Silicide coatings for carburization protection

Materials Science and Engineering, 88 (1987) 81-87

Silicide Coatings for Carburization

81

Protection

*

G. SOUTHWELL, S. MACALPINE and D. J. YOUNG

School of Chemical Engineering and Industrial Chemistry, University of New South Wales, P.O. Box 1, Kensington, New South Wales 2033 (Australia) (Received April 30, 1986)

ABSTRACT

Silicide coatings were f o r m e d on two austenitic heat-resistant steels by chemical vapour deposition at high temperatures from flowing H 2-silane mixtures. Coatings produced using tetrachlorosilane were grown to thicknesses o f up to 60 p m and contained nickel silicides, (Cr, Ni) silicides and iron silicides. These coatings were porous near the coating-alloy interface and were prone to separation when subject to temperature cycling. The effectiveness o f these coatings in reducing carburization when subsequently exposed to H2-hydrocarbon mixtures at 1273 K was negligible. Coatings were also produced using dichlorodimethylsilane which pyrolyses to produce both silicon and carbon. The resultant coatings were o f similar composition but w i t h o u t the porosity. These coatings also were liable to separation on temperature cycling. Examination o f the substrate alloys revealed that extensive carburization had occurred during formation o f the coating. Subsequent exposure tests in H2-hydrocarbon atmospheres showed that the coatings had no effect. It is concluded that silicide coatings are n o t very effective against carburization in non-oxidizing atmospheres.

1. INTRODUCTION Industrial pyrolysis of hydrocarbon feedstocks in steam-cracking furnaces results in the formation of carbon as an unavoidable byproduct. At the high temperatures of operation, carbon is able to dissolve and dif*Paper presented at the International Symposium on High Temperature Corrosion, Universitd de Provence, Marseille 13331, France, July 7-11, 1986. 0025-5416/87/$3.50

fuse into the Fe-Cr-Ni alloy constituting the furnace tube. This leads to extensive carburization and thus to premature furnace tube failure. Carburization kinetics for commercially employed alloys have been shown [1-5] to obey a parabolic rate law

X 2 = 2kpt

(1)

where X is the depth of carburization after time t. The inward diffusion of carbon to meet carbide-forming elements leads to the precipitation of chromium-rich carbides within the steel. It is the relatively rapid inward diffusion of carbon through the depleted matrix of the precipitation zone which controls the carburization rate. Under these conditions [6] the rate constant in eqn. (1) is given by

eDcNc kp - - VNM

(2)

Here D c is the diffusion coefficient of carbon in the depleted matrix of the precipitation zone and Nc is the concentration of carbon at the alloy surface; the quantity v is the stoichiometric ratio for the composition of the carbide MC,, NM is the concentration of carbide-forming elements (essentially chromium) and e is a "labyrinth f a c t o r " to take account of the diffusional blocking effect of the precipitates. Given that Ncr is essentially fixed by other considerations, it is apparent that the only means of decreasing the corrosion rate are by modifications to the diffusion coefficient and solubility of carbon. It has long been known [7] that minority alloy additions of silicon can substantially reduce carburization rates. T h e r m o d y n a m i c interactions between dissolved silicon and carbon reduce N c [8] and will also reduce the © Elsevier Sequoia/Printed in The Netherlands

82

effective value o f D c [2, 9]. A more important effect is found under oxidizing conditions, when an external film of SiO2 may form. If sufficient silicon is present, the SiO2 film will be continuous and act as a barrier to carbon dissolution. Thus, N c is decreased and the corrosion rate reduced. These effects have been reported in a number of studies [2-5, 8-12] which make it clear that a relatively high silicon content is required. However, high silicon levels are deleterious to alloy mechanical properties and weldability, and optimizing the silicon content has proved difficult in practice. In an alternative approach, the silicon may be applied in the form of a coating. The coating ideally acts as a barrier to carbon migration, thus reducing the effective alloy surface carbon activity. The application of such a coating is well documented [13-17] and involves high temperature chemical vapour deposition from a silicon-containing gas, usually a silane. However, layer effectiveness in reducing carburization rates has not yet been conclusively demonstrated. Corrosion tests on silicon-enriched layers of nickel, produced via solid state diffusion [ 14], were unsuccessful because the layers were brittle and adhesion to the base material was bad. In contrast, it has been demonstrated [13] that silicon coatings produced by solid state diffusion on Hastalloy X were effective in reducing carburization rates. Consequently, the effectiveness of such coatings formed on commercially used alloys was the primary concern of this investigation. Because separation between silicon-rich coatings and their substrate metals has frequently been noted, and because industrial reactors must withstand temperature cycling, it is evident that a means of repairing or renewing the coatings is essential. For this reason, it has been suggested [18] that small

amounts of a silane could be added to the hydrocarbon process stream. Thus, continuous pyrolysis of the silane would maintain a silicon activity which could, perhaps, facilitate the desired coating renewal process. A secondary aim of this work was to test this possibility.

2. E X P E R I M E N T A L D E T A I L S

The alloys selected for study were HK40 (supplied b y Quality Castings) and PG2535Nb (supplied b y Pose-Marre). These materials represent p o o r and good carburization resistance respectively [2]; their compositions appear in Table I. Samples of approximate dimensions 10 mm × 5 mm × 1.5 mm, were cut and ground to a 120 grit finish and subsequently ultrasonically cleaned in acetone. Silicon-enriched surface layers were produced using chemical vapour deposition from flowing silicon-containing gas streams. In the first set of experiments, the alloys were exposed to 0.87 mol.% tetrachlorosilane (SiCl4) in the H2 feed at 1273 K for 4 h. In the second set, 0.8 mol.% dichlorodimethylsilane (SiCI2(CHa)2) in the H 2 feed was used at either 1323 or 1123 K for 3 h to simulate a dual siliconizing-carburizing atmosphere. In each case a total gas pressure of close to 100 kPa and linear flow rates of 30-80 mm s -I were used. After reaction, the samples were cooled slowly under flowing argon at 100 K h -1 for 4 h, then 200 K h -1 for 1.5 h and then slowly to ambient temperatures in order to minimize thermal shock. X-ray diffraction, semiquantitative electron microprobe and energy-dispersive X-ray analysis were employed to identify the coating compositions, while metallographic examination was employed to reveal surface morphology and evidence of carbon penetration. Etching with

TABLE 1 Alloy compositions

Alloy

HK40 PG2535Nb

A m o u n t (wt.%) o f the following elements Fe

Ni

C

Cr

Nb

Si

Mn

Mo

50.9 34.6

21.6 35.2

0.4 0.4

24.5 25

-1.6

1.3 2.2

0.6 1.0

0.7 --

83

hot alkaline KMnO 4 solution served to highlight precipitated carbides. Silicided samples were subsequently exposed for various time intervals to a carburizing atmosphere at 1273 K. The atmosphere was provided b y flowing H2-11vol.%C3H6 which produced a gas phase carbon activity of unity. The oxygen potential of this atmosphere was insufficient to oxidize silicon. Uncoated alloy samples were also reacted at the same time in order to provide standards. After carburization, samples were m o u n t e d in cold-setting resin, sectioned and polished to a 1 pm finish. The precipitated carbides were etched, and the carburization depths were measured using an optical microscope.

L.~

IOOpm

3. RESULTS AND DISCUSSION

3.1. Siliconizing in carbon-free atmospheres Silicon vapour deposition from H2-SiC14 atmospheres at 1273 K led to the formation on the alloy surfaces of continuous reaction product layers. Despite the slow rate at which the coated samples were cooled, extensive lateral cracking and spallation of the scales were evident, as shown in Fig. l(a). Lateral cracking appeared to occur at an interface between a porous inner layer and a more or less compact outer layer. The product layers were quite thick, as shown in Table 2. These thicknesses may be compared with the expected depths of inward silicon diffusion within the course of the siliconizing treatment. On the reasonable assumption that the surface silicon activity is constant, then the resulting concentration profile of silicon in the subsurface diffusion zone is given as a function of position x and time t b y Cs - - C

~-C0

x

- erf/2(Dt) ~}

(3)

Here C is the concentration within the diffusion zone, Co is the initial concentration in the alloy, Cs is the concentration at the alloy surface, D is the silicon diffusion coefficient and erf denotes the error function. The position at which C s -- C

Cs -- Co

-0.9

Fig. 1. Silicon-enriched layers produced on PG2535Nb at 1273 K in (a) a siliconizing atmosphere and (b) a siliconizing atmosphere and subsequently carburized at 1273 K (etched).

after 4 h at 1273 K is calculated, using D = 2 × 10 -11 cm 2 s -1 [14, 19], to be at a d e p t h of 12/2m. Since this is very much smaller than the observed product layer thicknesses, it is concluded that layer growth is sustained b y outward diffusion of other metals. Semiquantitative electron microprobe analysis of the layers confirmed that they are rich in silicon and that a discontinuous change in silicon concentration occurs at the layeralloy interface. The porous inner layer was enriched in chromium and nickel whereas the outer layer was depleted in chromium b u t even more strongly enriched in nickel. These features have been observed before [ 14, 20] and presumably reflect the combined effects o f the relative stabilities of the chromium and nickel silicides, the relative volatilities of the metal chlorides and the diffusion coefficients within the product layer of the different metals. Analysis by X-ray diffraction of the product layers revealed the presence of a complex mixture of nickel silicides (Ni3Si and Ni2Si), chromium silicide (CrSi2) and (Cr, Ni) silicides ((Cr, Ni)gSi, CrsNisSi2 and (Cr, Si)sNi2Si) plus quantities of iron silicides

84 TABLE 2 Reaction z o n e w i d t h s

R e a c t i o n z o n e w i d t h (pro) HK40 4 h a t 1 2 7 3 K in H 2 - S i C l 4 Outer scale Inner scale

PG2535Nb

60 30

30 40

3 h at 1 3 2 3 K in H2-SiCl2(CH3) 2

Scale Carburized z o n e 3 h at 1 3 2 3 K i n H 2 - C 3 H 6 ( a c = 1) Carburized z o n e

(FeSi and FeSi2). These coatings therefore differ from those reported earlier on nickelbased alloys [13, 14, 16, 20] in that silicide formation is more extensive and iron silicide formation occurs. Exposure of silicon-coated alloys to carburizing conditions at 1273 K led to some disruption of the silicide layer and internal carbide precipitation, as shown in Fig. l(b). The carburization kinetics of coated and uncoated alloys are shown in Fig. 2. It is seen that no significant protection was afforded by the layers. This finding is in contrast with that of Wahl et al. [13], w h o reported that a silicide coating on Hastalloy X led to a substantial increase in carburization resistance at 1223 K. Other reports [3] have described improved high temperature oxidation and hot corrosion resistance resulting from silicon coatings on nickel-based superalloys. The reason for this difference in performance is the failure, in the present case, to form coatings on the cast steels capable of retaining adherence after temperature cycling. The damaged coating cannot prevent gas access to the underlying metal and carburization proceeds unabated. On the nickel-based materials, tightly adherent coatings were reported. A factor contributing to this failure is the relatively high coefficient of thermal expansion of a cast 25wt.%Cr-35wt.%Ni steel, which is typically 19 × 1 0 -6 K -1 compared with values of around 13 × 10 -6 K -1 for nickel-based superalloys. It is concluded therefore that siliconrich coatings of the t y p e described here are unsuitable for use on cast austenitic steels subject to temperature change.

40 120 + 30

10 100 + 30

280 -+ 10

210 + 10

ox

800

600 -i 13

0

2O0

0

2

(a)

I

I

&

6

t V2 ( h I12)

600

/

~00

~ o 200

/o 0 (b)

f

0

2

d

4 tu/2 (hU/2)

f

6

Fig. 2. Carburization kinetics o f (a) H K 4 0 and (b) P G 2 5 3 5 N b at 1 2 7 3 K and gas phase a c = 1 : o, no coating; a, siliconizing at 1 2 7 3 K for 4 h; ~ siliconizing-carburizing at 1 3 2 3 K for 3 h; , regression line for cvxburization o f uncoated samples.

85

3.2. Siliconizing in carbonaceous atmospheres Exposure of the alloys to H2-SiC12(CHs) 2 atmospheres led to the formation at 1323 K of an external reaction product layer and a subsurface zone of internally precipitated carbides. An example is shown in Fig. 3 where the external layer is seen to be compact b u t separated from the substrate alloy. The thicknesses of these layers grown in 3 h are shown in Table 2. It is clear that these layers are substantially thinner than those formed in H 2SiC1 a atmospheres at 1273 K. Analysis b y X-ray diffraction of the external layers revealed the presence of the phases Cr7C3, Cr3C2, FeSi, CraNi5Si2, Ni2Si and austenite on both alloys. Semiquantitative analysis revealed a step change in silicon concentration at the layer-alloy interface, and a concentration within the layer which was roughly uniform. Distributed throughout the layer were localized regions of high chromium concentration and low nickel and silicon concentrations. These were taken to be chromium carbides as etching revealed some carbide existence within portions of the scale (see Fig. 3). The phase CrsC 2 was essentially pure, b u t the phase CrTC a contained significant levels of iron and will be referred to as MTCs. Small deposits of approximate composition C r 3 N i s S i 2 were found in a matrix rich in nickel, iron and silicon. The differing patterns of reaction products found using the two silanes arises, of course, from the presence of carbon in one of them and the probable difference in silicon activities achieved. Equilibration of H2-1vol.%SiCl4 produces [14] a silicon activity of approximately 10 -5

Fig. 3. Silicon-enriched layers produced on HK40 at 1323 K in a siliconizing-carburizing atmosphere (etched).

Pa at 1273 K, the principal silicon-containing species being SIC12. The silicon activity resulting from the use of SiC12(CHa) 2 is unknown, as is also the carbon activity. However, the strength of the Si--C bond is considerably lower than that of Si--C1 [21], and cracking of the former bond is therefore expected to be favoured. Limits on the possible values of ac may be established as follows. If the CH 4 formation reaction SIC12(CH3)2 + H2 = SIC12 + 2CH 4

(4)

is essentially complete, then sufficient CH a is formed that the carbon activity due to the equilibrium CH4 = C + 2H2

(5)

would be unity. However, no solid carbon was found on the specimen surfaces. The formation of CrsC 2 as well as CrTC s in the product layer at 1323 K shows that the carbon activity, referred to pure solid graphite, must be at least 0.27, as may be calculated from the thermodynamics of the reaction acrTCa + C = 5ZCr3C2

(6)

for pure carbides. It is concluded then that the gaseous carbon activity has a value in the range 0.27 ~ a c ~ 1. The depths of the internal carburization zones formed at 1323 K during the coating treatment are listed in Table 2. Obviously the silicide coatings do not exclude carbon. However, comparison with the carburization depths found in uncoated alloys exposed to unit carbon activity at this temperature (Table 2 also) shows that slower rates were in effect for the coated samples. As argued earlier, a reduction in rate indicates a lowering in the surface carbon concentration, and the effect can be investigated using eqn. (2). In the present case, where a surface layer is formed, its effect is taken into account by replacing kp by (kp (kp -- k) }112 [22] where k is the parabolic rate constant for alloy surface recession. The correction turns out to be negligible. Using the measured carburization depths to evaluate kp, the right-hand side of eqn. (2) can be evaluated for b o t h the uncoated and the silicon-coated cases. If it is noted that the quantities e, D c, v and N M ar e fixed, the effect of the coating on the value of N c can be gauged. In this w a y it is found that the ratio of Nc (coated) to N c (uncoated)

86

has the value 0.17 for HK40 and 0.22 for PG2535Nb. Evidently the coating has a considerable b u t not a dramatic effect. When the alloys HK40 and PG2535Nb are exposed to unit carbon activity at these temperatures, they form non-protective external scales of MTCs, and consequently the value o f ac at the alloy surface is close to unity [23]. In the present case of simultaneous siliciding and carburization, the scale contains both Cr3C 2 and M7C 3 in addition to metal silicides and is compact. It may therefore be concluded that the carbon activity at the scale base is controlled by the carbides in the scale rather than b y the gaseous carbon activity. Since the carbide M7C 3 is found within the alloy as well as within the scale, the value of ac at the scalealloy interface must lie within the existence range of this carbide. At 1323 K, pure CrTC3 exists over the range 0.005 ~ ac ~ 0.27. The presence of iron dissolved in the carbide destabilizes the compound and shifts the equilibrium carbon activities to somewhat higher values. As there is a lack of knowledge of the precise carbide composition at the scalealloy interface, it is n o t possible to calculate the value of ac, and hence Nc, which will be found here. All that can be concluded is that the observed degree of protection is consistent with the lowering in carbon activity expected to result from the formation of an external carbide-rich scale. It should be noted that the same result can be achieved b y selecting a higher nickel content alloy [23] without the need for silicon deposition. High nickel alloys form protective layers of Cr3C 2 o n top of the CrTCs. Silicon deposition from carburizing-siliconizing atmospheres at 1123 K led to the formation of a very thin surface layer containing the phases FeSi, SiC, Cr7C3 and C r s C 2 o n both alloys. This layer was found to be severely detached from the alloy and highly convoluted, as shown in Fig. 4. No internal carburization was detected after exposure for 3 h to H 2 - S i C 1 2 ( C H s ) 2 at 1123 K. The maximum carbon activity which could be produced under reaction conditions is calculated from the thermodynamics of the gas phase to be 0.56. The carburization depths expected [2] for these alloys after a 3 h exposure to ac = 0.56 at 1123 K are extremely small and it is unlikely that this experiment could reveal any protective effect due to the coating.

100 JJm Fig. 4. Silicon-enriched layers produced on H K 4 0 at 1123 K in a silieonizing-e~burizing atmosphere.

The results of subsequently carburizing alloys at 1273 K after first coating them in a siliconizing-carburizing atmosphere at 1323 K are shown in Fig. 2. The data are plotted according to the relationship ( X 2 __ X0 2)1/2 = (2kp t) 1/2

(7)

where X0 is the carburization depth resulting from the coating pretreatment. It is seen that no protection was achieved. The same result was found for alloys coated at 1173 K. The coatings failed to provide protection because they did not retain adhesion to the alloys during temperature cycling.

4. CONCLUSIONS

Silicide coatings can be formed on cast heat-resistant steels b y chemical vapour deposition from an H2-silane atmosphere. However, coating adherence to the underlying alloy does not withstand temperature cycling. Consequently, they afford no protection against carburization. Silicide coatings can also be produced in atmospheres which are simultaneously sillconizing and carburizing. In this case the coating consists of chromium-rich carbides as well as silicides. These coatings provide some degree of protection against carburization during exposure to the mixed gas atmosphere. The degree of protection is consistent with the control of alloy surface carbon activity b y equilibrium between carbides in the scale and the substrate alloy. However, these coatings also are unable to withstand temperature cycling. Their physical disruption permits gas access to the metal, and carburization pro-

87

ceeds. The coatings w o u l d therefore be of little use in a process environment. 12 13 R EFERENCES 1 J. K. Stanley, in S. A. Jansen and Z. A. Foroulis (ed~), High Temperature Gas-Metal Reactions in Mixed Environments, AIME, New York, 1973, pp. 143-154. 2 G. Smith, D. J. Young and D. L. Trimm, Oxid. Me£, 18 (1982) 229. 3 H. W. Grunling and R. Bauer, Thin Solid Films 95 (1982) 3. 4 H. J. Grabke and A. Schnaas, in B. W. Betteridge, R. Krefeld, H. Krockel, S. J. Lloyd, M. Van de Voorde and C. Virante (eds.), Proc. Int. Conf on Alloy 800, Petten, March 14-16, 1978, NorthHolland, Amsterdam, 1978, pp. 195-211. 5 A. Schnaas and H. J. Grabke, Oxid. Met., 12 (1978) 387. 6 C. Wagner, Z. Elektrochem., 63 (1959) 772. 7 E. N. Skinner, J. F. Mason and J. J. Moran, Corrosion, 16 (1960) 593t. 8 R . P . Smith, J. A m . Chem. Soc., 70 (1948) 2724. 9 S . K . Roy, H. J. Grabke and W. Wepner, Arch. Eisenh~ttenwes., 15 (1980) 91. 10 K. Ledjeff, A. Rahmel and M. Schorr, Werkst. Korros., 31 ( 1 9 8 0 ) 8 3 . 11 M. Saori and S. Ohta, High Temperature Corro-

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