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Microstructure and local strains in GH3535 alloy heat affected zone and their influence on the mechanical properties
Materials Science & Engineering A 699 (2017) 48–54
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Microstructure and local strains in GH3535 alloy heat affected zone and their influence on the mechanical properties
MARK ⁎
Shuangjian Chena,b, D.K.L. Tsanga, Li Jianga,b, Kun Yua,b, Chaowen Lia, Zhong Lia, Zhijun Lia, , Xingtai Zhoua, Jianguo Yangc a b c
Center for Thorium Molten Salt Reactor System, Shanghai Institute of Applied Physics, Chinese Academy of Sciences, Shanghai 201800, PR China University of Chinese Academy of Sciences, Beijing 100049, PR China Zhejiang University of Technology, Hangzhou 310014, PR China
A R T I C L E I N F O
A B S T R A C T
Keywords: EBSD Twinning Dislocations Nickel based superalloys Welding Hardening
GH3535 alloy plates were welded by Gas Tungsten Arc Welding in order to study the evolution of microstructure and mechanical properties in the heat-affected zone (HAZ). Our results suggest that welding thermal cycles induced the morphology evolution of M6C carbides from block to eutectic near the fusion line in the HAZ. Electron backscatter diffraction (EBSD) results show that significant amounts of plastic strains occurred in the HAZ after welding. In addition, local coherent twin boundaries (Σ3) and dislocations were observed to decrease with the distance from the fusion line. Mechanical tests indicate that the hardness, yield strength and ultimate strength in HAZ are higher than those in base metal, and their values decrease with the distance from the fusion line. However, the elongation increases as the strengths decrease. The higher strength and lower elongation in the HAZ are mainly attributed to residual strains with the function of strain-hardening. Moreover, the change of Σ3 boundary which is in good agreement with that of elongation suggests a positive influence on the plastic deformation.
1. Introduction Ni-Mo-Cr alloys have been widely used in aerospace, chemical and nuclear industries due to their high corrosion resistance, superior strength at room and elevated temperatures [1–5]. To date, extensive studies have been carried out on Ni-Mo-Cr alloys. These studies are mainly focused on weld properties [6,7] in view of the fact that a welded joint is normally a weakest part in a welded component. The residual strains and changes of material properties such as local microstructure, chemical composition and carbide distribution may also occur in the heat-affected zone (HAZ) during the multiple welding thermal cycles [8,9]. Residual strain-hardening is considered to be a main factor to lead to heterogeneous mechanical performance and some potential problems [10–12]. Previous study on Alloy 690TT revealed that strain-hardening resulting from weld shrinkage increases the micro-hardness in the HAZ and sequentially promotes corrosion cracking sensitivity at high temperatures [11,13]. Ductility-dip crack can also easily occur in the HAZ for some Ni-Mo-Cr alloys with high concentration of carbon and alloy elements [14–16]. Uneven grain boundary is another important factor which can affect the mechanical properties in the HAZ in nickel base alloy, Σ3 boundaries is reported to
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Corresponding author. E-mail address: [email protected] (Z. Li).
http://dx.doi.org/10.1016/j.msea.2017.05.072 Received 27 January 2017; Received in revised form 18 May 2017; Accepted 19 May 2017 Available online 20 May 2017 0921-5093/ © 2017 Elsevier B.V. All rights reserved.
be very helpful to prevent intergranular cracking and expansion of elongation [17,18]. As a representative Ni–Mo–Cr solid solution strengthened alloy developed by Oak Ridge National Laboratory (ORNL), Hastelloy N alloy (UNS N10003) has been used as a main structural material for Molten Salt Reactor in 1960s [19] and Thorium Molten Salt Reactor (TMSR) in 2010s [20,21], which serves at elevated temperatures between 650 °C and 750 °C [22–24]. To our best knowledge, no works on detailed investigation of local strains and grain boundaries in the HAZ and its influence on the mechanical performance have been reported to date. In this work, we have investigated the effect of welding thermal cycle on the microstructure and mechanical properties of HAZ. Samples were cut from the HAZ to study the evolution process of microstructure and mechanical properties. The hardness, tensile properties, microstructure properties in terms of distribution of precipitations, grain sizes, dislocations and grain boundaries character parameters such as fractions of Σ3 boundary, distribution of kernel average misorientation (KAM) in Alloy GH3535 (UNS N10003) HAZ were characterized. Furthermore, the relationship between the microstructure properties and mechanical properties in both HAZ and base metal were studied.
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Table 1 Chemical compositions of GH3535 alloy and filler wire (wt%). Alloy
Mo
Cr
Fe
C
Si
Mn
Ni
GH3535 ERNiMo-2
16.5 16.4
7.0 8.0
4.0 5.0
0.06 0.05
0.27 1
0.5 1
Bal. Bal.
Fig. 2. Dimensions of the tensile specimens.
strain rates were set as 0.005/min and 0.05/min before and after the yield point, according to ASTM E8 [26] and ASTM E21 [27], respectively. 3. Results 3.1. Microstructure evolution and residual strains in the HAZ and base metal Microstructure of the HAZ near the fusion line was observed by SEM as shown in Fig. 3. The region near the fusion line as shown in Fig. 3(a) can be divided into three parts, namely, weld, HAZ(E) and HAZ(R). HAZ(E) stands for the region where primary M6C (Location A) carbides have transformed into eutectic carbides (Location B), and HAZ(R) stands for the rest of the HAZ without carbides transformation. In addition, Fig. 3(a) shows that the area with significant microstructure change is approximately in the range of 400 µm from the fusion line. Fig. 3(b) shows a greater magnification of HAZ(E) where two forms of eutectic carbides are observed. One is lamellar structure (Location B) as shown in Fig. 3(b) and (d). The other one close to the fusion line is spheroidized carbides (Location C) consisting of fine particles and small rod-like as shown in Fig. 3(e). The lamellar structure carbides (Location B) is the production of eutectic reaction of primary carbides which occurred when peak temperature of the thermal cycle over 1300 °C [28] during welding. The type of lamellar carbides seems to be stable when the heat input is small, however, in the cladding process, they can occur spheroidization and turn into carbides with the structure of fine particles and small rod-like under the influence of multiple repeated welding thermal cycles with high peak temperatures. To further investigate the influence of welding process on the microstructure of HAZ, several regions with different distances from the fusion line (X=0.5–20 mm) were characterized by SEM to investigate the microstructure evolution. As shown in Fig. 4(a)–(h), the eight local regions possess similar microstructure, the sizes of primary M6C in the HAZ are all in the range from 3 µm to 10 µm. Precipitates in the grain boundaries are observed in the whole HAZ and base metal. In general, there are no significant changes on the microstructure in the region within 0.5–20 mm from the fusion line. The grain boundary character maps, all-Euler maps, kernel average misorientation (KAM) distribution maps and the grain boundary character distribution of the HAZ analyzed by EBSD are shown in Figs. 5–7, respectively, as a function of distance from the fusion line. Eight samples were taken from the HAZ and base metal for the EBSD experiment, the distances of sample locations from the fusion line were 0 mm, 1 mm, 2 mm, 3 mm, 4 mm, 5 mm, 10 mm, 20 mm, respectively. Figs. 5(a) and 6(a) show that the grain size generally decreases with the distance from the fusion line. The average grain size near the fusion line (X=0 mm) is 38 µm, approximately 5 µm larger than at location X=20 mm. Figs. 5(b) and 6(b) show that Σ3 boundaries increase with the distance from the fusion line. And the frequency of Σ3 boundaries goes up from 19.2% (X=0 mm) to about 40% (X=20 mm). The KAM map can be used to access the residual strains [29]. Blue and red colors in Fig. 5(c) represent the minimum and the maximum KAM values, respectively. The color of the KAM near the fusion line (X=0 mm) is much brighter, indicating the residual strain in this region
Fig. 1. Weldment schematic and locations of tensile specimens in parent alloy.
2. Experimental procedure The weldment was prepared by Gas Tungsten Arc Welding by using two GH3535 alloy plates with each thickness of 20 mm and ERNiMo-2 filler wire with a diameter of 1.2 mm. The nominal chemical compositions of the parent alloy and the filler wire are listed in Table 1. 2% Cerium Tungsten electrode with a diameter of 2.4 mm and high purity Argon (99.99%) were applied as welding consumable. A cladding layer with a thickness of 30 mm was buttered on the surface of GH3535 alloy plate. The detailed welding parameters were: welding current 280 A, pulse frequency 2.5 Hz, peak pulse duration 50% and base value/peak value 50%, welding speed 90 mm/min. Fig. 1 shows the schematic of the weldment of GH3535 alloy. Microstructure was characterized on a Zeiss LEO 1530VP scanning electron microscope (SEM) equipped with energy dispersive spectroscopy (EDS). The grain boundary character and the local strains distributions in the HAZ were analyzed by an Oxford electron backscattered diffraction (EBSD) system with the Aztec software. KAM is the average misorientation between every pixel and its surrounding pixels in the EBSD measurements, which can estimate the local strains and represent the density of geometrically necessary dislocation (GND) in crystalline materials [9,25]. Samples for EBSD were cut from the HAZ and processed by vibration polishing for 2 h with 0.5 µm diamond paste to remove surface stress. EBSD data post-processing was done with Aztec software packages. The whole EBSD experiment process is as follows. (1) Samples for EBSD were processed by vibration polishing for 2 h with 0.5 µm diamond paste to remove surface stress. The acceleration voltage and current were set as 20 kV and 100 mA respectively. EBSD mapping were carried out within an area of 300*500 µm2 with a step size of 1 µm. (2) EBSD maps were processed by noise reduction after mapping. (3) Corresponding histograms and data of circle equivalent grain size distribution, KAM were calculated automatically from EBSD maps. Vickers hardness in the HAZ was measured on a ZHV 30 micro Vickers with a load of 500 gf and a holding time of 15 s. Eight tensile specimens with a thickness of 2 mm were sampled from both HAZ and base metal along the thickness direction as shown in Fig. 1. The geometry of tensile specimens is depicted in Fig. 2. The tensile tests with a set of 2 specimens were performed on a Zwick Z100 universal testing machine at two temperatures 25 °C and 700 °C. The 49
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Fig. 3. Microstructure of HAZ near the fusion line (a) HAZ near the fusion line, (b) morphology evolution of eutectic carbides, (c) primary granular-like carbides, (d) lamellar eutectic carbides, (e) spheroidized eutectic carbides.
X=15 mm. Fig. 8 illustrates the TEM results at a high magnification with 7000×. Fig. 8(a) shows the dislocation distribution in the HAZ near the fusion line. Plenty of parallel dislocation lines with two directions indicate that double slip of dislocation happened during welding, the angels between the two slip directions is about 70.5° which is consistent with that of direction along [1-1-1] and [1-11] on {111} plane with FCC crystal structure (GH3535 alloy is of FCC structure). In addition, dislocation networks formed by double slip of discation and tangle can also be observed. Those numerous dislocation slip and tangle indicate large strain and stress had been produced in this region during welding process since the dislocation with high density needs considerable driving force to multiplication and motion. The density of dislocations in the region with X=3 mm (Fig. 8(b)) decreases in comparison with the former region (X=0 mm), and the presence of double slip and tangle of dislocation is much less. In the range of 6–12 mm from the fusion line as shown in Fig. 8(c)–(e),
is high, especially at the grain boundaries and carbide-rich regions in the interior of the grains. With the distance away from the fusion line, the colors are getting darker and bluer. These results indicate the residual strain decreases as the distance from the fusion line. KAM vs relative frequency is also helpful to determine the residual strains. In Fig. 7 the four regions (0 mm, 1 mm, 2 mm, 3 mm) close to the fusion line have similar angles at 0.45°, it is higher than that of the last four regions (4 mm, 5 mm, 10 mm, 20 mm) which have similar KAM angles at 0.35°. As for the relative frequency, it increases with the distance from the fusion line. It can be draw that the residual strain in the region near the fusion line is much higher, which is in good agreement with the results in Fig. 5(c). TEM was employed to determine variation of dislocations density in the HAZ and base metal. The specimens were sampled in the regions with various distances (X) from the fusion line, as follows, (a) X=0 mm, (b) X=3 mm, (c) X=6 mm, (d) X=9 mm, (e) X=12 mm, (f)
Fig. 4. Carbides and microstructure with the distance (X) from the fusion line, (a) X=0.5 mm, (b) X=1 mm, (c) X=2 mm, (d) X=3 mm, (e) X=4 mm, (f) X=5 mm, (g) X=10 mm, (h) X=20 mm.
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Fig. 5. (a) Grain boundary character maps, (b) all-Euler and (c) KAM maps as a function of the distance from the fusion line. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
densities and tangle of dislocations with the distance decline slowly in addition to pile-up of dislocation and the less double slip with the angle about 70.5°. On contrary, double slip of dislocation cannot be observed in Fig. 8(f) with the distance X=15 mm from the fusion line, which indicates a smaller deformation occurred in this region. 3.2. Varieties of Hardness and tensile properties of the HAZ and base metal The Vickers hardness was measured along the line across the fusion boundary as shown in Fig. 9. According to the change of trend in hardness, the width of HAZ can be approximately determined as 9 mm since the hardness seems stable in the range of 10–20 mm. As indicated in Fig. 3, HAZ consists of HAZ(E) and HAZ(R). Fig. 9 confirms that the width of HAZ(E) is about 300 µm. The highest hardness in the whole HAZ is about 310 HV. HAZ(R) has the width of about 9 mm, in which the hardness values fluctuate in the range of 215-270HV. Further, the obviously hardened area in the HAZ is within about 4 mm from the fusion line, where the average hardness is about 275 HV, this is consistent with the KAM angles in Figs. 5(c) and 7 that the serious residual strains occurred in the zone within the range of 0–4 mm from the fusion line. As a comparison, the hardness in the base metal
Fig. 7. Kernel average misorientation angles with the distance from the fusion line in the HAZ and base metal.
fluctuating from 190 to 220HV, is less than those of HAZ and weld metal. Fig. 10 shows the tensile properties of different regions with the
Fig. 6. Grain sizes and boundaries calculated by Aztec software packages, (a) average grain sizes in the different regions in the HAZ and base metal, (b) percentage of Σ3 in the different regions of HAZ and base metal.
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Fig. 8. Densities of dislocations with the distance from the fusion line in the HAZ and base metal. The locations of TEM samples with distances (X) from the fusion line are: (a) X=0 mm, (b) X=3 mm, (c) X=6 mm, (d) X=9 mm, (e) X=12 mm, (f) X=15 mm.
from about 35% (2.5 mm) to 56% (10 mm), and then fluctuates at 55% for the last four specimens (12.5–20 mm). In the case of tensile tests at 700 °C as shown in Fig. 10(b), the σ0.2, UTS and elongation of specimens at different locations have similar trends as that of specimens at 25 °C. The σ0.2 and UTS decline from 380 MPa and 545 MPa (2.5 mm) to 265 MPa and 515 MPa (10 mm), then fall to about 250 MPa and 510 MPa for the last four specimens (12.5–20 mm), respectively. Moreover, the elongation goes up from 25% to 38% in the HAZ within the range of 0–10 mm and fluctuates around 35% in the base metal (12.5–20 mm). 4. Discussion Significant variations of mechanical performance in GH3535 alloy HAZ and base metal are detected as shown in Figs. 8 and 9. The σ0.2, UTS and hardness are found to decrease away from the fusion line. However, the elongation presents a contrary trend with those of the strength. The observed enhancement in strength and reduction plasticity in the HAZ is expected as the results of several factors, such as precipitates, grain sizes, grain boundary character distribution, residual strains. In this work, Fig. 3 indicates eutectic carbides transformation occurred near the fusion line in the HAZ and brought significant microstructure change, in addition, the eutectic carbides was reported to enhance the strength of the Ni-Mo-Cr matrix at elevated temperature to some extent [2]. Note that, the region of HAZ(E) with eutectic carbides transformation is only a narrow zone with a width of about 300 µm. This cannot explain the high strength and hardness in the HAZ (R) without eutectic carbides. Fig. 4 shows that morphologies of primary M6C and precipitates in the grain boundaries and/or grains do not vary within the range of 0.5–20 mm from the fusion line. Therefore the high strength in the HAZ could not be explained by the M6C carbides. As for the grain sizes, a slight change of the grain sizes in
Fig. 9. Distributions of Vickers hardness with the distance from the fusion line in the HAZ and base metal.
function of distance from the fusion line at 25 °C and 700 °C. The trends of strength shown in Fig. 10(a) and (b) are consistent with that of hardness in Fig. 9. Significant changes of the tensile properties can be seen in the range of 0–10 mm. However, there is only slight variation of strength and elongation in the range of 10–20 mm, confirming that the HAZ locates in the first 10 mm from the fusion line. In addition, the yield strength (σ0.2) and ultimate tensile strength (UTS) of the specimen near the fusion line are the highest at 25 °C as show in Fig. 10(a). As the distance from the fusion line, the σ0.2 and UTS decline from about 560 MPa and 860 MPa at 2.5 mm to about 400 MPa and 800 MPa at 10 mm, then remains almost stable at about 350 MPa and 800 MPa in the range of 12.5–20 mm, respectively. However the elongation presents an almost contrary trend with the strength. It increases rapidly 52
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Fig. 10. Tensile properties of the regions with distance from the fusion line in the HAZ and base metal at (a) 25 °C and (b) 700 °C.
elongation in the HAZ with comparison to the base metal. In addition, the lower elongation near the fusion line is also affected by the less Σ3 boundaries caused by welding thermal cycles. Furthermore, it should be a concern for such low ductility in the HAZ, especially in the HAZ(E) since the ductility-dip crack may occur in a high constraint state at elevated temperature, and related experiment of crack susceptibility will be necessary in future.
different regions of HAZ and base metal are observed in Figs. 5(a) and 6(a), and the average grain size near the fusion line is 38 µm and higher about 5 µm than that of base metal far away from the fusion line, the minor difference in grain size cannot be explained the observed higher strength in the HAZ as well. Previous studies [18,30,31] confirmed that CSL grain boundaries with Σ≤29 have special properties. The elongation in the 304L SS was improved due to the numerous presence of low Σ coincidence boundaries [18]. There are several merits of low Σ coincidence boundaries, including intrinsically low energy and high probability of slip continuity [32]. The EBSD study in Fig. 5 shows that the welding thermal cycle results in a decrease of Σ3 boundary as the distance close to the fusion line in welded GH3535 alloy, especially in the region near the fusion line. The change of trend in Σ3 boundaries is in good agreement with that of elongation as shown in Fig. 9(a) and (b). Thus, the rapidly reduction of Σ3 boundaries near the fusion line in the HAZ may resulted in a lower elongation compared with the region far away from the fusion line. In other words, the influence of Σ3 boundaries on tensile properties is evident with regard to the elongation. The residual strains and densities of dislocation near the fusion line in the HAZ as show in Fig. 5(a) and Fig. 8 demonstrate plastic strain had occurred during welding process. For the metallic material with residual strain, its plasticity will decreases and the strength increases, especially the yield strength when reloaded due to the effect of strainhardening [32,33]. The dislocation density has a positive correlation with local strain/deformation and any large local strain/deformation generally suggests a high dislocation density [34,35], which is confirmed by Fig. 8(a) that the region near the fusion line has the most severe strain and highest density of dislocation. In addition, the existed networks and tangle of dislocation will impose barriers to dislocation glide and increase the resistance of dislocation movement when samples get reloaded,, subsequently the plastic deformation becomes difficult, which leads the mechanical strength to be improved. Resistance of dislocation glide caused by the interaction between dislocations can be expressed by the following equation:
5. Conclusions Welding thermal cycles introduced eutectic transformation of carbides, various Σ3 boundaries, local strains and dislocations in the HAZ of GH3535 alloy. Within the range of 0.5–20 mm from the fusion line, the grains sizes and precipitates did not vary significantly. For the mechanical properties, the hardness, yield strength and ultimate strength decrease with the distance from the fusion line in the HAZ and higher than that in base metal. However, elongation presents a contrary trend with the strength. The local strain/deformation caused by welding shrinkage in the HAZ is thought to be the main reason for the high strength and low elongation due to the effect of strainhardening. Strain/deformation occurred during the welding process caused dislocation multiplication, slip and tangle, which increases the resistance of dislocation movement when reloaded. In addition, the change of Σ3 boundary which is in good agreement with that of elongation suggests a positive influence on the plastic deformation. Acknowledgement This work was supported by National Key Research and Development Program of China (2016YFB0700404), National Natural Science Foundation of China (Grant No. 51371188, 51671122, 51671154, 51601213, 51501216), Strategic Priority Research Program of the Chinese Academy of Sciences (Grant No. XDA02004210), and Talent Development Fund of Shanghai (201650). References
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Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Microstructure and local strains in GH3535 alloy heat affected zone and their influence on the mechanical properties
MARK ⁎
Shuangjian Chena,b, D.K.L. Tsanga, Li Jianga,b, Kun Yua,b, Chaowen Lia, Zhong Lia, Zhijun Lia, , Xingtai Zhoua, Jianguo Yangc a b c
Center for Thorium Molten Salt Reactor System, Shanghai Institute of Applied Physics, Chinese Academy of Sciences, Shanghai 201800, PR China University of Chinese Academy of Sciences, Beijing 100049, PR China Zhejiang University of Technology, Hangzhou 310014, PR China
A R T I C L E I N F O
A B S T R A C T
Keywords: EBSD Twinning Dislocations Nickel based superalloys Welding Hardening
GH3535 alloy plates were welded by Gas Tungsten Arc Welding in order to study the evolution of microstructure and mechanical properties in the heat-affected zone (HAZ). Our results suggest that welding thermal cycles induced the morphology evolution of M6C carbides from block to eutectic near the fusion line in the HAZ. Electron backscatter diffraction (EBSD) results show that significant amounts of plastic strains occurred in the HAZ after welding. In addition, local coherent twin boundaries (Σ3) and dislocations were observed to decrease with the distance from the fusion line. Mechanical tests indicate that the hardness, yield strength and ultimate strength in HAZ are higher than those in base metal, and their values decrease with the distance from the fusion line. However, the elongation increases as the strengths decrease. The higher strength and lower elongation in the HAZ are mainly attributed to residual strains with the function of strain-hardening. Moreover, the change of Σ3 boundary which is in good agreement with that of elongation suggests a positive influence on the plastic deformation.
1. Introduction Ni-Mo-Cr alloys have been widely used in aerospace, chemical and nuclear industries due to their high corrosion resistance, superior strength at room and elevated temperatures [1–5]. To date, extensive studies have been carried out on Ni-Mo-Cr alloys. These studies are mainly focused on weld properties [6,7] in view of the fact that a welded joint is normally a weakest part in a welded component. The residual strains and changes of material properties such as local microstructure, chemical composition and carbide distribution may also occur in the heat-affected zone (HAZ) during the multiple welding thermal cycles [8,9]. Residual strain-hardening is considered to be a main factor to lead to heterogeneous mechanical performance and some potential problems [10–12]. Previous study on Alloy 690TT revealed that strain-hardening resulting from weld shrinkage increases the micro-hardness in the HAZ and sequentially promotes corrosion cracking sensitivity at high temperatures [11,13]. Ductility-dip crack can also easily occur in the HAZ for some Ni-Mo-Cr alloys with high concentration of carbon and alloy elements [14–16]. Uneven grain boundary is another important factor which can affect the mechanical properties in the HAZ in nickel base alloy, Σ3 boundaries is reported to
⁎
Corresponding author. E-mail address: [email protected] (Z. Li).
http://dx.doi.org/10.1016/j.msea.2017.05.072 Received 27 January 2017; Received in revised form 18 May 2017; Accepted 19 May 2017 Available online 20 May 2017 0921-5093/ © 2017 Elsevier B.V. All rights reserved.
be very helpful to prevent intergranular cracking and expansion of elongation [17,18]. As a representative Ni–Mo–Cr solid solution strengthened alloy developed by Oak Ridge National Laboratory (ORNL), Hastelloy N alloy (UNS N10003) has been used as a main structural material for Molten Salt Reactor in 1960s [19] and Thorium Molten Salt Reactor (TMSR) in 2010s [20,21], which serves at elevated temperatures between 650 °C and 750 °C [22–24]. To our best knowledge, no works on detailed investigation of local strains and grain boundaries in the HAZ and its influence on the mechanical performance have been reported to date. In this work, we have investigated the effect of welding thermal cycle on the microstructure and mechanical properties of HAZ. Samples were cut from the HAZ to study the evolution process of microstructure and mechanical properties. The hardness, tensile properties, microstructure properties in terms of distribution of precipitations, grain sizes, dislocations and grain boundaries character parameters such as fractions of Σ3 boundary, distribution of kernel average misorientation (KAM) in Alloy GH3535 (UNS N10003) HAZ were characterized. Furthermore, the relationship between the microstructure properties and mechanical properties in both HAZ and base metal were studied.
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Table 1 Chemical compositions of GH3535 alloy and filler wire (wt%). Alloy
Mo
Cr
Fe
C
Si
Mn
Ni
GH3535 ERNiMo-2
16.5 16.4
7.0 8.0
4.0 5.0
0.06 0.05
0.27 1
0.5 1
Bal. Bal.
Fig. 2. Dimensions of the tensile specimens.
strain rates were set as 0.005/min and 0.05/min before and after the yield point, according to ASTM E8 [26] and ASTM E21 [27], respectively. 3. Results 3.1. Microstructure evolution and residual strains in the HAZ and base metal Microstructure of the HAZ near the fusion line was observed by SEM as shown in Fig. 3. The region near the fusion line as shown in Fig. 3(a) can be divided into three parts, namely, weld, HAZ(E) and HAZ(R). HAZ(E) stands for the region where primary M6C (Location A) carbides have transformed into eutectic carbides (Location B), and HAZ(R) stands for the rest of the HAZ without carbides transformation. In addition, Fig. 3(a) shows that the area with significant microstructure change is approximately in the range of 400 µm from the fusion line. Fig. 3(b) shows a greater magnification of HAZ(E) where two forms of eutectic carbides are observed. One is lamellar structure (Location B) as shown in Fig. 3(b) and (d). The other one close to the fusion line is spheroidized carbides (Location C) consisting of fine particles and small rod-like as shown in Fig. 3(e). The lamellar structure carbides (Location B) is the production of eutectic reaction of primary carbides which occurred when peak temperature of the thermal cycle over 1300 °C [28] during welding. The type of lamellar carbides seems to be stable when the heat input is small, however, in the cladding process, they can occur spheroidization and turn into carbides with the structure of fine particles and small rod-like under the influence of multiple repeated welding thermal cycles with high peak temperatures. To further investigate the influence of welding process on the microstructure of HAZ, several regions with different distances from the fusion line (X=0.5–20 mm) were characterized by SEM to investigate the microstructure evolution. As shown in Fig. 4(a)–(h), the eight local regions possess similar microstructure, the sizes of primary M6C in the HAZ are all in the range from 3 µm to 10 µm. Precipitates in the grain boundaries are observed in the whole HAZ and base metal. In general, there are no significant changes on the microstructure in the region within 0.5–20 mm from the fusion line. The grain boundary character maps, all-Euler maps, kernel average misorientation (KAM) distribution maps and the grain boundary character distribution of the HAZ analyzed by EBSD are shown in Figs. 5–7, respectively, as a function of distance from the fusion line. Eight samples were taken from the HAZ and base metal for the EBSD experiment, the distances of sample locations from the fusion line were 0 mm, 1 mm, 2 mm, 3 mm, 4 mm, 5 mm, 10 mm, 20 mm, respectively. Figs. 5(a) and 6(a) show that the grain size generally decreases with the distance from the fusion line. The average grain size near the fusion line (X=0 mm) is 38 µm, approximately 5 µm larger than at location X=20 mm. Figs. 5(b) and 6(b) show that Σ3 boundaries increase with the distance from the fusion line. And the frequency of Σ3 boundaries goes up from 19.2% (X=0 mm) to about 40% (X=20 mm). The KAM map can be used to access the residual strains [29]. Blue and red colors in Fig. 5(c) represent the minimum and the maximum KAM values, respectively. The color of the KAM near the fusion line (X=0 mm) is much brighter, indicating the residual strain in this region
Fig. 1. Weldment schematic and locations of tensile specimens in parent alloy.
2. Experimental procedure The weldment was prepared by Gas Tungsten Arc Welding by using two GH3535 alloy plates with each thickness of 20 mm and ERNiMo-2 filler wire with a diameter of 1.2 mm. The nominal chemical compositions of the parent alloy and the filler wire are listed in Table 1. 2% Cerium Tungsten electrode with a diameter of 2.4 mm and high purity Argon (99.99%) were applied as welding consumable. A cladding layer with a thickness of 30 mm was buttered on the surface of GH3535 alloy plate. The detailed welding parameters were: welding current 280 A, pulse frequency 2.5 Hz, peak pulse duration 50% and base value/peak value 50%, welding speed 90 mm/min. Fig. 1 shows the schematic of the weldment of GH3535 alloy. Microstructure was characterized on a Zeiss LEO 1530VP scanning electron microscope (SEM) equipped with energy dispersive spectroscopy (EDS). The grain boundary character and the local strains distributions in the HAZ were analyzed by an Oxford electron backscattered diffraction (EBSD) system with the Aztec software. KAM is the average misorientation between every pixel and its surrounding pixels in the EBSD measurements, which can estimate the local strains and represent the density of geometrically necessary dislocation (GND) in crystalline materials [9,25]. Samples for EBSD were cut from the HAZ and processed by vibration polishing for 2 h with 0.5 µm diamond paste to remove surface stress. EBSD data post-processing was done with Aztec software packages. The whole EBSD experiment process is as follows. (1) Samples for EBSD were processed by vibration polishing for 2 h with 0.5 µm diamond paste to remove surface stress. The acceleration voltage and current were set as 20 kV and 100 mA respectively. EBSD mapping were carried out within an area of 300*500 µm2 with a step size of 1 µm. (2) EBSD maps were processed by noise reduction after mapping. (3) Corresponding histograms and data of circle equivalent grain size distribution, KAM were calculated automatically from EBSD maps. Vickers hardness in the HAZ was measured on a ZHV 30 micro Vickers with a load of 500 gf and a holding time of 15 s. Eight tensile specimens with a thickness of 2 mm were sampled from both HAZ and base metal along the thickness direction as shown in Fig. 1. The geometry of tensile specimens is depicted in Fig. 2. The tensile tests with a set of 2 specimens were performed on a Zwick Z100 universal testing machine at two temperatures 25 °C and 700 °C. The 49
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Fig. 3. Microstructure of HAZ near the fusion line (a) HAZ near the fusion line, (b) morphology evolution of eutectic carbides, (c) primary granular-like carbides, (d) lamellar eutectic carbides, (e) spheroidized eutectic carbides.
X=15 mm. Fig. 8 illustrates the TEM results at a high magnification with 7000×. Fig. 8(a) shows the dislocation distribution in the HAZ near the fusion line. Plenty of parallel dislocation lines with two directions indicate that double slip of dislocation happened during welding, the angels between the two slip directions is about 70.5° which is consistent with that of direction along [1-1-1] and [1-11] on {111} plane with FCC crystal structure (GH3535 alloy is of FCC structure). In addition, dislocation networks formed by double slip of discation and tangle can also be observed. Those numerous dislocation slip and tangle indicate large strain and stress had been produced in this region during welding process since the dislocation with high density needs considerable driving force to multiplication and motion. The density of dislocations in the region with X=3 mm (Fig. 8(b)) decreases in comparison with the former region (X=0 mm), and the presence of double slip and tangle of dislocation is much less. In the range of 6–12 mm from the fusion line as shown in Fig. 8(c)–(e),
is high, especially at the grain boundaries and carbide-rich regions in the interior of the grains. With the distance away from the fusion line, the colors are getting darker and bluer. These results indicate the residual strain decreases as the distance from the fusion line. KAM vs relative frequency is also helpful to determine the residual strains. In Fig. 7 the four regions (0 mm, 1 mm, 2 mm, 3 mm) close to the fusion line have similar angles at 0.45°, it is higher than that of the last four regions (4 mm, 5 mm, 10 mm, 20 mm) which have similar KAM angles at 0.35°. As for the relative frequency, it increases with the distance from the fusion line. It can be draw that the residual strain in the region near the fusion line is much higher, which is in good agreement with the results in Fig. 5(c). TEM was employed to determine variation of dislocations density in the HAZ and base metal. The specimens were sampled in the regions with various distances (X) from the fusion line, as follows, (a) X=0 mm, (b) X=3 mm, (c) X=6 mm, (d) X=9 mm, (e) X=12 mm, (f)
Fig. 4. Carbides and microstructure with the distance (X) from the fusion line, (a) X=0.5 mm, (b) X=1 mm, (c) X=2 mm, (d) X=3 mm, (e) X=4 mm, (f) X=5 mm, (g) X=10 mm, (h) X=20 mm.
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Fig. 5. (a) Grain boundary character maps, (b) all-Euler and (c) KAM maps as a function of the distance from the fusion line. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
densities and tangle of dislocations with the distance decline slowly in addition to pile-up of dislocation and the less double slip with the angle about 70.5°. On contrary, double slip of dislocation cannot be observed in Fig. 8(f) with the distance X=15 mm from the fusion line, which indicates a smaller deformation occurred in this region. 3.2. Varieties of Hardness and tensile properties of the HAZ and base metal The Vickers hardness was measured along the line across the fusion boundary as shown in Fig. 9. According to the change of trend in hardness, the width of HAZ can be approximately determined as 9 mm since the hardness seems stable in the range of 10–20 mm. As indicated in Fig. 3, HAZ consists of HAZ(E) and HAZ(R). Fig. 9 confirms that the width of HAZ(E) is about 300 µm. The highest hardness in the whole HAZ is about 310 HV. HAZ(R) has the width of about 9 mm, in which the hardness values fluctuate in the range of 215-270HV. Further, the obviously hardened area in the HAZ is within about 4 mm from the fusion line, where the average hardness is about 275 HV, this is consistent with the KAM angles in Figs. 5(c) and 7 that the serious residual strains occurred in the zone within the range of 0–4 mm from the fusion line. As a comparison, the hardness in the base metal
Fig. 7. Kernel average misorientation angles with the distance from the fusion line in the HAZ and base metal.
fluctuating from 190 to 220HV, is less than those of HAZ and weld metal. Fig. 10 shows the tensile properties of different regions with the
Fig. 6. Grain sizes and boundaries calculated by Aztec software packages, (a) average grain sizes in the different regions in the HAZ and base metal, (b) percentage of Σ3 in the different regions of HAZ and base metal.
51
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Fig. 8. Densities of dislocations with the distance from the fusion line in the HAZ and base metal. The locations of TEM samples with distances (X) from the fusion line are: (a) X=0 mm, (b) X=3 mm, (c) X=6 mm, (d) X=9 mm, (e) X=12 mm, (f) X=15 mm.
from about 35% (2.5 mm) to 56% (10 mm), and then fluctuates at 55% for the last four specimens (12.5–20 mm). In the case of tensile tests at 700 °C as shown in Fig. 10(b), the σ0.2, UTS and elongation of specimens at different locations have similar trends as that of specimens at 25 °C. The σ0.2 and UTS decline from 380 MPa and 545 MPa (2.5 mm) to 265 MPa and 515 MPa (10 mm), then fall to about 250 MPa and 510 MPa for the last four specimens (12.5–20 mm), respectively. Moreover, the elongation goes up from 25% to 38% in the HAZ within the range of 0–10 mm and fluctuates around 35% in the base metal (12.5–20 mm). 4. Discussion Significant variations of mechanical performance in GH3535 alloy HAZ and base metal are detected as shown in Figs. 8 and 9. The σ0.2, UTS and hardness are found to decrease away from the fusion line. However, the elongation presents a contrary trend with those of the strength. The observed enhancement in strength and reduction plasticity in the HAZ is expected as the results of several factors, such as precipitates, grain sizes, grain boundary character distribution, residual strains. In this work, Fig. 3 indicates eutectic carbides transformation occurred near the fusion line in the HAZ and brought significant microstructure change, in addition, the eutectic carbides was reported to enhance the strength of the Ni-Mo-Cr matrix at elevated temperature to some extent [2]. Note that, the region of HAZ(E) with eutectic carbides transformation is only a narrow zone with a width of about 300 µm. This cannot explain the high strength and hardness in the HAZ (R) without eutectic carbides. Fig. 4 shows that morphologies of primary M6C and precipitates in the grain boundaries and/or grains do not vary within the range of 0.5–20 mm from the fusion line. Therefore the high strength in the HAZ could not be explained by the M6C carbides. As for the grain sizes, a slight change of the grain sizes in
Fig. 9. Distributions of Vickers hardness with the distance from the fusion line in the HAZ and base metal.
function of distance from the fusion line at 25 °C and 700 °C. The trends of strength shown in Fig. 10(a) and (b) are consistent with that of hardness in Fig. 9. Significant changes of the tensile properties can be seen in the range of 0–10 mm. However, there is only slight variation of strength and elongation in the range of 10–20 mm, confirming that the HAZ locates in the first 10 mm from the fusion line. In addition, the yield strength (σ0.2) and ultimate tensile strength (UTS) of the specimen near the fusion line are the highest at 25 °C as show in Fig. 10(a). As the distance from the fusion line, the σ0.2 and UTS decline from about 560 MPa and 860 MPa at 2.5 mm to about 400 MPa and 800 MPa at 10 mm, then remains almost stable at about 350 MPa and 800 MPa in the range of 12.5–20 mm, respectively. However the elongation presents an almost contrary trend with the strength. It increases rapidly 52
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Fig. 10. Tensile properties of the regions with distance from the fusion line in the HAZ and base metal at (a) 25 °C and (b) 700 °C.
elongation in the HAZ with comparison to the base metal. In addition, the lower elongation near the fusion line is also affected by the less Σ3 boundaries caused by welding thermal cycles. Furthermore, it should be a concern for such low ductility in the HAZ, especially in the HAZ(E) since the ductility-dip crack may occur in a high constraint state at elevated temperature, and related experiment of crack susceptibility will be necessary in future.
different regions of HAZ and base metal are observed in Figs. 5(a) and 6(a), and the average grain size near the fusion line is 38 µm and higher about 5 µm than that of base metal far away from the fusion line, the minor difference in grain size cannot be explained the observed higher strength in the HAZ as well. Previous studies [18,30,31] confirmed that CSL grain boundaries with Σ≤29 have special properties. The elongation in the 304L SS was improved due to the numerous presence of low Σ coincidence boundaries [18]. There are several merits of low Σ coincidence boundaries, including intrinsically low energy and high probability of slip continuity [32]. The EBSD study in Fig. 5 shows that the welding thermal cycle results in a decrease of Σ3 boundary as the distance close to the fusion line in welded GH3535 alloy, especially in the region near the fusion line. The change of trend in Σ3 boundaries is in good agreement with that of elongation as shown in Fig. 9(a) and (b). Thus, the rapidly reduction of Σ3 boundaries near the fusion line in the HAZ may resulted in a lower elongation compared with the region far away from the fusion line. In other words, the influence of Σ3 boundaries on tensile properties is evident with regard to the elongation. The residual strains and densities of dislocation near the fusion line in the HAZ as show in Fig. 5(a) and Fig. 8 demonstrate plastic strain had occurred during welding process. For the metallic material with residual strain, its plasticity will decreases and the strength increases, especially the yield strength when reloaded due to the effect of strainhardening [32,33]. The dislocation density has a positive correlation with local strain/deformation and any large local strain/deformation generally suggests a high dislocation density [34,35], which is confirmed by Fig. 8(a) that the region near the fusion line has the most severe strain and highest density of dislocation. In addition, the existed networks and tangle of dislocation will impose barriers to dislocation glide and increase the resistance of dislocation movement when samples get reloaded,, subsequently the plastic deformation becomes difficult, which leads the mechanical strength to be improved. Resistance of dislocation glide caused by the interaction between dislocations can be expressed by the following equation:
5. Conclusions Welding thermal cycles introduced eutectic transformation of carbides, various Σ3 boundaries, local strains and dislocations in the HAZ of GH3535 alloy. Within the range of 0.5–20 mm from the fusion line, the grains sizes and precipitates did not vary significantly. For the mechanical properties, the hardness, yield strength and ultimate strength decrease with the distance from the fusion line in the HAZ and higher than that in base metal. However, elongation presents a contrary trend with the strength. The local strain/deformation caused by welding shrinkage in the HAZ is thought to be the main reason for the high strength and low elongation due to the effect of strainhardening. Strain/deformation occurred during the welding process caused dislocation multiplication, slip and tangle, which increases the resistance of dislocation movement when reloaded. In addition, the change of Σ3 boundary which is in good agreement with that of elongation suggests a positive influence on the plastic deformation. Acknowledgement This work was supported by National Key Research and Development Program of China (2016YFB0700404), National Natural Science Foundation of China (Grant No. 51371188, 51671122, 51671154, 51601213, 51501216), Strategic Priority Research Program of the Chinese Academy of Sciences (Grant No. XDA02004210), and Talent Development Fund of Shanghai (201650). References
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where τ, b and ρ are the shear stress of dislocation glide, Burgers vector and dislocation density, respectively. The increasing dislocation density will cause a rise in the shear stress of dislocation glide. Obviously, this is unconducive to the operation of multiple slip systems and plasticity [33] but helpful to enhance the strength. Fig. 5(c) and Figs. 9 and 10 show a strong coherence between the strain changes and mechanical strength distributions in the HAZ and base metal, they both declines with the distance from the fusion line and remains almost stable within the range of 10–20 mm from the fusion line. Based on the above results, residual strain as the result of weldshrinkage plays an important role to the higher strength and lower 53
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