Effect of Ti and TiNb Microalloying on Microstructures and Mechanical Properties of 2200 MPa Low-Alloy Ultra-High-Strength Steels

Abstract

The 2200 MPa low-alloy ultra-high-strength steels microalloyed with Ti (Ti-steel) and Ti and Nb (TiNb-steel) were quenched at 860 °C and tempered at 180 °C, and the mechanical properties, microstructure, and precipitated phase were studied by SEM, TEM, and physicochemical phase analysis. The results of mechanical properties showed that the TiNb-steel had higher strength than the Ti-steel, with a yield strength of 1746 MPa and 1802 MPa and a tensile strength of 2198 MPa and 2232 MPa, respectively. The TiNb-steel had a finer structure than the Ti-steel, with the effective grain sizes being 0.86 μm and 1 μm, respectively. The TiNb-steel had more MC-type carbides than the Ti-steel, and the MC-type carbide contents were 0.24 wt.% and 0.19 wt.%, respectively. The reason for the higher yield strength of the TiNb-steel is that it has higher strengthening effects due to finer grains, a higher density of dislocation, and more precipitation. The reason for the higher tensile strength of the TiNb-steel is that it has a higher coefficient of variable hardening K and a higher strain hardening index n.

1. Introduction

Ultra-high-strength steels are widely used in the automobile and engineering machinery industries. As the demand for components with higher load-bearing capacity and wear-resistant performance, especially the increase in overall weight reduction, the development of materials with higher strength and toughness becomes more and more important [1,2,3]. To meet this requirement, AISI 4340 [4], 300M [5], AerMet 100 [6], and other ultra-high-strength steels with strengths of about 2000 MPa are widely employed. However, these materials are composed of a relatively large amount of precious metal elements such as Ni and Co, leading to high prices. Materials with better mechanical properties and lower cost are of great interest for industry application; thus, the advancement of ultra-high-strength steel has emerged as a hot topic of research [7,8,9].

Microalloying technology has been proved to be an effective method to improve the mechanical properties of steels [10,11]. Wu et al. [12] investigated the effect of Nb content on the corresponding microstructure and mechanical properties performance of 1800 MPa ultra-high-strength steel and found that 0.021 wt.% Nb addition retarded the formation of harmful iron carbides. Zhang et al. [13] studied the effect of Nb content on the mechanical properties and microstructure of a high-strength low-alloy steel and found significant improvement in strength and low-temperature toughness after the addition of 0.04% Nb, which could be ascribed to microstructure refinement and the precipitation of nanosized carbides. Qi et al. [14] investigated the impact of Nb content changes on the structure and mechanical characteristics of NiCrMo-welded joints and the results indicated that the increase in the quantity of Nb additions could result in the austenite grain deformation and the expansion of the precipitate area in the experimental steel, leading to the increase of tensile strength. Cao et al. [15] studied the effect of Nb addition on the high manganese steels and revealed that the microstructure of the steels was refined by the addition of Nb; moreover, the steels with Nb addition had higher stacking fault energy, which provided higher strength and higher ductility for the steels. Despite intensive research work, the effect of Nb addition on low-alloy ultra-high-strength steels has not been revealed clearly. Therefore, this study focuses on the effect of Nb element in addition to Ti microalloying on the microstructure and mechanical properties of 2200 MPa grade low-alloy ultra-high-strength steels. The microstructure and mechanical properties were investigated, and the strengthening mechanism was discussed. This research aims to provide references for the development and application of low-alloy ultra-high-strength steels.

2. Materials and Methods

2.1. Experimental Materials

The chemical compositions of the experimental steels are presented in

Table 1. Chemical compositions of the experimental steels (mass fraction, %).

Steel	C	Si	Mn	P	S	Mo	Ti	Nb
Ti-steel	0.48	0.40	1.21	0.011	0.003	0.30	0.065	-
TiNb-steel	0.47	0.41	1.19	0.010	0.003	0.31	0.068	0.035

, which were denoted as Ti-steel and TiNb-steel, respectively. A 50 kg vacuum induction furnace was employed for the preparation of steel ingots, then the ingots were forged into 80 mm thick billets, followed by heating at 1200 °C for 2 h, afterwards rolled into 6 mm thick plates, and air-cooled to room temperature. Heat treatment of the rolled plates was performed by heating at 860 °C for 30 min and then water-cooled to room temperature, followed by tempering at 180 °C for 2 h to obtain the final state for the experiments.

2.2. Experimental Methods

Plate tensile specimens and Charpy V-notch impact specimens were machined from the plates in the transverse orientation. Flat tensile specimens were cut directly from the steel billets without any surface treatment; the dimensions are illustrated in Figure 1. The Charpy impact specimens are in dimensions of 10 mm × 5 mm × 55 mm with a 2 mm V-notch. Tensile tests were conducted at room temperature at a strain rate of 10⁻² s⁻¹ using an AMSLER-50 testing machine (Rancho Palos Verdes, CA, USA) according to ISO 6892-1:2019 [16], and Charpy impact tests were performed at −40 °C on a standard impact machine according to ISO 148-1:2016 [17].

The microstructure of the experimental steels was observed using specimens with dimensions of 10 mm × 15 mm × 6 mm. The cross-section surfaces were ground and polished, followed by etching with a 3% nitric acid-alcohol solution. The microstructures of the specimens were observed using a Quanta 650 scanning microscope (SEM) (Hillsboro, OR, USA) operated at 15 kV and a FEI Tecnai G2 F20 field emission transmission electron microscope (TEM) (Hillsboro, OR, USA) operated at 200 kV using thin foil specimens prepared using a double jet polishing machine (Kejing, China) in a solution containing 5% perchloric acid and 95% methanol at −20 °C. The effective grain sizes were analyzed using electron backscatter diffraction (EBSD) (Ametek, PA, USA), and the compositions of the carbides were studied by SEM-EDS spectra. The dislocation densities of the experimental steels were measured using a Bruker D8 ADVANCA X-ray diffractometer (Karlsruhe, Germany) after the samples were electropolished with 10% oxalic acid after grinding and mechanical polishing to remove the influence of surface stress. The dislocation density was calculated using the corrected Williamson–Hall Equation [18].

The carbides in the experimental steels were analyzed by the physicochemical phase analysis method. Specimens with dimensions of 20 mm × 80 mm × 5 mm were used, and precipitated carbides were extracted by electrolysis in a methanol solution containing 5% HCl, 5% glycerol, and 1% citric acid. Thereafter, the precipitated carbides were analyzed by inductively coupled plasma-mass spectrometry (ICP-AES).

3. Results

3.1. Microstructure

Figure 2 shows the microstructure images of the experimental steels obtained by SEM and TEM. Both steels show lath martensite microstructures with fine grain sizes.

The EBSD microstructure results of the experimental steels are shown in Figure 3a,b. It was determined that both steels were divided by martensitic sub-structures (high- and low-angle grain boundaries) into a number of small and different regions [19,20]. The martensite block with high-angle grain boundaries was proven to be the smallest unit. Therefore, the martensite block sizes can be regarded as the effective grain sizes (EGS) of martensite [21]. AZtec Crystal (3.1.) software was used to calculate the block sizes of the martensite substructure in the grain boundary images obtained by EBSD, where the sizes are determined by calculation from the statistical measurement of grain areas. The measured grain size distributions of the experimental steels are shown in Figure 3c. It is obvious that TiNb-steel has many more small-sized grains than Ti-steel, especially for grains smaller than 1.5 μm. Moreover, the statistical data of grains larger than 5 μm also show that TiNb-steel has a decreased number compared to Ti-steel. The EGS of the martensite substructure was determined to be 1.00 μm for the Ti-steel and 0.86 μm for the TiNb-steel. This suggests that the addition of Nb can further refine the grains of the TiNb-steel.

The TEM images also show that the microstructure of both steels exhibited a high density of dislocations, as shown in Figure 4. The dislocation densities of the Ti-steel and TiNb-steel were further examined by XRD analysis, and the spectra are presented in Figure 5. The dislocation density can be calculated using the modified Williamson-Hall equation [18,22]. The calculated value for Ti-steel is 1.53 ± 0.44 × 10¹² cm⁻², and that for TiNb-steel is 1.68 ± 0.37 × 10¹² cm⁻². The higher value of the TiNb-steel is in accordance with the phenomenon observed in the TEM. It can be concluded that the addition of Nb to the TiNb-steel results in an augmentation of its dislocation density and enhances its strength.

3.2. Mechanical Properties

The mechanical properties of the Ti-steel and the TiNb-steel are illustrated in

Table 2. Mechanical properties of the experimental steels.

Steel	Tensile Strength, Rm/MPa	Yield Strength, Rp0.2/MPa	Elongation, A/%	Reduction of Area, Z/%	Impact Toughness, KV2 (−40 °C)/(J/cm2)
Ti-steel	2198 ± 4	1746 ± 3	10.5 ± 0.5	40.5 ± 0.5	20.0 ± 1.0
TiNb-steel	2232 ± 3	1802 ± 4	10.0 ± 0.5	41.5 ± 0.5	18.8 ± 0.5

. The TiNb-steel shows an enhanced strength compared to the Ti-steel with an increase of 34 MPa in tensile strength and 56 MPa in yield strength. The ductility results show slight differences between the experimental steels. Notably, both steels exhibit values of elongation and reduction of area higher than 10% and 40%, respectively, indicating good ductility in the strength level of 2200 MPa. The values of impact toughness are 20.0 and 18.9 J/cm² for the experimental steels, respectively, which are comparable to each other. In this sense, the addition of Nb to the steel benefits a better strength while slightly decreasing the toughness.

3.3. Precipitation Analysis

Precipitation strengthening is an important mechanism for martensite steels. The existence of ε-carbide precipitates can be observed through transmission electron microscopy [23]. These ε-carbides are distributed in large quantities in the interior of the lath martensite, as illustrated in Figure 6. A physicochemical phase analysis was used to qualitatively and quantitatively examine the ε-carbides of the Ti-steel and TiNb-steel, as illustrated in

Table 3. ε-carbide determination of the experimental steels.

Steel	The Mass Fraction of Each Element in ε-Carbides, wt.%
	Fe	C	Σ
Ti-steel	0.609	0.044	0.653
TiNb-steel	0.640	0.046	0.686

. The results showed that the ε-carbides of both steels were composed of Fe and C. Additionally, the amount of ε-carbide precipitates in the experimental steels shows little difference, indicating that the addition of Nb would not promote the precipitation of ε-carbides.

The precipitates of the experimental steels were subjected to further observation by employing a high-magnification scanning electron microscope. The typical morphologies of the precipitates are shown in Figure 7a,b. The figures revealed the presence of fine and diffusely distributed spherical precipitates in both steels. The EDS spectroscopy was employed to identify the main compositions of the precipitates, and the typical results are illustrated in Figure 7a1,b1, which demonstrate that the precipitates of the Ti-steel are primarily composed of Fe, Mn, Ti, Mo, and C, while those of the TiNb-steel are mainly composed of Fe, Mn, Ti, Nb, Mo, and C. The qualitative and quantitative detection of the MC carbides of the experimental steels was determined by physicochemical phase analysis, as illustrated in

Table 4. MC carbides of the experimental steels determined by physicochemical phase analysis.

Steel	The Mass Fraction of Each Element in MC Carbides, wt.%
	Ti	Nb	Mo	$\mathbf{C}$	Σ
Ti-steel	0.060		0.102	0.028	0.190
TiNb-steel	0.066	0.033	0.104	0.034	0.237

. The results showed that the MC carbides in the Ti-steel were composed of Ti, Mo, and C, while those in the TiNb-steel consisted of Ti, Nb, Mo, and C, which is in accordance with the results of the EDS spectroscopic detection. In addition, both Ti and Nb elements existed predominantly in the form of precipitation regarding the total element amount of the experimental steels. The precipitation of MC carbides in the TiNb-steel is more than that of the Ti-steel, which can be ascribed to the precipitation of NbC. The size distributions are illustrated in Figure 8. The MC carbides in the experimental steels show similar distribution. It is noteworthy that more than half of the MC carbides are smaller than 10 nm, and the amount of carbides smaller than 5 nm is 20% and 22% for Ti-steel and TiNb-steel, respectively. Moreover, it seems that the addition of Nb would promote the formation of large particles, as for TiNb-steel, more carbides have sizes larger than 200 nm, which might be related to element segregation and co-precipitation of Nb, Ti, and Mo elements.

4. Discussion

The experimental steels are typical martensitic steels. To reveal the strengthening mechanism of Ti and Nb elements, the yield strength is determined by considering the four primary strengthening components: fine grain strengthening, precipitation strengthening, dislocation strengthening, and solid solution strengthening. However, the strengthening components are somewhat interrelated rather than independent of each other. When the strength of the steel is relatively low, there is a reduction in the interaction between the strengthening components, and they can be linearly superimposed. When the strength of the steel is high enough, the interaction effect of the strengthening components cannot be ignored. In this case, a root-mean-square superposition is utilized to perform the requisite calculations, as illustrated in Equation (1) [24].

$${\sigma }_y={\sigma }_0+{\sigma }_{ss}+\sqrt{{\sigma }_d^2+{\sigma }_g^2+{\sigma }_p^2}$$

(1)

where ${\sigma }_0$ is Peierls–Nabarro force, which is taken as 39 MPa in this paper; ${\sigma }_{ss}$ is the solid solution strengthening component; ${\sigma }_d$ is the dislocation strengthening component; ${\sigma }_g$ is the fine grain strengthening component; ${\sigma }_p$ is the second phase precipitation strengthening component; and the strengthening unit of lath martensite is the lath block.

The solid solution strengthening component can be calculated by the following Equation (2) [25]:

$${\sigma }_{ss}={\mathrm{k}}_{\mathrm{C}}{\left[\mathrm{C}\right]}^{{\displaystyle \frac{1}{2}}}+{\mathrm{k}}_{\mathrm{M}}\left[\mathrm{M}\right]$$

(2)

where [C] represents the mass percentage of strong solid solution strengthening element C dissolved in the matrix, while k_c is a constant with a value of 1722.5 [26]. The carbon content of the TiNb-steel is greater than 0.20 wt.%. When the carbon content is greater than 0.20 wt.%, about 0.20 wt.% carbon element would segregate in the dislocations. Consequently, 0.20 wt.% of carbon should be deducted from the calculation of carbon solid solution strengthening. Moreover, the carbon element in carbides should be subtracted. [M] denotes the mass percentage of weakly solid solution-strengthened elements; k_M denotes the scale factor, where the values of Si, Mo, and Mn elements are 83, 11, and 37, respectively [27]. Thus, the solid solution-strengthening fractions were determined to be 866 MPa and 832 MPa for the Ti-steel and the TiNb-steel, respectively.

The dislocation strengthening component can be calculated by the following Equation (3) [28]:

$${\sigma }_d=\mathsf{\alpha }\mathrm{G}\mathrm{b}\sqrt{\rho }$$

(3)

where α is a constant related to the crystal structure with a value of 0.3; G is the shear modulus with a value of 81 GPa; $\mathrm{b}$ is the Burgers vector with a value of 0.248 nm [28]; and $\rho$ is the dislocation density determined by XRD spectra. Thus, the dislocation strengthening components of the Ti-steel and the TiNb-steel were calculated to be 751 MPa and 788 MPa, respectively.

The fine grain strengthening fraction can be calculated by the Hall–Petch Equation (4) [29,30]:

$${\sigma }_g={\mathrm{k}}_{\mathrm{y}}{\mathrm{d}}^{-{\displaystyle \frac{1}{2}}}$$

(4)

where ${\mathrm{k}}_{\mathrm{y}}$ is a scaling factor of 6.47 MPa·mm^1/2; d is the size of the martensitic block. Thus, the fine grain strengthening fractions for the Ti-steel and the TiNb-steel were calculated to be 204 MPa and 220 MPa, respectively.

The precipitation strengthening of steels depends on the size, composition, and quantity density of the second phase particles that are diffusely precipitated from the supersaturated solid solution, which are mainly utilized to strengthen steels by hindering the motion of dislocations. In this work, there are two types of particles that contribute to precipitation strengthening, i.e., ε-carbides and MC-type carbides. The precipitation strengthening fraction can be calculated by the following Equation (5) [25]:

$${\sigma }_p={\sigma }_p^{\epsilon -carbide}+{\sigma }_p^{MC}$$

(5)

where ${\sigma }_p^{\epsilon -carbide}$ is the precipitation strengthening increment of ε-carbide and ${\sigma }_p^{MC}$ is the precipitation strengthening increment of MC-type carbides.

The precipitation strengthening increment of ε-carbide can be calculated by the following Equation (6) [31]:

$${\sigma }_p^{\epsilon -carbide}={\displaystyle \frac{6.66}{\mathrm{D}}}\mathrm{I}\mathrm{n}{\displaystyle \frac{\mathrm{L}}{2.48\times {10}^{-4}}}$$

(6)

where L is the size of the planar intercept of the ε-carbides (length S, width W), while D is the spacing of the ε-carbides, as calculated in

Table 5. ε-carbide dimensions of the experimental steels.

Steel	ε-Carbide Dimensions							Precipitation Strengthening
	$\mathbf{Length}\left(\mathsf{\mu }\mathbf{m}\right)$			$\mathbf{Width}\left(\mathsf{\mu }\mathbf{m}\right)$			$\mathbf{D}\left(\mathsf{\mu }\mathbf{m}\right)$	According to Length (MPa)	According to Width (MPa)
	$\mathbf{Average}{\mathit{X}}_{\mathit{v}}$	$\mathbf{Standard}\mathbf{Deviation}{\mathit{\sigma }}_{\mathit{v}}$	S	$\mathbf{Average}{\mathit{X}}_{\mathit{v}}$	Standard $\mathbf{Deviation}{\mathit{\sigma }}_{\mathit{v}}$	W
Ti-steel	0.072	0.052	0.073	0.013	0.007	0.012	0.158	239	164
TiNb-steel	0.088	0.046	0.081	0.011	0.005	0.01	0.163	237	150

.

L can be calculated by the following Equation (7) [29]:

$$\mathrm{L}={\left[{\displaystyle \frac{2}{3}}\left(X_v^2+{\sigma }_v^2\right)\right]}^{{\displaystyle \frac{1}{2}}}$$

(7)

where $X_v$ is the average width or length of ε-carbides; ${\sigma }_v$ is the standard deviation of the width or length of ε-carbides, and the calculation results are presented in Table 5.

Assuming that the spacing of ε-carbides in both length and width is D, D can be calculated by the following Equation (8) [29]:

$$f={\displaystyle \frac{\mathsf{\pi }}{4}}{\displaystyle \frac{\mathrm{S}\mathrm{W}}{(\mathrm{S}+\mathrm{D})(\mathrm{W}+\mathrm{D})}}$$

(8)

where f is the volume percentage of ε-carbides, which can be obtained statistically from TEM photographs of ε-carbides, and the values are calculated to be 1.76% and 1.51% for the Ti-steel and the TiNb-steel, respectively.

The precipitation strengthening increment of MC-type carbides can be calculated by the following Equation (9) [32]:

$${\sigma }_p^{MC}=8995{\displaystyle \frac{f^{\frac{1}{2}}}{d}}\ln 2.417d$$

(9)

where d is the particle size of MC-type carbides; $f$ is the volume fraction of MC-type carbides. $f$ can be calculated from the results of the physicochemical analysis by the following Equation (10) [32]:

$$f=f_m{\displaystyle \frac{{\rho }_{Fe}}{100{\rho }_{MC}}}$$

(10)

where ${\rho }_{Fe}$ is the density of iron, which is 7.875 g/cm³; ${\rho }_{MC}$ is the density of MC-type carbides; $f_m$ is the mass fraction of MC-type carbides. The physicochemical phase analysis indicates that the average chemical formula of the Ti-steel is (Ti_0.54Mo_0.46)C, and that of the TiNb-steel is (Nb_0.12Ti_0.50Mo_0.38)C. Their densities were calculated using the weighted average density of the precipitates [33], which were NbC (7.82 g/cm³), TiC (4.93 g/cm³), and MoC (8.4 g/cm³). The calculated densities of the MC-type carbides were 6.54 g/cm³ and 6.6 g/cm³, respectively. The strengthening fractions were determined to be 250 MPa for the Ti-steel and 322 MPa for the TiNb-steel. The total precipitation strengthening could be calculated by Equation (5), and the results were 414 MPa and 472 MPa, respectively.

Based on the above-calculated strengthening components and Equation (1), the yield strength values of the Ti-steel and the TiNb-steel can be calculated as 1790 MPa and 1815 MPa, respectively, as shown in

Table 6. Calculated yield strength values of Ti-steel and TiNb-steel.

Steel	σ0 (MPa)	${\mathit{\sigma }}_{\mathit{s}\mathit{s}}$ (MPa)	${\mathit{\sigma }}_{\mathit{d}}$ (MPa)	${\mathit{\sigma }}_{\mathit{g}}$ (MPa)	${\mathit{\sigma }}_{\mathit{p}}$ (MPa)	Calculated Values (MPa)	Measured Values (MPa)
Ti-steel	39	866	751	204	414	1786	1746
TiNb-steel	39	832	788	220	472	1816	1802

. The calculated results are comparable to those measured ones with 40 MPa and 14 MPa positive deviations, respectively, which might be ascribed to the accuracy of measurements and the nonuniformity of the materials, indicating the relative effectiveness of the calculation. It is obvious that the strengthening of both the Ti-steel and the TiNb-steel are dominated by solid solution, dislocation, and precipitation strengthening. The higher yield strength of the TiNb-steel compared to the Ti-steel can be attributed to the presence of its higher levels of fine grain, dislocation, and precipitation strengthening effects, which can be attributed to the addition of the Nb element, leading to more MC-type carbide precipitates.

A number of mechanical parameters that are intrinsic to the material can be obtained from the uniaxial tensile stress (σ) strain (ε) curve. The true stress (S) and true strain (e) can be calculated by the following Equations (11) and (12), respectively [34]:

$$S=\sigma (1+\epsilon )$$

(11)

$$e=\ln (1+\epsilon )$$

(12)

The true stress and true strain for the Ti-steel and the TiNb-steel are illustrated in Figure 9. It can be observed that at the same true strain, the TiNb-steel exhibits a higher true stress than the Ti-steel.

In the uniform plastic deformation stage, the true stress (S) and the true strain (e) of the TiNb-steel conform to the following Zener–Hollomon Equation (13):

$$S=K{e}^n$$

(13)

In the equation, n is the strain hardening index between the conditional yield point and the maximum uniform strain point, and K is the strain hardening coefficient [34]. At the stage of uniform plastic deformation, the true stress and true strain were taken as natural logarithms, respectively, and the strain hardening coefficient K and strain hardening index n were obtained by linear regression. The fitting results are presented in

Table 7. Variable hardening coefficient K and strain hardening index n of the experimental steels.

Steel	Strain Hardening Coefficient K (MPa)	Strain Hardening Index n
Ti-steel	3878	0.162
TiNb-steel	4105	0.183

. It can be noticed that the TiNb-steel possesses a higher hardening coefficient K and strain hardening index n than the Ti-steel. This is associated with the higher quantity of MC-type carbides present in the TiNb-steel as the MC-type carbides have been reported to enhance the coefficient of variation hardening K and the strain hardening index n of steels [35,36]. Consequently, the tensile strength of the TiNb-steel is observed to exceed that of the Ti-steel.

Based on the above analysis, the addition of Nb to the ultra-high-strength steel leads to a higher yield strength and tensile strength, which could be attributed to the effect of more carbide precipitation, causing the results of grain refining and more dislocation. Compared to the Ti-steel, the TiNb-steel shows higher mechanical properties, providing a new method to develop ultra-high-strength steels.

5. Conclusions

In this paper, the microstructure, mechanical properties, and strengthening mechanism of two low-alloy ultra-high-strength steels microalloyed by Ti and TiNb were studied. The following conclusions can be drawn.

• Both steels showed refined microstructure with the effective grain sizes of lath martensite of 1 μm and 0.86 μm for the Ti-steel and the TiNb-steel, respectively. The finer grains of the TiNb-steel can be attributed to the addition of the Nb element, resulting in a higher concentration of MC-type carbides.
• Both steels exhibited good plasticity and relatively high toughness at the strength level of 2200 MPa. The yield strengths and tensile strengths of the Ti-steel and the TiNb-steel were tested to be 1746 MPa, 2198 MPa, and 1802 MPa, 2232 MPa, respectively. The higher yield strength of the TiNb-steel is attributed to its stronger strengthening effects due to finer grains and more precipitates, while the higher tensile strength is related to its higher coefficient of variable hardening and a higher work-hardening exponent.
• Microalloying of Ti and TiNb on the microstructure refinement and mechanical property enhancement of ultra-high-strength steels has been proved, and the addition of Nb exhibits higher strength without reducing the ductility and toughness, which provides data support for the development of new steels with improved properties.