Text|Yanagihachihara
Editor|Yanagihachihara
With the increasing demand for high-strength and lightweight materials in modern industry, ultra-high-strength steels have become one of the hot spots in materials research and application fields. To meet the requirements for high strength and durability, researchers have been exploring various ways to improve the properties of ultra-high strength steels.
Thermomechanical treatment and microalloying are widely used in the development of ultra-high strength steels and show the potential to have a significant impact on the hardenability of materials. By controlling the heating, deformation and cooling processes, thermomechanical treatment can significantly change the grain structure and phase change behavior of ultra-high strength steel, thereby adjusting its mechanical properties and hardenability.
On the other hand, microalloying technology, especially the addition of trace amounts of Nb elements, also plays an important role in improving the performance of ultra-high-strength steels. As a reinforcing element, NB refines the grains and improves the cold deformation and heat treatment response of the material, thereby enhancing the strength and toughness of the steel and improving its hardenability.
So let's investigate together what effect thermomechanical treatment and Nb microalloying have on the hardenability of ultra-high strength steel!
Investigative materials
In the study, we experimented on three different steels for the effect of TM machining on hardenability. The melt of these steels is produced in induction furnaces according to the components listed in Table I, and these steels represent hardened steels with a carbon content of 0.17 pct. Steel 1 and 2 delay the γ-α transition by alloying Mn, Si and B to ensure martensitic microstructure under corresponding cooling conditions. The only difference between the steels is the content of Nb in order to check the effect of Nb on the hardenability after TMP.
In steel 3, the content of Mn is reduced, but it is still alloyed using Microalloying Element (MAE) V to optimize the TM machining route. In addition, B is omitted from steel 3, but modified by using high content Cr, Ni, Mo and Cu to compensate for the softening effect during tempering.
These studies aim to understand the effect of TM machining on hardenability and the role of different alloying elements in this process. By adjusting the content and type of alloying elements, the phase change behavior and microstructure formation of the steel can be controlled, thus affecting its hardenability properties.
For Steel 1 and Steel 2, the alloying elements introduced help delay the γ-α transition, allowing more martensitic phases to form during cooling. In steel 3, by reducing the Mn content and alloying with MAEV, the TM machining route is optimized to further improve the hardenability performance.
In addition, B is omitted from steel 3 and modified with high content of Cr, Ni, Mo and Cu. The addition of these elements is intended to improve the tempering stability of the steel to avoid excessive softening. With these optimized alloy designs and treatment paths, we hope to achieve precise control of the hardenability of the steel.
Experimental procedures and sample preparation
To prepare the material for investigation, we pre-rolled the raw materials and extracted an expandometer sample measuring 10 mm long and 5 mm in diameter by wire etching. To reduce heat conduction between the sample and the deformation stamp, we used tungsten sheets and installed an "S-type" thermocouple on the sample to determine the sample temperature during deformation.
To reproduce the TM rolling scheme, we used the deformation expander Bähr805A/D and performed the deformation sequences listed in Table II. We investigated two different deformation temperature ranges (FRTs), and the deformation procedure included a deformation with an overall compression ratio of 1.0. After 1250 min of solution annealing at 5 °C, subsequent deformation is performed in 5 batches, starting at 1000 °C and ending at 875 °C (rolling scenario FRT1), or starting at 1075 °C and target temperature at 950 °C (rolling scenario FRT2).
During the deformation process, the sample is cooled and quenched for an additional 3 seconds to simulate the actual situation in the Industrial TM rolling process, as quenching usually takes place within seconds after finishing rolling. We performed five different cooling rates: λ=1, 3, 10, 30 and 100K/s.
To compare the difference in hardenability after re-austenitization (Q+T route) of steel samples, we performed austenitization at 930°C for a duration of 5 minutes for prerolled samples. The sample is then quenched at the same cooling rate as the deformed sample. Through the expansion data, we analyze the temperature of the phase transition using the length expansion ΔL/L that occurs at the phase transition, where we use the three-cut method to determine the beginning and end of the phase transition, which are 5% and 95%, respectively.
Before the hardness test, the samples were hot-embedded and ground for at least 30 seconds using SiC paper with 320 to 4000 particle sizes. Subsequently, polishing was performed for at least 3 minutes using 3 μm diamond slurry, followed by 30 sec polishing using 1 μm slurry. Five HV10 hardness measurements were performed on each sample and it was ensured that the hardness measurement point was centered on the sample.
To study residual austenite grains (PAGs) in microstructures, we etched the samples using picric acid etch. The composition and procedure of PAG are described in detail in Reference 34. Using the image analysis software MIPAR,™ we determined the equivalent grain diameter and aspect ratio of the PAG. For the microstructure analysis of the transformed microstructure, we polished the samples using silicate polishing (StruersOPS) for 10 min, followed by 35 sec using electrolytic ablation at 5 V.
Subsequently, we immersed the sample in Nital etching agent and by optical microscopy and FIBVersaFEI3D dual-beam scanning electron microscopy (SEM), we recorded images of the microstructure. The properties of the newly formed NbC precipitate were evaluated by performing EDS analysis using the EDAXOctanePlus detector and the TEAM4.3 software package.
The figure shows the effect of the process route on hardenability. For steel 1, FRT has no significant effect on hardening behavior, and the hardness progression is consistent with the cooling rate. However, the lower FRT (433HV10) and samples quenched without deformation at a maximum cooling rate of 435 K/s (10HV100) showed significantly increased hardness relative to hardness at 448°C (10HV950).
Steel 3 shows a similar trend. At a lower FRT (3°C), the hardness increases significantly at a high cooling rate (>875K/s), where the hardness of a sample with a cooling rate of 472K/s is 10HV100, while the hardness of a deformed sample with an FRT of 460°C is 10HV950, and the hardness of an undeformed ordinary quenched sample is 451HV10.
However, at a slower cooling rate of 1K/s, both the undeformed quenched sample and the compressed sample with increased FRT had higher hardness values of 413HV10 and 417HV10, respectively, while the deformed sample with an FRT of 372°C had a hardness of 10HV875.
These results show that by adjusting the FRT and cooling rate, the hardenability and hardness of the steel sample can be significantly affected. High cooling rates and lower FRTs tend to increase hardness, while slower cooling rates can result in relatively low hardness. These findings provide an important reference for further optimizing the process route to achieve the desired material properties.
The study compares the difference in hardening behavior between steel 1 and steel 2, and reveals the influence of Nb elements in steel on quenching behavior. The results show that steel 1 has a similar hardness level at different cooling rates, while steel (steel 3) with Nb alloy can only reach a lower hardness level at lower cooling rates. However, at high cooling rates, the hardness of Nb steel even exceeds that of steel samples without Nb. In addition, adding a higher content of Nb at higher FRTs can further improve hardness.
The phase transition start and end temperatures of Steel 1 and Steel 2 are shown, and the microstructure related to the process route and cooling rate is illustrated. Steel 1 exhibits balanced phase transition behavior at different cooling rates, and martensite formation becomes more complete as the cooling rate increases. However, at lower cooling rates, the microstructure contains different proportions of bainite and martensitic segments.
Steel 2 behaves slightly differently, with a slight decrease in martensitic transition temperature across all routes. At high cooling rates, a complete martensitic microstructure can be obtained, while at lower cooling rates, γα transitions occur at higher temperatures. Metallographic analysis showed that steel 2 formed ferritic elements after re-austenitization, which transformed into full martensite after TMP.
Phase change behavior, corresponding hardness value and microstructure of steel 3 under different process routes. At medium cooling rates, phase change temperatures occur at simulated temperatures for all three routes. At a cooling rate of 400K/s, the phase change temperature (MS) is about 370°C, and drops to about MS~370°C at 30K/s.
However, at low cooling rate, the phase change temperature shifts to high temperature, such as MS of about 484°C when FRT is 875°C, MS~417°C when cooling rate is 1K/s, MS of about 486°C when FRT is 950°C, and MS~403°C when cooling rate is 100K/s. However, achieving a complete martensitic microstructure at a lower cooling rate requires a lower cooling rate (about 10 K/s) compared to steel 2.
The results show that increasing the cooling rate to more than 10K/s does not necessarily increase the hardness under the corresponding process route. At cooling rates below 10 K/s, partial bainite phases or even ferrite phases may occur.
Selected PAG (a to c) and deformation microstructure (d to f) of steel 3 at a cooling rate of 1K/s. Steel 1 machined at 875°C by FRT still shows a rough spherical PAG structure, and its corresponding transformation microstructure is much rougher than that of Steel 2.
Steel 2 and 3 are re-austenitized and PAG is very small. According to the PAG size data listed in Table III, the finest grains can be obtained by finishing at 875 °C, while steel 2 and 3 show a tendency to form the finest grains at 930 °C and then austenitization.
By comparing the hardening behavior of the three steels, we highlight the importance of a combination of microalloying and thermomechanical treatment (TMP) in achieving high strength. Although the strength gain obtained by adding 0.04% Nb trace alloy to steel 2 is only slightly improved compared to the traditional hardening method (re-austenitization +3%), the use of 100K/s TMP in the following isothermal quenching (DQ) process can achieve a hardness increase of 9% (FRT 875°C) and 10% (FRT 950°C) compared to steel without Nb.
In addition, at lower cooling rates (< 30 K/s), the hardenability of steel 1 and steel 2 is significantly different. Steel 1 still has a martensitic structure at a cooling rate of 1K/s, while Steel 2 is mainly composed of ferrite and bainite with a cooling rate of 30K/s.
By comparing the hardening behavior of the three steels, we highlight the importance of a combination of microalloying and thermomechanical treatment (TMP) in achieving high strength.
Although the strength gain obtained by adding 0.04% Nb trace alloy to steel 2 is only slightly improved compared to the traditional hardening method (re-austenitization +3%), the use of 100K/s TMP in the following isothermal quenching (DQ) process can achieve a hardness increase of 9% (FRT 875°C) and 10% (FRT 950°C) compared to steel without Nb.
In addition, at lower cooling rates (< 30 K/s), the hardenability of steel 1 and steel 2 is significantly different. Steel 1 still has a martensitic structure at a cooling rate of 1K/s, while Steel 2 is mainly composed of ferrite and bainite with a cooling rate of 30K/s.
After FRT processing at 875°C, the decrease in hardness of the steel may be due to the assumption that Nb is completely precipitated, while in other treatment routes, Nb may still be present in solution, so the formation of ferrite is delayed in the dissolved state according to the performance of the hardness value.
However, the size of the NbC precipitate observed at FRT at 875 °C was too large to achieve an adequate precipitation hardening effect.
In contrast, when the FRT is 950°C and 930°C, the NbC particles in the reaustenitized steel have a smaller size and the ability to achieve high hardness at a low cooling rate of 1K/s.
Microalloying combined with TMP is important in achieving ultra-high strength. By adding trace amounts of the element Nb (steel 2), a certain degree of strength gain can be obtained. Compared to conventional hardening after re-austenitization, steels with Nb addition provide significant hardness gains at a higher cooling rate (100 K/s) after TMP treatment.
At low cooling rates (< 30 K/s), the hardenability of steel 1 and steel 2 showed a significant difference. Steel 1 still has martensitic composition at lower cooling rates, while steel 2 is mainly composed of ferrite and bainite, which indicates that Nb microalloying and deformation promote the formation of bainite.
Steel 3 achieves higher strength by adding γ stabilizers such as Cr, Ni, Mo, and V, and behaves differently than Steel 2. At high temperatures, steel 3 has relatively good hardenability and no ferritic components. However, further research is needed to clarify whether Nb eliminates the impact of the shift from γ to α.
The size of the NbC precipitate plays an important role in the hardness of the material. Smaller NbC particles (after re-austenitization at 950°C and 930°C) help achieve higher hardness at low cooling rates. In contrast, larger NbC particles (at FRT at 875°C) cannot achieve sufficient precipitation hardening.