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Arutyunyan, N. Low-Carbon Ti-Mo Microalloyed Hot Rolled Steels. Encyclopedia. Available online: https://encyclopedia.pub/entry/15497 (accessed on 18 June 2024).
Arutyunyan N. Low-Carbon Ti-Mo Microalloyed Hot Rolled Steels. Encyclopedia. Available at: https://encyclopedia.pub/entry/15497. Accessed June 18, 2024.
Arutyunyan, Nataliya. "Low-Carbon Ti-Mo Microalloyed Hot Rolled Steels" Encyclopedia, https://encyclopedia.pub/entry/15497 (accessed June 18, 2024).
Arutyunyan, N. (2021, October 28). Low-Carbon Ti-Mo Microalloyed Hot Rolled Steels. In Encyclopedia. https://encyclopedia.pub/entry/15497
Arutyunyan, Nataliya. "Low-Carbon Ti-Mo Microalloyed Hot Rolled Steels." Encyclopedia. Web. 28 October, 2021.
Low-Carbon Ti-Mo Microalloyed Hot Rolled Steels
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Low-carbon Ti-Mo microalloyed steels represent a new generation of high strength steels for automobile sheet. Excellent indicators of difficult-to-combine technological, strength, and other service properties are achieved due to the superposition of a dispersed ferrite matrix and a bulk system of nanoscale carbide precipitates.

Ti-Mo microalloyed steels low-carbon steels grain boundary strengthening precipitation strengthening interphase precipitates ferritic precipitates (Ti Mo)C complex carbide thermo-deformation treatment hot rolling parameters

1. Introduction

At present, one of the most priority areas of development of world science and technology is the problem of creating a new generation of various types of structural steels (first of all, for automobile sheets, pipes) with a fundamentally improved complex of difficult-to-combine indicators of strength, ductility, formability, corrosion resistance, operational reliability, and other service properties, while reducing production costs. This is due to the high rates of both production volume and the importance of such materials for the strategic, technical, economic, environmental, and social development of the world. One of the promising areas of research and development within the framework of the formulated problem is the creation of new high-strength low-carbon microalloyed steels with a homogeneous dispersed ferrite matrix. Due to the low carbon concentration, ferritic steels have good weldability, corrosion resistance, and other service properties. High strength is achieved due to the creation of a volumetric system of various types of phase precipitates and dispersed microstructure that leads to grain boundary and precipitation strengthening. The successful use of the contribution of nanoscale precipitates to the strengthening Ti-Mo microalloyed steels has allowed the development of steels with a tensile strength of 780 MPa [1], which are produced by JFE Steel.

These steels have received special attention in a number of recent detailed reviews on advanced high-strength steels [2], microalloyed steels [3], and Ti-microalloyed steels [4][5].

The improvement of Ti-Mo microalloyed steels is currently underway. At the same time, intensive researches are aimed to finding the possibility of purposeful formation of the most optimal structural state, including determining the role and patterns of the formation of phase precipitates at various stages of thermo-deformation processing of steel. Precipitates formed during hot deformation in austenite (so-called “austenitic” precipitates) are influencing on the dispersion of the microstructure. Precipitates formed as a result of random nucleation and growth in a ferrite matrix (so-called “ferritic” precipitates) contribute to precipitation strengthening. During the γ → α phase transformation of steel, the so-called interphase precipitates are formed. They are arranged in layers parallel to the moving front of the γ / α transformation [6]. It is essential that interphase precipitates increase the strength of the steel not only by the precipitation strengthening mechanism, but also contribute to the formation of a fine-grained structure by inhibiting grain growth.

In this entry, the main attention is paid to the studies of the effect of precipitation of excessive phases of various types and dispersion on the microstructure, the implementation of various strengthening mechanisms and, accordingly, mechanical properties depending on the chemical composition and parameters of hot deformation of low-carbon Ti-Mo microalloyed steels.

2. Implementation of the Main Strengthening Mechanisms

2.1. Grain-Boundary Strengthening

Grain-boundary strengthening is one of the main factors that determine high strength characteristics of microalloyed steels. The formation of fine grains of ferrite in a number of microalloyed steels is the result of the formation of ferrite from unrecrystallized austenite. Microalloying with niobium most effectively slows down the recrystallization of austenite during hot deformation, but the effect of titanium is also significant [7]. The main reason that retards recrystallization of austenite during hot deformation is its inhibition by the “austenitic” phase precipitates formed upon initiation of deformation [8].

During the process of γ → α-transformation, the final structure of the steel is formed. The influence of microalloying elements on this process, which is controlled by the diffusion of carbon in austenite, depends on their state: in a solid solution or in a carbonitride phase. Phase precipitates formed in the austenite region are the sites of ferrite nucleation and, thus, contribute to the refinement of the microstructure.

In studies of Ti-Mo microalloyed steels, “austenitic” carbonitride precipitates were not detected [9] or were observed in very small amounts [10]. This led to the intensive formation of interphase precipitates. The addition of molybdenum is favorable for the formation of the dispersed ferrite (acicular type) in Ti-Mo microalloyed steels [11][12].

2.2. Peculiarities of the Acicular Ferrite Formation

The microstructure of acicular ferrite (AF) is characterized by a chaotic arrangement of ultrathin (2–3 μm) [4] needle-like ferrite grains (Figure 1). Comparison of the mechanical properties of metal with the microstructure of various types of ferrite, summarized in [13], indicate an increased strength and good impact toughness of the AF [14][15][16][17][18].

Figure 1. Microstructure of acicular ferrite in steel containing (wt. %): 0.08C-0.21Mo-0.165Ti [9].

Detailed information and analysis of AF formation is presented in [4][13][19]. A necessary condition for AF formation is the presence of a certain amount of non-metallic inclusions in the metal. Several mechanisms have been proposed to explain the generation of AF on non-metallic inclusions:

– a reduction in the interfacial energy for simple heterogeneous nucleation on a surface of inclusions [20][21][22],

– an epitaxial nucleation on the inclusions, which have a good coherency with ferrite [23],

– a nucleation arising from the thermal strains at the inclusions, which is associated with the different thermal expansion coefficients of the inclusions and steel matrix [24],

– a nucleation arising from solute depletion of elements in the matrix near inclusions [15][16][25][26][27][28].

The implementation of each mechanism depends on the chemical composition of the steel, cooling rate, grain size of austenite, type, size, morphology and amount of non-metallic inclusions [13]. For Ti-Mo microalloyed steels, a positive factor is that various titanium-containing non-metallic inclusions play an essential role for the formation of AF [13][15][18][23][26][29][31][32][33]. The heterogeneity and / or complexity of some inclusions (for example, Ti2O3-based inclusions with the MnS and TiN formed on them) make it possible to achieve a favorable correspondence between the lattices of the substrate and the incipient phase [13]. In [33] it was shown that for steels of one grade a certain type of inclusions could be favorable, for another grade – these inclusions are inert.

The question of the critical size of inclusions and their volume fraction still remains open. Summarizing the data of various studies, in [19], the critical size of the inclusion is defined as 0.3 - 1 µm. According to [13], inclusions larger than 1 µm are more likely to nucleate ferrite. Although the amount of inclusions in steel is considered the main factor in the nucleation of acicular ferrite, only about 10–36% of the inclusions take part in the nucleation process [13]. Therefore, the vast majority of inclusions do not play any role in the nucleation of ferrite. It was concluded in [33] that the number of inclusions should be large enough to provide sufficient nucleation centers, but as low as possible to meet the requirements for high quality steels. In this case, the type of inclusions for the nucleation of acicular ferrite is more important than the number of effective particles.

2.3. Precipitation Strengthening

The prospects for the development of steels with a high complex of properties due to the implementation of various mechanisms of precipitation strengthening are analyzed and substantiated in the works [34][35][36][37]. As shown by the results of studies of ferritic steels containing titanium and molybdenum, including [1][38][39][40][41][42], the amount of precipitation strengthening is significant and increases with a decrease in the size of particles (up to 1 - 3 nm) and an increase in their number.

For example, the contribution to the yield strength of Ti – Mo steels, as shown by estimates, can reach 400 MPa and more [43][44]. At the same time, high strength corresponding to the strength class of 1180 MPa [45] is achieved due to an increase in precipitation strengthening while maintaining practically unchanged contributions corresponding to the basic level of strength of ferritic steels: solid solution hardening and hardening due to refining of ferrite grains.

It should be noted the advantage of precipitation strengthening during microalloying with titanium in comparison with niobium. According to estimates [7], the contribution of nanoscale NbC precipitates to the hardening of industrial Nb-containing steels is only 80-100 MPa. This conclusion was confirmed by the results of [46][47][48]. This is most likely due to kinetic constraints. According to the kinetic study results of [49], NbC formation, even during deformation, is kinetically inhibited.

Analysis of nanoscale carbide precipitates in low-carbon microalloyed steels shows that, as a rule, they are of the “ferritic” and / or interphase type [38][39][46][47][48][50][51][52][53].

The formation of interphase precipitates occurs during the γ → α transformation at the α / γ interface [54], they have a size of 3-10 nm and make up layers with a period of 10-20 nm (Figure 2). If, after cooling to the temperature of coiling the hot-rolled strip into a coil, the ferrite matrix remains supersaturated with microalloying elements, then during the subsequent relatively slow cooling, the precipitation of nanoscale carbides occurs. The arrangement of these particles is not ordered, as in the case of interphase precipitates [38][55]. The nucleation and growth of disordered nanoscale carbides in ferrite occurs on dislocations, their average size is 3-10 nm [38].

Figure 2. Interphase precipitates in steel containing (wt. %): 0.08C-0.21Mo-0.165Ti [9].

According to estimates based on the results of electron microscopy studies [1][38][43][45], the high contribution of precipitation strengthening can only be due to interphase precipitates. At the same time, the presence of random particles precipitated in the ferrite matrix also contributes to the strengthening [50].

3. Influence of Chemical Composition

3.1. Strengthening Effect due to (Ti, Mo)C Carbide Precipitation

The ultimate strength of steels with various basic microalloying components can be increased up to 800-1000 MPa due to molybdenum addition [1][35][36][38][39][40][56][57]. In this case, the strengthening effect, mainly achieved due to the precipitation of nanoscale molybdenum-containing carbides, is estimated at 300-400 MPa.

The chemical composition of (Ti, Mo)C carbides depends on the duration of their precipitation and, accordingly, the size [50][58]. The longer the holding time, the less molybdenum present in the carbide [50].  

A decrease in the size of carbides with molybdenum addition to steel microalloyed with titanium [40][49] and Ti-Nb [38][41][60] has been established. On the one hand, the substitution of molybdenum for titanium and niobium in carbide is constrained by the increase in Gibbs free energy. On the other hand, molybdenum incorporation into the carbide reduces the degree of lattice mismatch between the carbide and the matrix due to a decrease in the lattice period, and, consequently, the energy of the carbide / austenite interface [61]. The second effect is more pronounced during the initial stage of precipitation when carbide particles are relatively small and, therefore, their precipitation is accelerated by Mo addition. Smaller particles (Ti, Mo)C contain a higher fraction of Mo, as was found in [50].

In addition, it was found that Mo atoms tend to be in the outer region of complex carbide precipitates at an early stage of precipitation, which leads to a stronger de-crease in the energy of the interface [62]. Since the surface-to-volume ratio decreases with increasing particle size, the volume fraction of Mo enrichment decreases, which leads to a decrease in the Mo fraction in carbide precipitates with increasing time [50][58].

The possibility of the formation of a system of nanoscale precipitates, coherent with the matrix, is due to their high nucleation rate, but low growth rate, which gives them thermal stability [41][44][50][57][58][63]. According to [57][64], the thermal stability of Mo-containing carbides is associated with diffusion processes. For the growth of Mo-containing precipitates, diffusion of Mo and other microalloying elements from the ferrite matrix to the ferrite / carbide interface is required. Since the diffusion rate of Mo is lower than that of other microalloying elements, the growth rate of precipitates is controlled by the Mo diffusion and therefore slows down [64]. In [65] it is said that the Oswald ripening of Ti-containing carbides is controlled by titanium diffusion and is suppressed by reducing the solution of titanium.

3.2. Influence of Titanium Content

Table 1 summarizes information on the mechanical properties of various studied Ti-Mo microalloyed steels and the corresponding parameters of hot rolling: finishing temperature of rolling (Tf) and coiling temperature (Tc), as well as the cooling rate from Tf to Tc (Vc). For the compositions studied in [9][10][39][65][66][67], the values of Tf and Tc corresponding to the highest tensile and yield strengths were selected. It can be seen that the titanium content is from 0.09 to 0.233 wt. %, and molybdenum is usually 0.19-0.22 wt. %, in some cases up to 0.43 wt.%.

The effect of titanium content on the mechanical properties of hot rolled low-carbon molybdenum alloyed steels was investigated in [1]. It was shown that with an increase in the titanium content to 0.09 wt. %, the tensile and yield strengths of steel containing 0.2 wt. % Mo increases and exceeds 800 MPa and 700 MPa, respectively.

Table 1. Chemical composition, hot rolling parameters, and mechanical properties of low-carbon Ti-Mo microalloyed steels.

No.

Chemical composition

Hot rolling parameters

Mechanical properties

Ref.

C

Si

Mn

Mo

Ti

Cr

Ni

Al

P

S

N

Тf, °С

Тc, °С

Vс, °С/с

σ0.2, MPa

σB, MPa

δ, %

1

0.043

0.18

1.62

0.19

0.092

-

-

-

0.008

0.001

0.0032

900

620

10

~780

~800

~23

[1]

2

0.045

0.19

1.53

0.20

0.119

-

-

-

0.009

0.001

0.0031

900

620

10

~780

~820

~18

3

0.046

 

1.30

0.19

0.099

 

 

0.042

 

0.001

0.0028

950

600

10

683

786

 

[65]

4

0.07

0.32

1.34

0.2

0.09

 

 

0.04

 

 

 

900

650

500

10

642

770

16.6

[42]

5

0.079

0.17

1.22

0.43

0.19

 

 

0.03

0.0038

0.0052

0.0048

870

600

~30

912

971

16

[39]

6

0.048

0.20

1.61

0.20

0.09

 

 

0.02

0.020

0.006

0.0040

900

650

15

747

807

27

[68]

7

0.059

0.30

1.490

0.193

0.233

 

 

 

 

 

 

900

500

10

858

941

 

[41]

8

0.075

0.20

1.75

0.275

0.17

0.16

0.175

0.035

 

 

0.005

880

620

20

860

951

23.5

[66]

9

0.070

0.030

1.64

0.22

0.090

0.050

0.04

0.030

0.008

0.006

0.009

870

600

11

649

771

19.5

[10]

900

620

28

717

767

15

10

0.063

0.063

1.61

0.21

0.096

0.021

0.01

0.057

0.003

0.003

0.003

850

650

8

694

777

19.6

910

620

23

725

780

16.5

11

0.052

0.05

1.51

0.17

0.12

 

 

0.053

0.003

0.003

0.003

890

650

~20

814

897

19

12

0.08

0.15

1.24

0.21

0.092

0.03

0.01

0.11

0.003

0.005

0.010

900

650

10-15

590

670

17

[67]

30

720

810

 

13

0.08

0.16

1.39

0.21

0.165

0.03

0.01

0.13

0.003

0.005

0.009

900

650

10-15

770

840

15

30

890

950

 

In addition, the ratio of titanium and molybdenum concentrations is important [1][69], and the maximum values of strength characteristics are achieved when atomic concentrations are equal. In [67] it was also found that steel No. 13 (Table 1) with a higher titanium content has higher yield and tensile strengths than steel No. 12 (Table 1).

4. Influence of Thermo-Deformation Parameters

4.1. Influence of Temperatures of the Rolling End and Coiling

The cooling process after hot rolling and during coiling of the strip largely controls the main changes in the microstructure of the metal, γ → α transformation, and the phase precipitation. The influence of Tf, Tc, and the cooling rate of rolled products from Tf to Tc on the structural state and mechanical properties has been studied in [9][10][39][41][65][67].

Despite the established regular parabolic form of the dependence of the hardness and yield strength of microalloyed steels on the coiling temperature in the case of titanium [70][71], the dependences for Ti-Mo microalloyed steels are different (Figure 3).

Figure 3. Influence of the coiling temperature on the yield strength of Ti-Mo microalloyed steels (Table 1): 1 – No. 3 [65], 2 – No. 5 [39], 3 – No. 6 [68], 4 – No. 7 [41], 5 – No. 11 [10], 6 –No.12 [67], 7 – No. 13 [67].

In [10] was showed that when high temperatures of the rolling end and coiling (Tf = 900 ⁰C, Tc = 620 ⁰C) were used for steel No.9 (Table 1), most of precipitates presented were interphase, while the obtained samples had a finer-grained structure and the highest values of tensile and yield strengths. With a decrease in the Tf and Tc, the relative amount of interphase precipitates decreased, and when using the regime with Tf = 860 °C and Tc = 560 °C, all observed nanoscale carbide precipitates were “ferritic”. It was explained by a joint analysis of the formation of excessive phase precipitates and the occurrence of γ → α phase transformation in steel [72][73].

4.2. Influence of the Cooling Rate from Tf to Tc

In the works [9][67], the effect of the cooling rate from Tf to Tc on the formation of the structural state and mechanical properties of steels No. 12, 13 (Table 1) was investigated. The results showed that, upon slow cooling (10-15 ⁰C/c), a polygonal ferrite microstructure with a high amount of nanoscale interphase precipitates is formed. With an increase in the cooling rate (30 ⁰C/c), its microstructure is acicular ferrite with higher strength. However, the formation of nanoscale carbide precipitates in this case occurs in ferrite in the form of randomly spaced individual or small groups of precipitates, instead of the implementation of the mechanism of formation of regularly spaced interphase precipitates. This is due to the fact that the rate of the γ → α phase transformation of steel turns out to be too high and the formation of precipitates on the moving phase boundary does not have time to occur. Thus, with a change in the cooling rate, opposite tendencies are observed in the implementation of main strengthening mechanisms.

5. Impact Toughness

Recently, the range of technical requirements for high-strength steels has been expanding, and in many cases it includes an additional high level of cold brittleness and low-temperature toughness.

Generally, to avoid significant reduction in toughness, the main approach is to prevent the formation of coarse microstructure and large non-metallic inclusions, usually TiN [74][75][76][77]. To increase the impact toughness of low-carbon microalloyed steels, the most effective solution is to increase the dispersion of the ferritic microstructure [68]. Results of [78][79] showed a positive effect of low finishing rolling temperature on the impact toughness of microalloyed steels. At the same time, the use of low Tf is not always effective for achieving high strength of titanium microalloyed steels, since it prevents the formation of interphase precipitates [9].

Thus, in order to obtain simultaneously high strength and impact toughness of Ti-Mo microalloyed steels, further search for the optimal combination of the content of microalloying elements and the parameters of thermo-deformation processing is necessary for a balanced effect of microstructure characteristics and phase precipitates on the properties of steel.

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