Effect of Titanium Addition on High Temperature Workability of High Manganese Austenitic Steel

Abstract

In the present study, the high temperature workability of high manganese austenitic steel has been examined to prevent the grain boundary embrittlement cracking problems in the continuous casting process. As-cast Fe-22Mn-0.4C steel exhibited poor hot ductility behaviors at 900°C tensile test. Phosphorus segregation and BN precipitation at grain boundaries were mainly responsible for this deterioration. In order to enhance the hot ductility, titanium was added to this steel, and high temperature workability was compared in view of reduction of area in tensile test at 900°C. BN precipitation at grain boundaries was effectively suppressed by the formation of interior Ti(C,N) precipitates. Furthermore, phosphorus atoms, a grain boundary embrittlement element, were observed to segregate at Ti(C,N) interfaces in Auger electron spectroscopy and atom probe tomography. These results show that titanium addition in Fe-22Mn-0.4C steel can effectively improve the high temperature workability by decreasing the segregation of phosphorus at grain boundary.

1. Introduction

There have been many reports on surface cracks along grain boundaries of steel slabs during continuous casting, and these cracks seriously damage the production process. In order to suppress surface cracks, numerous attempts have been carried out using hot tensile test to understand the role of grain boundaries in the hot ductility of steels.^{1,2,3,4,5,6,7,8,9,10,11)} C. L. Briant and S. K. Banerji⁴⁾ showed that boundaries with a second phase or elemental segregation caused grain boundary fracture at a certain temperature condition because they required the lowest fracture energy. The secondary phase found in most hot-ductility studies was a ferrite film formed along austenite grain boundaries.^2,4,5,6,7,8) On the other hand, it is generally accepted that formation of precipitates or the segregation of impurities at grain boundaries is a major cause of the high temperature embrittlement of austenite steels with no phase transformation.^9,12) For example, a recent work done by Peng Lan⁹⁾ showed the improved hot ductility and the fracture mechanism of fully austenite Fe–Mn–C steel in which the Mn and C micro-segregation ratio was lowered. However, there are not many studies on other microstructural effects compared with the effects of precipitates. In particular, direct observations on phosphorus segregation at grain boundaries in austenite steel are very limited.

High manganese TWIP steel is recognized for its value in many industries due to its non-magnetic characteristic and outstanding balance of strength and elongation. However, this steel is known to have the high temperature embrittlement. Detailed reasons of this embrittlement have not been clearly understood. In particular, the formation of ferrite films, a widely accepted mechanism for common steels, cannot be employed here because the TWIP steel has no phase transition during continuous casting process.

In order to avoid the high temperature embrittlement in steels, the content of elements causing the precipitation or segregation should be controlled to a specific value or less. However, some elements are practically difficult to control in this regard. In this case, other alternative methods have been used, such as thermomechanical treatments for suppressing the formation of second phase and adding grain boundary cohesion strengthening elements. For example, phosphorus is one of the typical metalloid impurities causing intergranular brittle fracture in continuous casting because the grain boundary segregation of phosphorus reduces the cohesion strength between grains.^{15,16,17,18,19)} T. Ogura et al.¹²⁾ found that the amount of phosphorus segregation was larger at incoherent boundaries in Ni–Cr steel. On the other hand, S-H Song²¹⁾ showed the hot ductility was improved by the grain boundary segregation of boron and sulphur in 2.25Cr-1Mo steel. In addition, Suzuki²²⁾ insisted the phosphorus segregation could be suppressed by boron segregation at austenite grain boundaries in low-alloyed steel. Therefore, boron is generally considered to strengthen the grain boundary cohesion by providing covalent bonding with matrix atoms.

In high manganese TWIP steel, nitrogen is generally added to improve the mechanical and corrosion properties. As mentioned previously, boron is added to increase the cohesion strength.^20,21,22) However, boron has a strong tendency to form BN precipitates with the presence of nitrogen at high temperature. These precipitates can deteriorate the high temperature workability of the TWIP steel. Thus it is necessary to analyze the microstructures and mechanical properties of boron and nitrogen added TWIP steel. Based on this analysis, it is possible to develop a more effective method to improve the high temperature workability.

The addition of titanium is known to improve the high temperature workability of TWIP steel. This improvement should be strongly related to the formation of stable Ti(C,N) precipitates at high temperature. However, the effect of Ti(C,N) precipitates on BN precipitation and phosphorus segregation at grain boundaries has not been deeply investigated. Therefore, commercial TWIP steels with phosphorus, boron and nitrogen elements were used in this experiment to observe the effect of titanium on the workability. Changes in microstructures and segregation concentration at boundaries were investigated to understand the effect of titanium addition on the high temperature workability.

2. Experimental

Fe–Mn–C TWIP steels were produced by vacuum induction melting and subsequent air cooling. The chemical compositions of these steels are shown in Table 1. In this study, the steels are divided into two groups. Compared with group 2, the steels of group 1 included slightly higher concentrations of additive elements, such as nitrogen, phosphorus, and sulfur. In particular, the nitrogen contents of group 1 and group 2 were adjusted to 0.03 wt% and less than 0.01 wt%, respectively. In each group, specimen A refers to a steel without titanium and specimen B refers to the steel containing titanium.

Table 1. Chemical compositions (wt%) of steels.

Alloy	Mn	P	B	S	N	C	Ti
1-A	23	0.02	0.003	0.002	0.03	0.4	–
1-B	23	0.02	0.004	0.002	0.03	0.4	0.05
2-A	22	0.01	0.003	0.001	<0.01	0.4	–
2-B	22	0.01	0.003	0.001	<0.01	0.4	0.05

For tensile test, cylindrical specimens were cut from dendritic grain zones. The dimensions of the specimens with a thread length of 10 mm at both ends were 120 mm in total length and 10 mm in diameter. The high temperature tensile tests were carried out with a strain rate of 1 s⁻¹ using a Gleeble Thermo-mechanical simulator. The thermal cycle employed to investigate the high temperature workability is shown in Fig. 1.

Fig. 1.

Schematic thermomechanical process diagram.

The microstructures and the fracture morphologies of TWIP steels were observed by scanning electron microscopy (SEM, TESCAN MIRA3) and optical microscopy (OM, Olympus GX 51). Samples for microstructural analysis were polished and etched by a 5 g of picric acid in 100 ml of water containing 1 g of sodium dodecyl benzene sulfonate added as a wetting agent. Samples for OM microstructural analysis were color etched by with a solution of 20 g sodium metabisulfite dissolved in 100 ml of water. The average length of the austenite grain boundary in given area was obtained from optical microscopy images on randomly sectioned planes. The length was measured with image analyzer. As many as 20 fields were measured at ×50 magnification. The average length of the Ti(C,N) interface in given area was obtained from scanning electron microscopy images. As many as 50 fields were measured at ×1 k magnification.

In-situ Auger electron spectroscopy (AES) was employed to evaluate the concentration of phosphorus at grain boundaries. 16 mm×2 mm×2 mm AES specimens with a sharp notch were prepared. For the AES analysis, the specimens were mounted in an ultrahigh vacuum system operated at a base pressure of 1.3∙10⁻⁸ Pa. The specimens were maintained at the liquid nitrogen temperature in the AES chamber for 1 h and then fractured by impact to obtain intergranular fracture surfaces. For this study, an Auger electron microprobe PHI 700 (ULVAC-PHI INC) with a cylindrical energy analyzer (CMA) was used to analyze individual intergranular facets. An accelerating voltage of 5 kV and a target current of 10 nA were used. The amplitude of an Auger peak in a differentiated spectrum is considered to be proportional to the atomic concentration of an element producing the peak. Impurity element segregation can be expressed in terms of the Auger peak height ratio of phosphorus (P123) to iron (Fe705) and manganese (Mn545). For identifying, In addition, an atom probe tomography (APT, Cameca LEAP 4000 X HR) was used to determine the element distribution at the precipitate boundaries of samples prepared by Focused Ion Beam (FIB, FEI Helios NanoLab G3 UC).

3. Results and Discussion

3.1. Changes in High Temperature Workability by Titanium Addition

The results of tensile tests at a temperature of 900°C at a speed of 1 s⁻¹ showed that the reduction of area (RA) of each specimen was 31.0% for 1-A, 75.0% for 1-B, 50.1% for 2-A and 75.1% for 2-B. In the same group with the similar phosphorus and nitrogen concentrations, titanium addition effectively increased the RA values. The increase in the RA values between specimens A and B is more evident in Group 1 than Group 2. In addition, it is shown that the high additive concentration decreased the RA values in titanium-free specimens.

Figure 2 shows the fracture surfaces of tensile specimens at 900°C. Specimen 1-A and 2-A have flat, smooth intergranular facets as shown in Figs. 2(a) and 2(c). On the other hand, the fracture morphologies of specimen 1-B and 2-B changed into ductile fracture with a high proportion of dimples, although some intergranular cracks are observed. Figure 3 shows the cross sectional morphologies perpendicular to the fracture surface. As shown in the figure, in undeformed regions, the addition of titanium does not make any noticeable change in the microstructures. However, in deformed regions, grain boundary cracks are clearly observed in specimens 1-A and 2-A as shown in Figs. 3(a), 3(c). On the contrary, in specimen 1-B and 2-B, grain boundary cracks are not found but grains are deformed in the loading direction. Furthermore, there exist small recrystallized grains indicating that there is no high temperature grain boundary embrittlement. These results clearly show that the grain boundary embrittlement occurred in the specimens without titanium and the RA values were dependent on the compositional changes. Thus, in order to understand these behaviors, it is critical to observe the microstructural changes caused by the titanium addition and other compositional changes.

Fig. 2.

Fracture morphologies of high manganese austenitic steels at 900°C. (a) 1-A (b) 1-B (c) 2-A (d) 2-B.

Fig. 3.

Grain morphologies of high manganese austenitic steels observed by OM. The tensile loading direction is vertical for all specimens. (a) 1-A (b) 1-B (c) 2-A (d) 2-B.

3.2. Changes in Second Phase Precipitation by Titanium Addition

The phase stability diagrams of secondary phases in high manganese austenite steels were constructed using thermodynamic software, FACTSAGE 5.5 program (Fig. 4). In Group 1, the phase fraction of BN precipitate in specimen 1-B is not reduced than that of specimen 1-A under the equilibrium state at 900°C, as shown in Figs. 4(a) and 4(b). Therefore, it is highly likely that all of boron atoms in specimen 1-B formed BN precipitates due to the presence of sufficient amounts of nitrogen atoms. This indicates that boron atoms, expected to suppress the grain boundary embrittlement, do not play any significant role in improving the hot ductility when BN precipitates are formed. On the other hand, in the specimens of the group 2 with reduced nitrogen content (less than 0.01 wt%), the weight percent of BN phase was sharply decreased by the addition of titanium, as shown in Figs. 4(c) and 4(d). This indicates that specimen 2-B had free boron atoms. In high temperature deformation, it is well-known that grain boundaries play a critical role in the formation and expansion of the rupture. The free boron atoms in specimen 2-B could be segregated into grain boundaries to increase the grain boundary strength. In addition, it is shown that titanium is a powerful nitride formation element and Ti(C,N) precipitates remain insoluble above 1250°C. This means that the high temperature workability should be highly dependent on the morphology and distribution of Ti(C,N) precipitates.

Fig. 4.

Phase stability diagram constructed by using FACTSAGE 5.5 program. (a) 1-A, (b) 1-B, (c) 2-A, (d) 2-B.

Figure 5 shows fine pores formed in specimen 1-A. As the deformation progresses, the pores expanded and led to a break at the grain boundary. When second phases exist on the boundary, the grain boundary pores could be more easily to form due to relatively low boundary cohesion strength. The microstructures near the fracture surfaces were investigated to determine whether the titanium addition had any effect on the distribution of grain boundary precipitates (Fig. 6). There are large precipitates on the fracture boundary in specimen 1-A. The EDS observation verifies that these precipitates have a BN phase. On the other hand, as shown in Fig. 2(b), specimen 1-B has a typical ductile fracture surface and BN precipitates are not observed on the fracture surface. Figure 6(b) shows that intra granular cubic Ti(C,N) precipitates with average size of 2.3 μm instead of BN precipitates are formed in the austenite grains of specimen 1-B. This indicates that BN precipitates in specimen 1-B could not play a significant role in fracture although they might exist at grain boundaries based on the thermodynamic calculation.

Fig. 5.

Formation of fine pores in the vicinity of grain boundary fracture in 1-A.

Fig. 6.

SEM micrograph of precipitates (a) BN at the fracture surface of specimen 1-A (b) Ti(C,N) inside the austenite grain of specimen 1-B.

The above morphological difference between specimen 1-A and 1-B should stem from the fact that the thermal stability of Ti(C,N) is higher than that of BN. As shown in Fig. 4. the high temperature annealing at 1200°C causes a small change in the weight fraction of BN precipitates in specimen 1-A. In addition, the high diffusion rate of boron in austenite, similar to that of nitrogen, can enhance the Ostwald ripening of BN precipitates at the grain boundaries, as shown in Fig. 6(a). On the other hand, in specimen 1-B, the addition of titanium lowers the dissolution temperature of BN precipitates from 1310 to 1060°C and the high onset temperature and thermal stability of the Ti(C,N) precipitates causes them to be finely dispersed inside grains, as shown in Fig. 6(b). Thus boron atoms should be redissolved during the high temperature annealing and the formation of BN precipitates be suppressed during the subsequent treatment, as shown in Fig. 6. As mentioned earlier, grain boundary precipitates promote the formation and expansion of voids when grain boundary sliding occurs. Reduction of these grain boundary precipitates and free boron atoms may have contributed to the increase in the high temperature workability.

The above discussion shows that BN precipitates at grain boundaries can play an important role in high temperature workability. However, there remains unanswered questions that the RA of specimen 1-A is much less than that of specimen 2-A despite the same equilibrium amount of BN precipitates and the RA of specimen 1-B is similar to that of specimen 2-B despite the large equilibrium amount of BN precipitates. Accordingly, the structural analysis of another influence of titanium addition on grain boundaries is required to clarify these behaviors.

3.3. Changes in Phosphorus Segregation by Titanium Addition

It is well-known that the grain boundary segregation of phosphorus atoms reduces the cohesion strength between grains. Thus the increase in the high temperature workability due to the addition of titanium should be affected by the formation of new Ti(C,N) interfaces that can be sites for phosphorus segregation. In the AES analysis, phosphorus is observed on the fracture surface of specimens 1-A, 1-B and 2-A (Fig. 7). Since the concentration of phosphorus is proportional to the peak height,^13,14) the relationship between the phosphorus content and the peak height can be expressed in the following manner.¹⁴⁾   

$${\text{X}}_{\text{P}}=2t\frac{\lambda cos\theta }{\text{a}}\frac{{\text{H}}_{\text{P}}}{{\text{H}}_{\text{Fe}}+{\text{kH}}_{\text{Mn}}}$$(1)

  

$$\text{k}=\frac{{\text{H}}_{\text{Fe}}^0}{{\text{H}}_{\text{Mn}}^0}=1.02$$(2)

Where λ is the attenuation length of the Auger electrons emitted by the phosphorus atoms, θ is the angle of emission from the surface normal, a is the thickness of the segregated layer. HP, HFe, HMn are the peak height at 123, 705, 545 eV, respectively. t is a spectrometer dependent constant. ${\mathrm{H}}_{\text{F}e}^{\text{0}}$ is the iron peak height for pure iron at 705 eV and ${\mathrm{H}}_{\mathrm{M}\mathrm{n}}^0$ is the manganese peak height at 545 eV for pure manganese.²³⁾ The factor 2 appearing in equation account that phosphorus is randomly distributed on both side of the intergranular fracture.¹⁴⁾ On the basis of specimen α, the relative ratio of segregated phosphorus $\left({\mathrm{X}}_{\mathrm{P}}^{\beta }/{\mathrm{X}}_{\mathrm{P}}^{\alpha }\right)$ of specimens can also be calculated in the following manner.   

$$\frac{{\text{X}}_{\text{P}}^{\beta }}{{\text{X}}_{\text{P}}^{\alpha }}=\left(\frac{{\text{H}}_{\text{P}}^{\beta }}{{\text{H}}_{\text{Fe}}^{\beta }+{\text{kH}}_{\text{Mn}}^{\beta }}\right)/\left(\frac{{\text{H}}_{\text{P}}^{\alpha }}{{\text{H}}_{\text{Fe}}^{\alpha }+{\text{kH}}_{\text{Mn}}^{\alpha }}\right)$$(3)

The relative ratio of segregated phosphorus on the basis of specimen 1-A, $\frac{{\mathrm{X}}_{\mathrm{P}}^{\beta }}{{\mathrm{X}}_{\mathrm{P}}^{1\mathrm{A}}}$ is shown in Fig. 8 with the RA values. Comparing the specimens without titanium atoms, the segregation fraction of phosphorus in specimen 2-A is less than that of specimen 1-A, which is due to that the phosphorus concentration of Group 2 was reduced to the half. As expected, the RA value of specimen 2-A is much higher than that of specimen 1-A. Thus, it is possible to conclude that the degree of grain boundary embrittlement is dependent on the phosphorus concentration at grain boundaries in specimens without titanium atoms.

Fig. 7.

Auger spectra from (a) intergranular fracture of specimens and (b) TiN interface of specimen 1-B.

Fig. 8.

Relative P ratio and RA values of specimens.

In Group 1, both specimens show phosphorus segregation at grain boundaries, but the segregation amount is sharply reduced in specimen 1-B, as shown in the Fig. 8. The RA value of specimen 1-B is increased 2.5 times compared with specimen 1-A. These results suggest that the addition of titanium effectively suppresses phosphorus segregation at grain boundaries and then enhances the high temperature workability. As illustrated in Fig. 6(b), a large number of Ti(C,N) precipitates were formed inside grains in specimen 1-B. These new interfaces of Ti(C, N) precipitates in specimen 1-B could be segregation sites for phosphorus. Figure 7(b) shows the AES spectrum obtained from the Ti(C,N) precipitate. For identifying the element distribution, an APT was also carried out and is presented in Fig. 9. These results prove that phosphorus atoms were segregated at the Ti(C,N) interface. Furthermore, free boron atoms existed and were also segregated at the Ti(C,N) interface. Figure 10 shows the concentration profiles of phosphorous in the vicinity of two different boundaries in specimen 2-B. These profiles were obtained by atom probe tomography (APT). The peak atomic percent concentration of phosphorus is 2 at% at the Ti(C,N) interface, whereas it is 0.3 at% in austenite grain boundary. This clearly shows that phosphorus atoms are preferable to be segregated at the Ti(C,N) interfaces rather than at the austenite grain boundaries. Meanwhile, the extent of two-dimensional internal boundaries of a solid can be calculated from measurements of the boundary length on plane sections.²⁴⁾ The length of internal boundaries (austenite grain boundary and Ti(C,N) interface) in given area was directly measured with microscopy using image analyzer and the average line length in square millimeter areas was 5 mm for Ti(C,N) interfaces, and 2.5 mm for austenite grain boundaries. The calculated area of Ti(C,N) interfaces is twice the surface area of the austenite grain boundaries. This should indicate that the amount of grain boundary segregated phosphorous atoms in specimen 2-B is less than the amount in specimen 2-A. Thus the high temperature workability of the specimen should be improved due to the phosphorous segregation at the Ti(C,N) interfaces.

Fig. 9.

APT images at the interface between TiN and austenite matrix in specimen 2-B.

Fig. 10.

APT concentration profiles for phosphorus in specimen 2-B (a) at the Ti(C,N) interface, (b) at the austenite grain boundary.

In specimen 2-B, the segregation of phosphorus at grain boundary is not observed in AES analysis. This relatively low concentration should be due to a decrease in phosphorus concentration and an increase in interface areas by Ti(C,N) precipitates. Furthermore, as shown in Fig. 4(d), the amount of BN precipitates is also reduced, indicating that free boron atoms could segregate to grain boundaries. These results clearly indicate that the high temperature workability of specimen 2-B must be higher than that of specimen 1-B. However, the RA values of specimen 1-B and 2-B are identical. Thus there may be a limit of phosphorus concentration at grain boundaries that determines the high temperature workability in this experimental condition.

As shown in Figs. 3(b) and 3(d), recrystallized grains near the fracture surfaces are well developed in specimen 1-B and 2-B. These recrystallized grains resulted from highly deformed grains and the phosphorus concentration at these boundaries should be less than that of initial grain boundaries. Thus the RA values of specimen 1-B and 2-B were determined by the austenite grain strength, not the grain boundary strength that depends on the phosphorus concentration. It was reported that the continuous casting process can be processed without a surface cracking if the RA value of steel is 60% or higher.⁷⁾ Accordingly, the specimens that lowered phosphorus segregation by the formation of Ti(C,N) precipitates are able to be deformed by the continuous casting without surface cracks.

From the results of analyzing the relationship between the high temperature workability and the microstructures of high manganese steel, it was found that the segregation of phosphorus atoms at austenite grain boundaries played an important role in deteriorating the high temperature workability. The addition of titanium of the steel improved the high temperature workability because the interfaces of Ti(C,N) precipitates acted as additional segregation sites of phosphorus atoms.

4. Conclusion

This study showed that high temperature workability was dramatically improved by adding titanium elements in Fe-22 Mn-0.4 C steel. In order to understand this behavior, the microstructural analysis was carried out and following conclusions can be drawn.

(1) When only boron elements (0.003 wt%) were added, the fracture morphology of specimens showed intergranular facets. At these facets, BN precipitates were formed and phosphorus elements were also segregated.

(2) In only boron added specimens, the high temperature workability was enhanced by reducing the amount of phosphorus concentration. This showed that the phosphorus segregation at grain boundaries played an important role in high temperature embrittlement.

(3) When additional titanium elements (0.05 wt%) were added, the RA values of specimens were sharply increased and the fracture mode changed into ductile fracture. In these specimens, fine Ti(C,N) precipitates were formed inside grains and phosphorus elements were segregated at the interface of Ti(C,N) resulting in the reduction of phosphorus concentration at the austenite grain boundaries.

(4) The above results proved that the addition of titanium in high manganese steel is an effective way to prevent surface cracks during continuous casting.

Acknowledgement

This work was performed under the support from POSCO and the Human Resources Development program (No. 20154030200680) of the Korea Institute of Energy Technology Evaluation and Planning (KETEP) grant funded by the Korea government Ministry of Trade, Industry and Energy.

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