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Effect of annealing on mechanical properties and microstructure evolution of borated stainless steels
Journal of Nuclear Materials 515 (2019) 206e214
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Journal of Nuclear Materials journal homepage: www.elsevier.com/locate/jnucmat
Effect of annealing on mechanical properties and microstructure evolution of borated stainless steels Chi-Hyoung Won a, b, Jae Hoon Jang a, *, Sung-Dae Kim a, Joonoh Moon a, Heon-Young Ha a, Jun-Yun Kang a, Chang-Hoon Lee a, Tae-Ho Lee a, Namhyun Kang b, ** a b
Ferrous Alloy Department, Korea Institute of Materials Science, Changwon, 51508, Republic of Korea School of Materials Science of Engineering, Pusan National University, Busan, 46241, Republic of Korea
g r a p h i c a l a b s t r a c t
a r t i c l e i n f o
a b s t r a c t
Article history: Received 21 August 2018 Received in revised form 18 December 2018 Accepted 23 December 2018 Available online 26 December 2018
A borated stainless steel based on the composition of type 304 stainless steel containing different amount of boron was fabricated through a conventional ingot metallurgy-hot working process, and the relationship between the microstructure and mechanical properties was investigated. As the content of boron increases, the volume fraction of (Fe,Cr)2B increases, as predicted by thermodynamic equilibrium calculations. As the annealing proceeds at 1180 C, the plate-like shape of (Fe,Cr)2B turned into a spherical shape. In addition, brittle fracture of the boride was dominant before the heat-treatment, but ductile mode due to the formation of dimples in the austenite was dominant after heat-treatment. In the specimen containing 1.78 wt% of boron, the Cr and Ni contents of the austenite matrix were homogenized after annealing at 1180 C for 192 h, and the strength and ductility were simultaneously improved. © 2018 Elsevier B.V. All rights reserved.
Keywords: Borated stainless steel Mechanical properties Microstructure Boride Homogenization
1. Introduction As the use of nuclear power increases, the development of neutron-absorbing structural materials used for the storage and
* Corresponding author. ** Corresponding author. E-mail addresses: [email protected] (J.H. Jang), [email protected] (N. Kang). https://doi.org/10.1016/j.jnucmat.2018.12.039 0022-3115/© 2018 Elsevier B.V. All rights reserved.
transportation of spent nuclear fuel is one of the major issues in the nuclear industry [1]. Spent nuclear fuel, which is inevitably generated in nuclear power plants, releases a large number of thermal neutrons. In order to prevent the release of thermal neutrons to the environment, spent nuclear fuel rods are mostly stored in wet or dry storage containers. It is necessary to achieve high efficiency and density of structure to store more spent nuclear fuel in a limited space. In this respect, materials used in storage containers should have an excellent combination of thermal neutron
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absorbability, corrosion resistance, mechanical properties, processability and weldability. Borated stainless steel (BSS), a candidate material satisfying the requirements, has been used for decades as a thermal neutron absorbing material in transportation casks and storage racks [2,3]. To ensure high neutron absorbing performance, BSS contains 0:2 2:25 wt% of boron (B). Since10 B isotopes, comprising about 20% of natural B, have an excellent neutron absorption cross sectional area of 3840 barn, the addition of natural boron to stainless steel enhances neutron absorbability [4]. The most widely used product among BSS is a type 304 stainless steel to which a certain amount of B is added. ASTM standard A887 covers eight different types of BSS depending on the B content [5]. For each type, the specification contains Grade-A (mainly for powder metallurgyhot working) with fine and uniform boride, and Grade-B (mainly for ingot metallurgy-hot working) with relatively large and heterogeneous boride. For these products, microstructural change, mechanical property, corrosion resistance and impact fracture analyses have been carried out with respect to B content [6e9]. Generally, increasing B content increases thermal neutron absorption capacity, hardness, tensile strength and yield strength, but is known to reduce ductility, impact toughness and corrosion resistance [7,9]. Even for identical amounts of B, the dispersion uniformity and size difference of boride cause differences in ductility and toughness. The mechanical properties of these grades of material are known to be largely influenced by the amount, distribution and shape of precipitates. Products made via powder metallurgy are known to exhibit good mechanical properties and corrosion resistance due to the presence of fine and uniform borides [9]. However, such materials are expensive to produce and have limitations in the fabrication of large-scale products. Therefore, efforts should be made to improve the product mechanical properties through the ingot-hot working process. BSS samples containing B ranging from 0.2 to 2.0 wt% are initially solidified as pro-eutectic austenite during cooling from the liquid phase. Subsequently, the B-enriched liquid phase transforms to an eutectic lamellar mixture consisting of austenite and (Fe,Cr)2B type boride [10]. Since austenite is a ductile phase and a large amount of dispersed boride is brittle intermetallic, the mechanical properties are greatly influenced by the amount and distribution of borides. It is necessary to control the amount and distribution of (Fe,Cr)2B in order to obtain the desired properties of BSS, in particular a product containing a large amount of B in excess of 1 wt %. Previous studies have focused on controlling the boride size via powder metallurgy, not by ingot-hot working process. In this study, we investigated the microstructure of BSS with different B contents induced by ingot-hot working. The effects of B addition on tensile strength and elongation were investigated. Based on the microstructure observation, thermodynamic calculation and diffusion simulation, we analyzed the solidification process of BSS. The heat treatment conditions for the boride phase change were obtained through annealing at 1180 C for different times, and the effect of the annealing process on the tensile strength was also investigated. 2. Calculation method and experimental procedure Three alloys were designed based on type 304 stainless steel composition with different B contents; the chemical composition of the alloys is shown in Table 1. Alloys were labeled B02, B08 and B18 according to B contents. The alloys were fabricated through conventional ingot metallurgy. The 50 kg ingots were reheated at 1150 C for 1 h and hot-rolled into plates with thickness of 40 mm. Then, the as-received specimens were prepared by heat treatment at 1050 C for 30 min, followed by cooling in air. Cylindrical specimens with radius of 10 mm and length of 20 mm were cut and sealed in
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Table 1 Chemical compositions of investigated alloys in wt%. Alloy
C
Si
Mn
Cr
Ni
B
Al
B02 B08 B18
0.056 0.066 0.073
0.25 0.26 0.28
1.35 1.55 1.57
17.7 18.4 18.3
11.9 12.3 12.4
0.19 0.78 1.76
0.024 0.035 0.039
vacuum quartz tubes, which were air-cooled after annealing at 1180 for 48, 96 and 192 h using a box furnace. For the tensile test, a subsize tensile plate specimen of 2 mm thickness was prepared according to ASTM E8 test [11]. Microstructures were observed with optical micrograph (OM) and scanning electron micrograph (SEM) analyses. After polishing with a diamond suspension of 1 mm or less, backscattered electron (BSE) images were obtained using a JEOL IT-3000 SEM together with energy-dispersive X-ray spectroscopy (EDS) in large current mode. The polished specimens were etched in a nitric acid solution (30 ml of nitric acid þ 20 ml of hydrochloric acid þ 50 ml of ethanol), and the microstructure was observed. The density, average radius, and aspect ratio were quantified by Image-Pro Analyzer 7.0 software for more than 1500 precipitates from SEM images [12]. The grain size was measured according to ASTM E11213 [13]. Grain size number (G) was determined using a proportional formula and the average diameter of the corresponding grains was determined. Transmission electron microscopy (TEM) analysis was carried out using a JEOL JEM2100F with an acceleration voltage of 200 kV to examine the crystallography and morphology of borides. Specimens for TEM observation were prepared from mechanical-polished samples using focused-ion-beam micro-machining in combination with a lift-out method from the as-received specimen. The selected-area diffraction patterns obtained from the precipitates were analyzed to identify the crystal structure. For the analysis of the composition in the microstructure, EDS analysis was carried out at 7 points and averaged. In order to identify the distribution of alloying elements in the austenitic matrix, electron-probe micro-analyzer (EPMA) observation of specimens for as-received state and after 192 h of annealing were performed. EPMA analysis was performed using a JXA-8530F highresolution electron microscope, with an acceleration voltage of 15 kV, a resolution of 300 225 points and a dwell time of 20 msec. B, Ni and Cr were investigated, and fine compositional changes were observed by five overlapping analyses. Three tensile tests were performed on the as-received state samples and annealed specimen with the heat-treatment condition described above. Tensile tests were carried out at room temperature and at 2 mm/min using a 10-ton INSTRON 5882 tester and a 25 mm gauge length extensometer. SEM was used to examine the fracture surface and its surrounding area after the tensile test. Using ThermoCalc software, the thermodynamic equilibrium of the system was calculated based on the TCFE9.0 database [14]. The variation of the concentration profile of Cr and Ni during the annealing was calculated based on the MobFe4 database using multicomponent diffusional transformation (DICTRA) software [15]. The austenite phase is assumed to be a matrix, and a total length of 200 mm was simulated based on 200 points. The right half of the specimen is assumed to have a composition of pro-eutectic austenite, and the left half is assumed to be eutectic austenite. C
3. Results The equilibrium phase diagram of the BSS was calculated using ThermoCalc software, and is shown in Fig. 1. The figure shows phase
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Fig. 1. Equilibrium phase diagram for borated stainless steel based on TCFE9.0 database. Liquid, FCC, BCC and orthorhombic M2B phases are allowed in calculation. Diagram was established based on 18Cr-12Ni-1.5Mn wt%, corresponding to the composition of type 304 stainless steel. Blank squares and triangles show measured liquidus and solidus temperatures based on 20Cr-14Ni-1.8Mn wt% system [3].
change according to B content based on 18Cr-12Ni-1.5Mn wt%, which corresponds to the composition of type 304 stainless steel. Liquid, austenite (FCC), and ferrite (BCC) phases were allowed in the calculations. The (Fe,Cr)2B with orthorhombic structure was assumed to be the boride; its crystal structure was confirmed by TEM observation. The phase diagram shows reasonably good agreement with the measured solidus and liquidus temperatures of the previously reported study [3]. Differences in the measurements
and calculations for liquidus temperature are mainly due to different Cr and Ni contents. The amount of B corresponding to the eutectic point is 2.15 wt% and the temperature is 1256 C. Thus, the three alloys (B02, B08, B18) prepared were included in the hypoeutectic composition and a certain amount of pro-eutectic austenite was formed first during solidification. After pro-eutectic austenite formation, from the B enriched liquid, a eutectic mixture of austenite and (Fe,Cr)2B is expected to be generated. Figs. 2 and 3 show microstructures obtained using OM and SEM of the as-received specimens and those after annealing at 1180 C for 192 h, respectively. In Fig. 2 (a), (c) and (e), it can be seen that all three specimens are mixtures of pro-eutectic austenite, which is generated first during the cooling process in the liquid and eutectic microstructure. The eutectic microstructure generated from the B enriched liquid has an interconnected network structure. Table 2 summarizes the precipitate fraction, grain size, average size and aspect ratio of precipitates in each specimen. As the content of B increases, the volume fractions of boride and eutectic microstructure increase almost linearly and the fraction of pro-eutectic austenite decreases. After annealing, austenite grain size, the average size and aspect ratio of the precipitate changes significantly compared to those of the as-received specimens. During annealing at 1180 C, the plate shape (Fe,Cr)2B changes to a spherical shape. In addition, the average size also increases through the coarsening process, thereby decreasing the density of the precipitates and increasing the average distance between the precipitates. The pinning effect of precipitates leads to a significant difference in the austenite grain growth kinetics in each specimen. The
Fig. 2. Optical micrographs of (a) as-received B02, (b) annealed B02, (c) as-received B08, (d) annealed B08, (e) as-received B18, and (f) annealed B18 specimen. Annealing condition of specimen is 192 h at 1180 C.
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Fig. 3. Scanning electron micrographs of (a) as-received B02, (b) annealed B02, (c) as-received B08, (d) annealed B08, (e) as-received B18 and (f) annealed B18 specimen. Annealing condition of specimen is 192 h at 1180 C.
annealed B02 specimen has a small fraction and number of precipitates, such that austenite grain growth continues to occur and results in a grain size 5 times larger than that of the as-received specimen. This grain size affects the strength and elongation, which will be discussed later. On the contrary, the higher the content of B is, the larger the fraction and number of precipitates will be; due to the pinning effect, the grain size will also be smaller than that of the annealed B02 specimen. Characterization of the boride observed as the dark region in Fig. 3 was carried out using TEM. Several studies have been conducted on the boride crystal structure and the electronic structure of transition metal boride [16,17]. Havinga et al. reported that borides of Cr, Fe, Mn and Ni have tetragonal cell structures with a space group of I4/mcm [17]. The existence of orthorhombic Cr2B with space group Fddd was reported by Kotzott et al.; the refined cell constants were a ¼ 4:2752 Å, b ¼ 7:4523 Å and c ¼ 14:795 Å
[18]. Fig. 4 shows a bright field image and selected area diffraction pattern of (Fe,Cr)2B. The results of EDS measurements confirmed that these materials are borides containing large amounts of Cr and Fe. The diffraction pattern shows that the crystal structure of borides was confirmed to be orthorhombic and belongs to the space group Fddd, with a lattice parameter similar to reported values for Cr2B. Fig. 5 provides an engineering strain (ε)-stress (s) curves for the tensile tests. B02, B08 and B18 exhibited similar tensile strengths in a range from 550 to 580 MPa in the as-received state. However, since the volume fraction of (Fe,Cr)2B increases as the B content increases, the elongation dropped from 38% for the as-received B02 specimen to 5% for the as-received B18 specimen. Compared with the austenite matrix of metallic bonding, (Fe,Cr)2B with covalent bonding is brittle. A eutectic structure of (Fe,Cr)2B and austenite is formed from the B enriched liquid and eutectic structures are
Table 2 Phase fraction, mean precipitate size and aspect ratio of (Fe,Cr)2B and average austenite grain size of as-received specimens and specimens annealed for 192 h at 1180 C for B02, B08 and B18 alloys, respectively. Alloy
B02 B08 B18
precipitate size / mm
phase fraction / %
grain size / mm
aspect ratio
as-received
annealed
as-received
annealed
as-received
annealed
as-received
annealed
2.26 10.30 22.49
2.68 11.85 23.04
1.889 1.938 2.813
2.409 3.117 4.554
2.979 3.179 2.761
1.818 1.741 1.972
266 177 191
1449 363 293
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Fig. 4. Transmission electron micrograph of (Fe,Cr)2B precipitates: (a) Bright field image. Selected area diffraction patterns with (b) z ¼ [201] zone axis and (c) z ¼ [110] zone axis.
Fig. 5. Engineering stress (s)-strain (ε) curve for as-received and annealed specimens at 1180 C for (a) B02, (b) B08 and (c) B18 alloy. The black line is as-received specimen; the red, purple and blue lines represent specimens annealed for 48 h, 96 h and 192 h, respectively. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
connected to each other as shown in Figs. 2 and 3. Therefore, cracks initiated in the (Fe,Cr)2B or interface between (Fe,Cr)2B and austenite propagate easily along the eutectic structure. As the fraction of the eutectic structure increases, the ductility decreases dramatically. The mechanical properties of the alloys change markedly through annealing at 1180 C. In the B02 alloy, the elongation increases from 37% for the as-received specimen to over 62e66%. The elongations of the B08 and B18 specimens also increase with annealing, increasing from 24% to 39% for the B08 specimen and from 5% to 16% for the B18 specimen. The tensile strength decreases from 559 MPa to 541 MPa due to a decrease in the grain refining strengthening effect for the B02 specimen. On the contrary, in the case of the B08 specimen, the strength tends to be maintained; in
the case of the B18 specimen, the strength steadily increases from 534 MPa to 636 MPa. The work-hardening mechanism may be analyzed through the relationship between true-strain (εT ¼ lnð1 þ εÞ), true-stress (sT ¼ sð1 þ εÞ) and work hardening rate (ddεsTT ), as shown in Fig. 6. The work hardening rate can be expressed as an engineering strainstress derivative (ddεs) using the following equation
dsT ds dε d½sð1 þ εÞ dε ds dε ,ð1 þ εÞ þ s ¼ T ¼ ¼ : dεT dε dεT dεT dεT dε dε (1) In the plastic instability state in which local necking initiates during tensile testing, ddεs ¼ 0 because flow-stress is maximized.
Fig. 6. True stress (s)-strain (ε) curve and work hardening rate for as-received and annealed specimens at 1180 C for (a) B02, (b) B08 and (c) B18 alloy. The black line is as-received specimen; the red, purple and blue lines represent specimens annealed for 48 h, 96 h and 192 h, respectively. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
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Fig. 7. Microstructure near fracture surface of (a) as-received specimen and (b) annealed specimen for B18 alloy. Fracture surface of (c) as-received specimen and (d) annealed specimen for B18 alloy. Annealing condition of specimen is 192 h at 1180 C.
Therefore, the plastic instability condition is redefined as follows from the true-strain and true-stress relationship, which is Conre's criterion [19]. side
dsT dε ¼s ¼ sð1 þ εÞ ¼ sT : dεT dεT
(2)
For the B02 specimen, work hardening rate and true stress are encountered at about 40% strain for both as-received and postannealing samples. Because, compared to other alloys, the distribution of (Fe,Cr)2B is small, the curves are similar to that of austenitic stainless steel. As annealing progresses, austenite grain growth and precipitate density decrease, resulting in a decrease in work hardening rate and lower tensile strength. In the case of the B08 specimen, there is no intersection between work-hardening rate and true stress curve in the as-received state, so fracturing takes place before local necking occurs. However, after annealing, more strain is absorbed in the austenite matrix, local necking happens and elongation improvement can be obtained. In the case of the B18 specimen, both the as-received state and the annealing state fracture before reaching the plastic instability condition. Compared to the as-received state, the annealed state shows greatly improved elongation. Nevertheless, due to the fraction and distribution of (Fe,Cr)2B, fracturing of the specimen occurs due to cracks initiated from the particles before the austenite matrix absorbs the deformation. The interesting part of tensile testing of the B18 specimen is that, as the annealing progresses, both strength and elongation are improved. Comparing the results of the 48 h annealing to those of 192 h annealing, the work-hardening rate increases with the heat treatment time, tensile strength is improved by about 65 MPa and elongation can be increased from 14.6 to 16.4%. The simultaneous increase of the strength and elongation of the B18 specimen was attributed to the change in the concentration of the matrix; this was confirmed by EPMA analysis mentioned later. Fig. 7 shows the fracture surface and microstructure around the fracture surface after tensile testing of the B18 specimen. In the case
of the as-received specimen, most of the (Fe,Cr)2B precipitate has a plate-like shape with the austenite matrix, and the interface between the precipitate and matrix or precipitate itself is the crack initiation site. Cracks easily propagate along the interface; cracks easily connect to each other due to the short distance to surrounding precipitates. In the fracture surface observation shown in Fig. 7(c), the main fracture mechanism is brittle fracture inside the (Fe,Cr)2B precipitate. Even in the specimen after annealing, initial crack is also generated in the precipitate. However, unlike the asreceived specimen, the shape of the precipitate changes to spherical during the annealing, and the average precipitate size increases, such that the relative distance to adjacent cracks is increases. Therefore, a ductile fracture with dimple shape appears in the austenite matrix. The ductility improvement of the specimen after the annealing described above is due to this change in the fracture mechanism. The compositional distributions of pro-eutectic and eutectic austenite in the B18 specimen were confirmed using EPMA measurement. Fig. 8 shows the distributions of B, Ni and Cr of the proeutectic and eutectic areas of as-received B18 specimens. It can be confirmed that B and Cr are contained in the (Fe,Cr)2B precipitates, and Ni is barely present. In the austenite of the pro-eutectic and eutectic regions of the as-received specimens, the differences in concentration of Cr and Ni can be confirmed. Table 3 summarizes the results of analyzing the concentrations of each microstructure through EDS measurement. As-received specimens showed 12.8 wt % Cr and 14.9 wt% Ni in the pro-eutectic austenite region. In the eutectic region, the concentration of Cr was 2.2 wt% lower and that of Ni was 0.9 wt% higher. Fig. 9 shows the elemental distribution of the B18 specimen after annealing, indicating a uniform distribution throughout the austenite. The concentrations of regions are well matched with the thermodynamic calculations and show that amount of Cr is about 1.8 wt% lower (2.1 wt% lower in thermodynamic calculation) and amount of Ni is 1.2 wt% higher (1.5 wt% higher in thermodynamic calculation) compared to the proeutectic austenite in the as-received specimen.
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Fig. 8. EPMA observation showing distribution of Ni, Cr and B for as-received specimen of B18 alloy.
Table 3 EDS observations of each microstructure for as-received specimen and specimen annealed for 192 h at 1180 C for B18 alloy. Pro-eutectic austenite Contents / wt% As-received Annealed
Cr 12:8±0:6 e
Eutectic austenite Ni 14:9±0:6 e
Cr 10:6±0:2 11:0±0:1
(Fe,Cr)2B boride Ni 15:8±0:6 16:0±0:4
Cr 43:4±1:0 47:1±1:1
Fe 43:7±1:0 40:4±0:7
Fig. 9. EPMA observation showing distribution of Ni, Cr and B for specimen of B18 alloy annealed for 192 h at 1180 C.
4. Discussion The mechanical properties of BSS are mainly related to the shape and distribution of (Fe,Cr)2B. The composition corresponding
to the eutectic point of BSS is about 2.15 wt% B, and the fraction of the eutectic structure is proportional to the amount of B added. In addition, because it is transformed from B enriched liquid, the eutectic structure encloses pro-eutectic austenite. During the
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Fig. 10. Thermodynamic calculation results of (a) equilibrium phase fraction (b) alloy contents in austenite and (c) alloy contents in M2B phases, as a function of temperature for the B18 alloy. In (b) and (c), black, red, blue and green lines represent Fe, Cr, Ni and Mn contents, respectively. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
deformation process, initial cracks propagate primarily in the (Fe,Cr)2B or interface with the austenite; they are then connected to neighboring cracks. The plate-shape (Fe,Cr)2B produced by the eutectic process has a larger aspect ratio than the spherical particles and thus has a longer crack propagation length. When the aspect ratio is n times in the same volume of precipitate, the maximum propagation length in the precipitate can be increased nð2=3Þ times. During annealing at 1180 C, the plate-like (Fe,Cr)2B changes to a spherical shape, the aspect ratio decreases, the average precipitate size increases, the precipitate density decreases and the distance between the precipitate increases. As a result, dimple shape fracture surfaces in austenite are mainly observed in specimens after annealing, whereas the major fracture mechanism in as-received specimens is brittle fracture in the precipitate. Another important contributor to the mechanical properties of BSS steels is the compositional inhomogeneity of the austenite matrix. BSS has a hypoeutectic composition, and the matrix consists of, first, pro-eutectic austenite made in liquid during cooling and also eutectic austenite made from B enriched liquid. Fig. 10 shows the calculated equilibrium composition of each phase and phase fraction as a function of temperature for the B18 alloy. The concentrations of Cr and Ni in the pro-eutectic austenite at 1260 C are similar at about 11 wt%, but the concentration of Cr is lower and that of Ni is higher in eutectic austenite formed near 1200 C. These results match well with the inhomogeneity of Cr and Ni in the EPMA observation of an as-received specimen. After annealing, the equilibrium calculation results at 1180 C indicate that the
composition of the entire austenite region changes to high Ni and low Cr. Fig. 11 shows DICTRA simulation results of the homogenization of Cr and Ni in austenite over time. The as-received specimens of B18 alloy were simulated by assuming a composition of proeutectic and eutectic austenite. During the homogenization at 1180 C, there is a certain concentration gradient of Cr and Ni at 48 h and 96 h, and there is a difference in homogeneity during the heat treatment up to 192 h. As shown in Fig. 10(c), the equilibrium concentration of Cr in (Fe,Cr)2B at 1180 C is higher than that at eutectic temperature; this is also consistent with the EDS measurements shown in Table 3. Therefore, in the initial stage of annealing, elemental redistribution between austenite and (Fe,Cr)2B in the eutectic structure takes place before homogenization in the matrix. In order to maintain the local equilibrium in the eutectic region, Cr of austenite is partitioned into (Fe,Cr)2B, and the concentration of Cr is lower in eutectic austenite. When equilibrium is assumed in the B enriched eutectic region, the calculated equilibrium Cr concentration in the austenite at 1180 C is about 10.2 wt%, which is about 1.3 wt% lower than that of orthoequilibrium. That is, in the initial stage of annealing, before homogenization sufficiently progresses, partitioning of Cr from austenite to (Fe,Cr)2B occurs in the eutectic region, so that the concentration of Cr in the adjacent austenite drops. The effects of Cr and Ni on the strength and ductility of austenitic stainless steel have been studied in various ways [20e22]. Addition of Ni is known to increase toughness, but the
Fig. 11. Diffusion profile simulated using multicomponent diffusion transformation simulation for (a) Cr contents and (b) Ni contents at different annealing times. Black is initial profile. Red, blue and green lines represent profiles for the specimen after annealing for 48 h, 96 h and 192 h at 1180 C. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
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exact effect of strengthening is not clear. Depending on the alloy composition, strengthening and softening are observed with Ni addition, but the effect is smaller than those of other alloying elements [22,23]. In contrast, the addition of Cr is reported to improve work-hardening rate and strengthening effect by lowering the stacking fault energy [20,22]. In the early stage of the annealing process, the eutectic austenite is relatively soft because it contains less Cr than the pro-eutectic austenite region. Therefore, it mainly deforms in the eutectic austenite region during the tensile test. When annealing proceeds for a sufficient time at 1180 C, homogenization occurs between the pro-eutectic and eutectic austenite. In addition, the Cr concentration in the eutectic austenite region increases due to Cr migration from the pro-eutectic austenite to the eutectic austenite. As a result, homogeneous heat treatment of BSS can improve both the strength and ductility. Changes in size and shape of the boride can affect the neutron absorbing properties as well as the mechanical properties. The thermal neutron absorption cross section of boron is very large, so there could be a self-shielding effect that would increase as the size of boride increases. Strong absorbing elements could result in significant neutron flux depression in the materials, and lower effective neutron absorbability [24]. The thermal neutron mean free path length (l) in (Fe,Cr)2B can be approximated as the reciprocal of the macroscopic cross section (S),
S ¼ NB10 ,sB10
(3)
NB10 ¼ NA XB10 =Vm
(4)
where NB10 is the number of 10B isotopes per unit volume, sB10 ¼ 3840 barn is the microscopic cross section, NA is Avogadro's number, XB10 ¼ 1=3 0:198 is the mole fraction of 10B in (Fe,Cr)2B, and Vm ¼ 7:10 m3 =mol is the molar volume of (Fe,Cr)2B. Since the calculated mean free path length in boride is about 465 mm, which is much larger than the observed average boride size of 2.4e4.5 mm, the decrease in neutron absorbability due to self-shielding is expected to be insignificant. 5. Conclusions The evolution of the microstructure and mechanical properties during the annealing were analyzed for borated stainless steel with three different levels of B content. The three alloys, containing up to 1.8 wt% B, belong to the hypoeutectic composition, and have microstructures of pro-eutectic austenite and eutectic structure in asreceived state. As the content of B increases, the volume fraction of the eutectic structure, which is a mixture of austenite and (Fe,Cr)2B, increases and the mechanical properties degrade significantly. Plate-like (Fe,Cr)2B precipitates change to spherical shape during annealing at 1180 C and ductility improves. In the case of an alloy containing about 1.8 wt% B, strength and ductility improved simultaneously as the size of the precipitate increased and the Cr content of the austenite matrix became uniform. Therefore, sufficient annealing time is required to obtain improved mechanical properties of borated stainless steel containing a large amount of B. Acknowledgements This work was funded by the Fundamental R&D Program of the
Korea Institute of Materials Science, Grant No. PNK5850. This work was also supported by Basic Science Research Program through the National Research Foundation of Korea (NRF-2016R1C1B1011593) funded by the Ministry of Science, ICT and Future Planning. This study was also supported by the Ministry of Trade, Industry & Energy (MI, Republic of Korea) under Strategic Core Materials Technology Development Program (No. 10067375).
Appendix A. Supplementary data Supplementary data to this article can be found online at https://doi.org/10.1016/j.jnucmat.2018.12.039.
References [1] A. Machiels, R. Lambert, Handbook on Neutron Absorber Materials for Spent Nuclear Fuel Applications, Electric Power Research Institute, 2005. [2] S. Soliman, D. Youchison, A. Baratta, T. Balliett, Neutron effects on borated stainless steel, Nucl. Technol. 96 (3) (1991) 346e352. [3] C. Robino, M. Cieslak, High-temperature metallurgy of advanced borated stainless steels, Metall. Mater. Trans. 26 (7) (1995) 1673e1685. [4] Y. Choi, B.M. Moon, D.-S. Sohn, Fabrication of Gd containing duplex stainless steel sheet for neutron absorbing structural materials, Nucl. Eng. Technol. 45 (5) (2013) 689e694. [5] A. A887-89, Standard Specification for Borated Stainless Steel Plate, Sheet, and Strip for Nuclear Application, Tech. Rep., ASTM, 2014. [6] E.A. Loria, H.S. Isaacs, Type 304 stainless steel with 0.5% boron for storage of spent nuclear fuel, JOM 32 (12) (1980) 10e17. [7] J. Dilip, G.J. Ram, Microstructures and properties of friction freeform fabricated borated stainless steel, J. Mater. Eng. Perform. 22 (10) (2013) 3034e3042. [8] J. He, S.E. Soliman, A.J. Baratta, T.A. Balliett, Fracture mechanism of borated stainless steel, Nucl. Technol. 130 (2) (2000) 218e225. [9] X. He, T. Ahn, T. Sippel, et al., Corrosion of Borated Stainless Steel in Water and Humid Air, CORROSION, 2012. [10] H. Goldschmidt, Effect of boron additions to austenitic stainless steels, J. Iron Steel Inst. 209 (11) (1971) 900e909. [11] A. E8/E8M-16a, Standard Test Methods for Tension Testing of Metallic Materials, Tech. Rep., ASTM, 2016. [12] M. Cybernetics, Image-Pro Analyzer 7.0, http://www.mediacy.com/ imageproplus. [13] A. E112-13, Standard Test Methods for Determining Average Grain Size, Tech. Rep., ASTM, 2016. €glund, P. Shi, B. Sundman, Thermo-Calc & [14] J.-O. Andersson, T. Helander, L. Ho DICTRA, computational tools for materials science, Calphad 26 (2) (2002) 273e312. €glund, J. Ågren, A. Engstro €m, DICTRA, a tool for simulation [15] A. Borgenstam, L. Ho of diffusional transformations in alloys, J. Phase Equil. 21 (3) (2000) 269. [16] C. Zhou, J. Xing, B. Xiao, J. Feng, X. Xie, Y. Chen, First principles study on the structural properties and electronic structure of X2B (X¼ Cr, Mn, Fe, Co, Ni, Mo and W) compounds, Comput. Mater. Sci. 44 (4) (2009) 1056e1064. [17] E. Havinga, H. Damsma, P. Hokkeling, Compounds and pseudo-binary alloys with the CuAl2 (C16)-type structure I. Preparation and X-ray results, J. Less Common Met. 27 (2) (1972) 169e186. [18] D. Kotzott, M. Ade, H. Hillebrecht, Synthesis and crystal structures of a-and bmodifications of Cr2IrB2 containing 4-membered B4 chain fragments, the tboride Cr7. 9Ir14. 1B6 and orthorhombic Cr2B, Solid State Sci. 10 (3) (2008) 291e302. [19] R. E. Reed-Hill, R. Abbaschian, R. Abbaschian, Physical Metallurgy Principles. [20] L. Vitos, J.-O. Nilsson, B. Johansson, Alloying effects on the stacking fault energy in austenitic stainless steels from first-principles theory, Acta Mater. 54 (14) (2006) 3821e3826. [21] L. Vitos, P.A. Korzhavyi, B. Johansson, Stainless steel optimization from quantum mechanical calculations, Nat. Mater. 2 (1) (2003) 25. [22] D. Llewellyn, Work hardening effects in austenitic stainless steels, Mater. Sci. Technol. 13 (5) (1997) 389e400. € m, Proof strength values for austenitic stainless steels at [23] J. Eliasson, R. Sandstro elevated temperatures, Steel Res. 71 (6e7) (2000) 249e254. [24] J. Copley, Scattering effects within an absorbing sphere immersed in a field of neutrons, Nucl. Instrum. Methods Phys. Res. Sect. A Accel. Spectrom. Detect. Assoc. Equip. 307 (2e3) (1991) 389e397.
Contents lists available at ScienceDirect
Journal of Nuclear Materials journal homepage: www.elsevier.com/locate/jnucmat
Effect of annealing on mechanical properties and microstructure evolution of borated stainless steels Chi-Hyoung Won a, b, Jae Hoon Jang a, *, Sung-Dae Kim a, Joonoh Moon a, Heon-Young Ha a, Jun-Yun Kang a, Chang-Hoon Lee a, Tae-Ho Lee a, Namhyun Kang b, ** a b
Ferrous Alloy Department, Korea Institute of Materials Science, Changwon, 51508, Republic of Korea School of Materials Science of Engineering, Pusan National University, Busan, 46241, Republic of Korea
g r a p h i c a l a b s t r a c t
a r t i c l e i n f o
a b s t r a c t
Article history: Received 21 August 2018 Received in revised form 18 December 2018 Accepted 23 December 2018 Available online 26 December 2018
A borated stainless steel based on the composition of type 304 stainless steel containing different amount of boron was fabricated through a conventional ingot metallurgy-hot working process, and the relationship between the microstructure and mechanical properties was investigated. As the content of boron increases, the volume fraction of (Fe,Cr)2B increases, as predicted by thermodynamic equilibrium calculations. As the annealing proceeds at 1180 C, the plate-like shape of (Fe,Cr)2B turned into a spherical shape. In addition, brittle fracture of the boride was dominant before the heat-treatment, but ductile mode due to the formation of dimples in the austenite was dominant after heat-treatment. In the specimen containing 1.78 wt% of boron, the Cr and Ni contents of the austenite matrix were homogenized after annealing at 1180 C for 192 h, and the strength and ductility were simultaneously improved. © 2018 Elsevier B.V. All rights reserved.
Keywords: Borated stainless steel Mechanical properties Microstructure Boride Homogenization
1. Introduction As the use of nuclear power increases, the development of neutron-absorbing structural materials used for the storage and
* Corresponding author. ** Corresponding author. E-mail addresses: [email protected] (J.H. Jang), [email protected] (N. Kang). https://doi.org/10.1016/j.jnucmat.2018.12.039 0022-3115/© 2018 Elsevier B.V. All rights reserved.
transportation of spent nuclear fuel is one of the major issues in the nuclear industry [1]. Spent nuclear fuel, which is inevitably generated in nuclear power plants, releases a large number of thermal neutrons. In order to prevent the release of thermal neutrons to the environment, spent nuclear fuel rods are mostly stored in wet or dry storage containers. It is necessary to achieve high efficiency and density of structure to store more spent nuclear fuel in a limited space. In this respect, materials used in storage containers should have an excellent combination of thermal neutron
C.-H. Won et al. / Journal of Nuclear Materials 515 (2019) 206e214
absorbability, corrosion resistance, mechanical properties, processability and weldability. Borated stainless steel (BSS), a candidate material satisfying the requirements, has been used for decades as a thermal neutron absorbing material in transportation casks and storage racks [2,3]. To ensure high neutron absorbing performance, BSS contains 0:2 2:25 wt% of boron (B). Since10 B isotopes, comprising about 20% of natural B, have an excellent neutron absorption cross sectional area of 3840 barn, the addition of natural boron to stainless steel enhances neutron absorbability [4]. The most widely used product among BSS is a type 304 stainless steel to which a certain amount of B is added. ASTM standard A887 covers eight different types of BSS depending on the B content [5]. For each type, the specification contains Grade-A (mainly for powder metallurgyhot working) with fine and uniform boride, and Grade-B (mainly for ingot metallurgy-hot working) with relatively large and heterogeneous boride. For these products, microstructural change, mechanical property, corrosion resistance and impact fracture analyses have been carried out with respect to B content [6e9]. Generally, increasing B content increases thermal neutron absorption capacity, hardness, tensile strength and yield strength, but is known to reduce ductility, impact toughness and corrosion resistance [7,9]. Even for identical amounts of B, the dispersion uniformity and size difference of boride cause differences in ductility and toughness. The mechanical properties of these grades of material are known to be largely influenced by the amount, distribution and shape of precipitates. Products made via powder metallurgy are known to exhibit good mechanical properties and corrosion resistance due to the presence of fine and uniform borides [9]. However, such materials are expensive to produce and have limitations in the fabrication of large-scale products. Therefore, efforts should be made to improve the product mechanical properties through the ingot-hot working process. BSS samples containing B ranging from 0.2 to 2.0 wt% are initially solidified as pro-eutectic austenite during cooling from the liquid phase. Subsequently, the B-enriched liquid phase transforms to an eutectic lamellar mixture consisting of austenite and (Fe,Cr)2B type boride [10]. Since austenite is a ductile phase and a large amount of dispersed boride is brittle intermetallic, the mechanical properties are greatly influenced by the amount and distribution of borides. It is necessary to control the amount and distribution of (Fe,Cr)2B in order to obtain the desired properties of BSS, in particular a product containing a large amount of B in excess of 1 wt %. Previous studies have focused on controlling the boride size via powder metallurgy, not by ingot-hot working process. In this study, we investigated the microstructure of BSS with different B contents induced by ingot-hot working. The effects of B addition on tensile strength and elongation were investigated. Based on the microstructure observation, thermodynamic calculation and diffusion simulation, we analyzed the solidification process of BSS. The heat treatment conditions for the boride phase change were obtained through annealing at 1180 C for different times, and the effect of the annealing process on the tensile strength was also investigated. 2. Calculation method and experimental procedure Three alloys were designed based on type 304 stainless steel composition with different B contents; the chemical composition of the alloys is shown in Table 1. Alloys were labeled B02, B08 and B18 according to B contents. The alloys were fabricated through conventional ingot metallurgy. The 50 kg ingots were reheated at 1150 C for 1 h and hot-rolled into plates with thickness of 40 mm. Then, the as-received specimens were prepared by heat treatment at 1050 C for 30 min, followed by cooling in air. Cylindrical specimens with radius of 10 mm and length of 20 mm were cut and sealed in
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Table 1 Chemical compositions of investigated alloys in wt%. Alloy
C
Si
Mn
Cr
Ni
B
Al
B02 B08 B18
0.056 0.066 0.073
0.25 0.26 0.28
1.35 1.55 1.57
17.7 18.4 18.3
11.9 12.3 12.4
0.19 0.78 1.76
0.024 0.035 0.039
vacuum quartz tubes, which were air-cooled after annealing at 1180 for 48, 96 and 192 h using a box furnace. For the tensile test, a subsize tensile plate specimen of 2 mm thickness was prepared according to ASTM E8 test [11]. Microstructures were observed with optical micrograph (OM) and scanning electron micrograph (SEM) analyses. After polishing with a diamond suspension of 1 mm or less, backscattered electron (BSE) images were obtained using a JEOL IT-3000 SEM together with energy-dispersive X-ray spectroscopy (EDS) in large current mode. The polished specimens were etched in a nitric acid solution (30 ml of nitric acid þ 20 ml of hydrochloric acid þ 50 ml of ethanol), and the microstructure was observed. The density, average radius, and aspect ratio were quantified by Image-Pro Analyzer 7.0 software for more than 1500 precipitates from SEM images [12]. The grain size was measured according to ASTM E11213 [13]. Grain size number (G) was determined using a proportional formula and the average diameter of the corresponding grains was determined. Transmission electron microscopy (TEM) analysis was carried out using a JEOL JEM2100F with an acceleration voltage of 200 kV to examine the crystallography and morphology of borides. Specimens for TEM observation were prepared from mechanical-polished samples using focused-ion-beam micro-machining in combination with a lift-out method from the as-received specimen. The selected-area diffraction patterns obtained from the precipitates were analyzed to identify the crystal structure. For the analysis of the composition in the microstructure, EDS analysis was carried out at 7 points and averaged. In order to identify the distribution of alloying elements in the austenitic matrix, electron-probe micro-analyzer (EPMA) observation of specimens for as-received state and after 192 h of annealing were performed. EPMA analysis was performed using a JXA-8530F highresolution electron microscope, with an acceleration voltage of 15 kV, a resolution of 300 225 points and a dwell time of 20 msec. B, Ni and Cr were investigated, and fine compositional changes were observed by five overlapping analyses. Three tensile tests were performed on the as-received state samples and annealed specimen with the heat-treatment condition described above. Tensile tests were carried out at room temperature and at 2 mm/min using a 10-ton INSTRON 5882 tester and a 25 mm gauge length extensometer. SEM was used to examine the fracture surface and its surrounding area after the tensile test. Using ThermoCalc software, the thermodynamic equilibrium of the system was calculated based on the TCFE9.0 database [14]. The variation of the concentration profile of Cr and Ni during the annealing was calculated based on the MobFe4 database using multicomponent diffusional transformation (DICTRA) software [15]. The austenite phase is assumed to be a matrix, and a total length of 200 mm was simulated based on 200 points. The right half of the specimen is assumed to have a composition of pro-eutectic austenite, and the left half is assumed to be eutectic austenite. C
3. Results The equilibrium phase diagram of the BSS was calculated using ThermoCalc software, and is shown in Fig. 1. The figure shows phase
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Fig. 1. Equilibrium phase diagram for borated stainless steel based on TCFE9.0 database. Liquid, FCC, BCC and orthorhombic M2B phases are allowed in calculation. Diagram was established based on 18Cr-12Ni-1.5Mn wt%, corresponding to the composition of type 304 stainless steel. Blank squares and triangles show measured liquidus and solidus temperatures based on 20Cr-14Ni-1.8Mn wt% system [3].
change according to B content based on 18Cr-12Ni-1.5Mn wt%, which corresponds to the composition of type 304 stainless steel. Liquid, austenite (FCC), and ferrite (BCC) phases were allowed in the calculations. The (Fe,Cr)2B with orthorhombic structure was assumed to be the boride; its crystal structure was confirmed by TEM observation. The phase diagram shows reasonably good agreement with the measured solidus and liquidus temperatures of the previously reported study [3]. Differences in the measurements
and calculations for liquidus temperature are mainly due to different Cr and Ni contents. The amount of B corresponding to the eutectic point is 2.15 wt% and the temperature is 1256 C. Thus, the three alloys (B02, B08, B18) prepared were included in the hypoeutectic composition and a certain amount of pro-eutectic austenite was formed first during solidification. After pro-eutectic austenite formation, from the B enriched liquid, a eutectic mixture of austenite and (Fe,Cr)2B is expected to be generated. Figs. 2 and 3 show microstructures obtained using OM and SEM of the as-received specimens and those after annealing at 1180 C for 192 h, respectively. In Fig. 2 (a), (c) and (e), it can be seen that all three specimens are mixtures of pro-eutectic austenite, which is generated first during the cooling process in the liquid and eutectic microstructure. The eutectic microstructure generated from the B enriched liquid has an interconnected network structure. Table 2 summarizes the precipitate fraction, grain size, average size and aspect ratio of precipitates in each specimen. As the content of B increases, the volume fractions of boride and eutectic microstructure increase almost linearly and the fraction of pro-eutectic austenite decreases. After annealing, austenite grain size, the average size and aspect ratio of the precipitate changes significantly compared to those of the as-received specimens. During annealing at 1180 C, the plate shape (Fe,Cr)2B changes to a spherical shape. In addition, the average size also increases through the coarsening process, thereby decreasing the density of the precipitates and increasing the average distance between the precipitates. The pinning effect of precipitates leads to a significant difference in the austenite grain growth kinetics in each specimen. The
Fig. 2. Optical micrographs of (a) as-received B02, (b) annealed B02, (c) as-received B08, (d) annealed B08, (e) as-received B18, and (f) annealed B18 specimen. Annealing condition of specimen is 192 h at 1180 C.
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Fig. 3. Scanning electron micrographs of (a) as-received B02, (b) annealed B02, (c) as-received B08, (d) annealed B08, (e) as-received B18 and (f) annealed B18 specimen. Annealing condition of specimen is 192 h at 1180 C.
annealed B02 specimen has a small fraction and number of precipitates, such that austenite grain growth continues to occur and results in a grain size 5 times larger than that of the as-received specimen. This grain size affects the strength and elongation, which will be discussed later. On the contrary, the higher the content of B is, the larger the fraction and number of precipitates will be; due to the pinning effect, the grain size will also be smaller than that of the annealed B02 specimen. Characterization of the boride observed as the dark region in Fig. 3 was carried out using TEM. Several studies have been conducted on the boride crystal structure and the electronic structure of transition metal boride [16,17]. Havinga et al. reported that borides of Cr, Fe, Mn and Ni have tetragonal cell structures with a space group of I4/mcm [17]. The existence of orthorhombic Cr2B with space group Fddd was reported by Kotzott et al.; the refined cell constants were a ¼ 4:2752 Å, b ¼ 7:4523 Å and c ¼ 14:795 Å
[18]. Fig. 4 shows a bright field image and selected area diffraction pattern of (Fe,Cr)2B. The results of EDS measurements confirmed that these materials are borides containing large amounts of Cr and Fe. The diffraction pattern shows that the crystal structure of borides was confirmed to be orthorhombic and belongs to the space group Fddd, with a lattice parameter similar to reported values for Cr2B. Fig. 5 provides an engineering strain (ε)-stress (s) curves for the tensile tests. B02, B08 and B18 exhibited similar tensile strengths in a range from 550 to 580 MPa in the as-received state. However, since the volume fraction of (Fe,Cr)2B increases as the B content increases, the elongation dropped from 38% for the as-received B02 specimen to 5% for the as-received B18 specimen. Compared with the austenite matrix of metallic bonding, (Fe,Cr)2B with covalent bonding is brittle. A eutectic structure of (Fe,Cr)2B and austenite is formed from the B enriched liquid and eutectic structures are
Table 2 Phase fraction, mean precipitate size and aspect ratio of (Fe,Cr)2B and average austenite grain size of as-received specimens and specimens annealed for 192 h at 1180 C for B02, B08 and B18 alloys, respectively. Alloy
B02 B08 B18
precipitate size / mm
phase fraction / %
grain size / mm
aspect ratio
as-received
annealed
as-received
annealed
as-received
annealed
as-received
annealed
2.26 10.30 22.49
2.68 11.85 23.04
1.889 1.938 2.813
2.409 3.117 4.554
2.979 3.179 2.761
1.818 1.741 1.972
266 177 191
1449 363 293
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Fig. 4. Transmission electron micrograph of (Fe,Cr)2B precipitates: (a) Bright field image. Selected area diffraction patterns with (b) z ¼ [201] zone axis and (c) z ¼ [110] zone axis.
Fig. 5. Engineering stress (s)-strain (ε) curve for as-received and annealed specimens at 1180 C for (a) B02, (b) B08 and (c) B18 alloy. The black line is as-received specimen; the red, purple and blue lines represent specimens annealed for 48 h, 96 h and 192 h, respectively. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
connected to each other as shown in Figs. 2 and 3. Therefore, cracks initiated in the (Fe,Cr)2B or interface between (Fe,Cr)2B and austenite propagate easily along the eutectic structure. As the fraction of the eutectic structure increases, the ductility decreases dramatically. The mechanical properties of the alloys change markedly through annealing at 1180 C. In the B02 alloy, the elongation increases from 37% for the as-received specimen to over 62e66%. The elongations of the B08 and B18 specimens also increase with annealing, increasing from 24% to 39% for the B08 specimen and from 5% to 16% for the B18 specimen. The tensile strength decreases from 559 MPa to 541 MPa due to a decrease in the grain refining strengthening effect for the B02 specimen. On the contrary, in the case of the B08 specimen, the strength tends to be maintained; in
the case of the B18 specimen, the strength steadily increases from 534 MPa to 636 MPa. The work-hardening mechanism may be analyzed through the relationship between true-strain (εT ¼ lnð1 þ εÞ), true-stress (sT ¼ sð1 þ εÞ) and work hardening rate (ddεsTT ), as shown in Fig. 6. The work hardening rate can be expressed as an engineering strainstress derivative (ddεs) using the following equation
dsT ds dε d½sð1 þ εÞ dε ds dε ,ð1 þ εÞ þ s ¼ T ¼ ¼ : dεT dε dεT dεT dεT dε dε (1) In the plastic instability state in which local necking initiates during tensile testing, ddεs ¼ 0 because flow-stress is maximized.
Fig. 6. True stress (s)-strain (ε) curve and work hardening rate for as-received and annealed specimens at 1180 C for (a) B02, (b) B08 and (c) B18 alloy. The black line is as-received specimen; the red, purple and blue lines represent specimens annealed for 48 h, 96 h and 192 h, respectively. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
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Fig. 7. Microstructure near fracture surface of (a) as-received specimen and (b) annealed specimen for B18 alloy. Fracture surface of (c) as-received specimen and (d) annealed specimen for B18 alloy. Annealing condition of specimen is 192 h at 1180 C.
Therefore, the plastic instability condition is redefined as follows from the true-strain and true-stress relationship, which is Conre's criterion [19]. side
dsT dε ¼s ¼ sð1 þ εÞ ¼ sT : dεT dεT
(2)
For the B02 specimen, work hardening rate and true stress are encountered at about 40% strain for both as-received and postannealing samples. Because, compared to other alloys, the distribution of (Fe,Cr)2B is small, the curves are similar to that of austenitic stainless steel. As annealing progresses, austenite grain growth and precipitate density decrease, resulting in a decrease in work hardening rate and lower tensile strength. In the case of the B08 specimen, there is no intersection between work-hardening rate and true stress curve in the as-received state, so fracturing takes place before local necking occurs. However, after annealing, more strain is absorbed in the austenite matrix, local necking happens and elongation improvement can be obtained. In the case of the B18 specimen, both the as-received state and the annealing state fracture before reaching the plastic instability condition. Compared to the as-received state, the annealed state shows greatly improved elongation. Nevertheless, due to the fraction and distribution of (Fe,Cr)2B, fracturing of the specimen occurs due to cracks initiated from the particles before the austenite matrix absorbs the deformation. The interesting part of tensile testing of the B18 specimen is that, as the annealing progresses, both strength and elongation are improved. Comparing the results of the 48 h annealing to those of 192 h annealing, the work-hardening rate increases with the heat treatment time, tensile strength is improved by about 65 MPa and elongation can be increased from 14.6 to 16.4%. The simultaneous increase of the strength and elongation of the B18 specimen was attributed to the change in the concentration of the matrix; this was confirmed by EPMA analysis mentioned later. Fig. 7 shows the fracture surface and microstructure around the fracture surface after tensile testing of the B18 specimen. In the case
of the as-received specimen, most of the (Fe,Cr)2B precipitate has a plate-like shape with the austenite matrix, and the interface between the precipitate and matrix or precipitate itself is the crack initiation site. Cracks easily propagate along the interface; cracks easily connect to each other due to the short distance to surrounding precipitates. In the fracture surface observation shown in Fig. 7(c), the main fracture mechanism is brittle fracture inside the (Fe,Cr)2B precipitate. Even in the specimen after annealing, initial crack is also generated in the precipitate. However, unlike the asreceived specimen, the shape of the precipitate changes to spherical during the annealing, and the average precipitate size increases, such that the relative distance to adjacent cracks is increases. Therefore, a ductile fracture with dimple shape appears in the austenite matrix. The ductility improvement of the specimen after the annealing described above is due to this change in the fracture mechanism. The compositional distributions of pro-eutectic and eutectic austenite in the B18 specimen were confirmed using EPMA measurement. Fig. 8 shows the distributions of B, Ni and Cr of the proeutectic and eutectic areas of as-received B18 specimens. It can be confirmed that B and Cr are contained in the (Fe,Cr)2B precipitates, and Ni is barely present. In the austenite of the pro-eutectic and eutectic regions of the as-received specimens, the differences in concentration of Cr and Ni can be confirmed. Table 3 summarizes the results of analyzing the concentrations of each microstructure through EDS measurement. As-received specimens showed 12.8 wt % Cr and 14.9 wt% Ni in the pro-eutectic austenite region. In the eutectic region, the concentration of Cr was 2.2 wt% lower and that of Ni was 0.9 wt% higher. Fig. 9 shows the elemental distribution of the B18 specimen after annealing, indicating a uniform distribution throughout the austenite. The concentrations of regions are well matched with the thermodynamic calculations and show that amount of Cr is about 1.8 wt% lower (2.1 wt% lower in thermodynamic calculation) and amount of Ni is 1.2 wt% higher (1.5 wt% higher in thermodynamic calculation) compared to the proeutectic austenite in the as-received specimen.
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Fig. 8. EPMA observation showing distribution of Ni, Cr and B for as-received specimen of B18 alloy.
Table 3 EDS observations of each microstructure for as-received specimen and specimen annealed for 192 h at 1180 C for B18 alloy. Pro-eutectic austenite Contents / wt% As-received Annealed
Cr 12:8±0:6 e
Eutectic austenite Ni 14:9±0:6 e
Cr 10:6±0:2 11:0±0:1
(Fe,Cr)2B boride Ni 15:8±0:6 16:0±0:4
Cr 43:4±1:0 47:1±1:1
Fe 43:7±1:0 40:4±0:7
Fig. 9. EPMA observation showing distribution of Ni, Cr and B for specimen of B18 alloy annealed for 192 h at 1180 C.
4. Discussion The mechanical properties of BSS are mainly related to the shape and distribution of (Fe,Cr)2B. The composition corresponding
to the eutectic point of BSS is about 2.15 wt% B, and the fraction of the eutectic structure is proportional to the amount of B added. In addition, because it is transformed from B enriched liquid, the eutectic structure encloses pro-eutectic austenite. During the
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Fig. 10. Thermodynamic calculation results of (a) equilibrium phase fraction (b) alloy contents in austenite and (c) alloy contents in M2B phases, as a function of temperature for the B18 alloy. In (b) and (c), black, red, blue and green lines represent Fe, Cr, Ni and Mn contents, respectively. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
deformation process, initial cracks propagate primarily in the (Fe,Cr)2B or interface with the austenite; they are then connected to neighboring cracks. The plate-shape (Fe,Cr)2B produced by the eutectic process has a larger aspect ratio than the spherical particles and thus has a longer crack propagation length. When the aspect ratio is n times in the same volume of precipitate, the maximum propagation length in the precipitate can be increased nð2=3Þ times. During annealing at 1180 C, the plate-like (Fe,Cr)2B changes to a spherical shape, the aspect ratio decreases, the average precipitate size increases, the precipitate density decreases and the distance between the precipitate increases. As a result, dimple shape fracture surfaces in austenite are mainly observed in specimens after annealing, whereas the major fracture mechanism in as-received specimens is brittle fracture in the precipitate. Another important contributor to the mechanical properties of BSS steels is the compositional inhomogeneity of the austenite matrix. BSS has a hypoeutectic composition, and the matrix consists of, first, pro-eutectic austenite made in liquid during cooling and also eutectic austenite made from B enriched liquid. Fig. 10 shows the calculated equilibrium composition of each phase and phase fraction as a function of temperature for the B18 alloy. The concentrations of Cr and Ni in the pro-eutectic austenite at 1260 C are similar at about 11 wt%, but the concentration of Cr is lower and that of Ni is higher in eutectic austenite formed near 1200 C. These results match well with the inhomogeneity of Cr and Ni in the EPMA observation of an as-received specimen. After annealing, the equilibrium calculation results at 1180 C indicate that the
composition of the entire austenite region changes to high Ni and low Cr. Fig. 11 shows DICTRA simulation results of the homogenization of Cr and Ni in austenite over time. The as-received specimens of B18 alloy were simulated by assuming a composition of proeutectic and eutectic austenite. During the homogenization at 1180 C, there is a certain concentration gradient of Cr and Ni at 48 h and 96 h, and there is a difference in homogeneity during the heat treatment up to 192 h. As shown in Fig. 10(c), the equilibrium concentration of Cr in (Fe,Cr)2B at 1180 C is higher than that at eutectic temperature; this is also consistent with the EDS measurements shown in Table 3. Therefore, in the initial stage of annealing, elemental redistribution between austenite and (Fe,Cr)2B in the eutectic structure takes place before homogenization in the matrix. In order to maintain the local equilibrium in the eutectic region, Cr of austenite is partitioned into (Fe,Cr)2B, and the concentration of Cr is lower in eutectic austenite. When equilibrium is assumed in the B enriched eutectic region, the calculated equilibrium Cr concentration in the austenite at 1180 C is about 10.2 wt%, which is about 1.3 wt% lower than that of orthoequilibrium. That is, in the initial stage of annealing, before homogenization sufficiently progresses, partitioning of Cr from austenite to (Fe,Cr)2B occurs in the eutectic region, so that the concentration of Cr in the adjacent austenite drops. The effects of Cr and Ni on the strength and ductility of austenitic stainless steel have been studied in various ways [20e22]. Addition of Ni is known to increase toughness, but the
Fig. 11. Diffusion profile simulated using multicomponent diffusion transformation simulation for (a) Cr contents and (b) Ni contents at different annealing times. Black is initial profile. Red, blue and green lines represent profiles for the specimen after annealing for 48 h, 96 h and 192 h at 1180 C. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
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exact effect of strengthening is not clear. Depending on the alloy composition, strengthening and softening are observed with Ni addition, but the effect is smaller than those of other alloying elements [22,23]. In contrast, the addition of Cr is reported to improve work-hardening rate and strengthening effect by lowering the stacking fault energy [20,22]. In the early stage of the annealing process, the eutectic austenite is relatively soft because it contains less Cr than the pro-eutectic austenite region. Therefore, it mainly deforms in the eutectic austenite region during the tensile test. When annealing proceeds for a sufficient time at 1180 C, homogenization occurs between the pro-eutectic and eutectic austenite. In addition, the Cr concentration in the eutectic austenite region increases due to Cr migration from the pro-eutectic austenite to the eutectic austenite. As a result, homogeneous heat treatment of BSS can improve both the strength and ductility. Changes in size and shape of the boride can affect the neutron absorbing properties as well as the mechanical properties. The thermal neutron absorption cross section of boron is very large, so there could be a self-shielding effect that would increase as the size of boride increases. Strong absorbing elements could result in significant neutron flux depression in the materials, and lower effective neutron absorbability [24]. The thermal neutron mean free path length (l) in (Fe,Cr)2B can be approximated as the reciprocal of the macroscopic cross section (S),
S ¼ NB10 ,sB10
(3)
NB10 ¼ NA XB10 =Vm
(4)
where NB10 is the number of 10B isotopes per unit volume, sB10 ¼ 3840 barn is the microscopic cross section, NA is Avogadro's number, XB10 ¼ 1=3 0:198 is the mole fraction of 10B in (Fe,Cr)2B, and Vm ¼ 7:10 m3 =mol is the molar volume of (Fe,Cr)2B. Since the calculated mean free path length in boride is about 465 mm, which is much larger than the observed average boride size of 2.4e4.5 mm, the decrease in neutron absorbability due to self-shielding is expected to be insignificant. 5. Conclusions The evolution of the microstructure and mechanical properties during the annealing were analyzed for borated stainless steel with three different levels of B content. The three alloys, containing up to 1.8 wt% B, belong to the hypoeutectic composition, and have microstructures of pro-eutectic austenite and eutectic structure in asreceived state. As the content of B increases, the volume fraction of the eutectic structure, which is a mixture of austenite and (Fe,Cr)2B, increases and the mechanical properties degrade significantly. Plate-like (Fe,Cr)2B precipitates change to spherical shape during annealing at 1180 C and ductility improves. In the case of an alloy containing about 1.8 wt% B, strength and ductility improved simultaneously as the size of the precipitate increased and the Cr content of the austenite matrix became uniform. Therefore, sufficient annealing time is required to obtain improved mechanical properties of borated stainless steel containing a large amount of B. Acknowledgements This work was funded by the Fundamental R&D Program of the
Korea Institute of Materials Science, Grant No. PNK5850. This work was also supported by Basic Science Research Program through the National Research Foundation of Korea (NRF-2016R1C1B1011593) funded by the Ministry of Science, ICT and Future Planning. This study was also supported by the Ministry of Trade, Industry & Energy (MI, Republic of Korea) under Strategic Core Materials Technology Development Program (No. 10067375).
Appendix A. Supplementary data Supplementary data to this article can be found online at https://doi.org/10.1016/j.jnucmat.2018.12.039.
References [1] A. Machiels, R. Lambert, Handbook on Neutron Absorber Materials for Spent Nuclear Fuel Applications, Electric Power Research Institute, 2005. [2] S. Soliman, D. Youchison, A. Baratta, T. Balliett, Neutron effects on borated stainless steel, Nucl. Technol. 96 (3) (1991) 346e352. [3] C. Robino, M. Cieslak, High-temperature metallurgy of advanced borated stainless steels, Metall. Mater. Trans. 26 (7) (1995) 1673e1685. [4] Y. Choi, B.M. Moon, D.-S. Sohn, Fabrication of Gd containing duplex stainless steel sheet for neutron absorbing structural materials, Nucl. Eng. Technol. 45 (5) (2013) 689e694. [5] A. A887-89, Standard Specification for Borated Stainless Steel Plate, Sheet, and Strip for Nuclear Application, Tech. Rep., ASTM, 2014. [6] E.A. Loria, H.S. Isaacs, Type 304 stainless steel with 0.5% boron for storage of spent nuclear fuel, JOM 32 (12) (1980) 10e17. [7] J. Dilip, G.J. Ram, Microstructures and properties of friction freeform fabricated borated stainless steel, J. Mater. Eng. Perform. 22 (10) (2013) 3034e3042. [8] J. He, S.E. Soliman, A.J. Baratta, T.A. Balliett, Fracture mechanism of borated stainless steel, Nucl. Technol. 130 (2) (2000) 218e225. [9] X. He, T. Ahn, T. Sippel, et al., Corrosion of Borated Stainless Steel in Water and Humid Air, CORROSION, 2012. [10] H. Goldschmidt, Effect of boron additions to austenitic stainless steels, J. Iron Steel Inst. 209 (11) (1971) 900e909. [11] A. E8/E8M-16a, Standard Test Methods for Tension Testing of Metallic Materials, Tech. Rep., ASTM, 2016. [12] M. Cybernetics, Image-Pro Analyzer 7.0, http://www.mediacy.com/ imageproplus. [13] A. E112-13, Standard Test Methods for Determining Average Grain Size, Tech. Rep., ASTM, 2016. €glund, P. Shi, B. Sundman, Thermo-Calc & [14] J.-O. Andersson, T. Helander, L. Ho DICTRA, computational tools for materials science, Calphad 26 (2) (2002) 273e312. €glund, J. Ågren, A. Engstro €m, DICTRA, a tool for simulation [15] A. Borgenstam, L. Ho of diffusional transformations in alloys, J. Phase Equil. 21 (3) (2000) 269. [16] C. Zhou, J. Xing, B. Xiao, J. Feng, X. Xie, Y. Chen, First principles study on the structural properties and electronic structure of X2B (X¼ Cr, Mn, Fe, Co, Ni, Mo and W) compounds, Comput. Mater. Sci. 44 (4) (2009) 1056e1064. [17] E. Havinga, H. Damsma, P. Hokkeling, Compounds and pseudo-binary alloys with the CuAl2 (C16)-type structure I. Preparation and X-ray results, J. Less Common Met. 27 (2) (1972) 169e186. [18] D. Kotzott, M. Ade, H. Hillebrecht, Synthesis and crystal structures of a-and bmodifications of Cr2IrB2 containing 4-membered B4 chain fragments, the tboride Cr7. 9Ir14. 1B6 and orthorhombic Cr2B, Solid State Sci. 10 (3) (2008) 291e302. [19] R. E. Reed-Hill, R. Abbaschian, R. Abbaschian, Physical Metallurgy Principles. [20] L. Vitos, J.-O. Nilsson, B. Johansson, Alloying effects on the stacking fault energy in austenitic stainless steels from first-principles theory, Acta Mater. 54 (14) (2006) 3821e3826. [21] L. Vitos, P.A. Korzhavyi, B. Johansson, Stainless steel optimization from quantum mechanical calculations, Nat. Mater. 2 (1) (2003) 25. [22] D. Llewellyn, Work hardening effects in austenitic stainless steels, Mater. Sci. Technol. 13 (5) (1997) 389e400. € m, Proof strength values for austenitic stainless steels at [23] J. Eliasson, R. Sandstro elevated temperatures, Steel Res. 71 (6e7) (2000) 249e254. [24] J. Copley, Scattering effects within an absorbing sphere immersed in a field of neutrons, Nucl. Instrum. Methods Phys. Res. Sect. A Accel. Spectrom. Detect. Assoc. Equip. 307 (2e3) (1991) 389e397.