Study on the Heat Treatment Process of SXQ500/550DZ35 Steel QLT for the Seat Ring and Guide Vane of a Hydropower Station

Abstract

Because steel plates of 150~300 mm cannot reach the ideal cooling rate under the industrial cooling condition, the development of a super thick plate for the seat ring and guide vane of hydropower stations cannot be successful. In this paper, the influence of the original microstructure and the quenching temperature on the austenite grain size and the analysis of the quenching temperature on the austenite transformation amount were carried out through experiments. The effect of quenching temperature on the content of austenite carbon was calculated by a phase diagram, and then the effect of quenching temperature on the cooling characteristics of the steel was simulated by JMatPro, and the microstructure changes and performance advantages of the QLT heat treatment process were studied and analyzed. Finally, the QLT heat treatment process of SXQ500/550DZ35 steel for the seat ring and guide blade of hydropower stations was researched and formulated (925 °C + 830 °C + 640 °C), and the heat treatment process optimization of a thick steel plate under insufficient cooling conditions during industrial production was solved.

1. Introduction

SXQ500/550DZ35, an extra thick plate used to develop the seat ring and guide vane of a hydropower station (as shown in Figure 1), is a kind of high strength low alloy bainitic steel with a thickness of 150 mm to 300 mm, so the cooling rate from the surface to center is different under the conventional heat treatment process and the cooling rate decreases with the increase in thickness. Considering the carbon equivalent of weldability and the requirement of the low welding crack sensitivity index, the microstructure of the steel plate is mainly granular bainite, so the heat treatment process of quenching + lamellarizing at the dual-phase zone + tempering (Quench + Lamellarizing + tempering hereafter referred to as QLT) is adopted to obtain its high strength and high toughness [1,2].

Sub-temperature is a heat treatment method in which hypoeutectoid steel with an equilibrium or non-equilibrium original structure is heated to the two-phase region of ferrite and austenite and then quenched or isothermally quenched for a certain period of time. It is a strengthening and toughening heat treatment process which uses a ductile phase to refine the microstructure. Sub-temperature treatment can greatly improve the impact toughness, restrain the reversible tempering brittleness, reduce the ductile-brittle transition temperature, prevent cracking and deformation, and solve the problems of the hardenability of large workpieces [3,4].

Sub-temperature quenching can refine grains. The reasons for obvious grain refinement after subcritical quenching (as shown in Figure 2) are as follows. Firstly, there is undissolved ferrite in the two-phase zone, which can hinder the migration of austenite grain boundaries and inhibit its grain growth. According to related research, after subcritical quenching, the area of ferrite and austenite grain boundary is 10 to 50 times that of the austenite grain boundary obtained by conventional complete quenching [3]. Secondly, the quenching temperature in the two-phase region is relatively low, the atomic diffusion coefficient is small, and the grain boundary migration speed is slow. Thirdly, the original microstructure of the 920 °C quenched sample belongs to a non-equilibrium bainite structure, in which there are many substructures, including a sub-strip, subunit, ultra-fine subunit, and so on. Additionally, there are a considerable number of dislocations. When heated in the two-phase region, some substructures and dislocations are retained as the nucleation centers of austenite nucleation, which increases the rate of austenite nucleation [5].

The undissolved ferrite in the steel hinders crack propagation. The sub-temperature quenching temperature is lower than that of complete quenching, and a part of the fine undissolved ferrite is retained in the steel. Ferrite with low hardness and good plasticity can prevent stress concentration and hinder crack propagation, so it can improve the low-temperature toughness of steel. Before fracture, the crack propagates in the plastic zone at the tip of the material. In dual-phase steel, when the radius of the plastic zone is greater than the grain radius, the crack propagates along the softer phase, and when the ferrite and bainite/martensite are needle-like, the brittle phase bainite/martensite is separated to the maximum extent by the plastic phase ferrite, so the crack propagates not only through bainite/martensite but also through ferrite. Due to the large amount of plastic deformation of ferrite before fracture, which consumes more energy, the toughness increases. In addition, when the crack propagates to the ductile phase in undissolved ferrite, the crack propagation is blocked or forced to change to a direction with less resistance and less harm, such as delamination, to relax the energy and improve the toughness [6].

Improving the distribution of harmful impurity elements:

The results show that reversible tempering brittleness is caused by the depolymerization of harmful impurity elements (P, Sn, Sb, S, etc.) on the grain boundaries and microcracks of the original austenite. However, alloying elements such as Ni and Cr not only promote the segregation of impurity elements but also self-segregation, which reduces the fracture strength of grain boundary and produces temper brittleness. Sub-temperature quenching can improve the distribution of harmful impurity elements in steel. First, when the steel is heated in the two-phase zone, the grain size is refined, the grain volume decreases, the ratio of surface area to volume increases, the total grain boundary area increases, and the content of harmful impurity elements (P, Sn, Sb, S, etc.) in the unit area decreases, which effectively reduces the segregation of harmful impurity elements. Second, the microstructure after intercritical quenching is austenite and ferrite. The solubility of impurity elements in ferrite is much greater than that in austenite, which can enrich and purify impurity elements and effectively restrain the segregation of harmful impurity elements on austenite grain boundaries [5].

Some studies have shown that low-carbon granular bainitic steel has greater strength, toughness, and fatigue strength compared to ferrite + pearlite steel without significantly reducing plasticity. The tensile strength of granular bainitic steel (${\mathsf{\sigma }}_{\mathrm{B}}\mathrm{g}$) increases with the increase in the total amount of M/A island, and there is an empirical relationship as follows:

$${\mathsf{\sigma }}_{\mathrm{B}}\mathrm{g}=\mathrm{a}+\mathrm{b} T \%.$$

(1)

In the formula, a and b are constants. In addition, grain boundary, microstructure type in the “island”, and fine structure also have a certain influence on the strength of granular bainite steel. For the bainitic–ferrite matrix, the M-An island is a kind of low plastic strengthening phase. Properly reducing the total amount of island structure or the chord length of the island is beneficial to improve the toughness of granular bainitic steel. When the island size is less than 1 μ m, it has almost no effect on the low-temperature impact toughness of the steel. Increasing the cooling rate, reducing the austenite grain size, or increasing the Mn content in the steel can reduce the island chord length while reducing the C content can reduce the total island amount, and increasing the cooling rate is beneficial to the refinement of the island structure [7]. Therefore, reducing the C content and controlling the heat treatment process of the steel will be beneficial to the improvement of the toughness of the granular bainitic steel.

The parameters of granular bainite structure, that is, the shape, quantity, size, and distribution of the “island”, are the decisive factors for the strength and toughness of granular bainitic steel [8]. The main factors affecting the strength and toughness of granular bainite are the total number of islands and the island morphology (chord length, island spacing). With the increase in the total number of islands, the chord length and island spacing decrease, but the strength increases. With the decrease in the total number of islands, the chord length of islands decreases, the island spacing increases, and the toughness increases. In a certain range of components, microstructure with good strength and toughness can be obtained by properly controlling the number and size of islands.

Since 150~300 mm steel plates cannot reach the ideal cooling rate under industrial cooling conditions, the influence of the original structure and the sub-temperature quenching temperature on the austenite grain size, the analysis of the transformation amount of austenite by the sub-temperature quenching temperature, and the influence of the sub-temperature quenching temperature on the austenite carbon content were roughly calculated by the phase diagram. Furthermore, JMatPro was used to simulate its influence on the cooling characteristics of the steel, and the microstructure changes and performance advantages of the QLT heat treatment process were studied and analyzed. In particular, the heat treatment process was optimized to ensure the uniform microstructure refinement of bainite under the insufficient cooling rate. Finally, the QLT heat treatment process of SXQ500/550DZ35 steel used in the seat ring and guide vane of the hydropower station was studied and formulated, and heat treatment process optimization was solved under insufficient cooling conditions during the industrial production of thick plate steel plates.

2. Experimental Materials and Methods

2.1. Experimental Materials

The materials used are from a 270 mm thick SXQ500/550DZ35 hot rolled steel plate ((structure is granular bainite (GB) + lath bainite (LB) + less ferrite (FLess) + less pearls (PLess), hereinafter referred to as AR state) and a 925 °C quenched steel plate (structure is GB + LB, hereinafter referred to as Q state) produced by the Han Ye Special Steel Industry. The composition is shown in

Table 1. Chemical composition of SXQ500/550DZ35 steel (Wt %).

C	Si	Mn	P	S	V	Ti	Als
0.11	0.12	1.18	0.009	0.001	0.042	0.004	0.019
Cr	Mo	Nb	Ni	Cu	B	N	O
0.506	0.386	0.035	1.380	0.020	0.0003	0.0058	0.005

, and the samples are processed into Φ4 mm × 10 mm and 10 mm × 10 mm × 10 mm.

2.2. Experimental Methods

A DIL805L thermal expansion instrument (TA Instruments, Newcastle, DE, USA) is used to conduct thermal expansion tests on AR and Q samples. The heating rate of samples below 600 °C is 10 °C/s, and the heating rate of samples above 600 °C is 0.05 °C/s. The Ac1 = 744 °C (the temperature at which the pearlite transforms to austenite when heated) and Ac3 = 847 °C (the final temperature at which ferrite is completely transformed into austenite when heated) of the steel are obtained by heating to 770, 810, 830, 850, 870 and 920 °C, respectively. The heat preservation coefficient is 2.2 min/mm. There are two ways of cooling: water quenching to room temperature and cooling at 30 °C/min. Then, after grinding and polishing, the metallographic sample is prepared. The austenite grain boundaries of the samples quenched to room temperature are corroded by a certain proportion of saturated picric acid + seagull shampoo solution, and the austenite grains at different temperatures are observed by a metallographic microscope (Leica Microsystems Wetzlar, Germany). Finally, the effects of the original structure and subcritical quenching temperature on austenite grain size are analyzed. The sample cooled at 30 °C/min simulates the transformation amount of austenite during heating and the change in phase transformation point during cooling. The sample cooled at 30 °C/min simulates the transformation amount of austenite during heating and the change in phase transformation point during cooling. The change law of carbon content in austenite is calculated according to the lever law, and the effect of carbon content in austenite on the transformation products is analyzed. The average grain size of the original austenite is determined according to the average grain size of GB/T 6394-2002 metal. Five fields of view are randomly selected on the sample, and then three concentric equidistant circles are used as the measuring grid on the field of view photos. the field of view has 50 cut points, and the number of cut points of five field of view pictures is recorded to determine the error number. The relative error of counting, the average grain size, and the confidence interval are calculated. Photoshop software is used to determine the sum of the perimeter of the three circles and the average intercept number of the five field of view pictures and then calculate the average grain size. The results are required to meet the relative counting error RA% ≤ 10 to ensure the relative reliability of the results. ASTM grain size level calculation formula:

$$G=-6.643856 \mathrm{Lg} D - 3.288.$$

(2)

3. Effect of Original Structure and Intermediate Quenching Temperature on Austenite Grain Size

The microstructure of the sample in the AR state is GB + LB + less F + less P, and the microstructure of the Q state sample is GB + LB, as shown in Figure 3. Under the scanning electron microscope, the P in the AR sample is shown in Figure 2c, and its carbide distribution is a short rod or flake distribution on the ferrite matrix. Due to linear stack cooling at approximately 600 °C after AR, it shows a small amount of degenerated P. The carbides in degenerated P are precipitated on the F matrix in the form of large-size carbides with an irregular arrangement. According to the study, the definition of granular shellfish should be: GB is the mixture of bainitic ferrite matrix and “island” structures; the “island” structure may be two or more mixed structures of F, M, and M/A due to varying compositions and processes.

Austenite grain growth is a combination of thermal activation, diffusion, and interfacial reaction. Grain boundary migration is the main manifestation. The comprehensive effect of driving force and grain boundary movement resistance controls the grain growth trend. The higher the temperature of steel is, the greater the atomic diffusion coefficient is, and the carbon atom diffusion coefficient in austenite will also increase, which makes the austenitic nuclei grow faster. At the same time, the interphase free energy difference of the original microstructure increases and the driving force ∆Gv also increases, thus reducing the incubation period of austenitic nuclei and the time required for transformation. Under the action of a driving force, austenitic grain boundary migration occurs. When the holding time is the same, the growth rate of austenite grains can be expressed by the formula below. The higher the temperature of steel is, the greater the atomic diffusion coefficient is, and the carbon atom diffusion coefficient in austenite will also increase, which makes the austenitic nuclei grow faster. At the same time, the interphase free energy difference of the original microstructure increases and the driving force ∆Gv also increases, thus reducing the incubation period of austenitic nuclei and the time required for transformation. Under the action of a driving force, austenitic grain boundary migration occurs. When the holding time is the same, the growth rate of austenite grains can be expressed by the formula.

$$V=\mathrm{k}\exp \left(-\frac{\mathrm{Q}}{\mathrm{R}T}\right)\frac{\mathsf{\alpha }}{d},$$

(3)

where k is constant, Q (J/mol) is the activation energy of grain boundary migration, R (8.31 J/(mol·K)) is the gas constant, T (K) is absolute temperature, d (μ m) is the average diameter of austenite grains, and α (J/mol) is the interface energy.

According to Formula (3), the austenite grain growth rate has a proportional exponential relationship with quenching temperature. When the holding time is the same, the average grain diameter increases exponentially with the increase in heating temperature. Figure 4 shows the austenite grain morphology of quenched and rolled samples at sub-temperature intervals after holding the sample for 2.2 min/mm, which reflects the above law. According to the standard of GB/T 6394-2002 “Method for Determination of Average Grain Size of Metals”, the austenite grain size and grain size level of AR and Q state samples at each heating temperature are shown in

Table 2. Austenite grain sizes and grades of the steel at different heating temperatures.

State	Temperature/°C	D/μm	S	%RA	95%CI	G	95%CL	G ± 95%CL
AR state	770	15.15	5.89	3.57	7.31	8.79	0.10	8.79 ± 0.10
	810	15.50	4.71	2.92	5.85	8.72	0.08	8.72 ± 0.08
	830	16.58	2.86	1.90	3.56	8.53	0.06	8.53 ± 0.06
	850	17.38	3.21	2.23	3.98	8.40	0.06	8.40 ± 0.06
	870	18.32	11.61	8.51	14.41	8.24	0.25	8.24 ± 0.25
Q state	770	9.39	4.09	3.82	5.07	10.18	0.11	10.18 ± 0.11
	810	9.93	6.73	6.65	8.36	10.02	0.16	10.02 ± 0.16
	830	10.67	5.50	5.84	6.82	9.81	0.14	9.81 ± 0.14
	850	11.90	6.68	7.92	8.30	9.50	018	9.50 ± 0.18
	870	16.84	5.05	8.47	6.27	8.50	0.25	8.50 ± 0.25

Note: D is the average diameter of austenite grains, S is the cut-off count standard deviation, %RA is the relative error of measurement results, 95%CI is the 95% confidence interval, and G is the number of average grain size classes.

. Figure 4 and Figure 5 show the grain size comparison between rolled and hardened samples in the interval of intertemperature.

In summary, at the same temperature, the austenite grain size of the AR sample is larger than that of the Q state sample for the following reasons: The distribution of carbon elements in the R state is not as uniform as that in the Q state, the areas of high carbon content are easily steered toward austenitic nucleation, and the austenitizing state is early. However, during the heating process of GB, nucleation occurs at the GB grain boundary. With the increase in temperature, fine austenite is also formed in the grain, and there are more austenite nuclear particles. In addition, SXQ500/550DZ35 has low carbon content, and due to the influence of carbon atom diffusion and alloying element redistribution, the austenite advances to bainite relatively slowly, and the growth rate of austenite grains slows down accordingly.

When S = ΔD/ΔT for the characterization of grain size growth,

Table 3. Austenite grain sizes and grades of the steel at different heating temperatures of Q state.

Temperature range (°C)	770~810	810~830	830~850	850~870
The growth rate (μm/°C)	0.0135	0.037	0.0615	0.247

shows statistical results and Figure 5 shows the rate–temperature curve. It can be seen that the grain growth rate increases slightly with the increase in temperature. The main reason is: the higher the temperature, the easier to provide energy to the austenitic grain growth activation energy. In addition, there is still some ferrite in the sub-temperature range. As the temperature approaches the AC3 temperature, the lower the ferrite content, the lower the ability to hinder the growth of austenite grains, the higher the grain growth rate, and the weaker the grain refining effect.

As shown in Figure 6 and

Table 4. Statistics of ferrite ratios in different sub-temperature ranges of Q state.

Temperature/°C	770	810	830	850
Proportion of undissolved ferrite/%	14.13	10.4	8.67	0

, the quenching structures at 770, 810, 830, and 850 °C are M and F, and the proportion of F decreases with increasing temperature. At 850 °C, the microstructure is almost all M. Compared with 810 °C and 830 °C, the ferrite mass is relatively large and uneven at 770 °C.

4. Effect of Sub-Temperature Quenching Temperature on Austenite Conversion

The thermal expansion curves of holding for 2.2 min at different temperatures and cooling rates of 30 °C/min in the industrial production process are shown in Figure 7. It can be concluded that with the increase in temperature, the volume shrinkage of SXQ550/550DZ35 increases; that is, the austenite transformation increases. At temperatures below 740 °C, there is no volume shrinkage; that is, there is no α→γ transformation. At 847 °C, the volume shrinkage ends, and it is completely austenitized. In the cooling process, γ→α phase transition leads to volume expansion, and the phase transition point increases with the increase in quenching temperature.

According to the expansion curve in Figure 7, the shrinkage of the length from the beginning of austenite transformation to the end of heat preservation is determined, recorded as D770 °C, D810 °C, D830 °C, and D850 °C, respectively. Since the dimensional shrinkage of the sample is concluded when heated to 850 °C, the state of complete austenitization, the ratio of the size change at each temperature to D850 °C is the relative transformation of austenite and the relationship between the amount of austenite transformation and quenching heating temperature is obtained. The amount of austenite transformation of SXQ500/550Z35 at 770 °C, 810 °C, 830 °C, and 850 °C is 85%, 90%, 93%, and 100%, respectively.

5. Effect of Sub-Temperature Quenching Temperature on Carbon Content in Austenite

When SXQ500/550DZ35 is heated in the two-phase region, austenite nucleation and growth is a diffusion process. Since the diffusion coefficient of C is much larger than that of alloying elements, the content of C is the main factor affecting the formation of austenite. The carbon content of the austenite first formed near the carbide is close to the eutectoid point, and the eutectoid point of low alloy steel is lower than that of the binary Fe-C phase diagram (0.78%). The empirical formula summarized by Bain and Paxon is as follows [9]:

$$\mathrm{C}\left(\mathrm{wt}\%\right)=0.78 - 0.0623\mathrm{Mn} - 0.115\mathrm{Si} - 0.0313\mathrm{Ni} -1.0408\mathrm{Ti} - 0.0902\mathrm{Cr}+0{.0071\mathrm{Cr}}^2- 0.4456\mathrm{Mo}.$$

(4)

According to the formula, the eutectoid point carbon content of the steel is 0.429%.

According to the lever law in Figure 8, the carbon content in austenite at the temperature of the two-phase region is calculated, and the results are shown in

Table 5. Carbon content in austenite at different temperatures in the two-phase region.

Temperature (°C)	744	755	765	770	785	795	805	810	825	830	850
CA(Wt%)	0.429	0.395	0.364	0.348	0.302	0.271	0.240	0.225	0.178	0.163	0.110

.

$$C_A=C_O+\left(C_{eut}-C_O\right)\frac{A_{c3}-T}{A_{c3}-A_{c1}}.$$

When held at 770 °C, 810 °C, 830 °C, and 850 °C, the carbon content in austenite is 0.348%, 0.225%, 0.163%, and 0.11%, respectively. With the increase in temperature, the carbon content in austenite decreases gradually, and it is completely austenitized at 850 °C. The C element diffuses uniformly, and the carbon content reaches the equilibrium state of 0.110%. The lower the temperature in the two-phase region, the more uneven the components of austenite; carbon enrichment stabilizes the supercooled austenite. The CCT (Continuous cooling transition curve) shifts to the right so that the transition point decreases, resulting in different transition points and transformation products at the same cooling rate.

The CCT curve of SXQ500/550DZ35 with carbon content of 0.348%, 0.225%, 0.163%, and 0.110% under complete austenitization is simulated by JMatPro software. As shown in Figure 9, The simulated curve is basically consistent with the actual microstructure, which shows that the austenite carbon content has a great influence on the phase transformation products.

When the carbon content is 0.110%, the transformation temperature of bainite begins at 573.13 °C, and the critical cooling rate for producing F is 0.5 °C/s. When the carbon content is 0.163%, the transformation temperature of bainite begins at 563.38 °C, and the critical cooling rate for producing F is 0.2 °C/s. When the carbon content is 0.225%, the initial transformation point of bainite is 550.98 °C, and the critical cooling rate of producing F is 0.1 °C/s. When the carbon content is 0.348%, the initial transformation point of bainite is 527.35 °C, and F is not produced even if cooled at 0.01 °C/s. It can be seen that the temperature of the two-phase region affects the uniformity of various components in austenite and the enrichment degree of carbon and finally affects the level of the transformation point, resulting in different phase transformation points and transformation products at the same cooling rate.

6. Microstructure Change in SXQ500/550DZ35 Steel during the QLT Heat Treatment Process

SXQ500/550DZ35 is cooled after all austenitizing above Ac3 temperature, forming LB + Gb with high-density dislocation at the cooling rate (approximately 0.5 °C/s) of industrial production (Figure 10). After that, alloy elements such as Cr, Ni, Mo, and Cu slow down the decomposition and recrystallization of bainite in the heating process of subcritical quenching and retain the characteristics of lath. Under this microstructure condition, the austenite first nucleates on the lamellar boundary and grows along the lamellar boundary at sub-temperature. The nucleation and growth of the austenite is a diffusion process, and the diffusion coefficient of carbon is much larger than that of other alloying elements, so it is mainly affected by carbon content. During heating, the initially formed austenite is near the carbide, and the carbon content of the austenite is close to the eutectoid point. With the increase in temperature, the austenite content increases, but the carbon content decreases, while the alloying elements diffuse into the austenite, resulting in the composition fluctuation of the steel. In the process of secondary quenching, the higher the heating temperature is, the lower the carbon content in austenite is, and the worse the hardenability of undercooled austenite is. At the same time, with the carbon element diffusion of untransformed bainite during heating, the dislocation density decreases greatly, and F is formed. This kind of F retains the lath feature at the beginning and becomes polygonal with the increase in temperature, so the microstructure after secondary (sub-temperature) quenching is BF + “island” structure. The austenitizing temperature in the two-phase region is relatively low, which leads to the local fluctuation of carbon concentration, so the “island” structure is formed. Because of the different composition and technology, the “island” structure may be two or more mixed tissues of F, M, and M/A.

Figure 11 shows the scanning electron microscope organization chart of (Q) after quenching (L) at different sub-temperatures and cooling (L) at a rate of 30 °C/min. It can be seen that with the increase in temperature, the distribution of the “island” structure along the lath boundary gradually changes to that along the original austenite grain boundary.

Fang Hongsheng and others believe that according to the morphology and distribution characteristics of the island, granular bainite can be divided into two types. One is that the “island” is irregular and distributed irregularly on the massive ferrite matrix and does not produce needle-like relief during phase transformation, which is called “block granular bainite” [10]. The other is that the “small island” is discontinuous and distributed in parallel on the ferrite matrix with strip substructure, and there is typical needle-like relief during phase transformation, which is called “strip granular bainite”. The structures of these two morphologies can exist alone or at the same time. Both of them have the following common characteristics: (1) both of them are composed of ferrite + island second phase under a microscope and (2) both kinds of structures have experienced secondary phase transformation.

In the process of phase transformation, if the decomposition rate of the “carbon-rich” austenite region at the front of ferrite growth cannot keep up with the growth rate of ferrite, and the carbide is too late to precipitate, then when the adjacent ferrite lath merges, it will leave a “carbon-rich” austenite “island”. When the “island” is cooled to room temperature, it is partially decomposed, and the final structure is α + (M + γ), so-called granular bainite [11,12].

The observation results of the “island” structure in granular bainite by transmission electron microscope show that the structure in “island” can be martensite + retained austenite (M + residual γ), martensite + bainite + retained austenite (M + B + residual γ), or other metastable mixed microstructures, which depend on the composition of “island” undercooled retained austenite and its cooling rate during secondary transformation. Therefore, it is not comprehensive to call the “island” of granular bainite a “α/M or M/An island” [9].

When the test steel is heated in the two-phase zone, the percentage of austenite formed determines the amount of new martensite or bainite, while the carbon content in the new austenite determines the stability of austenite and the morphology of its evolution structure. Thus, it finally determines the mechanical properties of the steel. The finer the original structure is, the greater the dispersion of carbide is; the finer the austenite grain is, the microstructure is uniform and fine after repeated Q + nL + T, and the impact toughness is more stable.

7. Effect of Heat Treatment Process on Properties of QLT

In order to further study the effect of the sub-temperature quenching process on the microstructure and properties of experimental steel, and to guide the property control in practical production, the quenched sample of 920 °C, which is referred to as the Q state sample, is selected for heat treatment in different sub-temperature ranges of 770 °C, 810 °C, 830 °C, and 850 °C. The ultimate goal of the two-phase zone process is to further improve the low-temperature toughness of the steel by heating to the two-phase region, retaining part of the undissolved ferrite and making use of the ferrite toughness phase. The statistical results of mechanical properties after different heat treatments are shown in

Table 6. C Statistical table of mechanical properties of C quenched sample after sub-temperature treatment.

Q State Sample Number	Sub-Temperature Quenching Temperature/°C	Tempering Temperature/°C	Tensile Strength Rm/MPa	Yield Strength Re/MPa	Elongation After Fracture A/%	Impact Work/J (−20 °C)
1-1	770	610	704.1	365.3	21.0	107.2
1-2	770	640	695.4	516.1	23.7	97.28
1-3	770	670	684.8	447.5	24.2	55.1
2-1	810	610	754.0	619.5	21.4	183.2
2-2	810	640	717.2	587.1	21.5	197.4
2-3	810	670	688.5	519.6	23.2	177.3
3-1	830	610	758.3	623.0	19.8	206.4
3-2	830	640	715.0	592.1	21.2	250.3
3-3	830	670	673.9	542.0	21.4	202.2
4-1	850	610	772.4	649.4	20.0	160.7
4-2	850	640	743.2	631.5	19.4	211.1
4-3	850	670	703.3	561.6	24.2	161.4
5-1	-	610	747.3	639.6	19.5	117.6
5-2	-	640	726.2	625.9	22.2	158.6
5-3	-	670	687.9	590.4	21.7	196.1

.

It can be obtained from Table 6 that after sub-temperature treatment and tempering, the properties of almost all samples can meet the requirements. When the tempering temperature is constant, the tensile strength and yield strength increase with the increase in sub-temperature temperature, but the tensile strength is slightly lower than 810 °C at 830 °C and then continues to increase. However, the trend of elongation affected by sub-temperature temperature is not obvious. The impact work first increases with the increase in sub-temperature, reaching the maximum value of 250.3 J at 830 °C, and then begins to decrease and becomes unstable. When the sub-temperature quenching temperature is constant, the tensile strength and yield strength decrease with the increase in tempering temperature, the elongation at break increases with the increase in tempering temperature [13], while the impact work increases at first and then decreases, reaching the maximum value of 250.3 J at 640 °C. At the same time, compared with the samples without sub-temperature treatment, at the same tempering temperature, when the sub-temperature temperature is between 770 °C and 810 °C, the tensile strength and yield strength of the samples without sub-temperature treatment are relatively weaker than those without sub-temperature treatment. However, at 810 °C, the tensile strength of the sub-temperature treated sample is the same as that of the non-sub-temperature treated sample, the low-temperature impact energy begins to increase rapidly, and the impact energy increases from 158.6 J to 250.3 J when tempered at 640 °C; however, the strength is not greatly weakened. By comprehensive comparison, the best sub-temperature heat treatment process is 830 °C sub-temperature quenching + 640 °C tempering, which can greatly improve low-temperature toughness without too much loss of the original sample strength, and material properties with a good matching of strength and toughness can be obtained.

The two-phase zone heat treatment process is mainly used in dual-phase steel. The heat treatment process heats the steel plate to the austenite–ferrite (a + γ) two-phase zone and holds it for a period of time, then rapidly cools to obtain the required ferrite + martensite dual-phase structure. On the other hand, a dual-phase structure is a strengthening and toughening way to improve the toughness of steel by introducing a small number of ductile phases with good plasticity and toughness (austenite, lower bainite, and ferrite) based on martensite with good strength [14]. Because of its unique structural characteristics, dual-phase steel has high tensile strength, low yield strength, and high elongation compared with low alloy steel with similar strength, which is its application value [15]. For microalloyed high-strength steel, a composite structure containing a soft ferrite phase and hard phase can be obtained using a heat treatment process similar to the production of dual-phase steel to reduce the yield strength ratio and improve the impact toughness of the steel. The two-phase zone heat treatment process of microalloyed steel is to add (a + γ) two-phase zone quenching between normal quenching and tempering; that is, a quenching + lamellarizing + tempering process (QLT). The combination of soft ferrite and martensite can be obtained by the QLT process, and the ferrite–bainite structure can be obtained by controlling the cooling rate. Bainite replaces martensite as the strengthening phase in the structure [16]. Compared with traditional ferritic–martensitic steel, although the strength of Ferritic bainitic steel is slightly lower, it has a lower yield ratio and cold formability, better toughness, and weldability. The increase in soft ferrite content is beneficial to improve the toughness of ferritic–bainitic steel and the low-temperature impact toughness obtained by multiple sub-temperature quenching (Q + nL + T) processes is more stable.

The effect of austenitizing temperature on the microstructure and mechanical properties of bainite is as follows: with the increase in heating temperature and holding time, the austenite grains continue to grow. The existence of ferrite increases the C concentration of retained austenite, improves its stability, and makes it easier to obtain bainite. With the increase in austenitizing temperature, the grain size increases, the average chord length of the granular bainite island increases, the total amount of the island increases, the granular bainite increases, and the strength and toughness increase. It should be noted that under the condition of obtaining granular bainite, the higher the austenitization temperature, the better because too-coarse austenite grains will coarsen the bainitic ferrite lath, resulting in poor properties of the steel. The structural parameters of granular bainite, that is, the shape, quantity, size and distribution of “island”, are the decisive factors for the strength and toughness of granular bainitic steel. The main factors affecting the strength and toughness of granular bainite are the total amount and shape of islands (chord length and island spacing). With the increase in the total number of islands, the chord length and spacing of islands decrease, and the strength increases; with the decrease of the total number of islands, the chord length of islands decreases, the distance between islands increases, and the toughness increases. Within a certain range of composition, the number and size of islands are properly controlled, and a structure with a good combination of strength and toughness can be obtained.

Lower austenitizing temperature and shorter holding time can refine the austenite grain size, and fine grains can provide more nucleation positions, thus accelerating bainite transformation. The mixture of nanoscale bainitic ferrite and austenite lamellar can be obtained by bainitic transformation at low temperatures. Steel within this structure not only has high toughness but also has high strength and hardness. The effect of austenite deformation on the microstructure, strength, and toughness of grain boundary ferrite/granular bainite duplex steel is discussed below. The results show that austenite deformation not only refines the grain boundary ferrite but also promotes the nucleation of proeutectoid ferrite in the original austenite grain, which is beneficial to refine the granular bainite grain and its internal ferrite strip and MA island [17]. With the increase in deformation and the decrease in deformation temperature, the grain boundary ferrite is significantly refined. At the same time, a part of finer intragranular ferrite begins to appear in the original austenite grain. The bainite with high strength is obtained by water cooling, and the granular bainite is produced due to the limitation of the cooling rate of the extra thick plate. The granular bainite is formed at a higher temperature and has no obvious lath characteristics [18]. When the composition is constant, the cooling rate decreases, and the average chord length of the M/A island increases, but it has little effect on the total amount of the island. This is because if the cooling rate is slow, the Bs point is high, and the phase transformation rate is low, then the carbon atom has sufficient diffusion conditions, so it is obvious that the austenite is rich in carbon over a long distance. If the size of the island increases, the number decreases and the spacing increases.

8. Conclusions

(1)

The thermal expansion tests of AR and Q state samples were carried out using a DIL805L thermal dilatometer. The heating rate below 600 °C was 10 °C/s. The heating rate above 600 °C was 0.05 °C/s (simulating industrial production). The steel was heated to 770 °C, 810 °C, 830 °C, 850 °C, 870 °C, and 920 °C, respectively. The Ac1 = 744 °C and Ac3 = 847 °C of the steel were obtained, and the heat preservation coefficient of 2.2 min/mm was obtained according to the industrial production heat preservation coefficient. There are two ways of cooling: water quenching to room temperature, cooling at 30 °C/min, and grinding and polishing the metallographic samples. The comparison of austenite grain size at different temperatures was obtained under a LEICADMRX metallographic microscope to analyze the effect of the original structure and sub-temperature quenching temperature on austenite grain size. The sample cooled at 30 °C/min simulated the change in austenite transformation during the heating and cooling of SXQ500/550DZ35 in the two-phase zone, calculated the change law of carbon content in austenite according to lever law, and analyzed the influence of carbon content in austenite on phase transformation products. When the carbon content was 0.110%, the bainite transition temperature was 573.13 °C, and the critical cooling rate of F was 0.5 °C/s. When the carbon content was 0.163%, the bainite transition temperature was 563.38 °C, and the critical cooling rate of F was 0.2 °C/s. When the carbon content was 0.225%, the bainite transition point was 550.98 °C, and the critical cooling rate of F was 0.1 °C/s. When carbon content was 0.348%, bainite began to shift to 527.35 °C, and even 0.01 °C/s cooling did not produce F. The visible two-phase zone temperature affected the uniformity of the main components of the austenitic and carbon enrichment degree and influenced the high and low transformation points, resulting in the same cooling rate for different phase transformation points and transformation products.

(2)

In the original microstructure, the distribution of carbon elements in the rolled R state was not as uniform as that in the Q state. Under the condition of the same intertemperature quenching, the original austenite in the Q state had a smaller grain size, and the lower the intertemperature quenching, the more uneven the components in austenite were. Carbon enrichment made the supercooled austenite more stable, and the CCT curve shifted to the right, which reduced the transition point. Thus, different transformation points and transformation products were obtained at the same cooling rate, which effectively solved the problem of coarse and uneven bainite structure caused by the insufficient cooling rate of extra-thick plates. The microstructure after secondary quenching (subtemperature) was BF + “island” microstructure, and the “island” microstructure had a relatively low austenitization temperature in the two-phase zone, resulting in local carbon concentration fluctuations. Therefore, the formed “island” microstructure may be two or more mixed tissues of F, M, and M/A due to different components and processes. The finer the original microstructure was and the larger the carbide dispersion, the smaller the austenite grains obtained. Moreover, after repeating Q + nL + T, the microstructure was uniform and fine, and the impact toughness was more stable.

(3)

Finally, QLT (925 °C + 830 °C + 640 °C) steel had better strength, toughness, and lower yield ratio. The staggered microstructure of bainite and ferrite in QLT steel was similar to that of “fiber reinforced composites”. The bainite on the grain boundary in the structure played the role of strengthening the grain boundary and improving the properties. The interface between ferrite and bainite was highly coherent, there was no formation of brittle carbides on the interface, and it was not easy to produce local stress concentration. The soft ferrite phase not only reduced the yield–strength ratio but also prevented crack propagation and further improved the impact toughness. In conclusion, a large number of ferrite/bainite grain boundaries can effectively prevent crack propagation in steel.