Optimization of annealing process for low carbon ferritic stainless steel high frequency straight seam resistance welded pipe

With the rapid development of modern industry driven by the progress of science and technology, large-scale infrastructure construction leads to the increasing demand for stainless steel. Although austenitic stainless steel has absolute advantages in plasticity and corrosion resistance, its high cost limits its large-scale application. Ferritic stainless steel has low cost, low expansion coefficient and high thermal conductivity. Among them, 12% Cr ferritic stainless steel has become the representative of economic stainless steel because of its relatively low Cr content [1,2,3]. CR can combine with oxygen to form a layer of oxide Cr2O3 which is closely combined with the matrix on the surface of the material, thus playing the role of anti-corrosion and anti-oxidation. Therefore, 12% Cr stainless steel is widely used in railway, mining, chemical and automobile industries [4,5].

Contents:

In order to promote the application of ferritic stainless steel pipe in the field of mine belt conveyor and mine anti-corrosion gas pipe, it is urgent to prepare welded pipe with excellent service performance, high accuracy of appearance size and low product cost. In recent years, a lot of detailed work has been carried out and remarkable progress has been made in the field of microstructure control and advanced welding process evaluation of ferritic stainless steel welding joints.
In the aspect of microstructure control of ferritic stainless steel, although there is no phase transformation in the traditional single-phase ferritic stainless steel during welding, the grain in the heat affected zone of the joint is prone to coarsening, which worsens the mechanical properties of the joint [6,7,8]. However, 12% Cr stainless steel will form austenite structure at high temperature, which may change into martensite structure after cooling. Austenite can exist stably in the steel with high ferrite factor, and the content of high temperature residual ferrite in the heat affected zone of ferritic stainless steel joint can be predicted by calculating the ferrite factor of the material [9]. Warmelo et al. [10] systematically summarized the influence of multi pass welding process parameters on the sensitization effect of ferritic stainless steel. Further research showed that 12% Cr ferritic stainless steel has excellent welding performance. Zheng et al. [11] Based on the systematic analysis of the effect of C content on the microstructure and properties of hot-rolled low carbon 12Cr Ni stainless steel, studied the transformation behavior of the microstructure in the HAZ of 12% Cr ferritic stainless steel at high temperature by using a thermal simulation test machine, and analyzed the rule of grain growth [12]. Wang et al. [13] showed that the appearance of martensitic phase and the growth of grains would increase the strength and hardness of 12% Cr ferritic stainless steel, and worsen the plasticity and toughness of the material.
In terms of welding process evaluation, Yang et al. [14] selected ER309L as the welding material, and adopted the non melting inert gas arc welding (TIG) method to obtain a 12% Cr ferritic stainless steel joint with good quality. Zheng et al. [9] studied the rule of microstructure evolution of low carbon ferritic stainless steel in the welding process by means of thermal simulation. During the welding process, when the high temperature heat affected zone enters the δ + γ dual phase zone, austenite will form, part of austenite will change into α ferrite when the temperature drops, while the rest austenite will change into martensite, and the content of martensite formed in the process is related to the ferrite factor. Zhang et al. [15] used metal inert gas welding method to realize the effective welding of t4003 ferritic stainless steel. Taban et al. [16] realized the effective connection of 12% Cr stainless steel through the hybrid welding (plasma + tungsten gas shielded welding) process, and made an effective exploration of the high-efficiency laser welding process [17]. In order to further reduce the production cost and improve the accuracy of tube appearance size, high frequency electrical resistance welding (ERW) provides a new possibility. It is a welding method that uses the proximity effect and skin effect formed by high-frequency current to heat the edge of the rolling plate to the melting state, and then applies the upsetting pressure to make the plate reach the combination between metals [18,19,20]. High frequency straight seam resistance welding does not need filling material, it has the characteristics of fast welding speed and high dimensional accuracy, and can produce large diameter thin-walled steel pipe. In addition, the high-frequency straight seam resistance welding production line has not only high welding efficiency, but also environmental protection and energy saving, and has a broad application prospect in the stainless steel welded pipe industry [21, 22, 23]. Zhang et al. [24] studied the influence of the coil structure of the medium frequency heat treatment of high frequency resistance welding on the heating effect, and revealed the influence law of the angle and distance between the main coil and the side coil on the induction power and heating effect. Zuo and Hou [25] studied the cause of the low temperature Charpy impact toughness reduction in high frequency resistance joints. Yana et al. [26] found that there is a low toughness area in the high frequency welded joint, which needs to be improved by appropriate post weld heat treatment.

Because the low carbon ferritic stainless steel joint bears large compressive stress in the high frequency straight seam resistance welding process, and after rapid heating and cooling process, it forms large residual stress in the joint, which seriously endangers the connection quality of the welded pipe. It is necessary to anneal the joint to eliminate or reduce the residual stress and prevent the initiation of welding cracks. This work mainly studies the influence of annealing process on the structure and properties of low carbon ferritic stainless steel high frequency straight seam resistance welded pipe. By optimizing annealing process to adjust the structure of weld area and heat affected area, the purpose of eliminating residual stress and improving plastic toughness is achieved.

1. Experimental method

The actual composition (mass fraction,%) of the low carbon ferritic stainless steel used in the experiment is: C 0.01, Cr 11.54, Si 0.2, Ni 0.57, Mn 1.12, Nb 0.09, Ti 0.12, Fe margin. Based on tcfe6 database, thermocalc thermodynamic software was used to calculate the equilibrium phase diagram. The high frequency welding machine is arranged between the finishing machine and the extrusion roller. The edge of the V-shaped area to be welded is rapidly heated to the melting state by the skin effect, proximity effect and ring effect of the high frequency current. The edge of the plate is pressed together by a certain degree of compressive stress applied by the extruder. Most of the molten metal is extruded out of the weld with oxides and flows to the upper and lower surface shapes of the weld Forming internal and external burr. At last, the internal and external burr is scraped off with carbide tool to obtain the ERW welding joint of low carbon ferritic stainless steel. A 3 mm × 3 mm diameter rod-shaped sample was cut from the base metal area of low carbon ferritic stainless steel welded pipe. TGA / dsc1 thermal analyzer was used for thermal analysis and test. Firstly, the sample was heated to 1300 ℃ at the rate of 40 K / min, and then cooled to room temperature at the rate of 40 K / min after holding for 5 min. The curve was drawn by the data collected by differential scanning calorimeter (DSC). According to the measured phase transition temperature, a reasonable annealing system was established. The joints were annealed in muffle furnace. The furnace was heated to 650 ℃, 750 ℃, 850 and 950 ℃ at the rate of 15 ℃ / min. after that, the samples were put into the furnace for 3 min. then the samples were taken out and cooled in air.
The cross-section samples of weld area were taken by wire cutting. The phase analysis was carried out by D8 advanced X-ray diffractometer (XRD). The Cu target was tested with a voltage of 40 kV, a current of 40 Ma and a test angle of 30 ° ~ 90 °. Prepare the metallographic samples of each area of the joint, polish them with No. 280, 600, 1000, 1500 and 2000 sandpaper in turn, then mechanically polish them with a diamond polishing agent with a particle size of 2.5 μ m, and finally soak them in FeCl3 / HCl corrosion solution (20 ml HCl + 100 ml H2O + 5 g FeCl3) for 1 min. The microstructure of the joint was observed by using the model DMI 8 optical microscope (OM) and the model s4800 scanning electron microscope (SEM). Carbon was sprayed on the surface of the metallographic specimen in the weld area by magnetron sputtering. The prepared carbon film was immersed in FeCl3 / HCl solution for 30 min. after the carbon film fell off, it was fished out with Cu net and cleaned in distilled water and alcohol respectively. Jem-2100f transmission electron microscope (TEM) was used to observe the morphology of precipitates in the weld area. The working voltage was 200 kV. The types of precipitates were determined by EDS, HRTEM and saed.
The Vickers hardness tester duramin-a300 was used to test the microhardness of each area of the joint. The load was 200 g and the holding time was 10 s. The tensile specimen with a standard distance of 30 mm was processed with the weld as the center. The tensile test was carried out at room temperature with GNT 100 universal testing machine, and the tensile rate was 0.03 mm / s. In addition, 5 mm × 10 mm × 55 mm impact specimen is processed, and V-groove is opened at the weld position. The impact energy of the weld is tested at 0 ℃ by pit452c-1 pendulum impact tester.

2. Experimental results and discussion

2.1 Low carbon ferritic stainless steel ERW welded pipe
2.1.1 Low carbon ferritic stainless steel base metal
The equilibrium phase diagram of experimental low carbon ferritic stainless steel calculated by thermo Calc thermodynamic software is shown in Figure 1a. During the heating process, the low carbon ferritic stainless steel passes through α – ferrite single phase zone, α – ferrite and γ – austenite dual phase zone, γ – austenite single phase zone, γ – austenite and δ – ferrite dual phase zone and δ – ferrite single phase zone. In order to measure the transformation temperature of α – ferrite to γ – austenite more accurately, DSC was used to test the low carbon ferritic stainless steel. As shown in Fig. 1b, there are two peaks in the heating curve, the first peak is Curie point, and its temperature is about 720 ℃. Then the peak of transformation from α – ferrite to γ – austenite appeared, and the initial transformation temperature was about 830 ℃. The peak in the cooling curve is the martensitic transformation peak, and the martensitic transformation temperature is about 450 ℃.
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Fig.1 Calculation of phase diagram (a) and DSC curves (b) of the low-carbon ferritic stainless steel

2.1.2 Initial structure of joint
During the high frequency straight seam resistance welding, the welding interface of the pipe is heated to about 1420 ℃, which is in the delta ferrite phase region. Because δ ferrite is in viscoplastic state at high temperature, its fluidity is good. With the extrusion roller exerting greater pressure stress on the welding interface, δ ferrite is extruded out of the weld and flows to the upper and lower surfaces of the weld to form internal and external burrs, so the interface temperature retained in the weld is about 1100 ℃, which is in the austenite phase area. The burr was scraped by cemented carbide cutter in the following process, and it was found that the microstructure of the weld area of the joint was α ferrite and martensite formed by the transformation of austenite after air cooling. The microstructure of the low carbon ferritic stainless steel welded pipe joint is symmetrical on both sides with the weld as the center, and the macromorphology of one side is shown in Figure 2. According to its microstructure, it can be divided into weld zone, HAZ 1, HAZ 2 and base metal zone.
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Fig.2 OM image of welded joint (HAZ—heat affected zone)

The initial structure of low carbon ferritic stainless steel is α – ferrite. The welding interface undergoes many non-equilibrium phase transformation processes in the high frequency resistance welding process. The structure of the weld area of the joint is more complex. In order to determine the phase composition in the weld area of the joint and analyze whether there is residual austenite in the weld area, the cross-section sample of the weld area of the joint is taken for XRD test, and the results are shown in Figure 3. In the figure, the peaks at 44 °, 65 ° and 82 ° are typical α – ferrite three strong peaks. There are no other peaks within the test angle range. It can be seen that the weld area is still α – ferrite phase and there is no retained austenite. After annealing at different temperatures, the XRD spectrum of the annealed weld area is the same as that of the original weld area, which shows that the phase of the annealed weld area is still α ferrite.

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Fig.3 XRD spectra of weld zone after annealing at different temperatures

FIG. 4A shows the microstructure of the weld zone in the joint, which is composed of α ferrite and martensite. The weld zone is in the austenite phase during the heating process, and the austenite grains are coarsened seriously at higher temperature. Due to the rapid heating rate and short holding time, the element distribution in austenite is not uniform, and the transformation into a-ferrite and martensite structure after air cooling. The microstructure of the HAZ adjacent to the weld is also α ferrite and martensite, but its grain size is relatively small, as shown in Fig. 4B. The temperature of high temperature heat affected zone is lower than that of weld zone, so the grain size of austenite is smaller than that of weld zone. The welding in this area bears compressive stress, which leads to the irregular shape of α – ferrite and martensite. The microstructure of the low temperature heat affected zone is shown in Figure 4C, which is composed of α ferrite and a small amount of martensite. When welding is heated, the temperature in this region is low, only part of α ferrite is transformed into austenite, and then α ferrite and martensite are formed in the cooling process, and the rest of α ferrite remains at room temperature. The base metal area of the joint is a single α – ferrite structure (Fig. 4D), which has not changed during the welding process.

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Fig.4 OM images of microstructure of welded joint
(a) weld zone (b) HAZ 1 (c) HAZ 2 (d) base metal
2.2. Effect of annealing temperature on Microstructure and properties of joint
2.2.1 Microstructure of annealed joints at different temperatures
Fig. 5 shows the annealed structure of the weld area after holding for 3 min at different temperatures. When the annealing temperature (650 ℃) is lower than the austenite transformation temperature, there is no obvious change in α ferrite in the weld area, as shown in Fig. 5A. Due to the low annealing temperature and short holding time, only a small amount of martensite decomposes into a mechanical mixture of ferrite and carbide during annealing. Fig. 5B shows the microstructure of the weld area annealed at 750 ℃. With the increase of annealing temperature, the carbides formed by the decomposition of martensite coarsened, while the α – ferrite did not undergo phase transformation. When the annealing temperature is 850 ℃, more martensite will be decomposed, and the larger carbide will dissolve slowly when heated, so there are not fully dissolved carbide particles in the room temperature structure. At the same time, a small amount of acicular structure was observed in the polygonal α ferrite, as shown in Fig. 5C, which shows that part of α ferrite in the weld structure undergoes austenite transformation in the process of heat preservation at 850 ℃, and forms α ferrite and granular bainite structure after air cooling. When the annealing temperature is increased to 950 ℃, the condition of complete annealing can be met when the temperature above AC3 (austenite transformation end temperature) is reached. The microstructure in the weld area is all austenitized. The shorter holding time makes the equiaxed austenite grains keep smaller size. After rapid cooling, the microstructure is transformed into α ferrite and granular bainite (Fig. 5d).
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Fig.5 OM images of microstructures in weld zone after annealing at different temperatures for 3 min
(a) 650 ℃ (b) 750 ℃ (c) 850 ℃ (d) 950 ℃
The carbon film replica was prepared by magnetron sputtering, and the precipitated phase particles were analyzed by TEM. No precipitates were observed in the weld area samples without annealing treatment, but nano precipitates were found in the annealed samples, and their morphology was mostly square or nearly round, as shown in Fig. 6A. EDS results (Fig. 6b) show that the main alloy elements in the precipitates are Nb and Ti, except for C. A small size precipitate particle was selected for high-resolution characterization (Fig. 6c). The crystal surface spacing was measured to be 0.22 nm, which was consistent with the {200} crystal surface spacing of MX type precipitate. Because only one direction crystal surface of precipitate can be observed, and the type of precipitate can not be completely determined, a larger particle is selected for saed analysis (Fig. 6D). After the diffraction spot is calibrated, combined with the EDS spectrum and HRTEM results, it is determined that this kind of nano scale precipitate is rich in Nb and Ti MX phase.

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Fig.6 TEM (a), EDS (b), HRTEM (c) images and SAED pattern (d) of precipitated phase in as-annealed weld zone (d-interplanar spacing)

Figure 7 shows the TEM image of the second phase precipitated in the weld area after annealing at different temperatures. After annealing at 650 ℃, the distribution of MX phase particles in the weld area is relatively scattered, and the average size of particles is small, about 15 nm, as shown in Fig. 7a. After annealing at 750 ℃, the number and size of MX phase particles in the weld area increased, and the average size was about 21 nm (Fig. 7b). Figure 7C shows the morphology of MX precipitated in the weld area after annealing at 850 ℃. Compared with the sample annealed at 750 ℃, the number of MX phases is reduced and some particles are combined with each other. When the annealing temperature is 950 ℃, the density of MX phase particles in the weld area is the lowest and the size is the largest, the average size increases to about 46 nm, and the size of the largest particles can reach 91 nm, as shown in Fig. 7d. The size of MX phase in weld zone increases with the increase of annealing temperature, while the number of MX phase decreases gradually. Some small-sized particles dissolve at higher annealing temperature. MX phase particles need to reduce surface area to reduce energy in order to be stable at high temperature. Therefore, some particles combine with each other to form large-sized particles. It can be seen that with the increase of annealing temperature, the coarsening behavior of MX phase in the weld zone follows the Ostwald ripening mechanism [27,28].
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Fig.7 TEM images of precipitated phase in weld zone after annealing at different temperatures for 3 min (a) 650 ℃ (b) 750 ℃ (c) 850 ℃ (d) 950 ℃
The high temperature heat affected zone is close to the weld zone, the deformation degree of the welding process is large, and the grain is flattened. After rapid cooling, there is a high elastic distortion energy in the high temperature heat affected zone, which increases the driving force of recrystallization and reduces the recrystallization temperature. Fig. 8a is the OM image of the microstructure of the high temperature heat affected zone annealed at 650 ℃. Compared with the original high temperature heat affected zone, the number of irregular α ferrite grains is reduced, and the grain size is significantly reduced. It shows that some large deformation α ferrite grains recrystallize during annealing. Because of the short holding time, the recrystallized grains are not able to form uniform equiaxed grains, so the α ferrite grain size presents uneven distribution at room temperature. As the annealing temperature increases to 750 ℃, the number of recrystallized α ferrite grains increases, and the recrystallized grain size increases, as shown in Figure 8b. The microstructure of the high temperature heat affected zone annealed at 850 ℃ is shown in Fig. 8C. The increase of annealing temperature makes the martensite almost decompose into the mechanical mixture of α ferrite and carbide, so there are more α ferrite grains recrystallized, and the growth process of the recrystallized grains is restrained by the role of carbide particles pinning the grain boundaries. The high temperature heat affected zone is fully annealed at 950 ℃, and the grains are recrystallized into austenite grains. After cooling, α ferrite and granular bainite structures are formed, as shown in Fig. 8D.
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Fig.8 OM images of microstructures in HAZ 1 after annealing at different temperatures for 3 min (a) 650 ℃ (b) 750 ℃ (c) 850 ℃ (d) 950 ℃
The temperature of the HAZ is lower than that of the HAZ, and the content of martensite formed after cooling is less than that of the HAZ. When annealing treatment is carried out under austenite transformation temperature, α ferrite does not change and martensite decomposes, as shown in Fig. 9A. As the annealing temperature increases to 750 ℃, the martensite almost decomposes (Fig. 9b). When the annealing temperature is 850 ℃, the martensite will decompose into α ferrite and carbide during the heating process. After reaching a certain temperature, some α ferrite will undergo austenite transformation and form α ferrite and granular bainite structure after air cooling. Therefore, the room temperature structure of the annealed state at 850 ℃ is composed of α ferrite and a small amount of granular bainite, as shown in Fig. 9C. When the annealing temperature rises to 950 ℃, the carbide dissolves and α ferrite transforms into austenite. Due to the short holding time, the grains after complete austenization are too late to grow and keep smaller grain size. Then austenite transforms into α ferrite and granular bainite structure during the cooling process, as shown in Fig. 9D.
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Fig.9 OM images of microstructures in HAZ 2 after annealing at different temperatures for 3 min (a) 650 ℃ (b) 750 ℃ (c) 850 ℃ (d) 950 ℃
During the welding process, the base metal area is affected by a small amount of heat transferred from the weld area, and the structure is basically unchanged, which is composed of α ferrite. Fig. 10a and b show the microstructure of the base metal annealed at 650 ℃ and 750 ℃, respectively, which is α – ferrite. When the annealing temperature rises to 850 ℃, the grain boundary of α – ferrite is zigzag, which indicates that part of α – ferrite at the grain boundary undergoes austenite transformation during heating, and then forms granular bainite structure during cooling, as shown in Figure 10C. The transformation degree of austenite gradually increases with the increase of annealing temperature. When heated to 950 ℃, the structure in the base metal area is fully austenitized, and α ferrite and granular bainite structure are formed after cooling, as shown in Fig. 10d. Compared with 850 ℃, the grain size of the base metal annealed at 950 ℃ is smaller, and the content of granular bainite is higher.
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Fig.10 OM images of microstructures in base metal after annealing at different temperatures for 3 min (a) 650 ℃ (b) 750 ℃ (c) 850 ℃ (d) 950 ℃
2.2.2 Mechanical properties of annealed joints at different temperatures
Fig. 11 shows the microhardness changes of each region of the joint before and after annealing at different temperatures. The hardness of the weld area without annealing treatment is the highest, which can reach about 315 HV. With the increase of the distance from the weld, the hardness of the high temperature heat affected zone and the low temperature heat affected zone decreased in turn, and the hardness of the base metal zone decreased to the lowest, about 177hv. The microstructure in the weld zone, the high temperature heat affected zone and the low temperature heat affected zone are all composed of martensite and α – ferrite. The martensite content in the high temperature heat affected zone and the low temperature heat affected zone decreases gradually, while the microstructure in the base metal zone is all α – ferrite, so the hardness of each zone decreases gradually. After annealing at 650 ℃, the hardness of weld zone, high temperature heat affected zone and low temperature heat affected zone decreased slightly. The annealing temperature is low and the holding time is short, only a small amount of martensite decomposes, so the hardness of each zone is still high. However, when the annealing temperature is increased to 750 ℃, the hardness of the weld zone, the high temperature heat affected zone and the low temperature heat affected zone are all significantly reduced, which is due to the increase of annealing temperature that more martensite is decomposed. Among them, the microhardness of the weld zone and the high temperature heat affected zone is significantly higher than that of the base metal zone, which are 258 and 226 HV respectively, indicating that the martensite is not completely decomposed. However, the martensite in the low temperature heat affected zone is almost completely decomposed, and its hardness is equivalent to the base metal zone of the original joint. After annealing at 850 ℃, the hardness change trend of each zone in the joint is different from that after annealing. The hardness of the weld zone, the high temperature heat affected zone and the low temperature heat affected zone are basically the same, which shows that most of the martensite in the above three zones are decomposed during annealing, and the structure of the annealed zones is similar. The hardness of α – ferrite is higher than that of α – ferrite in the original base metal, which is due to the austenite transformation of some α – ferrite during annealing and the formation of a small amount of granular bainite after cooling [28]. When the annealing temperature reaches 950 ℃, the hardness of each zone in the joint is about 260 HV. In the process of heating, the martensite in each zone is completely decomposed, then the carbide is dissolved, α – ferrite is transformed into austenite, after holding for 3 minutes, each zone is fully austenitized, and after cooling, α – ferrite and granular bainite are formed.
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Fig.11 Microhardness distribution from weld zone to base metal after annealing at different temperatures for 3 min
The tensile strength and elongation of the joints treated at different annealing temperatures are shown in Figure 12. The tensile strength of the joint without annealing treatment is 546 MPa, the elongation after fracture is 15%, and the fracture position is in the base metal area, which indicates that the welding quality is good, and the strength of the weld area in the joint is higher than that in the base metal area. After annealing at 650 and 750 ℃, the fracture position of the tensile specimen is also in the base metal area, and the tensile strength and elongation of the joint have no obvious change compared with the original joint. However, the tensile strength of the joint annealed at 850 ℃ decreased to 508mpa, and the elongation increased to 23.5%. The tensile specimen also broke in the base metal area. In the process of annealing at 850 ℃, the α – ferrite grain boundary in the base metal area becomes zigzag, which improves the coordination of the deformation among the grains, so that it can bear a large amount of deformation and improve the plasticity of the base metal. When the annealing temperature is 950 ℃, the tensile strength of the joint is the highest (708 MPa), and the elongation is reduced to 12.5%. This is because the microstructure of annealed joints is composed of a large amount of granular bainite and a small amount of α – ferrite. The strength of granular bainite is higher than that of α – ferrite, so the formed granular bainite can greatly improve the tensile strength of joints, but at the same time, it will deteriorate the plasticity of joints.
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Fig.12 The tensile strength and elongation of welded joint after annealing at different temperatures for 3 min
Figure 13 shows the fracture morphology of the tensile specimen. After annealing at 650 ℃, the fracture surface of the joint is composed of dimple and many small planes of quasi cleavage, on which there are cleavage steps and river patterns. Such quasi cleavage fracture belongs to transgranular fracture, as shown in Fig. 13A. After annealing at 750 ℃, the number of dimples in the fracture increased (Fig. 13b), while the quasi cleavage plane decreased and the plasticity improved. When the annealing temperature is 850 ℃, the fracture surface of tensile specimen is composed of two kinds of dimples. There are large dimples scattered in the middle of small and dense dimples. As shown in Fig. 13C, this kind of dimple fracture belongs to ductile fracture with good plasticity. With the increase of annealing temperature to 950 ℃, the number of small and medium-sized dimples in the fracture is greatly reduced (Fig. 13D), and a large number of quasi cleavage planes appear at the same time, which is a typical quasi cleavage fracture in transgranular fracture. It can be seen that the granular bainite formed in the process of annealing at 950 ℃ reduces the plasticity of the joint.
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Fig.13 SEM images of tensile fracture surfaces of welded joint after annealing at different temperatures for 3 min (a) 650 ℃ (b) 750 ℃ (c) 850 ℃ (d) 950 ℃
In the tensile test, the fracture position of each sample is in the base metal area, which indicates that the welding quality of the joint is good. Because the strength of weld zone, high temperature heat affected zone and low temperature heat affected zone are higher than that of base metal zone, each group of tensile tests can only reflect the properties of base metal zone. Combined with the actual use of welded pipe, the toughness of weld is the most important index for its stable service at low temperature. The impact test is carried out on the weld area of the joint after different annealing temperature treatment at 0 ℃, and the impact value is shown in Figure 14. The impact value of the original joint weld at 0 ℃ is about 0 J, which is due to the high residual stress in the original weld and the large amount of brittle and hard martensite structure, resulting in the poor toughness. After annealing at 650, 750 and 850 ℃, the impact value of the weld still hasn’t been improved, which indicates that the residual stress in the weld area can’t be completely eliminated by annealing at the above temperature, and there may still be the undissolved martensite structure in the weld area. When the annealing temperature increased to 950 ℃, the impact value of the weld area increased to 23.2j. The martensite structure in the original weld area decomposes completely in the heating process, transforms into austenite in the heat preservation process at 950 ℃, and then forms α ferrite and granular bainite structure in the cooling process. At the same time, the residual stress is fully released, which improves the toughness of the weld joint.
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Fig.14 Charpy impact energy of welded joint after annealing at different temperatures for 3 min
The fracture morphology of the impact specimen annealed at 650 ℃, 750 ℃ and 850 ℃ is shown in Fig. 15A ~ C respectively. It can be seen that the sections are relatively flat, which shows that there is almost no plastic deformation before fracture. River pattern, tongue pattern and fan-shaped pattern can be observed in the section of three groups of samples, which are typical features of cleavage fracture, indicating that the three groups of samples have poor toughness and cleavage fracture. Fig. 15d shows the fracture morphology of the sample annealed at 950 ℃. The fracture surface is uneven and consists of dimples with different sizes and depths. Small and shallow dimples are densely distributed around the large and deep dimples, which belong to dimple fracture. The results show that the toughness of the mixed structure of α – ferrite and granular bainite is good, and the residual stress in the joint is basically eliminated.
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Fig.15 SEM images of impact fracture surfaces of welded joint after annealing at different temperatures for 3 min (a) 650 ℃ (b) 750 ℃ (c) 850 ℃ (d) 950 ℃
In conclusion, the tensile test results mainly reflect the properties of the base metal area of the joint, while the impact test results mainly reflect the properties of the weld area of the joint. After annealing at 850 ℃, the ductility of the base metal area is better than that of the weld area. After annealing at 950 ℃, the toughness of the weld zone is better, but the plasticity of the base metal is worse. The residual stress and martensite structure in the weld area can not be completely eliminated after annealing at 850 ℃. Only when the annealing temperature is raised to 950 ℃, the residual stress in the weld area can be completely eliminated, and the martensite structure in the weld area can be completely decomposed, and then transformed into granular bainite and α ferrite structure, so as to improve its toughness. After annealing at 850 ℃, the microstructure of the base metal is still α – ferrite. When the annealing temperature rises to 950 ℃, the microstructure of the base metal changes into granular bainite and α – ferrite, resulting in the decrease of plasticity.

3. Conclusion

The peak temperature of high frequency straight seam resistance welding process can reach 1420 ℃, which can make the microstructure of low carbon ferritic stainless steel weld area change into austenite completely, and the austenite grain is distorted under the action of compressive stress. After air cooling, there is a large residual stress. At room temperature, the weld area is composed of martensite and α ferrite, with poor toughness. After annealing at 650, 750 and 850 ℃, the toughness of the joint is still poor, but when the annealing temperature rises to 950 ℃, the martensite structure in the weld zone of the joint decomposes completely, and the microstructure in the annealed weld zone transforms into α ferrite and granular bainite, at the same time, the residual stress is fully released, which improves the toughness of the low carbon ferrite high frequency resistance welded pipe joint At 0 ℃, the impact value increased to 23.2j, but the plasticity of the base metal decreased.

Source: China Welded Pipe Manufacturer – Yaang Pipe Industry Co., Limited (www.steeljrv.com)

(Yaang Pipe Industry is a leading manufacturer and supplier of nickel alloy and stainless steel products, including Super Duplex Stainless Steel Flanges, Stainless Steel Flanges, Stainless Steel Pipe Fittings, Stainless Steel Pipe. Yaang products are widely used in Shipbuilding, Nuclear power, Marine engineering, Petroleum, Chemical, Mining, Sewage treatment, Natural gas and Pressure vessels and other industries.)

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Summary
optimization of annealing process for low carbon ferritic stainless steel high frequency straight seam resistance welded pipe - Optimization of annealing process for low carbon ferritic stainless steel high frequency straight seam resistance welded pipe
Article Name
Optimization of annealing process for low carbon ferritic stainless steel high frequency straight seam resistance welded pipe
Description
With the development of economy and technology, the application of ferritic stainless steel is becoming increasingly wider. 12Cr ferritic stainless steel has low carbon equivalent and good weldability, and it can not only be applied to a variety of conditions, but also reduce the production cost. High frequency longitudinal resistance welding is an advanced welding technology with high quality and efficiency. In this work, low-carbon ferritic stainless steel pipe has been joined successfully by high frequency longitudinal resistance welding. Microstructure characteristics and mechanical properties of the stainless steel pipe joint after annealing at different temperatures for 3 min were investigated by OM, SEM, TEM and mechanical testing. In the process of high frequency longitudinal resistance welding, the weld zone was heated quickly to a high austenization temperature which led to a coarse grain structure in this zone assisted by high pressure. The weld zone presented martenite and ferrite microstructure with irregular grain. As a result, the hardness of the weld zone reached 315 HV and the impact energy dropped to near zero. After annealing at 950 ℃ for 3 min, the decomposition of martensite was the main reason of the decrease of hardness (260 HV) in weld zone. The microstructure of weld zone was composed of ferrite and bainite, resulting in the increase of impact energy from 0 to 23 J.
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