) and STS 308L (285 µm). However, the main effect is due to the significant difference in the constituent components of each WM with HEA and STS 308L filler metal. As shown in Figs. 4 and 6, the HEA and STS WMs exhibited a different distribution and degree of segregation of the main components. The Co, Ni, and Mn components in the WM using the HEA filler were larger than that in the WM using the STS 308L filler metal. Amount and number of constituent components should be associated to the inclusion and the lattice distortion of the WMs [36,37]. Therefore, the average hardness of the HEA weld having a large amount and type of constituent components indicated a large hardness than that of the STS 308L weld.
Fig. 7. Hardness distribution of the GTA welds produced using HEA and STS 308L filler metals.
3.3.2 Tensile properties of welds using various filler metals tested at various temperatures
Figs. 8(a) and 8(b) show typical tensile stress–strain curves and fracture positions of cast BM and GTA welds produced using HEA and STS 308L filler metals. The tensile tests were conducted at different temperatures: 298 and 77 K. The yield strength (YS), tensile strength (TS), and elongation of the cast BM increased by approximately 52%, 44%, and 24%, respectively, as the temperature was decreased from 298 to 77 K. Regardless of the filler metals, the cryogenic tensile properties of all welds were superior compared with their tensile properties at 298 K. The YS and TS of all welds increased by ~100% and the elongation increased by ~20% when the temperature was decreased from 298 to 77 K. Regardless of the testing temperature and filler metals used, the YS and TS of each weld were mostly the same and they were larger than those of the cast BM. To clearly distinguish the tensile fracture position, the red lines are indicated to show the fusion lines of the welds (Fig. 8b). Among the welds, tensile fractures occurred only in the cast BM (red-dotted boxes). Among the welds, tensile fractures occurred only in the cast BM. Therefore, the cast HEA welds produced using HEA and STS 308L filler metals exhibited superior tensile properties compared with those of the cast BM. Furthermore, the STS 308L filler metal exhibited reasonable tensile property for the GTA welding of the cast CoCrFeMnNi HEA.
Fig. 8. Tensile properties of GTA welds with HEA and STS 308L filler metals tested at 298 and 77 K: (a) tensile stress–strain curves and (b) fracture positions in transverse welds.
Fig. 9 shows the microstructure near the crack tip of the tensile fracture position in GTA welds produced using HEA filler metal at 298 and 77 K. Among the welds, tensile fractures occurred only in the cast BM (red-dotted boxes). As the tensile fractures occurred in the cast BM with a coarse grain size, a microstructure analysis near the crack tip (red-dotted boxes in Fig. 8b) was performed in the cast BM of the transverse welds. Following the tensile testing, significant deformation twins and dislocation density were observed at 298 K. Deformation twins can be observed as the coincidence site lattice (CSL) boundary of Σ3 (red lines in the magnified figure in Fig. 9(a)). The fraction of the deformation twins and dislocation density tested at 77 K was higher than that at 298 K. Therefore, the cryogenic properties of all welds were superior compared with their properties at room temperature. This was due to the predominant formation of deformation twins and high dislocation densities at cryogenic temperature [19,38–40].
Fig. 9. Microstructures near the crack tip of the tensile fracture position tested at different temperatures: (a) 298 K and (b) 77 K.
As shown in Fig. 9, among the welds fractures occurred only in the cast BM (Fig. 8(b)). Fig. 10 shows the fracture morphology of the tensile specimens tested at 298 and 77 K. Typical ductile fractures appeared on the fracture surface. The tensile specimen tested at 298 K exhibited large dimples and the non-metallic inclusions in the dimples. The non-metallic inclusions were previously investigated as CrMn oxides [18,19]. However, the fine and coarse dimples coexisted on the fracture surface at 77 K. All welds exhibited significantly smaller dimples at 77 K than those observed at 298 K. As shown in Fig. 9, the deformation twins and tangled dislocations exist in the microstructure near the fracture at 77 K. Because these twins terminate the crack propagation, fine and coarse dimples coexist on the fracture surface at 77 K [19,35].
Fig. 10. Tensile fractographs of welds tested at different temperatures: (a) 298 K and (b) 77 K
4. Conclusions In this study, GTA weldability with HEA and STS 308L filler metal was investigated for cast CoCrFeMnNi HEAs. Based on the study of the microstructures and mechanical properties of the welds produced by GTA welding at 90 A and 12.5 cm/min at room and cryogenic temperatures, the following conclusions can be drawn: 1) No macro-defects, such as internal pores or cracks, were observed for all GTA welds. 2) Despite the application of various filler metals, an FCC solid solution phase was observed in all GTA welds. In the STS 308L WM, the FCC stabilising elements, such as Ni, Co and Mn of the cast HEA BM was diluted into the STS WM and the BCC phase (δ-ferrite) in the STS WM was disappeared. Furthermore, the columnar grains were observed growing epitaxially from the fusion line to the WM in the HEA welds, and the columnar grains exhibited incomplete epitaxial growth from the fusion line in the STS 308L weld. 3) The hardness distribution of the BM and WM in each weld exhibited a sharp variation, because the dendrite packet of the WM was forty times smaller than the grain size of the cast BM. In addition, the average hardness in the STS WM was lower than that of the HEA WM.
4) Regardless of temperature of tensile test (77 and 298 K), all GTA welds and cast BM showed mostly the same tensile properties, and tensile fracture occurred only at the cast BM among all welds. In addition, the cryogenic tensile properties of all welds were superior compared with their tensile properties at 298 K. This is due to the predominant formation of deformation twins and high dislocation densities at 77 K. Therefore, it is possible to apply the commercial STS 308L filler metal as welding material for CoCrFeMnNi HEA.
Acknowledgment This work was supported by the Future Material Discovery Project of the National Research Foundation of Korea (NRF) funded by the Ministry of Science, ICT and Future Planning (MSIP) of Korea [NRF-2016M3D1A1023534], and by the National Research Foundation of Korea funded by Ministry of Science and ICT [GCRC-SOP 2011-0030013].
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Figures and Tables
Fig. 1. Schematics of specimens for GTA welding: (a) cast BM and (b) V-groove. Fig. 2. Shapes of the weld pools of the GTA welds produced by various filler metals: (a) HEA and (b) STS 308L. Red- and blue-dotted lines indicate the location of the quantitative EPMA analysis and fusion lines, respectively. Fig. 3. XRD patterns of cast HEA BM and GTA welds using HEA and STS 308L filler metals. Fig. 4. Elemental behaviour of the GTA welds produced using various filler metals: (a) HEA and (b) STS 308L. Fig. 5. Microstructural behaviour of (a) cast BM and near the fusion line and WM centreline using various filler metals: (b and d) HEA and (c and e) STS 308L. Fig. 6. Elemental distribution map near the fusion line in the WM using various filler metals: (a) HEA and (b) STS 308L. Fig. 7 Hardness distribution of the GTA welds produced using HEA and STS 308L filler metals Fig. 8. Tensile properties of GTA welds with HEA and STS 308L filler metals tested at 298 K and 77 K: (a) tensile stress-strain curves and (b) fracture positions in the transverse welds. Fig. 9. Microstructures near the crack tip of the tensile fracture position tested at different temperatures: (a) 298 K and (b) 77 K. Fig. 10. Tensile fractographs of welds tested at different temperatures: (a) 298 K and (b) 77 K.
Highlights ▪ GTA welds using the HEA and STS 308L fillers had good weldability ▪ STS 308L welds had FCC phase as same as the phase of CoCrFeMnNi HEA ▪ Tensile properties of the GTA welds at 77 K were superior to those at 298 K ▪ Deformation twins and high dislocation densities were significantly formed at 77 K ▪ STS 308L filler produced the reasonable tensile property for welding of cast HEA.
Author Contributions Section Hyunbin Nam : Conceptualization, Writing - Original Draft, Data Curation, Sangwon Park : Investigation, Visualization Nokeun Park : Investigation, Validation Youngsang Na : Resources, Methodology Hyoungseop Kim : Resources Sun-Joon Yoo : Formal analysis Young-Hoon Moon : Formal analysis Namhyun Kang : Writing - Review & Editing, Supervision
Declaration of interests ☒ The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. ☒The authors declare the following financial interests/personal relationships which may be considered as potential competing interests:
No potential conflict of interest was reported by the authors.