Welding and Joining
Influence of Aging Treatment on the Microstructure Evolution and Mechanical Properties of 316H Weld Metals with Different C Contents
2021 Volume 61 Issue 3 Pages 911-918
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2021 Volume 61 Issue 3 Pages 911-918
The influence of 600°C aging treatment on the microstructure and mechanical properties of 316H stainless steel weld metals with different C contents have been studied. The results indicated that during the aging process, the rapid precipitation of M23C6 carbide occurred in δ-ferrite firstly owing to the high diffusion rate of C. After the C is depleted by precipitation of M23C6, the residual δ-ferrite transforms into σ phase through eutectoid transformation (δ-ferrite → σ phase + γ). Furthermore, after a long enough aging time, the transformation from M23C6 to σ phase occurred. The C content has a significant influence on the δ-ferrite transformation behavior, δ-ferrite in the low and medium C weld metals transforms into M23C6 and σ phase successively, while δ-ferrite in high C weld metal only transforms into M23C6 carbides. The variations of mechanical properties with aging conditions depended mainly on the microstructures at different aging conditions. For the low C weld metal, as the aging time increased, the increasing σ phase content improved the strength obviously. For the medium and high C weld metals, as the aging time increased, first the depletion of the solid solution C as a result the M23C6 precipitation deteriorated the strength, then the formation of σ phase improved the strength. Furthermore, with the increasing of the aging time, the precipitation of M23C6 and σ phase deteriorated the elongation and impact energy.
With the increasing shortage of energy supply around the world and the growing concern about the global warming and sustainable development, nuclear energy will play a vital role in the future development of the world. Nuclear power is a clean and efficient way to generate electricity.1) The development of nuclear power is a major strategic measure to meet power demand, optimize energy structure, ensure energy security and promote sustainable economic development. It is also an effective way to reduce environmental pollution and achieve ecological development.
During the nuclear reaction operation, the structural materials used in the reactor are exposed to a high temperature, high corrosion and high irradiation environment for a long time, and the service conditions are extremely harsh, and in some cases, they are subject to high temperature liquid metal corrosion. In addition, these critical components are not replaceable during their service life.1,2) The feedback experience of the in-service inspection for nuclear power plant shows that many problems are caused by the failure of welded joints, and many factors affect the quality of welded joints.3,4,5,6) The performance of joints is not only related to the mechanical properties and weldability of base metals, but also to the properties of weld metal and welding processes. The properties of the weld metal depend on the welding material. Therefore, improving the research and development capability and level of welding materials, strictly controlling the manufacturing process of welding materials, and ensuring the stable quality of each batch of welding materials are the premise and basis for guaranteeing the quality of nuclear power reactor manufacturing.
316 austenitic stainless steel is widely used in the manufacture of nuclear reactor components owing to its excellent comprehensive properties, such as main vessel, support assembly, core barrel, etc.7,8,9,10,11,12) During the fusion welding of austenitic stainless steel, the tendency of hot cracking tends to occur when the structural restraint is too large. In order to prevent cracking in the welding process, it is usually desirable to form a certain amount of δ-ferrite in the weld microstructure. However, the δ-ferrite is harmful to the mechanical and corrosion properties of the weld metal under the high temperature service environment, so the δ-ferrite content in weld metal must be strictly controlled.13,14,15,16,17,18,19,20,21,22) At present, there is still much controversy about the precipitation mechanisms of carbides and intermetallic phases in the austenitic stainless steel weld metal and the decomposition mechanism of δ-ferrite during the aging process.2,5,6,23,24,25,26,27,28,29) Rhouma et al.23) conducted the phase identification of the 316 L austenitic stainless steel with a small amount of δ-ferrite aged for up to 80000 hours at the temperature range of 550–700°C. They found that the δ-ferrite decomposed gradually into M23C6 at 550°C and decomposed totally into intermetallic phases (σ, η, χ, and R) and into secondary austenite (γ) at the temperatures equal to or higher than 650°C. Farrar5) analyzed the phase transformation mechanism of metastable δ-ferrite in 316 austenitic stainless steel weld metals with two different carbon contents under different aging temperature. He found that a rapid transformation to χ phase occurred in the lean weld metal under the aging temperatures of 700 and 850°C, while for the high C weld metal, only M23C6 carbides precipitated at 700°C aging and a small amount of χ phase coupled with a large amount of M23C6 carbides formed at 850°C aging. Sasikala et al.24) studied the δ-ferrite transformation behavior in the 316(N) stainless steel weld metal during the creep process at the temperature range of 550–650°C. They stated that the δ-ferrite transformation during the creep process is controlled initially by diffusion of Cr through δ/γ phase boundaries to form carbides, and then followed by matrix diffusion of Mo to form the intermetallic phases. Smith et al.6) investigated the δ-ferrite decomposition products of the 316 weld metal with different δ-ferrite contents aged between 600 and 850°C. They developed a δ-ferrite transformation model which indicated that solute diffusion via the δ/γ interface and enrichment were important in controlling the δ-ferrite decomposition process. Gill et al.25) analyzed the effects of Cr and δ-ferrite content on the transformation kinetics of δ-ferrite in the 316 L stainless steel weld metal. They developed an equation to predict the δ-ferrite transformation kinetics and established a nomograph to optimize the weld metal composition and to minimize σ phase formation from δ-ferrite at high temperatures. Gill et al.26) studied the role played by carbon in influencing the δ-ferrite transformation behavior in 316 weld metal in the temperature range 600 to 750°C. They found that increase in C content changed the high temperature transformation behavior of δ-ferrite. The depletion of Cr and Mo caused by the precipitation of M23C6 resulted in the slow transformation of δ-ferrite in high C 316 weld metal. Dutt et al.2) analyzed the mechanical properties of the 316L stainless steel weld after long term aging treatment. They found that both impact energy and J-R curves decreased after aging at 370, 475 and 550°C, and variations of the mechanical properties were related to the transformation of the δ-ferrite during aging process. The above findings indicate that the weld metal composition, δ-ferrite content, aging temperature and time have great influences on the precipitation of carbides and kinetics of δ-ferrite decomposition, which still needs to be studied in depth.
In the present investigation, the microstructure and mechanical properties of the 316H austenitic stainless steel weld metals with three different C contents have been studied. The transformation behavior of δ-ferrite and precipitation mechanisms of M23C6 and σ phase in the weld metals aged at 600°C for different aging times were investigated. The variations of tensile and impact properties of the weld metals under different aging times were analyzed. Furthermore, the relationship between microstructure evolution and mechanical properties of the weld metals with different carbon contents under different aging conditions were discussed. This research provides theoretical and practical guidance for the control of the chemical composition, δ-ferrite content, microstructure and mechanical properties of the 316H austenitic stainless steel weld metal for high temperature and long term service exposure.
The 16 mm thick 316H stainless steel plates were butt welded using the gas tungsten arc welding process with an inter-pass temperature lower than 50°C, arc current of 180 A, arc voltage of 14 V, and a welding speed of 0.1 m/min. The schematic diagram of the butt type joint is shown in Fig. 1. Three kinds of welding wires with different carbon contents were used in this experiment. Tables 1 and 2 show the chemical compositions of the welding wires and corresponding weld metals. The weld metals were aged at 600°C for various times from 0 to 3000 hours for analyzing the influences of aging conditions and carbon content on the microstructure, tensile and impact properties of the weld metals. The δ-ferrite contents of different weld metals were tested using a calibrated ferrite tester (Ferrite Determ SP10a). The values taken are the average of 10 readings scattered on the samples.
Schematic diagram of the butt type joint.
| Materials | C | Cr | Ni | Mo | Mn | Si | N |
|---|---|---|---|---|---|---|---|
| C1 | 0.015 | 19.0 | 12.7 | 2.53 | 1.59 | 0.41 | 0.065 |
| C2 | 0.042 | 18.8 | 12.6 | 2.49 | 1.59 | 0.42 | 0.063 |
| C3 | 0.064 | 18.8 | 12.7 | 2.45 | 1.59 | 0.47 | 0.069 |
| Materials | C | Cr | Ni | Mo | Mn | Si | N |
|---|---|---|---|---|---|---|---|
| C1 | 0.016 | 19.1 | 12.5 | 2.48 | 1.54 | 0.41 | 0.056 |
| C2 | 0.042 | 19.2 | 12.7 | 2.47 | 1.56 | 0.42 | 0.054 |
| C3 | 0.062 | 18.9 | 12.4 | 2.49 | 1.55 | 0.43 | 0.046 |
Specimens for microstructure examination by the optical microscope (OM) were grounded, polished up to 0.25 μm diamond paste and then electrolytically etched in a 10% oxalic acid under the voltage of 5 V and current of 0.3 A for 10 s. For TEM observation, slices were cut from specimens, then grounded to 0.03 mm in thickness and punched to disks in 3 mm diameter. Thin-foil specimens were made from the disks by the jet-electro-polishing technique in a 10% perchloric acid and 90% ethanol electrolytic solution. TEM examination was carried by the FEI Tecnai G2 F20 STEM. The electrochemical extraction of the precipitates in the as-weld and as-aged weld metals were conducted in the solution of 10% hydrochloric acid and 90% methanol solution at room temperature under the constant extraction current of 0.2 A. The residue was rinsed, dried and collected. The phase types and contents of the residual phases were analyzed by the X-ray diffraction analysis equipment with a Cu Kɑ source.
Tensile and impact tests were conducted at room temperature on the MTS tensile apparatus and SANS-ZBC2452-CC impact machine. The tensile strain rate was 2.5 × 10−4 s−1 and 2 × 10−2 s−1 before and after yielding. Tensile specimens were machined parallel to the welding direction from the center of the as-welded and as-aged weld metals, while the standard size V-notch impact specimens were machined perpendicular to welding direction of the weld metal and the V-notch located at the center of the weld cross section. The configurations of the tensile and impact specimens are shown in Figs. 2 and 3.
Configuration of the tensile sample.
Configuration of the impact sample.
The OM micrographs of the as-welded weld metals with different C contents are presented in Fig. 4. It indicated that all the three kinds of weld metals contained austenite and a certain amount of δ-ferrite, and with the increasing C content, the δ-ferrite content decreased. The δ-ferrite contents in the as-welded C1, C2 and C3 weld metals obtained by the ferrite tester were 7.2%, 5.7%, and 2.8%, respectively.
Optical micrographs of the as-welded weld metals: (a) C1; (b) C2; (c) C3. (Online version in color.)
For the as-aged weld metals with different C contents, the variations of δ-ferrite transformation fractions with aging time on a logarithmic coordinate at 600°C are shown in Fig. 5. It is clear that the weld metals with different C contents have different transformation mechanisms. During the aging time from 0 to 3000 h at 600°C, for the low carbon (C1) and medium carbon (C2) weld metals, the transformation of δ-ferrite finished after 3000 h aging, while the δ-ferrite transformation in high carbon (C3) weld metal finished quickly after aging for 150 hours. At the initial stage of the aging treatment, for the weld metals with different C contents, the δ-ferrite transformation rate decreased in the order of medium carbon (C2), high carbon (C3) and low carbon (C1). With the increase of the aging time, the δ-ferrite transformation rates of low carbon (C1) and high carbon (C3) first increased and then kept stable, while the δ-ferrite transformation rate of medium carbon (C2) first increased then decreased and finally increased again. Therefore, when the aging time was lower than 50 h, the δ-ferrite transformation fraction decreased in the order of medium carbon (C2), high carbon (C3) and low carbon (C1). When the aging time was higher than or equal to 50 h, and lower than 1000 h, the δ-ferrite transformation fraction decreased in the order of high carbon (C3), medium carbon (C2) and low carbon (C1). Furthermore, when the aging time was higher than or equal to 1000 h, the δ-ferrite transformation in high carbon (C3) weld metal had already completed, and the low carbon (C1) weld metal had higher δ-ferrite transformation fraction than the medium carbon (C2) weld metal until the δ-ferrite transformation finished at the aging time of 3000 h.
Effect of aging time on δ-ferrite transformation fraction.
The TEM micrographs of the C1, C2 and C3 weld metals aged at 600°C for different holding times are shown in Figs. 6, 7, 8. Figure 6 shows TEM micrographs of C1 weld metal under different aging conditions. At the aging temperature of 600°C for 500 h, a large number of σ phase formed in δ-ferrite along with small spherical M23C6 precipitates in the interior of δ-ferrite and at δ/γ boundaries, as shown in Fig. 6(a). Former researches indicated that the M23C6 precipitation occurred firstly at δ/γ boundaries or within the δ-ferrite relating to the segregation behavior of the alloying elements.26,30,31) With the increase of the aging time to 1000 h, the content of the M23C6 precipitates decreased, while the content of the σ phase increased, as shown in Fig. 6(b). When the aging time was further increased to 3000 h, M23C6 precipitate could hardly be observed, as shown in Fig. 6(c), almost only σ phase was detected in the TEM micrograph. For the low C weld metal (C1), at the first stage of the aging process, M23C6 precipitated. When the C element was depleted, the M23C6 precipitation ceased. Subsequently, the eutectoid transformation δ-ferrite → σ phase + γ occurred.31,32) Furthermore, after a long enough aging time, the transformation from M23C6 to σ phase occurred, which led to the decrease of M23C6 content and increase of σ phase content.8)
TEM micrographs of C1 weld metal aged at 600°C for different aging times: (a) 500 h; (b) 1000 h; (c) 3000 h.
TEM micrographs of C2 weld metal aged at 600°C for different aging times: (a) 10 h; (b) 300 h; (c) 500 h; (d) 1000 h; (e) 3000 h.
TEM micrographs of C3 weld metal aged at 600°C for different aging times: (a) 500 h; (b) 1000 h; (c) 3000 h.
Figure 7 shows TEM micrographs of C2 weld metals under different aging conditions. At the aging temperature of 600°C for 10 h, only fine M23C6 precipitated (Fig. 7(a)). When the aging time was increased to 300 h at 600°C, M23C6 particles coarsened, and no σ phase formed (Fig. 7(b)). When the aging time was prolonged to 500 h, both M23C6 and σ phase were observed (Fig. 7(c)). Further increasing the aging time to 1000 and 3000 h, the M23C6 content decreased and the σ phase content increased (Figs. 7(d) and 7(e)).
Figure 8 shows TEM micrographs of C3 weld metals aged at 600°C for different times. After aged at 600°C for 500 h, only M23C6 particles were observed in the weld metal (Fig. 8(a)). When the aging time was increased to 1000 and 3000 h at 600°C, M23C6 precipitates and small amount of σ phase were found (Figs. 8(b) and 8(c)).
3.2.3. Phase Identification by X-ray DiffractionThe X-ray diffraction analysis results of the residues of the as-aged weld metals after electrochemical extraction are shown in Fig. 9. These results are essentially in agreement with the TEM observation results. As shown in Fig. 9(a), for the C1 weld metal aged at 600°C, when the aging time was chosen to be 150 h, only M23C6 formed. When the aging time was higher than or equal to 500 h, both M23C6 and σ phase precipitated, and as the aging time increased, the M23C6 content decreased, while the σ phase content increased. Furthermore, when the aging time was higher than or equal to 500 h, the main part of the precipitates was σ phase.
Effect of aging time on phase types and contents of weld metals with different C contents: (a) C1; (b) C2; (c) C3.
As shown in Fig. 9(b), for the C2 weld metal aged at 600°C, when the aging time was lower than or equal to 300 h, only M23C6 carbides were detected. σ phase began to be found until the aging time reached 500 h. After that, as the aging time increased, the M23C6 precipitate content decreased, while the σ phase content increased. Finally the σ phase became the major precipitate after aged for 3000 hours. Compared with the C1 weld metals aged at 600°C, for obtained the σ phase, more aging time was needed for C2 weld metal at the same aging temperature.
For the C3 weld metal, as shown in Fig. 9(c), the formation of σ phase is inhibited remarkably. After aged at 600°C for 1000 hours, a small amount of σ phase was detected. With the increase of the aging time to 3000 h, the σ phase content increased slightly, correspondingly the M23C6 precipitate content decreased a little. The analysis results of the phase types and contents in the as-aged weld metals with different C contents indicated that the increase of the C content promoted the formation of M23C6 carbides and correspondingly inhibited the formation of σ phase.
3.3. Mechanical PropertiesThe influence of aging condition on the tensile and impact properties of the low C (C1), medium C (C2) and high C (C3) weld metals are shown in Figs. 10, 11, 12. For the low carbon weld metal (C1), as shown in Fig. 10, at the aging temperature of 600°C, with the increasing of the aging time, the yield and tensile strengths increased all the time. For the medium carbon and high carbon weld metal (C2 and C3), as shown in Figs. 11 and 12, with the increasing aging time, the yield and tensile strengths first decreased, and then increased. For all the weld metals with different C contents, after aging at 600°C for 500 h, the elongation and impact energy decreased sharply. Further increasing the aging time from 500 to 3000 hours, the elongation and impact energy decreased slowly.
Effect of aging time on the mechanical properties of C1 weld metal: (a) yield strength; (b) tensile strength; (c) elongation; (d) impact energy.
Effect of aging time on the mechanical properties of C2 weld metal: (a) yield strength; (b) tensile strength; (c) elongation; (d) impact energy.
Effect of aging time on the mechanical properties of C3 weld metal: (a) yield strength; (b) tensile strength; (c) elongation; (d) impact energy.
The above results of the microstructure analyses indicated that during the aging process, the rapid precipitation of M23C6 carbide occurred in δ-ferrite firstly owing to the high diffusion rate of C. Once the carbon is depleted by precipitation of M23C6, the residual δ-ferrite transforms into σ phase through eutectoid transformation (δ-ferrite → σ phase + γ) depending on the diffusion of Cr and Mo.3,4,6,24,32) Furthermore, after a long enough aging time, the transformation from M23C6 to σ phase occurred.
The C content has significant influence on the δ-ferrite transformation behavior, with the increase of the C content, after the same aging treatment, the content of M23C6 carbide increased, while the content of σ phase decreased. For the low C and medium C weld metals, δ-ferrite transformed into M23C6 and σ phase successively. However, for high C weld metal, δ-ferrite only transformed into M23C6 carbides owing to the high C content and low δ-ferrite content, and σ phase in high C weld metal was transformed fully from the M23C6 phase.
At the initial stage of the aging treatment, the rapid precipitation of M23C6 carbide occurred in δ-ferrite firstly owing to the high diffusion rate of C. Therefore, after the same aging time, the high C content in weld metal resulted in the high δ-ferrite transformation rate, and high δ-ferrite transformation fraction accordingly. However, for the as-welded high C (C3) weld metal, M23C6 carbides precipitated under the heat treatment effect of subsequent weld pass, which decreased the solid solution C content, and then decreased δ-ferrite transformation rate. Therefore, at the initial stage of the aging treatment, for the weld metals with different C contents, the δ-ferrite transformation rate decreased in the order of medium carbon (C2), high carbon (C3) and low carbon (C1). Correspondingly the δ-ferrite transformation fraction decreased in the same order when the aging time was lower than 50 h. With the increase of the aging time, in the low carbon (C1) and high carbon (C3) weld metals, the diffusion rates and quantities of C, Cr and Mo could ensure the transformation from δ-ferrite to M23C6 and σ phase (low carbon C1 weld metal) and M23C6 (high carbon C3 weld metal) proceeded continuously. Therefore, as the aging time increased, the transformation rates of low carbon (C1) and high carbon (C3) first increased, and then kept stable. Compared with the low carbon (C1) weld metal, the higher C content in the medium carbon (C2) weld metal resulted in more M23C6 carbide precipitation before the formation of σ phase. For the medium carbon (C2) weld metal, when the carbon is depleted by precipitation of M23C6, the consumption of Cr is more than the low carbon (C1) weld metal. Therefore, a certain amount of time was needed for reaching the Cr content critical value for the σ phase formation through diffusion of Cr, which resulted in the almost invariable δ-ferrite transformation fraction among the aging time from 50 to 300 h. Furthermore, the δ-ferrite transformation rate of medium carbon (C2) first increased then decreased and finally increased again. Therefore, when the aging time was higher than or equal to 50 h, and lower than 1000 h, the δ-ferrite transformation fraction decreased in the order of high carbon (C3), medium carbon (C2) and low carbon (C1). Furthermore, when the aging time was higher than or equal to 1000 h, the δ-ferrite transformation in high carbon (C3) weld metal had already completed, and the low carbon (C1) weld metal had higher δ-ferrite transformation fraction than the medium carbon (C2) weld metal until the δ-ferrite transformation finished at the aging time of 3000 h.
The variations of mechanical properties with aging time depended to a large degree on the microstructures at different aging conditions. For the low carbon (C1) weld metal, at the aging temperature of 600°C for 500 hours, the low C content resulted in the formation of a large number of σ phase, which increased the yield and tensile strengths. With the increase of the aging time, the increasing σ phase content improved the strength obviously. For the medium carbon (C2) and high carbon (C3) weld metals, the high C contents have strong solid solution strengthening effect, after aging at 600°C for 500 h, the depletion of the solid solution C due to the precipitation of M23C6 phases deteriorated the strength. When the aging time was increased to 1000 h, formation of σ phase improved the strength to a certain extent. Further increasing the aging time to 3000 h, the increasing σ phase content further raised the strength. Furthermore, the rapid precipitation of M23C6 phase at the initial stage of the aging process deteriorated the elongation and impact energy significantly. After that the slow formation of σ phase continued to decrease the elongation and impact energy slowly.
The above analyses show that during aging process the evolution of microstructure is consistent with the changes of mechanical properties. Furthermore, former researchers3,4,6,24,32) have obtained similar conclusions to the abovementioned results of the microstructure and mechanical properties analyses. All these could prove the validity of the findings in this research.
In this research, the 316H weld metals with three different C contents was aged at 600°C for different aging holding times, and microstructural evolution as well as mechanical properties variations of the weld metals during the aging process were evaluated and analyzed. The following conclusions can be drawn:
• During the aging process, the rapid precipitation of M23C6 carbide occurred in δ-ferrite firstly owing to the high diffusion rate of C. Once the carbon is depleted by precipitation of M23C6, the residual δ-ferrite transforms into σ phase through eutectoid transformation (δ-ferrite → σ phase + γ) depending on the diffusion of Cr and Mo. Furthermore, after a long enough aging time, the transformation from M23C6 to σ phase occurred, which led to the decrease of M23C6 content and increase of σ phase content.
• The C content has significant influence on the δ-ferrite transformation behavior, with the increase of the C content, after the same aging treatment, the content of M23C6 carbide increased, while the content of σ phase decreased. For the low C and medium C weld metals, δ-ferrite transformed into M23C6 and σ phase successively. However, for high C weld metal, δ-ferrite only transformed into M23C6 carbides, and σ phase in high C weld metal was transformed fully from the M23C6 phase.
• The variations of mechanical properties with aging conditions depended to a large degree on the microstructures at different aging conditions. For the low C weld metal aged at 600°C, the low C content resulted in the formation of a large number of σ phase, and as the aging time increased, the increasing σ phase content improved the strength obviously. For the medium C and high C weld metals aged at 600°C, as the aging time increased, first the depletion of the solid solution C as a result the M23C6 precipitation deteriorated the strength, then the formation of σ phase improved the strength. Furthermore, with the increasing of the aging time, the precipitation of M23C6 and σ phase deteriorated the elongation and impact energy.
• The abovementioned findings could provide theoretical and practical guidance for the control of the chemical composition, δ-ferrite content, microstructure and mechanical properties of the 316H austenitic stainless steel weld metal for high temperature and long term service exposure.
We would like to acknowledge the financial support from the Opening Foundation of Key Laboratory of Nuclear Materials and Safety Assessment, Institute of Metal Research, Chinese Academy of Sciences (2019NMSAKF04) and China Institute of Atomic Energy (2016-DGB-I-KYSC-0024).
