solid state sic/ni alloy reaction

Solid State SiC/Ni Alloy Reaction M.R. JACKSON, R.L. MEHAN, A. M. DAVIS, and E.L. HALL

The solid state reaction between silicon carbide and a model superalloy consisting of 70 at. pct Ni, 20 at. pct Cr, and 10 at. pet A1 was studied between 700 ~ and 1150 ~ for times ranging from "0" hours to 330 hours. Reaction couples consisting of SiC/Ni, SiC/Cr, and SiC/NiCr were also studied. The reactions were carded out in air with the materials, in the shape of discs, maintained in contact under a pressure of 7 MPa. A reaction was detected with SiC and the model alloy at all temperatures studied, and the reaction was diffusion controlled with an activation energy of 184 kJ/mole. In the ceramic the reaction was dominated by the diffusion of Ni into the ceramic forming a banded structure consisting of alternating layers of 8-Ni2Si and a two phase mixture of graphite and 8. On the metal side, the reaction was very dependent on the presence of alloying elements, with pure Ni reacting to the greatest extent, followed by the binary NiCr alloy, and finally by NiCrA1. The growth and presence of the phases detected in these reactions is consistent with phase equilibria concepts.

I. INTRODUCTION

SILICON-based structural ceramics such as SiC and Si3N4 are under active consideration for structural use in hot machinery such as gas turbines. In these applications, ceramics and metals will be in contact with each other at temperatures of at least 1000 ~ Because the ceramic and metal must be in intimate contact in order to prevent vi- bration induced stresses in the ceramic, oxygen will tend to be excluded at the interface. Under such conditions, the protective oxide films may not form, and the possibility exists for a chemical reaction between the metal and ceramic. This may lead to severe degradation of the strength of both materials in service.

Ceramic-metal reactions under inert or reducing conditions were the subject of intensive study some fifteen years ago in connection with the development of metal matrix composites, and much of this work has been summarized. 1.2 It has been known since 1966, for example, that SiC filaments react in a reducing environment with Fe, Co, and Ni at temperatures as low as 700 ~ 3 More recently, Comie et al reported on the reactivity of SiC filaments with a number of nickel-based superalloys, as well as with an Fe-Cr-AI-Y alloy. 4 All the alloys reacted to some extent with SiC during hot isostatic pressing in vacuum at 996 ~ forming silicides in both metal and filament. The existence of these and similar reactions has been the primary barrier to the development of metal matrix composites using boron, SiC, or A1203 filaments in any metal matrix with a melting point higher than Al.

More recently, the reactions of SiC, SiaN4, and reaction- bonded SiC with complex superalloys have been studied. 5'6"7 The intent of these studies was to assess directly the degree of chemical interaction between these materials under conditions similar to those that could be encountered in hot machinery applications. However, most of this past work, including that performed with reinforcing filaments, was phenomenological in nature because of the complex chemistry of superalloys. No real attempt has been made to study

M. R. JACKSON, R. L. MEHAN, A. M. DAVIS, and E. L. HALL are all Members, Technical Staff, General Electric Corporate Research and Development, Schenectady, NY 12345.

Manuscript submitted March 3, 1981.

METALLURGICAL TRANSACTIONS A

the fundamental nature of such reactions, nor have phase equilibria concepts been applied to the reactions.

The present work was undertaken to apply such tech- niques to reactions between silicon-based ceramics and metal. In order to simplify the chemistry of the metal/ ceramic reaction, and at the same time obtain results appli- cable to actual superalloys, a model superalloy consisting only of Ni, Cr, and A1 was used in these studies. Several ceramics have been used, but the present paper will consider only SiC in detail, although some results pertaining to reaction-bonded SiC will be discussed. More detailed discussions of the behavior of the reaction-bonded material, as well as Si3N4, will be presented elsewhere.

II. EXPERIMENTAL

A. Materials

The majority of the experiments were conducted using hot-pressed SiC manufactured by the Norton Company. Some work with sintered SiC fabricated in this laboratory 8 was also conducted. Behavior of these two materials in reaction was similar, except for minor phases and structural variations caused by the presence o f WC in the hot pressed material. The predominant phases were identical. The reaction-bonded ceramics containing free silicon used in some of the study came from two sources: NC-435 from the Norton Company, and Silcomp T M fabricated in this labora-

*TM--Trademark of the General Electric Company

tory. 9 No difference in reactivity between the latter two materials was noted.

The model superalloy was fabricated in this laboratory and had a composition of 70 at. pct Ni, 20 at. pct Cr, 10 at. petAl (75.8 wt pct Ni, 19.2 wt pct Cr, 5.0 wt petAl). In this model alloy, the complexity is lessened by the fact that only nickel and chromium are silicide formers. The rationale for this choice is as follows: In general, high strength nickel-based superalloys being considered for service with SiC have microstructures predominantly of 3 t + 7', where 7 is the face centered cubic nickel solid solution and 7' is an ordered structure based on NiaA1. The

ISSN 0360-2133/83/0311-0355500.75/0 �9 1983 AMERICAN SOCIETY FOR METALS AND VOLUME laA, MARCH 1983--355

THE METALLURGICAL SOCIETY OF AIME

remaining small volume fraction (<0.02) consists primarily of carbides and borides which are expected to have a minor role in reaction with SiC. One approach to developing a model composition, therefore, is to eliminate all phases except 7 and 7'. The elements frequently employed for solid solution strengthening of 7 are chromium, molybdenum, and tungsten at total levels of 14 to 30 at. pct. To approximate the solid solution elements and to provide oxidation resistance, the model alloy contains 20 at. pct chromium. The 7' phase in commercial materials is generally strength- ened by partial substitution of aluminum by titanium, tanta- lum, and niobium. The total of aluminum and these other elements is in the range of 6 to 12 at. pct; thus, the model alloy contains 10 at. pct aluminum to approximate the composition of the 7' forming elements. Exclusion of Mo, W, Ti, Ta, and Nb might be considered as too large a sim- plification. However, previous studies 5 on commercial Ni- base alloys did not indicate a dominance of the reactions by these elements. Several experiments with Ni-Cr at a composition of 80 at. pct Ni, 20 at. pct Cr (81.8 wt pct Ni, 18.2 wt pct Cr), and with pure Ni and Cr were also performed in order to gain further insight into the pertinent phase relationships.

B. Reaction Couples

The ceramic specimens were machined into coupons 6.35 mm in diameter and 3.18 mm thick. The metal specimens had the same thickness but were 12.7 mm in diameter. Dimensions were maintained to 0.25 mm. Prior to assembly, the metal specimens were polished on 400 and 600 mesh SiC paper, and the ceramics were used in the as- ground condition; both were cleaned with acetone prior to assembly in a reaction couple. The couple usually consisted of a ceramic/metal/ceramic sandwich and was placed into the apparatus shown schematically in Figure 1. In some cases the order was reversed (metal/ceramic/metal). Gener- ally, two different ceramics were run against the same metal piece. The couple was loaded to 222 N, leading to a stress of 7 MPa at the interface. The specimens were surrounded by a platinum wound tube furnace (not shown in Figure 1), brought to temperature in about 30 minutes, and held there for the required time, ranging from "0" time to 330 hours. The furnace ends were covered with insulation, but an inert atmosphere was not used. An air atmosphere was chosen deliberately as being representative of actual service conditions. This procedure does allow the possibility of oxygen entering into the reaction; however, once the reaction between SiC and the metal begins, oxygen is effectively excluded. Experimentation in a vacuum environment (10 -7 torr) and 900 ~ to duplicate otherwise the air experiments showed a metallographic reaction structure very similar in depth and detail to that obtained in the air environment.

After a reaction run had been completed, the ceramic and metal pieces were removed, cut in two, examined metallographically, and then subjected to other analytical investigations. Initially, the ceramic and metal coupons were separated and sectioned individually, but it was found that during this process some pieces of the reaction product were lost. The sectioning method adopted, at least in those cases where sufficient reaction occurred to cause the individual coupons to adhere to each other, was to encase the

entire reaction "sandwich" in clear plastic and then section the mount. A macrophotograph (Figure 2) showing such a section will be discussed subsequently.

C. Analytical Method

The microstructural and chemical analysis was performed using electron microchemical probe analysis (EMPA), X-ray diffraction (XRD), transmission electron microscopy (TEM), and X-ray spectroscopy in a scanning transmission electron microscope (STEM).

For the XRD analysis, phases and particles were removed from each reaction zone in the polished specimens using

Fig. 1--Schematic diagram of experimental apparatus.

I I 2 mm

Fig. 2--Macrophotograph of Si-SiC, SiC, and NiCrA1 alloy after 50 h at 1150 ~

356--VOLUME 14A, MARCH 1983 METALLURGICAL TRANSACTIONS A

fine-tipped dissecting probes, and the extricated material analyzed by the Debye-Scherrer technique. After the polished specimens were examined, selective etching was employed to separate phases in the various reaction zones, and to extract phases that might have been undetected previously. Three different etches were used in this in- vestigation, namely:

1. 20 pct H3PO4 in H20 (electrolytic) 2. 7 pct HC1 in methanol (electrolytic) 3. 5 pet bromine in methanol (chemical)

In cases where more than one etch was used in a single zone, the specimen was repolished between etches to insure that no residual effects of the preceding etchant would remain.

In preparation for examination using TEM/STEM, the samples were demounted from the epoxy and two types of sectioning employed. The first of these was sectioning parallel to the reaction zone interfaces (normal to the compression direction). In this way a planar section containing only one type of zone could be produced. To isolate the zones present on the metal side of the original metal/ceramic interface, the metal coupon was cut using a diamond wa- feting blade and metallographically polished to isolate a chosen reaction zone. The reaction zones in the ceramic coupon were isolated by either cleavage or wafering with a diamond blade, followed by metallographic polishing to a thickness of -25 /xm. Final preparation of the metal reaction zone samples was accomplished by electropolishing in a solution of 20 pct perchloric acid in methanol at - 4 0 ~ and 25 V. The ceramic reaction zone samples were prepared by ion milling using a beam of argon ions accelerated through a potential 6 kV. In the second case, the samples were sectioned parallel to the compression axis, producing a wafer with appearance similar to Figure 2. This wafer was mechanically polished and then ion milled until a large hole intersecting the areas of interest was formed. In this way all of the reaction zones could be studied in a single specimen with their spatial relationships preserved.

Transmission electron microscopy was performed on a JEOL JSEM 200 TEM/STEM operating at 200 kV, or an Hitachi H-600-1 TEM/STEM operating at 100 kV. Bright- field imaging in conjunction with selected-area electron diffraction was used to identify crystallographically the phases present in each reaction zone. In addition, the microscopes were equipped with STEM capabilities, so that a fine probe of electrons could be focused on phases or particles in the structure, and energy dispersive X-ray spectroscopy used to determine the chemistry of the features of interest. The Cliff-Lorimer technique 1~ was used for the chemical analysis. Absorption was taken into account using procedures which will be more fully described in a subsequent paper, H and which are identical to those discussed by Goldstein.t2

A molten Si-rich liquid was extruded at the Si-SiC/metal side, and obviously a great deal of liquid phase interdiffusion took place. By contrast, the hot pressed SiC/metal side showed a much smaller (but still severe) zone of reaction. From the uniformity of reaction depth in the macrophotograph in Figure 2, the apparatus shown in Figure 1 distributes the load uniformly.

A typical hot-pressed SiC/NiCrA1 reaction, in this case at 1150 ~ is shown in a montage microphotograph, Figure 3. The gap, presumably caused by a thermal expansion mis- match between the reacted SiC and the metal (see Figure 2) has been removed for clarity. The reaction consists of three distinct zones: a banded structure in the SiC at the top of the photograph, a transition reaction zone (TRZ) in the metal, and a metal reaction zone (MRZ). Not evident from the photomicrograph is an area beyond the metal reaction zone where Si has diffused into the NiCrA1 to a maximum concentration equal to the solubility limit. The original reaction couple interface is at the junction of the banded region and the transition zone. This was determined by dimensional measurements and by metallography at the edge of the ceramic disc, and was verified by the use of fine quartz filament markers in selected experiments. These metallographic features persist in decreasing quantities as the reaction temperature and time are reduced. Figure 4 shows the reaction between SiC and the model NiCrAI alloy at 900 ~ The same general features are seen to be present as at higher temperature, although the phase morphology in the metal reaction zone is different. Figures 5 and 6 show the results of typical EMPA traces through the reaction zones of the ceramic-metal couples depicted in the previous two micrographs. The similarity of the two analyses is apparent, except at 900 ~ there is only a very small Si diffusion zone, while at 1150 ~ Si has diffused more than 200/zm into the

HI. RESULTS

A. Nature of the Reaction

Metallographically, a reaction was detected in both the model alloy and SiC at all temperatures investigated (700 ~ to 1150 ~ More severe reactions were detected with the reaction-bonded (Si-SiC) material. This is illustrated in Figure 2, where 222N was applied at 1150 ~ for 50 hours.

Fig. 3--React ion zones in hot-pressed SiC and NiCrAI alloy after 100 h at 1150 ~ SiC is at top of figure.

METALLURGICAL TRANSACTIONS A VOLUME 14A, MARCH 1983--357

Fig. 6--Electron microprobe chemical analyses for Ni, Cr, Si, and A1 in a SiC/NiCrAI couple treated 220 h at 900 ~

Fig. 4--Reaction zones in hot-pressed SiC and NiCrAI alloy after 220 h at 900 ~ SiC is at top of figure.

Fig. 5--Electron microprobe chemical analyses for Ni, Cr, Si, and AI in SiC/NiCrA1 couple treated 100 h at 1150 ~

metal after all metallographic evidence of a chemical reaction has vanished. This is likely due to greater solubility and diffusivity of Si in NiCrAI at the higher temperature. In addition, a Cr-rich phase has formed in the MRZ at 900 ~ For the reaction at 1150 ~ (Figures 3 and 5), a carbon analysis was performed and the results are shown in Table I. As may be seen, a carbon gradient is present in the banded structure of the SiC reaction zone and is vanishingly small in the metal reaction zone.

The dark bands in the SiC reaction zone shown in Figures 3 and 4 contain graphite, 8-Ni2Si, and a small amount of a Cr, Si carbide (Crs_xSi3_yfx+y), while the light bands contain 8-Ni2Si and trace amounts of the Cr, Si carbide. The exact values of x and y are unknown, but XRD, TEM, and X-ray spectroscopy in STEM all indicate the

358--VOLUME 14A, MARCH 1983

Table I. Carbon Concentration in SiC/Metal Reaction Zone*

Microstructural Feature Wt Pet C

A. SiC 35 B. SiC Reaction Zone

1 st dark layer 31 2nd light layer 1.8 3rd dark layer 26 4th light layer 2.2 5th dark layer 30 6th light layer Not determined 7th dark layer 28 8th light layer Not determined 9th dark layer 10

10th light layer Not determined 11 th dark layer 9

C. Transition Zone (a) Adjacent to SiC reaction zone 4.5 (b) Center of zone 4.0 (c) Adjacent to metal 3.0

D. Metal reaction zone ~ 1.0 *Using diamond standard. Absolute accuracy estimated to be about

25 pet.

presence of this phase. The Cr, Si carbide will be identified as A carbide in the remainder of the discussion, to indicate the uncertainty.

Finer details of the structure are shown in the TEM micrographs of Figures 7 and 8. Some XRD, TEM, and STEM analysis was performed in the sintered SiC because the bands were broader than in the hot-pressed material. Chemi- cally, the two materials are identical except for the presence of WC (about 8 wt pet) in the hot-pressed material, presumably introduced during ball-milling the powder. In Figure 7 the graphite, which has a fibrous morphology, is light compared with the dark 8-Ni2Si due to preferential thinning of the graphite. Light microscopy and SEM performed on etched samples indicated that the graphite grains are equiaxed, and that the 8-Ni2Si forms a semicontinuous network structure around the graphite. The small amount of


] "1

0.5/zm Fig. 7--Transmission electron micrograph of the structure of the dark bands in the banded reaction zone in SiC, consisting of &Ni2Si (dark) and graphite (light).

Cr shown to be present in the dark bands from the electron microprobe trace (Figure 5) occurs as the A carbide. This latter compound is found in the 8-Ni2Si/graphite network as a discrete phase, with a grain size on the order of 1 to 2/xm. X-ray spectroscopy in STEM indicates that some Ni is soluble in this phase. Other unidentified phases are present in the bands and elsewhere in the reaction areas. The size and volume of these phases were generally too small to allow for analysis, an indication of the complexity of the mixed silicides that may form. It should be noted that even in the layers containing graphite, the predominant species in terms of weight fraction are Ni and Si, as can be seen in Figure 5.

The transition zone (arbitrarily named because it exists between a clearly defined ceramic and metal reaction area) consists primarily of a matrix of fcc 7/'-Cr3Ni2SiC. The lattice parameter (a0) of this phase was calculated to be 1.0625 nm - 0.0001 from X-ray diffraction studies. In this matrix, localized regions of a-A1203 are found, accounting for the Al-rich regions shown on the electron microprobe scan of this zone. The a-A1203 is most likely to have formed on the metal surface during heating to the reaction temperature, prior to the beginning of reaction. The distribution of these two phases is shown in Figure 9, which is an SEM micrograph of a sample which has been etched in bromine/ methanol. The a-A1203 is present in the cavities in the r/' matrix, and is distributed in the form of small particles, 0.05 to 0.2/zm matrix in diameter. It is unlikely, but some A1203 may have been introduced during metallographic preparation.

In Figure 3, this zone appears to consist of a two-phase columnar or lath structure. TEM/STEM analysis revealed that the phase which has contrast similar to the metal matrix in Figure 3 is y' , with approximate composition 79 wt pct Ni, 8 wt pct Cr, 5 wt pct Si, and 8 wt pct A1. In at. pct, this corresponds to 24 pct (AI + Si), indicating that Si is present on the A1 sublattice and Cr on the Ni sublattice in this METALLURGICAL TRANSACTIONS A

I I 0.2/xm

Fig. 8--Transmission electron micrograph of the structure of the light bands in the banded reaction zone in SiC, consisting of &Ni2Si.

I I 20/zm

Fig. 9 - - Scanning electron micrograph of the transition reaction zone after etching in bromine/methanol.

ordered structure. 13 The darker phase in the MRZ in Figure 3 consisted of a complex multiphase structure with a /3-NiA1 matrix and small grains of a-Cr and ~ , a complex Ni-Si-A1 compound, 14 distributed uniformly throughout. The Z phase of undetermined crystal structure was reported by Guard and Smith 14 to be approximately Ni2(Si,AI) with a narrow range of allowed AI and Si contents of about 9 to 11 wt pct each. Electron diffraction indicated that ~ is heavily faulted on parallel planes. A crystal structure similar to that of &Ni2Si is indicated by XRD and TEM for ~ , although the ~ and ~ phases are not miscible. The

VOLUME 14A, MARCH 1983--359

phase is unrelated to 0(Ni2Si), either in chemistry or in crystallography. Neither a-Cr nor ~ were observed in the MRZ by XRD, presumably due to the small volume fraction of these phases present in the metal reaction zone.

The phase composition of the MRZ appears to be very dependent upon reaction temperature. At 1150 ~ a region consisting of alternating y' and "-2-- plus Cr3Si areas was found in the MRZ near the TRZ, with the remainder of the MRZ as described above. At 900 ~ as previously men- tioned, the MRZ is composed of a Ni-rich phase and a Cr-rich phase. Based on the microprobe data and unpublished research on the Si/NiCrAI reaction, 15 these two phases may be y' and a-Cr. However, further work is necessary to describe fully the temperature effects on the MRZ.

The Si diffusion zone, as well as the model superalloy matrix, consisted of a uniform dispersion of y' precipitates in a y matrix, with the size and shape of the y' varying with Si content. A complete description of the phase distribution and chemistry in the MRZ, Si diffusion zone, and metal matrix will be presented in a subsequent paper.11

B. Effect of Alloy Chemistry

In order to gain additional understanding of the effect of the individual constituents in the overall SiC-model alloy reaction, reactivity experiments with Ni, Cr, and an alloy consisting of 80 at. pct Ni and 20 at. pct Cr were performed at several temperatures using only the hot pressed SiC; the results are discussed below.

With Ni alone at both 1000 ~ and 1150 ~ the same banded structure observed in the model superalloy is formed in the SiC reaction zone. This is seen clearly in Figure 10, which is a micrograph of the SiC-Ni interface after a 100 hour exposure at 1000 ~ The microprobe trace through this reacted area is shown in Figure 11. The bands again consist of alternating layers of &Ni2Si and 8-NizSi plus graphite. TEM verified that the structures in the bands were similar to that in Figures 7 and 8, but the graphite grains were much smaller. A layer of coarse-grained &Ni:Si plus lesser amounts of NiO, WC (from the SiC powder fabri- cation), and an unidentified phase were detected by XRD adjacent to the metal. At a reaction temperature of 1150 ~ the extent and severity of the reaction increases, but the basic chemical nature remains unchanged. In view of the

well-known ability of nickel to act as a graphitizer, as well as not forming a carbide, the above observations are not surprising. It seems clear that Ni is responsible for much of the reaction observed in the SiC by the model NiCrAI superalloy. However, the complete absence of the transition reaction zone and metal reaction zone in Figures 10 and 11 indicates that Cr and AI play a significant role in the overall reaction.

At 1000 ~ with a 100 hour reaction time, Cr and SiC react to a very small extent. Metallographically, a layer about 10/.tm was found on the SiC, and no detectable reaction was observed in the Cr. This represents less than one tenth the depth of Ni/SiC reaction. Microanalysis suggested the presence of a chromium carbide, and perhaps the A carbide. X-ray diffraction identified Cr~Si and several carbide phases including Cr7C3, Cr23C6, and A. Many phases form but the depth of the reaction is so limited that Cr by itself has much less effect on the severity of the reaction than does elemental Ni.

The reaction between a binary alloy consisting of 80 at. pct Ni and 20 at. pct Cr at 1150 ~ and 100 hours is

Fig. 11 --Electron microprobe chemical analyses for Ni and Si in a SiC/Ni couple treated 100 h at 1000 ~

Fig. 10--Reaction zones in hot-pressed SiC and Ni after 100 h at 1000 ~

360--VOLUME 14A, MARCH 1983

Fig. 12--Reaction zones in hot-pressed SiC and NiCr after 100 h at 1150 ~


shown in the micrograph of Figure 12. Although a small amount of Cr is observed by EMPA to have diffused into the SiC reaction zone, the banded structure in the SiC reaction zone is still dominated by the presence of graphite and 6-NizSi, caused by the diffusion of Ni into the SiC. However, in the metal, microanalysis and X-ray diffraction of the reaction zone are consistent with the formation of ~/'-Cr3Ni2SiC. The ~7' phase was positively identified for reactions involving NiCrA1, so its presence as a product of the SiC-NiCr reaction is not unexpected. In both types of couples, ~7' occurs adjacent to the 6 + graphite layered structures. Evidently, AI has the effect of adding an additional zone to the reaction of SiC and the NiCrA1 model superalloy.

C. Reaction Kinetics

The kinetics of the reaction between SiC and the model superalloy have been studied systematically. On the as- sumption that the reaction is diffusion controlled, and hence the reaction zone thickness follows a parabolic growth law, the thicknesses of the SiC reaction zone have been plotted at various temperatures as a function of the square root of reaction time in Figure 13. Similar comparisons were made for the thicknesses of reaction zones on the metal side of the couple (transition and metal reaction zones). As may be seen by the linear relationships, the reaction does seem to be diffusion controlled. By plotting the logarithm of the rate of reaction thickness, in cm 2 s-I, vs the reciprocal of the absolute temperature, the activation energy for the process can be calculated. The Arrhenius plot is shown in Figure 14, and the slope was found to be 184 U/mole with a correlation coefficient of 0.99. Reactions both on the SiC side and metal side can be described by the same activation energy.

Also shown in Figure 14 is the square root-time plot of the reaction between SiC and an actual superalloy, Ren6 77.* The reaction rates are similar, lending credence

*Ren6 77:14.5 Cr, 15.3 Co, 3.4 Ti, 4.3 A1, 4.2 Mo, 0.5 Fe, 0.07 C, 0.016 B, 0.04M Zr, balance Ni (wt pcts).

to the use of a simple alloy to model a superalloy. It also may be noted that these reaction rates are similar to those reported by Cornie et al for reactions involving SiC filaments and a number of different metal compositions fabricated in an inert environment. 4 These activation energies are similar to those found for M2Si silicides formed by reacting elemental Si and M (M = Ni, Co, Pd, Pt, and Mg), which are of the order of 146 kJ/mole at lower temperatures (100 ~ to 700 ~

IV. DISCUSSION

A. SiC-Ni Reaction

Although most of the experimental work reported dealt with SiC/NiCrAI, the reactions can best be understood by first considering the SiC/Ni couples. The variation of chemistry from the SiC terminal across the reaction zone to the Ni terminal can be plotted on a ternary phase diagram for the simplest case. The binary diagrams Ni-C, Ni-Si, and Si-C are well-characterized, ]7'1s'19 but the ternary Ni-Si-C system has not been treated. A pseudo-ternary diagram hypothe- sized previously by two of the present authors, 2~ based on

X I0 "2 4

3

i 2

X 10 -2

4

�9 3

z 2 0

I

0 _- ~O*C / ./,,oo.c

- o / j / / / ~ ~______e----8oo'c

o 200 ,oo ,oo 800 ,ooo TIME, SEC 1/2

(a)

- / 1 5 0 * C

_ IIO0*C

IO00*C

- - o v 900~

~ 700 ~

200 400 600 800 I000 TIME, SEC U2

(b) Fig. 13--Reaction depths in (a) SiC and (b) NiCrA1 v s square root of time for various times and temperatures.

reaction studies performed at that time, is a reasonable representation, but may contain some errors.

A diagram for Ni-Si-C that is consistent with the reaction path observed for SiC/Ni couples is shown in Figure 15, along with the reaction path. It is unclear whether a 0-Ni2Si + graphite (CG) phase field exists, and the diagram has been drawn without such a field. At high temperatures the 0 phase extends to the 6-Ni2Si composition, but at lower temperatures 6-Ni2Si and 0 are two distinct phases. The weight percents of Si in 6 and 0 are very close, however, and may not be distinguishable for the fine, complex microstructure such as shown in Figure 10. X-ray diffraction indicates 6-Ni2Si, with a trace determination of the 0 phase. However, it is possible that at reaction temperatures 0 was present, but was converted to 6 while the couples were cooling from the reaction temperatures.

The diagram of Figure 15 thus shows two-phase equilibria between 6 and SiC, between 6 and CG, and between 6 and Ni. The reaction path drawn on the diagram is simply a plot of the chemical composition across the reaction zone. The path begins at the SiC terminal, adjacent to which is a layer


I000 900 BOO 700% I ~ I I I I

I100

f o SiC SIDE

�9 ", NiCrAI ALLOY SIDE

T

10 "9 cO

e, l

0

,.e ~ io-lO

~ Io-n Ni-BASE SUPERALLOY/SiC (Q =227 KJ MOLE -l)

10-12~1 I 7

KJ MOLE -I FOR SiC AND NiCrAI

\ o :

I I o 8 9 I0 X 10 .4

RECIPROCAL.ABSOLUTE TEMPERATURE, ~

Fig. 14--Reaction rates in SiC and NiCrA1 alloy v s reciprocal absolute temperature. Reaction rates for SiC/superalloy from Ref. 5.

of 8 + CG. In a ternary system, three phases coexist only when the composition of each of the phases is invariant. In a reaction couple, then, a three-phase equilibrium can exist only at an interface, since there can be no chemical potential difference to act as a driving force for diffusional growth. 21 In the reaction path of Figure 15, the three phase field SiC + Co + 8 is crossed with a dashed line to indicate it represents no physical distance in the actual couple, but is the interface between SiC and the ~ + Co layer.

The microstructural and microchemical analyses of Fig- ures 10 and 11 show that the structure below that first

+ Co layer consists of alternating layers of &Ni2Si and + Co. This is plotted in Figure 15, where the reaction path

nearly follows tie-lines in the two-phase field of t~ + Co. The path first enters the 8 phase field (the clear white layers of Figure 10), then to 8 + C6, and back to 8. Each suc- ceeding layer of 8 + Cc is shown to be lower in carbon content (i.e., a lower volume fraction of C6). This has not been verified directly for SiC/Ni couples, since microprobe analysis for C was done by difference (100 pct-pct Si-pct Ni) for most reaction couples. However, the direct microanalysis of carbon reported in Table I for a SiC/NiCrA1 couple showed a decrease in carbon content in the 8 + CG layers from the SiC side to the metal side.

The layered structure ends with a broad layer of the clear 8-Ni2Si adjacent to the Ni terminal of the couple. This is shown in the reaction path of Figure 15, as the last layer is in the single-phase 8-Ni2Si field. The path leaves the t~ field and crosses the Ni + 8 field along a tie-line, since 8 and Ni coexist in the microstructure only at an interface.

,-, t f ~, ~ Ni2Si 0 LIQUID

Fig. 15 - - Hypothetical phase diagram at 1150 ~ for Ni-Si-C on which is superimposed the path followed by reaction of Ni with SiC.

This completes the description of the SiC/Ni reaction path at 1150 ~ The path shown is thermodynamically al- lowable since the chemical potential for each element varies monotonically from one terminal to the other. As noted earlier, less reaction occurs at lower temperature, but the microstructural features are quite similar and can be understood from the same reaction path.

B. SiC-NiCr Reac6on

The addition of Cr to the Ni terminal results in a new layer of reaction for the SiC/NiCr couples. The path could be plotted in a quaternary phase equilibria space, but the microstructure in Figure 12, together with microanalytical data, suggests that this may not be necessary. A layered structure, similar to that for SiC/Ni, is the dominant feature on the SiC side of the reaction, with the major new feature being the addition of a region of ~7'-Cr3Ni2SiC appearing between the banded layers and the NiCr terminal. The NiCr terminal contains 20 at. pct Cr for a Ni/Cr ratio of 4. In the layered zone of the reaction, however, the Ni/Cr ratio is generally greater than 50. At the original, SiC/NiCr interface, Ni diffuses across to react and form the alternating layers of 8 and of ~ + Co. Almost no Cr crosses the interface, so that a concentration of Cr builds up on the metal side of the interface due to the departure of Ni. As Si and C diffuse into the metal side, their reaction with the concentrated Cr and the remaining Ni forms t/' carbide.

The reaction path for the bulk of the couple is again as shown in Figure 15, with the variation of high Ni and Si to low Ni and Si as the structure progresses from 8 to 8 + Co. The small amount of Cr that does diffuse into SiC is present in this region primarily in the form of the A carbide as was observed in the more detailed study of SiC/NiCrAI couples. A quaternary phase diagram would need to include this phase in a three-phase field being traversed in the alternating layers, but the volume fraction of A carbide is quite small. The major difference, requiring quaternary equilibrium description, occurs in the concentrated Cr zone. The reaction path leaves the ~ and ~ + Co layers, possibly along a tie-line in the two-phase field between 8 and r / through the single phase 7/' field, then along a tie-line to single


phase NiCr solid solution, ending at the Ni 20 at. pct Cr composition.

Although Ni and Cr move through the reaction path in quite different fashion, the Si and C fluxes in the opposite direction are apparently nearly equal and little porosity de- velops. For example, the reaction of Ni and SiC would have mass balance if the reaction were:

2Ni + SiC ~ &Ni2Si + Cc

On an atomic basis, there would be three times as much formed as graphite, due to the addition of two Ni atoms for each Si atom. Because of density differences, this would require 80 vol pct 8 and 20 vol pct C~ if mass balance is maintained. A rough estimate of the microstructures indicates the graphite-free 8 regions are - 5 0 vol pct of the structure, so that the fraction of the ~ + Cc layer that is graphite, on average, will be 40 pct. This will vary along the reaction path, since the carbon content of the (5 + C6 layer decreases with increasing distance from the SiC terminal. In the rl' carbide zone, Si and C masses are balanced naturally by the carbide formula, Cr3Ni2SiC. In the Ni or NiCr terminal, Si leads C diffusion because of the greater solubility of Si in the y lattice.

C. SiC-NiCrA1 Reaction

Reaction paths for SiC/NiCrAI could be plotted in all their complexity on a quinary representation of the equilibrium phase diagram. Microchemical analyses such as shown in Figure 5, however, indicate that A1 is nearly totally excluded from the SiC side of the reaction, appearing occasionally in the transition reaction zone as isolated A1203 particles, and being more concentrated in the metal reaction zone as a more homogeneous structure. More Cr is present in the SiC reaction zone than was the case for SiC/NiCr couples. However, the Ni/Cr ratio is generally greater than 20 in the SiC reaction zone, compared to 3.5 in the initial metal terminal. Again, the initial portion of the reaction path in Figure 15 is characteristic of the layered structure of SiC/NiCrA1 couples. Some small amount of )t carbide was observed in the layers, but the alternation of ~ and ~ + C~ is the dominant feature. Beyond the layers is the r/' carbide zone resulting from the concentration of Cr, as noted for NiCr. The isolated particles of A1203 probably formed during heating to the reaction temperature, before reaction occurred. Otherwise, even this small amount of A1 would probably have been rejected from the growing A carbide zone, but otherwise A1 is excluded from the 7?' carbide zone. Essentially, the reaction path passes from the &Ni2Si single- phase field along a tie-line to the 71' carbide single-phase field, behaving as though AI were absent and the couple were between NiCr and SiC.

At this point in the couple, an Al-rich metal reaction zone is interjected between 7/' and the NiCrAI terminal. The original couple interface was at the layered structure/r/' interface. Therefore, the AI concentration in the metal reaction zone is due in part to outward diffusion of Ni and Cr, but in larger part to actual rejection of A1 by the growth of the 7?' zone. As TEM and X-ray diffraction observations indicate, the structure in the metal reaction zone is ex- tremely fine and complex. Details of the structure vary considerably with temperature as seen in Figures 3 and 4. At 1150 ~ y' , /3, a, and Y~. were found where Y~. is the Ni2

(Si,A1) ternary phase identified by Guard and Smith. t4 An accurate quinary description, if one existed, would be useful to describe these structures since they involve all five elements. The phase relations in the MRZ will be discussed in a subsequent paper.ll

D. Rates of Reaction

The data of Figures 13 and 14 indicate the reactions are governed by parabolic growth rate dependency for the SiC/NiCrA1 couples. The activation energy determined from the data is 184 kJ/mole, but the question remains whether this is due to interdiffusion of the elements involved, or whether the reaction is governed by formation of oxides at the original interface, prior to reaction, since the experiments were performed in air. In air experiments, NiO was occasionally observed in the SiC/Ni couples, and A1203 was occasionally observed in the SiC/NiCrAI couples. Activation energies are unavailable for Si or C diffusion through these oxides or for Ni diffusion through A1203. For Ni diffusion in NiO, as well as for O and AI diffusion through A1203, activation energies of 255 kJ/mole are commonly reported. 22 These energies are much higher than observed here.

In addition, the amount of A1 found locally in the transition reaction zone was not strongly dependent on reaction time or temperature (see, for example, Figures 5 and "6). This indicates that A1203 formation is a transient, rather than an on-going, occurrence. During heating to the reaction temperature, some oxidation of the metal and ceramic surfaces occurs. However, at relatively low temperatures the metal is deformed elastically and plastically, conforming to the shape of the SiC wafer. Although the compressive stress is only 7 MPa when averaged over the contact area, it would be considerably greater locally for any asperities on the two surfaces, so that deformation to produce exactly mating surfaces occurs, as shown in Figure 2. At this point, easy access of oxygen to the reaction surfaces is eliminated, and further reaction occurs in essentially inert conditions. As discussed previously, an experiment was conducted to duplicate the temperature and loading conditions in vacuum environment. As noted, the depths of reactivity on both the metal and ceramic sides, as well as the major features of the reaction, were very similar to those observed in air reactions of SiC/NiCrAI couples. The activation energies determined are believed to represent the actual reaction between SiC and NiCrA1, and not to be governed by oxidation or the presence of the isolated oxide particles. This is supported by similar results reported by Cornie 4 for SiC filaments reacting with different metal matrices in inert environments.

Reaction rate data for SiC/Ni and SiC/NiCr couples are not sufficiently complete to allow determination of activation energies for those cases. However, comparisons can be made relative to the degree of reaction for constant reaction time and temperature experiments. Reaction at 1150 ~ for 100 hours produces a total reaction depth (metal and ceramic) of 1.7 nun for SiC/Ni, 0.8 mm for SiC/NiCr, and 0.5 mm for SiC/NiCrA1 couples. Reaction on the SiC end of the couple is dependent on Ni diffusion, since Cr and AI are not strongly involved in the ceramic. When Cr is present in the alloy, 71' carbide is formed, interjecting a barrier that effectively reduces the Ni source strength to which SiC is exposed. With addition of AI to NiCr, not only does 7?'


carbide form, but a complex metal reaction occurs, creating another barrier to diffusion of Ni to SiC.

For these simple model alloys, both Cr and A1 act to retard the reaction of Ni and SiC. As shown in Figure 14, a Ni base superalloy, with all its elemental additions of strong carbide and silicide formers, shows still lower reaction rates with SiC than does NiCrA1. As the alloy becomes more complex, the number of reaction products increases. The form the microstructure takes on reaction appears to present a barrier to rapid reaction by slowing the diffusive processes and by reducing the driving forces for reaction.

V. SUMMARY

Reaction of SiC with Ni-base alloys occurs readily over the temperature range of 700 to 1150 ~ Kinetics of growth in thickness of reaction products on either side of the original metal/ceramic interface follow a parabolic relation with time for a Ni 20 at. pct Cr 10 at. pct AI alloy, and the activation energy for the process was found to be 184 kJ/mole.

The dominant structural feature formed in the SiC is a banded structure. Alternate bands are first &Ni2Si + Cc (graphite) followed by &Ni2Si. Depending on composition of the substrate, other phases may be present in minor amounts within the banded structure, but the ~ and Cc phases dominated, independent of the metal composition. Reaction products on the metal side were very dependent on metal composition. In pure Ni, a thick 8 layer formed. When NiCr reacted with SiC, Cr was concentrated on the metal side, being nearly excluded from reaction in the SiC. With Si and C diffusing from the ceramic, the Cr concentration resulted in formation of ~7'-Cr3Ni2SiC. With A1 present in the metal, it was also excluded from the SiC. Rejection of AI by the growth of t/' carbide resulted in a concentration of AI in a zone containing y', /3-NiA1, and several other phases, depending on reaction temperature.

Reaction was most rapid for pure Ni. Reaction depth was decreased when Cr was present, and was further decreased when AI was present as well. Compared to complex superalloys, with Ti, Mo, W, Co, C, B, and Zr, reaction was more rapid for the model NiCrA1 alloy. Additional complexity of the metal composition apparently causes more complex reaction products, but this in turn leads to less overall reaction.

VI. CONCLUSIONS

Based on the work presented in this study, the following conclusions may be drawn:

1. SiC reacts with nickel and nickel-base alloys under inert conditions at temperatures of 700 ~ and above.

2. The reaction is quite complex even in relatively simple alloy systems, such as NiCrAI. At least 14 individual phases have been identified, and more are known to be present.

3. The formation of many and complex phases reduces the severity and extent of the SiC/metal reaction. Reaction

was most rapid with pure nickel, and least for the ternary NiCrAI alloy.

4. Phase equilibria concepts aid in the understanding of the gross features of such reactions, but work is needed to understand the detailed mechanisms of the formation of the fine-scale structures observed.

ACKNOWLEDGMENTS

Contributing to the results of the TEM and STEM sample preparation and micrographs were E.F. Koch and J. Dunlap. C.D. Robertson prepared the SEM micrographs. This work was performed under the sponsorship of the Materials Science Division of the United States Depart- ment of Energy on Contract DE-AC02-79ERI0413 with Dr. Robert J. Gottschall acting as Project Monitor.

REFERENCES 1. A.G. Metcalfe: in Interfaces in Metal Matrix Composites, A.G.

Metcalfe, ed., Academic Press, New York, NY, 1974, pp. 65o123. 2. K.G. Kreider: in Metallic Matrix Composites, V.G. Kreider, ed.,

Academic Press, New York, NY, 1974. 3. R. E Schneidmiller and J. E. White: SAMPE, 1966, vol. 10, p. E-53. 4. J.A. Cornie, C.S. Cook, and C.A. Anderson: NASA-CR-134956,

1976. 5. R.L. Mehan and R. B. Bolon: J. Mater. Sci., 1979, vol. 14, p. 2471. 6. M.J. Bennett and M.R. Houlton: J. Mater. Sci., 1979, vol. 14,

p. 184. 7. R.L. Mehan and D. W. McKee: J. Mater. Sci., 1976, vol. 11, p. 1009. 8. S. Prochazka: in Ceramics for High Performance Applications,

J.J. Burke, A.E. Gormin, and R.N. Katz, eds., Brook Hill, MA, 1974, pp. 239-52.

9. W.B. Hillig, R.L. Mehan, C.R. Morelock, V.J. DeCarlo, and W. Laskow: Am. Ceram. Soc. Bulletin, 1975, vol. 54, p. 1054.

10. G. Cliff and G.W. Lorimer: Proceedings of the Fifth European Congress on Electron Microscopy, Institute of Physics, London, England, 1972, pp. 140-41.

11. E.L. Hall, Y.M. Kouh, M.R. Jackson, and R.L. Mehan: General Electric Company, Schenectady, NY, unpublished research, 1982.

12. J.I. Goldstein: Introduction to Analytical Electron Microscopy, J.J. Hren, J.l . Goldstein, and D.C. Joy, eds., Plenum Press, New York, NY, 1979, pp. 83-120.

13. R.W. Guard and J. H. Westbrook: Trans. TMS-AIME, 1959, vol. 215, pp. 807-14.

14. R.W. Guard and E. A. Smith: J. Inst. Metals,,, 1959, vol. 88, p. 369. 15. E.L. Hall, R.L. Mehan, and M.R. Jackson: General Electric

Company, Schenectady, NY, unpublished research, 1981. 16. K.N. Tu and J. W. Mayer: Thin Films--lnterdiffusion and Reactions,

J. M. Poate, K. N. Tu, and J. W. Mayer, eds., John Wiley, New York, NY, 1979, p. 359.

17. M. Hansen and K. Anderko: Constitution of Binary Alloys, 2nd edition, McGraw-Hill, New York, NY, 1958, pp. 374-76, 378-80, 1039-42.

18. R.P. Elliott: Constitution of Binary Alloys, First Supplement, McGraw-Hill, New York, NY, 1965, pp. 227-29.

19. E A. Shunk: Constitution of Binary Alloys, Second Supplement, McGraw-Hill, New York, NY, 1969, p. 555.

20. R.L. Mehan and M. R. Jackson: Surfaces and Interfaces in Ceramic and Ceramic-Metal Systems, Materials Science Research, Plenum Press, 1980, vol. 14, pp. 513-24.

21. J.S. Kirkaldy and L.C. Brown: Can. Met. Quart., 1963, vol. 2, p. 89.

22. C.E. BirchenalI: Diffusion, Am. Soc. for Metals, Metals Park, OH, 1973, pp. 309-32.


solid state sic/ni alloy reaction

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