Effect of microstructural evolution on high-temperature strengthof 9Cr–3W–3Co martensitic heat resistant steel under differentaging conditions

Peng Yan a,b, Zhengdong Liu b, Hansheng Bao b, Yuqing Weng a, Wei Liu a,n

a Key Laboratory of Advanced Materials of Ministry of Education of China, School of Materials Science and Engineering, Tsinghua University, Haidian District,Beijing 100084, Chinab Institute for Special Steels, China Iron and Steel Research Institute, Beijing 100081, China

a r t i c l e i n f o

Article history:Received 21 February 2013Received in revised form8 September 2013Accepted 10 September 2013Available online 18 September 2013

Keywords:9Cr–3W–3Co martensitic heat resistantsteelAgingMicrostructural evolutionHigh-temperature strengthAthermal yield stress

a b s t r a c t

Evolution of microstructures and high-temperature strength at 650 1C of 9Cr–3W–3Co martensitic heatresistant steel after aging at 650 1C and 700 1C for different time durations have been experimentallyinvestigated using field emission scanning electron microscopy (FESEM), X-ray diffraction (XRD), fieldemission transmission electron microscopy (FETEM) and post-aged tensile tests. The results show thatafter aging at 650 1C, the high-temperature strength and the microstructures of 9Cr–3W–3Co steel keepalmost stable with increasing aging time from 300 h to 3000 h. In comparison, after aging at 700 1C,there are obvious changes in the high-temperature strength and the microstructures. The strengtheningmechanisms of the 9Cr–3W–3Co steel were also discussed and the athermal yield stresses werecalculated. The change of the high-temperature strength is mainly affected by the evolution ofdislocations and laths. The precipitates mainly act as obstacles against motion of dislocations and lathboundaries.

& 2013 Elsevier B.V. All rights reserved.

1. Introduction

For the purposes of reducing the CO2 emission, improving thethermal efficiency of traditional coal-fired plants as well as ful-filling the requirement of saving energy sources and protecting theenvironment, many efforts have been done to expand the steamtemperature from conventional 600 1C up to 650 1C and higher inultra-supercritical (USC) power plants [1–5]. The increase in steamtemperatures exceeding 600 1C requires extensive research anddevelopment of advanced steels with long-term creep rupturestrength higher than that of conventional steels [1]. The targetcreep rupture strength for main pipe steam steels used in USCpower plants is 100 MPa for 100,000 h rupture life at the operatingtemperatures [6].

Nowadays there are still great difficulties in the choice of thesteels for main steam pipe at 650 1C. Nickel-base superalloyswhich can fulfill the requirement of creep rupture strength [1]are too expensive for commercial use. Austenitic steels with lowthermal conductivity coefficient and high thermal expansioncoefficient [4,7] are unsuitable for the main steam pipes.

Therefore, development of martensitic steels is highly desirable.Martensitic steels with low cost possess high thermal conductivitycoefficient and low thermal expansion coefficient, and they alsopossess less susceptibility to thermal fatigue than austenitic steels[7]. However, a major shortcoming of the martensitic steels is theirlow creep rupture strength [4], which places a limit on themaximum service temperature. The upper-use temperatures ofthe existing martensitic steels like P92, E911 and P122 are about600–620 1C [3,5,8,9]. In recent years, attention has been paid on9Cr–3W–3Co steels which were firstly developed in Japan by NIMS[1]. Compared with P92 and P122 steel, W was added instead ofMo, Co was added and contents of B and N were optimized in 9Cr–3W–3Co steels. Completely tempered martensite can be obtainedby adding 3% Co which can also strengthen the matrix by acting asa solute element. Coarsening of the M23C6 precipitates is expectedto be suppressed by adding W and B. Therefore, 9Cr–3W–3Costeels are expected to be used in USC power plants up to 650 1C.Lots of work has been done on 9Cr–3W–3Co steels, such as thealloy design [10–13], the effect of chemical composition onmicrostructural aspects and creep behavior [14–18], the micro-structural stabilization during creep [19–21] and so on. However,there is still little report on the effect of microstructural evolutionon the high-temperature tensile and yield strength in 9Cr–3W–

3Co steels. During long-time service, to maintain high level ofstrength of heat resistant steels is the key point for developing

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Materials Science & Engineering A

0921-5093/$ - see front matter & 2013 Elsevier B.V. All rights reserved.http://dx.doi.org/10.1016/j.msea.2013.09.033

n Corresponding author. Tel.: þ86 10 62772853; fax: þ86 10 62771160.E-mail addresses: maxeric@163.com (P. Yan),

liuw@mail.tsinghua.edu.cn (W. Liu).

Materials Science & Engineering A 588 (2013) 22–28

new steels. In the present paper, high-temperature strength at650 1C and the corresponding quantitative microstructure char-acterization have been performed in 9Cr–3W–3Co heat resistantsteel under different aging conditions. Furthermore, relationshipbetween high-temperature strength and microstructures of 9Cr–3W–3Co heat resistant steel has also been clarified.

2. Experimental

The test steel used in this study was named G115 steel and itwas melted by a vacuum induction furnace and then casted into aningot at Bao-Shan Iron and Steel Company (BaoSteel). The mainchemical composition of the steel is given in Table 1. The ingot wasthen forged and hot extruded into a pipe with an outside diameterof 254 mm, an inside diameter of 204 mm, and a length of 3.5 m.

Tensile and impact specimens were all cut from the pipe alonglongitudinal axial direction. They were normalized at 1100 1C for1 h, followed by tempering at 780 1C for 3 h. The normalizingtemperature is higher than the classical ones. Because grain-boundary strength is weaker than matrix strength in the servicetemperature and too much grain boundaries will reduce the creepstrength of the material in heat resistant steels, the austenite grainsize is usually not expected to be fine. Meanwhile, there are somecoarse particles after forging and hot processing. In order to makethese particles re-dissolved as many as possible during thenormalizing process and then precipitate finely and denselyduring the tempering process, a high normalizing temperature isexpected and 1100 1C is just enough to make the particles re-dissolved in G115 steel. It is generally acknowledged that a coarseparent austenite grain size is detrimental to toughness. As it ismentioned in Section 3.1, the Charpy impact energy values are stillhigh after this heat treatment. Therefore, 1100 1C is considered tobe a suitable normalizing temperature for G115 steel. This tem-perature or even higher temperatures are also used in somesimilar steels [1,18,22]. After heat treatments, blanks were agedat 650 1C and 700 1C for 300 h, 1000 h and 3000 h, respectively.Then the blanks were machined into standard tensile test speci-mens of 5 mm in gage diameter and 25 mm in gage length andstandard impact specimens of 10�10�55 mm3 containing a2 mm long 451 V-notch with a root radius of 0.25 mm. The post-aged tensile tests were carried out at 650 1C using INSTRON 5582tensile testing machine [23]. Crosshead displacement in tensiletesting machine was measured in real time. The rate of displace-ment was 0.25 mm/min before Rp0.2 and 2.5 mm/min after Rp0.2.The tests were carried out in a furnace. The furnace was connectedwith the laboratory. There are three thermocouples in the top,middle and bottom of the furnace, respectively, and the uniformtemperature zone is big enough to make sure the specimens at therequired temperature. The Charpy impact tests were carried out atroom temperature.

Precipitates in the as-treated samples were characterized byHitachi S-5500 and TESCAN MIRA 3 LMH field emission scanningelectron microscopy (FESEM). The chemical composition of theprecipitates was measured using energy dispersive spectrometry(EDS) in the FESEM. Small angle X-ray scattering (SAXS) was usedfor the size measurement of M23C6 and Laves phase precipitates.Detailed information of SAXS is as follows: the mixed powder wasobtained by electrolyzing the specimens in 3.6% ZnCl2þ5%

HClþ1% citric acidþ90.4% methanol at �5 to 0 1C using electriccurrent density of 0.05 A/cm2. The mixed powder was measured tobe MX, M23C6 and Laves phases by using XRD. Laves phasespowder could be obtained after the mixed powder was furthermixed with 20% HClþ80% CH3CH2OH and boiled for 2 h. M23C6

powder could be obtained after the mixed powder was mixed with6% H2SO4þ20% H2O2þ2% citric acidþ72% H2O and boiled for 1 h.After boiling treatments, the powder was separated by suctionfiltration and the apertures of the filter paper were about 200 nm.Since the precipitates electrolyzed from the specimens clustered,they could not be filtered through the filter paper. Therefore,nearly all of the precipitates no matter which size would beremained after separation. After the above steps, pure Lavesphases, and M23C6 could be separated. Then the separated powderwas measured using Rigaku 3014 X-ray diffracto-spectrometer[24]. The radiation was Co Kα and the load was 35 kV and20 mA. The SAXS method is based on the fact that SAXS effectresults from the difference of electron density between particlesand their surroundings. The mean size was the average size of thediameter of equivalent scattering ball. The sizes of each precipi-tates but not clusters could be easily obtained by the SAXSmethod, so the SAXS method was considered to be more precisethan SEM images method. The expected bias of this method onsize measurement was below 10%.

The lath width test was carried out on thin foils using TECNAIG220 field emission transmission electron microscopy (FETEM) at200 kV, and the lath width was statistically evaluated (measuredperpendicularly to the long axis of the martensite laths) in severalprior austenite grains covering a total area 4400 μm2. More than100 laths were measured and then the average values of themwere taken. Dislocation density in this paper was obtained byusing XRD. From previous research [25,26], dislocation densitiescalculated from XRD line profiles agreed well with those measuredusing TEM. Therefore, XRD is also a believable method to measuredislocation density and the value of the experimental uncertaintycan be considered similar with the uncertainty of the TEMmethod. The radiation was Co Kα, and the voltage was 30 kV.

3. Results

3.1. Evolution of microstructures of G115 steel under as-receivedand different aging conditions

Under as-received condition, the average parent austenite grainsize is about 70 μm, as shown in Fig. 1. It was reported [1] that the

Table 1Main chemical composition of G115 steel (wt%).

C Cr W Co V Nb N B Fe

0.076 8.83 3.11 2.99 0.19 0.042 0.014 0.013 Bal.Fig. 1. Optical micrograph showing prior austenite grains of G115 steel.

P. Yan et al. / Materials Science & Engineering A 588 (2013) 22–28 23

peak time to rupture and the minimum creep rate were located atabout 80–100 ppm nitrogen, which corresponded to the max-imum solid solubility of nitrogen in equilibriumwith boron nitridein the 9Cr–3W–3Co steel with 139 ppm boron at a normalizingtemperature of 1100 1C. The content of nitrogen in the presentpaper is a little higher than 100 ppm but the average Charpyimpact energy value is 114 J which stands for nice toughness,indicating that there is no precipitation of large boron nitridesdetrimental to the toughness of the steel and the content ofnitrogen can be considered optimum. Fig. 2 shows the precipitatesand the laths of G115 steel under as-received condition. It is atypical tempered martensitic microstructure with plenty of parti-cle precipitated in the steel. There are nearly no Laves phases inthe material since Laves phases are gradually precipitated duringthe aging process.

Back-scattered electron (BSE) images of G115 steel after agingat 650 1C and 700 1C are shown in Fig. 3. There are three maintypes of precipitates (M23C6, Laves phase and MX) in the 9–12% Crmartensitic heat resistant steel. The main composition of MX is(V, Nb)(C, N) and their coarsening rate is usually very low. Themain composition of M23C6 is Cr23C6 while Laves phase is mainly

Fe2W. Compared with MX precipitates, M23C6 and Laves phasecoarsen fast and they can obviously affect the high-temperaturestrength of the steel [1,27]. As W is heavier than Cr in atomweight,Laves phase is brighter than M23C6 [27,28] in the BSE imageswhich are used to distinguish Laves phase from M23C6 precipitatesin this paper. From Figs. 2 and 3 and the statistical results obtainedfrom SAXS, the average sizes of M23C6 and Laves precipitates underdifferent aging conditions are shown in Fig. 4. During the first300 h aging, M23C6 particles grow slowly at 650 1C from 92 nm to108 nm while they coarsen quickly from 92 nm to 152 nm at700 1C. Additionally, a large number of precipitation of Lavesphase particles occurs during this period. With increasing agingtime from 300 h to 3000 h, the average size of M23C6 increasesslowly from 108 nm to 153 nm and the average size of Laves phaseincreases gradually from 129 nm to 184 nm after aging at 650 1C,while the average size of M23C6 increases from 152 nm to 200 nmand the average size of Laves phase increases obviously from238 nm to 300 nm and then slightly to 315 nm after aging at700 1C. With increase in aging temperature from 650 1C to 700 1C,the size of the precipitates increases obviously, especially for theLaves phase. Thus, the coarsening of the precipitates is promoted

Fig. 2. (a) SEM and (b) TEM images of G115 steel under as-received condition.

Fig. 3. BSE images of G115 steel after aging at 650 1C for (a) 300 h, (b) 1000 h, and (c) 3000 h and at 700 1C for (d) 300 h, (e) 1000 h, and (f) 3000 h.

P. Yan et al. / Materials Science & Engineering A 588 (2013) 22–2824

by the aging time and temperature, especially by the elevatedtemperature.

TEM images of G115 steel after aging at 650 1C and 700 1C areshown in Fig. 5 and the average lath width of G115 steel under as-received and different aging conditions is shown in Fig. 6. TheM23C6 and Laves phase particles are mainly precipitated along theprior austenite grain boundaries and the lath boundaries. Duringthe first 300 h aging, the width of lath increases slowly from330 nm to 350 nm at 650 1C while it increases obviously from330 nm to 382 nm at 700 1C. After aging for 300 h at bothtemperatures, it is still regular martensite lath microstructure inthe steel. With increasing aging time from 300 h to 3000 h, thelath width increases obviously from 350 nm to 392 nm and thenslightly to 403 nm after aging at 650 1C, while it increasesobviously from 382 nm to 430 nm and then slowly to 473 nmafter aging at 700 1C. After aging at 650 1C, the martensite laths aremainly in shape of regular laths. In comparison, after aging at700 1C, breakup of the martensite laths and the gradual develop-ment of an equiaxed substructure obviously occur, resulting fromsome small-scale changes, such as dislocation annihilation andrearrangement. The lath morphology is transformed to dislocationarrays lying transverse to longitudinal axes of the laths (Fig. 5(d)),followed by the formation of an equiaxed grain structure with alow dislocation density (Fig. 5(f)). With increase in aging tem-perature from 650 1C to 700 1C, the lath width increases greatly,especially after 3000 h. Therefore, it is summarized that therecovery and coarsening of the martensite laths in G115 steel are

promoted by the aging time and temperature, especially by theelevated temperature.

The dislocation densities of G115 steel under as-received anddifferent aging conditions are shown in Fig. 7. During the first300 h aging, the dislocation density decreases slowly from2.65�1014 /m2 to 2.16�1014 /m2 at 650 1C while it decreases

Fig. 4. Average size values of (a) M23C6 and (b) Laves phase in G115 steel as a function of aging time at 650 1C and 700 1C.

Fig. 5. TEM images of G115 steel after aging at 650 1C for (a) 300 h, (b) 1000 h, and (c) 3000 h and at 700 1C for (d) 300 h, (e) 1000 h, and (f) 3000 h.

Fig. 6. Lath width values of G115 steel as a function of aging time at 650 1C and700 1C.

P. Yan et al. / Materials Science & Engineering A 588 (2013) 22–28 25

drastically from 2.65�1014 /m2 to 5.49�1014 /m2 at 700 1C. Withincreasing aging time from 300 h to 3000 h, the dislocationdensity decreases slightly from 2.16�1014 /m2 to 1.88�1014 /m2

while it decreases sharply from 5.49�1013 /m2 to 1.16�1013 /m2

and then slowly to 8.5�1012 /m2. The sharp decrease in disloca-tion density during aging at 700 1C also implies that the disloca-tion annihilation and rearrangement drastically occur, as shown inFig. 5. With increase in aging temperature from 650 1C to 700 1C,the dislocation density decreases obviously, especially after morethan 1000 h. Therefore, the recovery of the dislocations is greatlypromoted after aging at 700 1C, while the dislocation densitykeeps almost the same after aging at 650 1C.

3.2. Evolution of high-temperature strength of G115 steel underas-received and different aging conditions

The high-temperature tensile strength (Rm) and the high-temperature yield strength (Rp0.2) of G115 steel measured at650 1C under as-received and different aging conditions are shownin Fig. 8. During the first 300 h aging, the tensile strengthdecreases from 380 MPa to 340 MPa at 650 1C while it decreasesdrastically from 380 MPa to 303 MPa at 700 1C. With increasingaging time from 300 h to 3000 h, the tensile strength slightlydecreases from 357 MPa to 343 MPa after aging at 650 1C, while it

decreases sharply from 340 MPa to 320 MPa and then to 297 MPaafter aging at 700 1C. With increase in aging temperature from650 1C to 700 1C, the tensile strength decreases greatly, especiallyafter 3000 h. The high-temperature yield strength under differentaging conditions shows the similar tendency with the high-temperature tensile strength.

4. Discussion

It is well known that the basic methods by which martensiticheat resistant steels can be strengthened are solution hardening,precipitation hardening, dislocation hardening and lath hardening.According to previous report [29], solution hardening is small insteels with precipitates and dislocation substructure. In order tofurther clarify the mechanism of high-temperature strength withthe corresponding microstructural evolution in G115 steel, ather-mal yield stress is used to describe the influence of different agingconditions on precipitation hardening, dislocation hardening andlath substructure hardening [29]:

sp ¼ 0:8 MGb=λi ð1Þ

sρ ¼ 0:5 MGbffiffiffi

ρp ð2Þ

sl ¼ 10 Gb=λl ð3Þwhere M is the Taylor factor and a value of 3 is usually taken intempered martensitic heat resistant steels [1,29], G is the shearmodulus (64 GPa at 650 1C), b is the length of Burgers vector(0.25 nm), λi is the mean spacing of the precipitates, ρ is thedislocation density and λl is the average lath width.

Athermal yield stresses of precipitates, dislocations and laths ofG115 steel under as-received condition and after different agingtimes at 650 1C and 700 1C are listed in Table 2. It is difficult how toconsider the contribution of different strengthening mechanisms.In this paper, athermal yield stress is used to semi-quantitativelydescribe influence of the microstructures under different agingconditions, and we do not pay much attention to the absolutevalues of the athermal yield stresses, but the tendency of them.Therefore, we consider the contribution of the different strength-ening mechanisms as additive for convenience. By summing upthe three types of athermal yield stresses sp, sρ and sl calculated ineach specimen, athermal yield stresses under as-received anddifferent aging conditions are shown in Fig. 9, which are similarwith the tendency of the high-temperature tensile and yieldstresses as shown in Fig. 8. From Table 2, the precipitationhardening is obviously weaker than the other two hardeningmechanisms. Therefore, it is concluded that the change in high-temperature strength is mainly affected by the evolution of thedislocations and the laths.

Though the precipitation hardening is smaller than the dis-location hardening and the lath hardening in G115 steel, theprecipitates play an important role in pinning and prohibitingthe movement of the dislocations and the lath boundaries. After

Fig. 7. Dislocation density values of G115 steel as a function of aging time at 650 1Cand 700 1C.

Fig. 8. High-temperature strength values of G115 steel as a function of aging timeat 650 1C and 700 1C.

Table 2Athermal yield stresses of precipitates (sp), laths (sl) and dislocations (sρ) of G115steel under as-received and different aging conditions.

Aging temperature Aging time (h) sp (MPa) sl (MPa) sρ (MPa)As-received 90 485 391

650 1C 300 96 457 3531000 91 408 3333000 85 397 329

700 1C 300 82 419 1781000 72 372 823000 65 338 70

P. Yan et al. / Materials Science & Engineering A 588 (2013) 22–2826

aging at 650 1C, M23C6 and Laves phases are fine and they can pinthe dislocations and lath boundaries effectively, resulting that thedislocation density keeps almost the same and the laths coarsenslightly. In comparison, after aging at 700 1C, sizes of M23C6 andLaves phases increase greatly from 300 h to 1000 h and thenslowly from 1000 h to 3000 h, resulting that the lath widthincreases greatly from 300 h to 1000 h and then slowly from1000 h to 3000 h, and the dislocation density decreases drasticallyfrom 300 h to 1000 h and then slightly from 1000 h to 3000 h.Therefore, to maintain high level of high-temperature strength inmartensitic heat resistant steels, it is important to suppress thecoarsening of the precipitates during long-term service.

From the results and analyses above, after aging at 700 1C forjust 3000 h, the obvious change in microstructures and the sharpdecrease of the high-temperature strength of G115 steel imply that700 1C is the limited operating temperature for G115 martensiticheat resistant steels. However, compared with P92 [30] and T122[27] steel, precipitates in G115 steel are finer and coarsen moreslowly after aging at 650 1C until 3000 h, as listed in Table 3. Thismainly owes to the containing of boron. According to the pub-lished results on steels of the 9Cr–3W–3Co family [1], in the steelwithout boron, a fine distribution of M23C6 carbides was observedafter tempering but extensive coarsening took place in the vicinityof the PAGBs during aging at 650 1C. With the increasing of boroncontent in the steel, the size of M23C6 carbides increased moreslowly after the same aging time. In the present study, M23C6

carbides in G115 steel with 130 ppm boron also had a lowcoarsening rate during aging at 650 1C, indicating that the additionof boron reduced the rate of the coarsening of M23C6 carbides inthe vicinity of PAGBs during exposure at 650 1C. As the coarseningrate of M23C6 reduces, migration of the lath boundaries andrecovery of the free dislocations can be suppressed by the pinning

effect of the fine and dense precipitates, resulting in a more stablemicrostructure in G115 steel than that in P92 and T122 steel.Therefore, G115 steel is better than P92 and T122 steel during thefirst 3000 h aging at 650 1C. Whether G115 steel can be used at650 1C, which is a significant improvement over the upper-usetemperature of about 620 1C for martensitic steels now available,long-term aging and creep are needed. More aspects of G115 steelinvolving the precipitates stability and the microstructural evolu-tion during creep test will be discussed in the following works.

5. Conclusion

Evolution of microstructures and high-temperature strength ofG115 martensitic heat resistant steel after aging at differenttemperatures have been experimentally studied. The importantconclusions of the study are listed below:

(1) After aging at 650 1C, the high-temperature strength and themicrostructures of G115 steel keep almost stable with increas-ing aging time from 300 h to 3000 h. And G115 steel is betterthan P92 and T122 steel during the first 3000 h aging at650 1C. In comparison, after aging at 700 1C, there are obviouschanges in the high-temperature strength and the microstruc-tures implying 700 1C is the limited operating temperature forG115 steel.

(2) The change of the high-temperature strength in G115 steel ismainly affected by the evolution of the dislocations and thelaths. Precipitation hardening is smaller than dislocation hard-ening and lath hardening. The precipitates mainly act asobstacles against motion of dislocations and lath boundaries.

Acknowledgments

The present study was financially supported by the National BasicResearch Program of China (973 Program, No. 2010CB630804), andthe National High Technology Research & Development Program ofChina (863 Program, No. 2012AA03A501).

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Fig. 9. Athermal yield stress values of G115 steel as a function of aging time at650 1C and 700 1C.

Table 3Average size values of the precipitates in some 9–12% Cr martensitic heat resistantsteel under different aging conditions.

Agingtime (h)

G115,650 1C

T122, 650 1C[27]

P92,600 1C [30]

M23C6 300 108 nm – 145 nm1000 116 nm 310 nm 160 nm3000 153 nm 380 nm –

Laves phase 300 129 nm – –

1000 143 nm 410 nm –

3000 184 nm 485 nm –

P. Yan et al. / Materials Science & Engineering A 588 (2013) 22–28 27

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