# Dynamic Recrystallization and Hot Workability of 316LN Stainless Steel

^{1}

^{2}

^{3}

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## Abstract

**:**

^{−1}using a Gleeble-1500 thermal-mechanical simulator. The microstructural evolution during deformation processes was investigated by studying the constitutive law and dynamic recrystallization behaviors. Dynamic recrystallization volume fraction was introduced to reveal the power dissipation during the microstructural evolution. Processing maps were developed based on the effects of various temperatures, strain rates, and strains, which suggests that power dissipation efficiency increases gradually with increasing temperature and decreasing stain rate. Optimum regimes for the hot deformation of 316LN stainless steel were revealed on conventional hot processing maps and verified effectively through the examination of the microstructure. In addition, the regimes for defects of the product were also interpreted on the conventional hot processing maps. The developed power dissipation efficiency maps allow optimized processing routes to be selected, thus enabling industry producers to effectively control forming variables to enhance practical production process efficiency.

## 1. Introduction

^{−1}. However, the study neglected the effects of deformation processes on the processability of the material. Zhang [10] investigated the high temperature behavior of 316LN and simulated the microstructural evolution during deformation without an explanation of the microstructural evolution in locating the optimal regions for large plastic deformation processes. Bai [11] also constructed the constitutive equation for 316LN in a relatively narrow strain range, but did not explain the means in which microstructural evolution influences the macroscopic deformation. Zhang [12] studied the dynamic and post deformation recrystallization of 316LN stainless steel without clearly identifying the proper processing conditions for the material.

## 2. Materials and Methods

^{−1}, 0.01 s

^{−1}, 0.1 s

^{−1}, 1 s

^{−1}, and 10 s

^{−1}up to the true (logarithmic) strain of 0.916 (the maximum compression ratio of all specimens was 60%), followed by water quenching. For microstructural investigations, the compressed specimens were then sectioned parallel to the deformation axis, mounted and then polished in aqua regia (HNO

_{3}20 mL, HCl 60 mL).

## 3. Results and Discussion

#### 3.1. Constitutive Analysis

^{−1}are shown in Figure 1. As expected, the flow stress is dependent on both the deformation temperature and strain rate. Meanwhile, the flow stress increased with increase of the strain rate and decrease of the temperature. Furthermore, all the flow stress curves increased significantly at the initial stage of hot deformation, attributed to the work hardening, and then followed by a relatively steady state which indicates the occurrence of DRX [20,21].

_{act}is determined as the activation energy of deformation which is the threshold of the dynamic recrystallization; σ

_{p}is the first peak stress of each curve; $\dot{\mathsf{\epsilon}}$ is the strain rate; R is the universal gas constant (8.314 J·mol

^{−1}·°C

^{−1}); T is the temperature; A

_{1}, A

_{2}, A

_{3}, n

_{1}, n, β, and α (≈β/n

_{1}) are material constants. Equation (2) is a transformation of Equation (1). So the variables can be defined as:

#### 3.2. Microstructure Evolution

_{c}has been attained. It is the core characteristic of the softening mechanism of the materials with low and medium stacking fault energy (SFE). When the softening process is governed by dynamic recrystallization, the flow stress passes through a peak σ

_{p}and drops to a steady state regime, as shown in Figure 3. Generally, the critical strain ε

_{c}symbolizing the start of DRX can be obtained either by a direct microstructure observation or through an analysis of the flow stress curve. However, microstructure observation is a more complicated and time-consuming method compared to the flow stress curve analysis, as it requires a large number of samples for examination. The flow stress curve analysis method, firstly proposed by Kocks and Mecking [23], and then further developed by McQueen and Ryan [24], can be used to emphasize the point where DRX occurs on the flow stress curve.

_{sat}. During the entire process of deformation, work hardening occurs as a result of the refinement of the microstructure caused by the increased strain, and thus the flow stress can be described by:

_{1}is dislocation interaction constant relating to dislocation spacing and stress component, G is the shear modulus, and b is the Burgers vector. Based on the Kocks-Mecking model, the apparent hardening depends on the dislocation density which is the result of interplay between storage and annihilation of dislocations as follows:

_{1}and k

_{2}are constants related to microscopic parameters. Consequently, the Kocks-Mecking model assumes that hardening is caused by the increase of the average value of dislocation density, which is directly proportional to the square root of the dislocation density. Therefore, when combined with Equation (7), this can be written as:

_{0}is the initial real stress, C

_{1}is the integration constant and ρ

_{0}is the initial dislocation density. The equation can also be expressed as:

_{c}is the stress at critical strain ε

_{c}. Figure 4a shows the relationship between ε

_{c}and ε

_{p}, and it is similar to the result studied by Ji [26]. Figure 4b shows the method to identify the saturation stress σ

_{sat}and the steady flow stress σ

_{ss}due to the dynamic recrystallization: the intersection point where a tangent line at the critical strain point cuts the x-axis (θ = 0) is the σ

_{sat}; the intersection point of the lower value where the σ-θ curve cuts the x-axis (θ = 0) is the σ

_{ss}[27,28].

_{sat}− σ

_{ss}. Therefore, the conditions in Equation (16) are only fulfilled if$\text{}X\left({\mathsf{\epsilon}}_{\mathrm{c}}\right)={X}^{\prime}\left({\mathsf{\epsilon}}_{\mathrm{c}}\right)={X}^{\u2033}\left({\mathsf{\epsilon}}_{\mathrm{c}}\right)={X}^{\u2034}\left({\mathsf{\epsilon}}_{\mathrm{c}}\right)=0$, so as to guarantee that κ

_{ε}is continuous which is defined as:

_{p}is the first peak strain, q is the constant corresponding to nucleation mode, and α

_{2}is the term associated with the nucleation and growth rates. These parameters can be used to represent the flow behavior of the material under consideration, although some rate equations must also be provided, as will be shown later in the paper. The reader must bear in mind that the latter approach neglects any possible hardening taking place concurrently with the deformation during the dynamic recrystallization regime. This approximation is valid when the deformation process is mainly governed by softening due to dynamic recrystallization, as is generally accepted once the peak stress has been attained. The experimental values of the Avrami exponents reported later also support the validity of this approach. Obviously, at the critical strain, X(ε

_{c}) = 0.

_{DRX}is the real stress derived from isothermal compression experimental data, σ

_{DRV}is the predicted stress without the recrystallization when ε > ε

_{c}. By plotting $\mathrm{ln}\left(-\mathrm{ln}\left(1-X\left(\mathsf{\epsilon}\right)\right)\right)$ against $\text{ln}\left((\mathsf{\epsilon}-{\mathsf{\epsilon}}_{\mathrm{c}})/{\mathsf{\epsilon}}_{\mathrm{p}}\right)$, the material constant α

_{2}and q can be calculated by the slope and the intercept of the linear regression line, as shown in Figure 5. The recrystallization fraction of the 316LN can be given by:

^{−1}. Figure 6a displays the initial microstructure of the material, in which the average size of the initial grains is measured to be about 75 μm. Figure 6b displays microstructure after deformation to a strain of 0.1, and because of the orientation difference between adjacent grains, smaller grains emerge at the grain boundaries. The mixed partially recrystallized structure observed in Figure 6c was attained after being compressed to a strain of 0.4, and the occurrence of dynamic recrystallization (finer equiaxed grains) during deformation can be clearly recognized, showing that the average size of the newly formed grains is measured to be about 45 μm. But there are still some unexpected grains (grain diameters above 45 μm at spot A). Figure 6d displays recrystallized grains obtained after being compressed to a strain of 0.9, in which grains are 75% (average size is measured to be about 20 μm) smaller than that of the initial microstructure state, corresponding to steady state of stresses.

#### 3.3. Hot Deformation Behavior of 316LN Stainless Steel

_{act}in the Equation (1) is replaced by the deformation energy Q and the σ

_{p}is replaced by the real stress σ, the strain rate sensitivity m can be described as:

_{2}are the functions of strain ε, and it follows that m, η and ζ are also the functions against strain ε. Figure 8 demonstrates the hot processing map of 316LN stainless steel at ε = 0.4 and ε = 0.9. The number against each contour represents the power dissipation efficiency in the processing map, and the shaded area represents the flow instability (unstable) regime.

^{−1}, with a peak efficiency of about 57% occurring at 1050 °C/0.001 s

^{−1}; (A(2)) 1110–1150 °C and 0.001–0.03 s

^{−1}, with a peak efficiency of about 53% occurring at 1150 °C/0.001 s

^{−1}; (B) 950–1010 °C and 3.25–10 s

^{−1}, with a peak efficiency of about 10%; as well as (C) 1100–1150 °C and 1–10 s

^{−1}, with a peak efficiency of about 31%.

^{−1}(Figure 6c) corresponds to the domain A(2) in Figure 8a, representing dynamically recrystallized microstructure with typical wavy grain boundaries and showing fine grains (average grain diameter is about 45 μm), while there are still some unexpected grains (grain diameters above 45 μm at spot A) due to the insufficient deformation. Domains A(1) and A(2) in Figure 8a transform into one domain in Figure 8b owing to the abundant dynamic recrystallization in the temperature range of 990–1150 °C and the strain rate range of 0.001–0.01 s

^{−1}with the peak efficiency of 0.57 occurring at 1150 °C/0.001 s

^{−1}(domain A in Figure 8b). The domain B in Figure 8a is the regime assumed to be inappropriate for hot deformation because its power dissipation efficiency is very low and its instability parameter ξ < 0. Along with the increase of strain, the power dissipation efficiency becomes smaller (shown as the domain B in Figure 8b), which means it may cause tearing, surface cracking, inter-crystalline cracks, or inhomogeneous microstructure. The microstructure of the specimen deformed at the temperature range of 1100–1150 °C and strain rate range of 1–10 s

^{−1}(shown as the domain C in Figure 8a,b) corresponds to the flow instability regime exhibiting flow localization. In the localized regions, the material has undergone static recrystallization during cooling resulting in fine grains along the localized bands [3,5].

## 4. Conclusions

^{−1}and (b) 1110–1150 °C/0.001–0.03 s

^{−1}, in which dynamic recrystallization is apt to occur. The dynamic recrystallization domain at ε = 0.9 in the temperature range of 990–1150 °C and the strain rate range of 0.001–0.01 s

^{−1}with the peak efficiency of 0.57 occurring at 1150 °C/0.001 s

^{−1}is illustrated on the processing map.

## Acknowledgments

## Author Contributions

## Conflicts of Interest

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**Figure 1.**Flow curves for the 316LN stainless steel at: (

**a**) $\dot{\mathsf{\epsilon}}$ = 0.001 s

^{−1}; (

**b**) T = 1150 °C.

**Figure 2.**Relationships between: (

**a**) peak stress and ln(strain rate); (

**b**) ln(peak stress)and ln(strain rate); (

**c**) ln[sinh(ασ

_{p})] and ln(strain rate); (

**d**) 10,000/(T + 273.15) and ln[sinh(ασ

_{p})].

**Figure 3.**Schematic description of the flow behavior of 316LN at high temperature T = 1150 °C/$\dot{\mathsf{\epsilon}}$ = 0.001 s

^{−1}(σ

_{sat}is the saturation stress; σ

_{ss}is the steady state stress of the experimental data; σ

_{c}is the critical stress; σ

_{p}is the peak stress; ε

_{c}is the critical strain; ε

_{p}is the peak strain; Δσ = σ

_{sat}− σ

_{ss}).

**Figure 4.**(

**a**) Relationship between ε

_{c}and ε

_{p}; (

**b**) the true stress-strain hardening rate curve at T = 1050 °C/$\dot{\mathsf{\epsilon}}$ = 0.001 s

^{−1}.

**Figure 6.**Microstructures of 316LN stainless steel at T = 1150 °C and $\dot{\mathsf{\epsilon}}$ = 0.01 s

^{−1}: (

**a**) initial microstructure; (

**b**) ε = 0.1; (

**c**) ε = 0.4; (

**d**) ε = 0.9.

**Figure 8.**Hot processing map of 316LN stainless steel at: (

**a**) ε = 0.4; (

**b**) ε = 0.9 (The legend scale represents the power dissipation efficiency).

**Figure 9.**Updated hot processing map of 316LN stainless steel at (

**a**) $\dot{\mathsf{\epsilon}}$ = 0.01 s

^{−1}; and (

**b**) $\dot{\mathsf{\epsilon}}$ = 0.1 s

^{−1}(The legend scale represents the power dissipation efficiency. Critical strain curve is believed to be the start of the dynamic recrystallization).

Component | C | Cr | Ni | Mo | Mn | P | S | Si | N | Fe |
---|---|---|---|---|---|---|---|---|---|---|

wt. % | 0.017 | 17.03 | 12.71 | 2.53 | 1.29 | 0.020 | 0.001 | 0.34 | 0.12 | Balance |

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**MDPI and ACS Style**

Sun, C.; Xiang, Y.; Zhou, Q.; Politis, D.J.; Sun, Z.; Wang, M.
Dynamic Recrystallization and Hot Workability of 316LN Stainless Steel. *Metals* **2016**, *6*, 152.
https://doi.org/10.3390/met6070152

**AMA Style**

Sun C, Xiang Y, Zhou Q, Politis DJ, Sun Z, Wang M.
Dynamic Recrystallization and Hot Workability of 316LN Stainless Steel. *Metals*. 2016; 6(7):152.
https://doi.org/10.3390/met6070152

**Chicago/Turabian Style**

Sun, Chaoyang, Yu Xiang, Qingjun Zhou, Denis J. Politis, Zhihui Sun, and Mengqi Wang.
2016. "Dynamic Recrystallization and Hot Workability of 316LN Stainless Steel" *Metals* 6, no. 7: 152.
https://doi.org/10.3390/met6070152