Effects of High-Temperature Tempering on Mechanical Properties and Microstructure of SA738 Gr.B Steel

metals

Article Effects of High-Temperature Tempering on Mechanical Properties and Microstructure of SA738 Gr.B Steel

Yanmei Li 1,*, Shuzhan Zhang 1,*, Chunyao Zhao 1, Minghui Song 1 and Zaiwei Jiang 2

1 The State Key Laboratory of Rolling and Automation, Northeastern University, Shenyang 110819, China; [email protected] (C.Z.); [email protected] (M.S.) 2 Nanjing Iron and Steel Co., Ltd., Nanjing 210035, China; [email protected] * Correspondence: [email protected] (Y.L.); [email protected] (S.Z.)

Received: 4 August 2020; Accepted: 3 September 2020; Published: 9 September 2020

Abstract: In this paper, JMatPro thermodynamic software, OM, SEM, TEM, and EPMA were used to study the microstructure and mechanical properties of SA738 Gr.B nuclear power steel after tempering at 630–710 ◦C. When tempered within the range of 630–670 ◦C, a huge amount of M3C and MC carbides were dispersed and precipitated in the ferrite matrix, and the strength and toughness matched well; when the tempering temperature rose above 670 ◦C, hard and brittle plate-martensites formed at the grain boundary, leading to the tensile strength of the experimental steel increased, while the low-temperature impact toughness signiﬁcantly decreased and the yield strength also declined due to the disappearance of the ﬁnely dispersed second phase particles in the matrix.

Keywords: SA738 Gr.B; high-temperature tempering; mechanical properties; microstructure

1. Introduction In recent years, nuclear energy, as a clean energy source, has received considerable attention from countries in the context of increasingly tight world energy supplies and worsening environmental problems. Nuclear safety is an integral part of national security [1]. The typical nuclear power reactor is AP1000, designed by Westinghouse [2]. The structure of the AP1000 nuclear reactor is a double-layer envelope structure composed of steel and reinforced concrete, which is different from the traditional prestressed concrete envelope structure [3]. Therefore, the unconventional design of the AP1000 containment has put forward higher requirements of strength and low-temperature toughness. To meet the higher standard [4], the primary material of its containment is an SA738 Gr.B steel plate in the American Society of Mechanical Engineers (ASME) SA738/SA738 M standard [5], according to the material specification, SA738 Gr.B steel plates shall have a minimum impact energy of 27 J either at 45 °C in the quarter of the thickness direction. To obtain a better combination of strength and − low-temperature impact toughness, the typical heat treatment process of SA738 Gr.B is quenching and tempering [5]. For the quenching and tempering process of SA738 Gr.B steel, relevant scholars have carried out related research: Han et al. [6] and Sun [7] studied the ferrite matrix after quenching and tempering and found that the increase in tempering temperature coarsened the grains, and with the increase in the soft and ductile phase ferrite content, the strength of SA738 Gr.B steel decreased and the impact energy increased. Bi et al. [8] observed the precipitated phases after tempering at 630 ◦C by TEM; it was found that a large number of carbides with various shapes were dispersed and precipitated between the laths, hindering the growth of grains and inhibited the movement of dislocations. Zhang [9] et al. used the master curve (MC) method to modify the ductile–brittle transition zone and fracture toughness of the tempered SA738 Gr.B steel, which proves that it has sufficient toughness reserve at a low temperature. However, a systematic and comprehensive study on

Metals 2020, 10, 1207; doi:10.3390/met10091207 www.mdpi.com/journal/metals MetalsMetals 20202020, 10,,10 x ,FOR 1207 PEER REVIEW 2 of2 11 of 12 and comprehensive study on the brittle temperature range and the mechanisms of strengthening and the brittle temperature range and the mechanisms of strengthening and toughening of SA738 Gr.B toughening of SA738 Gr.B steel is still lacking. steel is still lacking. In this paper, the effects of tempering temperature on the strength and toughness of SA738 Gr.B In this paper, the eﬀects of tempering temperature on the strength and toughness of SA738 nuclear power steel are systematically studied by comparing the strength and low-temperature Gr.B nuclear power steel are systematically studied by comparing the strength and low-temperature impact energy at different tempering temperatures and clarifying the evolution of the microstructure impact energy at diﬀerent tempering temperatures and clarifying the evolution of the microstructure combined with structural characterization methods such as SEM and TEM, and thermodynamic combined with structural characterization methods such as SEM and TEM, and thermodynamic software calculations. software calculations.

2. 2.Material Material and and Methods Methods TheThe material material used used in the in theresearch research was was rolled rolled SA738Gr.B SA738Gr.B steel, steel, and the and thickness the thickness was 46mm. was 46mm. Its chemicalIts chemical composition composition is shown is shown in Table in 1. Table Table1. 1 Table also 1lists also the lists chemical the chemical composition composition range of range SA738 of Gr.BSA738 steel Gr.B in the steel ASME in the standard ASME [5]. standard The steel [5]. plat Thee was steel cut plate in the was quarter cut in theof the quarter thickness of the direction, thickness 3 anddirection, several samples and several with samples dimensions with of dimensions 12 × 40 × 120 of mm 12 were40 120taken mm for3 heatwere treatment. taken for heatThe samples treatment. were austenitized at 900 °C for 30 min, followed by water× quenching× to room temperature, and then The samples were austenitized at 900 ◦C for 30 min, followed by water quenching to room temperature, tempered at 630, 650, 670, 690, and 710 °C for 60 min, respectively. Figure 1 is the SEM micrograph and then tempered at 630, 650, 670, 690, and 710 ◦C for 60 min, respectively. Figure1 is the SEM of micrographSA738 Gr.B of steel SA738 after Gr.B quenching, steel after through quenching, which through we can which conclude we can that conclude the microstructure that the microstructure of as- quenchedof as-quenched steel is lath steel martensite. is lath martensite.

TableTable 1. 1.ChemicalChemical compositions compositions ofof the the steel steel SA738 SA738 Gr.B Gr.B in ASME in ASME standard standard and this and experiment this experiment (mass fraction, (mass %). fraction,%). Element C Mn P S Si Ni Cr Mo Nb V Ti Fe Element C Mn P S Si Ni Cr Mo Nb V Ti Fe Standard 0.2 0.9–1.5 0.008 0.005 0.15–0.55 0.6 0.3 0.3 Total 0.08 Bal. ≤ ≤ ≤ ≤ ≤ ExperimentalStandard ≤0.2 0.14 0.9–1.5 1.55 0.008 0.008 0.005 0.001 0.15–0.55 0.25 ≤ 0.550.6 ≤0.230.3 ≤0.280.3 Total Total ≤= 0.080.077 Bal. Bal. Experimental 0.14 1.55 0.008 0.001 0.25 0.55 0.23 0.28 0.02 0.04 0.017 Bal.

FigureFigure 1. SEM 1. SEM image image of as-quenched of as-quenched SA738 SA738 Gr.B Gr.B steel. steel.

AfterAfter heat heat treatment, treatment, Charpy Charpy V-notch V-notch specimen specimenss with with dimensions dimensions of of 10 10 mm mm × 1010 mm mm × 5555 mm mm3 3 × × werewere prepared. prepared. The The V-notch V-notch was was parallel parallel to tothe the rolling rolling direction. direction. Charpy Charpy impact impact tests tests were were performed performed at at−45 45°C ◦onC onan anInstron Instron Dynatup Dynatup 9250 9250 instrumented instrumented impact impact machine machine (Norwood, (Norwood, MA, MA, USA) USA) according according − to tothe the working working condition condition of ofthe the steel. steel. Three Three specimen specimenss were were tested tested for for each each heat heat treatment treatment process process in in orderorder to toguarantee guarantee reproducibility. reproducibility. Specimens Specimens for for the the strength strength tests tests with with a 5 a mm 5 mm diameter diameter and and 50 50mm mm gaugegauge length length were were prepared prepared according according to toGB/T GB/ T228–2010 228–2010 [10]. [10 ]. SamplesSamples for for metallographic metallographic analysis analysis were were prepared prepared by bywire-cut wire-cut electrical electrical discharge discharge machining. machining. TheThe microstructures microstructures were observedobserved by by ZEISS ZEISS ULTRA ULTRA 55 Field 55 EmissionField Emission Scanning Scanning Electron MicroscopyElectron Microscopy(FESEM, Carl (FESEM, Zeiss Carl AG, Jena,Zeiss Germany)AG, Jena, afterGermany) grinding, after polishing, grinding, andpolishing, etching and with etching 4% nital with solution. 4% nitalA Tecnaisolution. G2 A F20 Tecnai transmission G2 F20 transmission electron microscope electron microscope (TEM, FEI, (TEM, Hillsboro, FEI, OR,Hillsboro, USA) wasOR, USA) also used was to alsoobserve used to the observe dislocations the dislocations and phases. and The phases. distribution The distribution of alloy elements of alloy waselements analyzed was byanalyzed using JEOLby usingJXA-8530F JEOL JXA-8530F field emission field electronemission probe electron (EPMA, probe Jeol (EPMA, Co. Ltd., Jeol Kyoto, Co. Ltd., Japan). Kyoto, The Japan). crystal The structures crystal of structuresthe particles of the were particles identified were by Digitalidentified Micrograph by Digital software Micrograph (Gatan, software Pleasanton, (Gatan, CA, USA).Pleasanton, The diagram CA, USA). The diagram between mass fraction and temperature of all the phases and the trend of the Metals 2020, 10, x FOR PEER REVIEW 3 of 11 Metals 2020, 10, 1207 3 of 12 precipitation type and quantity of the carbides were calculated by JMatPro software (Version 6.0, Generalbetween Steel mass module, fraction Sente and temperatureSoftware Ltd., of Guildford, all the phases UK) and according the trend to the of the chemical precipitation composition type and of SA738quantity Gr.B. of the carbides were calculated by JMatPro software (Version 6.0, General Steel module, Sente Software Ltd., Guildford, UK) according to the chemical composition of SA738 Gr.B. 3. Results 3. Results 3.1. Mechanical Properties 3.1. Mechanical Properties Figure 2a shows the strength and elongation change of SA738 Gr.B steel after tempering at differentFigure temperatures2a shows the for strength 60 min. and It elongation can be know changen that of SA738 the tensile Gr.B steel strength after temperingof SA738 atGr.B di ﬀ steelerent decreasestemperatures from for 659 60 MPa min. at It 630 can ℃ be to known 611 MPa that at the 670 tensile °C; when strength tempered of SA738 at 690 Gr.B and steel 710 decreases°C, the tensile from strength659 MPa increased, at 630 °C to reaching 611 MPa a at maximum 670 ◦C; when of 698 tempered MPa. The at 690 yield and strength 710 ◦C, the gradually tensile strength decreased increased, from 581reaching MPa at a maximum630 °C to 498 of 698 MPa MPa. at 710 The °C. yield The strength elongation gradually reached decreased the maximum from 581 of 20.8% MPa at at 630 650◦ C°C, to and498 gradually MPa at 710 decreased◦C. The elongation with the increase reached thein tempering maximum temperature, of 20.8% at 650 and◦C, the and final gradually is 17.66% decreased at 710 °C.with Figure the increase 2b shows in the tempering impact temperature,toughness of andSA738 the Gr.B ﬁnal steel is 17.66% at −45 at°C. 710 The◦C. impact Figure energy2b shows of the steelimpact maintained toughness at of 250 SA738 J at 630 Gr.B to steel 650 at°C;45 the◦C. im Thepact impact energy energy dropped of the to steel203 J maintainedat 670 °C; when at 250 the J at − tempering630 to 650 ◦temperatureC; the impact is energy 690 °C, dropped the impact to 203 energy J at 670 decreased◦C; when sharply the tempering to 83 J, and temperature 42 J at 710 is 690°C. ◦InC, summary,the impact when energy the decreased tempering sharply temperature to 83 J, andis lo 42wer J atthan 710 670◦C. In°C, summary, SA738 Gr.B when steel the has tempering better comprehensivetemperature is lowermechanical than 670properties.◦C, SA738 Gr.B steel has better comprehensive mechanical properties.

300 750 22 Tensile strength (a) (b)

Yield strength /J 250

700 V Elongation 21 650 200 20 600 150

550 19 Elongation/%

Strength/MPa 100 500 18 50 45℃ Impact energy AK 450 - 17 630 650 670 690 710 630 650 670 690 710 Tempering Temperature/℃, 1×1h Temperature/℃ (a) (b)

FigureFigure 2. 2. Strength,Strength, elongation elongation (a (a),), and and impact impact energy energy (b (b) )of of tempered tempered steel steel at at different diﬀerent temperatures. temperatures.

3.2.3.2. Microstructure Microstructure of of the the Tempered Tempered Steel Steel FigureFigure 33 showsshows thethe SEMSEM morphologymorphology imagesimages ofof SA738SA738 Gr.BGr.B steelsteel atat didifferentﬀerent tempering tempering temperatures.temperatures. After After tempering tempering at at630 630 °C,◦ theC, themicros microstructuretructure of the of thesteel steel was wastempered tempered sorbite, sorbite, the ferritethe ferrite matrix matrix had hada polygonal a polygonal or a or lath a lath structure structure (Figure (Figure 3a),3a), and and the the width width of of the the lath lath structure structure was was lessless than than 1 1 μµm.m. A A large large number number of of second-phase second-phase particles particles pr precipitatedecipitated on on the the matrix matrix or or laths laths diffusely diﬀusely (Figure(Figure 33b),b), with a size of less than than 200 200 nm. nm. When When tempered tempered at at 650 650 °C,◦C, the the microstructure microstructure was was roughly roughly thethe same same as as 630 630 °C.◦C. The The size size of of the the second-phase second-phase particles particles increased, increased, with with a a maximum maximum size size of of more more thanthan 200 200 nm nm (Figure (Figure 33c,d).c,d). WhenWhen temperedtempered at at 670 670◦ °C,C, aasmall small amount amount of of island-like island-like structures structures could could be beobserved observed at at the the grain grain boundaries, boundaries, with with the the size size betweenbetween 0.20.2 andand 1 µμm (Figure 33e),e), andand the number ofof second-phase second-phase particles particles on on the the ferrite ferrite matrix matrix was was reduced reduced (Figure 33f).f). The island-like structure continuedcontinued to to precipitate precipitate along along the the grain grain boundary boundary when when tempered tempered above above 690 690 °C;◦C; the the volume volume fraction fraction graduallygradually increased, andand interconnectedinterconnected to to become become the the main main phase phase of theof microstructurethe microstructure (Figure (Figure3g,h). 3g,h).The average The average proportion proportion of the of island-like the island-like structure structure in the in tempered the tempered structure structure at 670–710 at 670–710◦C was °C 5.1%,was 5.1%,18.6%, 18.6%, and 34.9%, and 34.9%, respectively. respectively. Metals 2020, 10, 1207 4 of 12 Metals 2020, 10, x FOR PEER REVIEW 4 of 11

Figure 3. SEM images ofof SA738SA738 Gr.BGr.B steel steel tempered tempered at at di differentﬀerent temperatures: temperatures: (a, b(a), 630b) 630◦C; °C; (c,d ()c, 650d) 650◦C; °C;(e,f )(e 670,f) 670◦C; °C; (g) 690(g) 690◦C; °C; (h) ( 710h) 710◦C. °C.

The fracture morphology of the SA738 Gr.B steel impact specimen is shown in Figure4 4.. ItIt couldcould be seenseen thatthat when when the the tempering tempering temperature temperature was wa 630–670s 630–670◦C, the°C, corresponding the corresponding fracture fracture was a typicalwas a typicaldimple-type dimple-type fracture fracture morphology, morphology, as shown as in shown Figure in4a–c. Figure The 4a–c. sections The were sections all composed were all composed of plastic ofpits. plastic A large pits. plastic A large deformation plastic deformation occurred occurred before the before fracture, the fracture, and the and corresponding the corresponding impact impact energy energywas relatively was relatively high; when high; thewhen tempering the tempering temperature temperature was 690 was and 690 710 and◦C, 710 as °C, shown as shown in Figure in Figure4d,e, 4d,e,the fracture the fracture morphology morphology was a cleavagewas a cleavage plane, whoseplane, typicalwhose rivertypical pattern river could pattern be observed,could be observed, and there andwere there block were craters block (red craters dashed (red lines) dashed and secondary lines) and cracks; secondary that is,cracks; the original that is, block the original structure block was structurethe source was of cracking,the source corresponding of cracking, corresponding to the island to structure the island in structure the SEM (Figurein the SEM3). The (Figure size 3). of The the sizeplane of wasthe plane smaller was at 690smaller◦C, andat 690 the °C, river and pattern the river was pattern zigzagged was zigzagged and intertwined, and intertwined, which indicated which indicatedthat the unit that expansion the unit expansion path was small.path was After small. the microcracksAfter the microcracks blocked along blocked the cleavagealong the plane cleavage and planethe expansion and the directionexpansion changed, direction it requiredchanged, more it requ energy.ired more The crack energy. initiation The crack area wasinitiation large atarea 710 was◦C, largeso the at corresponding 710 °C, so the impactcorresponding energy was impact the lowest.energy was the lowest. Metals 2020, 10, 1207 5 of 12 Metals 2020, 10, x FOR PEER REVIEW 5 of 11

FigureFigure 4. 4. FractureFracture surface surface morphologies morphologies of of SA SA738738 Gr.B Gr.B steel tempered at 630 ◦°CC( (a), 650 °C◦C( (bb),), 670 670 °C◦C ((cc),690), 690 °CC( (dd),), 710 710 °CC( (ee) )after after a aCharpy Charpy impact impact test test at at −4545 °CC.. ◦ ◦ − ◦

FigureFigure 55 isis TEMTEM micrographsmicrographs ofof SA738SA738 Gr.BGr.B steelsteel afterafter temperingtempering atat 630630 andand 710710 °C.◦C. When When the the steelsteel tempered tempered at at 630 630 °C,◦C, part part of of the the laths laths could could be be observed. observed. The The second second phase phase particles particles precipitated precipitated alongalong with with favourable favourable positions positions such such as as grain grain bounda boundaries,ries, lath lath boundaries, boundaries, or or the the matrix matrix as as shown shown in in FigureFigure 55a;a; atat 650 650◦ °C,C, the the morphology morphology of of the the lath lath disappeared, disappeared, and and the carbidethe carbide distribution distribution was thewas same the sameas 630 as◦ C.630 The °C. sizeThe ofsize the of large-sized the large-sized precipitates precipitates ranged ranged from from 100 to 100 300 tonm, 300 asnm, shown as shown in Figure in Figure5a,b. 5a,b.Diﬀerently, Differently, it was it was strip-like strip-like or spherical; or spherical; the the small-size small-size carbides carbides were were less less than than 100 100 nm,nm, as shown shown inin Figure Figure 55c,c, andand itit waswas squaresquare andand distributeddistributed ininthe the matrix. matrix. ByBy usingusing JmatProJmatPro thermodynamicthermodynamic softwaresoftware (Figure (Figure 66)) andand EnergyEnergy DispersiveDispersive SpectrometerSpectrometer (Figure7 7),), thethe large-sizedlarge-sized precipitatesprecipitates were were MM33C-typeC-type alloy alloy carbides carbides (M (M == Fe,Fe, Mn, Mn, Cr), Cr), and and the the small-sized small-sized precipitates precipitates were were MC-type MC-type alloy alloy carbidescarbides (M(M= =Ti, Ti, Nb, Nb, Cr, Cr, V). BecauseV). Because of the of low the content low content and small and size small of MC size carbides, of MC the carbides, precipitation the precipitationstrengthening strengthening and the damage and to the toughness damage areto limited,toughness so weare couldlimited, ignore so we the could discussion ignore of MCthe discussioncarbides in of this MC paper. carbides in this paper. At 690 and 710 °C, as shown in Figure 5d, almost no precipitates could be observed, but a large area of black island-like structure could be observed, and its size was more than 500 nm. Corresponding to the island-like structure, the surrounding stress field was strong due to the phase change, so high-density dislocation networks distributed on the surrounding ferrite matrix; Figure 5e is the morphology of the island-like structure and the corresponding selected area electron diffraction spectrum. The diffraction pattern determined that it was plate martensite with a bct-twin structure. The thickness of the twin layer was about a few nanometers and was relatively uniform. Metals 2020, 10, 1207 6 of 12 Metals 2020, 10, x FOR PEER REVIEW 6 of 11

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FigureFigure 5. TEM 5. TEM images images and and selected selected didiffractionﬀraction spots spots of of SA738 SA738 Gr.B Gr.B steel steel after after tempering tempering at different at diﬀ erent Figure 5. TEM images and selected diffraction spots of SA738 Gr.B steel after tempering at different temperatures:temperatures: (a) ( 630a) 630◦C; °C; (b (,bc,)c 650) 650◦ °C;C; ((dd) 690 °C;◦C; (e (e) )710 710 °C.◦C. temperatures: (a) 630 °C; (b,c) 650 °C; (d) 690 °C; (e) 710 °C.

80 80 1.5 M3C Austenite Ccarbides/% 3 1.5％ M C 60 Fe 7 3 M3C Austenite Ccarbides/% 3 ％ M C 60 Fe 1.0 7 3 40 1.0

0.5 40 C Mass Fraction/ 20 Mn 0.5 MC C

Mass Fraction/ Cr 20 Mn 600MC 625 650 675 700

Mole fraction of alloying elements in M in elements alloying of fraction Mole 580 600 620 640 660 680 700 720 Temperature/℃ Cr Temperature/℃ 600 625 650 675 700

(a) M in elements alloying of fraction Mole 580 600 620(b) 640 660 680 700 720 Temperature/℃ 60 Temperature/℃ (a) C (b) 60 40 C

Ti 40 20 Nb

Ti Cr V 20 Nb Molealloying of fraction elements in MC carbides/% 580 600 620 640 660 680 700 720 Temperature/℃ Cr (c) V

FigureFigure 6. Thermodynamic 6. Thermodynamic calculation calculationMolealloying of fraction elements in MC carbides/% 580 ofof 600 thethe 620second 640 phase phase 660 particles. particles. 680 700 (a) (Thea 720) The second second phase phase precipitates precipitates in thein temperingthe tempering temperature temperature range; range; ( b(b)) thethe compositioncompositionTemperature/℃ of of the the M3C; M3C; (c) (thec) the composition composition of the of MC. the MC. (c)

Figure 6. Thermodynamic calculation of the second phase particles. (a) The second phase precipitates in the tempering temperature range; (b) the composition of the M3C; (c) the composition of the MC. Metals 2020, 10, 1207 7 of 12

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Figure 7. Energy Dispersive Spectrometer(EDS) analysis of precipitates in the tempered SA738 Gr.B Figure 7. Energy Dispersive Spectrometer(EDS) analysis of precipitates in the tempered SA738 Gr.B steel. steel. (a) M3C; (b) MC. (a)M3C; (b) MC. 4. Discussion At 690 and 710 ◦C, as shown in Figure5d, almost no precipitates could be observed, but a large area of black island-like structure could be observed, and its size was more than 500 nm. 4.1. Three Different Tempering Temperature Ranges Corresponding to the island-like structure, the surrounding stress field was strong due to the phase change,According so high-density to the dislocationmicrostructure networks of SA738 distributed Gr.B steel on after the surrounding high-temperature ferrite tempering, matrix; Figure the 5e is thetempering morphology temperature of the island-likecan be divided structure into three and different the corresponding ranges: selected area electron diffraction spectrum.1. Tempering The diffraction at 630–650 pattern °C: There determined was little that chan it wasge in platethe strength martensite and toughness. with a bct-twin The changes structure. Thein thicknessthe microstructure of the twin mainly layerwas included about athe few precipitation nanometers and and growth was relatively of M3C uniform. carbides and the polygonalization of ferrite. 4. Discussion2. Tempering at 670 °C: The strength and toughness were reduced. A small amount of island- like plate-martensite began to form on the grain boundary, and the size of M3C carbides in the matrix 4.1.increased. Three Di ffInerent this Temperingprocess, the Temperature larger M3C Ranges particles grow at the expense of the smaller particles. This further reduces the surface energy and is accomplished by the diffusion of carbon atoms through the According to the microstructure of SA738 Gr.B steel after high-temperature tempering, the tempering iron matrix [11]. temperature can be divided into three different ranges: This range continued the changes of the ferrite matrix and carbides from the previous stage. Although1. Tempering a small at amount 630–650 of◦C: hard There and was brittle little phas changees generated, in the strength but andthe amount toughness. was The not changes enough into the microstructureincrease the strength; mainly included so the overall the precipitation strength of the and steel growth was reduced. of M3C carbides Correspondingly, and the polygonalization the softened ofmatrix ferrite. interacted with the hard and brittle plate-martensite, and the impact toughness did not decrease2. Tempering much. The at 670 transition◦C: The point strength of tempering and toughness temperature were reduced.of 670 °C Acan small be used. amount of island-like plate-martensite3. Tempering began at 690–710 to form °C: on During the grain this boundary, range, the andferrite the was size completely of M3C carbides polygonal, in theand matrixthe increased.lath morphology In this completely process, the disappeared. larger M3 CA large particles amou grownt of plate-martensite at the expense ofprecipitated the smaller and particles. grew Thisalong further the grain reduces boundaries the surface duri energyng cooling and after is accomplished tempering. M by3C carbides the diffusion may precipitate of carbon atomsin the initial through thestage, iron matrixbut as the [11 ].tempering temperature increased, carbon and alloys elements would gradually be enrichedThis range in austenite continued with a the higher changes solid ofsolubility. the ferrite This matrix can also and be carbidesfully proved from in the calculation previous stage. of Althoughthermodynamic a small amountsoftware, of as hard shown and in brittle Figure phases 6a. Fi generated,nally, ferrite, but plate-ma the amountrtensite, was and not retained enough to increaseaustenite the structures strength; can so thebe obtained. overall strength of the steel was reduced. Correspondingly, the softened matrix interacted with the hard and brittle plate-martensite, and the impact toughness did not decrease 4.2. Mechanism of Effect of Plate Martensite on Properties much. The transition point of tempering temperature of 670 ◦C can be used. 3.During Tempering the transformation at 690–710 ◦C: of During austenite this to range, plate-martensite, the ferrite wasthe volume completely expansion polygonal, due to and the the lathlattice morphology expansion; completely the expanded disappeared. plate-martensite A large co amountmpressed of the plate-martensite ferrite matrix; so precipitated the stress-induced and grew alonghigh the density grain dislocation boundaries effectively during cooling hinders after slippage tempering. [12], resulting M3C carbides in significantly may precipitate improved in thetensile initial stage,strength. but as the tempering temperature increased, carbon and alloys elements would gradually be enrichedAdditionally, in austenite according with a higher to the solidAshby–Orow solubility.an Thismechanism can also [13], be the fully yield proved strength in the σppt calculation due to ofprecipitation thermodynamic strengthening software, can as shownbe expressed in Figure as follows:6a. Finally, ferrite, plate-martensite, and retained austenite structures can be obtained. 0.8MGb x (1) σppt= ln 2πL√1-v 2b where Metals 2020, 10, 1207 8 of 12

4.2. Mechanism of Effect of Plate Martensite on Properties During the transformation of austenite to plate-martensite, the volume expansion due to the lattice expansion; the expanded plate-martensite compressed the ferrite matrix; so the stress-induced high density dislocation effectively hinders slippage [12], resulting in significantly improved tensile strength. Additionally, according to the Ashby–Orowan mechanism [13], the yield strength σppt due to precipitation strengthening can be expressed as follows:

0.8MGb x σppt = ln (1) 2πL √1 v 2b − where s r ! 2 π L = 2 r (2) 3 f − r 2 x = 2 r (3) 3 so the σppt can be expressed as another form:

a + clnr σppt = p (4) π 2 r f − Here, M is the Taylor constant, G is the shear modulus, v is the Poisson’s ratio, b is the Berkeley vector mode, r is the size of the precipitated phase, f is the volume fraction of the precipitated phase, a and c are constants. From Formula (4), it can be seen that the precipitation strengthening effect is mainly related to f and r; that is, the larger the volume fraction of the precipitated phase, the smaller the r, the higher the strength. The amount of M3C carbides in the matrix decreased sharply; almost no carbide precipitated at 710 °C, and the precipitation strengthening effect in the ferrite matrix was greatly reduced, resulting in a continuous decrease in the yield strength of the steel. When there are twin substructures in martensite, the number of slip systems reduces, and the dislocations need to be in a “Z” shape when passing through the twins, which increases the resistance to deformation and leads to stress concentration. So large-sized martensite twins can promote crack growth and severely deteriorate the low-temperature toughness of steel [14]. At the same time, EPMA was used to analyze the plate-martensite. As shown in Figure8, it was found that the phosphorus element segregated on the prior austenite grain boundary. Takayama et al. [15] proposed a method for calculating the ductile–brittle transition temperature of steel using experiments and Taylor series expansion. The results showed that the ductile–brittle transition temperature of steel increased with the increase in the degree of phosphorus segregation, so the phosphorus segregation seriously affects the low-temperature toughness of steel. When the matrix was subjected to impact load, plate-martensite hindered the movement of dislocations in the matrix, causing dislocation plugging, as shown in Figure5. The plugged dislocation, in turn, exerted stress on the plate-martensite [16]. The stress interaction caused the plate-martensite to break or caused the surrounding ferrite to form microcracks due to stress concentration. Therefore, microcracks that met the Griffith condition nucleated on the plate-martensite or the interface of the martensite-matrix [17], which caused cleavage fracture. Observation of secondary cracks on the fracture side section of the impact specimen by SEM also verified this view, as shown in Figure9. In summary, plate-martensite is the preferred location for crack initiation and propagation [12] and is the direct cause of the low-temperature impact toughness reduction. The schematic diagrams of the microstructure evolution, precipitation behaviour, and crack propagation path of SA738 Gr.B steel at different tempering temperatures are shown in Figure 10. Metals 2020, 10, x FOR PEER REVIEW 8 of 11

2 π - (2) L= 3 f 2 r

2 (3) x=2 3 r so the σppt can be expressed as another form:

a+clnr σppt= π (4) -2 r f Here, M is the Taylor constant, G is the shear modulus, v is the Poisson’s ratio, b is the Berkeley vector mode, r is the size of the precipitated phase, f is the volume fraction of the precipitated phase, a and c are constants. From Formula (4), it can be seen that the precipitation strengthening effect is mainly related to f and r; that is, the larger the volume fraction of the precipitated phase, the smaller the r, the higher the strength. The amount of M3C carbides in the matrix decreased sharply; almost no carbide precipitated at 710 ℃, and the precipitation strengthening effect in the ferrite matrix was greatly reduced, resulting in a continuous decrease in the yield strength of the steel. When there are twin substructures in martensite, the number of slip systems reduces, and the dislocations need to be in a “Z” shape when passing through the twins, which increases the resistance to deformation and leads to stress concentration. So large-sized martensite twins can promote crack growth and severely deteriorate the low-temperature toughness of steel [14]. At the same time, EPMA was used to analyze the plate-martensite. As shown in Figure 8, it was found that the phosphorus element segregated on the prior austenite grain boundary. Takayama et al. [15] proposed a method for calculating the ductile–brittle transition temperature of steel using experiments and Taylor series expansion. The results showed that the ductile–brittle transition temperature of steel increased with the increase in the degree of phosphorus segregation, so the phosphorus segregation seriously affects the low-temperature toughness of steel. When the matrix was subjected to impact load, plate-martensite hindered the movement of dislocations in the matrix, causing dislocation plugging, as shown in Figure 5. The plugged dislocation, in turn, exerted stress on the plate-martensite [16]. The stress interaction caused the plate- martensite to break or caused the surrounding ferrite to form microcracks due to stress concentration. Therefore, microcracks that met the Griffith condition nucleated on the plate-martensite or the interface of the martensite-matrix [17], which caused cleavage fracture. Observation of secondary Metalscracks2020 on, 10the, 1207 fracture side section of the impact specimen by SEM also verified this view, as shown9 of 12 in Figure 9.

Metals 2020, 10, x FOR PEER REVIEW 9 of 11

(a) (b) Figure 8. EPMA analysisFigure of the 8. segregationEPMA analysis of phosphorus. of the segregation (a) Microstruture; of phosphorus. (b) segregation of phosphorus. Metals 2020, 10, x FOR PEER REVIEW 9 of 11

(a) (b)

1μm 1μm

Figure 9. SEM images of the cross-sectional area of the Charpy impact specimens fractured at −45 °C for the SA738 Gr.B steel tempered at 690 °C (a), 710 °C (b). 1μm 1μm In summary, plate-martensite is the preferred location for crack initiation and propagation [12] and is the direct cause of the low-temperature impact toughness reduction. The schematic diagrams FigureFigure 9.9. SEM images ofof thethe cross-sectionalcross-sectional areaarea ofof thethe CharpyCharpy impactimpact specimensspecimens fracturedfracturedat at −4545 °CC of the microstructure evolution, precipitation behaviour, and crack propagation path of SA738− ◦ Gr.B for the SA738 Gr.B steel tempered at 690 °C (a), 710 °C (b). steel forat different the SA738 tempering Gr.B steel tempered temperatures at 690 ◦areC(a shown), 710 ◦C( inb Figure). 10. In summary, plate-martensite is the preferred location for crack initiation and propagation [12] and is the direct cause(a) of the low-temperature impact(b) toughness reduction. The schematic diagrams M3C M3C of the microstructure evolution, precipitation MCbehaviour, and crack propagationMC path of SA738 Gr.B steel at different tempering temperatures are shown in Figure 10. Block boundary

Prior austenite Prior austenite (a) grain boundary (b) grain boundary M3C M3C

MC MC (c) (d)

M3C Block boundary Crack Crack Dislocation PriorDislocation austenite Prior austeniteRA grain boundary grain boundary RA Twin martensite Twin martensite Segregation of P Segregation of P (c) (d) M C 3 Crack Figure 10. Schematic representation for microstructure Crack an andd precipitates of SA738 Dislocation Gr.B under different different Dislocation temperingtempering temperatures:temperatures: (a) 630 ◦°C;C; (b) 650 ◦°C;C; ( c) 690 °C;◦C; ( d)) 710 710 °C.◦C. RA RA Twin martensite Twin martensite Segregation of P 4.3. Formation Mechanism of Plate MartensiteSegregation of P 4.3. Formation Mechanism of Plate Martensite Regarding the formation mechanism of twin-martensite in steel, it is generally believed that the RegardingFigure 10. Schematic the formation representation mechanism for microstructure of twin-martensite and precipitates in steel, ofit SA738is generally Gr.B under believed different that the martensite twins are caused by the shear transformation of the austenite [18]. As carbon content in martensitetempering twins temperatures: are caused ( aby) 630 the °C; shear (b) 650 transforma °C; (c) 690tion °C; (ofd) 710the °C.austenite [18]. As carbon content in austenite increases, the substructure of martensite transitions from dislocation type to twin type [19].Additionally,4.3. Formation Mechanism Huang ofet Plateal. [20] Martensite believed that the substructure of martensite also evolved with the decreaseRegarding in the the phase formation transition mechanism temperature. of twin-martensite Dislocation-type in steel, martensite it is generally generated believed at thathigher the temperaturesmartensite twins and are twin-type caused atby lower the shear temperatures. transforma tion of the austenite [18]. As carbon content in The SA738 Gr.B steel used in this experiment was a type of low-carbon steel: plate martensite austenite increases, the substructure of martensite transitions from dislocation type to twin type will not form. In fact, through the calculation of JMatPro software, the AC1 temperature of the [19].Additionally, Huang et al. [20] believed that the substructure of martensite also evolved with the experimental steel was 675 °C. Heat treatment above 675 °C did not temper in the traditional sense; decrease in the phase transition temperature. Dislocation-type martensite generated at higher so a certain amount of austenite was formed,and during the cooling, only the austenite had phase temperatures and twin-type at lower temperatures. transformation.The SA738 Gr.BJMatPro steel thermodynamic used in this experiment software wawass a also type used of low-carbon to calculate steel: the plate main martensite chemical component content and Ms temperature of the austenite at various temperatures, as shown in Table will not form. In fact, through the calculation of JMatPro software, the AC1 temperature of the 2.experimental It can be seen steel that was the 675 austenite °C. Heat phase treatment formed above in the 675 tempering °C did not stage temper could in the be traditionalconsidered sense; to be high-carbon steel, ensuring the composition fluctuation to form the plate-martensite. Besides, Kelly so a certain amount of austenite was formed,and during the cooling, only the austenite had phase et al. [21] discovered that in iron-carbon alloys, phase transition twins easily formed when the Ms transformation. JMatPro thermodynamic software was also used to calculate the main chemical point was lower than 300 °C due to the increase in the mass fraction of carbon. Mn, Ni, Mo, and Cr component content and Ms temperature of the austenite at various temperatures, as shown in Table can not only improve the hardenability of steel but also increase the tendency to form plate- 2. It can be seen that the austenite phase formed in the tempering stage could be considered to be high-carbon steel, ensuring the composition fluctuation to form the plate-martensite. Besides, Kelly et al. [21] discovered that in iron-carbon alloys, phase transition twins easily formed when the Ms point was lower than 300 °C due to the increase in the mass fraction of carbon. Mn, Ni, Mo, and Cr can not only improve the hardenability of steel but also increase the tendency to form plate- Metals 2020, 10, 1207 10 of 12 austenite increases, the substructure of martensite transitions from dislocation type to twin type [19]. Additionally, Huang et al. [20] believed that the substructure of martensite also evolved with the decrease in the phase transition temperature. Dislocation-type martensite generated at higher temperatures and twin-type at lower temperatures. The SA738 Gr.B steel used in this experiment was a type of low-carbon steel: plate martensite will not form. In fact, through the calculation of JMatPro software, the AC1 temperature of the experimental steel was 675 ◦C. Heat treatment above 675 ◦C did not temper in the traditional sense; so a certain amount of austenite was formed, and during the cooling, only the austenite had phase transformation. JMatPro thermodynamic software was also used to calculate the main chemical component content and Ms temperature of the austenite at various temperatures, as shown in Table2. It can be seen that the austenite phase formed in the tempering stage could be considered to be high-carbon steel, ensuring the composition fluctuation to form the plate-martensite. Besides, Kelly et al. [21] discovered that in iron-carbon alloys, phase transition twins easily formed when the Ms point was lower than 300 ◦C due to the increase in the mass fraction of carbon. Mn, Ni, Mo, and Cr can not only improve the hardenability of steel but also increase the tendency to form plate-martensite [22]. As can be seen from Table2, the austenite under this composition had a lower M s point and a higher alloy content. Under the continuous cooling condition, due to the small size of the sample, the cooling rate was fast in the air, and the distortion energy was constantly increasing. If martensite grows in the direction of symmetry of the crystal plane during the growth process, the strain energy can be reduced and adjusted to meet the energy fluctuation. This is also an essential reason for the formation of martensite twins [15].

Table 2. Elemental mass fraction (%) for equilibrium austenite phase of SA738 Gr.B steel at diﬀerent

tempering temperatures and Ms (◦C) point.

Temperature/◦C C Mn Ni Cr Mo Ms 690 0.52 3.80 1.15 0.36 0.19 161.6 710 0.42 3.23 1.03 0.34 0.21 175.7

Ahn et al. [23] found that adding intercritical heat treatment which performed in the ferrite+austenite (α + γ) temperature region between the quenching and tempering processes of SA508-3, which is also a type of nuclear pressure vessel steel, can effectively improve the toughness, and its matrix composition (Fe-0.21C-1.24Mn-0.25Si-0.88Ni-0.21Cr-0.47Mo) (wt%) is similar to the SA738 Gr.B steel. However, according to the analysis of the current results, tempering at 690 and 710 ◦C is equivalent to a intercritical heat treatment. Carbon atom distribution of ferrite-martensite was uneven compared with that of lath martensite after quenching. After tempering again, carbides were unevenly precipitated near the island-like structure, which inevitably affected the properties of the steel. Wang et al. [24] also proved that their performance was poor after tempering SA738 Gr.B steel with ferrite-martensite dual-phase structure. Therefore, in industrial production, the tempering temperature should be strictly controlled to achieve a well-fit strength and toughness.

5. Conclusions In this paper, in order to achieve suitable strength and toughness for SA738 Gr.B nuclear power steel, five tempering experiments with different processes were conducted, and the mechanical properties and microstructure were tested and characterized. The following conclusions were drawn from the test results: 1. The best tempering temperature range of SA738 Gr.B steel is 630–670 ◦C, and it has a better combination of strength and low-temperature impact toughness. 2. When the tempering temperature of SA738 Gr.B steel is lower than 670 ◦C, the microstructure of SA738 Gr.B steel still maintains the lath morphology. The precipitated particles of the second phase are fine M3C and MC which are dispersed; so the steel has a better combination of strength and Metals 2020, 10, 1207 11 of 12

toughness. When the tempering temperature was higher than 670 ◦C, a large amount of carbon-rich austenite formed during the tempering process, and most of them transformed into hard and brittle plate martensite which is favourable for crack growth, leading to the sharp decrease in impact energy at 45 C; the yield strength decreased due to the disappearance of the ﬁne dispersed second particles − ◦ in the matrix.

Author Contributions: Conceptualization, Y.L. and S.Z.; investigation, S.Z.; data curation, C.Z. and M.S.; writing—original draft preparation, S.Z.; writing—review and editing, Y.L., C.Z., M.S., and Z.J.; funding acquisition, Y.L. All authors have read and agreed to the published version of the manuscript. Funding: This research was supported by the China National Key R&D Program during the 13th Five-year Plan Period (Grant No. 2017YFB0305301). Conﬂicts of Interest: The authors declare no conﬂict of interest.

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