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Interdiffusion behaviour in aluminide-coated René 80H at 1150°C
~
~
MATERIALS ENGINEERING
E L SEVI E R
Materials Science and Engineering A224 (I 997) 27- 32
Interdiffusion behaviour in aluminide-coated Ren6 80H at 1150°C E. Basuki, A. Crosky, B. Gleeson * School of Materials Science and Engineering, The Unit'ersity of New South II,'ales, Sydney, NSIV 2052. Australia
Recetved 23 August 1996; revised 9 October 1996
Abstract
The hot-section components in commercial aero gas turbines are typicaIly made of a nickel-base superatloy and are protected by a diffusion-aluminide coating. At elevated temperatures this alloy/coating system ts metastable and. as a consequence, microstructural changes occur. The extent to which these changes occur is critically dependent upon temperature and time. In the case of overheating, where the component temperature exceeds about 1100°C, the microstructural changes are often extensive enough to cause a decrease in both the strength and protectiveness of the alloy/coating system in a relatively short period of time. This paper reports the effects of overheating on the microstructural changes that occur in an aluminide-coated, nickel-base superalIoy, Ren6 80H. The overheating condition was simulated by isothermally heating at I150°C for up to 167 h. Both the alloy/coating interdiffusion kinetics and the time-dependent phase changes resulting from interdiffusion are discussed. The interdiffusion behaviour is complex and requires the application of diffusion paths for proper intet~pretation. © 1997 Elsevier Science S.A. Keywords: Aluminide; Diffusion coating; Interdiffusion: Superalloy
1. Introduction
Materials used in high-temperature appIications must possess both strength and corrosion resistance. These requirements, however, are often incompatible and the solution is to apply a corrosion-resistant coating to a strong material. An example of this is provided by the high strength, nickel-base superalloy components used in aero gas turbine engines. Such components are usually protected by a coating based on the intermetallic compound, fl-NiA1. This type of coating is very resistant to oxidation owing to its ability to form exclusively the slow-growing oxide, A1203. At elevated temperatures, the structural stability of an aluminide coating is primarily affected by (i) selective oxidation and hot-corrosion processes, and (ii) interdiffusion between the coating and substrate. Each of these processes has a profound influence on the coating structure during its service life and, hence, on its ability to provide continued protection to the alloy component. The oxidation and hot-corrosion behaviour of aluminide coatings have been extensively studied [1-3] and * Corresponding author. 0921-5093/97/$17.00 © 1997 Elsevier Science S.A. All rights reserved. PIIS0921-5093(96)10556-6
the mechanisms are largely understood. By contrast, studies pertaining to coating/substrate interdiffusion are limited and the phenomenon is consequently less well understood. Those studies which have been conducted [4-8] have shown that coating/substrate interdiffusion in nickel-base superalloys can occur to a significant extent under overheating conditions in which temperatures exceed about 1100°C. Such interdiffusion affects both the mechanical properties and surface stability of the alloy. The aim of this study was to investigate the effects of overheating to 1150°C on the extent of coating/substrate interdiffusion and the resulting microstructurai changes in an aluminide-coated, nickel-base superalloy, Ren~ 80H. 2. Experimental procedures Substrate samples were cut from rods of directionally solidified Ren~ 80H having a principaI phase constitution },-Ni + ; -Ni3(AI, Ti). The nominal chemical composition of Ren~ 80H is given in Table 1. The substrate coupons were approximately 1 cm-" in surface area and were ground to a 1200-grit finish and then ultrasonically cleaned in acetone prior to coating deposition.
E. Basuk~ et a l . / Materials Science and Engineering A224 (1997) 27-32
28
Table l Nominal compomion of Ren6 80H (wt.%) Ni
Cr
AI
Co
Ti
Nio
W
Hf
Zr
B
C
Balance
14.0
3.0
9.5
5.0
4.0
4.0
0.8
0.02
0.01
0.2
An aluminium-rich fl-NiA1 diffusion coating was applied to the substrate coupons using a halide-activated pack cementation process. The pack consisted of a powder mixture of 6 wt.% A1 (master alloy), 0.5 wt.% NHgCI (activator) and 93.5 wt.% A1203 (inert filler). A cylindrical alumina container, about 50 c m 3 in volume, was filled with the pack mixture into which four substrate coupons were embedded. The spacing between each coupon was about 1.5 cm. Coating depositions were carried out by heating a given pack in a resistance tube furnace to 800°C under a flowing argon atmosphere, and holding for a deposition time of 4 h (excluding heat-up time). The pack was then heated to 1050°C and diffusion annealed at this temperature for 1 h in order to obtain primarily a fl-NiA1 coating. The pack was furnace-cooled to room temperature at the completion of the diffusion-annealing treatment. Overheating was simulated by heating the coated coupons to 1150°C in a resistance tube furnace for up to 167 h. The coupons were contained in evacuated silica tubes to preclude unwanted oxidation. The encapsulated samples were rapidly quenched in water at the completion of a given overheating experiment. Cross-sections of both the as-deposited and overheated coatings were prepared using standard metallographic techniques. The polished cross-sections were analysed using optical microscopy, scanning electron microscopy (SEM) and electron probe microanalysis (EPMA). The various phases in the coatings were identified on the basis of the EPMA results. The grain structure and interphase boundaries in the coatings were revealed by etching in a 50 ml H 2 0 + 150 ml HC1+25 g C r O s solution for about 3 s at room temperature. Carbides were revealed by a stain-etching technique using Murakami's reagent, which consists of 10 g KOH + 10 g KaFe(CN)6 + 100 ml H~O.
layer boundary; and an inner layer of aluminium-lean fl + ~ + MC (M = Ti and Mo) + M23C 6 ( M = Cr, M o ) . The average composition of the cr phase, which is known to be a brittle N i - C r - M o - C o compound [9], was found by EPMA to be N i - 2 1 C r - 2 3 M o - g C o 13W-6Ti-4.5A1. (Unless stated otherwise, all compositions are in at.%.) The ~-Cr particles in the outer layer were very small in size ( ~ 0.2 pro) and were typically located along the boundaries of the relatively small /? grains ( ~ 0.5 I;m). The fl grain size in the intermediate layer was comparatively larger ( ~ 8.5 ~tm) and essentially no c~-Cr particles were contained in this layer. The very low volume fraction of c~-Cr particles in the intermediate layer of the coating can be attributed to the lower aluminium content in the/? phase of this layer. This is because the solubility of chromium in fl increases with decreasing aluminium content [10]. The small c-phase precipitates in the inner layer were located along the 13 grain boundaries, and some ¢ phase had penetrated into the substrate. The inner layer in the as-coated
]-
N
,,gh; lf +
~ " High-Al 13 ~ L~IcA+[N~I+~ + , -'-
.
+Ma~Ca
[ /
(a)
.
70-
:.
,
,
Low-AI ~3+ ct-Cr
20 lam . ~
Low-AI !3+ o'+ MC + M23C6
Low-AI p
60"
y+7'+MC + M23C6
50.
3. Results o~
A representative microstructure of the as-coated Ren~ 80H, together with a corresponding concentration profile obtained from EPMA is shown in Fig. 1. The as-deposited coating was about 60 ktm thick and consisted of three layers: an outer layer of aluminium-rich fl-NiA1 plus numerous fine ~-Cr precipitates; an intermediate fl layer in which the aluminium content ranged from about 51 at.% at the outer/intermediate-layer boundary to about 35 at.% at the intermediate/inner-
g 0
40. 30, 20"
0 10
(b)
Surface
20
30
Distance 0am)
40
50
Substrate
Fig. 1. (a) Optical micrograph of the as-coated aluminide coating; and (b) corresponding concentration profiles.
E. Basukz et al./ ?,IateHals Science a~d E , gineeri~zg ,4224 (I997) 27-32
29
80~
zone) formed in the substrate at the coating/substrate interface• The microstructure of the sample after overheating for 16.7 h is given in Fig. 3(a), with the corresponding concentration profiles given in Fig. 3(b). The total thickness of the coating is less than that of the coating overheated for 8.3 h. This recession of the coating commenced at the fl/?, interface. It is also seen in Fig. 3(a) that the ~-phase particles had coarsened and the ?, layer had widened. A relatively large number of MC carbide particles were distributed in the }. layer. The interdiffusion zone contained a mixture of the carbides M23C~, MC and M6C, together with ~. The M~C carbides were of the composition 2 1 M o - 1 2 W - 1 9 N i 1 9 C r - 8 C o - 5 T i - 1 6 C and are believed to be primarily the consequence of a reaction between M 2 3 C 6 and ~, i.e.
40-
~
(Mo + W),-~o~~, +
a0:
o
20-
It is seen in Fig. 3(b) that the outer portion of the coating has completely homogenised to a Ni-rich fl composition of Ni-28A1- 8Cr- 7Co-4Ti-0.6Mo, which very likely corresponds to the solubility limit of Ni ( + Co) in fl-NiAl. Thus, either nickel enrichment into, or aluminium depletion from this composition would result in a destabilisation of the ft. Fig. 4(a) shows the microstructure of the sample overheated for 167 h. The corresponding concentration
","' "'.
" 7 ...-.
.':,''
"'.-.'.,".'.
" ,;;'¢-,- |
Low-AlIS+e
.' ':,~-''; ::" a .... " . . ' .... " " ":'''' 1. i'~' ; ' ' ~-- +~*C+M,,C, ,.
!,,,,., ; . ' . , . . . . . -,: ...... ...., '..,,~1 it . / + M A C • " - ' " , .':. " "'.:.< " : . ' " ~ " ~ " ' " " " ~ :'~" . "; ~,_,-r+MC
Low-Alp
70-
Low-AIp+~+ i~' MC+M23C6 ; +
i .
"f+'/'+MC
60.
20
(o)
surface
40
60 Distance
80
(~rrl)
100
120
~40
Substrate
Fig. 2. (a) O p t i c a l m i c r o g r a p h o f t h e c o a t i n g / s u b s t r a t e r e g i o n a f t e r o v e r h e a t i n g for 8.3 h; a n d (b) c o r r e s p o n d i n g c o n c e n t r a t i o n profiles.
structure contained mainly fl and ~r. Formation of the ~r phase is attributable to the low solubilities of Cr, Mo, W and Co in both fl and 7' [11]. Carbide particles identified as MC-type of approximate stoichiometry (Ti0.77Mo0.12W0.1i)C , were contained in the different layers of the coating• Similar carbides were also present in the substrate. The coating grew primarily by the inward diffusion of aluminium, as evidenced by the absence of pack inclusions in the coating• Such an inward grown coating is typically referred to as a 'high activity' coating [1% Fig. 2(a) shows the microstructural changes that occurred in the coating and underlying substrate region after 8.3 h at 1150°C, with corresponding concentration profiles given in Fig. 2(b). The fine ~-Cr precipitates were almost completely absent from the outer layer of the coating and the grain size of the fi phase increased significantly in all coating layers. Moreover, the total coating thickness increased by about 80% of its ascoated value. The concentration profiles in Fig. 2(b) show that the outer and intermediate layers of the coating had partially homogenised as a result of both aluminium depletion and nickel enrichment. The o-phase particles in the interdiffusion zone lengthened and blocky particles of composition corresponding to the M 2 3 C e carbide ( 6 1 C r - 8 M o - 2 C o - 4 N i - 3 W - 2 1 C ) developed in this zone. A y-Ni layer (i.e. a ?'-denuded
M~C + Cr
M 2 3 C 6 --+
(1)
Low-A113
Low-At [3 + c~ biC + M ~ C ~ +MoC
•+
-7+MC 'g-7'-'
"--_ v," .,~,,e,~< . ~: "' ; -" ,
.
.
,2~ ~4"~, <., ~N/" - ,¢.2" =° ---
-y+y'+MC (a) .k'. "-',<'." b,,,.,
',', :
Low-A113
70-
",",,"
.... ,
~-oW-AI [3i
"f ÷ MC
"
~m _
Y+'r+MC
]
!-~+MC+: ~M23C6+ :
}~6c
,~
at% A] .~s- at% Cr •-x,- at% W at% Ti -c-- at% Co at% Mo
40 ,
8
0Su rfac2(] e
(b)
4(]
60
80
~00
Distance (;tm)
120
140 "~60 Substrate
Fig. 3. (a) O p t i c a l m i c r o g r a p h o f t h e c o a t i n g / s u b s t r a t e r e g i o n a f t e r o v e r h e a t i n g for 16.7 h; a n d (b) c o r r e s p o n d i n g c o n c e n t r a t i o n profiles.
30
E. Basuki et al. /'Materials Science and Engineering A224 (1997) 27-32
it can be inferred that the growth of the 7' layer was due to the counter migration of the /3/7' and 7'/7 interfaces. A comparison between the slopes of the /3 and/3 + 7' curves at any given time further reveals that the migration rate of the/3/7' interface was faster than that of the 7'/7 interface. ,,L
•
~
~'
'
"
et'
>
~L"
7+MC
4. D i s c u s s i o n ,a~ 7"-~' t ~ ~ v "
70
~ .~ "~ 1-':::1"-: -~-'~:'~'_O~m
Low-A113
i Y'
Y+M6C
Y+MC
Overheating of aluminide-coated Ren6 80H at 1150°C resulted in various microstructural changes from the as-coated structure. Of particular interest are the initial homogenisation and consequent thickening of the coating/3 layer, followed by the recession of this layer.
Y+Y+MC
60 ~ " 5O i
-~- at% AI at% Cr .-,4- aft,; W ~ a[~,~TI ~ at% Co
~, 40 i i i
i/
(3 20
............
i
L..
I0
,.'2",:!22;1'2,'~22.......
--,... ....... T~,'r,,~v,
0 (b)
0 20 Surface
40
60
80
100
Distance (gin)
120
140
160 180 Substrate
Fig. 4. (a). Optical micrograph of the coating/substrate region after overheating for 167 h: and (b) corresponding concentration profiles.
profiles are given in Fig. 4(b), which shows that the outer/3-layer composition is unchanged from that after 16.7 h overheating. The extent of coating (i.e. /3) recession was significant, and resulted in the formation of a 7' layer between the/3 and 7 layers (as indicated in Fig. 4(a)). The M6C particles are globular in morphology and are contained within the 7 layer. The M6C volume fraction was greater than that in the sample overheated for 16.7 h, and the MC carbide volume fraction in the }' layer has decreased. The kinetics of coating (i.e. /3 layer) enlargement during homogenisation and subsequent coating recession are shown in Fig. 5. Also included in this figure are the thickening kinetics of the }, and 7' layers. From Fig. 5,/3-NiAI enlargement occurred relatively rapidly at the beginning of overheating, but then slowed upon approaching the transition from enlargement to recession at about 14 h. The instantaneous rate of/3-NiA1 recession (e.g. ~ 0.45 l~m h - t at 20 h) was comparatively slower than that of enlargement (e.g. ~ 7.8 pm h - 1 at 5 h). When the thickness of the }/layer is added to that of /3, as indicated by the dashed-line in Fig. 5, the resulting curve is indicative of the migration rate of the 7'/7 interface. This migration rate was parabolic (i.e., diffusion controlled) with a rate constant of about 3.8 lama h - 2. Thus, on the basis of the kinetics curves in Fig. 5,
4.I. Homogenisation of coating fl layer The driving force for coating/substrate interdiffusion is the differences in thermodynamic activities of the elements in the coating and substrate. Unfortunately, activity data are not available for the multicomponent systems relevant to the present investigation. If, however, concentration differences are approximately proportional to the activity differences, then the diffusion of a given element can be assumed to occur down its concentration gradient. However, equating activities to concentrations is a rather crude approximation for a multicomponent system, since chemical interactions between the various components in the system are disregarded. Notwithstanding, on the basis of the aforementioned approximation and the measured concentration profiles (e.g. Fig. 2(b), Fig. 3(b), and Fig. 4(b)), the elements which diffused from the coating to the substrate during overheating were A1, Cr, Mo and W; whilst Ni, Co, and Ti diffused from the substrate to the coating.
120
100
................1
~- s0 =o e0
Y
-8 z: F-
~
/
40
2O
0
I
]
I
20
40
60
I*'~E 80
100
~
i
l
120
140
160
Time (hour)
Fig. 5. Kinetics of microstructural changes in coating/substrate regions after overheating for up to 167 h.
E. Basuki et al./lkraterials Science alzd Engineer#zg A224 (1997) 27-32
where J, is the flux of component i, the conditions were such that the flux of chromium was greater than the flux of aluminium. In addition, the solubility limit of o--phase forming elements apparently increases as the aluminium content in the fi phase decreases, as indicated by comparing the results in Figs. 1 and 3. Consequently, a certain amount of o--phase dissolution occurred during homogenisation of the coating fl layer which, in turn, resulted in the release of chromium from the o- phase into the surrounding ~ phase.
A1 40
60
~.. so/)
I;1 'x~° %
Zo • ."'" 3
Cr
2o
40 6o Nickel (at.%)
7
7 ,', \
8o
\
Ni
Fig. 6. Diffusion paths for coating/substrate regions after overheating for times of up to 167 h.
Interdiffusion caused the outer and intermediate layers to homogenise to a composition corresponding to Ni-rich fl-NiA1. The time for this homogenisation to occur was approximately 14 h. The homogenisation process coincided with a thickening of the /?-layer, which can be attributed to the transformations ~ + ~' --+ f l ( + o) at the substrate/interdiffusion-zone interface, and fl-NiA1,, + G-+ fl-NiAly (where x > y) at the interdiffusion-zone/coating interface. As shown in Fig. 6, and disregarding the presence of the ¢ phase, these two transformations can be represented as diffusion paths in the 1150°C Ni-Cr-A1 phase diagram. The transition of fl from an Al-rich composition to a more Cr-soluble, Ni-rich composition resulted in the dissolution of o- and any c~-Cr originally located in the interdiffusion zone. The diffusion path for this aluminium-depletion process is represented by arrow I in Fig. 6. The 7'+ ~ ' ~ f l ( ~ ) transformation is represented by the arrow 2 in Fig. 6. Such a diffusion path corresponds primarily to the depletion of nickel from the 7 + Y' substrate. In other words, the rate of nickel diffusion from the substrate to the fl phase of the interdiffusion zone was greater than the rate of aluminium diffusion from the fl phase to the substrate. Formation of the y layer (or 7'-denuded zone) at the coating/substrate interface is believed to be due primarily to the enrichment of chromium into the substrate, as indicated by the arrow 3 in Fig. 6. This enrichment of chromium resulted from the relatively high concentration of chromium in the interdiffusion zone. Thus, even though the diffusivity of aluminium in y-Ni (D21) is greater than that of chromium (D~r) [13], the extent of chromium enrichment into the ~ + 7' subsurface region was greater than that of aluminium. This is because the concentration gradient of chromium (dCcffdx) was much greater than that of aluminium (dCal/dx), as indicated by the concentration profiles of Cr and A1 in the 7 layer (see Fig. 2(b)). In other words, on the basis of Fick's first law for diffusion,
di= Di(dCffdx)
31
(2)
4.2. fl recession After the fl layer in the coating had homogenised to a maximum level of nickel enrichment (or, conversely, atuminium depletion), the fl layer began to recede by the transformation fi ~ 7. This transformation was due to the diffusion of aluminium from the fi/Y interface to the ?,-layer, and nickel from the y-layer to the fl/?, interface. A fl-~ 7' transformation did not occur in the early stages of the fl recession due to the relatively high chromium content in the fi + o- + M•3C 6 + M6C zone. However, with continued recession, together with continued reaction M 2 3 C 6 -+-o---+M 6 C - 1 - C r , the fl phase was no longer in contact with the Cr-enriched region, thus causing the chromium content in the fl phase to decrease and the fl ~ ?,' transformation to be possible. According to the N i - C r - A 1 phase diagram in Fig. 6, this transformation could occur only when the fl composition at the interface corresponded closely to fl + ~ + ?" equilibrium (i.e. N i - 2 8 A 1 - 8 C r 7Co-4Ti-0.6Mo).
5. Conclusions The investigation into the effects of overheating aluminide-coated Ren6 80H at 1150°C has revealed that several microstructural changes occur. The coating enlarged until it had entirely homogenised to a fl-NiA1 composition of maximum nickel content. Recession of the fl-phase started from dissolution of fl in the/3 + ointerdiffusion zone, followed by a /3--, 7' transformation. The phase transformations, as well as the microstructural changes, are best explained through the application of diffusion paths in the N i - C r - A I phase diagram.
Acknowledgements This work was supported by an Australian Research Council Small Grant. The authors wish to acknowledge the support of E.B. through an AusAID Fellowship.
32
E. Basukl et al. /Materials Science and Eugineerit~g A224 (1997) 27-32
References [1] B Gleeson, W.H. Cheung, W. Da Costa, and D.J. Young, Oxid. Met., 38 (1992) 407. [2] G.H. Meier and F.S. Pettit, S m f Coat. Tec/mol., 39/40 (1989) 1. [3] T.N. Rhys-Jones, Mater. Sct. Technol., 4 (1988) 421. [4] K.L. Luthra and M.R. Jackson, in M. Khobaib and R.C. Kruenat (eds.), Higk Temperature Coatings, The Minerals, Metals and Materials Society, Warrendale, PA, 1987, p. 85. [5] C.W.M. Lee, E. Basuki, A. Crosky, and B. Gleeson, Co@ Proc. 5th Austrahan Aeronautical Conference, The Institution of Engineers, Australia, National Conference Publication No. 93/6,
1993, p. 369. [6] J.E. Mortal and M.S. Thompson, Stof Coat. Tec/mol., 4.3/44 (1990) 371. [7] J.E. Mortal and R.H. Barkalow, Scr. Metall., 16 (I982) 593. [8] T.K. Redden, Trans. Metal/. Soc. AIME, 242 (I968) 1695. [9] E.W. Ross, J. Met., 19 (1967) ili9. [10] S.M. Merchant and M.R. Notis, Mater. Sei. Eng., 66 (I984) 47. [11] C.C. Jia, K. Ishida, and T. Nishizawa, Metall. Trans. A, 25A (1994) 473. [12] G.W. Goward and D.H. Boone, Oxid. ?v[et., 3 (1971) 475. [13] J.A. Nesbitt and R.W. Heckel, Metall. Trans. A, I8,4 (I987) 2061.
~
MATERIALS ENGINEERING
E L SEVI E R
Materials Science and Engineering A224 (I 997) 27- 32
Interdiffusion behaviour in aluminide-coated Ren6 80H at 1150°C E. Basuki, A. Crosky, B. Gleeson * School of Materials Science and Engineering, The Unit'ersity of New South II,'ales, Sydney, NSIV 2052. Australia
Recetved 23 August 1996; revised 9 October 1996
Abstract
The hot-section components in commercial aero gas turbines are typicaIly made of a nickel-base superatloy and are protected by a diffusion-aluminide coating. At elevated temperatures this alloy/coating system ts metastable and. as a consequence, microstructural changes occur. The extent to which these changes occur is critically dependent upon temperature and time. In the case of overheating, where the component temperature exceeds about 1100°C, the microstructural changes are often extensive enough to cause a decrease in both the strength and protectiveness of the alloy/coating system in a relatively short period of time. This paper reports the effects of overheating on the microstructural changes that occur in an aluminide-coated, nickel-base superalIoy, Ren6 80H. The overheating condition was simulated by isothermally heating at I150°C for up to 167 h. Both the alloy/coating interdiffusion kinetics and the time-dependent phase changes resulting from interdiffusion are discussed. The interdiffusion behaviour is complex and requires the application of diffusion paths for proper intet~pretation. © 1997 Elsevier Science S.A. Keywords: Aluminide; Diffusion coating; Interdiffusion: Superalloy
1. Introduction
Materials used in high-temperature appIications must possess both strength and corrosion resistance. These requirements, however, are often incompatible and the solution is to apply a corrosion-resistant coating to a strong material. An example of this is provided by the high strength, nickel-base superalloy components used in aero gas turbine engines. Such components are usually protected by a coating based on the intermetallic compound, fl-NiA1. This type of coating is very resistant to oxidation owing to its ability to form exclusively the slow-growing oxide, A1203. At elevated temperatures, the structural stability of an aluminide coating is primarily affected by (i) selective oxidation and hot-corrosion processes, and (ii) interdiffusion between the coating and substrate. Each of these processes has a profound influence on the coating structure during its service life and, hence, on its ability to provide continued protection to the alloy component. The oxidation and hot-corrosion behaviour of aluminide coatings have been extensively studied [1-3] and * Corresponding author. 0921-5093/97/$17.00 © 1997 Elsevier Science S.A. All rights reserved. PIIS0921-5093(96)10556-6
the mechanisms are largely understood. By contrast, studies pertaining to coating/substrate interdiffusion are limited and the phenomenon is consequently less well understood. Those studies which have been conducted [4-8] have shown that coating/substrate interdiffusion in nickel-base superalloys can occur to a significant extent under overheating conditions in which temperatures exceed about 1100°C. Such interdiffusion affects both the mechanical properties and surface stability of the alloy. The aim of this study was to investigate the effects of overheating to 1150°C on the extent of coating/substrate interdiffusion and the resulting microstructurai changes in an aluminide-coated, nickel-base superalloy, Ren~ 80H. 2. Experimental procedures Substrate samples were cut from rods of directionally solidified Ren~ 80H having a principaI phase constitution },-Ni + ; -Ni3(AI, Ti). The nominal chemical composition of Ren~ 80H is given in Table 1. The substrate coupons were approximately 1 cm-" in surface area and were ground to a 1200-grit finish and then ultrasonically cleaned in acetone prior to coating deposition.
E. Basuk~ et a l . / Materials Science and Engineering A224 (1997) 27-32
28
Table l Nominal compomion of Ren6 80H (wt.%) Ni
Cr
AI
Co
Ti
Nio
W
Hf
Zr
B
C
Balance
14.0
3.0
9.5
5.0
4.0
4.0
0.8
0.02
0.01
0.2
An aluminium-rich fl-NiA1 diffusion coating was applied to the substrate coupons using a halide-activated pack cementation process. The pack consisted of a powder mixture of 6 wt.% A1 (master alloy), 0.5 wt.% NHgCI (activator) and 93.5 wt.% A1203 (inert filler). A cylindrical alumina container, about 50 c m 3 in volume, was filled with the pack mixture into which four substrate coupons were embedded. The spacing between each coupon was about 1.5 cm. Coating depositions were carried out by heating a given pack in a resistance tube furnace to 800°C under a flowing argon atmosphere, and holding for a deposition time of 4 h (excluding heat-up time). The pack was then heated to 1050°C and diffusion annealed at this temperature for 1 h in order to obtain primarily a fl-NiA1 coating. The pack was furnace-cooled to room temperature at the completion of the diffusion-annealing treatment. Overheating was simulated by heating the coated coupons to 1150°C in a resistance tube furnace for up to 167 h. The coupons were contained in evacuated silica tubes to preclude unwanted oxidation. The encapsulated samples were rapidly quenched in water at the completion of a given overheating experiment. Cross-sections of both the as-deposited and overheated coatings were prepared using standard metallographic techniques. The polished cross-sections were analysed using optical microscopy, scanning electron microscopy (SEM) and electron probe microanalysis (EPMA). The various phases in the coatings were identified on the basis of the EPMA results. The grain structure and interphase boundaries in the coatings were revealed by etching in a 50 ml H 2 0 + 150 ml HC1+25 g C r O s solution for about 3 s at room temperature. Carbides were revealed by a stain-etching technique using Murakami's reagent, which consists of 10 g KOH + 10 g KaFe(CN)6 + 100 ml H~O.
layer boundary; and an inner layer of aluminium-lean fl + ~ + MC (M = Ti and Mo) + M23C 6 ( M = Cr, M o ) . The average composition of the cr phase, which is known to be a brittle N i - C r - M o - C o compound [9], was found by EPMA to be N i - 2 1 C r - 2 3 M o - g C o 13W-6Ti-4.5A1. (Unless stated otherwise, all compositions are in at.%.) The ~-Cr particles in the outer layer were very small in size ( ~ 0.2 pro) and were typically located along the boundaries of the relatively small /? grains ( ~ 0.5 I;m). The fl grain size in the intermediate layer was comparatively larger ( ~ 8.5 ~tm) and essentially no c~-Cr particles were contained in this layer. The very low volume fraction of c~-Cr particles in the intermediate layer of the coating can be attributed to the lower aluminium content in the/? phase of this layer. This is because the solubility of chromium in fl increases with decreasing aluminium content [10]. The small c-phase precipitates in the inner layer were located along the 13 grain boundaries, and some ¢ phase had penetrated into the substrate. The inner layer in the as-coated
]-
N
,,gh; lf +
~ " High-Al 13 ~ L~IcA+[N~I+~ + , -'-
.
+Ma~Ca
[ /
(a)
.
70-
:.
,
,
Low-AI ~3+ ct-Cr
20 lam . ~
Low-AI !3+ o'+ MC + M23C6
Low-AI p
60"
y+7'+MC + M23C6
50.
3. Results o~
A representative microstructure of the as-coated Ren~ 80H, together with a corresponding concentration profile obtained from EPMA is shown in Fig. 1. The as-deposited coating was about 60 ktm thick and consisted of three layers: an outer layer of aluminium-rich fl-NiA1 plus numerous fine ~-Cr precipitates; an intermediate fl layer in which the aluminium content ranged from about 51 at.% at the outer/intermediate-layer boundary to about 35 at.% at the intermediate/inner-
g 0
40. 30, 20"
0 10
(b)
Surface
20
30
Distance 0am)
40
50
Substrate
Fig. 1. (a) Optical micrograph of the as-coated aluminide coating; and (b) corresponding concentration profiles.
E. Basukz et al./ ?,IateHals Science a~d E , gineeri~zg ,4224 (I997) 27-32
29
80~
zone) formed in the substrate at the coating/substrate interface• The microstructure of the sample after overheating for 16.7 h is given in Fig. 3(a), with the corresponding concentration profiles given in Fig. 3(b). The total thickness of the coating is less than that of the coating overheated for 8.3 h. This recession of the coating commenced at the fl/?, interface. It is also seen in Fig. 3(a) that the ~-phase particles had coarsened and the ?, layer had widened. A relatively large number of MC carbide particles were distributed in the }. layer. The interdiffusion zone contained a mixture of the carbides M23C~, MC and M6C, together with ~. The M~C carbides were of the composition 2 1 M o - 1 2 W - 1 9 N i 1 9 C r - 8 C o - 5 T i - 1 6 C and are believed to be primarily the consequence of a reaction between M 2 3 C 6 and ~, i.e.
40-
~
(Mo + W),-~o~~, +
a0:
o
20-
It is seen in Fig. 3(b) that the outer portion of the coating has completely homogenised to a Ni-rich fl composition of Ni-28A1- 8Cr- 7Co-4Ti-0.6Mo, which very likely corresponds to the solubility limit of Ni ( + Co) in fl-NiAl. Thus, either nickel enrichment into, or aluminium depletion from this composition would result in a destabilisation of the ft. Fig. 4(a) shows the microstructure of the sample overheated for 167 h. The corresponding concentration
","' "'.
" 7 ...-.
.':,''
"'.-.'.,".'.
" ,;;'¢-,- |
Low-AlIS+e
.' ':,~-''; ::" a .... " . . ' .... " " ":'''' 1. i'~' ; ' ' ~-- +~*C+M,,C, ,.
!,,,,., ; . ' . , . . . . . -,: ...... ...., '..,,~1 it . / + M A C • " - ' " , .':. " "'.:.< " : . ' " ~ " ~ " ' " " " ~ :'~" . "; ~,_,-r+MC
Low-Alp
70-
Low-AIp+~+ i~' MC+M23C6 ; +
i .
"f+'/'+MC
60.
20
(o)
surface
40
60 Distance
80
(~rrl)
100
120
~40
Substrate
Fig. 2. (a) O p t i c a l m i c r o g r a p h o f t h e c o a t i n g / s u b s t r a t e r e g i o n a f t e r o v e r h e a t i n g for 8.3 h; a n d (b) c o r r e s p o n d i n g c o n c e n t r a t i o n profiles.
structure contained mainly fl and ~r. Formation of the ~r phase is attributable to the low solubilities of Cr, Mo, W and Co in both fl and 7' [11]. Carbide particles identified as MC-type of approximate stoichiometry (Ti0.77Mo0.12W0.1i)C , were contained in the different layers of the coating• Similar carbides were also present in the substrate. The coating grew primarily by the inward diffusion of aluminium, as evidenced by the absence of pack inclusions in the coating• Such an inward grown coating is typically referred to as a 'high activity' coating [1% Fig. 2(a) shows the microstructural changes that occurred in the coating and underlying substrate region after 8.3 h at 1150°C, with corresponding concentration profiles given in Fig. 2(b). The fine ~-Cr precipitates were almost completely absent from the outer layer of the coating and the grain size of the fi phase increased significantly in all coating layers. Moreover, the total coating thickness increased by about 80% of its ascoated value. The concentration profiles in Fig. 2(b) show that the outer and intermediate layers of the coating had partially homogenised as a result of both aluminium depletion and nickel enrichment. The o-phase particles in the interdiffusion zone lengthened and blocky particles of composition corresponding to the M 2 3 C e carbide ( 6 1 C r - 8 M o - 2 C o - 4 N i - 3 W - 2 1 C ) developed in this zone. A y-Ni layer (i.e. a ?'-denuded
M~C + Cr
M 2 3 C 6 --+
(1)
Low-A113
Low-At [3 + c~ biC + M ~ C ~ +MoC
•+
-7+MC 'g-7'-'
"--_ v," .,~,,e,~< . ~: "' ; -" ,
.
.
,2~ ~4"~, <., ~N/" - ,¢.2" =° ---
-y+y'+MC (a) .k'. "-',<'." b,,,.,
',', :
Low-A113
70-
",",,"
.... ,
~-oW-AI [3i
"f ÷ MC
"
~m _
Y+'r+MC
]
!-~+MC+: ~M23C6+ :
}~6c
,~
at% A] .~s- at% Cr •-x,- at% W at% Ti -c-- at% Co at% Mo
40 ,
8
0Su rfac2(] e
(b)
4(]
60
80
~00
Distance (;tm)
120
140 "~60 Substrate
Fig. 3. (a) O p t i c a l m i c r o g r a p h o f t h e c o a t i n g / s u b s t r a t e r e g i o n a f t e r o v e r h e a t i n g for 16.7 h; a n d (b) c o r r e s p o n d i n g c o n c e n t r a t i o n profiles.
30
E. Basuki et al. /'Materials Science and Engineering A224 (1997) 27-32
it can be inferred that the growth of the 7' layer was due to the counter migration of the /3/7' and 7'/7 interfaces. A comparison between the slopes of the /3 and/3 + 7' curves at any given time further reveals that the migration rate of the/3/7' interface was faster than that of the 7'/7 interface. ,,L
•
~
~'
'
"
et'
>
~L"
7+MC
4. D i s c u s s i o n ,a~ 7"-~' t ~ ~ v "
70
~ .~ "~ 1-':::1"-: -~-'~:'~'_O~m
Low-A113
i Y'
Y+M6C
Y+MC
Overheating of aluminide-coated Ren6 80H at 1150°C resulted in various microstructural changes from the as-coated structure. Of particular interest are the initial homogenisation and consequent thickening of the coating/3 layer, followed by the recession of this layer.
Y+Y+MC
60 ~ " 5O i
-~- at% AI at% Cr .-,4- aft,; W ~ a[~,~TI ~ at% Co
~, 40 i i i
i/
(3 20
............
i
L..
I0
,.'2",:!22;1'2,'~22.......
--,... ....... T~,'r,,~v,
0 (b)
0 20 Surface
40
60
80
100
Distance (gin)
120
140
160 180 Substrate
Fig. 4. (a). Optical micrograph of the coating/substrate region after overheating for 167 h: and (b) corresponding concentration profiles.
profiles are given in Fig. 4(b), which shows that the outer/3-layer composition is unchanged from that after 16.7 h overheating. The extent of coating (i.e. /3) recession was significant, and resulted in the formation of a 7' layer between the/3 and 7 layers (as indicated in Fig. 4(a)). The M6C particles are globular in morphology and are contained within the 7 layer. The M6C volume fraction was greater than that in the sample overheated for 16.7 h, and the MC carbide volume fraction in the }' layer has decreased. The kinetics of coating (i.e. /3 layer) enlargement during homogenisation and subsequent coating recession are shown in Fig. 5. Also included in this figure are the thickening kinetics of the }, and 7' layers. From Fig. 5,/3-NiAI enlargement occurred relatively rapidly at the beginning of overheating, but then slowed upon approaching the transition from enlargement to recession at about 14 h. The instantaneous rate of/3-NiA1 recession (e.g. ~ 0.45 l~m h - t at 20 h) was comparatively slower than that of enlargement (e.g. ~ 7.8 pm h - 1 at 5 h). When the thickness of the }/layer is added to that of /3, as indicated by the dashed-line in Fig. 5, the resulting curve is indicative of the migration rate of the 7'/7 interface. This migration rate was parabolic (i.e., diffusion controlled) with a rate constant of about 3.8 lama h - 2. Thus, on the basis of the kinetics curves in Fig. 5,
4.I. Homogenisation of coating fl layer The driving force for coating/substrate interdiffusion is the differences in thermodynamic activities of the elements in the coating and substrate. Unfortunately, activity data are not available for the multicomponent systems relevant to the present investigation. If, however, concentration differences are approximately proportional to the activity differences, then the diffusion of a given element can be assumed to occur down its concentration gradient. However, equating activities to concentrations is a rather crude approximation for a multicomponent system, since chemical interactions between the various components in the system are disregarded. Notwithstanding, on the basis of the aforementioned approximation and the measured concentration profiles (e.g. Fig. 2(b), Fig. 3(b), and Fig. 4(b)), the elements which diffused from the coating to the substrate during overheating were A1, Cr, Mo and W; whilst Ni, Co, and Ti diffused from the substrate to the coating.
120
100
................1
~- s0 =o e0
Y
-8 z: F-
~
/
40
2O
0
I
]
I
20
40
60
I*'~E 80
100
~
i
l
120
140
160
Time (hour)
Fig. 5. Kinetics of microstructural changes in coating/substrate regions after overheating for up to 167 h.
E. Basuki et al./lkraterials Science alzd Engineer#zg A224 (1997) 27-32
where J, is the flux of component i, the conditions were such that the flux of chromium was greater than the flux of aluminium. In addition, the solubility limit of o--phase forming elements apparently increases as the aluminium content in the fi phase decreases, as indicated by comparing the results in Figs. 1 and 3. Consequently, a certain amount of o--phase dissolution occurred during homogenisation of the coating fl layer which, in turn, resulted in the release of chromium from the o- phase into the surrounding ~ phase.
A1 40
60
~.. so/)
I;1 'x~° %
Zo • ."'" 3
Cr
2o
40 6o Nickel (at.%)
7
7 ,', \
8o
\
Ni
Fig. 6. Diffusion paths for coating/substrate regions after overheating for times of up to 167 h.
Interdiffusion caused the outer and intermediate layers to homogenise to a composition corresponding to Ni-rich fl-NiA1. The time for this homogenisation to occur was approximately 14 h. The homogenisation process coincided with a thickening of the /?-layer, which can be attributed to the transformations ~ + ~' --+ f l ( + o) at the substrate/interdiffusion-zone interface, and fl-NiA1,, + G-+ fl-NiAly (where x > y) at the interdiffusion-zone/coating interface. As shown in Fig. 6, and disregarding the presence of the ¢ phase, these two transformations can be represented as diffusion paths in the 1150°C Ni-Cr-A1 phase diagram. The transition of fl from an Al-rich composition to a more Cr-soluble, Ni-rich composition resulted in the dissolution of o- and any c~-Cr originally located in the interdiffusion zone. The diffusion path for this aluminium-depletion process is represented by arrow I in Fig. 6. The 7'+ ~ ' ~ f l ( ~ ) transformation is represented by the arrow 2 in Fig. 6. Such a diffusion path corresponds primarily to the depletion of nickel from the 7 + Y' substrate. In other words, the rate of nickel diffusion from the substrate to the fl phase of the interdiffusion zone was greater than the rate of aluminium diffusion from the fl phase to the substrate. Formation of the y layer (or 7'-denuded zone) at the coating/substrate interface is believed to be due primarily to the enrichment of chromium into the substrate, as indicated by the arrow 3 in Fig. 6. This enrichment of chromium resulted from the relatively high concentration of chromium in the interdiffusion zone. Thus, even though the diffusivity of aluminium in y-Ni (D21) is greater than that of chromium (D~r) [13], the extent of chromium enrichment into the ~ + 7' subsurface region was greater than that of aluminium. This is because the concentration gradient of chromium (dCcffdx) was much greater than that of aluminium (dCal/dx), as indicated by the concentration profiles of Cr and A1 in the 7 layer (see Fig. 2(b)). In other words, on the basis of Fick's first law for diffusion,
di= Di(dCffdx)
31
(2)
4.2. fl recession After the fl layer in the coating had homogenised to a maximum level of nickel enrichment (or, conversely, atuminium depletion), the fl layer began to recede by the transformation fi ~ 7. This transformation was due to the diffusion of aluminium from the fi/Y interface to the ?,-layer, and nickel from the y-layer to the fl/?, interface. A fl-~ 7' transformation did not occur in the early stages of the fl recession due to the relatively high chromium content in the fi + o- + M•3C 6 + M6C zone. However, with continued recession, together with continued reaction M 2 3 C 6 -+-o---+M 6 C - 1 - C r , the fl phase was no longer in contact with the Cr-enriched region, thus causing the chromium content in the fl phase to decrease and the fl ~ ?,' transformation to be possible. According to the N i - C r - A 1 phase diagram in Fig. 6, this transformation could occur only when the fl composition at the interface corresponded closely to fl + ~ + ?" equilibrium (i.e. N i - 2 8 A 1 - 8 C r 7Co-4Ti-0.6Mo).
5. Conclusions The investigation into the effects of overheating aluminide-coated Ren6 80H at 1150°C has revealed that several microstructural changes occur. The coating enlarged until it had entirely homogenised to a fl-NiA1 composition of maximum nickel content. Recession of the fl-phase started from dissolution of fl in the/3 + ointerdiffusion zone, followed by a /3--, 7' transformation. The phase transformations, as well as the microstructural changes, are best explained through the application of diffusion paths in the N i - C r - A I phase diagram.
Acknowledgements This work was supported by an Australian Research Council Small Grant. The authors wish to acknowledge the support of E.B. through an AusAID Fellowship.
32
E. Basukl et al. /Materials Science and Eugineerit~g A224 (1997) 27-32
References [1] B Gleeson, W.H. Cheung, W. Da Costa, and D.J. Young, Oxid. Met., 38 (1992) 407. [2] G.H. Meier and F.S. Pettit, S m f Coat. Tec/mol., 39/40 (1989) 1. [3] T.N. Rhys-Jones, Mater. Sct. Technol., 4 (1988) 421. [4] K.L. Luthra and M.R. Jackson, in M. Khobaib and R.C. Kruenat (eds.), Higk Temperature Coatings, The Minerals, Metals and Materials Society, Warrendale, PA, 1987, p. 85. [5] C.W.M. Lee, E. Basuki, A. Crosky, and B. Gleeson, Co@ Proc. 5th Austrahan Aeronautical Conference, The Institution of Engineers, Australia, National Conference Publication No. 93/6,
1993, p. 369. [6] J.E. Mortal and M.S. Thompson, Stof Coat. Tec/mol., 4.3/44 (1990) 371. [7] J.E. Mortal and R.H. Barkalow, Scr. Metall., 16 (I982) 593. [8] T.K. Redden, Trans. Metal/. Soc. AIME, 242 (I968) 1695. [9] E.W. Ross, J. Met., 19 (1967) ili9. [10] S.M. Merchant and M.R. Notis, Mater. Sei. Eng., 66 (I984) 47. [11] C.C. Jia, K. Ishida, and T. Nishizawa, Metall. Trans. A, 25A (1994) 473. [12] G.W. Goward and D.H. Boone, Oxid. ?v[et., 3 (1971) 475. [13] J.A. Nesbitt and R.W. Heckel, Metall. Trans. A, I8,4 (I987) 2061.