Microstructure evolution and mechanical properties of Inconel 740Hduring aging at 750 1C

Chong Yan a,b,n, Liu Zhengdong b, Andy Godfrey a, Liu Wei a, Weng Yuqing b

a School of Materials Science and Engineering, Tsinghua University, Beijing 100084, Chinab Central Iron&Steel Research Institute, Beijing 100084, China

a r t i c l e i n f o

Article history:Received 6 May 2013Received in revised form16 September 2013Accepted 18 September 2013Available online 1 October 2013

Keywords:Inconel 740HMicrostructure evolutionγ′ precipitatesMechanical propertiesThree-dimensional atom probe

a b s t r a c t

The microstructure evolution of Inconel 740H during aging at 750 1C for up to 3000 h was investigated bymeans of optical microscopy (OM), scanning electron microscopy (SEM), transmission electron micro-scopy (TEM), small-angle X-ray scattering (SAXS), three-dimensional atom probe (3DAP) analysis andmicro-phase analysis. The mechanical properties of samples after aging were also studied. The grain sizeincreased substantially during the whole aging period. Two different types of grain boundary carbideswere observed; block-shaped and needle-shaped. Both were identified to be M23C6 by selected areadiffraction measurements. The grain boundary carbides did not coarsen significantly during aging, withthe weight fraction increasing only from 0.20% to 0.28%. In contrast, a much higher coarsening rate ofγ′ precipitates was observed, as evidenced from both TEM observations and SAXS analysis. 3DAP wasused to study the elemental partitioning behavior between γ′ precipitates and γ matrix as well as theevolution in width of the γ/γ′ interface. A large increase in the width of γ/γ′ interfaces was seen between1000 h and 3000 h aging. In addition, for the sample aged at 750 1C for 3000 h, Cr enrichment on theγmatrix side of the γ/γ′ interface was found. Tensile tests at 750 1C of the aged samples showed a gradualdecrease in elevated-temperature yield strength after 500 h, when this alloy was over-aged. The criticalprecipitate size for the transition from precipitate cutting by weakly coupled dislocations to stronglycoupled dislocations for Inconel 740H was calculated to be approximately 50 nm, which agrees well withthe experiment measurements of elevated-temperature yield strength. The room temperature impacttoughness of all samples decreased during aging as the grain size kept growing.

& 2013 Elsevier B.V. All rights reserved.

1. Introduction

As a result of increasing energy demands and acceleratedenvironmental problems, there is an urgent need to improve thethermal efficiency of coal-fired power plants. To achieve this goal,next generation ultra-supercritical power plants, operating atsteam temperatures of 700–750 1C and operating pressure up to37.5 MPa, are under development [1–3]. However, the increasedoperating parameters place more stringent requirements on theproperties of candidate materials and cannot be met by conven-tional ferritic and austenitic steels. Consequently, there are effortsto replace these materials by Ni-base superalloys, which show alonger creep rupture life and higher corrosion resistance.

Ni-based superalloys have already found widespread applica-tions in a number of critical technological areas, especially thoseinvolving high temperatures, such as jet-engine turbines and coal-fired power plants. These alloys typically exhibit an excellent

balance of properties, including good mechanical properties andductility (both at room temperature and elevated temperatures),improved fracture toughness and fatigue resistance, as well asenhanced creep and oxidation resistance at high temperatures [4].One of the most promising candidate Ni-base superalloys for themain steam pipe of 700 1C ultra-supercritical coal-fired powerplants is Inconel 740H, which is a modified version of Inconel 740developed by Special Metals Corp, Huntington, West Virginia, USA.The nominal composition of Inconel 740H is 25Cr, 20Co, 0.5Mo,1.5Nb, 1.4Ti, 1.4Al, 0.3Mn, 0.03Fe, 0.03C and Ni balance (wt%).Compared with Inconel 740, the ratio of Ti to Al in Inconel 740H islowered in order to eliminate microstructure instabilities foundin Inconel 740 during prolonged thermal aging at 750 1C [5–11].In addition, the Nb content is also reduced to improve the weldabilityof the alloy. Since Inconel 740H has only recently been developed,few studies concerning the microstructure evolution and mechanicalproperties have been carried out so far on this material. Moreover, noreports exist on the detailed microstructure evolution during aging ofInconel 740H at the atomic scale.

Microstructural stability is of prime importance for alloys to beused at high temperatures for long periods of time. Inconel 740H

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n Corresponding author. Tel.: þ86 10 6277 3460; fax: þ86 10 6277 2853.E-mail address: chongyan871210@sina.com (C. Yan).

Materials Science & Engineering A 589 (2014) 153–164

consists of a γ matrix with a coarse grain size, containing adispersion of γ′ precipitates inside the grains, as well as carbideslocated at the grain boundaries. The microstructural evolutionduring long time aging, including grain size, γ′ precipitates andgrain boundary carbides, should be studied at different scales indetail. In addition to traditional optical microscopy and electronmicroscopy observation, three-dimensional atom probe (3DAP)has been used in this study for the first time. In the last decade,there have been large numbers of studies employing 3DAP tocharacterize the microstructure of Ni-base superalloys at theatomic scale, including the size, morphology and composition ofγ′ precipitates within the γ matrix as a function of varied heattreatments [4,12–15]. In particular, the capability of 3DAP toprovide information on elemental partitioning between phases,and on segregation to certain material defects, lend it a

competitive edge in research related to interfaces and grainboundaries in Ni-base superalloys.

The change of mechanical properties in Ni-base superalloys withaging time is also an important issue that needs to be investigated.The room temperature impact toughness and microhardness ofInconel 740H have already been studied after aging at 750 1C for1000 h at different temperatures [6]. Also, the elevated-temperaturemicrohardness of the alloy in the standard heat treatment conditionhas been studied, and the results compared with room temperaturemicrohardness [6]. However, studies of the evolution of elevated-temperature strength and of room temperature toughness ofInconel 740H with aging time, as well as their relationship withmicrostructure have not been reported yet. In this work, tensiletests at 750 1C and impact toughness measurements at roomtemperature have been carried out for samples aged up to 3000 hat 750 1C. The results are analyzed with respect to the observedchanges in the size of the γ′precipitates and the grain size.

2. Experimental procedure

The bulk chemical composition of the Inconel 740H sample wasanalyzed to be 49.27Ni–25.57Cr–19.92Co–1.48Nb–0.50Mo–1.38Ti–1.47Al–0.34Mn–0.03Fe–0.02C wt%. Samples cut from the forgedbars were solution treated at 1150 1C for 30 min and then waterquenched. These samples were subsequently aged at 750 1C for100, 300, 500, 1000 and 3000 h and then air quenched. Forconvenience, these samples will be subsequently referred to asWQ100, WQ300, WQ500, WQ1000 and WQ3000.

After aging, the specimens were ground and polished followingstandard metallographical methods and then chemically etched in afresh solution of HCl, HNO3 and H2O volume proportions of 10:1:10at 50 1C for 15 min. Optical microscope (OM) investigations wereconducted using an Olympus GX51 microscope. Scanning electronmicroscope (SEM) investigations were conducted using a Hitachi4500 scanning electron microscope. Samples for transmissionelectron microscope (TEM) investigations were prepared by astandard electro-polishing technique using a double-jet device.

Fig. 1. The geometries of the mechanical test samples (a) elevated-temperaturetensile test and (b) room temperature impact toughness test.

Fig. 2. Optical micrographs of the samples aged at 750 1C for different times.

C. Yan et al. / Materials Science & Engineering A 589 (2014) 153–164154

Slices of 0.3 mm thickness were cut out of the bulk material andmechanically ground to 40–50 μm in thickness. Discs of 3 mm indiameter were then punched from the thinned slice and thenelectro-polished at �30 1C at a voltage of 16 V, using a solution of85% ethanol and 15% perchloric acid. TEM investigations wereconducted using a JEOL 2010 TEM operated at 200 kV. The weightfractions of the γ′ precipitates and carbides were ascertainedthrough micro-chemical phase analysis methods. The γ′ precipitateswere electrolytically extracted in an aqueous solution containing 1%(NH4)2SO4 and 1% C6H8O7.H2O from 0 to 5 1C. The carbides wereelectrolytically extracted in a methanol solution containing 5% HCl,5% C3H8O3 and 1% C6H8O7 �H2O from �10 to �5 1C. In both cases acurrent density of 0.05 A/cm2 was used. The particle size distribu-tion and mean radius of the extracted γ′ precipitates were deter-mined using an X-ray diffractro-spectrometer equipped with aKrathy small angle scattering (SAXS) goniometer.

Samples for 3DAP tomography studies in the LEAP microscopywere prepared by an electro-polishing method. For this purpose,

samples of different ageing conditions were first electro-dischargemachined into thin wires with a square cross section of0.7�0.7 mm2. These wires were mechanically ground and subse-quently electro-polished in two steps; first with a 75% aceticacidþ25% perchloric acid solution for the coarse polish, and finallywith a 98% butyl celluloseþ2% perchloric acid solution at 12–14 Vfor the final polish. The 3DAP tomography experiments, wereperformed using an Imago (now Camera Instruments) localelectrode atom probe (LEAP 3000 h) at a residual pressure of5�10–9 Pa and a specimen temperature of 50 K, and with a pulserepetition frequency of 200 kHz and a pulse-voltage to dc-voltageratio of 15%. Data analysis was performed using the IVAS 3.6.0software. Impact testing of Charpy V-notch samples cut from theaged samples, was performed at room temperature. Tensile tests at750 1C were carried out on aged samples in a hydraulic test systemequipped with a temperature-controlled furnace. The temperaturefluctuations were kept within 3 1C. The load rate before yieldingwas 0.5 mm/min and increased to 2.5 mm/min after yielding.

Fig. 3. SEM micrographs and the typical EDS of the grain boundary carbides aged at 750 1C for different times.

C. Yan et al. / Materials Science & Engineering A 589 (2014) 153–164 155

The geometries of the Charpy V-notch samples and tensile testsamples are provided in Fig. 1.

3. Results and discussion

3.1. Microstructure evolution

3.1.1. Grain sizeThe microstructures of the WQ100, 300, 500, 1000, 3000

samples as seen in the optical microscope are shown in Fig. 2.The microstructure consists of a coarse grain structure containingannealing twins. It is evident that grain growth took place duringaging. The average grain sizes, determined by an intercept method,for each aging time were 108, 112, 120, 137, 151 μm respectively.

3.1.2. Grain boundary carbidesThe evolution of the grain boundary carbides with aging time is

illustrated in Fig. 3. The carbides were identified to be Cr richM23C6 according to energy dispersive spectrometer (EDS) mea-surements of the carbides in Fig. 3, which is consistent withprevious studies of Inconel 740 by Xie et [5]. There was onlymarginal coarsening of the carbides at prolonged aging time, withthe weight fraction increasing just from 0.20% to 0.28% (Fig. 4(a)).TEM images and selected area diffraction (SAD) patterns of thegrain boundary carbides are shown in Fig. 5. In general, the grainboundary carbides exhibited two different shapes; block-shapedmost common, and needle-shaped rarely. The needle-shapedcarbides appear to be nucleated from the grain boundaries andto grow perpendicularly into the grains, similar to the morphologyof η phase reported by Zhao on the microstructure evolution of

Fig. 4. Evolution of phase fraction (wt%) with aging time (a) carbides, (b) γ′ phase.

Fig. 5. TEM micrographs and corresponding SAD patterns of different shape grain boundary M23C6. (a) and (c) block-shaped M23C6 and corresponding SAD pattern, (b) and(d) needle-shaped M23C6 and corresponding SAD pattern.

C. Yan et al. / Materials Science & Engineering A 589 (2014) 153–164156

Inconel 740 aged at 750 1C [6,7]. However, SAD patterns of theneedle-shaped phase confirm that it is Cr rich M23C6 carbidesrather than η phase.

3.1.3. γ′ Phase3.1.3.1. Temporal evolution of γ′ morphology. As the maincontributor to high temperature strength in Ni-base superalloys,the morphology, size and chemical composition of γ′ precipitatesare of primary importance in an investigation of Inconel 740H.TEM images, together with diffraction pattern showing typicalsuperlattice reflections from γ′ precipitates for the WQ100, 300,1000, 3000 samples are given in Fig. 6. No γ″ precipitates weredetected from either the TEM images or the SAD patterns. The γ′precipitates exhibited a spheroidal morphology despite ofbecoming slightly cuboidal after long time aging. To determineprecisely the variations in size distribution and mean radius of γ′precipitates with aging time, SAXS experiments were also carriedout. The results are shown in Fig. 7. In general the γ′ precipitatesexhibit a one-modal size distribution, with evident coarseningduring aging, from a mean precipitate radius increasing of 22.4–69.2 nm. The γ′ precipitate weight fraction increased from 14.0% to16.5%, as shown in Fig. 4(b). In the first 100 h, all γ′ precipitatesremained below 30 nm in radius, whereas after that γ′ precipitatesof size 30–60 nm were detected. In the samples aged longer than1000 h, even larger γ′ precipitates of size 60–90 nmwere detected.It should be noted however that, these larger precipitate sizeswere not included in the calculation of the mean precipitate size,as it was assumed these were in fact clusters of several small γ′precipitates together. At the same time, the number density ofprecipitates decreased with aging time (Fig. 6). Although the

morphology of γ′ precipitates can be clearly seen in TEM foils,the elemental distribution between γ′ precipitates and γmatrix, aswell as the compositional profile across the γ/γ′ interfaces cannotbe precisely obtained due to the small size of the γ′ precipitates.For this purpose, 3DAP studies were carried out to study themicrostructure at the atomic scale.

3.1.3.2. Chemical analysis of this alloy. Examples of 3DAP recon-structions from the WQ100 and 3000 samples are shown in Fig. 8(a) and (b). The size of each reconstruction volume is shown in thefigure. The Co rich regions (blue) corresponded to the γ matrix,whereas the red Cr¼25 at% isosurfaces are used to delineate theoutlines of the γ′ precipitates. According to previous work [14],the concentration value used to delineate the Cr isosurfaceswas selected to be the average value of the concentrations in theγ matrix and γ′ precipitates (in this case E25%). Owing to thelimited size of the reconstruction volume, only part of eachγ′ precipitate can be seen in the reconstructed volume, althougha few complete small precipitates are seen in the WQ100 sample.Nevertheless, the 3-D morphology of the γ′ precipitates is seen tobe near spherical, in agreement with the TEM investigations.

We take here the WQ3000 sample as an example to investigatethe partitioning behavior of different alloying elements in Inconel740H. The results are illustrated in Fig. 9. By comparing Figs. 8 and9, it is seen that different elements have a different partitioningbehavior between the γ matrix and the γ′ precipitates. Cr, Co andMo tend to distribute in the γ matrix, while Al, Ti and Nbpreferentially segregate to the γ′ precipitates. This segregationtendency is consistent with other studies of Ni-base superalloys[13,16–18]. It is worth noting that the level of segregation for

Fig. 6. TEM images of γ′ precipitates in samples aged for different times at 750 1C.

C. Yan et al. / Materials Science & Engineering A 589 (2014) 153–164 157

each element was not the same. For example, segregation of Mo inγ matrix was rather low compared with that of Cr and Co.Likewise, the segregation levels for different γ′ segregating ele-ments were also different. A lever rule diagram, first implementedby Blavette et al. [19], was constructed to determine in detail thepartitioning behavior as well as partitioning intensity of thesolutes (Fig. 10). In the diagram, the nominal concentration foreach alloying element minus the solute content in the γ matrix isplotted versus the difference in composition between γ′ precipi-tates and γ matrix. The slope of the best-fit line passing throughthe origin for all elements gives the volume fraction of γ′precipitates. The elements segregating to the γ′ precipitates areseen in the first quadrant of the coordinate system. Points to theright of the diagram correspond to stronger segregation in γ′precipitates (vice versa for the γ matrix segregating elements). Inagreement with the results shown in Fig. 9, the lever rule diagramshows that Al, Ti and Nb tend to segregate in γ′ precipitates, withsegregation intensity in descending order. Similarly in the γmatrixthe Cr, Co and Mo segregate in descending order.

In order to quantify the partitioning behavior of differentelements during aging, we define the solute partitioning ratio asKγ′=γ ¼ Cγ′i =C

γi , where Cγ′i and Cγi are the solute concentrations of

element i in the γ′ precipitates and in the γ matrix, respectively.

Values of Kγ′=γ above 1 indicates γ′ segregating elements, and for γsegregating elements, the value is below 1. Moreover an increasingvalue of Kγ′=γ corresponds to an increasing tendency for segrega-tion to the γ′ phase. The values of Kγ′=γ for different elementsduring aging are summarized in Fig. 11. For the γ segregatingelements, an interesting phenomenon was observed, namely thatthe values of Kγ′=γ for Cr, Co and Mo all reached a peak value at300 h, which means that, the γ′ precipitates contained highestcontents of Cr, Co, Mo after aging at 750 1C for 300 h. The reasonfor this phenomenon might be as follows. Elements such as Cr, Coand Mo that segregate to γ matrix diffuse from γ′ precipitates intothe γmatrix during growth of γ′ precipitates. As observed in Fig. 7,a sudden growth in size of γ′ precipitates was seen between 100 hand 300 h. This growth may leave insufficient time for diffusion ofCr, Co and Mo into the γmatrix. After 300 h aging, the growth rateof the γ′ precipitates slowed down, leaving sufficient time fordiffusion. It is suggested therefore that, the combined effects ofsize growth and repartitioning by diffusion lead to an enrichmentof Cr, Co and Mo inside the γ′ precipitates at 300 h. The trends ofKγ′=γ for Ti and Al with aging time were opposite, indicating thatwith prolonged aging time, Ti in the γ′ phase was graduallyreplaced by Al. The trend of Kγ′=γ with aging time, however, hadnothing in common with other elements.

Fig. 7. Size distribution of γ′ measured by SAXS in Inconel 740H at different aging times.

C. Yan et al. / Materials Science & Engineering A 589 (2014) 153–164158

3.1.3.3. Temporal evolution of the γ/γ′ interface. The compositionalprofiles across the γ/γ′ interfaces for the WQ100, 300, 1000, 3000samples were determined using a powerful analysis tool, theproxigram, for the various γ′ precipitates shown in the reconstruc-tion volumes in Fig. 8. Widely employed in other studies using 3DAP,proxigrams have been shown to be useful for the study of any internalsurfaces, even those involving a complex topology [14,15]. In thisstudy proxigrams for the γ/γ′interfaces were constructed based onisosurfaces with a threshold value of 25 at% Cr. The results are shownin Fig. 12, where only the main alloying elements (Cr, Co, Mo, Ti,Al, Nb) are shown for simplicity. The proxigrams show that thecompositional profile across the γ/γ′ interface is not abrupt but

transitionary. The width of the transitionary region is defined as theinterface width. According to Yoon [19], the choice of Cr thresholdvalue used to fix the isosurfaces only influences the location of theinterface, but not the width of the interface. The width of the interfacein each proxigram was determined from the Al composition profile,between the steady-state γ′ and γ. The width of the γ/γ′ interfaces wasdetermined using a 10–90% of the plateaumethod (long-range γ and γ′compositions) [14,15]. The results are shown in Fig. 12. The interfacewidth in theWQ100 samplewas about 2.0 nm, similar to the results ofRene88DT [14]. The interface width was relatively stable over the first1000 h with a large increase between 1000 h and 3000 h, which isbelieved to be associated with the growth of the precipitates and themisfit strain around the interface.

It is worth noting that in the WQ3000 sample, an enrichmentof Cr was found on the γ matrix side of the γ/γ′ interface, (Fig. 13).This phenomenon was not found in the other aged samples. Thewidth of the enriched region was about 2.5 nm. Yoon observed asimilar phenomenon in the commercial Ni-base superalloy RENEN6, but the reason for this enrichment was not discussed [19].In addition, the enrichment of Co at the interface was also detectedin RENE N6. In contrast this was not found in the WQ3000 sampleinvestigated in the present study.

Two possible explanations have been proposed for the enrich-ment of Cr at the γ/γ′ interfaces. The first is based on a diffusionmechanism. As shown in Fig. 10, Cr segregates strongly to the γmatrix. Therefore, with the increase of γ′ precipitate size, excess Crinside the γ′ precipitates must be excluded and must diffuse acrossthe γ/γ′ interface into the γmatrix. A low diffusion coefficient of Cracross the interface might therefore lead to the enrichment of Cr atthe interface. However, this explanation cannot account for theabsence of Cr enrichment at earlier aging times. The otherexplanation is associated with the suggestion that a new phaseCr-rich may form at the γ/γ′ interface. It has been reported that aNi2Cr-type long-range ordered phase can be expected in Ni–Alalloys with high Cr and Mo content. In addition to the binary Ni–Crsystem such a phase has also been found in ternary Ni–Cr–Mo[20,21] and Ni–Cr–Al alloys [22]. The Ni2Cr superstructure (Pt2Motype), with orthorhombic unit cell, forms in the Ni–Cr systembelow 590 1C. It has been shown that addition of Mo to binaryNi2Cr significantly increases the critical temperature for the for-mation of this phase, e.g. in Ni–25Mo–8Cr, the critical temperatureof Ni2(Cr,Mo) phase is 750 1C [23,24]. Additionally Svoboda and hisco-workers found the existence of Ni2Cr at the γ/γ′ interface in aNi–Cr–Al–Mo system aged at 600 1C for 3000 h [25]. Sundararamanin Alloy 625 also found similar results [26]. Intermetallic phaseNi2(Cr,Mo) precipitates with Pt2Mo-type structure have also beenobserved, in addition to that of the γ″ phase, in Alloy 625 afterprolonged (�70,000 h) service at temperatures close to but lessthan 600 1C. Identification of the nature of the Cr enrichment at theγ/γ′ interface will therefore be the subject of further research.

3.2. Mechanical properties

3.2.1. Tensile strength at 750 1CThe results of tensile tests at 750 1C of Inconel 740H aged at

750 1C for 100, 300, 500, 1000 and 3000 h are shown in Fig. 14. Theyield strength of Inconel 740H at 750 1C increased in the first500 h and then decreased gradually with further increase in agingtime. The elongation increased during the whole aging. Thechanges in yield strength can be rationalized on the basis of thegrowth of γ′ precipitates based on two previously reported models[27,28]. In both cases, it is believed that the critical resolved shearstress (CRSS) is determined by the force necessary to move twocoupled edge dislocations in the ⟨110⟩ direction on the {111} planethrough the γ′ precipitates. The first model [27] describes the CRSSfor cutting through small γ′ precipitates in a disordered matrix by

Fig. 8. 3-D reconstruction showing Co atoms in blue and isosurfaces for Cr¼25 at%in red for more clear visualization of the γ′ precipitates for (a) WQ100 and(b) WQ3000. (For interpretation of references to color in this figure legend, thereader is referred to the web version of this article.)

C. Yan et al. / Materials Science & Engineering A 589 (2014) 153–164 159

two weakly coupled dislocations while the second [28] describesthe same motion for cutting of larger γ′ precipitates by stronglycoupled dislocations. In the first model, the CRSS for two weaklycoupled dislocations is given by

Δτ0 ¼12A

Γb

� �32 bdf

T

� �12

�12

Γb

� �f ð1Þ

where Γ is the anti-phase boundary energy of the γ′ precipitates inthe {111} plane, b is the Burgers vector of the edge dislocation inthe γ matrix, d is the γ′ precipitates diameter, f is the volumefraction of the γ′ precipitates, T is the line tension of the disloca-tion and A is a numerical factor depending on the morphology ofthe particles. For spherical particles A¼0.72 [29]. The line tensionis given by

T ¼ Gb2

2ð2Þ

where G is the elastic shear modulus and is taken as 60 GPa inthis study.

For cutting of larger precipitates by strongly coupled disloca-tion pairs, the CRSS is given by

Δτ0 ¼ 0:86Tf

12ωbd

!1:28dΓωT

� �12

ð3Þ

where w is a constant accounting for the elastic repulsion betweenthe strongly paired dislocations, and which is of the order of unity.For quantitative calculation, f is taken as the average phase fractionof γ′ precipitates (15%), w is taken as 1 for simplicity, and Γ istaken to be 0.42 J/m2.

Fig. 15 is a plot of γ′ precipitate diameter d versus CRSS Δτ0 forthe two dislocation mechanisms. For precipitate cutting by weaklycoupled dislocations, Δτ0 monotonically increases with increasingprecipitate diameter. In contrast Δτ0 gradually decreases withincreasing precipitate diameter for the mechanism of precipitatescutting by strongly coupled dislocations. The resultant Δτ0 for agiven d is the lower value of the two, since dislocation activityfollows the mechanism that provides the least resistance to glide.

Fig. 9. Elemental partition between γ and γ′ phase of sample aged at 750 1C for 3000 h.

C. Yan et al. / Materials Science & Engineering A 589 (2014) 153–164160

For any given alloy, there exists a critical value of d that gives thehighest flow stress. For Inconel 740H, the optimum value of d wascalculated to be around 50 nm (shown in Fig. 15), which coincideswell with the experimental results of yield strength shown inFig. 14. In the first 500 h aging, when the diameter of γ′ pre-cipitates is smaller than 50 nm, the precipitates cutting by weak-coupled pairs of dislocations is the dominant mechanism and theyield strength increases with the increasing precipitate diameter.However, after 500 h as the diameter of γ′ precipitates exceeds thecritical value of d, the mechanism changes to cutting by strong-coupled pairs of dislocations. As a result, the yield strength ofInconel 740H decreases monotonically with longer aging time.

3.2.2. Impact toughness at room temperatureThe results of the impact toughness measurements at room

temperature of Inconel 740H aged at 750 1C for 100, 300, 500, 1000and 3000 h are shown in Fig. 16. A loss of impact toughness with aging

time is evident. For the sample aged at 750 1C for 3000 h, the valuedecreased to less than 20 J, which would present a huge threat to thesafety of use in operation. As reported by Zhao [6], the fractographs ofInconel 740 samples heat-treated at 750 1C indicated a clearly brittlefracture with a localized mixed-mode behavior and a clear intergra-nular fracture behavior. Therefore, the impact toughness of the agedsamples will be influence by the change in grain size and by the size ofthe grain boundary carbides. As mentioned in Sections 3.1.1 and 3.1.2,during aging the size of grain boundary carbides remained unchanged,whereas the grain size increased. Therefore, it can be concluded thatthe drastic loss of impact toughness at room temperature for agedsamples is mainly due to the growth in grain size.

4. Conclusions

The microstructure evolution and mechanical propertiesof Inconel 740H during aging at 750 1C up to 3000 h have

Fig. 11. Partitioning ratios, Kγ′=γ , for main elements in Inconel 740H plotted as a function of aging time.

Fig. 10. Lever rule diagram of Inconel 740H(WQ1000). The nominal concentration (Cn) for each element (Cr, Al, Ti, etc.) minus the solute content in the γ matrix (Cγ) is plottedas a function of the difference in composition between γ and γ′ (Cγ�Cγ′). The slope gives the molar fraction of γ′ present in Inconel 740H.

C. Yan et al. / Materials Science & Engineering A 589 (2014) 153–164 161

been investigated with a combination of OM, SEM, TEM, 3DAP,SAXS, micro-chemical analysis and elevated-temperature ten-sile tests, as well as room temperature impact toughness

measurements. The main conclusions of this study are asfollows:

(1) The average initial grain size of Inconel 740H is above 100 μmand evident grain growth takes place during aging. The grainboundaries are fully decorated with Cr-rich M23C6 carbides,which does not exhibit significant coarsening during aging up

Fig. 12. Proxigrams showing the compositions profiles of Cr, Co, Mo, Ti, Al and Nb atoms across the γ/γ′ interfaces for (a) WQ100, (b) WQ300, (c) WQ1000 and (d) WQ3000conditions.

Fig. 13. Proxigram across the γ/γ′ interface of WQ3000, note that γ is on the rightand γ′ is on the left.

1000100

18

20

22

24

26

28

30

32

34

Aging time/h

Elon

gatio

n/%

450

500

550

600

650

700

750

800

850

900

Stre

ngth

/MPa

A

Fig. 14. The change of mechanical properities of Inconel 740H with aging time.

C. Yan et al. / Materials Science & Engineering A 589 (2014) 153–164162

to 3000 h, with the carbide weight fraction increasingfrom 0.20% to 0.28%. Two different forms of grain boundarycarbides, namely block-shaped and needle-shaped, are observed.Both are identified to be M23C6 by selected area diffractionmeasurements.

(2) TEM and 3DAP image data, together with the results of SAXSexperiments, indicated that size distribution of the spheroidalγ′ precipitates inside each grain is one-modal and that themean precipitate radius increases rapidly during aging at750 1C. The elemental partitioning behaviors have been stu-died with the help of 3DAP measurements and the resultsanalyzed using lever rule diagrams and the solute partitioningratio Kγ′=γ . It was found that, Cr, Co, Mo tend to distribute inthe γ matrix, while Ti, Al, Nb preferentially segregate to the γ′precipitates. The partitioning intensities for different elementsare not the same and vary with aging time.

(3) The compositional profiles across the γ/γ′ interfaces wereinvestigated using proxigrams and the compositional widthsof the interface determined for different aging times. Anincrease in the γ/γ′ interface width was found between1000 h and 3000 h aging, which was believed to be associatedwith the growth of the γ′ precipitates and the misfit strain

around the γ/γ′ interface. Enrichment of Cr at the γmatrix sideof γ/γ′ interface is found for the sample aged for 3000 h.

(4) There exists a critical value for the diameter of γ′ precipitatesthat gives optimum yield strength of Inconel 740H. When thediameter of γ′ precipitates is below the critical value, the yieldstrength increases monotonically with increasing precipitatesize, while when the size is above the critical value, the yieldstrength decreases with the increase of precipitate size. Thecritical value for Inconel 740H was calculated to be around50 nm, which agrees well with the experimental results. A lossof room temperature impact toughness of Inconel 740H withaging time was evident. The main reason for this was attrib-uted to the increase in grain size during aging.

Acknowledgments

The authors are grateful to the financial support from the High-Tech Research and Development Project (No. 2102AA03A501)supported by the National Ministry of Science and Technology ofChina.

Appendix A. Supplementary materials

Supplementary data associated with this article can be found inthe online version at http://dx.doi.org/10.1016/j.msea.2013.09.076.

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