Effect of heat treatment and cooling rate on microstructure and properties of T92 welded joint

The mixed corrosion liquid could lead to the martensite color becoming darker than the ferrite coloring. Figure 5 presents the deposited metal microstructure. It can be observed from Figure 5 that the tempering microstructure of the deposited metal microstructure was uniform. The microstructure of martensite in the weld metal was relatively clear. A high amount of martensite lath presented which was consistent with the crystalline direction.

Fig. 5.

The phase change points A_c1 and A_c3 of the deposited materials were 773°C and 850°C, as measured through thermal expansion testing. The sample heat treatment holding temperatures were 740°C and 760°C, which were below the A_c1 point. The heat treatment holding temperature was 800°C, which was between A_c1 and A_C3. Figures 6, 7, and 8 present the typical metallurgical structures of the deposited metal under various holding temperatures and times. Following lapping and polishing, the deposited metal was corroded with a mixed solution of 5 g CuCl₂, 30 ml HCl, 25 ml alcohol, and 30 ml H₂O. The ferrite was colored, and the martensite color became deep. It can be observed in Figure 6 that the tempering time of 2 h could not guarantee the sufficient tempering of the deposited metal. Due to the lower temperatures, the incompletely tempered martensite structure existed in the deposited metals, such as the black area in Figure 6a. The room temperature microstructure was composed of tempered martensite, non-tempered martensite, and the second phase formed during tempering. As the tempering time increased, the tempering microstructure of the deposited metal became uniform. Its microstructure at room temperature was a single-tempered martensite and the uniformly distributed second phase [14].

Fig. 6.

Fig. 7.

Fig. 8.

As seen in Figure 7, as the heat treatment temperature increased, the tempering microstructure time of the deposited metal to reach the uniformity time decreased. This meant that the tempering at low temperature required a longer time to temper the homogenized structure, while the required tempering time decreased accordingly as the tempering temperature increased. As can be observed in Figure 8, the martensite lath began to decompose when the tempering temperature was 800°C. Simultaneously, the segregation and dissolution of the carbonitride decreased the deposited metal strength.

The tempering-induced fine dispersed second phases precipitated in the deposited metal. The second phase precipitated with unequal size distribu-tion in the original austenite grain boundaries, as well as in the lath bundle boundaries and in the crystal. However, the partition of the non-tempered precipitated phases was small. Figure 9 presents the distribution of the second phases in the tempered and non-tempered martensite microstructure. It can be seen in Figure 9 that dispersed fine second-phase precipitates existed in the tempering zone of martensite, while the second phase was mainly composed of the M23C6 and MX phases. The carbonitrides in the non-tempering zone (black areas in Figure 6) were the solid solution treatment phases.

Fig. 9.

Figure 10 presents the impact absorption work of different tempering temperatures and holding times. With the increase of heat treatment temperature, the toughness of the same deposited metal increases, obviously. With the decrease of heat treatment temperature, the average impact toughness shows an increasing trend, but the increasing trend slows down gradually. At the same temperature, with the increase of time, the impact energy at the three temperatures shows three obviously different trends. At low temperature, it increases slowly and then rapidly. The change trend is not obvious at medium temperature. At high temperature, it shows a gradual increasing trend.

Fig. 10.

The toughness value of the deposited metal after heat insulation for 2 h at 760°C was equal to the toughness value after heat insulation for 8 h at 740°C.

It could be observed that the heat treatment temperature increase could improve the working efficiency. The excess temperature in heat treatment was also unfavorable. However, the excess temperature in heat treatment could increase the weld metal toughness and severely favor the martensite decomposition. The carbonitrides segregation in the deposited metal highly reduced the weld metal strength. The aging temperature of the deposited metal must be strictly controlled to ensure the high temperature strength of the deposited metal.

The deposited metal hardness under different heat treatments was measured through the Vickers hardness measurement, and the load was 5000 g. For each heat treatment process, the hardness of the deposited metal sample is shown in Figure 11. The hardness acquired from different heat treatment processes had a certain corresponding relationship with toughness. With the increase of heat treatment temperature, the hardness of the same deposited metal is obviously reduced. With the decrease of heat treatment temperature, the decrease of hardness value is more obvious. When the temperature is reduced to a certain extent, the decrease of the hardness value is no longer obvious.

Fig. 11.

The impact fracture of the deposited metal under different heat treatment processes was measured and the results are presented in Table 4 and Figure 12. It could be observed that the lower the fibrous fracture rate was, the higher the brittleness tendency was, as well as the higher the toughness was and the better the average value was. The fibrous area proportion had a good correspondence with the impact work height, while its fibrous fracture rate increased as both the heat treatment temperature and the holding time increased. Under the condition of 740°C tempering temperature, the fibrous fracture rate was below 40% and the brittleness tendency was quite severe, when the samples were heat treated for at least 2 h and approximately 4 h. When the holding time increased to 8 h, the toughness apparently increased. At the tempering temperature of 760°C and 800°C, the fiber fracture rate changed slightly. It could be observed that under the tempering temperature of 760°C and 800^°C, the toughness increase was not apparent as the holding time increased.

Fig. 12.

Figures 13, 14, and 15 show the impact fracture morphologies of each heat treatment process. Under different heat treatment conditions, the impact fracture was composed of three zones: the fiber zone, the radiation zone, and the shear lip zone. Due to the differences in fracture properties and impact absorption work values, the proportion of each region differed. As can be seen, under the tempering temperature of 740°C, the proportion of the radiation zone in the impact fracture of the thermal insulation for 2 h and 4 h reached beyond 80% and 60%, respectively, which was attributed to brittle fracture. As can be seen in Figure 13a and 13b, the fracture patterns were mainly river and scallop patterns, accompanied by a torn edge. The fracture mechanism was quasi-cleavage transgranular. When the holding time increased to 8 h, the proportion of fiber area reached beyond 80%, while the fracture mechanism changed from the quasi-cleavage fracture to the ductile fracture. The fracture morphology, presented in Figure 13c, was mainly dimple morphology. It could be observed from Figures 14 and 15 that the fiber fracture rate exceeded 80%, at the tempering temperature of 760°C, while the ductile fracture mode was mainly manifested as dimple type.

Fig. 13.

Fig. 14.

Fig. 15.

Figure 16 presents the metallographic structure of the weld, when the simulated t_8/5 times were 5 s, 10 s, 20 s, and 70 s. The heat treatment was not considered on the simulated welded joints. The corrosion solution mixed with 5 g FeCl₃, 15 ml HCI, and 80 ml H₂O was utilized to corrode the microstructure of the deposited metal. As can be observed in Figure 16, the organizations were all martensite tissues, when the simulated t_8/5 cooling rate changed from 5 s to 70 s, which demonstrated that the weld metal could obtain a martensitic structure under a wide heat input. Generally, as the welding heat input increased, the grains would become coarse, the toughness would decreased, and the hardness would increase.

Fig. 16.

The curves of toughness and hardness with the t_8/5 cooling time are presented in Figures 17 and 18, respectively. It can be observed that when the t_8/5 was lower than or equal to 20 s, the toughness of the weld joint decreased as the t_8/5 increased. However, the weld toughness substantially increased when the t_8/5 reached up to 70 s. The trend of hardness variety was more consistent with the trend of toughness variety. When the t_8/5 was lower than or equal to 20 s, the hardness increased, while the weld hardness decreased significantly at t_8/5=70 s. The occurrence of this phenomenon and the Precipitation Law of the second phase in the weld had a certain relation to the welding heat input change. The dissolution and precipitation of the second phase in the iron matrix was a reversible chemical reaction process [15]. The number of elements involved in the second phase in a solid solution matrix and the amount present in the second phase tend to be balanced subsequently to prolonged temperature preservation. The temperature change under certain chemical compositions of the steel, the amount of solid solution, and the content of the second phase will also change [16].

Fig. 17.

Fig. 18.

The main secondary phase precipitation strengthening phases in the T92 steel were the M₂₃C₆ and MX phases, and the stable MX phase (M=V, Nb; X=C, N) was mainly distributed within the sub-grain, while the M₂₃C₆ phase enriched in Cr mainly precipitated at the grain and sub-grain boundaries. The MX phase was relatively stable, while the low growth rates of the MX phases maintained even at high temperature for a long time, which played an important role in the high temperature strength durability of the T/P92 steel. In contrast, the M₂₃C₆ growth under high temperature was apparent, while B was added to restrain the M23C6 growth rate at high temperature. In addition, the T/P92 steel would produce a brittle Laves phase at high temperature for a long tempering time, which precipitated along the original austenite grain boundaries, the martensite blocks, and the lath boundaries [17]. The effect of the Laves phases on the high-temperature performance of the T/P92 steel depends on their distribution and size. When the coarsening rate of the Laves phase is low, it has a beneficial effect on the high-temperature creep strength improvement. When the high-temperature aging and creep deformation increased in duration, the coarsening rate would increase, while the relative creep strength effect of the Laves phases would change to the harmful [18]. For the aforementioned three types of main precipitates in the welding cooling process, the precipitation of the M₂₃C₆ and MX phase contents increased as the temperature decreased under the equilibrium state, as presented in Figure 16b. The welding cooling process was a non-equilibrium state. When the cooling rate was higher, the precipitation of the second phases were inhibited, while most second phases became solid-solutes in the matrix. When the cooling rate was lower, the precipitation time of the second phase was longer and the amount of precipitation was relatively higher.

In Figure 16, when the t_8/5 was lower than or equal to 20 s, the cooling rate increased, the second phase precipitation was inhibited, so the second phase could be ignored. As the heat input increased, the toughness decreased, and the hardness increased. When the t_8/5 was equal to 70 s, the second phase of precipitation was relatively increased in amount and the role of second phases in the matrix had the following two points [19]. Firstly, the C and N interstitial solution strengthening elements with relatively high brittleness vector would precipitate in the relatively low-sized MX phase of the brittle medium, which reduced the matrix stress. Secondly, the precipitated second phase had the grain boundaries pinning role, which could inhibit the grain growth and alleviate the heat input caused by the grain growth to a certain extent. Therefore, these two points would present the rebound phenomenon of toughness, while the hardness value decreased when the t_8/5 was equal to 70 s.

Figure 19 presented the fracture at different cooling speeds. The fracture was composed of the cleavage fracture mode under the different cooling rates in the thermal simulation state, which was mainly related to the welding joints without heat treatment subsequent to welding. The C and N interstitial solution strengthening elements with relatively high brittleness vector and the forming elements of strong and medium carbon nitrides almost completely dissolved in the matrix in the state of the solid solution under the faster cooling rate following a thermal simulation, while the fracture modes were brittle fractures.

Fig. 19.