Materials and Design 47 (2013) 218–226

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Materials and Design

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Microstructure evolution and stress-rupture properties of Nimonic 80A aftervarious heat treatments

Yulai Xu a, Caixiong Yang a, Qingxuan Ran a, Pengfei Hu a, Xueshan Xiao a,⇑, Xiuli Cao b, Guoqing Jia c

a Laboratory for Microstructures, Institute of Materials, Shanghai University, Shanghai 200072, Chinab Baoshan Iron & Steel Co., Ltd., Shanghai 200940, Chinac Shanghai Electric Power Generation Equipment Co., Ltd., Turbine Works, Shanghai 200240, China

a r t i c l e i n f o

Article history:Received 29 September 2012Accepted 23 November 2012Available online 5 December 2012

Keywords:Nickel-based superalloyHeat treatmentStress-rupture propertiesMicrostructureTransmission electron microscopy

0261-3069/$ - see front matter � 2012 Elsevier Ltd. Ahttp://dx.doi.org/10.1016/j.matdes.2012.11.043

⇑ Corresponding author. Tel./fax: +86 21 56331484E-mail address: xsxiao@mail.shu.edu.cn (X. Xiao).

a b s t r a c t

Four heat treatments have been designed to produce multi-modal size distributions of c0 phase in cmatrix, and stress-rupture properties of Nimonic 80A have been investigated. The c0 phase exhibitstwo typical coherent orientation relationships with c matrix. The fine c0 particles are in spherical shapewhen the average size is about 20 nm. However, the coarse c0 particles present a cuboidal shape withround corners when they are larger than 75 nm. The Al and Ti elements diffuse from c matrix to c0 phase,and about 80% Al and Ti are used to form c0 phase after various heat treatments. The increased volumefraction of the fine c0 particles with average size of 15.8–20.3 nm increases the stress-rupture life ofNimonic 80A at 750 �C/310 MPa. The rod-like Cr23C6 carbide at grain boundary can suppress grain bound-ary sliding at high temperature and benefit the stress-rupture life. The coarse c0 particles facilitate themovement of dislocations and contribute to the ductility. The precipitate of multi-modal size distribu-tions of c0 phase is beneficial to the stress-rupture ductility.

� 2012 Elsevier Ltd. All rights reserved.

1. Introduction

Nimonic 80A has been widely employed for high temperaturehigh strength applications due to the excellent castability, highcreep and rupture strength by the addition of Al and Ti elements[1–4]. The mechanical properties of wrought nickel-based superal-loys are strongly dependent on the microstructural features suchas grain size, grain boundary precipitate, the volume fraction, sizeand distribution of c0 phase [5,6]. The most significant contributionto the strength of Nimonic 80A is the coherent precipitate of c0

phase [7,8]. In particular, creep strength has been found to behighly sensitive to the volume fraction and size of the smaller c0

phase [9]. The c0 phase which retains its high degree of order upto the melting point has a cube-on-cube orientation relationshipwith c matrix and hinders dislocation motion [10–12]. Sharmaet al. [13] found that the reverse aging heat treatment 1080 �Cfor 8 h, AC + 700 �C for 12 h, AC + 800 �C for 2 h, AC gave signifi-cantly improved stress-rupture life and ductility of alloyNi–19Cr–1.2Al–2Ti–2Co–0.1C–2Fe at 750 �C/300 MPa. Due to theincreased volume fraction of c0 phase and the inhibition of the dis-location shearing of the precipitate, a final aging heat treatment of870 �C for 16 h, AC gave rise to an increased yield stress and pro-moted the operation of {111} h112i slip systems in single crystal

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superalloy SRR99 [14]. Varying precipitate sizes resulted in differ-ent slip modes and different rates of fatigue crack propagation. Itwas found that small c0 particles resulted in greatly reducedfatigue crack propagation rates in Waspaloy (Ni–19.2Cr–1.3Al–3.0Ti–13.9Co–0.05C–1.2Fe) [15].

Therefore, information on the influence of heat treatment onthe morphology and coarsening of c0 phase would be invaluablein the design of heat treatment of nickel-based superalloys.However, hardly any papers have systematically investigated theevolution of c0 phase in Nimonic 80A after various heat treatmentsand the imaging of c0 phase especially the fine c0 phase using trans-mission electron microscope is both difficult and time consuming.The purpose of the present study was to quantitatively analyze themulti-modal size distributions of c0 phases and the chemicalcompositions of Al and Ti elements in c0/c phases. Stress-ruptureproperties of Nimonic 80A after different heat treatments havebeen correlated with the microstructure.

2. Experimental details

The nickel-based superalloy Nimonic 80A (Ni–19.5Cr–1.8Al–2.25Ti–0.06C wt.%) was vacuum induction melted. Due to the meltloss of metal elements during vacuum induction melting, therecovery rates for Al, Cr and Ti were about 95%, 96% and 97%,respectively. The casting ingots with a diameter of 80 mm werehot forged into sticks of 30 mm in diameter at temperature no less

Y. Xu et al. / Materials and Design 47 (2013) 218–226 219

than 1100 �C. The sticks were given the following four heattreatments:

T1: 1070 �C for 8 h, air cooling (AC) + 700 �C for 16 h, AC.T2: 1070 �C for 8 h, AC + 980 �C for 4 h, AC + 700 �C for 16 h, AC.T3: 1070 �C for 8 h, AC + 845 �C for 24 h, AC + 700 �C for 16 h,AC.T4: 1070 �C for 8 h, AC + 980 �C for 4 h, AC + 845 �C for 24 h,AC + 700 �C for 16 h, AC.

The full heat-treated bars were machined into stress-rupturespecimens with a gauge length of 25 mm and gauge diameter of5 mm. The specimens for the stress-rupture tests were preparedaccording to the National Standard of the PR China, GB/T 2039-1997 [16]. The machining tolerance on the nominal dimensionwas ±0.02 mm for the stress-rupture specimens. The heatingdevice heated the stress-rupture specimens to the specified tem-perature 750 �C. The indicated temperature was the temperaturemeasured at the surface of the parallel length of the specimens.The permitted deviations between the indicated temperature andthe specified temperature should be ±4 �C. Stress-rupture testswere carried out at 750 �C/310 MPa. At least three samples wereprepared for each alloy, the time for fracture to occur was mea-sured and the average value was taken as the stress-rupture life.The stress-rupture properties including the stress-rupture life,elongation and reduction in area were determined.

The microstructures of the samples after full heat treatment andstress-rupture tests were investigated by KEYENCE VH-Z100 opti-cal microscope (OM), HITACHI SU-1510 scanning electron micro-scope (SEM) and JEM 2010F transmission electron microscope(TEM) equipped with energy dispersive spectroscopy. TEM foilswere prepared by conventional twin-jet electro-polishing machineat the voltage of 50 V in the solution composed of 10% HClO4 and90% C2H5OH at �30 �C. Since the fine c0 phase presented a spheri-cal shape and the coarse c0 phase presented a cuboidal shape withround corners, the average diameter of the fine spherical c0 phaseand edge length of the coarse cubic c0 phase were measured usinga linear intercept method. In order to reduce the measuring error,at least 200 c0 particles distinguished in several TEM micrographswere analyzed.

3. Results and discussion

After heat treatments T1–T4, stress-rupture tests were carriedout under the applied stress of 310 MPa at 750 �C. The variationsof stress-rupture life, elongation (e) and reduction in area (w) areshown in Fig. 1. The stress-rupture life was about 283 h for thealloy after heat treatment T1. The stress-rupture life slightly

Fig. 1. Variations of stress-rupture life, elongation (e) and reduction in area (w) ofthe alloys after heat treatments T1–T4.

decreased after heat treatment T2, but both the e and w were im-proved. The stress-rupture life dramatically decreased to about189 h after heat treatment T3, but the e and w values were slightlyhigher than those of the alloys after heat treatment T1. The stress-rupture life was about 202 h after heat treatment T4, and the e andw values remarkably increased to about 13% and 18.4%,respectively.

The optical micrographs and fracture morphologies of the alloysafter full heat treatment and stress-rupture tests are shown inFig. 2. The optical micrographs show that the average grain sizevaries slightly but has a wide size distribution in all the alloys afterfull heat treatments T1–T4 (Fig. 2a–d). The linear intercept methodwas employed to measure the average grain size subsequent to thefour heat treatments. The average grain size is found to be 112 lmfor heat treatment T1, 119 lm for T2 and T3, 116 lm for T4. Carbideis found in the grain interior and at the grain boundaries (GBs) afterfull heat treatment T1 (Fig. 2a). However, hardly any carbide can beseen from the optical micrographs after full heat treatments T2–T4

(Fig. 2b and c). After stress-rupture tests, grain boundary (GB)cracking can be clearly seen from the corresponding optical micro-graphs near the fracture surface. Small grain boundary crackingcan be found in the alloy after heat treatments T1–T3, and the num-ber of cracking increases from T1–T3 (Fig. 2e–g). The large crackingvertical to the stress direction can be seen in the alloys after heattreatment T4 (Fig. 2h). Moreover, the grains in the alloy after heattreatment T4 are significantly elongated after stress-rupture tests.The fracture morphologies after stress-rupture tests indicate thatthe fracture presents in the mode of intergranular cracking(Fig. 2i–l). The fine and equiaxial ductile dimples on the fracturesurface confirm the improved stress-rupture ductility of the alloyafter heat treatment T4 (Fig. 2l).

SEM and TEM observations were carried out after full heattreatment (Fig. 3). The SEM micrograph (Fig. 3a) and correspondingEDS (Fig. 3b) show that rod-like Cr23C6 carbide precipitates in thegrain interior and at the grain boundary after full heat treatmentT1. In addition, blocky Cr23C6 carbide with average size about220 nm is also found at grain boundary by TEM observation(Fig. 3c). However, blocky Cr23C6 carbide rather than rod-likecarbide is identified both in the grain interior and at the grainboundary after full heat treatments T2 (Fig. 3d and e), and the cor-responding selected area electron diffraction patterns (SADP) indi-cates that the grain boundary Cr23C6 carbide presents anorientation relationship with c0/c phases (Fig. 3f). The blockyCr23C6 carbide precipitates in the alloys after heat treatments T3

and T4 is similar to that in the alloy after T2, and the average sizeof grain boundary Cr23C6 carbide significantly increases to about410 nm after heat treatments T2–T4 (Fig. 3e).

Fig. 4 shows the typical TEM morphologies and SADP of c0/cphases after each step of heat treatments. The volume fraction(Vol.%) and average diameter (dc0) of the fine (about 20 nm) andcoarse (larger than 75 nm) c0 phase have been calculated basedon the TEM observations (Table 1). The microstructures which con-sist of multi-modal sizes and volume fractions of c0 phase afterheat treatments are complicated. So c0 phases which precipitateafter 980 �C for 4 h, AC, 845 �C for 24 h, AC and 700 �C for 16 h,AC are respectively classified into three distinct groups: theprimary c01, secondary c02 and tertiary c03. After heat treatment T1,the spherical c03 phase distributes in c matrix homogeneously(Fig. 4a). The volume fraction and average diameter of c03 phaseare about 67.1% and 20.3 nm, respectively. What is more, the SAD-Ps indicate that c03 phase exhibits two typical coherent orientationrelationships with c matrix after heat treatment T1: (100)c0 ||(100)c & [001]c0 || [001]c (Fig. 4b) and (100)c0 || (100)c &[011c0 ] || [011c] (Fig. 4c). As to heat treatment T2, small amountof c01 phase precipitates in c matrix with the volume fraction about7.1% and average size about 149.5 nm after aging at 980 �C for 4 h,

Fig. 2. Optical micrographs and fracture morphologies of the alloys after full heat treatment and stress-rupture tests.

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AC (Fig. 4d). When further aging at 700 �C for 16 h, AC, the volumefraction of c01 phase slightly increases to about 7.6%, while theaverage diameter of c01 phase grows up to about 151.5 nm(Fig. 4e). At the same time, the fine c03 phase is identified in c ma-trix with the volume fraction about 63.8% and average diameterabout 16.6 nm. Fig. 4f shows the enlarged morphology of the finec03 phase at the region marked with rectangle in Fig. 4e. As to heattreatment T3, a lot of c02 particles precipitate in c matrix after agingat 845 �C for 24 h, AC. The volume fraction and average size of c02phase are calculated to about 46.4% and 82.8 nm, respectively(Fig. 4g). When further aging at 700 �C for 16 h, AC, the volumefraction slightly increases to about 51.3% and correspondinglythe average size of c02 phase grows up to about 86.9 nm (Fig. 4h).Fig. 4i shows the morphology of the enlarged rectangular region

marked in Fig. 4h, and c03 phase precipitates in c matrix after agingat 700 �C for 16 h, AC in addition to c02 phase. The volume fractionand average diameter of c03 phase are calculated to about 22.0% and15.8 nm, respectively. The c01, c01 and c03 particles are all identifiedin c matrix after full heat treatment T4. Fig. 4d is also given inFig. 3j in order to show the precipitate behavior of c0 phase moreclearly after each step of heat treatment T4. As mentioned above,small amount of c01 phase precipitates in c matrix with the volumefraction about 7.1% and average size about 149.5 nm after aging at980 �C for 4 h, AC (Fig. 4j). When further aging at 845 �C for 24 h,AC, the volume fraction increases significantly to about 17.2%and c01 phase grows up to about 187.9 nm (Fig. 4k). In addition,the c02 phase also precipitates in c matrix. The volume fractionand average size of c02 phase are calculated to about 17.3% and

Fig. 3. Micrographs after full heat treatment: (a) SEM micrograph and (b) corresponding EDS and (c) TEM micrograph of Cr23C6 carbide after full heat treatment T1, Cr23C6

carbide (d) in the grain interior and (e) at the grain boundary and (f) corresponding SADP after heat treatment T2.

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75.9 nm, respectively (Fig. 4k). When further aging at 700 �C for16 h, AC, both c01 and c02 phases grow up (Fig. 4l). The volume frac-tion of c01 increases to about 19.7% and the volume fraction of c02increases to about 20.7%. The average size of c01 and c02 phases in-creases to about 205.6 nm and 90.7 nm, respectively. Besides, veryfine c03 phase is also identified in c matrix. The volume fraction ofc03 phase is about 25% and the average diameter is about 19.3 nm.Although different volume fraction and average size of c0 phasesare found after various heat treatments T1–T4, the c01, c02 and c03particles all exhibit two kinds of coherent orientation relationshipswith c matrix as shown in Fig. 4b and c.

The typical TEM morphologies of c0/c phases after stress-rup-ture tests are shown in Fig. 5. The volume fraction and averagediameter of c0 phase are also shown in Table 1. The SADPs of thesamples after stress-rupture tests are the same with those after fullheat treatment, and c0 phase still exhibits two typical coherent ori-entation relationships with c matrix: (100)c0 || (100)c & [001]c0 ||[001]c, and (100)c0 || (100)c & [011c0 ] || [011c]. After stress-rup-ture tests, the volume fraction and average size of c03 phase afterheat treatment T1 are about 57.2% and 58.2 nm, respectively. Thec03 phase has grown up significantly from about 20.3 nm to58.2 nm after stress-rupture tests compared with the morphologyof c03 phase after full heat treatment T1. But the volume fraction hasslightly decreased by about 10% (Fig. 5a). As to heat treatment T2,the volume fraction and average size of c01 phase after stress-rup-ture tests have increased from about 7.6% and 151.5 nm to 14.4%and 177.4 nm, respectively. The volume fraction of c03 phase afterstress-rupture tests has decreased significantly from about 63.8%to 20.1%. However, the average size of c03 phase after stress-rupturetests has increased significantly from about 16.6 nm to 53.7 nm(Fig. 5b). As to heat treatment T3, the volume fraction of c02 phaseafter stress-rupture tests slightly decreases to about 48.4% andthe average size of c02 phase slightly increases to about 88.6 nm.However, no c03 phase is identified in c matrix after stress-rupture

tests (Fig. 5c). As to heat treatment T4, both the volume fractionand average size of c01 phase after stress-rupture tests have in-creased to about 19.8% and 236.3 nm, respectively. Although theaverage size of c02 phase after stress-rupture tests is nearly thesame with that after full heat treatment T4, the volume fractionof c02 phase has decreased to about 10.8%. What is more, no c03phase precipitates in c matrix after stress-rupture tests (Fig. 5d).

Comparing both heat treatments T1 and T2 or T3 and T4, the heattreatment step of 980 �C for 4 h, AC seems to have very little influ-ence on the volume fraction of c03 phase because the volumefraction of c03 phase is constant at about 63.8–67.1% after full heattreatments T1 and T2, while the volume fraction of c03 phase is con-stant at about 22–25% after full heat treatments T3 and T4 (Table 1).However, comparing both heat treatments T1 and T3 or T2 and T4,the heat treatment step of 845 �C for 24 h, AC has a great influenceon the volume fraction of c03 phase due to the significant decreaseof the volume fraction of c03 phase after full heat treatments T3 andT4. As a result, the heat treatment step of 845 �C for 24 h, AC cansignificantly influence the volume fraction of c03 phase. But boththe heat treatment steps of 980 �C for 4 h, AC and 845 �C for24 h, AC have slight influence on the average size of the fine c03phase which is stable at about 20 nm (Table 1). Although the c03phase is identified after stress-rupture tests in the alloys after heattreatment T2, the amount of c01 phase (Vol.% = 7.6%) after heattreatment T2 is less than that after heat treatment T4

(Vol.% = 19.7%). So the disappearance of c03 phase after stress-rupture tests for the alloys after heat treatments T3 and T4 can beascribed to the precipitate of large amount of c01 and c02 phases dur-ing the isothermal aging heat treatment. In the system with multi-modal size distributions of coherent precipitates, coarsening pro-cess is driven by the reduction in total interfacial energy and largeprecipitates grow at the expense of small ones [6]. What is more,Fig. 4 shows that c01 and c02 particles in the alloys are in cuboidalshape with round corners when they are larger than about

Fig. 4. TEM micrographs and corresponding SADPs show the coherent precipitate of c0 phase for the alloys after heat treatments: (a) the c03 phase after heat treatment T1, (b)and (c) corresponding SADPs, (d) the c01 phase after heat treatment 1070 �C for 8 h, AC + 980 �C for 4 h, AC, (e) the c01 and c03 phase after heat treatment T2, (f) enlarged c03phase after heat treatment T2, (g) the c02 phase after heat treatment 1070 �C for 8 h, AC + 845 �C for 24 h, AC, (h) the c02 and c03 phase after heat treatment T3, (i) enlarged c02 andc03 phase after heat treatment T3, (j) the c01 phase after heat treatment 1070 �C for 8 h, AC + 980 �C for 4 h, AC, (k) the c01 and c02 phase after heat treatment 1070 �C for 8 h,AC + 980 �C for 4 h, AC + 845 �C for 24 h, AC and (l) the c01, c02 and c03 phases after heat treatment T4.

Table 1The volume fraction and average size of c0 phase after heat treatments and stress-rupture tests.

Heat treatment Types 980 �C � 4 h, AC 845 �C � 24 h, AC 700 �C � 16 h, AC After stress-rupture tests

Vol.% dc (nm) Vol.% dc0 (nm) Vol.% dc0 (nm) Vol.% dc0 (nm)

T1 c03 – – – – 67.1 20.3 57.2 58.2

T2 c01 7.1 149.5 – – 7.6 151.5 14.4 177.4c03 – – – – 63.8 16.6 20.1 53.7

T3 c02 – – 46.4 82.8 51.3 86.9 48.4 88.6c03 – – – – 22.0 15.8 – –

T4 c01 7.1 149.5 17.2 187.9 19.7 205.6 19.8 236.3c02 – – 17.3 75.9 20.7 90.7 10.8 92.8c03 – – – – 25.0 19.3 – –

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75 nm after heat treatments T2–T4. Sharma et al. [13] found that c0

particles are in spherical shape when they are about 20 nm andthat c0 particles are in cuboidal shape when they are about60 nm. The c03 particles with size less than about 60 nm in thealloys after stress-rupture tests for the alloys after heat treatmentsT1 and T2 are approximately in spherical shape (Fig. 5a and b).

However, the c01 and c02 particles in the alloys after stress-rupturetests for the alloys after heat treatments T3 and T4 are approxi-mately in cubic shape (Fig. 5c and d). The morphology change ofc0 particles at high temperature is found to follow the diffusioncontrolled model, and many factors such as the volume fraction,lattice misfit, applied stress and temperature can influence the

Fig. 5. TEM micrographs of the alloys after stress-rupture tests at 750 �C/310 MPa: (a) T1, (b) T2, (c) T3 and (d) T4.

Table 2Chemical compositions of Al and Ti elements in c0/c phases after heat treatment T4 and stress-rupture tests (wt.%).

Heat treatment Types 980 �C � 4 h, AC 845 �C � 24 h, AC 700 �C � 16 h, AC After stress-rupture tests

Al Ti Al Ti Al Ti Al Ti

T4 c01 1.92 3.27 3.75 4.16 4.62 5.19 5.98 6.78c02 – – 1.20 1.52 1.45 1.96 1.68 2.34c03 – – – – 1.06 1.96 – –c matrix 1.70 2.02 1.29 1.71 0.78 0.92 0.53 0.89d 7.9% 11.0% 50.2% 46.6% 84.5% 85.8% 78.8% 72.1%

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shape change [17–19]. Ricks et al. [20] found that development ofthe cubic corners can be enhanced by the diffusion of solute tothese regions in Nimonic 115 after heat treatment. The lattice mis-fit may be the driving force for the changes of c0 phase shape fromspherical to cubic at high temperature. The c0 phase observed inUDIMET 720 Li appears in shapes ranging from spheres to cubesand the structural transition is mainly driven by the lattice misfit[21]. An alternative explanation for the morphology change canbe that the morphology development is highly sensitive to theimpingement of adjacent particles. A white dashed circle is markedin Fig. 5d which indicates that the adjacent particles may influencethe morphology evolution of c0 phase at high temperature. Glatzeland Feller-Kniepmeier [22] found the interaction between adjacentparticles increases which leads to a more cubic morphology of c0

phase when the volume fraction is about 33–40%.In addition to the quantitative analysis of the volume fraction

and size of c0 phase, TEM–EDS was employed to analyze the chem-ical compositions of Al and Ti among the primary c01, secondary c02,tertiary c03 and c phases. Table 2 shows the chemical compositionsof Al and Ti elements in the multi-modal c0 particles and c matrix

after heat treatment T4. The weight percent of Al and Ti elements inc01 phase is respectively about 1.92 wt.% and 3.27 wt.% after heattreatment at 980 �C for 4 h, AC. The weight percent of Al and Ti ele-ments respectively increases to about 3.75 wt.% and 4.16 wt.% afterfollowing aging at 845 �C for 24 h, AC. At the same time, the c02phase precipitates in c matrix with the Al and Ti contents about1.20 wt.% and 1.52 wt.%, respectively. After full heat treatment T4,the weight percent of Al and Ti elements in c01 phase respectivelyincreases to about 4.62 wt.% and 5.19 wt.%, while the weight per-cent of Al and Ti elements in c02 phase respectively increases toabout 1.45 wt.% and 1.96 wt.%. The weight percent of Al and Ti ele-ments in c03 phase is respectively about 1.06 wt.% and 1.96 wt.%.The weight percent of Al and Ti elements in c matrix is respectivelyabout 1.70 wt.% and 2.02 wt.% after heat treatment 1070 �C for 8 h,AC + 980 �C for 4 h, AC. The weight percent of Al and Ti elements inc matrix respectively decreases to about 1.29 wt.% and 1.71 wt.%after heat treatment 1070 �C for 8 h, AC + 980 �C for 4 h,AC + 845 �C for 24 h, AC. The weight percent of Al and Ti elementsin c matrix after full heat treatment further decreases to about0.78 wt.% and 0.92 wt.%, respectively. After stress-rupture tests,

Fig. 6. Schematic morphologies of c0 phase and carbide after heat treatments: (a) T1, (b) T2, (c) T3 and (d) T4.

Fig. 7. Thermodynamic calculation of the equilibrium phase weight fractions as afunction of temperature.

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the weight percent of Al and Ti elements in c01 phase respectivelyincreases to about 5.98 wt.% and 6.78 wt.%, and the weight percentof Al and Ti elements in c02 phase respectively increases to about1.68 wt.% and 2.34 wt.%. While the weight percent of Al and Tielements in c matrix is respectively about 0.53 wt.% and0.89 wt.% after stress-rupture tests.

The weight percent ratio (d) of Al and Ti element in c0 phase tothat in c0 and c phases after heat treatment T4 was calculatedaccording to the equation:

di ¼Wi

c0

Wic0 þWi

c

ð1Þ

where i represents Al and Ti elements, Wic0 and Wi

c respectively rep-resent the weight percent of i element in c0 (includes c01, c02 and c03)and c phases. Wi

c0 and Wic are calculated according the volume frac-

tion of c0/c phases (Table 1) and weight percent of i element in c0/cphases (Table 2). After 1070 �C for 8 h, AC + 980 �C for 4 h, AC, thedAl and dTi values are about 7.9% and 11.0% in c0 phase, respectively.After 1070 �C for 8 h, AC + 980 �C for 4 h, AC + 845 �C for 24 h, AC,the dAl and dTi values respectively increase to about 50.2% and46.6% in c0 phase. After full heat treatment, the dAl and dTi valuesfurther increase to about 84.5% and 85.8% in c0 phase. The aboveresults indicate that more Al and Ti elements diffuse from c matrixto c0 phase and that at least 80% Al and Ti are used to form thestrengthening c0 phase after heat treatment T4. The chemical com-positions of Al and Ti elements in c0 particles after heat treatment T4

also indicate that most of the Al and Ti elements diffuse into c01phase and the contents of the Al and Ti elements in c matrix de-crease. After stress-rupture tests at 750 �C/310 MPa, Table 2 showsthat the dAl and dTi values slightly decrease to about 78.8% and72.1%, respectively. The decrease is primary due to the decreasedvolume fraction of c02 phase and disappearance of c03 phase afterstress-rupture tests in the alloy after heat treatment T4. Singhet al. [23] also found that the primary c0 phase contains a higheramount of Al and Ti elements, whereas the secondary or tertiaryc0 phase has lower Al and Ti elements. It is also found that the ter-tiary c0 phase disappears and the secondary c0 phase coarsens,which acts as the main reason for the decreasing of strength bothat room temperature and high temperature [24].

The schematic morphologies of c0 phase and carbide after heattreatments T1–T4 are summarized in Fig. 6. The above quantitativeanalytical results indicate that the total volume fraction of c0 phase(includes c01, c02 and c03 phases) is about 67.1%, 71.4%, 73.3% and65.4% for the alloys after heat treatments T1–T4, respectively. Soit is interesting that the approximate spherical c03 phase (withaverage diameter of 15.8–20.3 nm) rather than the total volume

fraction of c0 phase has a stronger influence on the high tempera-ture stress-rupture properties. The strengthening phase is mainlycomposed of c03 phase with the volume fraction about 67.1% andthe stress-rupture life is the longest (about 283 h) for the alloyafter heat treatment T1. As to heat treatment T2, the stress-rupturelife slightly decreases to about 277 h due to the slightly decrease ofthe volume fraction of the strengthening c03 phase to about 63.8%. Alarge number of c02 phase (about 46.4%) precipitates in c matrixafter heat treatment 1070 �C for 8 h, AC + 845 �C for 24 h, AC. Sothe volume fraction of c03 phase is very small (about 22.0%) afterheat treatment T3 compared with that after heat treatments T1

and T2. As a result, the stress-rupture life significantly decreasesto about 189 h after heat treatment T3. Although quite a lot of c01and c02 particles precipitate during heat treatment T4, the volumefraction of c03 phase increases compared with that after heat treat-ment T3. Hence, the stress-rupture life increases to about 202 h.Coakley et al. [25] suggested that the coarsening of secondary par-ticles may be the life limiting factor for the Nimonic 115. Yang et al.[8] investigated the effects of heat treatments on mechanical prop-erties of Rene 80 and also found that precipitate of the fine c0 phaseimproved stress-rupture life. So the higher volume fraction of c03phase can lead to the longer stress-rupture life.

Thermodynamic calculation software is employed to calculatethe equilibrium phase. The equilibrium phase weight fractions asa function of temperature are presented in Fig. 7. Only the liquidphase forms above 1384 �C. The c phase and liquid phase coexistin temperature range of 1348–1384 �C, and their weight fractionschange with the temperature. The equilibrium phases indicate that

Fig. 8. Schematic deformation mechanisms of the alloys after heat treatments: (a) T1 and (b) T4.

Y. Xu et al. / Materials and Design 47 (2013) 218–226 225

no c0 phase precipitates in c matrix when the temperature is high-er than 948 �C, but Fig. 4d shows that c01 phase can precipitate in cmatrix after heat treatment at 980 �C for 4 h. The equilibriumphases are c0 phase and M23C6 carbides at 845 and 700 �C andthe weight fraction of c0 phase increases with the decrease oftemperature, which is consistent with the TEM observations.Fig. 8 shows the schematic deformation mechanisms of the alloysafter heat treatments T1 and T4. The grain boundary sliding at hightemperature is suppressed after full heat treatment T1 (Fig. 8a),which is primary due to the precipitate of rod-like Cr23C6 carbideat grain boundary (Fig. 6a). Hence, hardly any large grain boundarycracking is observed from the optical micrograph. However, theblocky Cr23C6 carbide at grain boundary is observed after full heattreatment T4 (Fig. 6d) and the grain boundary sliding is easier. Sothe grain boundary cracking can easily form during stress-rupturetests at high temperature (the red region marked on Fig. 8b). Com-paring with the microstructure after stress-rupture tests, it can beknown that the high temperature ductility of the alloys after heattreatment T2 and T4 is better than that after heat treatment T3,although blocky Cr23C6 carbide precipitates at grain boundary(Fig. 6b–d). The increased e and w values indicate that the precip-itate of c0 phase with multi-modal size is beneficial to the stress-rupture ductility. And it seems that the high temperature ductilityhas great relation with the large size c0 particles because bothFig. 6b and d show that the large c0 particles with size larger than150 nm are found in c matrix. The coarsening of c0 particles facil-itates the movement of mobile dislocations and contributes to theductility, which is consistent with the recent investigation byWang et al. [24]. The small sizes of c0 phase usually promote dislo-cation shearing while larger particles usually promote Orowanlooping mechanism [13,15]. So the dislocation in c matrix wouldeither slip away or would not penetrate through c0 phase whenthe size of c0 phase was large. Fujita et al. [26] found that theanti-phase boundary energy on the {111} slip planes was higherthan that on {100} planes, and several steps of aging heat treat-ments were useful to improve stress-rupture ductility in alloyNi–19.5Cr–1.4Al–2.3Ti–0.07C–0.8Fe because dislocations couldcross slip from {111} planes to {100} planes more likely whenthe alloy was used at temperatures above 700 �C. Sharma et al.[13] suggested that slip became more homogeneous when thematerials subjected to reverse aging heat treatment, thus the duc-tility was promoted. So the improved stress-rupture ductility hasgreat relation with both Cr23C6 carbide and the strengthening c0

particles. The above results indicate that the microstructuresconsisting of multi-modal size distributions of c0 phase in c matrixprovide the best combination of high temperature stress-rupturelife and ductility.

4. Conclusions

Quantitative analysis of c0 phase was carried out after variousheat treatments and the stress-rupture properties were correlatedwith the microstructure evolution. The results indicate that c0

phase exhibits two typical coherent orientation relationships withc matrix: (100)c0 || (100)c & [001]c0 || [001]c, and (100)c0 || (100)c& [011]c0 || [011]c. The fine c0 particles are in spherical shapewhen the average size is about 20 nm, but c0 particles in the alloysare in cuboidal shape with round corners when they are larger than75 nm. The Al and Ti elements diffuse from c matrix to c0 phase,and at least 80% Al and Ti elements are used to form the strength-ening c0 phase after various heat treatments. The increased volumefraction of the fine c0 particles with average size of 15.8–20.3 nmincreases the stress-rupture life of Nimonic 80A at 750 �C/310 MPa. The precipitate of rod-like rather than blocky Cr23C6 car-bide at grain boundary can greatly suppress grain boundary slidingduring stress-rupture tests at high temperature and improve thestress-rupture life. The coarsening of c0 particles facilitates themovement of mobile dislocations and contributes to the stress-rupture ductility. The microstructures consisting of multi-modalsize distributions of c0 phase in c matrix after various heattreatments provide the best combination of high temperaturestress-rupture life and ductility.

Acknowledgements

The authors would like to thank Peng JC for the excellent sup-port during JEM 2010F TEM observations in Instrumental Analysis& Research Center of Shanghai University. And we also extend ourthanks to Dr. Chen K during the thermodynamic calculation inSchool of Materials Science and Engineering of Shanghai Jiao TongUniversity. This research was sponsored by the key project ofScience and Technology Commission of Shanghai Municipality un-der Contract No. 09521101404.

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