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Article

Effect of Thermomechanical Processing on Grain Boundary Character Distribution and Creep Properties of SP2215 Heat-Resistant Steel

by
Wen Feng
1,2,*,
Ting Sun
1,
Tianyu Zhao
1,
Junjie Zhou
1 and
Zhengyu Han
1
1
School of Mechanical Engineering, Jiangsu Ocean University, Lianyungang 222005, China
2
Jiangsu Institute of Marine Resources Development, Lianyungang 222005, China
*
Author to whom correspondence should be addressed.
Crystals 2026, 16(5), 282; https://doi.org/10.3390/cryst16050282
Submission received: 30 March 2026 / Revised: 21 April 2026 / Accepted: 22 April 2026 / Published: 24 April 2026
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Versions Notes

Abstract

This study presented an application of thermomechanical processing consisting of cold rolling and subsequent annealing in SP2215 heat-resistant steel to investigate the effects of thermomechanical processing parameters on the evolution of grain boundary character distribution (GBCD) and to elucidate the relationship between GBCD and creep properties. The experimental results show that the optimal process, characterized by 10% cold rolling reduction followed by annealing at 1100 °C for 10 min, was determined to significantly increase the fraction of low-Σ coincidence site lattice (CSL) boundaries up to 74.27%, and effectively disrupt the connectivity of the random boundary network, as corroborated by the highest average twin-related domain (TRD) size of 42.58 μm and average number of grains per TRD of 7.28. Such a modified GBCD leads to a notable enhancement in creep performance, resulting from the induction of a high fraction of low-Σ CSL boundaries and the disruption of the random boundary network, which effectively inhibits intergranular crack initiation and propagation during creep deformation.

1. Introduction

With the goals of mitigating pollutant emissions and reducing coal consumption, ultra-supercritical (USC) coal-fired power plants are regarded as a pivotal technical strategy for upgrading thermal power generation [1,2]. To enhance the thermal efficiency and operational safety of USC coal-fired power units, heat-resistant steels are crucial, particularly for key components such as reheaters and superheaters. Currently, Super304H and HR3C heat-resistant steels are the mainstream materials for the heating surfaces of high-temperature superheaters and reheaters [3,4,5]. However, Super304H steel has poor oxidation resistance with a high tendency for oxide scale spallation [6,7]. HR3C steel exhibits insufficient high-temperature strength and is prone to embrittlement after long-term high-temperature aging [8].
To meet the requirements for the design of USC power plants with elevated operating parameters and long-term safe service, a novel austenitic heat-resistant steel, designated SP2215 (22Cr15Ni3.5CuNbN), has been developed and is a promising substitute for traditional heat-resistant steel grades. This newly developed steel demonstrates exceptional high-temperature creep resistance, and its toughness remains nearly unaffected even after prolonged thermal aging. Compared with traditional heat-resistant steels, SP2215 steel offers superior mechanical properties and oxidation resistance with reduced chromium and nickel additions. Extensive prior studies [9,10,11] indicate that the superior high-temperature service properties of SP2215 steel stem from the synergistic stabilization and precipitation-strengthening effects of niobium and nitrogen, coupled with the formation of Cu-rich precipitates within the steel matrix. However, there remains a significant challenge to the service safety and lifespan of SP2215 steel due to intergranular degradation under high-temperature conditions. Hence, strengthening the intergranular degradation resistance of SP2215 steel is critical to further expanding its engineering applications, enhancing service safety and operational durability, and ultimately achieving cost and resource savings.
Grain boundary engineering (GBE), founded on the principle of “grain boundary control and design”, was initially proposed by Watanabe in 1984 [12]. It has been widely verified as a highly prospective strategy for alleviating a series of intergranular degradation behaviors in metallic materials, such as intergranular corrosion [13,14,15,16], intergranular stress corrosion [17,18], embrittlement [19,20], oxidation [21], creep [22], and fatigue [23,24]. The core idea of GBE is to optimize the grain boundary character distribution (GBCD) in low to medium stacking fault energy face-centered cubic metallic materials by increasing the fraction of low-Σ coincidence site lattice (CSL) boundaries and disrupting the connectivity of the random boundary network, mainly through appropriate thermomechanical processing (TMP). Despite the successful implementation of GBE across diverse alloy systems, its utilization in heat-resistant steels, especially in SP2215 steel, remains rarely reported. Therefore, this study focuses on two core research objectives. Firstly, to investigate the effects of TMP parameters, including deformation amount and annealing temperature, on the evolution of GBCD in SP2215 steel. Secondly, to assess the intrinsic correlation between GBCD optimization and alloy creep properties.

2. Experimental Procedures

2.1. Materials Preparation

In the present study, an SP2215 heat-resistant steel billet with a diameter of ϕ110 mm was used, supplied by Jiangsu Wujin Stainless Steel Co., Ltd. (Changzhou, China). The chemical composition of SP2215 heat-resistant steel is shown in Table 1. The as-received billet was cut into rectangular strip specimens by means of the wire-cutting electric discharge machining for cold rolling deformation. To eliminate carbide pinning during grain boundary migration, rectangular strip specimens were solution-treated at 1100 °C for 1 h, followed by water quenching. This processed specimen condition was referred to as the base material (BM). Afterward, various TMP routes consisting of cold-rolling deformation followed by annealing were performed on BM specimens to investigate the effects of deformation amount and annealing temperature on the evolution of GBCD. The rectangular strip specimens were rolled by multi-pass rolling to different final thicknesses, and the average strain rates were between 0.43 s−1 and 1.10 s−1. The surface roughness of working rolls was 0.2–0.3 μm. Detailed TMP parameters and corresponding specimen designations are summarized in Table 2. Furthermore, the specimen fabricated under the optimal parameters was designated as the grain boundary engineering material (GBEM).

2.2. Microstructural Characterization

Microstructural characterization was carried out via a scanning electron microscope (SEM) (Zeiss Merlin Compact, Oberkochen, Germany) integrated with an electron backscattered diffraction (EBSD) detector (Oxford Instruments, Abingdon, Oxfordshire, UK). To observe grain boundaries by EBSD, each specimen was electrochemically polished in an aqueous 10% oxalic acid solution at 8 V for 40 s at ambient temperature. The average grain size was measured using the standard linear intercept method. EBSD-based orientation imaging microscopy (OIM) scans were performed to analyze the GBCD in all specimens. Data acquisition was performed at an accelerating voltage of 20 kV with a step size of 2 μm, and subsequent data analysis was conducted using the Channel 5 software package. For statistical significance, at least three EBSD scans were acquired from various locations on each specimen (each 600 × 600 μm2), and the number of grains analyzed ranged from 1200 to 18,000, depending on grain size. In the present study, grain boundaries with 3 ≤ Σ ≤ 29 were defined as low-Σ CSL boundaries, while the remaining grain boundaries were classified as random boundaries. Grain boundary classification was performed in accordance with Brandon’s criterion [25]. The fraction of low-Σ CSL boundaries reported in this study specifically corresponds to the grain boundary length fraction. In addition to OIM analysis, the ARPGE version 2.3 software (Automatic Reconstruction of Parent Grains from EBSD data) was used to reconstruct parent grains from the original EBSD data [26]. This parent grain reconstruction procedure is grounded in groupoid theory, which enables the extraction of parent-phase grain parameters, including twin-related domain (TRD) size and the number of grains in each TRD.

2.3. Creep Test

The creep testing was performed using an RDL100 creep testing machine (Sinotest Equipment Co., Ltd., Changchun, China) equipped with an extensometer assembly to monitor specimen deformation continuously. Bilateral extensometers (ST 1288, Heidenhain, Ulm, Germany) with a resolution of 1 μm, equipped with tension rods and thermocouples, were adopted to monitor creep strain. The accuracy of the creep strain measurement was ±0.1%. These devices were mounted onto the specimen shoulders and extended to a dedicated support structure located outside the furnace chamber. Creep tests were conducted at a constant applied stress of 325 MPa and a test temperature of 650 °C for each specimen. Deformation data were recorded at 72 s intervals throughout the test, which was terminated automatically upon specimen rupture. Each creep test was performed three times to ensure repeatability and accuracy, and the final properties were the average values. Following the completion of creep testing, each tested specimen was air-cooled to ambient temperature, and the metallographic structures were observed via optical microscopy, while the surface and cross-sectional morphologies after fracture were characterized using SEM.

3. Results and Discussion

3.1. Microstructural Characteristics of BM Specimen

The inverse pole figure (IPF) map of the BM specimen is shown in Figure 1a, where a standard stereographic triangle represents various crystallographic directions indicated by different colors. The grains are clearly oriented in different directions, showing the absence of strong texture in the BM specimen. Additionally, the microstructure of the BM specimen consists mainly of nearly equiaxed grains, with an average grain size (excluding twins) of about 15.11 μm. As shown in Figure 1b, the local misorientation map overlaid with different types of grain boundaries exhibits very low local misorientation values, indicating that the matrix is strain-free after solution treatment. Figure 1c displays the color-coded grain boundary maps, with gray, red, green, blue, yellow, and black lines denoting Σ1, Σ3, Σ9, Σ27, other low-Σ CSL, and random boundaries, respectively. Abundant Σ3 twin boundaries are observed in the microstructure owing to its relatively low stacking fault energy. Quantitative statistical analysis indicates that the fraction of low-Σ CSL boundaries was 50.21%, with 43.32% being Σ3 boundaries and 3.97% consisting of (Σ9 + Σ27) boundaries. Nevertheless, the majority of Σ3 boundaries either extend across an entire grain or end within a single grain, and the random boundary network connectivity in the BM specimen remains largely intact, as seen in Figure 1c. The TRD map for the BM specimen is shown in Figure 1d. Detailed quantitative data reveal that the BM specimen exhibits insufficient development of multiple twinning events, as evidenced by an average TRD size of 16.11 μm and an average number of grains per TRD of 1.91. Many studies [27,28,29] have demonstrated that this nearly complete network can lead to failure during intergranular corrosion and high-temperature applications, where cracks propagate through the well-connected random boundary network throughout the microstructure. Therefore, further GBCD optimization is essential to enhance the intergranular degradation resistance of the BM specimen.

3.2. Annealed Microstructures of SP2215 Steel Under Different Cold Rolling Reductions

Figure 2 displays the IPF, grain boundary reconstruction, and TRD maps for specimens that underwent different cold rolling reductions after annealing at 1100 °C for 10 min. The effect of cold-rolling reduction on the evolution of GBCD after annealing is shown in Figure 3. As observed from the IPF maps depicted in Figure 2a–d, the crystallographic texture of each specimen was characterized by nearly random orientation distribution, which indicates that the formation of low-Σ CSL boundaries exhibits no correlation with texture. The average grain size initially increased from 28.71 μm to 29.99 μm, 32.82 μm, and then decreased to 18.51 μm, while the grain size distribution shifted from unimodal to bimodal and then back to unimodal, as the strain level increased from 3% to 20% (see Figure S1). These phenomena can be attributed to strain-induced boundary migration (SIBM) occurring in low-strain deformed specimens, such as A-1, A-2, and A-3 specimens, while static recrystallization (SRX) was induced in the specimen with a high amount of deformation, specifically the A-4 specimen, during annealing. Specimens A-1 and A-2 exhibited relatively uniform grain growth, as the low stored energy limited the number of grains that grew rapidly via SIBM. For specimen A-3, increasing the stored energy led to a substantial increase in the number of grains undergoing rapid growth driven by SIBM. Consequently, these rapidly growing grains formed large grain clusters during annealing, resulting in a microstructure with a bimodal grain size distribution. Since the stored energy in specimen A-4 exceeded the critical threshold for recrystallization, extensive recrystallization occurred during annealing, resulting in a unimodal grain size distribution.
Figure 2a′,b′,d′ illustrate that in specimens A-1, A-2, and A-4, the majority of Σ3 boundaries, characterized by straight or parallel linear features, were confined within individual grains. Consequently, the microstructure of each specimen displayed a continuous random boundary network. Conversely, specimen A-3 exhibited complex and interconnected Σ3 boundaries morphologies, and the Σ3n (n = 1, 2, 3) boundaries constituted integral components of the random boundary network, as illustrated in Figure 2c′. The fraction of low-Σ CSL boundaries, Σ3 boundaries, and (Σ9 + Σ27) boundaries initially increased within the 3–10% deformation range and then decreased up to a 20% rolling reduction. Furthermore, specimen A-3 exhibited the highest fractions of low-Σ CSL boundaries, Σ3 boundaries, and (Σ9 + Σ27) boundaries (Figure 3a).
Figure 2a″–d″ show the reconstructed TRD maps obtained from TMP specimens subjected to various deformation amounts. It is clear from Figure 2a″–d″ that specimen A-3 exhibited significantly larger TRD size, whereas the other TMP specimens displayed comparatively finer TRD size. Figure 3b presents quantitative statistics for TRD in terms of the average TRD size and the average number of grains per TRD under various specimen conditions. As shown in Figure 3b, both the average TRD size and the average number of grains per TRD initially increased and then decreased with increasing deformation level. The A-3 specimen exhibits the peak values for both the average TRD size and the average number of grains per TRD. Specifically, the average TRD size in the A-3 specimen was approximately 62% larger than in the BM specimen. The average number of grains per TRD for the specimen A-3 (7.28) was significantly higher than that of the BM specimen (1.91). The increase in both the average number of grains per TRD and the TRD size in microstructure results from multiple twinning events during SIBM, which lead to a GBE microstructure.
Previous studies [30,31,32,33] have established that GBCD optimization is more efficiently achieved through recovery than recrystallization, which depends on the stored energy in the microstructure. As mentioned earlier, grain growth was triggered in specimens A-1, A-2, and A-3, while extensive recrystallization took place in specimen A-4 during annealing. Thus, specimens A-1, A-2, and A-3 should possess higher fractions of low-Σ CSL boundaries compared to specimen A-4. This phenomenon is attributed to the activation of SIBM during recovery, which facilitated interactions among pre-existing Σ3n (n = 1, 2, 3) boundaries. These interactions generated multiple twinning events, including Σ3 + Σ3 → Σ9, Σ3 + Σ9 → Σ3, or Σ3 + Σ9 → Σ27. This process serves as the primary mechanism for increasing the fraction of low-Σ CSL boundaries and interrupting the connectivity of the random boundary network.
Given that the grain boundaries within the TRD are linked via Σ3n (n = 1, 2, 3) boundaries, specimen A-3 presented the maximum average TRD size and average number of grains per TRD among all specimens (Figure 3b). However, specimens A-1 and A-2 exhibited relatively low fractions of low-Σ CSL boundaries. This is attributed to the fact that 3% and 7% rolling reductions provide insufficient driving force for long-distance grain boundary migration, thereby suppressing multiple twinning events during annealing. Multiple twinning behind migrating grain boundaries is attributed to the growth of TRDs within the microstructure, driven by SIBM. When SIBM is insufficient, the fraction of low-Σ CSL boundaries is low, and thus the connectivity of the random boundary network is hardly disrupted. Conversely, for specimen A-4, the stored energy exceeds the critical value, triggering static recrystallization during subsequent annealing. Such recrystallization gives rise to new strain-free grains with random boundaries and eliminates the pre-existing Σ3n (n = 1, 2, 3) boundaries. Owing to the preferential normal grain growth of these newly recrystallized grains, the migration distances of random boundaries are relatively limited, which ultimately hinders the formation and interaction of low-Σ CSL boundaries. Therefore, the product microstructure in specimen A-4 yields a lower fraction of low-Σ CSL boundaries, a smaller average TRD size, and a lower average number of grains per TRD, along with an increased connectivity of the random boundary network. These results align with the data reported in the literature [16]. Therefore, precise control of the stored energy is essential during GBE-type TMP to trigger extensive SIBM and suppress recrystallization. Specifically, the stored energy should be maximized to activate sufficient SIBM while remaining below the critical level required for recrystallization.

3.3. Microstructures of Cold-Rolled SP2215 Steel at Different Annealing Treatments

Figure 4 displays the IPD maps, along with the reconstructed grain boundary and TRD maps, for specimens with a 10% reduction that underwent various temperatures for 10 min. Figure 5 illustrates the corresponding characteristics. As shown in Figure 4a–d, the microstructural evolution was highly sensitive to the annealing temperature. The average grain size initially increased, then decreased, as the annealing temperature increased. For the specimens annealed at 1000 °C, 1050 °C, 1100 °C, and 1150 °C for 10 min, the average grain sizes were 17.24 μm, 17.36 μm, 32.82 μm, and 25.39 μm, respectively. Furthermore, the grain size distribution shifted from unimodal at 1000 °C and 1050 °C to bimodal at 1100 °C, then reverted to unimodal at 1150 °C (as shown in Figure S2).
The morphological evolution of Σ3 boundaries can be clearly seen in Figure 4a′–d′. At annealing temperatures of 1000 °C and 1050 °C, the majority of Σ3 boundaries manifested as straight twin pairs, typically regarded as coherent twin boundaries with limited migration ability, which made it difficult to promote grain boundary interactions effectively. As the annealing temperature was elevated to 1100 °C, an increasing fraction of Σ3 boundaries transformed into curved morphologies, signifying the formation of a larger fraction of incoherent Σ3 boundaries. Upon further increasing the annealing temperature to 1150 °C, some incoherent Σ3 boundaries were eliminated. Regarding the GBCD statistics presented in Figure 5a, the fraction of low-Σ CSL boundaries increased sharply from 35.21% to 74.27% as the annealing temperature increased from 1000 °C to 1100 °C. However, the fraction of low-Σ CSL boundaries dropped to 67.30% when the annealing temperature was further increased to 1150 °C.
It is apparent from Figure 4a″–d″ that the average TRD size and the average number of grains per TRD in the B-3 specimen were significantly larger, while the same were relatively finer in other TMP specimens. The average TRD size and the average number of grains per TRD have been determined for specimens with 10% reduction annealed at different annealing temperatures for 10 min, as shown in Figure 5b. The average TRD size and the average number of grains per TRD increased, then decreased, with increasing annealing temperature. Specimen B-3 had the largest average TRD size and the highest number of grains per TRD. A larger TRD size coupled with a higher number of grains per TRD typically signifies a greater extent of multiple twinning within the microstructure. Therefore, it can be inferred that specimen B-3 exhibits outstanding resistance against intergranular failure.
The occurrence of SIBM during TMP is identified as the core process responsible for GBCD optimization. By triggering extensive multiple twinning events, this process forms an abundant fraction of low-Σ CSL boundaries and disrupts the connectivity of the random boundary network. At low annealing temperatures of 1000 °C and 1050 °C, the limited stored energy induced by cold rolling deformation hardly promotes the long-distance migration of grain boundaries, which results in a limited formation of low-Σ CSL boundaries (Figure 5a). Increasing the annealing temperature to 1100 °C enables sufficient SIBM activation in regions with comparatively high stored energy. Subsequently, the migrating grain boundaries sweep through the adjacent deformed microstructure until they come into contact with other migrating grain boundaries. The rapid grain boundary sweep process facilitates the formation of numerous low-Σ CSL boundaries, which can encounter and interact with surrounding grain boundaries, thereby forming new low-Σ CSL boundaries. Accordingly, as illustrated in Figure 4 and Figure 5, the GBCD of specimen B-3 was optimized, accompanied by a larger average TRD size and a higher average number of grains per TRD, and the connectivity of the random boundary network was disrupted. When annealed at 1150 °C, a relatively high number of grains exhibit preferential growth due to the sufficient thermal activation energy to promote rapid grain growth. However, as the number of grains increases, the migration distances of grain boundaries are fairly restricted, which hinders the formation and mutual interaction of low-Σ CSL boundaries. It can explain the slight decline in the fraction of low-Σ CSL boundaries as the annealing temperature was further increased from 1100 °C to 1150 °C (see Figure 5a).

3.4. Creep Properties

To study the effect of GBCD on the creep properties of SP2215 steel, creep tests were performed, and the creep life curves and the corresponding creep test results for BM and GBEM specimens are presented in Figure 6 and Table 3, respectively. In general, creep deformation can be divided into three stages: an initial stage characterized by a decreasing creep rate, a subsequent steady-state stage with a comparatively constant creep rate, and a final accelerated stage with an increasing creep rate, ultimately leading to specimen fracture. As shown in Figure 6 and Table 3, the creep lives of the BM and GBEM specimens were 78 ± 7 h and 103 ± 3 h, respectively. The creep performance of the GBEM specimen was superior to that of the BM specimen. For the BM specimen, the steady-state creep stage lasted nearly 50 ± 5 h, with a relatively high minimum creep rate of (5.12 ± 0.3) × 10−6 s−1. By comparison, the creep curve of the GBEM specimen showed a longer steady-state creep stage that lasted approximately 80 ± 2 h and a lower minimum creep rate of (2.94 ± 0.2) × 10−6 s−1.
Figure 7 shows the different fracture morphologies of BM and GBEM. A relatively flat fracture surface of the BM specimen, characterized by typical cleavage features with river-like patterns, is observed in Figure 7a. However, the fracture surface of the GBEM specimen was uneven and rough, with no obvious flat area (Figure 7c). Further observation of Figure 7b,d reveals that both the BM and GBEM specimens exhibit characteristics of grain contours, with the presence of numerous dimples and creep cavities, indicating that both specimens belong to a mixed fracture mode dominated by intergranular fracture. However, the dimples in the BM specimen were relatively small and shallow, whereas those in the GBEM specimen were larger and deeper. Additionally, the creep cavities in the BM specimen were more numerous and exhibited a chain-like distribution. In contrast, the number of creep cavities in the GBEM specimen was significantly reduced, and they were distributed more evenly.
Figure 8 provides the OM and grain boundary reconstruction maps of crack morphologies near the fracture surface for the BM and GBEM specimens. As shown in Figure 8a,c, intergranular cracks perpendicular to the stress direction were observed in both BM and GBEM specimens after creep tests. However, the cracks formed in the BM sample were longer and wider, whereas those in the GBEM specimen were mostly fine and small. Furthermore, it can be observed from Figure 8b,d that the cracks in both BM and GBEM specimens nucleated primarily at random boundaries and propagated along them under the applied creep stress.
Currently, it is widely believed that stress concentrations caused by the obstruction of intergranular slip at grain-boundary regions during creep exposure are the predominant cause of the nucleation and growth of creep cavities and cracks [34,35,36]. However, the nucleation and growth of cavities and cracks strongly depend on the nature of grain boundaries. For random boundaries, dislocation multiplication and slip induce a pronounced dislocation pile-up, which in turn inhibits further multiplication and eventually deteriorates ductility [37,38]. In contrast, interactions between dislocations and low-Σ CSL boundaries are relatively limited. When dislocations encounter low-Σ CSL boundaries, some dislocations may be arrested, yet the majority can penetrate Σ3 boundaries and slip continuously along them [39,40]. This behavior promotes the propagation of dislocations, ultimately enhancing ductility. Compared with the BM specimens, the GBEM specimen exhibited a higher fraction of low-Σ CSL boundaries, which can effectively alleviate grain-boundary stress concentrations and thereby inhibit the formation of creep cavities and cracks. It is worth noting that having a high fraction of low-Σ CSL boundaries in the microstructure is insufficient to guarantee enhanced creep performance, as the connectivity of the random boundary network also plays a crucial role. Cracks tend to propagate preferentially along a continuous random boundary network. By employing TRD analysis, the connectivity of the random boundary network can be evaluated. Compared with the BM specimen, the GBEM specimen exhibited a larger average TRD size and a higher average number of grains per TRD, indicating a discontinuous random boundary network in the GBEM specimen. The reason lies in the Σ3n (n = 1, 2, 3) mutual misorientations that existed among all grains within the TRD, and all grain boundaries are interconnected to each other, forming a large number of percolation-resistant Σ3n-type triple junctions. As shown in Figure 8, the GBEM specimen exhibits tiny discontinuous cracks while the BM specimen presents coarse cracks. Hence, the inducements of a high fraction of low-Σ CSL boundaries coupled with a significantly fragmented random boundary network in the GBEM specimen contribute to the superior creep performance.

4. Conclusions

(1)
A GBE microstructure in SP2215 steel can be achieved through an optimal TMP consisting of 10% cold rolling deformation followed by annealing at 1100 °C for 10 min. In the GBEM specimen, the fraction of low-Σ CSL boundaries increased markedly from 50.21% (BM) to 74.27%, coupled with substantial disruption to the connectivity of the random boundary network.
(2)
Specimens processed through 10% rolling deformation followed by annealing at the temperature range of 1000 °C to 1100 °C for 10 min led to the occurrence of significant grain growth and an increase in the fraction of low-Σ CSL boundaries. A further increment in annealing temperature up to 1150 °C resulted in a decrease in the fraction of low-Σ CSL boundaries.
(3)
The optimized GBCD contributed to a significant enhancement in the creep property of the alloy. Compared with the BM specimen, the creep duration of the GBEM specimen was increased by 32%, and the minimum creep rate of the GBEM specimen was (2.94 ± 0.2) × 10−6 s−1 compared to (5.12 ± 0.3) × 10−6 s−1 for the BM specimen. The improvement is primarily attributed to a high fraction of low-Σ CSL boundaries induced by GBE treatment, which effectively disrupts the connectivity of the random boundary network, greatly suppressing intergranular crack initiation and propagation.

Supplementary Materials

The following supporting information can be downloaded at: https://www.mdpi.com/article/10.3390/cryst16050282/s1, Figure S1: Grain size distribution of A-1, A-2, A-3 and A-4 specimens; Figure S2: Grain size distribution of B-1, B-2, B-3 and B-4 specimens.

Author Contributions

Conceptualization, T.S. and W.F.; Methodology, T.S. and J.Z.; Validation, T.S. and T.Z.; Formal analysis, T.S. and T.Z.; Investigation, T.S. and J.Z.; Resources, W.F.; Data curation, Z.H.; Writing—original draft preparation, T.S.; Writing—review and editing, T.S. and W.F.; Supervision, W.F.; funding acquisition, W.F. All authors have read and agreed to the published version of the manuscript.

Funding

This research was funded by the Natural Science Foundation of the Jiangsu Higher Education Institutions of China (Grant No.: 18KJB430008), the Open-end Funds of Jiangsu Key Laboratory of Function Control Technology for Advanced Materials (Grant No.: JSKLFCTAM202202), the Research and Practice Innovation for Postgraduate in Jiangsu Ocean University (Grant No.: KYCX2023-06) and Lianyungang Postdoctoral Fund (Grant No.: LYGBSH2025014).

Data Availability Statement

All data is available in the main text and is available upon request.

Conflicts of Interest

The authors declare no conflicts of interest.

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Crystals 16 00282 g001
Figure 1. Microstructure of the BM specimen: (a) IPF map, (b) local misorientation map, (c) grain boundary reconstruction map, and (d) TRD map.
Figure 1. Microstructure of the BM specimen: (a) IPF map, (b) local misorientation map, (c) grain boundary reconstruction map, and (d) TRD map.
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Crystals 16 00282 g002
Figure 2. IPF, grain boundary reconstruction, and TRD maps of (a–a″) A-1, (b–b″) A-2, (c–c″) A-3, (d–d″) A-4 specimens.
Figure 2. IPF, grain boundary reconstruction, and TRD maps of (a–a″) A-1, (b–b″) A-2, (c–c″) A-3, (d–d″) A-4 specimens.
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Crystals 16 00282 g003
Figure 3. (a) The GBCD statistics, (b) average TRD size, and average number of grains per TRD of the TMP specimen with different strain levels.
Figure 3. (a) The GBCD statistics, (b) average TRD size, and average number of grains per TRD of the TMP specimen with different strain levels.
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Crystals 16 00282 g004
Figure 4. IPF, grain boundary reconstruction, and TRD maps of (a–a″) B-1, (b–b″) B-2, (c–c″) B-3, (d–d″) B-4 specimens.
Figure 4. IPF, grain boundary reconstruction, and TRD maps of (a–a″) B-1, (b–b″) B-2, (c–c″) B-3, (d–d″) B-4 specimens.
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Crystals 16 00282 g005
Figure 5. (a) The GBCD statistics, (b) average TRD size, and average number of grains per TRD of the TMP specimen with various annealing temperatures.
Figure 5. (a) The GBCD statistics, (b) average TRD size, and average number of grains per TRD of the TMP specimen with various annealing temperatures.
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Crystals 16 00282 g006
Figure 6. Creep strain versus creep time curves of BM and GBEM specimens.
Figure 6. Creep strain versus creep time curves of BM and GBEM specimens.
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Crystals 16 00282 g007
Figure 7. SEM images of the fracture morphologies: (a,b) BM and (c,d) GBEM specimens.
Figure 7. SEM images of the fracture morphologies: (a,b) BM and (c,d) GBEM specimens.
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Crystals 16 00282 g008
Figure 8. OM and grain boundary reconstruction maps of crack morphologies near the fracture surface: (a,b) BM and (c,d) GBEM specimens.
Figure 8. OM and grain boundary reconstruction maps of crack morphologies near the fracture surface: (a,b) BM and (c,d) GBEM specimens.
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Table 1. Chemical composition of SP2215 heat-resistant steel (wt%).
Table 1. Chemical composition of SP2215 heat-resistant steel (wt%).
CSiMnPSCrNiCuNbMoCoWTiNBFe
0.060.590.650.0160.00222.6715.533.580.510.340.060.0080.0060.410.003Bal.
Table 2. TMP parameters for different GBE conditions.
Table 2. TMP parameters for different GBE conditions.
GroupName IDDeformation/%Annealing Temperature/°CAnnealing Time/min
Type AA-13110010
A-27
A-310
A-420
Type BB-110100010
B-21050
B-31100
B-41150
Table 3. The creep performance of BM and GBEM specimens.
Table 3. The creep performance of BM and GBEM specimens.
SpecimenCreep Time/hSteady-State Creep Time/hMinimum Creep Rate/s−1
BM78 ± 750 ± 5(5.12 ± 0.3) × 10−6
GBEM103 ± 380 ± 2(2.94 ± 0.2) × 10−6
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MDPI and ACS Style

Feng, W.; Sun, T.; Zhao, T.; Zhou, J.; Han, Z. Effect of Thermomechanical Processing on Grain Boundary Character Distribution and Creep Properties of SP2215 Heat-Resistant Steel. Crystals 2026, 16, 282. https://doi.org/10.3390/cryst16050282

AMA Style

Feng W, Sun T, Zhao T, Zhou J, Han Z. Effect of Thermomechanical Processing on Grain Boundary Character Distribution and Creep Properties of SP2215 Heat-Resistant Steel. Crystals. 2026; 16(5):282. https://doi.org/10.3390/cryst16050282

Chicago/Turabian Style

Feng, Wen, Ting Sun, Tianyu Zhao, Junjie Zhou, and Zhengyu Han. 2026. "Effect of Thermomechanical Processing on Grain Boundary Character Distribution and Creep Properties of SP2215 Heat-Resistant Steel" Crystals 16, no. 5: 282. https://doi.org/10.3390/cryst16050282

APA Style

Feng, W., Sun, T., Zhao, T., Zhou, J., & Han, Z. (2026). Effect of Thermomechanical Processing on Grain Boundary Character Distribution and Creep Properties of SP2215 Heat-Resistant Steel. Crystals, 16(5), 282. https://doi.org/10.3390/cryst16050282

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