State Key Laboratory of Solidification Processing, Northwestern Polytechnical University, Xi’an 710072, PR ChinaSuzhou Nuclear Power Research Institute, Suzhou 215004, PR China
r t i c l e i n f o
rticle history:eceived 15 February 2012eceived in revised form 16 March 2012ccepted 26 March 2012vailable online 2 April 2012
a b s t r a c t
Precipitation behavior of grain boundary (GB) M23C6 and its effect on tensile properties at elevated tem-perature were investigated systematically in a Ni–Cr–W based superalloy. The results show that theM23C6 precipitation behavior is influenced obviously by grain boundary character (GBC) and interfacialenergy. The �≤9 GBs and low angle GBs have low interfacial energy, and no M23C6 carbide precipitatesat these GBs. Plenty of M23C6 carbide particles precipitate at the large angle GBs with high interfacialenergy. The coherent orientation relationship between M23C6 and the matrix plays an important role on
eywords:i–Cr–W superalloy23C6
recipitationSL
the precipitation morphology of M23C6. M23C6 carbides with four typical morphologies distribute at thelarge angle GBs, including lamellar carbide which grows into the matrix near one side or both sides of theGBs, rod-like carbide and small lamellar carbide both of which grow along GBs. Moreover, the decrease ofboth tensile and yield strength of the aged alloy is mainly caused by the lamellar M23C6 carbide breaking.
y irre
isorientationensile properties
The tensile properties var
. Introduction
Under the conditions of heat treatment or during high-emperature services, M23C6 as typical carbide commonlyrecipitates at the grain boundary (GB) in superalloys [1,2]. It
s indicated that interfacial energy and grain boundary characterGBC) plays an important role on M23C6 precipitation behavior3,4], and much attention has been focused on the special effectf coincidence site lattice (CSL) GBs [5–8]. The results showed that23C6 can hardly precipitate at the �3c GBs which have low inter-
acial energy in the 304 stainless steel with different carbon content3]. When M23C6 precipitates at the CSL GBs, the precipitation mor-hology of M23C6 is seriously affected by the CSL GBs [5,6]. Theendritic M23C6 is observed at �3i type GBs in the alloys such as 316tainless steel [7] and Inconel 690 [8]. However, plate-like M23C6an form at the �3 GBs in the 304 stainless steel during heat treat-ent [6]. Moreover, bar-like and needle-like M23C6 carbide can
recipitate near only one side of �9 GBs [8,9], and carbide particlessually precipitate at the �>9 CSL GBs [8]. The random GBs withigh interfacial energy are suitable for the formation of M23C6 andarbide particles usually form at these GBs [4,9]. Up to now, it is
till unclear why the morphology of M23C6 is different at differentypes of GBs in various alloys. The precipitation behavior of M23C6arbide remains to be further investigated.
The mechanical properties of superalloys are affected seriouslyby M23C6 carbides [10–13]. Koul and Castillo [11] investigated thecreep behavior of IN738LC alloy and pointed out that continuousM23C6 carbide films distributed along GBs reduce creep rupture lifeof the alloy. And the creep ductility can be improved by the discreteM23C6 carbide through hindering GB sliding [12,13]. The discreteM23C6 carbide can also have coherent strain strengthening and pre-cipitate strengthening effects, and they are considered necessaryfor obtaining optimum creep properties of alloys [14]. The role ofdifferent morphology of M23C6 in different alloys is very complex.Hence, it is important to study M23C6 precipitation behavior andits influence on mechanical properties in the given alloys.
Recently, much effort has been devoted to develop Ni–Cr–Wbased superalloy as potential materials for the applications suchas nuclear power, chemical processing and aerospace industry. It isreported that numbers of coarse lamellar M23C6 carbide can precip-itate at GBs in Ni–Cr–W superalloy [15]. However, it is unclear howthe precipitation behavior of M23C6 is affected by GBC. The purposeof this paper is to study the effect of GBC on M23C6 precipitationbehavior in Ni–Cr–W based superalloy. The influence of M23C6 ontensile properties at elevated temperature is also investigated.
2. Experimental method
The chemical composition in weight percent (wt.%) of thewrought Ni–Cr–W based superalloy used in this work was: Cr,19.82; W, 18.48; Mo, 1.24; Al, 0.46; C, 0.11; B, 0.0028; La, 0.026;P, S < 0.004, Bal. Ni. The experimental alloy was initially vacuum
84 R. Hu et al. / Materials Science and Engineering A 548 (2012) 83– 88
t 700
itdsaieNE1mc
tsEaata
Fig. 1. Microstructures of CSL GBs in Ni–Cr–W based superalloy aged a
nduction melted (VIM) and vacuum arc remelted (VAR). And thenhe ingot was homogenized at 1200 ◦C for 24 h under vacuum con-ition. Subsequently, it was hot forged and rolled at 1150 ◦C into aheet with a thickness of 9 mm. Samples were cut from the sheetnd solution treated at 1260 ◦C for 0.5 h followed by water quench-ng. The original microstructure of the alloy has been discussedlsewhere [16]. The aging treatment at 700 ◦C for 1–5 h on thei–Cr–W sample was performed, followed by water quenching.very value of tensile properties was a mean of 3 tests using a Zink50 mechanical testing machine at 900 ◦C in air, where a strain rateaintained 8.3 × 10−5 s−1 up to yield and 1.25 × 10−3 s−1 after yield
ondition was employed.Prior to the scanning electron microscopy (SEM) and the elec-
ron backscatter diffraction (EBSD) experiments, the polishedpecimens were etched with aqua regia (HCl:HNO3 = 3:1) for 30 s.BSD was employed for the determination of GBC. The CSL bound-
ries were defined by Brandon criterion (��max = 15�−1/2) [17]nd the low angle GB was discerned by criterion � < 15◦. Afterhe EBSD experiment, the microstructure of the EBSD scannedreas in the alloy was analyzed by JSM-6700 SEM. The volume
◦C for 5 h: (a) �3i and �3c; (b) �5; (c) �9; (d) �11; (e) �23; (f) �41c.
fraction of carbide was measured using an image analysis systemtype Image Pro-Plus, values were expressed as mean ± standarderror of the mean. The structure and morphology of precipitatedcarbide in the aged alloy were also analyzed by Tecnai G2 F30 trans-mission electron microscopy (TEM). TEM samples were prepared byelectrochemical polishing at −30 ◦C in 8% perchloric acid carbinolsolution.
3. Experimental results
3.1. Effect of CSL GBs on M23C6 precipitation behavior
The typical CSL GB morphologies of the aged Ni–Cr–W basedsuperalloy are shown in Fig. 1. There is no carbide precipitating at�3i, �3c, �5 and �9 CSL GBs (Fig. 1a–c). It can be seen that �≤9GBs exhibit good corrosion resistance and intragranular corrosion
happens. The CSL GBs have lower interfacial energy compared withthat of the random GBs, and �3 GBs have the lowest interfacialenergy [18–20]. M23C6 carbide does not precipitate at the �3 GBsdue to its low interfacial energy, and the carbide can precipitate
R. Hu et al. / Materials Science and Engineering A 548 (2012) 83– 88 85
based
aapcp[i
3
N4[ttd
Fl
Fig. 2. Microstructures of low angle GBs in Ni–Cr–W
t the �>3 GBs. In addition, a small amount of M23C6 particlesre observed at �11, �23 and �41c GBs (Fig. 1d–f). When M23C6recipitates at �>9 GBs, much Cr atoms are consumed and Croncentration near these boundaries decreases seriously. Cr atomslay an important role on corrosion resistance properties of alloys21]. Hence, the GBs containing M23C6 are heavily corroded due tots low Cr concentration (Fig. 1d–f).
.2. Effect of GB misorientation on M23C6 precipitation behavior
Fig. 2 shows the low angle GB microstructures of the agedi–Cr–W based superalloy. The GB misorientations are 3.73◦ and.03◦, respectively. The low angle GBs have low interfacial energy
3] and M23C6 cannot precipitate at these boundaries. It can be seenhat the low angel GBs have excellent corrosion resistance. The par-icles distributed at the GBs are primary M6C carbide which formeduring solidification process.
ig. 3. Microstructures of large angle GBs in Ni–Cr–W based superalloy aged at 700 ◦C foramellar M23C6 grows from the GB into both sides matrix of that GB; (c) rod-like M23C6; (
superalloy aged at 700 ◦C for 5 h: (a) 3.73◦; (b) 4.03◦ .
The large angle GB microstructures of the aged alloy are shownin Fig. 3. These GBs have high interfacial energy and plenty of M23C6carbide precipitates at the boundaries. It can be seen from Fig. 3a–dthat the M23C6 carbide with four typical morphologies distributesat the large angle GBs, including lamellar carbide which grow intomatrix near one side or both sides of GBs, rod-like carbide and smalllamellar carbide. The volume fraction of the lamellar carbide whichgrows into only one side matrix of the GBs is as high as 85 ± 1.7% andthe average width of the lamella is about 2–3 �m (Fig. 3a). The car-bide lamella can also grow into the austenite grains on both sidesof the GBs and the average width is about 3–4 �m (Fig. 3b). Therod-like and small lamellar M23C6 carbide grow along GBs (Fig. 3cand d). It can be seen from Fig. 3c that the growth of the rod-like
carbide into the austenite grains is inhibited. The complex mor-phologies of M23C6 in the aged Ni–Cr–W based superalloy have notbeen observed in other alloys. When M23C6 precipitate at the GBs,lots of Cr atoms are consumed and these GBs are corroded seriously.
5 h: (a) lamellar M23C6 grows from the GB into only one side matrix of that GB; (b)d) small lamellar M23C6.
86 R. Hu et al. / Materials Science and Engineering A 548 (2012) 83– 88
F y agee
3t
tTtTaMtctmii
4
4
oGat
Fe
ig. 4. Effect of aging treatment on tensile properties of Ni–Cr–W based superallolongation and reduction in cross-section vs. aging time.
.3. Effect of aging treatment on tensile properties at elevatedemperature
The effect of aging treatment on tensile properties at elevatedemperature of the Ni–Cr–W based superalloy is shown in Fig. 4.he results show that the tensile strength and yield strength ofhe solution treated alloy are 345 MPa and 285 MPa, respectively.he average tensile strength and yield strength of the aged alloyre obviously lower than those of solution treated alloy (Fig. 4a).oreover, the elongation of solution treated alloy is 66.7% and
he average elongation of the aged alloy is 76.1%. The reduction inross-section remains almost constant before and after the agingreatment (Fig. 4b). The test results show that the aging treat-
ent decreases both of the tensile strength and yield strength, butmproves the elongation of the alloy. The tensile properties varyrregularly with increasing aging time.
. Discussion
.1. GB M23C6 precipitation behavior during the aging treatment
The lamellar M23C6 carbide which grows into the matrix near
ne side or both sides of GBs has been found only at �3i and �9Bs in many other alloys such as 304, 316 stainless steels [5,7]nd Inconel 690 [8,9]. There are several mechanisms to explainhe formation of lamellar M23C6. Early studies show that the
ig. 5. TEM analysis of M23C6 in Ni–Cr–W based superalloy aged at 700 ◦C for 5 h: (a) blectron microscopy of M23C6/� interface.
d at 700 ◦C and tested at 900 ◦C: (a) tensile and yield strength vs. aging time; (b)
precipitation of lamellar M23C6 can be attributed to dislocationmovement, stacking faults formation or residual stress [7,22,23].However, the formation of the lamellar M23C6 in Ni–Cr–W basedsuperalloy cannot be explained perfectly by these mechanisms.Recent investigations indicate that when M23C6 nucleuses have acoherent orientation relationship with matrix, they would growon a certain plane and along a certain direction in the matrix. Asa result, the lamellar carbide forms [5,9]. So the coherent orien-tation relationship between the carbide and the matrix plays animportant role on the precipitation morphology of M23C6 carbide.
Fig. 5 shows bright-field TEM image, selected area electrondiffraction and high resolution electron microscopy of M23C6 inthe Ni–Cr–W based superalloy aged at 700 ◦C for 5 h. The resultsshow that M23C6 grows on the {1 1 1} plane and along 〈1 1 0〉direction of the matrix. The crystallography relationship betweenM23C6 and matrix can be expressed as 〈1 1 1〉M23C6//〈1 1 1〉matrix,{1 1 1}M23C6//{1 1 1}matrix. M23C6 carbides precipitate directly fromthe matrix [13]: 23M + 6C → M23C6.
According to the above analysis, the formation mechanism ofvarious morphologies of M23C6 at the large angle GBs in the agedalloy can be explained as follows. When M23C6 carbide nucleateson GBs and has {1 1 1}M23C6//{1 1 1}matrix coherent orientation
relationship with the matrix on one side of the GB, they grow intothe matrix near only one side of that GB (Fig. 6a). As a result, thelamellar carbide which grows into one of the two grains formson the GBs as shown in Fig. 3a. If the carbide particle nuclei have
right-field TEM image and selected area electron diffraction; (b) high resolution
R. Hu et al. / Materials Science and Engineering A 548 (2012) 83– 88 87
F clei which have orientation relationship with the matrix on one side of the GB grow intot ntation relationship with the matrix on either side of the GB grow into the matrix nearb lationship with either side of the matrix grow along the GB.
ctGIwrGae
4a
iaesatat
tnMcGtMaahiaHat
bde
Fig. 7. SEM microstructure near the fracture of the tested alloy aged at 700 ◦C for5 h.
Fig. 8. Orientation imaging microscopy (OIM) map shows the distribution of differ-ent types of GBs in Ni–Cr–W based superalloy aged at 700 ◦C for 5 h. The colorized
ig. 6. Schematic representation of M23C6 growth mechanism: (a) M23C6 carbide nuhe matrix near only one side of that GB; (b) M23C6 carbide nuclei which have orieoth sides of that GB; (c) and (d) M23C6 carbide nuclei which have no orientation re
oherent orientation relationship with the matrix on either side ofhe GB, the nucleus grow into the matrix near both sides of thatBs (Fig. 6b) and the morphology of the carbide is shown in Fig. 3b.
f the carbide nucleus have no coherent orientation relationshipith the matrix, they would grow along GBs (Fig. 6c and d), and the
od-like carbide and small lamellar carbide could precipitate at theBs (Fig. 3c and d). The different morphologies of M23C6 carbideslong GBs (Fig. 3c and d) might be controlled by the activationnergy for M23C6 nucleation at different GBs [6,24–26].
.2. Effect of M23C6 precipitation on the rupture behaviors of theged alloy
When M23C6 carbide precipitates at the GBs, lots of Cr atomsn the matrix are consumed and Cr concentration decreases. As
result, the crystal lattice distortion of the matrix is reduced. Crlement plays an important role on the strength of alloy, and thetrength decreases with the decreasing Cr concentration in super-lloys [27]. Hence, the strength of the alloy is reduced partly byhe GB M23C6 precipitation. The increasing elongation of the agedlloy can also be attributed to the decrease of Cr concentration inhe matrix.
Microstructure of the tested sample (aged at 700 ◦C for 5 h) nearhe fracture is shown in Fig. 7. Due to plenty of GB M23C6 carbide,umbers of cracks generate at the GBs. The size of the cracks nearby6C carbide is much smaller than that at the GBs. It can be con-
luded that the rupture of the aged alloy is mainly influenced by theB M23C6 breaking. GB M23C6 plays an important role during the
ensile deformation process of the aged Ni–Cr–W superalloy [16].ost of the matrix grains which are connected through GBs before
ging treatment are separated by the lamellar M23C6 carbide afterging treatment. It is well known that M23C6 carbide is brittle andas low M23C6-matrix bond strength [28]. When the tensile stress
s large enough, the carbide breaks firstly and cracks are generatedt the GBs. So the strength of the aged alloy decreases obviously.ence, the decreasing tensile and yield strength of the aged alloyre caused by the decrease of Cr concentration in the matrix andhe GB lamellar M23C6 carbide breaking.
Fig. 8 shows the distribution of different types of GBs in Ni–Cr–Wased superalloy aged at 700 ◦C for 5 h. Different types of GBsistribute non-uniformly in the aged alloy, and this results in het-rogeneous precipitation of M23C6. During the tensile deformation
lines represent CSL GBs, and the red lines �3 GBs. The low angle GB is marked asLA, and the rest large angle GBs (not including CSL GBs). (For interpretation of thereferences to color in this figure legend, the reader is referred to the web version ofthe article.)
8 nd En
pltmica
5
es
(
(
(
(
A
P
[
[[[
[
[
[
[
[[[
[[[[[25] M.A. Mangan, M.V. Kral, G. Spanos, Acta Metall. 47 (1999) 4263–4274.[26] K. Kaneko, T. Fukunaga, K. Yamada, N. Nakada, M. Kikuchi, Z. Saghi, J.S. Barnard,
P.A. Midgley, Scripta Mater. 65 (2011) 509–512.[27] A.K. Jena, M.C. Chaturvedi, J. Mater. Sci. 19 (1984) 3121–3139.[28] J. Larson, Metall. Mater. Trans. A 7 (1976) 1497–1502.
8 R. Hu et al. / Materials Science a
rocess, more cracks are generated in the area containing a mass ofarge angle GBs, as a result the tensile sample ruptures easily. Hence,he variation of the tensile properties with increasing aging time is
ainly attributed to heterogeneous precipitation of M23C6 carbiden different specimens. Besides, the volume fraction of GB M23C6ould also cause the variation of tensile properties of Ni–Cr–W alloyt different aging time.
. Conclusions
In the present study, M23C6 precipitation behavior and its influ-nce on tensile properties of Ni–Cr–W based superalloy have beentudied in details. The results are as follows:
1) M23C6 precipitation can be inhibited by the �≤9 GBs and thelow angle GBs in the aged Ni–Cr–W based superalloy and theseGBs have very good corrosion resistance.
2) M23C6 carbides with four typical morphologies precipitate atthe large angle GBs which have high interfacial energy. M23C6grows on the {1 1 1} plane and along 〈1 1 0〉 direction of thematrix. The coherent orientation relationship between M23C6and the matrix plays an important role on the precipitationmorphology of M23C6.
3) Aging treatment decreases both tensile and yield strength, butimproves the elongation of the alloy. The tensile properties varyirregularly with increasing aging time.
4) The decreasing tensile strength and yield strength of the agedalloy are mainly caused by the lamellar M23C6 carbide breaking.Otherwise, the decrease of Cr concentration in the matrix alsodecreases the strength of the alloy.
cknowledgement
The research was supported by NSFC (51171150) and the 111roject (B08040).
gineering A 548 (2012) 83– 88
References
[1] J.C. Zhao, M. Larsen, V. Ravikumar, Mater. Sci. Eng. A 293 (2000) 112–119.[2] M.D. Mathew, K. Bhanu Sankara Rao, S.L. Mannan, Mater. Sci. Eng. A 372 (2004)
327–333.[3] E.A. Trillo, L.E. Murr, Acta Mater. 47 (1998) 235–245.[4] H.Y. Bi, Z.J. Wang, M. Shimada, H. Kokawa, Mater. Lett. 57 (2003) 2803–2806.[5] R.J. Romero, L.E. Murr, Acta Metall. Mater. 43 (1995) 461–469.[6] H.U. Hong, B.S. Rho, S.W. Nam, Mater. Sci. Eng. A 318 (2001) 285–292.[7] B. Sasmal, Metall. Mater. Trans. A 30 (1999) 2791–2801.[8] Y.S. Lim, J.S. Kim, H.P. Kim, H.D. Cho, J. Nucl. Mater. 335 (2004) 108–114.[9] H. Li, S. Xia, B.X. Zhou, W.J. Chen, C.L. Hu, J. Nucl. Mater. 399 (2010)
108–113.10] S. Kihara, J. Newkirk, A. Ohtomo, Y. Saiga, Metall. Mater. Trans. A 11 (1980)
1019–1031.11] A. Koul, R. Castillo, Metall. Mater. Trans. A 19 (1988) 2049–2066.12] T. Garosshen, G. McCarthy, Metall. Mater. Trans. A 16 (1985) 1213–1223.13] L.Z. He, Q. Zheng, X.F. Sun, G.C. Hou, H.R. Guan, Z.Q. Hu, J. Mater. Sci. 40 (2005)
