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Materials Science and Engineering A 528 (2011) 5164–5176

Contents lists available at ScienceDirect

Materials Science and Engineering A

journa l homepage: www.e lsev ier .com/ locate /msea

%Cr heat resistant steels: Alloy design, microstructure evolution and creepesponse at 650 ◦C

. Rojasa,b, J. Garciab,∗, O. Prata,c, G. Sauthoff c, A.R. Kaysser-Pyzallab

Universidad de Concepción, Departamento de Ingeniería de Materiales, Edmundo Larenas 270, Concepción, ChileHelmholtz-Zentrum Berlin für Materialien und Energie GmbH, Hahn-Meitner-Platz 1, 14109 Berlin, GermanyMax-Planck-Institut für Eisenforschung GmbH, Max-Planck-Str. 1, 40237 Düsseldorf, Germany

r t i c l e i n f o

rticle history:eceived 12 November 2010eceived in revised form 1 March 2011ccepted 7 March 2011vailable online 11 March 2011

a b s t r a c t

In this work 9%Cr alloys were designed supported by computational thermodynamic methods. Two sets ofalloys were produced: 9%Cr alloys with 0.1%C and 0.05%C and 9%Cr alloys containing ∼0.03% Ti with 0.1%Cand 0.05%C (always wt%). Microstructure investigations showed good agreement with the predictedphases of the thermodynamic modeling. The volume fraction of precipitated M23C6 carbides is directlyrelated to the carbon content of the alloys. For Ti-containing alloys the precipitation of nano-sized Ti-rich

eywords:%Cr steelslloy designicrostructure evolution

TEMreep

MX carbonitrides was observed. The microstructure evolution (sub-grain and particle size) during creepat 650 ◦C/100 MPa was investigated by STEM-HAADF. The sub-grain size evolution and the coarsening ofprecipitates (MX carbonitrides, M23C6 and Laves phase) were more pronounced for Ti-containing alloys.9Cr alloys without Ti and with low carbon content presented the highest creep strength of all investigatedalloys.

© 2011 Elsevier B.V. All rights reserved.

omputational thermodynamics

. Introduction

Martensitic/ferritic 9–12%Cr steels (all values referred in thisork are in wt%) represent important high temperature materi-

ls for fossil-fired power plants operating at temperatures up to00 ◦C [1]. Martensitic/ferritic 9–12%Cr steels are used as struc-ural materials such as steam headers, steam lines, pipes or highemperature boilers and turbine parts. They combine high creeptrength, good ductility and good thermal cycling properties asell as acceptable room temperature properties [2,3]. Their goodroperties at high temperatures are related to their characteristicicrostructures. The martensitic/ferritic matrix of 9–12%Cr steels

ontains high dislocation density and several different internalnterfaces, such as prior austenite grain boundaries, prior packet orlock boundaries, prior martensite lath boundaries and sub-grainoundaries. Depending on the alloy composition different kinds ofrecipitates are present in the microstructure, e.g. M23C6 carbides,X carbonitrides and Laves phase [4]. Due to new environmental

egulations, energy saving requirements and commercial demands,here is a permanent driving force to increase the working temper-ture of fossil-fired power plants in order to improve their thermalfficiency [5]. The increment of the working temperature (up to

∗ Corresponding author. Tel.: +49 30 8062 42764; fax: +49 30 8062 42047.E-mail address: jose.garcia@helmholtz-berlin.de (J. Garcia).

921-5093/$ – see front matter © 2011 Elsevier B.V. All rights reserved.oi:10.1016/j.msea.2011.03.037

650 ◦C) combined with the exposure at stress during service pro-motes microstructural changes in the microstructure of 9–12% Crheat resistant steels reducing their creep strength [6]. Coarseningof precipitates (e.g. Laves, M23C6) and precipitation of undesirablephases (e.g. Z-phase) decrease the strength of the steels duringcreep. The particle coarsening process follows the Ostwald ripen-ing mechanism [7]. During ripening, the average precipitate sizeincreases with the time at elevated temperatures. As a result, thespacing of the precipitates – in particular on dislocations – increasesand the particle hardening is diminished [8].

According to Blum and Eisenlohr [9] the recovery of free dislo-cations may control the creep rate during the primary creep, butthat control probably shifts during creep to sub-grain boundaryprocesses. In particular, the migration of sub-grain boundaries maycontrol the recovery of both free and boundary dislocations. Conse-quently sub-grain boundary hardening is one of the most importantstrengthening mechanisms in 9–12%Cr martensitic/ferritic steels,which is enhanced by fine dispersions of precipitates along theinternal interfaces [10].

In previous investigations, the authors designed and producedtwo 12%Cr heat resistant steels for creep conditions at 650 ◦C and

100 MPa [4]. The microstructure evolution of these 12%Cr steelswas investigated during creep at 100 MPa/650 ◦C/8000 h focusingon quantitative investigations of precipitates by scanning electrontransmission microscopy (STEM). In a further work the influence ofhot-deformation and tempering temperature was studied by quan-

D. Rojas et al. / Materials Science and Engineering A 528 (2011) 5164–5176 5165

Table 1Nominal chemical composition of the master alloy and reference steel P92 (wt%) [17].

N

0.000.05

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d6cp

caatrt

onocbabc

mtwam

2

2

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in TEM were carried out in the bright field (BF) and scanning mode(STEM). Precipitates were identified by a combination of electrondiffraction patterns (DP) and energy dispersive spectroscopy (EDS)analysis, to avoid ambiguous identification of similar precipitates.

TA

Alloy B C Co Cr Mn

Master alloy 0.008 0.08 1.00 9.00 0.50P92 0.001 0.1 – 9.0 0.45

itative determination of dislocation density and sub-grain sizeSTEM-HAADF) at the initial stage (after tempering) and after short-erm creep at 80–250 MPa/650 ◦C [11]. The experimental STEMuantification was compared to thermodynamic modeling of dif-usion controlled transformations (DICTRA) in order to investigatehe coarsening on the MX and M23C6 precipitates at long-term con-itions [12] as well as to study the growth kinetic of the Laves phase13].

In these works, creep tests of 12%Cr martensitic/ferritic steelst temperatures above 600 ◦C showed increased creep strengthn the first few thousand hours [7]. However, prolonged testinghowed unexpected breakdowns on creep strength. Cipolla et al.14] explained this behaviour by transformation of fine MX in aast growing Z-phase. Coarse Z-phase provides fewer obstacles forinning of dislocations reducing severely the creep strength of thelloys. Danielsen and Hald [15] also observed that the driving forceor the nucleation of Z-phase is lower for Cr contents below 10 wt%.

Recently many investigations have been carried out on theevelopment of new 9%Cr steels for working temperatures above00 ◦C [16]. Abe et al. designed martensitic 9Cr–3W–3Co–B–Nreep steels [17], which showed enhanced creep properties com-ared to the well known P91 and P92 steels [18–20].

Although the oxidation resistance is reduced by the lower Crontent, the steam oxidation resistance of the 9–10%Cr steels at 600nd 650 ◦C is enough for large dimensional component parts, suchs rotors, which present a small surface to volume ratio [21]. Forhin-walled components, such as boiler tubes, the steam oxidationesistance of the 9–10%Cr steels may be disadvantageous, due toheir high oxidation rates [22].

In the present work a 9%Cr master alloy has been designed basedn a combination of physical metallurgy principles and thermody-amic modeling (Table 1). The novelty of this alloy compared tother alloys of previous works is the reduction of the W and Coontent and the achievement of a good balance between the car-ide and carbonitride former elements (e.g. Nb, V and Ti) as wells a good balance between B and N to avoid the formation of largeorides. In addition the effect of Ti additions for different carbonontents (between 0.05 and 0.1%) was investigated.

The main objective of this investigation was to study theicrostructure after tempering and the microstructure evolu-

ion during creep at 650 ◦C by STEM. The microstructure featuresere compared with the macro-mechanical properties (creep

nd hardness) in order to determine correlations between theicrostructure evolution and the observed creep strength.

. Experimental procedure

.1. Thermodynamic modeling

The Thermocalc software has been employed for the design ofhe alloys [23]. The software is linked with various databases and

able 2nalysed chemical composition of the produced alloys (wt%).

Alloy B C Co Cr Mn

9CrTi-H 0.007 0.106 1.01 9.08 0.539CrTi-L 0.008 0.047 1.01 8.90 0.539Cr-H 0.009 0.108 0.99 8.81 0.499Cr-L 0.013 0.051 1.01 8.97 0.48

Nb Si Ti V W Mo

6 0.030 0.35 0.04 0.15 2.00 –0.07 0.04 – 0.20 1.84 0.47

interfaces, where all the thermodynamic information such as Gibbsenergy is stored. Thermocalc has been used for calculations of thephase equilibria and the evaluation of phase stabilities, so that theinfluence of composition (addition of elements) and heat treat-ments on the 9%Cr martensitic/ferritic steels was modeled. Uponmodeling, the time and costs of trial-and-error of conventional alloydevelopment can be reduced [24,25]. All calculations were carriedout with the Thermocalc database TCFe6 [23].

2.2. Alloy production

Four alloys were prepared by vacuum induction melting withmasses of about 4 kg. The final chemical composition of the alloysis shown in Table 2. The samples were hot-rolled, austenitised andtempered with following parameters:

• Hot-rolling at 1150 ◦C with posterior air cooling (66% final defor-mation).

• Austenisation heat treatment for 0.5 h followed by air-cooling.The austenisation temperatures are shown in Table 3 and indi-cated in Figs. 1 and 2.

• Tempering for 2 h at 780 ◦C with subsequent air-cooling (indi-cated in Figs. 1 and 2).

2.3. Microstructure investigation

The microstructure of the alloys was analysed in the initial state(after tempering at 780 ◦C/2 h) and after different creep times torupture by light optical microscopy (LM) and by transmission elec-tron microscopy (TEM, TECNAI Supertwin F20 operating at 200 kV).The microstructure investigations after creep were carried out tak-ing samples about 15 mm from the fracture zone avoiding thenecking area.

For LM investigations, the samples were prepared by mechan-ical polishing, followed by etching with V2A (47.5 vol.% distilledwater, 47.5 vol.% hydrochloric acid, 4.8 vol.% nitric acid and 0.2 vol.%Vogel’s reagent). To prepare the TEM specimens mechanical polish-ing was followed by twin-jet-polishing (TenuPol-5 of Struers) witha solution of acetic acid and perchloric acid as electrolyte (95 vol.%acetic acid, 5 vol.% perchloric acid) at 15 ◦C and 43 V. Observations

Table 3Austenitisation temperatures from Thermocalc.

Alloy 9CrTi-H 9CrTi-L 9Cr-H 9Cr-LAustenitisation T (◦C) 1120 1080 1120 1080

N Nb Si Ti V W

0.008 0.030 0.36 0.030 0.15 1.930.007 0.030 0.36 0.035 0.15 1.900.005 0.030 0.38 – 0.15 2.000.005 0.035 0.37 – 0.15 1.98

5166 D. Rojas et al. / Materials Science and Eng

Fig. 1. Thermocalc phase diagram for alloys 9CrTi-H and 9CrTi-L (F = ferrite andA = austenite, Thermocalc TCFe6). The austenitisation temperature and the temper-ing temperature are indicated in the phase diagram by black circles for each alloy.Ti-MX denotes the Ti-rich phase which contains N, C and few Nb, whereas Nb-MXare Nb-rich particles with C and N and also few amounts of Ti and Cr.

Fig. 2. Thermocalc phase diagram for alloys 9Cr-H and 9Cr-L (F = ferrite andA = austenite, Thermocalc TCFe6). The austenitisation temperature and the temper-ing temperature are indicated in the phase diagram by black circles for each alloy.V-MX is V-rich phase containing Nb, N and C and few Fe and Cr, whereas Nb-MXdenotes Nb-rich particles with C, Cr and N and also few amounts of V.

ineering A 528 (2011) 5164–5176

Equivalent circle diameters of the particles were calculated withan image editor in order to carry out a quantitative analysis of pre-cipitates [26]. More than 120 particles for each type of precipitateswere quantified to ensure reliability of the measurements. In thecase that the carbides present a non-spherical form, two perpendic-ular axes were measured (a and b) and an average diameter (a + b)/2was calculated. Several micrographs from the sample section wereanalysed for each alloy in order to ensure that measurements wererepresentative for the whole material.

For quantitative analysis the error was determined as follows:

error = d ± k1S,

where S is the standard deviation k1 = 1.96/√

n and n is the numberof measured precipitates.

The determination of the sub-grain size at initial state and aftercreep was carried out using the line intersection method (the lathsand blocks can be regarded as elongated sub-grains [27]). SeveralSTEM micrographs were taken through the sample to obtain rep-resentative measurements [28,29]. An array of 6 reference lines,perpendicular to the direction of the elongated sub-grains was setfor each micrograph to measure the sub-grain widths using theAnalySIS 5.0/Olympus soft imaging editor software. The resultswere reported as average values from all measurements togetherwith the error of the average values.

2.4. Creep tests

Tensile creep tests in air at constant temperature of 650 ◦C (±5 K)with constant load between 80 and 175 MPa were carried out todetermine the creep rupture times. Standard cylindrical samplesaccording to DIN 50125 B 4x20 were used with 40 mm gauge lengthand 4 mm diameter.

2.5. Hardness

Vickers hardness measurements (HV10) according to DIN EN ISO6507-1 were carried out using a Universal Wolpert macrohardnessDIA Testor 2n. At least 10 measurements were performed on thesamples in the initial state (after tempering) and after creep (torupture).

3. Results and discussion

3.1. Alloy design

A 9%Cr master alloy to withstand high working conditions (creeprupture lives of 100,000 h under stress of 100 MPa at temperaturesof 650 ◦C) was designed considering basic metallurgical principles(Table 1). A brief description of the effect of the different elementadditions seems to be necessary in order to explain its relevancefor the design of the master alloy.

• 9%Cr was added in order to reduce the driving force for Z-phaseprecipitation. Several authors have demonstrated that high Crcontents increase the driving force for the precipitation of Z-phase [14,15,30]. The Cr content provides the necessary oxidationand corrosion resistance, as well as the strengthening of the mate-rial by precipitation of the M23C6 carbides.

• Between 0.3% and 0.4% of Si has been added to provide furthercorrosion resistance [31].

• To achieve sufficient creep strength at service conditions it isnecessary to improve the microstructural stability and providesufficient solid solution and precipitate strengthening. W hasproved to have a beneficial effect, suppressing the recovery ofthe martensitic matrix and increasing the stability of the precip-

d Engineering A 528 (2011) 5164–5176 5167

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Table 4Volume fractions of precipitates calculated with Thermocalc for alloys 9CrTi-H and9CrTi-L at 780 ◦C and 650 ◦C.

T (◦C) Phases 9CrTi-HVolume fraction (%)

9CrTi-LVolume fraction (%)

780

Ferrite 97.93 99.12M23C6 1.98 0.78Ti-MX 0.06 0.06Nb-MX 0.03 0.04

Ferrite 97.26 98.21

D. Rojas et al. / Materials Science an

itates by decreasing the self-diffusion rate [32,33]. In addition Wis the most potent Laves phase former [4,34] for strengtheningby Laves phase. The W content was set at 2%.V, Nb and Ti in combination with C and N are stable carbide andcarbonitride formers, which provide the necessary precipitationstrengthening due to the slow coarsening rate of the precipitates,hence improving the creep strength [17,35].Many investigations have demonstrated the relevant contribu-tion of Co in martensitic/ferritics steels [4,30,36]. In order tostabilise the austenitic field and avoid the formation of �-ferriteat the austenitisation temperature, the alloy must be balancedwith elements such as Co and Mn. 1% Co was added to the alloy.B stabilises the lath martensitic microstructure by stabilisationof fine M23C6 carbides and by reduction of the Ostwald ripeningrate in the vicinity of prior austenite boundaries [37]. This retardsthe tertiary creep and increases the time to rupture [38]. B andN additions should be balanced carefully to avoid the formationof large BN particles, which offset the beneficial effects of B [39].The B content was adjusted considering B/N ratios which avoidthe formation of BN particles based on previous observations ofAbe et al. [40]. The B content added was 0.008%. According to Ref.[39] the maximum N content in order to avoid the formation ofBN is 0.0125% N.N was added in order to form MX carbonitrides. The N contentwas set at 0.006% to avoid BN formation.

.2. Thermodynamic modeling results

The influence of Ti addition and carbon content was modeledy using Thermocalc. In particular the phase fields at the austeniti-ation, tempering (780 ◦C) and creep temperatures (650 ◦C) are ofnterest. Based on this information the production of the samples

as carried out. The Ti containing alloys are referred to as 9CrTi.he denomination for low and high carbon content is indicated by= low and H = high.

.2.1. 9CrTi-H and 9CrTi-L designThe chemical compositions of the alloys 9CrTi-H and 9CrTi-L

Table 2) were designed to obtain ferrite, Ti-MX and Nb-MX parti-les, M23C6 carbides and Laves phase at a temperature of 650 ◦C.

Ti addition combined with C and N promotes the precipitationf Ti-MX particles. Ti-MX particles have an extremely high stabilitynd thus a high potential for strengthening of martensitic/ferriticteels. As shown in the phase diagram (Fig. 1), Ti-MX precipitatesre more stable than the detrimental Z-phase. They suppress therecipitation of Z-phase above 650 ◦C, e.g. addition of 0.03% Ti isnough to decrease the precipitation temperature of Z-phase below50 ◦C for the range of compositions investigated.

The phase diagram in Fig. 1 indicates that the precipitation of Ti-X particles starts already in the liquid (around 1500 ◦C), even for

mall additions of Ti, and there is nearly no change in the amount ofhis phase with decreasing temperature. Therefore it seemed to beifficult to control the size and distribution of the Ti-MX particles

n the microstructure. Inspite of this, the austenitisation tempera-ures were fixed about 50 ◦C above the precipitation temperature ofhe phase field containing austenite and Nb-MX particles for bothlloys avoiding regions where �-ferrite is a stable phase at higheremperatures (Fig. 1).

The tempering temperature (780 ◦C) was chosen to ductilise theard and brittle martensite transformed during the air-cooling afterempering. The phase diagram shows phase fields with ferrite, Ti-

X, Nb-MX precipitates and M23C6 carbides in both alloys. Nb-MXenotes Nb-rich particles with C and N and also few amounts ofi and Cr. These particles show a high stability and thus a highotential for strengthening. No Laves phase is expected in the initialicrostructure of both alloys.

650M23C6 2.03 0.82Ti-MX 0.06 0.06Nb-MX 0.03 0.03Laves phase 0.62 0.88

The main difference between 9CrTi-H and 9CrTi-L is the amountof C added (0.1% and 0.05%, respectively) which influences the vol-ume fraction of M23C6 precipitates and Laves phase. In Table 4the calculated volume fractions of all precipitates in both alloys atthe tempering temperature (780 ◦C) and at the creep temperature(650 ◦C) are shown.

3.2.2. 9Cr-H and 9Cr-L designAlloys 9Cr-H and 9Cr-L were designed to have ferrite, V-MX,

Nb-MX precipitates, M23C6 carbides and Laves phase at 650 ◦C.Thermocalc calculations (see Fig. 2) show that alloys 9Cr-H and 9Cr-L indeed contain ferrite, V-MX, Nb-MX precipitates, M23C6 carbidesand Laves phase at 650 ◦C. V-MX corresponds to V-rich precipitateswhich contain Nb, N and C and few Fe and Cr, whereas Nb-MX refersto Nb-rich particles with C, Cr and N and also few amounts of V. For-mation of Nb-MX and V-MX particles improves the creep strength[35] due to the slow coarsening rate of these precipitates.

The phase diagram in Fig. 2 shows that Z-phase is more sta-ble compared to V-MX particles at 650 ◦C. In Ref. [14] it wasdemonstrated that V-MX precipitates are gradually transformedinto Z-phase, which leads to an early consumption of the V-MX par-ticles in the region adjacent to the prior austenite grain boundaries,decreasing the creep strength. Inspite of the calculations, severalstudies had reported that 9%Cr steels do not suffer from abundantformation of Z-phase after long-term creep [15,41] due to the slowkinetics of the Z-phase precipitation.

The austenitisation temperatures were chosen about 50 ◦Cabove the phase field containing Nb-MX particles (1120 ◦C and1080 ◦C for the 9Cr-H and 9Cr-L alloy, respectively), in order toobtain a fully austenitic field and ensure a completely martensitictransformation (avoiding �-ferrite see Fig. 2).

The tempering temperature (780 ◦C) was chosen to temper themartensite transformed during air-cooling after the austenisationtreatment. The stable phases present in both alloys were ferrite,M23C6 carbides, V-MX and Nb-MX precipitates. No Laves phase isexpected in the initial microstructure of the 9Cr alloys.

The main difference between 9Cr-H and 9Cr-L lies in the amountof C added (0.1% and 0.05%, respectively) which influences the vol-ume fraction of M23C6 precipitates and Laves phase. The volumefractions for all precipitates were calculated at the tempering tem-perature 780 ◦C and at the creep temperature 650 ◦C (Table 5).

3.3. Microstructure evolution

Light microscopy of the investigated alloys after heat treatmentshows a martensitic/ferritic matrix and precipitates. STEM investi-gations were carried out to quantify the microstructure features in

the initial stage and after creep (to rupture).

3.3.1. Alloys 9CrTi-H and 9CrTi-L3.3.1.1. Initial microstructure. In the initial condition (after temper-ing at 780 ◦C/2 h) both alloys presented a martensitic/ferritic matrix

5168 D. Rojas et al. / Materials Science and Engineering A 528 (2011) 5164–5176

Table 5Volume fractions of precipitates calculated with Thermocalc for alloys 9Cr-H and9Cr-L at 780 ◦C and 650 ◦C.

T (◦C) Phases 9Cr-HVolume fraction (%)

9Cr-LVolume fraction (%)

780

Ferrite 97.89 99.02M23C6 2.05 0.92V-MX 0.03 0.02Nb-MX 0.03 0.04

650

Ferrite 97.14 98.02M23C6 2.11 0.96Z-phase 0.06 0.06Nb-MX 0.02 0.03Laves phase 0.67 0.93

Table 6Quantitative determination of sub-grain size and hardness at initial stage.

Alloy Sub-grain size (nm) HV10

9CrTi-H initial 401 ± 44 246 ± 2

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aat

Table 7Average size of precipitates in alloys 9CrTi-H and 9CrTi-L (time in hours and size innanometers) under creep condition (650 ◦C/101 MPa).

9CrTi-H 0 h 7253 h

M23C6 79 ± 5 97 ± 5Ti-MX 30 ± 1 34 ± 2Nb-MX 29 ± 2 33 ± 3Laves 0 411 ± 32

9CrTi-L 0 h 2154 h

M23C6 89 ± 4 106 ± 8

F9

9CrTi-L initial 433 ± 25 223 ± 19Cr-H initial 426 ± 27 248 ± 39Cr-L initial 403 ± 33 221 ± 2

ith high density of interfaces, such as prior austenite grain bound-ries, prior lath boundaries and sub-grain boundaries, as well asigh dislocation density. The high dislocation density is producedhen martensite forms during air-cooling after the austenisationeat treatment. During tempering, the precipitation of solute atomsccurs, as well as recovery of the dislocation cell structure, resultingn a sub-grain structure [11].

STEM investigations of the initial microstructure were carriedut in order to measure the sub-grain size (see Table 6) because sub-rain boundary hardening is one of the most important strength-ning mechanisms for this kind of materials. The average measuredub-grain size for alloy 9CrTi-H was slightly smaller than that forlloy 9CrTi-L. Hence the sub-grain formation is related to both the

artensitic transformation and the tempering temperature.The initial microstructure of both alloys showed M23C6 carbides

nd MX particles (Nb-rich and Ti-rich precipitates) which are ingreement with the thermodynamic equilibrium calculation at theempering temperature (Fig. 1, 780 ◦C).

ig. 3. STEM-HAADF micrographs of initial microstructure of alloy 9CrTi-H (A) and alloyCrTi-L shows some large particles rich in W and Fe (possibly FeW2B).

Ti-MX 57 ± 6 65 ± 4Nb-MX 27 ± 2 28 ± 2Laves 0 347 ± 40

The M23C6 carbides were the most abundant precipitates in themicrostructure (Fig. 3). The M23C6 precipitates were mostly placedon prior austenite grain boundaries and lath or block boundaries.At such preferential sites, the effective surface energy is lower, thusthe free energy barrier is diminished and nucleation is facilitated[42].

The quantity of M23C6 carbides (white arrows, Fig. 3a and b)observed in alloy 9CrTi-L was much smaller than in alloy 9CrTi-H,which clearly reflects the effect of the carbon content in this alloy.The microstructure observations are consistent with the thermo-dynamic equilibrium calculations (Table 4). The average size of theM23C6 precipitates was slightly smaller in alloy 9CrTi-H than inalloy 9CrTi-L (Table 7).

Few large W and Fe rich particles (about 500 nm) were foundin alloy 9CrTi-L (Fig. 3b). According to Ref. [33] these particles mayprobably correspond to undissolved borides (FeW2B). Such parti-cles were not found in alloy 9CrTi-H probably due to the higheraustenitisation temperature (1120 ◦C), which allows further disso-lution of borides.

Nano-sized Nb-MX particles with spheroidal shape and Ti-MXprecipitates with rhomboidal shape were detected in both alloys

(Figs. 4 and 5). TEM-EDS measurements showed that the Nb-MXprecipitates are Nb-rich particles which contain C and N and fewamounts of Ti and Cr, whereas Ti-MX are Ti-rich precipitates whichcontain N, C, Nb and Cr. The average size of Nb-MX and Ti-MXparticles for alloy 9CrTi-H and 9CrTi-L are shown in Table 7.

9CrTi-L (B). White arrows indicate M23C6 precipitates in both micrographs. Alloy

D. Rojas et al. / Materials Science and Engineering A 528 (2011) 5164–5176 5169

Fig. 4. STEM-HAADF micrograph of the initial microstructure of alloy 9CrTi-H (A) Nano-sized Nb-MX precipitates are indicated by white arrows; black arrows indicate Ti-MXparticles. EDS spectrum of the encircled Nb-MX precipitate (B).

F winga

genb

sThqfiulTrTstttsp

ture (Table 8). However, the martensite lath structure could still bedistinguished (Fig. 7). In both cases the sub-grain sizes were morethan twice the size of the sub-grain size in the initial microstruc-ture.

Table 8Sub-grain size and hardness of studied alloys after creep at 650 ◦C.

ig. 5. STEM-HAADF micrograph of initial microstructure of alloy 9CrTi-L (A) shorrows. EDS spectrum of the encircled Ti-MX particle (B).

The Nb-MX and Ti-MX particles were mostly formed on sub-rain boundaries and within the sub-grains. According to Taneiket al. [43] the nano-sized MX particles precipitate more homoge-eously than M23C6 carbides or Laves phase due to the small misfitetween the crystallographic structure of the MX with the matrix.

The Ti-MX particles in alloy 9CrTi-L showed a larger particleize compared to the Ti-MX particle size in alloy 9CrTi-H (Table 7).he difference in the mean particle size may be related to theigher austenitisation temperature of alloy 9CrTi-H. As a conse-uence higher amounts of large Ti particles are dissolved and aner precipitation of Ti-MX carbonitrides due to higher supersat-ration during subsequent tempering is reached. Consequently no

arge TiN particles were found in alloy 9CrTi-H, whereas a few largeiN precipitates (700 nm) were observed in alloy 9CrTi-L (Fig. 6). Aseported by Gustafson [44] there are two size distributions of thei-MX in the steels, the primary and the secondary. In this work theizes of the small particles were measured, under the assumption

hat the (large) primary particles are so sparsely distributed thathey will not affect the coarsening of the secondary ones, becausehe coarsening of the large particles is expected to follow a muchlower process and has no important influence on the mechanicalroperties [44].

Ti-rich MX precipitates (white arrows) M23C6 precipitates are indicated by black

The average hardness for both alloys is shown in Table 6. Themean hardness value for alloy 9CrTi-H was slightly larger than thehardness measured for alloy 9CrTi-L, which is consistent with thecarbon content variation, corresponding to secondary hardening(precipitation hardening).

3.3.1.2. Microstructure after creep. For both alloys the microstruc-ture after creep showed large sub-grains as a result of theextensive deformation and to the recovery of the sub-grain struc-

Alloy Stress (MPa) Rupture time (h) Sub-grain size (nm) HV10

9CrTi-H 101 7,253 821 ± 69 224 ± 29CrTi-L 101 2,154 1011 ± 94 194 ± 39Cr-H 101 7,987 689 ± 68 244 ± 29Cr-L 125 10,168 647 ± 46 220 ± 3

5170 D. Rojas et al. / Materials Science and Engineering A 528 (2011) 5164–5176

Fig. 6. STEM-HAADF micrograph of the initial microstructure of alloy 9CrTi-L (A) showing a large Ti-rich precipitate and the M23C6 precipitates (white arrows). EDS spectrumof the Ti-rich particle (B).

F 101 M( black

ssao

l

wbsvkott

choice of the εf value could lead to overestimation of the calcu-lated sub-grain size compared to the experimental results. Thehardness values were consistent with the enlargement of the sub-grain size showing a decrease in the average hardness in both cases(Table 8).

Table 9Results of calculated evolution law of sub-grain size [44].

Alloy Rupture strain, εf Steady-statesub-grain size, �∞

Calculatedsub-grain size (nm)

ig. 7. STEM-HAADF micrograph of alloy 9CrTi-H (A) after creep (7253 h/2154 h/101 MPa/650 ◦C) with M23C6 precipitates (white arrows) and Laves phase (

The measured sub-grain sizes were compared to the predictedub-grain size determined using the “evolution law of sub-grainize” equation suggested by Blum and Götz [45]. Several authorspplied the resulting phenomenological equation to the evolutionf sub-grain sizes in 9–12%Cr steels [6,18,19,27].

As reported in Ref. [45] sub-grain sizes, �, evolve with strain:

og � = log �∞ + log(

�0

�∞

)exp

(−ε

klog �

)(1)

here � represents the sub-grain size and �0 is the sub-grain sizeefore creep exposure. �∞ is defined as the steady-state sub-grainize (�∞ = 10 [Gb/�] where G is the shear modulus, b is the burger

ector and � is the applied stress). ε is the accumulate strain andlog� is a constant with a value of 0.12 [45]. Table 9 shows the resultsf Eq. (1) considering values of G = 58 GPa and b = 0.248 nm [46] forhe investigated alloys. Good agreement was found by comparinghe measured and calculated values (compare Tables 8 and 9).

Pa/650 ◦C) and STEM-HAADF micrograph of alloy 9CrTi-L (B) after creeparrows).

It is important to note that the value of ε used for the calcu-lation in Eq. (1) is the rupture strain (εf), hence the strain in thehomogeneously deformed section of the gauge is not known. The

(nm)

9CrTi-H 0.116 1424 8799CrTi-L 0.254 1424 12349Cr-H 0.680 1424 7189Cr-L 0.119 1147 778

D. Rojas et al. / Materials Science and Engineering A 528 (2011) 5164–5176 5171

F 50 ◦CL

wfiLmsppdtmaT

abTcbi

tcbdpTpLa

ig. 8. STEM-HAADF micrograph of alloy 9CrTi-L (A) after creep (2154 h/101 MPa/6aves phase particle (C).

For both alloys the observed phases are in good agreementith the Thermocalc calculations at 650 ◦C, with stable phaseelds containing M23C6 carbides, Ti-MX and Nb-MX particles andaves phase. Inspite of V additions, no V-MX was observed in theicrostructure (in the initial state and after creep). This result

uggests, together with the Thermocalc calculations, that Ti-MXrecipitates are more stable than V-MX at 650 ◦C. Regarding to therecipitation of Z-phase, Cipolla et al. [14] demonstrated that theevelopment of the Z-phase is associated with the consumption ofhe finely dispersed small V-MX particles by uptake of Cr from the

atrix. The absents of V-MX and Z-phase in the initial conditionnd after creep together with Cipolla’s affirmation suggest that thei-MX suppress the formation of Z-phase at 650 ◦C.

M23C6 precipitates (Fig. 7, white arrows) were still the mostbundant particles in both alloys. The average size of M23C6 car-ides in alloy 9CrTi-H is smaller than in alloy 9CrTi-L (compareable 7). The lath or block boundaries were often pinned by M23C6arbides. In Fig. 7b it was observed that the recombination of lathoundaries caused the disappearance of some lath boundaries leav-

ng a row of M23C6 carbides in the matrix (see encircled area).Larger Laves phase particles (black arrows) were detected near

he M23C6 carbides, as shown in Fig. 7a and b. Laves phase parti-les were observed on prior austenite grain boundaries and lath orlock boundaries. Laves phase formed and grew under creep con-ition after several hundred hours. Fig. 8 shows an example of the

article identification procedure. The main elements detected byEM-EDS in the Laves phase were W, Fe and some Cr. The averagearticle sizes of Laves phase for alloy 9CrTi-H and for alloy 9CrTi-are shown in Table 7. The larger size of Laves phase particles forlloy 9CrTi-H may be related to the longer creep times (7253 h creep

Fig. 9. Tensile creep curves comparing the creep strength of alloys 9CrTi-H an

), diffraction pattern of the encircled Laves phase particle (B), and EDS spectrum of

at 650 ◦C/101 MPa compared to 2154 h creep at 650 ◦C/101 MPa forsample 9CrTi-L).

Ti-MX particles in alloy 9CrTi-H showed a low coarsening rate(13% of growth after 7253 h) compared to the Ti-MX in alloy 9CrTi-L (14% of growth after 2,154 h). An explanation for the differentcoarsening behaviour of Ti-MX precipitates may be related to a lessuniform particle size of such precipitates in alloy 9CrTi-L (Fig. 5)compared to the Ti-MX particles in 9CrTi-H (Table 7), which favoursthe coarsening process. The smaller particles have a higher surfaceto volume ratio than the larger particles, thus smaller particles areless stable than larger particles of the same phase. An increase inthe mean particle size will thus reduce the total free energy of thesystem and this reduction in free energy is the driving force for thecoarsening process.

The sub-grain enlargement as well as the coarsening of the pre-cipitates is directly related to the creep strength. Under long-termconditions the precipitates coarsen with decreasing particle num-ber during creep deformation leading to a decrease in the pinningof sub-grain boundaries.

Alloy 9CrTi-H shows decreased minimum creep rate andincreased time to rupture compared to 9CrTi-L, as shown in Fig. 9.This result suggests an enhanced stabilisation of fine M23C6 car-bides in alloy 9CrTi-H, hence a large pinning force for boundarymigration is maintained up to long times so that the onset of ter-tiary creep is retarded to longer times. This effectively decreases the

minimum creep rate and increases the time to rupture as shownin Fig. 9b. Creep tested samples showed a reduction of hardness,which correlates with the observed enlargement of sub-grain sizeand the decrease of dislocation density; both softening the alloy(Tables 6 and 8).

d 9CrTi-L at 101 MPa/650 ◦C: (A) strain vs. time and (B) creep vs. strain.

5172 D. Rojas et al. / Materials Science and Engineering A 528 (2011) 5164–5176

y 9Cr-H (A) and alloy 9Cr-L (B) with M23C6 precipitates (white arrows).

33aoapliafh

Mf

gomicha

Table 10Average size of precipitates in alloys 9Cr-H and 9Cr-L (time in hours andsize in nanometers) under creep condition (9Cr-H at 650 ◦C/101 MPa; 9Cr-L at650 ◦C/125 MPa).

9Cr-H 0 h 7,987 h

M23C6 78 ± 3 103 ± 6V-MX 30 ± 1 31 ± 1Nb-MX 29 ± 2 31 ± 1Laves 0 379 ± 48

9Cr-L 0 h 10,168 h

M23C6 97 ± 4 112 ± 8

Fig. 10. STEM-HAADF micrographs of initial microstructure of allo

.3.2. Alloy 9Cr-H and 9Cr-L

.3.2.1. Initial microstructure. A martensitic/ferritic matrix withhigh density of interfaces and a high dislocation density was

bserved in the initial microstructure of both alloys. Sub-grain sizend hardness values are shown in Table 6. During tempering, therecipitation of solute atoms occurs, as well as recovery of the dis-

ocation cell structure, resulting in a sub-grain structure [20]. Thenitial sub-grain size for both samples is very similar, though theustenitisation temperature was 40 ◦C higher for alloy 9Cr-H thanor alloy 9Cr-L. This suggests that the austenitisation temperatureas not a dominant effect on the sub-grain formation.

The observed precipitates (Nb-MX, V-MX precipitates and23C6 carbides) are in good agreement with the Thermocalc results

or the tempering temperature (Fig. 2, 780 ◦C).The M23C6 precipitates were mostly placed on prior austenite

rain boundaries and lath or block boundaries. The average sizesf the M23C6 carbides for both alloys are shown in Table 10. The

easured average size of the M23C6 carbides was slightly smaller

n alloy 9Cr-H than in alloy 9Cr-L. In agreement with the Thermocalcalculations (Table 5) alloy 9Cr-H with high C content presented aigher volume fraction of M23C6 carbides compared to the low Clloy (Fig. 10a and b).

Fig. 11. STEM-HAADF micrograph of the initial microstructure of alloy 9Cr-L (A) sho

V-MX 30 ± 2 31 ± 2Nb-MX 25 ± 2 29 ± 3Laves 0 354 ± 42

Nano-sized Nb-MX particles with a spheroidal shape and V-MXprecipitates with a plate-like shape were identified in both alloys

(see Figs. 11 and 12). TEM-EDS measurements showed that the Nb-MX precipitates are Nb-rich particles which contain C, N, V andCr. V-MX correspond to V-rich precipitates which contain N, C, Nband Cr. The mean particle size of the Nb-MX and V-MX precip-

wing Nb-MX particles and EDS spectrum of the encircled Nb-MX particle (B).

D. Rojas et al. / Materials Science and Engineering A 528 (2011) 5164–5176 5173

F ith V-ME

iatamta

3tiwaastta(

(tttd

ppewtp

iTatbMFEt

n

aries and within the sub-grains. V and Nb MX precipitates showedvery slow coarsening rates in both alloys (Table 10). This obser-vation suggests that both MX precipitates act as obstacles for thesub-grain boundaries, thus the fine distribution of V and Nb MX

ig. 12. STEM-HAADF micrographs of the initial microstructure of alloy 9Cr-L (A) wDS spectrum of the encircled V-MX particle (B).

tates in alloys 9Cr-H and 9Cr-L is shown in Table 10. The Nb-MXnd V-MX particles were mostly located within the sub-grains or athe sub-grain boundaries pinning dislocations or sub-grain bound-ries (Fig. 11). As explained in Section 3.3.1, MX particles precipitateore homogeneously than M23C6 carbides or Laves phase due to

he small misfit between the crystallographic structures of the MXnd the matrix.

.3.2.2. Microstructure after creep. For both alloys the microstruc-ure after creep showed large sub-grains. For alloy 9Cr-L thencrease in sub-grain size after 7,987 h creep at 650 ◦C/100 MPa

as slightly smaller compared to the sub-grain development inlloy 9Cr-H after 10,168 h creep at 650 ◦C/125 MPa (Table 7). Goodgreement was found by comparing the experimental results ofub-grain size (Table 8) to the calculated values obtained by usinghe evolution law of sub-grain size (Table 9). It is important to notehat the steady-state sub-grain size (�∞) is ca. 19.5% smaller forlloy 9Cr-L than for alloy 9Cr-H, due to the higher applied stress�).

The hardness values in both cases remain almost constantTable 8). As mentioned before the Vickers hardness correlates withhe sub-grain size and the dislocation density in the material. Forhe 9Cr alloys the increase in sub-grain size was smaller than forhe 9CrTi alloys. This effect would suggest an effective pinning ofislocations by precipitates for these alloys (Fig. 13).

Both crept alloys contained M23C6 carbides, V-MX and Nb-MXarticles and Laves phase after creep (Table 10). The observedhases after creep are in good agreement with the thermodynamicquilibrium calculation except for the Z-phase phase. The Z-phaseas not detected in the microstructure after creep probably due to

he slow precipitation kinetics of the Z-phase in the 9%Cr steels, asreviously described by Danielsen and Hald in Ref. [15].

The M23C6 particles were mostly placed on the prior austen-te boundaries and on lath or block boundaries (Fig. 14a and b).he particle size of M23C6 carbides in alloy 9Cr-L showed a rel-tively slow coarsening rate after 10,168 h/125 MPa compared tohe M23C6 carbides of the other alloys. As example the M23C6 car-ides showed 19% growth in alloy 9Cr-L, whereas in alloy 9Cr-H the23C6 particles showed 32% growth after creep (7987 h/101 MPa).

ig. 15 shows an example of the M23C6 carbides identification. TheDS measurement showed Cr, W and Fe as main components ofhis carbide.

Large Laves phase particles (see Fig. 14) were often observedear the M23C6 carbides (black arrows) which are placed on prior

X particles (white arrows) and Nb-MX particles (black arrows) and corresponding

austenite grain boundaries and lath or block boundaries. Measure-ments of Laves phase particle sizes after creep suggest a lowergrowth of Laves phase in alloy 9Cr-L than in alloy 9Cr-H. An expla-nation for this behaviour may be related to the competitive growthbetween M23C6 carbides and the Laves phase. Alloy 9Cr-H pre-sented a high amount of M23C6 carbides which nucleate and growthon prior austenite grain boundaries and lath or block boundaries,which are the same preferred nucleation sites as for the Lavesphase. This suggests that M23C6 carbides nucleate and growth firston this low energetic sites and restrict the nucleation of Lavesphase. Less nucleus of Laves phase particles with a fast growthbehaviour are expected for high carbon alloys such as the 9Cr-H.

V and Nb MX particles were mostly placed on sub-grain bound-

Fig. 13. STEM-HAADF micrograph of sample 9Cr-L after 10,168 h creep at650 ◦C/125 MPa (inversed contrast). Black arrows indicate Laves phase particles,white arrows indicate the M23C6 carbides. Sub-grains and dislocations are oftenpinned by the M23C6 carbides.

5174 D. Rojas et al. / Materials Science and Engineering A 528 (2011) 5164–5176

Fig. 14. STEM-HAADF micrograph of alloy 9Cr-H (A) after creep (7987 h/101 MPa/650 ◦C) and STEM-HAADF micrograph of alloy 9Cr-L (B) after creep(10,168 h/125 MPa/650 ◦C). White arrows indicate M23C6 precipitates; black arrows indicate Laves phase.

Fig. 15. (A) STEM-HAADF micrograph of alloy 9Cr-L after creep (10,168 h/125 MPa/650 ◦C), (B) diffraction pattern of the encircled M23C6 particle and (C) EDS spectrum ofthe encircled M23C6 particle.

Fig. 16. Tensile creep curves of alloy 9Cr-H at 101 MPa/650 ◦C and alloy 9CrTi-L at 125 MPa/650 ◦C: (A) strain vs. time and (B) creep vs. strain.

D. Rojas et al. / Materials Science and Eng

Fftr

ps

1st

p1

3

it

soc

tmlobop

(irtt

4

9tswc

ig. 17. Results of the tensile creep tests at 650 ◦C showing time to rupture as aunction of applied stress for the four investigated alloys. The alloy 9Cr-L showshe highest creep strength. Corresponding data for the P92 steel [17] are shown aseference.

articles may effectively exerts pinning force for the migration ofub-grain boundaries up to long times during creep.

Fig. 16 shows tensile creep curves of alloy 9Cr-H at01 MPa/650 ◦C and alloy 9CrTi-L at 125 MPa/650 ◦C. The figurehows that the minimum creep rate in both alloys is similar, despitehe higher tensile load in sample 9CrTi-L.

Alloy 9Cr-L presented fine and stable M23C6 carbides and MXrecipitates, as well as the lowest Laves phase particle size after0,168 h at 650 ◦C/125 MPa.

.4. Creep tests

Creep test results are shown in Fig. 17 for all investigated alloys.Alloy 9Cr-L showed the best creep performance of all alloys

nvestigated (10,168 h at 650 ◦C/125 MPa). Alloy 9CrTi-L showedhe lowest time to rupture (2154 h at 650 ◦C/101 MPa).

The microstructure investigations revealed that the 9Cr-L pre-ented the lowest sub-grain growth (Table 8), very slow coarseningf MX carbonitrides and the lowest Laves phase coarsening rateompared to all other alloys investigated (Tables 7 and 10).

The measured small sub-grain sizes suggest that the precipi-ates may provide a large pinning force that reduced the boundary

igration up to longer times, thus the tertiary creep is retarded toonger times. This may be an explanation for the high creep strengthf alloy 9Cr-L (Fig. 17), together with the relatively high amount oforon compared to the other investigated alloys. The effect of boronn 9–12%Cr alloys has been already discussed by the authors in arevious work [4].

The 9CrTi-L alloy presented the largest sub-grain size growthTables 6 and 8), large TiN particles (700 nm) and some FeW2B-likenclusions. The Ti-MX particles showed a relatively high coarseningate after only 2154 h (Table 7) compared to the other alloys inves-igated. Both effects reduce the effective pinning of dislocations byhe precipitates for the 9CrTi-L alloy.

. Conclusions

In the present work the microstructure of four newly designed

%Cr heat resistant alloys before and after creep was inves-igated by STEM-HAADF with respect to the evolution of theub-grains and their precipitate distribution. Two sets of alloysere studied: 9Cr alloys with high (9Cr-H/0.1%C) and low carbon

ontents (9Cr-L/0.05%C) and 9Cr alloys containing ∼0.03Ti% also

ineering A 528 (2011) 5164–5176 5175

with high (9CrTi-H/0.1%C) and low (9CrTi-L/0.05%C) carbon con-tents. Correlations between the microstructure evolution and themacro-mechanical properties were studied. The conclusions of thestudy are summarised as follows:

• Thermocalc calculations showed to be a reliable tool for alloydevelopment of heat resistant steels. Processing parameters(austenitisation and tempering temperatures) were definedbased on the phase diagram information. Investigations of themicrostructure showed good agreement with the predictedphases of the thermodynamic modeling.

• As predicted by the thermodynamic modeling, no Laves phaseprecipitates were found in the initial microstructure for the fourinvestigated alloys.

• Precipitation hardening of the Ti-containing alloys was achievedby precipitation of fine dispersed Ti-based MX particles withrhomboidal shape (alloy 9CrTi-H).

• Few large TiN precipitates (700 nm) were observed in the 9CrTi-Lalloy. For the high carbon 9CrTi alloy no large TiN particles wereobserved. The large Ti-MX formed in the melt must be avoidedbecause they reduced considerably the creep strength.

• The 9Cr alloy with reduced carbon content (∼0.05%) showed bet-ter distribution of M23C6 and MX precipitates, as well as minimalcoarsening of Laves phase after creep. Reducing the C contentreduced the volume fraction of M23C6, so that more nucleationsites for a fine dispersion of Laves phase were present. This sit-uation may be favourable to avoid the formation of large Lavesphases which are detrimental for the creep strength of the alloy.

• In 9CrTi alloys the average sub-grain size after creep was morethan twice the size in the initial microstructure. The hardnessvalues were consistent with the increase of the sub-grain sizeand experienced a decrease after creep exposure.

• For all investigated alloys, good agreement was found betweenthe measured sub-grain sizes and the calculated sub-grain sizes(evolution law of sub-grain size equation [45]).

• Nano-sized Nb-MX particles with a spheroidal shape and V-MXparticles with a plate-like shape were observed in alloy 9Cr-Hand alloy 9Cr-L at the initial stage and after creep. Precipitatesshowed low coarsening rates and were mostly placed within thesub-grains or at the sub-grain boundaries frequently pinning thedislocations or sub-grains boundaries.

• Alloy 9Cr-L showed fine and stable M23C6 carbides and MX pre-cipitates, as well as the minimal coarsening of Laves phase after10,168 h at 650 ◦C/125 MPa.

• Alloy 9Cr-L showed the highest creep strength of all the investi-gated alloys.

Acknowledgements

The authors thank Mr. G. Bialkowski (Max Planck Institute fürEisenforschung GmbH) for carrying out creep test experiments.Dr. J. Garcia gratefully acknowledges financial support from thejoint research group Microstructural Analysis (Helmholtz-ZentrumBerlin für Materialien und Energie/Ruhr Universität Bochum).

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