1. Introduction
Over the years, different pressure melting techniques have been proposed to produce high nitrogen steels. Pressurized electroslag remelting (PESR) has emerged as the most robust method and attained increasing attention.1) Figure 1 shows the schematic sketch of PESR furnace. The main feature is the closed pressure vessel that allows adjusting the pressure and composition of gas. The PESR process not only inherits the advantages of convectional electroslag remelting process, such as the compact structure, homogeneous composition and controlled solidification, but also enables to introduce sufficiently high amount of nitrogen into the molten steel beyond the solubility limit at atmosphere pressure, and reduces the detrimental effects of oxygen or hydrogen.2,3) The products manufactured by PESR are widely used in aerospace, aviation, energy, petrochemical industry and so on.
Currently, most literatures concerning the PESR process focus on the chemical composition homogeneity and mechanical property. G. Stain et al.3,4) reported the production of nitrogen alloyed steels using PESR with nitrides as nitrogen source and discussed the effect of nitrogen on the property of materials. M. Bartosinki et al.5) compared effects of the types of nitrogen bearing additives on the chemical homogeneity of austenitic and martensitic stainless steels. Besides, they also studied the deoxidation during the production of titanium alloys using PESR.6) A. Patel et al.1) investigated the effect of pressure and types of nitrides on nitrogen distribution in the stainless steel. F. Takahashi et al.7) studied the effect of nitrogen content on deformability of Mn–Cr–N steels manufactured using a lab-scale PESR.
Pressure, as an important thermodynamic parameter, shows significant effects on chemical reaction involving gas phase reactant or product, such as increasing reaction rate and solubility of gas or volatile element.8,9) In addition, pressure also affects the solidification process.10) On the one hand, increasing pressure could decrease the heat resistance at ingot/mould interface and enhance the cooling rate of ingots, providing a fine solidification structure.11,12,13,14) On the other hand, increasing pressure could affect the thermodynamic and kinetic parameters, and change the solidification mode.15,16) Zhu et al.16) pointed that the transformation from FA to A mode may occur in the solidification of 19Cr14Mn4Mo0.8N stainless steel with the increase of pressure from 0.1 MPa to1 GPa.
The above researches about the influence of pressure on solidification process concentrate on the pressurized induction melting and squeeze casting process. However, it has not been reported that the effect of nitrogen gas pressure on as-cast microstructure revolution and solidification parameters (pool depth, mushy zone width, cooling rate and SDAS) during PESR process so far. In this work, three AISI 304 stainless steel electrodes were remelted under 0.1, 0.8, and 1.2 MPa using a lab-scale PESR furnace. The considerable attention was paid to the influence of gas pressure on pool depth, SDAS, mushy zone width and cooling rate. In addition, the microstructure revolution was studied with the optimal microscope (OM) and electron probe microanalyzer (EPMA).
3. Results
3.1. Pool Profile
Figure 3 demonstrates the pool profile under different gas pressure. The pool profile was shallow at 0.1 MPa, with a depth of 28 mm. The pool profile displayed a “V” shape at 0.8 MPa, with a depth of 34 mm. When the gas pressure increased up to 1.2 MPa, the pool depth decreased to 30 mm. At the gas pressure of 0.1, 0.8 and 1.2 MPa, the measured electrode melting rate was 42, 38 and 38 kg/h, respectively. In general, the pool depth linearly increases with the melting rate.20) But, in this study, the pool depth increased with the decreasing melting rate. This indicates that the pool depth is not only determined by the melting rate but also by the gas pressure. The detailed discussion is given in part 4.1.
3.2. Macrostructure of Ingots
Figure 4 shows the macrostructure of ingots under different gas pressure. These figures showed that the overall shape of grains were columnar in PESR-1, PESR-2 and PESR-3. Columnar grains vertically grew at the ingot bottom where the cooling from baseplate was dominant. With the increase of ingot height, columnar grains were well developed and grew continuously from the ingot surface to the center. Columnar grains grew parallel to the direction of heat flow.21) When the gas pressure was 0.1, 0.8 and 1.2 MPa, at the ingot center, the angle of grains with respect to the vertical axis was 36°, 45° and 40°, respectively. A larger angle implies a deeper metal pool. At the same power input, the angle firstly increased and then decreased with the gas pressure increasing, which was consistent with the tendency of pool depth.
3.3. Microstructure and EPMA Analysis
Figure 5 shows the microstructure of the samples. Lichtenegger and Bloesch etchant makes austenite phase show different colors depending on the Cr microsegregation. The color changes in the following order as the etching time increases: Yellow→Red→Violet→Blue→Green.18) The ferrite is not attacked and remains white. The etchant could clearly distinguish between primary dendrites and interdendritic regions.22) These microstructures consisted of austenitic matrix and residual ferrite under current pressure range. The ferrite included the skeletal ferrite and lathy ferrite. The skeletal ferrite was presented in the dendrite core, and the lathy ferrite was ranged with entangled or parallel form.23)
Figure 6 illustrates the distribution of SDAS under different pressure. The SDAS firstly increased from the ingot surface to the mid-radius and then decreased toward the center. In the outer region of ingot, gas pressure had a slight influence on SDAS. However, in the inner region of ingot, the SDAS showed an evident difference under different gas pressure conditions. When the pressure was 0.1, 0.8 and 1.2 MPa, the maximum SDAS was 67, 88 and 81 um, respectively. The maximum SDAS of ingot firstly increased and then decreased with the increase of pressure from 0.1 to 1.2 MPa, which displayed a similar tendency as the pool depth. In addition, with the increase of pool depth, the position of maximum SDAS moved toward the ingot center.
Figure 7 shows the composition profiles obtained from EPMA between the lathy ferrite for PESR-1 ingot. The residual ferrite enriched in Cr and depleted in Ni because Cr is ferrite stabilizer and Ni is austenitic stabilizer. The Cr content of the austenite was rather uniform at 17.4 pct and increased to 19.1 pct within 8 μm of the interphase boundary. The Ni content of the austenite increased continually from the interphase boundary and reached a maximum of 9.5 pct between the lathy ferrite. Figure 8 shows the composition profiles obtained from EPMA between the secondary dendrite arms for PESR-2 ingot. The Cr content of the austenite reached maximum of 19.4 pct near interphase boundary and decreased to a uniform measured value between the secondary dendrite arms. The Ni content of the austenite was a minimum near the interphase boundary and increased to 9.9 pct at the center. Figure 9 shows the composition profiles obtained from EPMA between the primary dendrite arms for PESR-3 ingot. The Cr and Ni content between the primary dendrite arms were similar to those between the secondary dendrite arms. The chemical composition of residual ferrite is listed in Table 3. At 0.8 MPa, the residual ferrite showed a most serious enrichment of Cr and depletion of Ni.
Table 3. The Cr and Ni content in residual ferrite on the basis of the EPMA analysis.
| No. | Cr, wt % | Ni wt % |
|---|
| PESR-1 | 25.7 | 2.1 |
| PESR-2 | 26.9 | 1.9 |
| PESR-3 | 26.2 | 2.0 |
4. Discussions
4.1. Effects of Nitrogen Gas Pressure on Solidification Parameters
The ingot structure is rather important in determining the properties of the material. To a great extent, the structure will be determined by the pool profile and depth.24) Under the current experiment conditions, when the gas pressure was 0.1, 0.8 and 1.2 MPa, the melting rate was 42, 38 and 38 kg/h, respectively. Generally, the metal pool depth linearly increases with an increase in melting rate.20) However, in this work, with the increase of pressure from 0.1 to 1.2 MPa, the melting rate had a slight decrease, but the pool depth firstly increased and then decreased. The detailed discussion is given in the following from the view of heat balance.
The heat source for metal pool includes two parts: (1) the heat of metallic droplets; (2) and the heat transferred by slag bath at the slag/pool interface. L. G. Hosaman et al.25) proposed that the size and shape of the metal pool is the result of the rate of heat input into metal pool, and the rate of heat extraction from the ingot. When the pool profile reaches the steady state, the heat balance in metal pool is given as:
|
Q
droplet
+
Q
c
=
Q
loss
| (1) |
Where
Qdroplet represents the rate of heat input into metal pool by droplets, J/s;
Qc is the rate of heat input into metal pool by convective heat transfer at the slag/pool interface, J/s;
Qloss is the rate of heat extraction from ingot, J/s.
With the increase of pressure from 0.1 to 0.8 MPa, the melting rate slightly decreased due to the increase of heat loss in slag bath, resulting in the reduction of Qdroplet. Meanwhile, the Qloss increases with an increase in pressure because increasing pressure could decrease the air gap.10,11,12,13,14) However, the pool depth increased from 28 to 34 mm. This indicated that there was also an increase in the heat transfer rate at the slag/pool interface. Moreover, the heat transfer rate at the slag/pool interface had a larger increase than that at the ingot/mould interface, because the pressure of 0.8 MPa was still small and the air gap only had a slight decrease. The increase of heat transfer rate at the slag/pool interface may be attributed to the increase of effective thermal conductivity of molten slag in PESR process. During the electroslag remelting process, the volatile gas such as AlF3, SiF4 and HF could be generated via the reaction [2]–[4], and then it is escaped from slag bath in the form of gas bubble. The existence of gas bubble in slag bath could decrease the effective thermal conductivity of molten slag.
|
3(Ca
F
2
)+(A
l
2
O
3
)=3(CaO)+2{Al
F
3
}
| (2) |
|
2(Ca
F
2
)+(Si
O
2
)=2(CaO)+{Si
F
4
}
| (3) |
|
(Ca
F
2
)+(
H
2
O)=(CaO)+2{HF}
| (4) |
Where { } indicates the gas phase, ( ) indicates the liquid phase.
The formation of gas bubble in molten slag consists of two steps: (1) nucleation of the gas bubble; (2) growth of the gas bubble. Considering the first of these steps, a gas bubble may be nucleated in a liquid slag when the pressure created by ex-solution of gas is greater than the sum of new gas bubble’s surface energy and liquid pressure:26,27)
Where
Pg is the inner pressure of gas bubble,
σ is the liquid/gas interfacial energy,
r is the radius of the nucleated gas bubble,
Pliq is the pressure in liquid phase.
During the PESR, gas pressure is loaded on the slag surface, resulting in the increase of Pliq. The nucleation of gas bubble could be suppressed by increasing the pressure, which is benefit to improve the effective thermal conductivity of molten slag.
When the pressure further increased up to 1.2 MPa, the melting rate remained the same with that at 0.8 MPa, which indicated that the heat transfer rate at the slag/pool interface had no further increase. However, the air gap had a further decrease due to the increase of pressure, resulting in the further increase of heat transfer rate at the ingot/mould interface. Finally, there was an apparent decrease in the metal pool depth. Based on the current experiment results, the remelting should be avoided to carry out at low pressure range from the view of controlling metal pool depth and SDAS.
The mushy zone width is given as:25)
|
h
l
-
h
s
=
V
c
⋅
(
λ
2
/k)
-n
| (6) |
Where h
l represents the height of liquidus, m; h
s is the height of solidus, m; V
c is the casting velocity, m/s;
λ2 is the SDAS,
μm; k and n are constant, for 304 stainless steel, k=68, n=0.45.
28)
Figure 10 displays the mushy zone width under different gas pressure conditions. With the increase of pressure from 0.1 to 0.8 MPa, the pool depth increased from 28 to 34 mm, and the mushy zone width also increased from 5.7 to 9.6 mm. When the pressure further increased to 1.2 MPa, the pool depth decreased to 30 mm, and the mushy zone width was 7.9 mm. Mitchell and Ballantyne found that a deeper metal pool is representive of a wider mushy zone.29) The calculated mushy zone width is in accordance with Mithchell’s results.
The local cooling rate of ingot is related to the SDAS:
Where
λ2 is the SDAS,
μm;
ε is the local cooling rate, K/s; k and n are constant, for 304 stainless steel, k=68, n=0.45.
28)
Figure 11 shows the calculated cooling rate using the measured SDAS. With the increase of pressure from 0.1 to 0.8 MPa, the cooling rate decreased from 1.04 to 0.56 K/s because the heat transfer rate at slag/pool interface had a larger increase than that at the ingot/mould interface. When the pressure further increased to 1.2 MPa, the heat transfer rate at the ingot/mould interface had a further increase, while the heat transfer rate at the slag/pool interface had no evident change, and hence the cooling rate increased to 0.68 K/s.
4.2. Effects of Nitrogen Gas Pressure on the Microstructure Evolution
Metallurgists have been very active in investigating the solidification mode of stainless steel, since these determine the castability, the hot workability and the room temperature structure.22) The solidification mode of austenite stainless steel could be classified into the following four types:
A mode: L→L+γ→γ
AF mode: L→L+γ→L+γ+δ→γ+δ→γ
FA mode: L→L+δ→L+γ+δ→γ+δ→γ
F mode: L→L+δ→δ→γ+δ→γ
Where L, δ and γ represent liquid, ferrite and austenite, respectively.
Generally, the solidification mode of austenite stainless steel is strongly dependent on the chemical composition and cooling rate. Under equilibrium solidification conditions, the AISI 304 stainless steel falls into FA mode. However, increasing the cooling rate may lead to the transformation from FA to AF.30) In addition, pressure also affects the solidification mode.16)
In order to explore the effect of nitrogen gas pressure on solidification mode during PESR process, it was identified by the combination of the color metallography and EPMA.
Ferrite present in the dendrite cores is the FA mode, whereas the interdendritic ferrite represents the AF mode.18) As seen from Fig. 5, the solidification mode is FA mode because the ferrite presents in the dendrite cores. However, both AF and FA mode contain some ferrite in the final microstructure and require some care when identifying the AF and FA mode.18) Hence, the solidification mode is further confirmed using the EPMA. The Ni concentration profile was examined with EPMA shown in Figs. 7, 8, 9. J. A. Brooks et al.31) proposed that the Ni concentration profile of austenite between ferrite could be used to distinguish solidification mode. The schematic sketch of AF and FA mode is shown in Fig. 12. In the FA mode, the Ni content is a minimum in the austenite adjacent to the ferrite and increases to a maximum between the ferrite (at the cell boundary). In the AF mode, Ni is a maximum in the austenite adjacent to the ferrite (at the cell boundary) and is a minimum between the ferrite. According to the Ni concentration profile in Figs. 7, 8, 9, it is confirmed that the solidification mode remains the FA mode. Under the current experiment conditions, the solidification mode is not affected by the nitrogen gas pressure.
When the undercooling reaches the critical value, the primary ferrite dendrites firstly precipitate from the melt. Cr, as the ferrite stabilizer, is absorbed by ferrite, and Ni, as austenite stabilizer, is rejected to the remaining liquid. When the undercooling of liquid exceeds the nucleation undercooling of austenite, the austenite becomes more stable and envelops the solidified ferrite. After the liquid disappears, the solid state transformation δ→γ occurs on cooling. The primary ferrite, without complete transformation, retains down to room temperature, and forms a variety of morphology. The skeletal and lathy ferrites are frequently observed in austenite steels. Both the skeletal ferrite and lathy ferrite result from the primary ferrite solidification and the subsequent diffusion-controlled solid state transformation on cooling.32)
The solute partitioning occurred in the solid state transformation changes the concentration profile of alloy elements between dendrites and interdendritic regions at the end of solidification. It’s difficult to study the effect of nitrogen gas pressure on microsegregation during solidification via experiment. Hence, the microsegregation model was employed to investigate this problem. As much as 75 to 80 pct of the liquid in AISI 304 stainless steel may solidify by primary ferrite before austenite starts to form from the remaining liquid.32) Hence, it is reasonable to disregard the phase change at the end of solidification and subsequent solid state transformation, allowing the complete liquid to solidify as ferrite.33) The complete mixing in the liquid is assumed and solute diffusion in the solid is partially considered. The Cr and Ni concentration profile in solid could be predicted with the Clyne-Kurz equations.34)
|
C
s
=k
C
0
[1-(1-2Ωk)
f
s
]
(k-1)/(1-2Ωk)
| (8) |
|
Ω=α[1-exp(-1/α)]-1/2exp(-1/2α)
| (10) |
Where
D represents the diffusion coefficient in solid, m
2/s;
ts is the local solidification time (LST), s;
C0 is the nominal concentration, wt%;
λ2 is the SDAS, m;
k is the equilibrium partition coefficient;
fs is the solid fraction.
The equilibrium partition coefficient of Cr and Ni is 1.014 and 0.779,35) respectively. The diffusivities of Cr and Ni in the ferrite are D0Cr=7.8 ×10−5 m2/s and QCr=51240 Cal/mole and D0Ni=9.7×10−4 m2/s and QNi=62700 cal/mole.33)
Figure 13 shows the calculated Cr and Ni concentration profile with the maximum LST and SDAS. Cr is positive segregation in ferrite, whereas Ni is negative segregation. As the kCr>1 and kNi<1, the firstly solidified ferrite core enriches in Cr and depletes in Ni. The content of Cr and Ni within the ferrite core is 18.24 pct and 6.23 pct, respectively. As the solidification proceeds, Cr content gradually decreases, whereas Ni content gradually increases. In this study, due to the quick cooling rate and small ingot size, the effect of nitrogen gas pressure on Cr and Ni microsegregation is nearly negligible during solidification.
Figure 14 shows the content of residual ferrite under different nitrogen gas pressures. The content of residual ferrite is dependent on the chemical composition and cooling rate. The distribution of residual ferrite content across the width of ingot was M-shaped, which was consistent with the result of M. Saied.19) This may be due to the variation of cooling rate. Martorano et al. proposed that an increase in cooling rate caused a decrease in the content of residual ferrite.36)
The increase of nitrogen gas pressure not only affects the cooling rate in ingots, but also affects the nitrogen content in ingots. The content of residual ferrite firstly decreased and then increased with the increase of nitrogen gas pressure from 0.1 to 1.2 MPa. This may be the result of a combination of nitrogen content and cooling rate. The mass fraction of nitrogen in electrodes and ingots is shown in Table 4. The nitrogen pickup in ingot could be improved by increasing the nitrogen gas pressure during PESR process. When the nitrogen gas pressure was 0.1, 0.8 and 1.2 MPa, due to the difference of initial nitrogen content in electrodes, the final nitrogen content in ingots was 0.043%, 0.05% and 0.0484%, respectively. Then, the content of residual ferrite was calculated with the Eqs. (11), (12), (13) considering the difference in nitrogen content, which was developed by De Long.37) At 0.1, 0.8 and 1.2 MPa, the calculated content of residual ferrite was 5.93%, 5.37% and 5.5%, respectively, and the variation tendency was in agreement well with the measured results.
Table 4. The mass fraction of nitrogen in electrodes and ingots.
| Sample | The mass fraction of nitrogen, wt % |
|---|
| PESR-1 | PESR-2 | PESR-3 |
|---|
| Electrode | 0.0440 | 0.0450 | 0.0425 |
| Ingot | 0.043 | 0.05 | 0.0484 |
With the increase of gas pressure from 0.1 to 1.2 MPa, the cooling rate firstly decreased and then increased. The hypothetical content of residual ferrite should firstly increase and then decrease with the gas pressure increasing, however, which is contradictory with the experiment results. This contradiction indicates that the variation of nitrogen content plays a more important role in the content of residual ferrite than the cooling rate.
|
Ferrite%=166.66⋅(C
r
eq
/N
i
eq
-0.738)
| (11) |
|
C
r
eq
=[Cr%]+[Mo%]+1.5[Si%]+2.5[Al%]+2.5[Ti%]+18
| (12) |
|
N
i
eq
=[Ni%]+30[C%]+30[N%]+0.5[Mn%]+36
| (13) |
The measured mass fraction of Cr and Ni in residual ferrite is 25–27 pct and 1.8–2.0 pct, respectively, which is far deviated from the calculated value with the Clyne-Kurz equation. The solid state transformation is responsible for the deviation. As seen from Table 3, the residual ferrite had a minimum enrichment of Cr and depletion of Ni at 0.1 MPa, whereas the maximum enrichment of Cr and depletion of Ni occurred at 0.8 MPa. During the solid state transformation from ferrite to austenite, the additional partitioning takes place, resulting in the absorption of Cr and rejection of Ni within the ferrite. The more complete the transformation is, the more serious the enrichment of Cr and depletion of Ni within the ferrite becomes.33)
