The Nb-V-Ti-N-C System Microelements Coupling Precipitation Behavior and Its Effects on Properties in High Strength Naval Steel

Abstract

The Nb-V-Ti-N-C system microelements coupling precipitation behavior in high strength naval steel was thermodynamically analyzed. The effects of micron/nano particles on the microstructure, mechanical properties, and corrosion resistance were also studied by an in situ scanning electron microscopy (SEM) tensile test, transmission electron microscopy (TEM) analysis, and electrochemical polarization measurements. The results show that the solid solution amount of Nb, V, Ti, N, or C decreases in the steels as the temperature decreases. Carbonitrides begin to precipitate at 1506.39 °C in N1 steel, and the carbonitrides are nano-scale. Meanwhile, carbonitrides begin to precipitate at 1628.74 °C in N2 steel, which is 116.69 °C higher than the corresponding liquidus temperature of 1512.05 °C; carbonitrides with micron scale are formed in the metal melt. The tensile test revealed that with the increase in titanium content from 0.05% to 0.1%, the strength increases while the elongation decreases. The in situ SEM test results indicated that lower plasticity are associated with the carbonitrides of micron-scale, which are the micro crack sources under stress. Polarization test results indicated that pitting corrosion may easily occur at the abnormally large-sized (Nb, V, Ti)(C, N) carbonitrides.

1. Introduction

Because the service environment of naval ships is getting worse and worse, there is a great demand for improving the service life of naval ships. Naval ships demand a relatively high requirement for steels’ comprehensive performance; they require good weldability, high strength, and good plasticity–toughness, along with excellent corrosion resistance due to their service condition in the marine environment. Therefore, the main development direction of naval steel aims to ensure high strength and toughness while also meeting the extreme service environment. The design of typical traditional naval steel, such as HY-80 and HY-130 steel developed in the 1950s and 1960s, respectively, is primarily created through high carbon content, high alloy content, and quenching and tempering treatments to improve the performance of the steel [1]. However, the cost is high, and the processing cycle is too long. On the other hand, because of the high carbon content or high carbon equivalent, the welding difficulty and sensitivity of stress corrosion of these HY series steels in marine environment obviously increase with the improvement of strength. After the 1980s, low-carbon Cu-bearing HSLA steels have been developed, which can improve the weldability and lower the risk of corrosion cracking [2,3]. Nano-sized copper can precipitate within ferrite in copper-containing iron-based alloys and steels during the aging treatment to compensate for the strength loss caused by the decreased carbon content. However, these nano-sized Cu precipitations present low thermal stability and easily coarsen with over-aging at temperatures above 550 °C, imposing a negative impact on the ductility and toughness of the steel [4,5,6,7].

The scientific design of the secondary phase in steel is the main method for improving its performance; to obtain the best comprehensive performance, the volume fraction, size, shape, and distribution of the secondary phase in steels must be accurately controlled as required. However, no matter what type of secondary phase is formed, it is generated by the mutual coupling reaction among multiple microelements in steel. Therefore, in order to improve the comprehensive effect of the secondary phase in steels, it is necessary to comprehensively consider the design of the multiple microelements system. Microelements are the key factors that determine the properties of high strength and toughness steels; Pouraliakbar et al. [8] and Khalaj et al. [9] developed artificial neural network models to predict the toughness of high strength low alloy steels or to predict the bainite fraction. The inputs of the neural network included the weight percentage of 15 alloying elements, and for the tensile test results, 118 different steels from API X52 to X70 grades were used; good practical results were acquired. In terms of the latest research, Tang et al. have adopted an endogenous method of controlling the oxide reaction and solute concentration distribution in the process of deoxidization to obtain a high density of in situ, homogeneously dispersed nano titanium oxide particles to compensate for the strength loss caused by the decreased carbon content [10,11,12]. These in situ titanium oxide nanoparticles have been proven to exhibit a good precipitation strengthening effect, heterogeneous nucleation effect, and thermal stability. The pivotal issue for the above-mentioned results lies in the appropriate quantitative addition of titanium, which determines the reaction of titanium with oxygen in the steel to produce nano-scale titanium oxide particles. Too low of a Ti content provokes inadequate precipitation strengthening, whereas too much Ti will result in the redundant Ti reacting with C or N, generating large-sized inclusions, i.e., TiN, Ti(N, C), or (Ti, V, Nb)(N, C), etc. It is well-known that these large-sized inclusions present irregular shapes and easily become the crack sources under stress [13,14]. This will severely weaken the plasticity, toughness, and fatigue properties of the steels. In the meantime, pitting corrosion would easily occur at the abnormal large-volume inclusions in an environment containing Cl⁻, reducing the corrosion resistance of the steel. Recently, Xu et al. studied the optimum content of titanium in micro-alloyed steel and concluded that the optimum range is 0.04–0.1 wt.%. The authors only studied the change of strength in this range, and no attention was paid to the reason of the plasticity decrease or the damage to other performance factors [6]. However, as it is well-known that plasticity and toughness change with strength, it is necessary to investigate the effect of the content of microelements on the plasticity and toughness, among other properties. Therefore, the objective of this study was to examine the Nb-V-Ti-N-C system microelements coupling the precipitation behavior in high strength naval steels. The thermodynamic problems of the equilibrium solution for multivariate secondary phases was researched, and the effects of micron/nano-sized particles on the microstructure, mechanical properties, and corrosion resistance for the naval steels with different microelement contents were also systematically researched. It is sensible to scientifically optimize the content of microelements and then to improve the product’s comprehensive performance for naval steels.

2. Materials and Methods

2.1. Preparation of the Naval Steel

The chemical compositions of the compared microparticles/nanoparticles-strengthened naval steels are shown in

Table 1. Chemical compositions of the tested high strength naval steel (Wt %).

Naval Steel	C	Si	Mn	Cr	O	V	Ti	Nb	N	P	S	Fe
N1 steel	0.05	0.03	0.85	0.45	0.0018	0.030	0.05	0.06	0.0012	0.0075	0.0018	98.46
N2 steel	0.05	0.028	0.86	0.46	0.0016	0.032	0.10	0.06	0.0014	0.007	0.0018	98.39

. These two steels with 0.05 and 0.1 wt.% Ti are named N1 steel and N2 steel, respectively. Firstly, the base metal was put in a vacuum melting furnace. After the metal was fully melted and the solidification temperature of molten steel was in the range of 1530–1580 °C, the pure Ti wires were added by multi-point regional micro-supply processing. The detailed fabrication process and relevant information on the materials fabrication can also be found elsewhere [10,15]. Electromagnetic stirring (≈4 kHz frequency) was always present during the whole melting process. Under the effect of electromagnetic stirring, the flowing linear velocity of the molten metal was about 10–20 ms⁻¹. After all of the pure Ti wires were completely melted, the thermal insulation held for about 1 min, which is conducive to the uniform distribution of this element in melt. In the casting process, the flowing linear velocity of the molten metal was no less than 1.7 ms⁻¹. Secondly, the melt was cooled down inside the crucible in air at a cooling rate of about 500 °C min⁻¹ in the solidification process. Subsequently, two ingots with dimensions of 135 mm × 75 mm × 96 mm were obtained. Then, the thermo-mechanical control process (TMCP) was conducted. Two ingots were austenitized at 1200 °C for 2 h, then hot rolled into 11 mm thick plates by seven passes. The initial and final rolling temperatures were 1150 °C and 850 °C, respectively. Finally, the rolled plate was water-cooled to a temperature of ~750 °C and then air-cooled to the ambient temperature; the preparation process flow chart is as shown in Figure 1.

2.2. In Situ SEM Tensile Tests

Samples for tensile tests were prepared according to Figure 2, which is the schematic drawing for in situ SEM tensile testing; the samples were designed to fit in the tensile machine. First, dog bone-shaped tensile specimens were cut using electro-discharge machining followed by mechanical polishing. Then, the specimens were further polished to a mirror finish with a final thickness of about 0.8 mm, followed by etching. Specimens were put into an SM-TS40 tensile test apparatus in an JXA-840 SEM with a maximum load of 1 KN. All samples were tested at room temperature. However, it was not possible to determine the strain rate as only a small section of the sample was elongated and the strain rate was not invariable throughout the test.

2.3. Corrosion Resistance Experiments

To compare the electrochemical polarization resistance of the in situ nanoparticle-strengthened naval steels, electrochemical polarization measurements were performed using a PS-268A potentiate and a three-electrode cell. The counter electrode used here was graphite, and the reference electrode was a saturated calomel electrode (SCE). All potentials cited in this study will henceforth be in reference to the SCE. The electrolyte used was 3.5 wt.% NaCl solution, which was prepared from a reagent-grade chemical and distilled water. The samples were immersed in 3.5 wt.% NaCl solution for at least 20 min to attain a stable open circuit. The polarization curves were then measured with the cell open to air at room temperature with a 20 mv min⁻¹ potential sweep rate. An 9XB-PC optical microscope and Zeiss Supra 55 field emission scanning electron microscope (SEM) were used to characterize the microstructure of the investigated steels.

2.4. TEM Experiments

The microstructure of the naval steel was further investigated using a JEOL JEM-2100 TEM microscope operated at 120 kV. TEM foil was prepared by cutting thin wafer from the steel samples and was mechanically thinned to be ~35 μm in thickness. Three-millimeter discs were punched from the foils and electrochemical-polished using a solution of 10 vol.% HClO₄ methanol electrolyte at low temperature, followed by further ion-thinning for obtaining an electron-transparent area.

2.5. Method of Thermodynamics Analysis in the Nb-V-Ti-N-C System

The microalloying elements in combination with C or N can form multiple secondary phases. These multiple secondary phases are composed of carbides and nitrides with similar crystal structures, exhibiting continuous or extended mutual solubility. Therefore, the resulting multivariable secondary phases can be described by the chemical formula (M₁, M₂, M₃) (C, N). Valid for small element concentrations, the microalloying elements M₁, M₂, and M₃, as well as the interstitial elements C and N, form dilute solutions in the matrix, and their activities obey Henry’s law. In this case, the effective activity coefficients of the components M₁C, M₂C, M₃C, M₁N, M₂N, and M₃N were assumed to be k₁, k₂, k₃, m₁, m₂ and m₃, respectively, and the total molar fraction of the M_1(k1+m1)M_2(k2+m2)M_3(k3+m3)C_(k1+k2+k3)N_(m1+m2+m3) carbonitride formed in steel was t moles [16]. This carbonitride can be seen as a mixture of the following amounts of pure binary carbides and nitrides: k₁t mole M₁C, k₂t mole M₂C, k₃t mole M₃C, m₁t mole M₁N, m₂t mole M₂N, and m₃t mole M₃N. Therefore, based on the chemical equilibrium, the thermodynamics analysis model and computing method of equilibrium solution for the multivariable secondary phase in steels were developed according to the mass balance and solubility product equations for the quaternary secondary phase or more. Therefore, the solid solution precipitation of the secondary phases formed in steel is Nb_{(k1 + m1)}V(_{k2 + m2)}Ti_{(k3 + m3})C_{(k1 + k2 + k3)}N_{(m1 + m2 + m3)} in the Fe-M′(Nb, V, Ti)-N-C system micro-alloy steel, and the coefficient of solid solubility of the product is taken as in reference [16]. Therefore,

$$\lg \left\{\frac{\left[M_1\right]\left[C\right]}{k_1}\right\}=\lg K_{M_{1}C}-{\displaystyle \sum _{i=1}^n{e}_{M_1}^i\left[i\right]}-{\displaystyle \sum _{i=1}^n{r}_{M_1}^i{\left[i\right]}^2}-{\displaystyle \sum _{i=1}^n{r}_{M_1}^{i,M_1}\left[i\right]\left[M_1\right]}-{\displaystyle \sum _{i=1}^n{e}_C^i\left[i\right]}-{\displaystyle \sum _{i=1}^n{r}_C^i{\left[i\right]}^2}-{\displaystyle \sum _{i=1}^n{r}_C^{i,C}\left[i\right]\left[C\right]}$$

(1)

$$\lg \left\{\frac{\left[M_2\right]\left[C\right]}{k_2}\right\}=\lg K_{M_{2}C}-{\displaystyle \sum _{i=1}^n{e}_{M_2}^i\left[i\right]}-{\displaystyle \sum _{i=1}^n{r}_{M_2}^i{\left[i\right]}^2}-{\displaystyle \sum _{i=1}^n{r}_{M_2}^{i,M_2}\left[i\right]\left[M_2\right]}-{\displaystyle \sum _{i=1}^n{e}_C^i\left[i\right]}-{\displaystyle \sum _{i=1}^n{r}_C^i{\left[i\right]}^2}-{\displaystyle \sum _{i=1}^n{r}_C^{i,C}\left[i\right]\left[C\right]}$$

(2)

$$\lg \left\{\frac{\left[M_3\right]\left[C\right]}{k_3}\right\}=\lg K_{M_{3}C}-{\displaystyle \sum _{i=1}^n{e}_{M_3}^i\left[i\right]}-{\displaystyle \sum _{i=1}^n{r}_{M_3}^i{\left[i\right]}^2}-{\displaystyle \sum _{i=1}^n{r}_{M_3}^{i,M_3}\left[i\right]\left[M_3\right]}-{\displaystyle \sum _{i=1}^n{e}_C^i\left[i\right]}-{\displaystyle \sum _{i=1}^n{r}_C^i{\left[i\right]}^2}-{\displaystyle \sum _{i=1}^n{r}_C^{i,C}\left[i\right]\left[C\right]}$$

(3)

$$\lg \left\{\frac{\left[M_1\right]\left[N\right]}{m_1}\right\}=\lg K_{M_{1}N}-{\displaystyle \sum _{i=1}^n{e}_{M_1}^i\left[i\right]}-{\displaystyle \sum _{i=1}^n{r}_{M_1}^i{\left[i\right]}^2}-{\displaystyle \sum _{i=1}^n{r}_{M_1}^{i,M_1}\left[i\right]\left[M_1\right]}-{\displaystyle \sum _{i=1}^n{e}_N^i\left[i\right]}-{\displaystyle \sum _{i=1}^n{r}_N^i{\left[i\right]}^2}-{\displaystyle \sum _{i=1}^n{r}_N^{i,N}\left[i\right]\left[N\right]}$$

(4)

$$\lg \left\{\frac{\left[M_2\right]\left[N\right]}{m_2}\right\}=\lg K_{M_{2}N}-{\displaystyle \sum _{i=1}^n{e}_{M_2}^i\left[i\right]}-{\displaystyle \sum _{i=1}^n{r}_{M_2}^i{\left[i\right]}^2}-{\displaystyle \sum _{i=1}^n{r}_{M_2}^{i,M_2}\left[i\right]\left[M_2\right]}-{\displaystyle \sum _{i=1}^n{e}_N^i\left[i\right]}-{\displaystyle \sum _{i=1}^n{r}_N^i{\left[i\right]}^2}-{\displaystyle \sum _{i=1}^n{r}_N^{i,N}\left[i\right]\left[N\right]}$$

(5)

$$\lg \left\{\frac{\left[M_3\right]\left[N\right]}{m_3}\right\}=\lg K_{M_{3}N}-{\displaystyle \sum _{i=1}^n{e}_{M_3}^i\left[i\right]}-{\displaystyle \sum _{i=1}^n{r}_{M_3}^i{\left[i\right]}^2}-{\displaystyle \sum _{i=1}^n{r}_{M_3}^{i,M_3}\left[i\right]\left[M_3\right]}-{\displaystyle \sum _{i=1}^n{e}_N^i\left[i\right]}-{\displaystyle \sum _{i=1}^n{r}_N^i{\left[i\right]}^2}-{\displaystyle \sum _{i=1}^n{r}_N^{i,N}\left[i\right]\left[N\right]}$$

(6)

$$\frac{M_1}{A_{M_1}}-\frac{\left[M_1\right]}{A_{m_1}}=\left(k_1+m_1\right)t$$

(7)

$$\frac{M_2}{A_{M_2}}-\frac{\left[M_2\right]}{A_{m_2}}=\left(k_2+m_2\right)t$$

(8)

$$\frac{M_3}{A_{M_3}}-\frac{\left[M_3\right]}{A_{m_3}}=\left(k_3+m_3\right)t$$

(9)

$$\frac{C}{A_C}-\frac{\left[C\right]}{A_C}=\left(k_1+k_2+k_3\right)t$$

(10)

$$\frac{M_1}{A_{m_1}}+\frac{M_2}{A_{m_2}}+\frac{M_3}{A_{m_3}}+\frac{C}{A_C}+\frac{N}{A_N}=\frac{\left[M_1\right]}{A_{m_1}}+\frac{\left[M_2\right]}{A_{m_2}}+\frac{\left[M_3\right]}{A_{m_3}}+\frac{\left[C\right]}{A_C}+\frac{\left[N\right]}{A_N}+2t$$

(11)

k₁ + k₂ + k₃ + m₁ + m₂ + m₃ = 1

(12)

where M₁, M₂, M₃, C, and N are the mass percentages of Nb, V, Ti, C, and N, respectively, Am₁, Am₂, Am₃, A_C, and A_N are the atomic weights of Nb, V, Ti, C, and N, respectively, and A_Nb = 92.9, A_V = 50.9, A_Ti = 47.9, A_N = 14, A_C = 12, [M₁], [M₂], [M₃], [C], and [N] are the concentrations (in wt.%) of the respective elements dissolved in the solution. $e_i^j$, $r_i^j$, $r_i^{j,i}$ and $\left[i\right]$, denote the first order and second order interaction parameters between i and j, the cross-product term, and the concentration of i in liquid iron (in mass %), respectively. It was assumed that the second order parameter and cross-product term could be ignored in the present work. T is the temperature, and t is the total molar fraction of the carbonitride Nb_{(k1 + m1)}V_{(k2 + m2)}Ti_{(k3 + m3)}C_{(k1 + k2 + k3)}N_{(m1 + m2 + m3)} formed in the steel; thus, these multivariate secondary phases can be seen as a mixture of the following amounts of pure carbides and nitrides: k₁t mole NbC, k₂t mole VC, k₃t mole TiC, m₁t mole NbN, m₂t mole VN, and m₃t mole TiN. Equations (1)–(12) have twelve unknowns, which are numerically solved for to determine the equilibrium state. For a given steel at any appropriate temperature, the equilibrium matrix composition, precipitate composition, and precipitate volume fraction can be determined (i.e., [M₁], [M₂], [M₃], [C], [N], k₁, k₂, k₃, m₁, m₂, m₃, and t). Using [17],

$$\log K_{\mathrm{Nbc}}=2.96-7510/T$$

(13)

$$\log K_{\mathrm{NbN}}=2.80-8500/T$$

(14)

$$\log K_{\mathrm{VC}}=6.72-9500/T$$

(15)

$$\log K_{\mathrm{VN}}=3.63-8700/T$$

(16)

$$\log K_{\mathrm{TiC}}=2.75-7000/T$$

(17)

$$\log K_{\mathrm{TiN}}=0.32-8000/T$$

(18)

The main elements interaction activity coefficient of the solid solubility of the product is taken as in

Table 2. Selected interaction activity coefficient in dilute solutions of steel.

Element j	C	N	O	Mn	Cr	S	P	Si	Ti	V	Nb
$e_{Nb}^j$	−791.66/T + 0.1142 [18,19]	−1720/T + 0.503 [19]	−22066/T + 11.01 [19]	698/T − 0.37 [19]	−1007.94/T + 0.6566 [19,20]	−24.349/T [19]	−22.476/T [19]	666.79/T − 0.3876 [19,20]	30.53/T − 0.006 [18,19]	−0.0039 + 23.71/T [18,19]	16.74/T - 0.0071 [18,19]
$e_{Ti}^j$	−221/T − 0.072 [21]	−19500/T + 8.37 [22]	−2098/T + 0.0943 [23]	−0.043 [21]	−0.016 [19]	−0.27 [21]	−74.92/T [19]	177.5/T − 0.12 [24]	212/T − 0.0640 [25]	28.416/T + 0.0032 [18,19]	15.74/T - 0.00314 [18,19]
$e_V^j$	−571.75/T + 0.0644 [18,19]	−1270/T + 0.33 [19]	−7950/T + 3.20 [19]	6.2427/T + 0.000146 [18,19]	—	−29.968/T [19]	−43.079/T [19]	162.74/T − 0.0385 [18,19]	30.196/T + 0.00313 [18,19]	470/T − 0.22 [26]	13.28/T - 0.0022 [18,19]
$e_N^j$	0.06 [27]	109/T [19,28]	0.05 [27]	−36.75/T − 0.123 + 0.016 In T [19,20]	−303.8/T + 0.112 [19,20]	0.007 [27]	167/T − 0.038 [19]	−286/T + 0.202 [24]	−4070/T + 1.643 [19]	−356/T + 0.0973 [29]	−260/T + 0.0796 [19]
$e_C^j$	158/T + 0.0581 [19]	100.29/T [17,19]	−0.34 [27]	−300/T + 0.154 [19]	−102/T + 0.0327 [19,30]	0.046 [27]	1190/T − 0.608 [19]	162/T − 0.008 [31]	−55/T − 0.015 [21]	−134.79/T + 0.0185 [18,19]	−176.76/T [19,20]

—: Values not found in the literature; they are assumed to be zero in the current calculation.

. It was known that, for a given multicomponent steel Nb-V-Ti-N-C system’s microelements at any appropriate temperature, the equilibrium matrix composition, precipitate composition, and precipitate total molar fraction could be determined (i.e., [Ti], [Nb], [V], [N], [C], and t). Furthermore, the numerical iteration calculation process was carried out in MATLAB 9.0 software based on the fixed-point iteration method.

3. Results

3.1. Microstructure Transformation and Micro-Cracks Sources

Figure 3 and 4 show the microstructure of the N1 and N2 steels at different tempering temperatures ranging from 450–700 °C for 0.5 h. The results indicate that the microstructure of the specimens were mainly lath bainite, granular bainite, and polygonal ferrites. However, with the increase in tempering temperature, the bainite lath gradually coarsened. As shown in Figure 3e and Figure 4e, the lath boundaries blurred and merged together, and the amount of filminess retained austenite that was between the laths also gradually decreased when the tempering temperature was 650 °C and 700 °C; the boundary between the strips became blurred and gradually merged, and its microstructure changed into a granular or island structure, as shown in Figure 3f and Figure 4f.

The microstructure transformations for the two steels during the tensile tests under different stresses are shown in Figure 5 and Figure 6, respectively. There were many micron-scale inclusions, and the inclusions spanned about 2–4 μm across in N2 steel; however, there were minimal micron-scale inclusions in N1 steel. Therefore, according to many data statistics and the visual view of the graph, the quantity of inclusions in N2 steel was observed to significantly higher than that of the inclusions in N1 steel. As for N1 steel, Figure 5 shows that the morphology changed upon tensile loadings under the stress of 879 N. Necking has occurred prior to the fracture (Figure 5a). Thin straight slip lines initiate in the matrix and start to coarsen to form slip bands and steps. New multiple slip lines and cross-slip lines form gradually as the tensile load increases. At the initial stage of tension, slipping is along in essentially only one direction; that is, only one slip system is dominant [32,33]. As the tensile test proceeds, additional slip systems become active and space between the slip bands become larger, as shown in Figure 5b–d. Micro-cracks initiate at different locations while the slip line appears. There are three main sites of crack initiation: the surface defect or stress focus sites of the irregular inclusion, the matrix near the inclusion, and the interface between the inclusion and the matrix. It can be seen from Figure 5b,d that micro-cracks mainly developed between the slip planes. Beyond that, there were a few voids initiated at inclusions, as Figure 5c shows. As for the ellipsoidal inclusions, the void mainly initiated at the interface. The reason for this kind of crack initiation is that the interface bonding force between this kind of inclusion and the matrix is weak. In addition, the interface stress concentration also contributed to the crack initiation.

The stress concentration existed before loading and became higher after loading. This phenomenon is also reflected in Figure 6. Through further observation, it is obvious to see that the surface defect or stress focus part of the irregular inclusion were the main forms of crack initiation in N2 steel, (as Figure 6b,d show). Along with the energy dispersive spectroscopy (EDS) results presented in Figure 6e, the EDS result indicates that the secondary phases were mainly titanium carbonitride. Figure 6d shows that primarily titanium carbonitrides with irregular shape can be cracked during the tensile deformation. Defects with this type of irregular shape will result in a non-uniform distribution of stress and strain in the process of tensile deformation and make the deformation of matrix and inclusion inharmonious. Finally, the stress concentration is too high on the surface of carbonitride inclusion, which makes those inclusions particle crack [34].

3.2. Mechanical Properties

Figure 7 shows the variations of the mechanical properties of N1 and N2 steels under different tempering temperatures. The tensile strength and yield strength of N2 steel was higher than that of N1 steel. However, the elongation of N2 steel was much lower than that of N1 steel; the strength increase was attributed to the increase in particle density and the interactions between the moving dislocations and these particles. The elongation decreases were caused by the emergence of large-sized irregular inclusions because of excess titanium content. With the increase in tempering temperature, the tensile strength of both steels first decreased, then increased, and finally decreased, and the yield strength and elongation basically first increased and then decrease. Especially when the tempering temperatures was 700 °C, the tensile strength, yield strength, and elongation greatly decreased, which may be related to the precipitation behavior of the microelements. So, in order to pursue a better comprehensive mechanical property, the content of each microelement should be scientifically optimized, especially Ti, N, and C to avoid the occurrence of large-sized irregular inclusions in the steels.

3.3. Corrosion Resistance

The polarization curves for the N1 and N2 steels obtained in 3.5 wt.% NaCl solutions are presented in Figure 8. The corresponding corrosion current (I_corr) and open circuit potential (E_corr) of the corresponding steels were determined by Tafel slope extrapolation, and the values are listed in

Table 3. Selected interaction activity coefficient in dilute solutions of the steels.

Sample	Ecorr (mV vs. Ag/AgCl)	Icorr (mA/cm2)
N1	−0.383 ± 0.005	6.49 ± 0.1
N2	−0.592 ± 0.005	8.09 ± 0.1

. The low-magnified optical micrographs of the investigated steel after the polarization experiments were conducted. It is obvious that both groups of contrast steel showed pitting corrosion, which is consistent with the polarization curve result. It is obvious that the polarization curves of N1 and N2 steels show a turning point in the polarization potential range between −850 mv and 800 mv. The potential at the turning points is the pitting potential. The polarization current slowly changes with the increase of polarization voltage in the inception particles, and then quickly increases as the polarization voltage exceeds the pitting potential.

The I_corr of N1 and N2 steels were 6.49 mA/cm² and 8.09 mA/cm², respectively. Apparently, the I_corr of N1 steel was lower than that of N2 steel. In terms of voltage, N2 steel showed an E_corr shift of 35 mv towards the cathodic direction compared with N1 steel. I_corr and E_corr are kinetic and thermodynamic parameters; they represent the rate of corrosion and the degree of corrosion, respectively. Here, I_corr indicates that the corrosion rate of N1 steel was obviously slower than that of N2 steel. Meanwhile, E_corr suggests that it became more difficult for N1 steel to be corroded than N2 steel. The low-magnified optical micrographs were also consistent with the corrosion measurement data, where it is obvious that the pitting holes were small and shallow and the corrosion depth was relatively evenly distributed in the matrix for N1 steel. However, for N2 steel, those pitting holes were big and deep, and the corrosion depth was unevenly distributed. The relatively better corrosion resistance of N1 steel was mainly attributed to two aspects. First, Ti is a corrosion-resistant element. Xu. et al. have reported the effect of Ti on improving the corrosion resistance in their study [6]. Titanium is evenly distributed in the matrix and rust layer and thus prevents the entry of oxygen. Furthermore, the in situ nano dispersion particles can absorb aggressive chloride anions and improve corrosion resistance in the aggressive environment.

For the Fe-M′ (Nb, V, Ti)-N-C system multivariate secondary phase in high strength naval steels, the multivariate secondary phase precipitates are (Nb, V, Ti) (C, N), and the atomic model of the Fe-Nb-V-Ti-N-C system coupling precipitation is shown in Figure 9. Based on the thermodynamics analysis model, the thermodynamic problems of equilibrium solution for the Fe-M′(Nb, V, Ti)-N-C system micro-alloyed steels were systematically researched to calculate and analyze the equilibrium thermodynamic state, including the concentrations of the respective elements in solution ([Nb], [V], [Ti], [N], and [C]) from 750 °C to complete dissolution temperature; the total molar amount (t) of the multivariate secondary phases was also calculated and analyzed. In this formula, [Nb], [V], [Ti], [N], and [C] are the solid-dissolved contents of Nb, V, Ti, N, and C in the steels at different temperatures [16]. The N1 steel was thermodynamically analyzed; its complete solution temperature T_AN was 1506.39 °C, and the solid-dissolved contents of different elements of the steel at different temperatures and the total molar amount (t) of the multivariate secondary phases at different temperatures are shown in Figure 10. The solid-dissolved contents of different elements decreased with the decrease in temperature, especially with respect to the change in N and Ti contents. The total molar amount (t) gradually increased with the temperature decreasing, and the [Nb] was 4.9399 × 10⁻⁴%, the [V] was 0.01486752%, the [Ti] was 0.00139658%, the [N] was 9.65 × 10⁻⁷%, the [C] was 0.02759748%, and the total molar amount t was 0.00195252 mol at 750 °C.

The multiple microelements coupling precipitation behavior was thermodynamically analyzed in N2 steel, and its complete solution temperature T_AN was 1628.74 °C, which is 116.69 °C higher than the corresponding liquidus temperature of 1512.05 °C [16,35]. As is known for micro-alloy steels, if the complete dissolution temperature is above the liquidus temperature, carbonitrides constitutional liquation will occur, so it is obvious that a strong carbonitrides constitutional liquation has already occurred in N2 steel, and the micron-scale carbonitrides begin to precipitate in the metal melt, which is consistent with the Figure 6d. In contrast, for N1 steel, the complete dissolution temperature was 1506.39 °C by calculation, which is lower than the liquidus temperature. Therefore, no carbonitrides constitutional liquation would occur in N1 steel, and the nano-scale carbonitrides. The solid-dissolved contents of different elements in the steel at different temperatures and the total molar amount (t) of the multivariate secondary phases at different temperatures are shown in Figure 11. The solid-dissolved contents of different elements also decreased with the decrease in temperature, especially with respect to the change in N contents at high temperature, the change in V at low temperature. Meanwhile, there was a change in Nb, Ti, and C contents at 1200–800 °C, and the [Nb] was 5.724 × 10⁻⁴%, the [V] was 0.01600606%, the [Ti] was 0.00336181%, the [N] was 2.713 × 10⁻⁷%, the [C] was 0.01584288%, and the total molar amount t was 0.00293212 mol at 750 °C.

Many researchers have observed hydrogen trapping sites in precipitation-hardening steel containing nano-sized precipitates and have concluded that the deuterium atoms are located on the boundary surface of these nano precipitates, which indicates that the broad interface between the matrix and the precipitates is the main trapping site [10,32]. We believe that the adsorption of Cl⁻ by nanoparticles occurs under a similar reaction. There is a large number of in situ nanoparticles formed in N1 steel, and they are homogeneously dispersed in the matrix [11,12], just as Figure 12a shows. As the N1 steel is in the chloride-containing (3.5 wt.% NaCl solution) environments, the location of in situ nanoparticles will be trapping sites for the Cl⁻ infiltrating into the matrix, and those Cl⁻ will be absorbed by in situ nanoparticles [11,12,32]. Due to the dispersion distribution of the in situ nanoparticles in the matrix, the Cl⁻ infiltrating into the matrix will also become scattered, just as Figure 12b shows. It follows that the corrosion effect of Cl⁻ is weakened, which is due to the dispersion effect of nanoparticles to Cl⁻. Therefore, the combination of these two factors increases the corrosion resistance of steel containing in situ nanoparticles.

However, if there is too much Ti content in steel, small occlusion areas will be formed between these large inclusions and the matrix boundary. Thus, an occluded cell will form at these small occlusion areas, which will accelerate the corrosion rate in these areas. This is why the corrosion resistance of N2 steel is decreased. Therefore, it is important to select the content of each microelement, accurately control the size and shape of the multivariable secondary phase, and then improve the product comprehensive performance.

The thermodynamic analysis results show that the solid solution amount of each microelement decreased as the temperature decreased in high strength naval steel, and the complete solution temperature T_AN was 1506.39 °C for N1 steel; thus, the carbonitrides began to precipitate at this temperature, and the carbonitrides were nano-scaled in N1 steel. As the complete solution temperature T_AN was 1628.74 °C for N2 steel, which is 116.69 °C higher than the corresponding liquidus temperature of 1512.05 °C, and carbonitrides constitutional liquation has occurred in this steel, the performance of N1 steel was better than that of N2 steel, which had micron-scale carbonitrides. This is significant to provide a scientific basis for optimizing the chemical composition during the micro-alloy composition design through thermodynamic analyses.

4. Conclusions

The Nb-V-Ti-N-C system microelements coupling precipitation behavior was thermodynamically analyzed in high strength naval steel. It was observed that the solid solution amount of Nb, V, Ti, N, or C in the steels decreased as the temperature decreased, the carbonitrides began to precipitate at 1506.39 °C in N1 steel, and the carbonitrides began to precipitate at 1628.74 °C in N2 steel, which is 116.69 °C higher than corresponding liquidus temperature of 1512.05 °C; micron-scale carbonitrides formed in the metal melt. The strength increased while the elongation decreased with the increase in titanium content from 0.05% to 0.1%. The in situ SEM results indicated that the elements Nb, V, and Ti reacted with the C and N in steel, and a lower plasticity was associated with the micron-scale carbonitrides. Many of the surface micro-cracks were associated with the large-sized and irregularly shaped inclusions, which was the main cause of the elongation decreases. Therefore, it is logical to scientifically optimize the content of each microelement, consider the low cost, and accurately control the size and shape of the multivariable secondary phase through thermodynamic analyses, thus improving the product comprehensive performance for micro-alloyed steels.