Effect of carbon on microstructure and mechanical properties of HR3C type heat resistant steels

Journal Pre-proof Effect of carbon on microstructure and mechanical properties of HR3C type heat resistant steels J.M. Bai, Y. Yuan, P. Zhang, J.B. Yan PII:

S0921-5093(20)30034-4

DOI:

https://doi.org/10.1016/j.msea.2020.138943

Reference:

MSA 138943

To appear in:

Materials Science & Engineering A

Received Date: 9 October 2019 Revised Date:

19 December 2019

Accepted Date: 12 January 2020

Please cite this article as: J.M. Bai, Y. Yuan, P. Zhang, J.B. Yan, Effect of carbon on microstructure and mechanical properties of HR3C type heat resistant steels, Materials Science & Engineering A (2020), doi: https://doi.org/10.1016/j.msea.2020.138943. This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. Please note that, during the production process, errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain. © 2020 Published by Elsevier B.V.

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Effect of Carbon on Microstructure and Mechanical Properties of

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HR3C Type Heat Resistant Steels

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J.M. Baia,b,c, Y. Yuana,*, P. Zhanga, J.B. Yana

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a

R&D Center, Xi'an Thermal Power Research Institute Co. Ltd., Xingqing Road 136, Xi'an, Shanxi 710032, China b

School of Materials Science and Engineering, Northeastern University, Shenyang, Liaoning 110819, China c

High Temperature Material Institute, Central Iron and Steel Research Institute, Beijing 100081, China

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Abstract

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The modified HR3C austenitic heat-resistant steels for applications of ultra-

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supercritical (USC) power plants were developed and investigated. As the C content

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breaks through the limitation of the HR3C composition range (>0.1 wt.%), the evolution

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of carbides and mechanical properties after isothermal aging at 700 °C have not been

13

understood. In this study, two modified HR3C with different C content were studied by

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tensile test and Charpy impact test at room temperature after long-term aging up to

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10,000 h. The M23C6 carbide is rapidly precipitated at the interface (GBs, TBs or

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NbC/γ), and these carbides at grain boundaries (GBs) are gradually changed from a

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continuous distribution to a semi-continuous distribution, then finally agglomerate and

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coarsen near the GBs. Moreover, a new morphology type lamellar M23C6 carbide forms

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in the grain. The effect of lamellar carbides on mechanical properties at room

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temperature (RT) is not obvious because the strength of GB is rapidly weakened.

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Besides, the Larson-Miller parameter was obtained and the creep strength of the

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modified HR3C was extrapolated, by conducting creep rupture experiments on un-aged

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sample. The creep tests reveal that the rupture lives decrease with increasing C content.

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Lamellar carbides are precipitated and weaken the strength of grain during high

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temperature creep.

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Keywords: Ultra-supercritical power plant; HR3C; long-term aging; lamellar carbide;

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mechanical properties; creep *

Corresponding author. Tel.: +86 29 8210 2452; fax: +86 29 8210 2090. E-mail address: [email protected]; [email protected]

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1. Introduction

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25Cr-20Ni type heat-resistant steels are considered as a superior oxidation-resistant

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austenite steel at high temperatures due to higher Cr content [1, 2]. It has been

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developed for high temperature applications in USC power plants [3, 4], such as HR3C

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(25Cr-20Ni-Nb-N). The precipitation of M23C6 in heat-resistant steels always plays a

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significant role in mechanical properties. Yang et al. [5] studied the precipitation

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behavior of M23C6 in P92 steel aged at 800 °C, it was revealed that the ripening of the

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M23C6 carbides was controlled by grain boundary (GB) diffusion. The metal elements

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ratio in M23C6 may slightly effect on the coarsening rate [6]. Kaneko et al. [7] and Yan

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et al. [8] proposed a model describing the precipitation of M23C6 and the formation of

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the Cr-depleted zone, and believed that this area promotes the crack propagation. Owing

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to the higher Cr content in 25Cr-20Ni-Nb-N austenite steels, M23C6 carbide is easy to

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precipitate from the matrix during aging [8-11]. The GB is found to be beneficial to the

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nucleation of M23C6 because GB has higher free energy and diffusion rate of atoms [12].

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Most reports [13, 14] focus on the morphology of M23C6 at GB to illustrate the

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change of mechanical properties of alloys. Peng et al. [14] observed the sharp decrease

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in Charpy impact energy of HR3C after aging treatment for 500 h at 700 °C, which is

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attributed to the lower GB hardness with respect to the intragranular hardness. Zieliński

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et al. [15] investigated the microstructures of HR3C after aging at 650 °C, 700 °C and

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750 °C up to 10,000 h, it was found that the discontinuous and continuous of M23C6

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carbides were easily formed at the GB. Wang et al. [16] proposed that the M23C6 had

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different morphologies (semi-continuous and continuous) associated with different

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aging time, and the morphology of M23C6 at GBs affected the ductility of the steels.

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Recently, modified HR3C steels with higher creep strength have been developed for

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applications in 650 °C-class coal-fired power plants [13]. In this case, the maximum

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metal temperature of super-heater and re-heater tubes may reach 700 °C. Usually, the C

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content in HR3C steels is 0.04-0.10 wt.%. As the C content in HR3C type steels is

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further increased beyond 0.10 wt.%, the evolution of carbides and mechanical properties

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after isothermal aging at 700 °C have not been understood. In this study, the

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microstructure and mechanical properties of two modified HR3C steels with different C

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addition after aging at 700 °C up to 10,000 h have been investigated. The effects of C

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content on the microstructure evolution, impact energy, tensile and creep rupture

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properties are discussed.

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2. Experimental Procedure

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2.1 Materials

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The chemical composition in mass percent (wt.%) of experimental austenitic steels

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with different C content are listed in Tab.1, named as high-C content (HC) steel and

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low-C content (LC) steel, respectively. The C content of HC is as high as 0.17%, which

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is about twice that of LC. Co element was added in two steels in order to increase the

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creep strength [17] and inhibit the formation of σ phase [18]. Two test steel ingots, each

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30 kg, were produced in a vacuum induction furnace and hot rolled into plate with a

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thickness of 15 mm. After the solution treatment at 1230 °C for 30 minutes (the grain

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size is about 60 µm).

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Tab. 1 The chemical compositions of modified HR3C steels (wt.%, Fe: Bal.) Alloy HC LC

Ni 18.9 19.2

Cr 23.3 24

Nb 0.57 0.59

Mn 0.03 0.03

Si 0.44 0.31

74 75

2.2 Tensile and Charpy impact tests

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N 0.16 0.18

C 0.17 0.088

Co 4.81 4.76

V 0.1 0.1

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Tensile samples were cut from plate steels after solution treatment, and then aged at

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700 °C for 0, 500, 1000, 3000 and 10, 000 h. The cross-section of fracture is

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schematically shown in Fig. 1(a). The tensile tests were carried out at room temperature

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(RT) with a strain rate of 2.5 × 10-4 s-1. According to the standard ISO 148: 2016[19],

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the Charpy impact tests were performed on the HC and LC at RT. The size of tensile

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samples and impact samples is displayed in Fig. 1(b) and (d).

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2.3 Creep tests

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Creep rupture tests were carried out at 700 °C/200 MPa, 725 °C/120 MPa and

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750 °C/130 MPa, respectively. The experimental conditions were selected in order to

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obtain the trend of the Larson-Miller curve and predict the creep strength of the material

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at the service temperature. All samples are solution treated and un-aged states, the

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longitudinal direction was along the RD direction. Samples were heated from RT to the

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target temperature after 6.5 hours and held at the test temperature for 30 minutes before

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the test. Three sets of thermocouples were attached to the surface of the sample to

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control the temperature. The test is terminated when the sample fractured. Fig 1(c)

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shows the size of the creep sample.

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2.4 Microstructure Characterization

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The microstructure was observed using a Zeiss-Sigma HD scanning electron

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microscope (SEM) and a JEOL-2100 Plus transmission electron microscope (TEM).

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The metallographic samples were manually ground, then etched in a solution consisting

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of 25 ml HCl+5 g CuSO4·5 H2O+25 ml H2O. TEM samples were prepared by a twin-jet

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electro-polisher at below -30 °C and 25 V.

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98 99

Fig. 1 Dimensions of specimens. (a) The position of the specimens; (b) tensile

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specimens; (c) creep specimens; (d) impact specimens.

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3. Results

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3.1 Microstructure evolution during 700 °C long-term aging

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Thermodynamic equilibrium phase diagrams of both HC and LC steels were

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calculated by JMatPro, as shown in Fig. 2. These diagrams indicate that the matrix of

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two steels is a single-phase γ-Fe austenite at 700 °C. Since the C content is nearly

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doubled, the mass fraction of M23C6 in HC steel is twice as high as that of LC. However,

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the mass fraction of the Z phase in two steels is almost the same at 700 °C.

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109 110

Fig. 2 The thermodynamic phase diagrams calculated using JMatPro. (a) and (b) HC

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steel; (c) and (d) LC steel.

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Fig. 3 shows the morphology of primary precipitates in HC and LC after solution

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treatment. Energy dispersive spectrum (EDS) results indicate that these primary

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precipitates are Nb-rich carbonitrides (MX phase), The volume fraction (VF) of MX in

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HC and LC was 1.15% and 0.69%, respectively. HC steel has higher VF of MX than

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LC steel. It indicates that a higher C content promotes the formation of primary MX

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phase. The bulk primary MX was randomly distributed in grains and at GBs. Moreover,

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some annealed twins can be clearly observed within the grains.

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Fig. 3 Microstructure of HC and LC after 1230 °C/30 min + water quenching. (a) HC;

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(b) LC.

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3.1.1 Evolution of carbide at the GB

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Under the service temperature, the precipitates are formed gradually from the

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supersaturated solid solution. The morphology of GBs after aging at 700 °C for 200 h is

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shown in Fig. 4. The precipitates appeared along the GB at the early stage of aging.

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EDS mapping results indicated that the precipitates were Cr-rich M23C6 carbides, which

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were continuously distributed at the GBs. In addition, for HC and LC steels, some

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regular tetragonal M23C6 carbides were found around the GBs, and the size of this type

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M23C6 is larger in HC (about 200~300 nm). This implies that precipitation kinetics of

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M23C6 is accelerated as the C content increases. After aging for 500 h, the carbides at

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the GBs of HC and LC further coarsened, the width of M23C6 at the GBs is about 500

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nm (Fig. 5(a) and (b)). However, the M23C6 gradually spheroidized and semi-

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continuously (chain-like structures) distributed along the GB after aged for 1000 h, as

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shown in Fig. 5(c) and (d), which is consistent with the research of Wang [16]. At the

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same time, from 200 h to 1000 h, the tetragonal M23C6 near the GBs gradually

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combined and coarsened. After aging for 10,000 h, the morphologies of GB carbides of

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HC and LC are shown in Fig. 5(e) and (f). The M23C6 at GBs further coarsened with

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increasing aging time, which combined and agglomerated with carbides near the GBs.

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Fig. 4 Morphology of M23C6 at GBs after aging at 700 °C for 200 h. (a) HC; (b) LC.

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143 144

Fig. 5 Morphology of M23C6 at GBs after aging at 700 °C. HC: (a) 500 h; (c) 1000 h; (e)

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10,000 h. LC: (b) 500 h; (d) 1000 h; (f) 10,000 h.

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3.1.2 Evolution of precipitates in the grain

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As shown in Fig. 6(a) and (b), it can be observed that the precipitates are formed in

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the grains of HC and LC after aging for 200 h. In the high magnification SEM image,

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Fig. 6(c) and (d), some tetragonal M23C6 carbides were found. Moreover, a large

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number of tetragonal M23C6 carbides around the primary MX phase in HC were

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observed, as seen in Fig. 6(e) and (f). It was worth noting that the HC steel began to

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form the lamellar and long-chain shaped precipitates in the grain after aged at 700 °C

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for 1000 h, as shown in Fig. 7(a). Selected area electron diffraction (SAED) and EDS

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confirmed that both precipitates were M23C6 (see Fig. 7(c) and (d)), and such

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morphologies have not been reported before in HR3C. The intragranular long-chain

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shape M23C6 retained a coherent relationship with γ-Fe, i.e. 220

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Therefore, increasing C content results in the complex morphology of M23C6 carbide

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formed in the grain of HR3C type austenitic heat resistant steel. In addition, the Z phase

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was found in the grains of HC and LC, which is consistent with the thermodynamic

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calculation by JMat Pro (Fig. 2).

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// 220

.

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Fig. 6 The M23C6 precipitated in the grains after aged at 700 ° C for 200 h. (a) and (c)

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HC; (b) and (d) LC. The M23C6 around the primary MX phase in HC: (e) SEM image; (f)

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TEM image and inset SAED.

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Fig. 7 The M23C6 precipitated in the grains after aged at 700 °C for 1000 h. (a) HC, (b)

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LC, (c) lamellar M23C6, (d) long-chain shaped M23C6.

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After aging for 3000 h, the microstructures of HC were shown in Fig. 8(a), (b), (c)

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and (d). As the aging time further extended, a large amount of lamellar M23C6 was

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precipitated in the grains. These lamellar M23C6 carbides could be up to 4-5 µm in

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length and only a few tens of nanometers in thickness. In contrast, the lamellar M23C6

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was not found in the grains of LC steel (see Fig. 8(e)), and some long-chain shaped

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M23C6 was observed near the GBs and within the grains (see Fig. 8(f)). After aging for

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10,000 h, the intragranular carbides of HC and LC continued to grow and coarsen, as

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shown in Fig. 8(g) and (h). It can be inferred that increasing C content from 0.088 to

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0.17 wt.% significantly affects the VF and morphology of intragranular M23C6 after

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long term aging at 700 °C.

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Fig. 8 The M23C6 precipitated in the grains after aged at 700 ° C. (a), (b), (c) and (d) are

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HC aged for 3000 h; (e) and (f) are LC aged for 3000 h; (g) and (h) are HC and LC aged

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for 10,000 h, respectively

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3.2 The tensile and impact properties at RT

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Fig. 9 shows the engineering stress-strain curves for the tensile tests of HC and LC at

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RT after long-term aging. The results reveal that the engineering strain at RT of both

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steels decreases with the aging time and two steels have similar tensile properties at RT.

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Fig. 10 shows the variation of strength and elongation of two steels with aging time. It

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can be found that the yield strength (YS) and ultimate tensile strength (UTS) of HC and

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LC slightly increase compared with the solution state (0 h), and the elongation drops

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significantly, from 51% and 48% in solution state to 20% and 29% after aging for 10,

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000 h, respectively.

197 198

Fig. 9 Engineering stress-strain curve of two steels at RT after different aging time.

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199 200

Fig. 10 Yield strength, ultimate tensile strength and elongation of two experimental

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steels at RT after aging at 700 ºC for 0-10000 h.

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Fig. 11 shows the RT Charpy impact energy of two steels aged up to 10,000 h at 700

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°C. In the initial solution state (0 h), the impact energy of HC and LC is almost the

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same, 153 J/cm2 and 155 J/cm2, respectively. The rapid decline in impact energy

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occurred in the early stages of aging, and the impact energy of HC and LC decreased to

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65 J/cm2 and 55 J/cm2 after aging for 200 h. Subsequently, up to 10,000 h, and the

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impact energy of both steels continued to drop slowly. However, the impact energy of

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LC is about 46% higher than that of HC in the later aging period (3000-10,000 h).

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Fig. 11 The impact energy of two experimental steels at RT after aging at 700 ºC for 0-

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10000 h

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3.3 Creep rupture properties

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The creep rupture properties of HC and LC at various conditions are listed in Tab. 2.

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The LC steel has a longer rupture life under the same conditions. At 725 °C/120MPa,

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the rupture life of LC is longer at least 44% than that of HC. The difference in

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elongation of two steels is not obvious after rupture. The Larson-Miller (L-M)

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parameters of the HR3C (Ref. [13]), HC and LC were fitted in Fig. 12. It can be seen

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that HC and LC have better creep rupture properties than HR3C under stresses 175 MPa

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or less. For the HC and LC, serving for 100,000 h at the required stress of 100 MPa, it is

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estimated from the extrapolated results that the operating temperature of LC is increased

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by approximately 12 °C with respect to HC. For serving 100,000 h at 650 °C, the

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rupture strength of LC is 10 MPa higher than that of HC.

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Tab. 2 Creep properties of the modified HR3C steels under different conditions

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Test Condition 700 ºC /200 MPa 725 ºC /120 MPa 750 ºC /130 MPa

224 225 226

Rupture Life (h) HC LC 302 315 2255 3311 399 492

Elongation (%) HC LC 24 22 7 10 11 9

Fig. 12 L-M parameters of the HC, LC and HR3C. 4. Discussion

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According to the SEM observation from Fig. 5-8, it can be found that the evolution of

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M23C6 at the GBs of HC and LC is basically similar, while lamellar carbides are formed

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in grains of HC steel. Schematic diagram Fig. 13 shows the microstructure evolution of

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two steels after long-term aging at 700 °C. It can be determined that increasing the C

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content to exceed the HR3C standard limit (0.1 wt.%), the width and morphology of the

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GB carbides hardly changed during aging, and the new morphology carbides (lamellar

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M23C6) are formed within the grain in high-C modified HR3C steel. In addition, a few

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long-chain M23C6 and tetragonal M23C6 clustered around the primary MX phase were

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also found.

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Fig. 13 The schematic diagram of carbides evolution of HC and LC during aging at 700

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°C.

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The most significant difference in the microstructure of the HC and LC after thermal

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exposure at 700 °C is reflected in the number and morphology of the intragranular

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carbide. Some reports were mentioned that the lamellar M23C6 has been observed at

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GBs or non-coherent twin boundaries in austenitic steels [20-22] or superalloys [23]. It

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is generally believed that the discovered lamellar carbides may grow extensively along

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such non-coherent twin boundaries (TBs), which nucleate on such non-coherent steps

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and grow parallel to the coherent TBs [24]. Terao and Sasmal [25] found the M23C6

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carbide plate was formed on the {110} plane, close to the undissolved NbC particles.

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However, in this study, the lamellar M23C6 precipitated in HC steel was formed

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preferentially in the grain without relying on any interfaces (see Fig. 7 and 8). This

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phenomenon is very rare in austenitic steels, especially in HR3C type steels, which has

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not been reported. This may indicate that as the C content breaks through the limitation

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of the HR3C composition range (>0.1 wt.%), in addition to the M23C6 carbide are

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rapidly precipitated at the interface (GBs, TBs or NbC), a new morphology type of

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lamellar M23C6 form in the grain.

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It is evident that HC and LC exhibit similar mechanical properties at RT after aging

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(see Fig. 10 and 11). The fracture mechanism of the steels is inferred by observing the

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fracture morphology. Fig. 14 summarizes that the tensile fracture morphology at RT

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after various aging time. Clearly, the fracture mode of two steels gradually changed

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from transgranular fracture to intergranular fracture, some intergranular cracks were

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observed in HC and LC after 500 h. Then, the complete intergranular fracture

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characteristics were found in HC and LC after 3000 h. Subsequently, until 10,000 h, the

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fracture mode of two steels remained intergranular fracture. The impact fracture surface

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was observed by SEM, as shown in Fig. 15. The solution treated specimens (0 h)

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present transgranular ductile fracture characteristics with large dimples, which reveals

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excellent toughness performance of HC and LC in solution state. After aging for 200 h,

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the impact value of the two steels decreases sharply (see Fig. 11), and the fractures

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show brittle characteristics with intergranular crack. After aging for 3000 to 10,000 h,

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the micropores on the fracture surface gradually disappear, and the surface of the grains

268

becomes smoother. Therefore, this certifies that the deformation mainly occurs near the

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GBs after aging, and for two steels, the strength of GBs (GBσ) is lower than that of the

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intragranular strength (Gσ) after aging. According to the process of M23C6 evolution at

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the GBs (see Fig. 13), it is known that the coarsened M23C6 directly affects the GBσ at

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RT. The crack propagates along the GBs, and the deformation occurs in the grains is not

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sufficient, which might be the main factor responsible for the plasticity drop of two

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steels at RT after aging treatment. In addition, from the analysis of fracture mode, it is

275

inferred that the lamellar M23C6 in the grains of HC after aging has no obvious influence

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on the mechanical properties at RT.

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277 278

Fig. 14 Fracture surfaces of tensile specimens at RT.

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279 280

Fig. 15 Fracture surfaces of impact tests under RT

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It has been reported [26] that in HR3C the harmful phase (i.e. σ phase) precipitated at

282

GBs during long-term service can deteriorate its RT ductility and toughness. It is

283

assumed that the brittle phase is not precipitated at the GB at the isothermal aging

284

temperature, and the decrease in the strength of GBs at RT is mainly attributed to the

285

precipitation of M23C6. In this study, thermodynamic calculations (see Fig. 2) show that

286

the σ phase may be precipitated in LC, and the mass fractions of σ phase is 2.8 wt.% at

287

700 °C. Actually, σ phase was not observed in two steels aged up to 10,000 h, which

20

288

indicates that the M23C6 precipitated at the GBs is the main reason why the GB at RT is

289

weakened.

290

For the HC and LC subjected to solution treatment (0 h), the M23C6 was not

291

precipitated at the GBs, and there are only randomly distributed primary MX phases in

292

the grains and GBs. As shown in Fig 16(a) and (b), owing to the large size of the bulk

293

primary MX phase, the micropores are generated at the interface during the deformation

294

process. The MX phase was also observed in the dimple of fracture, as shown in Fig. 14

295

and 15. In addition, although both steels are mainly transgranular fractures, a large

296

number of secondary intergranular cracks are found in HC. Moreover, in Fig. 16(c) and

297

(d), lots of slip bands were found in the grains of HC and LC, which demonstrates that

298

significant deformation occurs within the grains. It is seen that the higher C content in

299

the matrix increases the Gσ, which can be confirmed from the higher UST of HC than

300

that of LC, i.e. 756MPa and 728MPa, respectively (see Fig. 10.).

301

302

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303

Fig. 16 The microstructure of two experimental steels in solution state. (a) and (b) are

304

cross-section of HC and LC after tensile test, respectively; (c) and (d) are section of HC

305

and LC after tensile test, respectively.

306

After aging at 700 °C, continuously or semi-continuously distributed M23C6 are

307

precipitated rapidly at the GBs of the two steels. The GBσ is deteriorated, and the

308

fracture morphology is characterized by intergranular fracture, the impact toughness and

309

tensile plasticity of the steels drop sharply at RT as well. Peng et al. [14] believed that

310

the bonding of the M23C6 and the grains is weak and cause the reduction of impact

311

toughness. The grain orientation changes and promotes the grain rotation during the

312

deformation of HR3C at RT after service, which has been reported by Yan et al. [8]. On

313

the other hand, Fig. 17 shows the interaction of Z phase and dislocation during

314

deformation. Some dislocation loops are observed around the Z phase and dislocation

315

lines are strongly curved near the Z phase. This implies that the fine Z phase

316

precipitated in the grains also contributes to enhance the Gσ by pinning the dislocation.

317 318

Fig. 17 TEM images showing the interaction of the Z phase with the dislocations in the

319

grains. The sample was tensile tested at RT after aging for 3000 h. (a) HC and (b) LC.

22

320

The mechanism of high temperature creep rupture is different from that of RT

321

deformation. The fracture surface and microstructure of the creep fracture specimens

322

under the three conditions are observed and summarized in Fig. 18. Lamellar M23C6

323

carbides were found in HC at 725 °C/120MPa and 750 °C/130MPa. Although the

324

fracture modes of HC and LC are intergranular and transgranular mixed fractures, as the

325

lamellar carbide is precipitated, HC is characterized by more transgranular fracture. A

326

large amount of lamellar M23C6 was detected on the fracture surface (HC:

327

750 °C/130MPa). The cross-section observation certifies that the transgranular crack

328

dominates the fracture of HC, and the fracture is located at the interface of M23C6.

329

Creep cavities are nucleated at the interface between the γ-Fe matrix and lamellar M23C6,

330

which cause cracks to grow along the grain. Or cracks propagate along the lamellar

331

M23C6 in the later stages of the rupture. It is clear that the lamellar M23C6 precipitated in

332

HC reduces the Gσ of steels during creep test, which indicates the creep strength of HC

333

is weakened compared to LC. The lamellar M23C6 was not observed in LC under all

334

three experimental conditions. Under conditions of 725 °C/120 MPa and 750 °C/130

335

MPa, in addition to M23C6 precipitated at the GBs, the irregular massive M23C6

336

precipitates concentrated around the GBs of LC. Intergranular fracture characteristics of

337

LC exhibit at 750 °C/130 MPa compared to HC, the back scattered electron (BSE)

338

image of the cross-section of LC clearly shows that the M23C6 mainly precipitated at the

339

GB and around the GB, these carbides are hardly found in the grains. For 700 °C/200

340

MPa fracture specimens, due to the short fracture time, carbides only precipitated at the

341

GBs, and HC and LC mainly exhibit intergranular fracture characteristics.

23

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24

343

Fig. 18 Summarize fracture and cross-section morphologies of HC and LC under

344

various creep conditions

345

4. Summary and Conclusion

346

In this study, the microstructure evolution and mechanical properties of two modified

347

HR3C steels with different C-content after long-term aging were investigated. The

348

following conclusions were obtained:

349

(1) After solution heat treatment at 1230 °C, some residual primary MX phase was

350

found in two steels. The VF of primary MX phase in modified HR3C increases with

351

increasing C content. And these primary MX particles become a source initiating cracks

352

during the deformation at RT.

353

(2) After long-term aging at 700 °C, the continuously distributed M23C6 precipitated

354

at the modified HR3C steel in the early stage of aging (200 h). With extended aging

355

time, the M23C6 gradually spheroidized and semi-continuously distributed at the GB.

356

After 10,000 h, the M23C6 coarsened and concentrated at the GBs. For the modified

357

HR3C steel, increasing the C content to 0.17 wt.% has a slight effect on the carbide

358

morphologies at the GBs. However, the lamellar M23C6 was precipitated only in steels

359

with higher C content.

360

(3) After long-term aging at 700 °C, the tensile elongation and impact energy of two

361

steels drops rapidly at RT, owing to the weakened GBσ by M23C6. The effect of lamellar

362

M23C6 in the grains on mechanical properties at RT is not obvious because the

363

deformation mainly occurs at the GBs.

364

(4) Under various creep conditions, LC has longer rupture lives than HC. Based on

365

the L-M parameters, the extrapolated operating temperature of LC is about 12 °C higher

366

than that of HC for 100,000 h/100 MPa. For the 0.17 wt.% C content steel, under

25

367

725 °C/120MPa and 750 °C/130MPa, the lamellar M23C6 was observed in the fracture

368

surface and grains, and cracks propagated along with the carbide-substrate interface.

369

Acknowledgments

370 371

This work has been financially supported by the China Huaneng Group Co. Ltd.: [grant No. HNKJ18-H08].

372

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CRediT author statement:

Jiaming Bai: Writing- Original draft preparation, Data curation, Methodology, Investigation. Yong Yuan: Conceptualization, Writing-Reviewing and Editing, Funding acquisition, Supervision. Peng Zhang: Resources. Jingbo Yan: Resources.

Declaration of interests ☒ The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. ☐The authors declare the following financial interests/personal relationships which may be considered as potential competing interests: