Microstructural Evolution of Cold-Rolled Type 347H Austenitic Heat-Resistant Steel

Abstract

The influence of cold rolling deformation degree (15%, 30%, 45%, 60%, 75%, and 90%) on the microstructural evolution and the mechanical properties of type 347H austenitic heat-resistant steel was investigated using optical microscopy, X-ray diffraction, magnetic hysteresis loop measurement, transmission electron microscopy, and a hardness test. Two types of martensite formed in the deformed specimens, as thin ε-martensite in the cold-rolled steels when the deformation degree was less than 60%, and α′-martensite in the heavily cold-rolled steels when the deformation degree ranged from 60% to 90%. Furthermore, the amount of α′-martensite increases rapidly with the increase in the cold rolling deformation degree. Hence, 60% is considered as the critical point of cold rolling reduction for the formation of α′-martensite. If the specimen experienced a cold rolling reduction of 90%, ε-martensite was hardly observed, while the volume faction of the α′-martensite amounts to 25%. It is verified by the TEM observations that the α′-martensite is transformed from the austenitic matrix as well as the preformed ε-martensite.

1. Introduction

Austenitic heat-resistant steel has excellent corrosion resistance and good ductility [1,2,3,4]; for instance, the 18Cr-8Ni type has been widely used in the boiler, petroleum, and other industries. It is imperative to develop super-critical and ultra-super-critical generators to reduce pollution and cost; therefore, austenitic heat-resistant steel with an excellent combination of high ductility, strength, and steam oxidation resistance was required. Most austenitic heat-resistant steels are unstable during cold deformation, and the strength of the steel can be enhanced by converting austenite (face-centered cubic, fcc) into deformation-induced martensite [5,6]. Two types of martensite, ε-martensite and α′-martensite, are commonly obtained during cold rolling reduction. The ε phase has a hexagonal close-packed structure and exhibits paramagnetism, while the α′-martensite has a body-centered cubic structure and exhibits ferromagnetism [7,8,9,10]. In austenitic heat-resistant steel, a certain amount of high-strength and high-stiffness martensite induced by deformation is expected to achieve a perfect combination of strength and toughness [11].

During the early stage of deformation study in AISI 304LN stainless steel [12], stacking faults, mechanical twins, and ε-martensite are formed on shear bands in parent austenite, and γ-austenite and the ε-martensite follow a crystallographic orientation relationship given by {111}γ//{0001}ε, <110>γ// $<11\stackrel{-}{2}0>$ε [13]. When the cold rolling deformation degree increased, the α′-martensite formed at the intersections of the ε-martensite plates, and the α′-martensite grew by consuming the preformed ε-martensite and the parent austenite phase. Both the intersection of slip lines and the isolated micro shear bands of the austenite served as nucleation sites of the α′-martensite [14]. Apart from this, the α′-martensite also nucleated at either the twinning edge or the austenite grain boundaries. During the cold rolling process, the formation of deformation-induced martensite involves a variety of mechanisms, such as γ (fcc) → ε (hcp), γ (fcc) → ε (hcp) → α′ (bcc), γ (fcc) → α′ (bcc), and γ (fcc) → deformation twin → α′ (bcc). These mechanisms depend on the thermodynamic energy of martensitic transformation [15]. It has been reported that the stacking fault energy is an essential influence parameter for the formation of the deformation twin, the ε-martensite, and the α′-martensite [16,17].

Although the evolution of the deformation-induced martensite in cold-rolled AISI 304 and 304L steels has been studied, the type 347H austenitic heat-resistant steel is barely addressed. In order to improve the mechanical property of the type 347H austenitic heat-resistant steel by martensitic strengthening, cold rolling was employed to obtain a certain content of martensite in this study. It is therefore necessary to investigate the influence of the cold rolling deformation on the microstructure of the type 347H austenitic heat-resistant steel, and to clarify the relationship between the cold rolling reduction and the α′-martensite content.

2. Experimental Procedure

The chemical composition of the used type 347H austenitic heat-resistant steel specimens is shown in

Table A1. Chemical composition of the investigated type 347H austenitic heat-resistant steel (in weight percent).

Element	C	Cr	Ni	Nb	N	Mn	P	Mo	S	Fe
wt.%	0.059	17.60	10.71	0.54	0.013	1.59	0.024	0.116	0.0008	Balanced

. The specimens with a size of 65 mm × 20 mm × 10 mm were cut from a φ120 mm bar. The specimens were solution heat-treated at 1150 °C for 30 min followed by water quenching. Different thickness reductions from 15% to 90% (coded as 15%, 30%, 45%, 60%, 75%, and 90%) were rolled by a laboratory rolling mill at 25 °C. X-ray diffraction (Bruker D8 Advanced, Ettlingen, Germany) measurements of the cold-rolled samples, which are denoted as 15%CR, 30%CR, 45%CR, 60%CR, 75%CR and 90%CR, for phase identification was made using Cu Kα radiation in the range of 40~100°.

In order to represent the magnetic characterization, the saturation magnetization of each sample was obtained by using a vibrating sample magnetometer—VSM (LakeShore 7400, Lake Shore Cryotronics, Ltd., Westerville, OH, USA). To evaluate the volume fraction of α′-martensite, an equation was applied as follows [18]:

$$V_{{\mathsf{\alpha }}^{\prime }}={\displaystyle \frac{I_{{\mathrm{S}}^{\prime }}}{I_{\mathrm{S}}}}$$

(1)

where I_S′ is the magnetic saturation of the cold-rolled samples tested in VSM and I_S is the magnetic saturation of a 100% martensitic sample. It has to be noted that I_S was determined as 163 Am²·kg⁻¹, indicating the intrinsic magnetization saturation of the martensite phase [19].

Microstructural examinations of the cold-rolled specimens were carried out by using an optical microscope (DMI 8, Leica Microsystems, Jena, Germany). The hardness of each sample was determined on a Vickers hardness test (0.2 kg load) device (Duramin-A300, Beijing Shoufeng united measuring equipment Co., Ltd.,  Beijing, China). The thin foils for the analysis were prepared by a double-jet electrolytic polisher (MTP-1A, CASIO, Tokyo, Japan) at a voltage of 50 V at a temperature of −25 °C. The electrolyte contained 5 vol% of perchloric acid and 95 vol% of alcohol. Transmission electron microscope (TEM, JEM 2100f, JEOL, Tokyo, Japan) operated at 200 kV was used to examine the martensite formation.

3. Results and Discussion

Figure 1 shows the optical microscopic images of the 347H austenitic heat-resistant steel after solution treatment at 30%CR, 60%CR, and 90%CR. Through the comparison of Figure 1a and b–d, respectively, it can be seen that, different from the solid solution-treated sample with a complete austenite microstructure, parallel shear bands can be seen inside the deformed grains for the 30% deformation amount, and some needle-like morphology phases appear. For the 60% deformation sample, the number of intersecting needle-like structures in the grains is significantly increased compared with the 30% sample. When the cold deformation reaches 90%, the large deformation makes it almost impossible to observe any obvious shear bands. In order to further understand the dislocation and microstructure changes during the cold rolling process, TEM observations were performed for each sample.

Figure 2 shows the deformation microstructure transmission electron microscope image of the sample with 30% deformation. Through the comprehensive observation of Figure 1b and Figure 2a, it can be seen that there should be shear bands observed in the parallel shear bands observed in the metallographic phase. In particular, the slip line cluster in Figure 2a corresponds to the {111} crystal plane of austenite, which is confirmed by the diffraction pattern of Figure 2d. Moreover, previous studies have also shown that when the deformation reaches a critical value, part of the Shockley dislocations will nucleate on the {111} crystal plane, and the structural transformation is dominated by the slip of the parent austenite along the {111} crystal plane [20]. Therefore, these shear bands in Figure 1b are considered as the slip lines of the {111} crystal plane of the parent austenite [21,22]. A stacking fault formed due to the motion of the Shockley partial dislocations along the {111} plane in Figure 2b, as lamellar lines are the typical features of a stacking fault in fcc materials [15]. With the movement of the Shockley dislocations, the ε-martensite with a hexagonal close-packed structure formed by the overlapping of the stacking faults on the alternately close packed {111} planes of austenite, as shown in Figure 2c. Based on the diffraction patterns presented in Figure 2f, two shear bands with a different orientation can be identified. The bright one is the ε-martensite band, and the other is the austenitic slip line. The ε-martensite (hcp) with the [0001] zone axis nucleated in the austenitic matrix (fcc) with the $[\overline{1}12]$ zone axis.

Figure 3 shows the TEM images of the 60% cold-rolled type 347H austenitic heat-resistant steel sample, and that the lamellar ε-martensite was represented, accompanied by the mechanical twin. The presence of the mechanical twinning (see Figure 3a) indicates that the free energy required for the nucleation of the mechanical twinning is higher than that of the ε-martensite [16]. In view of the similar morphology of the slip lines, the mechanical twinning, the ε-martensite, the electron diffraction patterns were employed to distinguish them. In Figure 3a,b, a large number of slip lines with a width of 100 nm existed within a grain. The 20 nm-width thin bands represent mechanical twinning, and the 10 nm-width ones are the ε-martensite. Based on the electron diffraction patterns in Figure 3d, the orientation relationship between the ε-martensite and the austenite is identified as (402)γ// $(21\overline{3}0)$ε and [ $\stackrel{-}{1}22$]γ//[0001]ε [12]. It is worth noting that although the nucleation process of α′-martensite is not obvious at the junction of the ε-martensite bands and the isolated austenite band, it can still indicate that the rough edge may become a potential nucleation site of α′-martensite.

In the 90%CR type 347H austenitic heat-resistant steel sample (see Figure 4), due to the high deformation level, one could hardly observe the shear band containing the slip line, stacking fault, mechanical twinning, or ε-martensite. Instead, numerous dislocations are located in the austenitic matrix (see Figure 4a). Such a considerable amount of dislocations promoted the nucleation of the α′-martensite, and the high density of the dislocations accelerated the γ → α′ phase transformation rate. By analyzing the diffraction rings in Figure 4c, the α′-martensite is identified, and is proved to nucleate within the austenite matrix [23]. In addition, the diffraction pattern (Figure 4d) for the border between the α′-martensite and the austenite (Figure 4b) revealed that the α′(bcc) with the zone axis of [ $01\stackrel{-}{1}$] nucleated in the austenitic matrix along the direction of [ $\stackrel{-}{1}13$]. Since no significant evidence indicates the existence of the ε-martensite in the 90%CR sample, the α′-martensite possibly formed at the expense of the ε-martensite [24].

The plasticity mechanism of metals depends largely on the stacking fault energy (SFE). The SFE for the formation of mechanical twins is between 12 mJ/m² and 35 mJ/m² in the Fe-Mn-C system, while both the strain-induced martensite and mechanical twins can form at an SFE below 18 mJ/m² [13]. The deformation mechanisms were found to be dislocation slip, mechanical twins, ε-martensite, and α′-martensite in type 347H austenitic heat-resistant steel. The SFE could be estimated by the main composition, according to the following equation proposed by Schramm and Reed [25]:

$$\mathrm{SFE}({\mathrm{mJ}/\mathrm{m}}^2)=-53+0.7(\%\mathrm{Cr})+6.2(\%\mathrm{Ni})+3.2(\%\mathrm{Mn})+9.3(\%\mathrm{Mo})$$

(2)

Accordingly, the approximate SFE of the type 347H austenitic heat-resistant steel is calculated to be 32 mJ/m², which is not as low as expected for low-SFE material. This suggests that the SFE value varies in different austenitic systems, which affects the formation of the mechanical twins and the strain-induced martensite. In the same Fe-Cr-Ni system, the SFE of AISI 304LN stainless steel was calculated by A. Hedayati et al. [26] to be 19 mJ/m², which is lower than that of type 347H austenitic heat-resistant steel. Hence, the stability of austenite in the type 347H austenitic heat-resistant steel is improved, and the formation of α′-martensite is therefore suppressed. It turns out to be the increasing deformation that induced the formation of α′-martensite.

Figure 5 shows the XRD patterns of the cold-rolled type 347H austenitic heat-resistant steel specimens from 0% up to 90%CR. Prior to the cold rolling, the steel is approximately complete austenite. The α′-martensite is found to exist under 60% cold rolling, which disagrees with the TEM observation in Figure 3. This inconsistence can be interpreted by the limited content of the α′-martensite under small deformation. As the cold rolling deformation is increased from 60% to 90%, the intensity of the α′-martensite peaks rises rapidly, indicating that 60%CR is the critical deformation degree, upon which the α′-martensite began to form at a large rate. Since the diffraction peak of (002)ε is extremely close to that of (111)γ, the peak of ε-martensite failed to be separated. The peak position shifts from (111)γ to the direction of the (002)ε peak with the increasing deformation degree, which is attributed to the increasing amount of the ε-martensite and the simultaneous decrease in the austenite phase content. This on the other hand indicated that a certain proportion of the austenite might transform into ε-martensite in the process of cold rolling.

With the increasing cold rolling reduction, the (111)γ diffraction peak evidently shifts to the high-angle direction [27]. The peak intensity of the austenite and the ε-martensite is gradually decreased, and the peak corresponding to the α′-martensite emerges as the cold deformation is increased from 60% to 90%, and accordingly the peak intensity increases. Thus, the sequence of the martensitic transformation is summarized to be γ (fcc) → ε (hcp) → α′ (bcc). Both the transformation from the ε (hcp) to the α′ (bcc) and that from the γ (fcc) to the α′ (bcc) are likely to take place.

The magnetic hysteresis loops of the samples under different cold rolling deformations are presented in Figure 6. The saturation magnetization of 90%CR sample is up to 38.86 Am²·kg⁻¹, while the saturation magnetization of the sample comprised with complete austenite is nearly 0 Am²·kg⁻¹. The ratio of ferromagnetic and non-ferromagnetic phases varied with cold rolling in the explored type 347H austenitic heat-resistant steel [10], which should be responsible for the increasing saturation magnetization with the cold rolling deformation. The magnetization of the 0–60%CR samples enhanced slowly and was saturated as the magnetic field increased, which is characteristic of paramagnetism (austenite and ε-martensite). In contrast, the magnetization of the 75%CR and the 90%CR samples saturated in a faster manner, followed by the ferromagnetic characteristic (α′-martensite). Hence, the increase in cold rolling deformation induces a transition from the paramagnetism type to the ferromagnetic type. The samples containing a higher volume fraction of the ferromagnetic phase exhibit a higher saturation magnetization. As a result, when the cold rolling reduction is increased from 60% up to 90%, a considerable amount of α′-martensite is formed.

The volume fraction of the α′-martensite, evaluated by magnetic measurement, is displayed in

Table A2. The α′-martensite contents measured by the saturation magnetization.

CR (%)	0	15	30	45	60	75	90
Vα′ (%)	0	2.56	4.52	5.52	7.43	11.42	23.84

. The specific value of the α′-martensite volume fraction under the same deformation exhibits a slight distinction by applying the two methods, yet the α′-martensite volume fraction presents a similar variation tendency to the cold rolling deformation, as shown in Figure 7. With the increasing deformation degree from 0 to 60%, the volume faction of the α′-martensite is slowly enhanced to be about 10%. When the cold rolling reduction is further increased, the growth of α′-martensite is accelerated. Under the 90% cold rolling reduction, the volume faction of the α′-martensite reaches 25% or above.

Figure 8 collects the measured Vicker hardnesses of the type 347H austenitic heat-resistant steel under different cold reductions, which increase with the increasing cold rolling deformation from 178 HV to 473 HV. The non-linear relation for hardness values with the cold deformation degree is divided into two segments. The hardness increases proportionally in a linear mode during the early stage, and then the increase slows down after a 30% cold rolling reduction. The change in hardness with cold rolling deformation is attributed to the extensive dislocation tangles, and the formation of the deformation-induced martensite. At the initial stage, the α′-martensite plays a predominant role in the increase in the hardness, while the α′-martensite and the dislocations contribute more at the final stage.

4. Conclusions

This study systematically analyzed the microstructure evolution process of 347H austenitic heat-resistant steel during the cold rolling process. The formation mechanism was found to include the following:

(1)

Cold rolling processing can induce a transition from metastable austenite to deformation-induced martensite. The formation mechanisms include γ (fcc) → ε (hcp), γ (fcc) → mechanical twins, γ (fcc) → ε (hcp) → α ‘(bcc), and γ (fcc) → α′ (bcc).

(2)

In the sample with a cold rolling deformation rate of 30%, the austenite { 111 } crystal plane promotes the formation of stacking dislocations through the slip of Shockley’s dislocations. These overlapping stacking dislocation regions undergo partial austenite to ∂ martensite transformation. When the cold rolling deformation reaches 60%, the critical deformation of α′-martensite formation is triggered, and the mechanical twin structure begins to appear in the matrix austenite and a small amount of α′-martensite is formed. When the cold rolling deformation reaches 90%, a significant amount of α′-martensite is produced due to the consumption of ɛ-martensite and the interaction of retained austenite.

(3)

It was confirmed by TEM observation that high-density dislocations significantly promoted the formation of α′-martensite. When the cold rolling deformation reached 90%, the content of α′-martensite was the highest and the hardness of the cold-rolled steel sample was increased to 473 HV.