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Influence of Ti on the Hot Ductility of High-manganese Austenitic Steels

  • Hongbo Liu , Jianhua Liu EMAIL logo , Bowei Wu , Xiaofeng Su , Shiqi Li and Hao Ding
Published/Copyright: September 20, 2016
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High Temperature Materials and Processes
From the journal High Temperature Materials and Processes Volume 36 Issue 7

Abstract

The influence of Ti addition (~0.10 wt%) on hot ductility of as-cast high-manganese austenitic steels has been examined over the temperature range 650–1,250 °C under a constant strain rate of 10−3 s−1 using Gleeble3500 thermal simulation testing machine. The fracture surfaces and particles precipitated at different tensile temperatures were characterized by means of scanning electron microscope and X-ray energy dispersive spectrometry (SEM–EDS). Hot ductility as a function of reduction curves shows that adding 0.10 wt% Ti made the ductility worse in the almost entire range of testing temperatures. The phases’ equilibrium diagrams of precipitates in Ti-bearing high-Mn austenitic steel were calculated by the Thermo-Calc software. The calculation result shows that 0.1 wt% Ti addition would cause Ti(C,N) precipitated at 1,499 °C, which is higher than the liquidus temperature of high-Mn austenitic steel. It indicated that Ti(C,N) particles start forming in the liquid high-Mn austenitic steel. The SEM–EDS results show that Ti(C,N) and TiC particles could be found along the austenite grain boundaries or at triple junction, and they would accelerate the extension of the cracks along the grain boundaries.

Keywords: high-Mn austenitic steels; Ti-microalloying elements; hot ductility; grain boundary sliding

Introduction

High-manganese austenitic steel, which is also called Hadfield steel [1], is well known for its unique combination of an extraordinary high strain hardening potential with excellent toughness properties. The chemical composition of high-Mn austenitic steel is specified with 1–1.4 wt% C, 10–14 wt% Mn, 0.3–0.6 wt% Si, and balanced Fe [2, 3]. Xu et al. [4] reported that the wear resistance of 0.1 wt% Ti and 0.41 wt% V microalloyed high-Mn austenitic steel would be increased, because the dispersive carbonitride could be obtained through precipitation heat treatment. Su et al. [5] suggested that the wear extent of high-Mn austenitic steel is ~60.6 mg without Ti and ~32.1 mg with 0.109 wt% Ti, which indicated that the wear resistance increased nearly 100 % alloying with Ti. Hu et al. [6] reported that alloying 0.06–0.12 wt% Ti can improve the tensile strength and impact toughness of high-Mn austenitic steel.

However, Ti additions generally give rise to very poor ductility [7, 8, 9, 10]. In the current work, the hot ductility of as-cast non- and Ti-microalloyed high-Mn austenitic steels has been examined over the temperature range 650–1,250 °C under a constant strain rate of 10−3 s−1 using Gleeble 3500 thermal simulation testing machine. Additionally, the equilibrium precipitated phases in Ti-bearing high-Mn austenitic steel in the temperature range of 700–1,600 °C were calculated by using the thermodynamics software Thermo-Calc.

Experimental and methodology

Two high-Mn austenitic steels, i.e., one not microalloyed (steel 1) and other single microalloyed with Ti (steel 2), were evaluated. Their composition is specified in Table 1. Steel 1 has a composition which combines 1.10 % C and 0.49 % Si in addition to 12.72 % Mn. However, it contains no microalloying additions. On the other hand, steel 2 has a similar main composition but it is further microalloyed with 0.1 % Ti. The steels were melted in a vacuum induction furnace in Engineering Research Institute, Beijing, China. It is worth noting that the molten steel was poured into the mold over a period of 3 min after deoxidation with aluminum. The contents of C and S were measured by infrared absorptiometric method after combustion, and Mn concentration was measured through perchloric acid oxidation trivalent manganese titrimetric method. The insert gas fusion-infrared absorptiometry was used to determine O and N contents in steel based on ASTME11019-2011 ASTM standard. The concentrations of other elements were determined by ICP-OES (Inductively coupled plasma optical emission spectrometric) method.

Table 1:

Composition of steels examined (wt%).

SteelCSiMnPSAlTiNO
11.100.4912.720.00540.00570.03–0.00900.0069
21.190.4512.290.00520.00590.050.100.01400.0073

Cylindrical tensile specimens of 120 mm in length and 10.0 mm in diameter were machined from the ingots. A Gleeble3500 machine was used to carry out the tests. The hot ductility experiments were carried out in argon atmosphere to protect the samples from oxidation. Meanwhile, the samples were heated at 10 °C/s to 1,200 °C, where they were held for 3 min. They were then cooled or heated at 3 °C/s to each testing temperatures with an interval of 50 °C in the range 650–1,250 °C as schematically shown in Figure 1. Once the sample had reached the test temperature, it was held for 3 min before straining to failure using a strain rate of 1 × 10−3 s−1. It should be noted that two samples were tested at each temperature in order to get an accurate hot tensile performance of steel. The experiment results would be accepted if the relative difference of these two values was less than 0.1. Once a tensile sample had failed, it was immediately gas quenched with argon to preserve the grain structure and the fracture surface present at the time of failure. After measurement of the reduction of area (RA), a sample taken close to the point of failure was prepared for metallographic examination.

Figure 1: Schematic diagram for processing of the tensile specimens.
Figure 1:

Schematic diagram for processing of the tensile specimens.

To investigate the microstructure, the specimen was etched in a mixture of 5 % nitric acid alcohol and 1.5 % saturation picric acid for about 120 s. The fractured surfaces were examined by SEM and the chemical nature of particles was determined by EDS analysis using the Zeiss Utra55 microscope operated with an accelerating 15 kV.

In the present study, equilibrium precipitated phases in Ti containing high-Mn austenitic steel in the temperature of 700–1,600 °C and the influence of alloying elements on the amount of Ti(C,N) precipitation was calculated by using the Thermo-Calc 5.0 Software. TCFE7 database was adopted in the calculation, and the amount of substance of alloy system was set as 1 mol. The temperature of reference state was 298.15 K, and the pressure was 105 Pa.

Results and discussion

Tensile behavior

Figure 2: True stress-true strain curves as a function of temperature (a) steel 1, (b) steel 2.
Figure 2:

True stress-true strain curves as a function of temperature (a) steel 1, (b) steel 2.

The results indicate that strength decreases as the temperature increases. All the high-Mn austenitic steels tested at 700 °C exhibit a single peak stress, the value of which is close to 182 MPa for the steel 1. This single peak can be associated to the necking phenomenon [11]. The flow curves in the range of 800–1,200 °C of two high-Mn austenitic steels include a short work hardening region up to the ultimate tensile strength followed by a long post-ultimate tensile strength region. The lower rate of the stress drop can be directly associated to the occurrence of dynamic recrystallization (DRX) [12]. As can be seen in Figure 2, DRX occurs at 800 °C during tensile test of steel 1, while 900 °C of steel 2. This indicated that the addition of 0.10 % Ti could retard the DRX, and thus deteriorating the hot ductility of high-Mn austenitic steel.

Figure 3: Peak stress as a function of the temperature of high-Mn austenitic steels.
Figure 3:

Peak stress as a function of the temperature of high-Mn austenitic steels.

A lot of peak stress versus temperature is given in Figure 3. It can be seen that the peak stress decreases as the temperature increases for two kinds of high-Mn austenitic steels. For example, the values of peak stress for steel 2 are close to 294 MPa at 650 °C and decrease to values close to 10 MPa at 1,200 °C. In addition, there is a strong effect of Ti on the maximum high-Mn steel strength at the lower temperatures. That may be related to the solid solution or second-phase particles strengthening [11]. The peak stress tend to have the same value for two high-Mn steel, which indicated the weaker effect of solid solution and second-phase particles strengthening. Diffusion is mainly controlled by the climb of dislocations at the higher temperatures.

Hot ductility

Figure 4: Hot ductility curves as a function of temperature of high-Mn austenitic steels.
Figure 4:

Hot ductility curves as a function of temperature of high-Mn austenitic steels.

The hot ductility curves showing RA as a function of the tensile test temperature for the two studied as-cast high-Mn austenitic steels are given in Figure 4. In steel 1, the RA value is 62 % at the lowest tensile testing temperature and shows a little change during 650–850 °C. The ductility improves at higher temperature reaching a maximum value of 80 % at 900 °C. Then, the RA values drop until reaching the minimum value of 29 % at 1,250 °C. The steel 2 presents the worse RA values in the almost entire range of testing temperatures. Moreover, Mintz [7] reported that the RA value which can be used to assess the susceptibility to transverse cracking of steel is 60 %. The crack sensitivity would be greatly increased when RA value is less than 40 % [7, 13]. Thus, the high-temperature brittle zone for steel 1 and steel 2 is in the range from the melting point to 1,250  and 1,200 °C, respectively.

Fractography

Fracture surfaces of the studied high-Mn austenitic steels tested at different temperature are shown in Figure 4. For steel 1, the character of the fracture is ductile at 700–900 °C as can be seen from the existence of dimple having various sizes, as shown in Figure 5(a)–(c). Crowther and Mintz [14] reported that the large voids were originally intergranular cracks, which were formed at the early stage of deformation due to the grain boundary sliding. And then, the cracks would get distorted into elongated voids until final failure, which occurs by necking between these voids.

Figure 5: SEM images from the fracture surface.
Figure 5:

SEM images from the fracture surface.

Figure 5(a)–(c) shows the shallow dimples with some flat grain surface, which indicated the characteristics of failure due to grain boundary sliding. At 900 °C, the fracture surface presents bigger dimples, which shows that steel 1 has a better ductile than that tested at other temperatures. Carpenter [15, 16] suggested that the fracture tested in single phase austenite region exhibited the intergranular decohesion, displaying flat, and featureless facets, indicating failure due to grain boundary sliding. An example of such a failure is shown in Figure 5(d). The fracture surface displays as flat grain surface. It indicated that intergranular cracking as a result of grain boundary sliding is the main fracture mechanism during hot tensile tests for non-microalloyed high-Mn austenitic steel. For steel 2, the fracture surface changes showing the flat smooth facets at almost the entire range of testing temperature, where intergranular-type fracture is obtained. This phenomenon shows that 0.10 % Ti addition seriously deteriorates the hot ductility of high-Mn austenitic steel.

Phases’ equilibrium diagrams and analysis of particles

Generally, a small addition of Ti, 0.02 %, could result in a small improvement of the hot ductility after the slow cooling [17]. Mintz [17] also found that 0.02 % Ti could slightly improve the hot ductility of the as-cast 0.05/0.15 %C–Mn–Al steels when the cooling rate from 1,540 °C was 60 K/min. Banks and Mintz [18] examined the influence of a low Ti addition, ~0.01 %, on the ductility of Nb-containing steels and found that the hot ductility decreased sharply with the addition of Ti for conventional cooling conditions. The tensile tests in the present study adopted a relatively fast cooling rate, ~3 °C/s, to test the temperature. In addition, the addition of Ti is higher, ~0.10 %. As a consequence, steel 2 presents the worse RA value in almost the entire range of testing temperature.

Figure 6: Phase transformation behavior and equilibrium phases in steel 2 predicted by Thermo-Calc.
Figure 6:

Phase transformation behavior and equilibrium phases in steel 2 predicted by Thermo-Calc.

Figure 6(a) shows the phases’ equilibrium diagrams in steel 2 calculated by Thermo-Calc commercial software. It should be noted that P, S, and O elements were not considered during the Thermo-Calc calculation, because the content of these elements is relatively low (P ~0.0052 %, S ~0.0059 %, and O ~0.0073 %).

According to the thermodynamic calculation, the main phases are the austenite solid face-centred cubic (FCC) phase and the Ti(C,N) compound. Ti(C,N) starts forming at 1,499 °C, which is much higher than its liquidus temperature, Tliquidus = 1,386 °C (calculated by Thermo-Calc). It indicated that Ti(C,N) starts forming in the liquid high-Mn steel, and the mole fraction of this phase would increase as solidification proceeds. The final mole fraction of Ti(C,N) is 2.22e−3. It can be seen from Figure 6(b) that the mass that fraction of Ti, C, and N is 77.7 %, 3.1 %, and 19.2 % at 1,499 °C, respectively. As the temperature decreased, the N content decreased gradually, while the C content increased.

SEM characterization of steel 2 taken from sectioned tensile specimens close to the fracture tip showed the particles at the austenitic grain boundaries. The SEM results of the sample tested at 800 °C give us the important information about nature of particles precipitated in high-Mn steel (Figure 7(a)–(e)). As can be seen from Figure 7(a), plenty of cracks and voids were found along the austenitic grain boundaries or located at triple junction. In an enlarged view of the cracks and voids indicated by Figure 7(b)–(e), some polygonal particles with a size of 3–8 μm were situated at the austenite grain boundaries. The number of analysis spectra of the particles could be divided in to two types: TiC, Ti(C,N). Charleux [19] reported that the nucleation of TiC or Ti(C,N) in austenite requires heterogeneous nucleation site such as grain boundaries, dislocations, or second-phase particles. According to the results of Thermo-Calc calculation, Ti(C,N) particles start forming in the liquid steel (Figure 6(a)), and they would grow to be large-size particles during the cooling process. As a consequence, these particles weaken the grains cohesion and accelerate the crack growth.

Figure 7: Microstructures and Ti(C,N) particles present in steel 2: tensile specimens were given test cycle shown in Figure 1 and tested at 800 °C.
Figure 7:

Microstructures and Ti(C,N) particles present in steel 2: tensile specimens were given test cycle shown in Figure 1 and tested at 800 °C.

Influence of alloying elements on the amount of Ti(C,N) precipitation

The steel 2, 1.19 %C–12.29 %Mn–0.45 %Si–0.05 %Al–0.1 %Ti–0.014 %Ti, was used as the base steel. The detailed studies on the influence of alloying elements on the precipitation of Ti(C,N) particles were carried out using Thermo-Calc Software, and the calculation results are shown in Figure 8.

Figure 8: Effect of Ti, C, and N contents on Ti(C,N) phases.
Figure 8:

Effect of Ti, C, and N contents on Ti(C,N) phases.

It can be seen from Figure 8 that the amount of Ti(C,N) is mainly influenced by Ti element, and the precipitation temperature is influenced by Ti and N elements. As the mass fraction of Ti is increased from 0.005 % to 0.1 %, the precipitation temperature of Ti(C,N) is 1,499, 1,457, 1,389, and 1,300 °C, respectively. Moreover, the final mole fraction of Ti(C,N) increases substantially from 1.16e−4 to 2.2e−3. C element has a little influence on both the amount and the precipitation of Ti(C,N). This is because the high C content in steel made the precipitation of solid solution element, Ti, was almost complete at a certain temperature. The precipitation temperature decreases from 1,499 to 1,388 °C with N content decreases from 0.014 % to 0.002 %.

Conclusions

  1. A unit of 0.1 % Ti addition to the as-cast high-Mn austenitic steel is deleteriated, and ductility would be worse than that obtained with a Ti-free steel of similar composition. Ti-containing steel presents the worse RA values in the almost entire range of testing temperatures. The high-temperature brittle zone for Ti-free and Ti-bearing high-Mn austenitic steels are in the range from the melting point to 1,250 °C and 1,200 °C using a strain rate of 1 × 10−3 s−1, respectively.

  2. The studied high-Mn austenitic steel microalloyed with 0.1 % Ti exhibits higher peak stress value than the Ti-free high-Mn austenitic steel, which is associated with solid-solution and precipitation hardening effects. The values of peak stress for Ti-free and Ti-bearing high-Mn austenitic steel are close to 424 and 293 MPa at 650 °C, respectively.

  3. Ti-free high-Mn austenitic steel shows fracture surface composed of many dimples, resulting in a predominantly ductile fracture, and while 0.1 % Ti-containing steel exhibits faceted fracture surfaces. The poor ductility in the Ti-bearing steel is more of a consequence of the Ti(C,N) particles which are situated at the grain boundary and triple junction. These Ti(C,N) particles weaken the grains cohesion and accelerate the crack growth.

  4. The results of Thermo-Calc calculation show that Ti(C,N) starts forming at 1,499 °C, which is much higher than its liquidus temperature. It indicated that Ti(C,N) starts forming in the liquid high-Mn steel.

  5. The precipitation temperature and the amount of Ti(C,N) are mainly influenced by Ti element. As the mass fraction of Ti is increased from 0.005 % to 0.1 %, the precipitation temperature of Ti(C,N) is increased from 1,300 to 1,499 °C, and the final mole fraction of Ti(C,N) increases substantially from 1.16e−4 to 2.2e−3. In order to reduce the incidence of transverse cracking, it is recommended that Ti content should not be too high in high-Mn austenitic steel.

Funding statement: This work was supported by the National Natural Science Foundation of China (grant nos. 51574022, 51374023).

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Received: 2016-1-6
Accepted: 2016-5-17
Published Online: 2016-9-20
Published in Print: 2017-7-26

© 2017 Walter de Gruyter GmbH, Berlin/Boston

This article is distributed under the terms of the Creative Commons Attribution Non-Commercial License, which permits unrestricted non-commercial use, distribution, and reproduction in any medium, provided the original work is properly cited.

Articles in the same Issue

  1. Frontmatter
  2. Research Articles
  3. Experimental Study on Application of Boron Mud Secondary Resource to Oxidized Pellets Production
  4. A Study at the Workability of Ultra-High Strength Steel Sheet by Processing Maps on the Basis of DMM
  5. Oxidation Behavior of TiAl-Based Alloy Modified by Double-Glow Plasma Surface Alloying with Cr–Mo
  6. Transient Liquid Phase Bonding of Nickel-Base Single Crystal Alloy with a Novel Ni-Cr-Co-Mo-W-Ta-Re-B Amorphous Interlayer
  7. Effects of Mn and Al on the Intragranular Acicular Ferrite Formation in Rare Earth Treated C–Mn Steel
  8. Effect of Plate Thickness on Tensile Property of Ti–6Al–4V Alloy Joint Friction Stir Welded Below β-Transus Temperature
  9. Characterization of High Temperature Deformation Behavior of BFe10-1-2 Cupronickel Alloy Using Orthogonal Analysis
  10. Influence of Ni Additions on the Viscosity of Liquid Al2Cu
  11. Corrosion Process of Stainless Steel 441 with Heated Steam at 1,000 °C
  12. Influence of Ti on the Hot Ductility of High-manganese Austenitic Steels
  13. Effect of Temperature Field on Formation of Friction Stir Welding Joints of Ti–6Al–4V Titanium Alloy
  14. Influence of Secondary Cooling Mode on Solidification Structure and Macro-segregation Behavior for High-carbon Continuous Casting Bloom
https://doi.org/10.1515/htmp-2016-0005
Keywords for this article
high-Mn austenitic steels; Ti-microalloying elements; hot ductility; grain boundary sliding
Creative Commons
BY-NC-ND 3.0

Articles in the same Issue

  1. Frontmatter
  2. Research Articles
  3. Experimental Study on Application of Boron Mud Secondary Resource to Oxidized Pellets Production
  4. A Study at the Workability of Ultra-High Strength Steel Sheet by Processing Maps on the Basis of DMM
  5. Oxidation Behavior of TiAl-Based Alloy Modified by Double-Glow Plasma Surface Alloying with Cr–Mo
  6. Transient Liquid Phase Bonding of Nickel-Base Single Crystal Alloy with a Novel Ni-Cr-Co-Mo-W-Ta-Re-B Amorphous Interlayer
  7. Effects of Mn and Al on the Intragranular Acicular Ferrite Formation in Rare Earth Treated C–Mn Steel
  8. Effect of Plate Thickness on Tensile Property of Ti–6Al–4V Alloy Joint Friction Stir Welded Below β-Transus Temperature
  9. Characterization of High Temperature Deformation Behavior of BFe10-1-2 Cupronickel Alloy Using Orthogonal Analysis
  10. Influence of Ni Additions on the Viscosity of Liquid Al2Cu
  11. Corrosion Process of Stainless Steel 441 with Heated Steam at 1,000 °C
  12. Influence of Ti on the Hot Ductility of High-manganese Austenitic Steels
  13. Effect of Temperature Field on Formation of Friction Stir Welding Joints of Ti–6Al–4V Titanium Alloy
  14. Influence of Secondary Cooling Mode on Solidification Structure and Macro-segregation Behavior for High-carbon Continuous Casting Bloom
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