New corrosion-resistant high temperature heat exchanger materials

Corrosion Science, 1968, Vol. 8, pp. 603 to 622. Pergamon Press. Printed in Great Britain

NEW CORROSION-RESISTANT HIGH TEMPERATURE HEAT EXCHANGER MATERIALS* D. R. HOLMES Central Electricity Research Laboratories, Leatherhead, Surrey, England Abstract--The paper examines the material properties required to provide an economic increase in the efficiency of electrical power generation by increasing steam temperatures from 560°C-660°C. An indication is given of the cost, strength and corrosion resistance required and it is concluded that none of the currently available materials will readily permit this advance at an economic cost. Possible directions for research and development to provide satisfactory new materials are outlined. Composite and coated alloys and non-metallic materials are discussed and it is concluded that the former offer the most immediate prospect for advances. R6sum6--I1 est question de matgriaux destin6s ~ augmenter 6conomiquement le rendement de la production d'61ectricit6 en 6levant la temp6rature de service de 100°C. On y renseigne le coot, la charge de rupture et la r6sistance/t la corrosion requis et on y r6partit les mat6riaux existant selon ces trois crit&es. On sugg&e des voies pour le recherche et le dgveloppement de mat6riaux nouveaux satisfaisants. On passe en revue les alliages composites et/l rev6tements ainsi que les mat&iaux non m6talliques et on conclut h de belles promesses d'avenir pour les premiers. Zusammenfassung--Es werden 0berlegungen tiber die Werkstoffe angestellt, die eine Steigerung des Wirkungsgrades von Energieumwandlungsanlagen dutch Steigerung der Arbeitstempera(ur um etwa 100°C zulassen. Die erforderliche Festigkeit und der notwendige Korrosionswiderstand sowie die zulfissigen Kosten ftir derartige Werkstoffe werden er6rtert und Kennwerte der bekannten Werkstoffe werden hiermit verglichen. Es werden Richtlinien ftir die Untersuchung und Entwicklung geeigneter Werkstoffe angegeben. Hierbei werden sowohl Verbundwerkstoffe als auch Plattierlegierungen und nichtmetallische Werkstoffe er6rtert, und es wird geschlossen, daf3 die Verbundwerkstoffe die besten Aussichten ftir einen Erfolg haben. INTRODUCTION THE PURPOSE of this paper is to consider materials to make possible a n economic increase i n the efficiency of electrical power generation. Over the last 10yr 1,2 the most advanced steam c o n d i t i o n s i n the power stations operated by the Central Electricity G e n e r a t i n g Board have changed only from 1500 p.s.i, at 563°C to 2300 p.s.i, at 566°C. F o r good e c o n o m i c reasons the latter are the steam c o n d i t i o n s chosen for most of the 2 0 , 0 0 0 M W of additional c o n v e n t i o n a l p l a n t to be installed by 1975, b u t because of superheater corrosion, the steam a n d metal temperatures i n oil-fired p l a n t 2 are restricted to 30°C lower t h a n c o r r e s p o n d i n g coal-fired boilers. Two boiler-turbine units with more advanced steam c o n d i t i o n s (3500 p.s.i, at 593°C) will be operated from 1968 o n w a r d s b u t because of the more expensive materials required for high-temperature strength a n d corrosion resistance, they are expected to have marginally higher generat i o n costs. The very small rises i n steam temperatures over the last few years are reflected i n the g e n e r a t i o n efficiency, which appears to be levelling off at a b o u t 36 per cent 3 (Fig. I). F o r t u n a t e l y the cost i n real terms per k W h of electricity generated (made up, as the A p p e n d i x shows, of both fuel a n d capital charges) has fallen as advantage has been t a k e n of the r e d u c t i o n i n capital costs afforded by the c o n s t r u c t i o n of very large generating units. *Manuscript received 1 February 1968. 603

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The experience in other countries is broadly similar although in Germany and the United States there are a few power stations with more advanced steam conditions 4 and slightly higher efficiencies. This difficulty in attaining economically efliciencies beyond about 36 per cent is partly responsible for the wide interest in novel supplementary and alternative methods of power generation, such as magnetohydrodynamic and thermionic generation and fuel cells. 5 In the present review an attempt is made to suggest possible ways of escaping from the present situation of stalemate. First, reasonable levels of the three primary parameters, cost, strength and corrosion resistance are established, and the deficiencies of currently available materials are reviewed. A fundamental discussion is given of the desirability of the various alloying elements and new alloys, composite and coated materials based on these considerations are explored. The main intention is to provide an advance in the steam conditions and eJiiciencies obtainable from fossil fuel fired power stations, but many of the arguments apply to a considerable extent to nuclear power stations employing an oxidizing gas as the heat exchange medium.

STRENGTH, COST AND EFFICIENCY In the past, high temperature alloys for superheater tubes have been developed by the steel makers primarily for their mechanical properties, in particular, for their high temperature creep resistance. Corrosion resistance, except for special applications such as valve steels or electrical resistance heaters, has been of secondary importance. A strong indication of the need for a fresh approach came from the work of Cain and Nelson, e Wyatt and Evans, 7 and Edwards, Jackson and Howes,s who showed that with current steam conditions and superheater metal temperatures, the corrosion rates are rising steeply with temperature with little prospect of improving steam conditions economically with present-day materials. For this reason recent

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research in the C.E.G.B. has concentrated on obtaining a fundamental understanding o f the factors responsible for corrosion resistance and on applying this understanding to the development of new materials for use at higher temperatures. It should be emphasized that the figures provided can only b e taken as guide lines because the permissible increases in cost obviously depend on the rates o f interest and capital depreciation assumed and on the load factor of the plant in which the new materials are to be used. Similarly, the exact strength levels required will be determined by the pipe diameters and wall thicknesses assumed; the latter, of course, are themselves determined by considerations o f heat transfer, thermal stresses and packing. As a first a p p r o a c h to finding the strength and cost levels we require, let us consider a modest target, an increase in steam temperature from 560°C to 660°C. This will imply metal surface temperatures of 700 ° to 720°C. If we assume that steam pressures will remain at 2300 p.s.i, the h o o p stress in a 2in. dia. tube of 0-2in. wall thickness will be about 10,000 p.s.i. If supercritical fluid at 3400 p.s.i, were to become the working fluid a reduced tube diameter would probably be used and the working stress would

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D.R. HOLMES

remain about the same. The ASME Boiler Design Code 9 lays down various criteria for materials to be used in boilers working at these temperatures and stresses. The most realistic criteria are: (1) The working stress should not produce a creep rate exceeding 10~/1000h at temperature. (2) A stress of 1.6 times the working stress should not produce rupture in 100,000h at temperature. Existing data relevant to these two criteria have recently been collated by Broom 1° and Baker. u These authors gave graphs showing the estimated temperature dependence of the stresses required to give a strain rate of 10-6/h and rupture in 100,000h, respectively. Unfortunately for the purpose of this review the creep data given did not cover a wide enough range of materials containing additions of specifically corrosionresistant elements, and therefore it has been necessary to use a less rigorous criterion for the level of mechanical strength required. We suggest (following a less applicable ASME code) 9 that the corrosion resistant materials we are seeking will require an U.T.S. of about 40,000 p.s.i, at temperature. This critical level is plotted in Fig. 2 which also gives the strength data for three classes of high temperature materials. The permissible increase in cost for new materials is a more speculative calculation ; it would be most unfortunate if premature concern with costs were to inhibit the development of promising new materials but finally the return on capital employed will determine whether or not a new material is introduced. The proposed increase in steam temperature would provide an optimized increase in overall cycle efficiency of a little over 2 per cent for a supercritical plant operating at 3400 p.s.i, and a similar increase for a sub-critical plant. With a load factor of 75 per cent and capital charges of 12½ per cent p.a. this increase in efficiency would permit an increase in capital costs (see Appendix) of about £4/kW of generation capacity. A suitable division might be to spend one quarter of this on trte boiler and one quarter on the turbine and put half to reduce specific capital charges, implying that £0.5m would be available for more corrosion-resistant materials in a 500MW boiler-turbine unit. In existing boilers of this size there are 100,000ft of austenitic stainless steel tubing and it has been suggested 12 that an additional 150,000ft of extra duty superheater tubing might be required for the more advanced steam conditions. This implies that the permissible price for the further 150,000ft of materials should not exceed £4/ft or about 20/per lb, assuming that no increased costs would be incurred for boiler and economizer tubing and steam piping. This is a very small margin when we recall that advanced austenitic materials of the 316 type are currently available at about 10-12/- per lb fabricated. We have therefore defined two broad objectives, a corrosion-resistant material costing not more than 20/- per lb with an U.T.S. of about 40,000 p.s.i, at 700-750°C. These must only be regarded as indications rather than firm figures because material costs vary widely with fabrication methods and production scale; some latitude in the U.T.S. required may be permitted by variations in tube diameter and wall thickness. D E F I C I E N C I E S OF CURRENTLY AVAILABLE MATERIALS Corrosion resistance--effects of Cr Boiler trialse, 7 and laboratory studies have shown that none of the currently

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FIO. 3. C o r r o s i o n rates as a function o f C r content. a. C o m m e r c i a l alloys in power stations (Marsdent4). b. Pure F e - C r alloys in pure O ( M o r t i m e r et al.~°).

used alloys have adequate corrosion resistance for use above 600°C. We deft ne adequate corrosion resistance in an arbitrary way as a metal wastage rate of 10-ein./h in severe boiler conditions. For a 200,000h life this gives a wastage of 0.2in. of wall thickness of an alloy tube, often rather more than could be tolerated in practice, so lower rates are higbly desirable. Several authors 1~,14 agree in showing that a practical method of obtaining adequate corrosion resistance in aggressive conditions is to use alloys with Cr contents above 20 per cent (see Fig. 3a). This is confirmed by the laboratory work of Greenert 15 who used crucible tests to simulate oil-fired conditions and more importantly by the electrochemical work of Cutler, Hart and Holland. 18 Cutler immersed alloys in a molten sulphate eutectic and used the c.d. at fixed time and potential as a measure of the overall corrosion rate. His curves of corrosion rate vs. Cr content bear a remarkable similarity to the oxidation rates in pure oxygen (Fig.

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Fio. 5. Parabolic rate constants for oxidation of Ni-Cr and Fe.--Cr alloys at 900°C (after Barrett et a1.17). minimum corrosion and oxidation rates we require a free Cr level of about 20 per cent. In complex steels, up to 10 per cent of the available Cr may be locked up in carbides and other compounds and in addition the Cr level at the oxidizing surface may be reduced24.25 due to the rate of its supply by diffusion from the matrix being inadequate to match its rate of incorporation in the oxide. For this reason it will be desirable to have Cr levels in commercial steels up to about 30-35 per cent by weight to ensure that a level of 20 per cent is maintained at the reaction interface.

Effects of Ni Many present-day corrosion-resistant alloys for heat exchangers also contain Ni to give added strength and workability by stabilizing the austenite phase. Because of the lower affinity of Ni than Fe or Or for oxygen, the Cr is eventually oxidized preferentially and a complex scale usually containing a layer of Cr2Oz with rather less than 1% of Ni in solution is formed, ze The mechanism of oxidation of N i - C r binary alloys is by no means clear 26 but if diffusion through such a layer is the rate controlling step then the solution of increasing but small amounts of Ni would give a reduction in diffusion and oxidation rates. This suggested mechanism is supported by the electrical conductivity measurements z7 which show an increase in electronic conductivity as the p-type NiO is added to the p-type Cr208, presumably associated with a decrease in cationic vacancy content. Broadly, the effect of adding small amounts of Ni to pure Cr might be expected to be similar to that of adding Fe (and indeed Co as CoO is also p-type). Figure 517 shows that it makes only a small difference to oxidation rates

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whether the Cr is diluted with Ni or Fe. However, there are obvious basic differences between the oxidation of the Fe-Cr, Ni-Cr and Co-Cr systems as Fe208 and Cr203 form a complete range of solid solutions 2s,z9 at all temperatures of interest, whereas only very small amounts of NiO (1-2%) and CoO dissolve in Cr203. At higher concentrations of Ni, the Ni-Cr spinels with increased diffusion rates are formed (Fig. 6) and this may explain the higher minimum oxidation rates of the Cr-Ni alloys compared with the Fe-Cr alloys (Fig. 5). The effects of Ni certainly require more study as preliminary results in power stations 23 and in the laboratory 3° show that small Ni additions reduce the corrosion and oxidation resistance of Fe-Cr alloys. Corti, Mortimer and Sharp 31 also reported that large Ni additions (35%) only slightly reduced the oxidation resistance of a 20Cr-Fe alloy in the range 650-850°C, and that improved resistance was obtained at higher temperatures. This alloy would already have a composition producing Cr2Oa films doped with the optimum amount of Fe to produce minimum transport rates; further doping with Ni (if this were possible on solubility grounds) might well have increased the ionic diffusion rate through the Cr203 film. There is, of course, the alternative possibility that no continuous layer of CrzO3 formed and that the higher oxidation rates were due to the formation of less protective Ni-Cr-Fe spinels. Similar considerations almost certainly apply to the presence of Mn as an alloying element; Caplan and his colleagues 32,33 and Francis and Whitlow ~ have already shown that the oxidation resistance of pure Fe-Cr alloys and commercial austenitic steels are impaired by the presence of Mn due to the formation

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of the less protective M n O . Cr203 spinel in place of the rhombohedral M203 oxide. The spinel films were blistered and much less coherent. To summarize, there seems to be no doubt that very few of the commonly used austenitic and ferritic alloys with Cr contents near the optimum level of 20--25 per cent (e.g. 18Cr-12Ni, 25Cr-20Ni, 20Cr-25Ni and 25Cr types) have minor element compositions designed to give maximum corrosion resistance.

Oxide .film integrity and adhesion Other points that have not previously received adequate consideration are the mechanical stability and integrity of the film. Parabolic diffusion controlled growth is, of course, an unstable situation. If cation migration is occurring then voids at the metal-oxide interface may lead to scale detachment. 3~ If anionic migration occurs then compressive stresses are likely to be built into the oxide (at least in the majority of practically important cases where the Pilling-Bedworth ratio > 1) and scale buckling and detachment may occur due to this cause. In both cases there are possibilities of some relief of the situation, the vacancies may diffuse away harmlessly into the metal or the compressive stresses may be relieved by plastic deformation. Another more remote possibility might be to choose alloy compositions giving oxides with minimum built-in stress; it may be possible to achieve this by having balanced anion and cation counter-currents. This desirable situation which occurs in the Potter and Mann 36 type of magnetite film growth and in some films grown in steam ~7 usually requires that permeable (but not necessarily unprotective) 2-layer films should be formed: There is also an increasing number of situations in which the alloy's resistance to oxidation must be maintained even under thermal cycling conditions. The main requirement here is a good match between the expansion coefficients of the oxide and the underlying metal; this point has recently been examined in some detail by Tedmon 21 in his attempt to explain the spalling characteristics of Fe-Cr alloys. He concluded that metal plasticity effects were more important than expansion coefficient mismatch, but had to make the very plausible assumption that the expansion coefficients and plasticities of the oxides would vary little with very slight changes in their composition. Comprehensive investigations of this kind are required if the spalling behaviour of the more complex alloys is to be understood. There will only be very limited scope for varying the expansion coefficient of the oxides formed in a way analogous to that used for producing successful ceramic enamel and crystallized glass coatings, by incorporation of suitable alloying elements in the metal. In fact this approach appears so difficult that it would almost certainly be more profitable to take the more direct approach of studying synthetic oxide coatings. In spite of the dearth of fundamental studies of oxide spalling remarkable empirical advances have been made by the addition of small amounts of rare earth metals? T M The mode of action of these is uncertain, although Felton and Francis and Whitlow 4° favour a keying mechanism by small Y203 or YCrO~ particles, formed by internal oxidation, penetrating the oxide-metal interface. Hage139 could find no metallographic evidence for the keying mechanism. Seybolt41 put forward two other hypotheses; his first suggestion that the mechanical properties of the Cr2Oa were radically altered by solution or dispersion of a few percent of Y~O3 was considered less likely. He tended to favour Stringer's suggestion 41 that the surface and sub-surface particles of Y20a might

612

D.R. HOLMES

be acting as vacancy sinks, thus reducing void formation at the metal-metal oxide interface 35 and improving scale adhesion. If this suggestion were the correct one it should be readily verifiable by the scanning electron microscope technique used by Howes; z5 in addition, cold working the alloy before oxidation might be expected to have similar or even more pronounced effects at low temperatures. The elucidation of the mechanism is of great scientific interest, but a possible bar to the development of these materials may be the high cost of the alloy elements. Yttrium, one of the most effective additives, costs £76/1b and therefore additions of even 1 per cent by weight can more than double the price of an alloy. But once the mechanism is understood it may be possible to achieve similar effects by the addition of cheaper elements such as A1 or Ce. At this stage it is constructive to summarize the deficiencies of present-day superheater alloys from the point of view of their corrosion resistance alone. (1) Many do not contain adequate coneentrations of Cr. In the matrix 20-30°,~Cr is required for maximum corrosion resistance in an Fe-based alloy, and possibly rather more with Ni additions. (2) Small Ni and Mn concentrations appear to reduce the protection provided by the Cr. (3) Few attempts have been made to incorporate constituents other than Cr to produce oxide films with minimum ionic transport properties. (4) Little attention has been paid to minimizing the built-in stresses in the oxide films or to choosing systems which give good film adherence under both isothermal and thermal cycling conditions. (5) No attempts have been made specifically to find alloys producing films with low dissolution rates in complex slags. PROSPECTS FOR NEW MATERIALS Alloy additions for corrosion resistance The previous section has pointed to some of the deficiencies of present-day superheater alloys; here we deal with possible ways of overcoming them and of developing new alloys. Figure 6 shows self-diffusion coefficients42 in a number of oxides as a function of 1/T. The data, while useful for presenting a broad indication of desirable types of oxide, require more detailed interpretation before they can be used to predict oxidation resistance of particular alloys and to determine compositions for new materials. Firstly, it is necessary to know the type and composition of the oxide forming and, in cases where complex multilayer oxides are formed, the oxide layer that is rate determining. Secondly, self-diffusion rates determined under conditions appropriate to the oxidizing conditions 43 must be used; for scales growing by anion transport, self-diffusion coefficients measured in oxides in equilibrium with the metal are required while for growth by cationic diffusion, coefficients determined on oxides in equilibrium with a partial pressure of oxygen equal to that used for the oxidation studies should be employed. In general this information will not be available but should be calculable if the conditions of the experiment were specified. Thirdly, an assessment of oxidation rates requires a knowledge of the maximum possible range of stoichiometries over which the oxide under consideration can exist. This last factor is often of great importance and explains some rather obvious anomalies. For

Materials for electrical power generation

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example, according to Fig. 6 the rate of oxidation of Fe to FezOa (by anion or cation diffusion might be expected to be slower than the rate of oxidation of Cr to Cr2Oa (by anion or cation diffusion), but in fact this is known to be untrue. The explanation for this is due to the possibility of Fe20 ~ showing much larger departures from stoichiometry. In spite of the complexities of interpretation of Fig. 6 it does show that diffusion rates are generally lower in MaO4 spinel oxides than in MO oxides and that the M203 oxides are even more resistant to ionic transport. Another point of particular interest is the low diffusion rates in the refractory oxides and rare earths, e.g. A1 and O in A12Oa, Mg in MgO and O in SiO2. "[he high temperature deformation of these materials will almost certainly be controlled by diffusion processes and thus their high refractoriness-under-load provides further confirmation that the diffusion rates in these high melting point oxides are exceptionally low. "I here is thus a strong indication that future corrosion-resistant alloys will have to contain the elements forming the most refractory oxides, in particular AI and Si as these are the cheapest. Already work by Hagel a4 and Mortimer, Sharp and Holmes 19 has shown that AI additions to an Fe-Cr alloy of about the most favourable composition (20Cr) provide a 3 or 4-fold decrease in oxidation rates and similar effects were obtained with small Si additions by Mortimer et al., ~ but not by Caplan and Cohen. 27 In support of this contention of the desirability of AI and Si as alloying elements the well-known oxidation resistance of the commercial alloy Kanthal* and the Sicromal* alloys (Fe 24Cr 1½AI 1½Si) can be cited. Besides the simple criterion given above of choosing alloying elements from those which give highly refractory oxides we can obtain further guidance from research on slag attack on ceramic materials. This has recently been given added impetus by the current interest in magnetohydrodynamic generation. Work in the British programme 4s has shown that AI2Oa, MgO and ZrO 2 have superior resistance to slags containing mixtures of alkali salts and vanadates, the aggressive products formed under oil and coal-firing conditions. There seems to be little prospect of incorporating either Mg or Zr into high temperature steels because of their very low solubilities 46 (Mg 0-1°o). Indeed, A1 and Si are the only alloying elements with adequate solubilities in Fe (up to 30at.% AI and 10at.°/oSi) 46 so again we are led to these two elements. In the case of elements forming very refractory oxides with strong affinities for oxygen, alloying levels of less than 1 per cent would be adequate on thermodynamic grounds, but in practice it is necessary to use higher levels to ensure that the alloying element can diffuse to the surface rapidly enough to prevent appreciable oxidation of the base element. The requirement for solubility should not be allowed to inhibit empirical investigation of all refractory oxide forming alloying elements, as very beneficial effects are sometimes obtained by mechanisms at present unexplained. (As an example less than 0 "'7°/-/oof Y is soluble in an Fe-250/oCr alloy at 1260cC ~7 but the addition of 1%Y forms Y(Fe, Cr) 9 and reduces the oxidation parabolic rate constant a° by a factor of 10.) It is not yet known whether the addition of Y lowers the permeability of the film to Cr 3+ ions by forming a YCrOz perovskite structure or by forming a dilute solid solution of Y~O3 in Cr203. 41 The further development of alloys with rare earth metal additions may well be limited by their high costs, but this limitation certainly does not apply to Al-containing *Commercial alloys from Aktiebolaget Kanthal, Sweden and Bohler Bros., Austria, respectively.

614

D.R. HOLM~

alloys. Already Chubb et al? 8 have shown that Al additions are compatible with Cr additions and produce alloys which can be both hot and cold worked. The oxides Cr203 and A1203 are also strictly compatible and form a complete range of solid solutions. 49 Measurements of high temperature electrical conductivity of AI203Cr203 solid solutions 5° suggest that doping of A12Oa by small amounts of Cr2Os may reduce the already very low diffusion rates in Al2Oa. "[his explanation could account for the oxidation rates below those predicted by extrapolation of the data for A1203 growing on A1 already noted. 19 Hensler and Henry's work 5° also suggests that A1 additions even at very low levels may have beneficial effects as doping of Cr2Oz with levels of A1208 up to 10 per cent appears to give reduced defect concentrations and transport rates in an analogous way to the effects of Fe203 on Cr203. 2"Other possible advantages of A1 additions are suggested by preliminary observations51.52 that improved spalling resistance may occur in oxides grown on these alloys possibly because of a more favourable balance of anionic and cationic fluxes. In fact there seem to be overwhelming reasons in favour of continued development of Fe-Cr-AI alloys, but first it is essential to prove that their superior corrosion resistance is maintained in complex liquid slag environments, is Strength

A major obstacle to the use of high Cr ferritic materials with and without A1 additions is that they do not have adequate high temperature strength above 600°C (Fig. 2). To obtain the required strength of 16-30 t.s.i, at 700°C it is necessary to form an austenitic matrix by adding Ni, an element which in small amounts impairs high temperature oxidation and corrosion resistance. It seems inescapable that f.c.c. matrices, being closer packed and having lower diffusion rates than b.c.c., are basically stronger. 53 For example, on passing through the ct-y transition in Fe, the creep rates 54 are reduced by a factor of 200 and the self-diffusion rates 55 by a factor of 350. Sherby 53 presents further data showing the superiority of the close-packed f.c.c, and h.c.p. structures in conferring high temperature strength and creep resistance. His evidence suggests the rather obvious course of seeking austenite formers other than Ni or Mn, preferably elements giving refractory oxides, as these two elements are known to impair oxidation and corrosion resistance. But no clear choices present themselves as the only strong austenite formers are Ni, Mn, C and N, and the oxides of the last two elements are by no means refractory. Examination of Fig. 2 shows that a relatively small increase in high temperature strength (by a factor of 2 or 3, equivalent to about 100°C) would make the F e - C r - A l alloys attractive, so it is essential to continue investigation of methods of strengthening them. Simple dispersion hardening methods may not be adequate because at these high temperatures precipitate barriers to dislocation movement can be readily by-passed by climb. A more profitable approach to the development of strong corrosion resistant materials may be on the lines of the sintered aluminium powder materials, where the possible existence of a continuous refractory oxide skeleton may explain the very high activation energy of creep. ~6,57 In principle it should be possible to use the same strengthening mechanism for the Fe-Cr-A1 alloys which also form predominantly a-A1203 oxide films, x9,4° The powder metallurgical methods used for fabrication of S.A.P. components would be both unsuitable and too costly for the manufacture of

Materials for electrical power generation

615

large complex superheater assemblies, but it may be possible to develop extrusion techniques for tube manufacture. Even so the very low ductilities of these materials will probably introduce great difficulties in fabrication and operation. Possible solutions If no satisfactory way of strengthening the Fe-Cr-AI alloys can be found it may still be possible to use them by an engineering design approach. Two methods of doing this suggest themselves: (1) By a redesign of the superheater tubes with narrower bores and thicker walls to give them adequate strength at temperature. This would incur penalties in both capital cost and heat transfer. (2) By the use of composite tubes with the corrosion-resistant alloys bonded to or coated on the outside of a stronger material. This approach is considered in more detail later. If the Fe-Cr-AI alloys cannot be brought to a successful realization then the much more expensive Ni-Cr alloys seem to be the only system capable of providing adequate corrosion resistance and strength at temperatures up to 750°C. Laboratory work showslS, 58 that the compositions 60Cr--40Ni and 50Cr-50Ni have the best performance in simulated corrosion environments containing vanadates, alkali sulphates and chlorides. Wickert s9 examined the behaviour of the above alloys and 35Cr-65Ni in natural and synthetic oil ashes and concluded that the latter alloy was the best in SO2-rich atmospheres. These results are broadly in accordance with expectations from the oxidation rates in air given by Barrett et al. 17 who suggest that the corrosion resistance is marginally superior to the best of the Fe-Cr-AI alloys, probably because the good resistance of Ni and NiO to vanadium compounds is combined with the high resistance of Cr to attack by alkali sulphates and S. Few attempts have been made to optimize their corrosion resistance by Si or A1 additions and, indeed, it is not certain whether AI will have its usual beneficial effects on this alloy (although it certainly does on a mixed iron-nickel base material of composition 20Cr-35Ni-45Fe). 31 The main obstacle to their use for superheaters in power stations is their high cost of £3-£4/lb in fabricated form. To summarize, alloys for use at 700-750°C must possess the correct levels of three attributes: A---cost, B---corrosion resistance and C--strength. Class I alloys (using the classification of Fig. 2) comprising the most advanced austenitic steels are adequate in A and C, class lI, the best ferritic alloys are adequate in A and B while class III, the Ni-Cr alloys are adequate in B and C. It is suggested that this unsatisfactory situation is the reason why steam conditions and generation efficiency have advanced so little over the past ten years. The most attractive prospects for overcoming this barrier seem to be in attempting to strengthen the class II alloys, in optimizing the corrosion resistance of class I alloys or by combining the properties of the two types of alloy in composite bodies.

ALTERNATIVE MATERIALS AND METHODS OF PROTECTION Composite alloy tubes and metallic coatings It was seen above that none of the three classes of materials possess all three desirable attributes. This has led to an examination by Asbury e° of the possibility of

616

D.R. HOLMES

combining the corrosion resistance of the ferritic materials with the high temperature strength of the austenitic steels. Asbury considered two factors of possible significance: (1) Interdiffusion between the two alloys. (2) Fatigue effects due to difference in expansion coefficients of the two materials. He concluded that (1) was of negligible importance but that (2) could be significant with some pairs of alloys ; however, he was able to specify five possible coating alloys on two strong base alloys (Table 1) worthy of further investigation. Asbury favoured co-ext usion as the most economic method of production and concluded that the cost increases should be within the permissible level specified earlier. TABLE 1.

CORROSION-RESISTANT/STRONG ALLOY PAIRS

Base material

Coating material

16Cr-12Ni-1Mn-1Mo

25Cr-20Ni 24Cr-12Ni 25Cr-20Ni 24Cr-12Ni 24Cr-2Co-6AI

15Cr-I 5Ni-2Mn-I

Mo

Further examination of the problem will probably reveal many other compatible corrosion-resistant/strong alloy pairs, but before they can be introduced into general boiler service other technological factors such as welding will require evaluation. To summarize, the prospects for wrought composite alloy tubes look encouraging and development work and laboratory studies should be intensified; already the commercial availability of 20Cr-32Ni clad with the highly corrosion-resistant 50Cr-50Ni has been announced. An obvious parallel development to that of wrought composite tubes is the use of sprayed, dipped, electroplated or diffused metal coatings. The most suitable materials for use as diffusion coatings or electroplates are Cr, Ni-Cr alloys and A1. Methods for applying coatings of these metals have recently been reviewed by Huminik 61 and the lower porosities of plasma sprayed coatings compared with other types have been emphasized. All sprayed coatings have some porosity (and so for that matter do most oxide films)but an improvement in their protectiveness can be obtained by multicoating procedures and by infiltrating second coats of glassy material. However, the complex procedures involved are more adaptable to the treatment of small finished components rather than to the mass production of tubes which have later to be bent and welded into large heat exchangers. As a result few trials of diffusion coatings in conventional power stations have been reported and even fewer details of application methods have been given. Harlow 4 observed that 17Cr-14Ni-Cu-Mo alloy tubes protected by electroplated and plasma sprayed Cr coatings were in good condition after 5 weeks' service at an estimated metal temperature of 700°C in coal combustion products. Foster and Tort 13 reported failures of chromized low alloy steel tubes after 1000h and 2500h operation at 630°C but apparently the failures were caused by poor mechanical properties rather than by corrosion, as an uncoated tube survived. This experience emphasizes one of the difficulties of metallic coatings: because of diffusion effects the coated alloy often cannot be given its correct heat treatment. Alternatively,

Materials for electrical power generation

617

if the correct heat treatment is given the protectiveness of the coating or the mechanical properties of the base material may be impaired through the formation of brittle intermetallic phases at the coating-alloy interface. These possibilities suggest that the most attractive prospects for composite tubes lie in clad tubes or in non-metallic coatings.

Non-metallic coatings The use of non-metallic coatings on alloy tubes appears to be a logical development; in principle they would allow the composition of the alloy tube to be optimized for strength and cheapness while corrosion resistance would be imparted by a ceramic coating of carefully designed composition. The coating could thus have the correct chemical composition and physical texture from the start of service instead of relying on diffusion from the underlying alloy and oxidation to produce protective oxide layers of unspecified and often variable composition, containing more or less built-in mechanical stresses and material discontinuities. This separation of duties is commonplace in many other fields of technology ; for example, in electrical transformer steel, a highly sophisticated alloy has been evolved with the correct directional electrical, magnetic and fabrication properties, but electrical insulation and prestressing is obtained by the high temperature application of a magnesium phosphate coating. Similarly with structural steel it is customary to separate the strength and corrosion resistance by the use of paints or other protective coatings. The first question to be decided is the choice of coating material; basic-guidance can be obtained from Fig. 6 and from a consideration of slag attack on refractory materials, 45 although in many cases experimental data exist only for higher temperatures. In practically all cases non-oxide refractory materials owe their corrosion resistance to a thin protective layer of oxide and for this reason it is logical to concentrate on oxide coatings, although the superior mechanical and thermal shockresistance properties of carbides, borides and nitrides may be attractive at very high temperatures. A recent survey of commercial refractories 6z confirms that oxide materials are the most useful for the range of temperatures encountered in boiler combustion products. Most of the coatings under commercial development consist of simple or complex oxides. It is also necessary to consider whether the combustion product environment will be predominantly acidic or alkaline; this will depend to some extent on the nature of the fuel and the combustion conditions but in the majority of cases (excess combustion air and sulphur bearing fuels) acidic deposits will be formed and therefore amphoteric or acidic oxide coatings will be more suitable. These considerations, coupled with diffusion data (Fig. 6), suggest that desirable coatings might consist of SiO2, AIzOa, ZrO2, BeO, MgO, Cr2Oa, Y203 and rare earth oxides, or compounds of these oxides, but many of these can be eliminated on the grounds of cost, toxicity or undesirable crystal transitions. With oxide coatings the problem of incorporating the parent elements into the alloy matrix is eliminated and instead the problem of matching the expansion coefficients of oxide and base metal is substituted. Almost the only oxides used extensively are A1203 and ZrO2, which can be applied conveniently by flame or plasma spraying.rZ These have expansion coefficients of 9.0 and 12 × 10-6, respectively against the 13 and 18 x 10-6 required to match ferritic and austenitic steels, respectively. With this mismatch the resistance

618

D.R. HOLMES

of oxide-coated materials to thermal cycling would be very poor and this may account for their poor performance in practice. Graded layers or combined metal-oxide layers may show useful advantages. The other major drawback of ceramic coatings is their porosity, up to 20 per cent for flame-sprayed coatings and rather less, about 10 per cent for plasma-sprayed material. The porosity provides a useful improvement in thermal shock resistance, but at the expense of reduced protection. Attempts have been made to reduce the porosity by infiltration with a glassy material as a second coat and to enhance the corrosion resistance and keying of the substrate metal by prior spraying 64.65 with a Ni-Cr alloy. The C.E.G.B. 6e and the British Coal Utilization Research Association e7 have carried out laboratory work and rig trials on alloys coated with proprietary enamel frits. Good results were obtained in short-term rig trials of enamels on an 18Cr-12Ni-Nb steel but these were not maintained in longer term laboratory crucible tests in Na2SO4-NaC1 mixtures. Extensive work on the development of new enamels and on their efficiency in protecting high temperature alloys from oxidation has also been carried out at the National Bureau of Standards. 68,69 This work showed that Cr2Oa and CeO2 were the most useful additions for reducing oxidation rates of 25Cr-20Ni steel at 1000°C but in most cases the reductions were marginal--by only a factor of 2 compared with uncoated material. An unexpected bonus was the finding 7° that tensile creep rates at 1000°C of 25Cr-20Ni steels and 80Ni-20Cr alloys could be reduced by 50 per cent by enamel coatings. Similar beneficial effects on mechanical properties have since been reported by other authors. 71 The new boiler trials carried out 7~,73 have tended to confirm the results of the laboratory and small-scale work. In no cases were successful results obtained and the performance of ceramic-coated tubes was inferior to that of Cr-coated tubes. 4 Finally, it is necessary to point out that as the coatings are not in thermodynamic equilibrium with the substrate alloy there is always the possibility of interdiffusion and even of enhanced attack. An example of this is provided by Richmond's observation e9 that a coating frit consisting of a complex oxide mixture (SiO~., TiO~, BaO, CaO, ZnO, BeO, P205, Cr203 and clay) actually gave accelerated attack on a 25Cr-20Ni steel compared with uncoated material. Huminik's el summary of reactions observed between metals and oxide coatings unfortunately deals only with refractory metals and temperatures beyond 1500°C. To summarize the prospects for ceramic-coated materials, it is clear that at the present stage of coating development we are far from a practical realization of their very attractive possibilities. At present no coating capable of providing protection for 10,000h in combustion products at 700°C exists. But this is a field in which new developments are occurring continually (for example the crystallized glass coatings and the fully dense transparent oxides of alumina, magnesia and yttria) and it may eventually be possible to apply similar methods to obtain fully dense impermeable coatings. Finally, the possibility of introducing additives into the fuel to deposit protective coatings on the heat exchanger surfaces must be mentioned. ALTERNATIVE MATERIALS At present no replacement materials for metals in heat exchangers employing

Materials for electrical power generation

619

5

'E

:> v-

aoo

§ d l.--

.5

.

~

~

..,..........~- --" "" "" ~

~.~.."'-

~,,.,.~.

10'

AUSTENITIC(MEANVALUE)

FERRITIC(MEANVALUE)

I

I

I

I

r

I

200

400

600

B00

I000

1200

TEMPERATURE°C.

F,G. 7. Comparison of thermal conductivities of high temperature austenitic and ferritic steels with those of selected ceramics. steam as the heat transfer medium can be foreseen. Ceramic tubular heat exchangers have been considered 45 for high temperature, high pressure air heaters for magnetohydrodynamic generation but the complexities of their construction have led to their abandonment; only crude simple devices with pressure differences of less than onethird of an atmosphere and none of the complexities of a modern power station superheater tube are in use. For working fluids operating beyond steam temperatures nonmetallic materials will almost certainly be used. Ceramic and cermet nuclear fuels are already in use and SiC fuel element cans are under active development; 74 SiC, BN, BeO and C seem the most immediately promising materials for use in tubular form as methods have already been developed for tube fabrication and joining. In addition their thermal conductivities are well above those of other ceramics and exceed the values for austenitic and ferritic steels (Fig. 7). The oxidation resistance of SiC and BN up to 1000°C is good but less is known about their performance in more corrosive environments. 75 CONCLUSIONS (1) The materials properties and costs needed to provide an economic increase in steam temperatures from 560-660°C have been examined. With a reasonable division of costs between boiler, turbine and capital savings requirements for new corrosion resistant materials are:

620

D.R. HOLMES

Cost--less than 20/- per lb fabricated. Strength--U.T.S. greater than 40,000 p.s.i, at 700°C or strain less than 1 per cent in 105h at 700°C. It appears that no existing materials meet these requirements. (2) The apparent excellent corrosion resistance of the Fe-Cr ferritic alloys with AI and other additions should be confirmed in longer term trials. At present their high temperature mechanical properties are inadequate so methods of employing them by increasing their strength, by design changes or in composite materials must be sought. (3) As an alternative approach to that outlined in (2) the corrosion resistance of existing high strength austenitic materials should be optimized by changes in the minor element composition. The effect of Ni on corrosion resistance is by no means clear and requires further research. (4) Ceramic coatings for high strength materials, although logically very attractive, are far from long-term practical application. Metallic coatings provide more immediate practical prospects. (5) There appear to be no immediate prospects of non-metallic materials replacing metals in heat exchangers using steam as the working fluid, but their use in novel generation methods using other working fluids will increase.

Acknowledgements--The author wishes to thank his colleagues at C.E.R.L., especially Dr. Mortimer and Mr. Asbury, for much useful discussion and permission to use unpublished results. The paper is published by permission of the Central Electricity Generating Board. APPENDIX The generation cost per kWh of electricity (U) in pence is given by 240Pr U = C---t 11 365 L × 24

(I)

where C is the cost per kWh of chemical or nuclear energy, r I is the overall generation efficiency, L is the load factor defined as kWh actually generated per year/maximum number of k w h that could be generated, P is the cost of the plant in £/kW of generation capacity, r is the fractional annual return on capital required. Differentiating to find the relation between increments in capital costs and efficiencies we obtain

CA N re

-

-

--

-

APr 36"5L -

(2)

or AP = 36.5C AqL/rrl 2. With Arl = 2%, C = 0.141d./kWh, r I = 40%, L = 75% and r = 12½% we obtain AP = £3.86/kW ~ £4/kW, It is interesting to note from (1) that one-third of the generation cost of a kWh electricity arises from capital costs and two-thirds from fuel costs Equation (2) shows clearly that the current high rates of capital return and low load factors (now approaching 50 per cent) weigh heavily against the introduction of expensive new materials.

Materials for electrical power generation

621

REFERENCES Central Electricity Authority Annual Report and Accounts 1956-7. H.M.S.O., London. Central Electricity Generating Board Annual Report and Accounts 1966-67. H.M.S.O., London. Engineering and Boiler House Review, 313 (Nov. 1967). J. H. HARLOW, Combustion 33, 37 (1962). K. H. SPRING (Ed.), Direct Generation o f Electricity. Academic Press, London (1965). C. CAIN and W. NELSON, Trans. Am. Soc. mech. Engrs. 83A, 468 (1961). L. M. WYATT and G. J. EVANS, The Mechanism o f Corrosion by Faellmpurities, p. 612 (Ed. H. R. JOHNSON and D. J. LI'VrLER). Butterworths, London (1963). 8. A. M. EDWARDS, P. J. JACKSON and S. HOWLS, J. htst. Fuel 35, 1 (1962). 9. ASME Boiler and Pressure Vessel Code, Section 1, Power Boilers, 152 (1962). 10. T. BROOM, Contemp. Phys. 8, 213 (1967). 11. D. W. C. BAKER, Advances in Materials, pp. 45-56. Pergamon Press, Oxford (1966). 12. M. G. GEMMILL and P. S. K. EAVES, private communication (1966). 13. G. G. FOSTER and L. H. TOFT, J. Inst. Fuel 35, 28 (1962). 14. K. A. MARSDEN, The Mechanism of Corrosion by Fael hnpurities, p. 632 (Ed. H. R. JOHNSON and D. J. LrrTLER). Butterworths, London (1963). 15. W. J. GREENERT, Corrosion 18, 57t, 9It and 95t (1962). 16. A. J. B. CURLER, A. B. HART and N. H. HOLLAND, Paper 67-WA,/CB-4 at the Winter Annual Meeting of ASME, Pittsburgh (1967). 17. C. A. BARRE'rr, E. B. EVANS and W. M. BALDWIN, Armed Services Technical Information Agency Document No. AD 80836 (1955). 18. G. C. WOOD and D. P. WHI-rrLE, Corros. Sei. 4, 263, 293 (1964). 19. D. MORTIMER, W. B. h . SHARP and D. R. HOLMES, Paper at the 3rd International Congress on Metallic Corrosion, Moscow (1966). 20. D. MORTXMERand W. B. A. SHARP, Dr. Corros. J. 3, 61 (1968). 21. C. TEDMON, J. electrochem. Soc. 114, 788 (1967). 22. P. K. FOOTNER, D. R. HOLMES and D. MORTIMER, Nature, Lotut. 216, 54 (1967). 23. N. H. HOLLAND and J. G. PARKER, unpublished results (1967). 24. G. C. WOOD, T. HODGKIESS and D. P. WHITTLE,Corr. Sci. 6, 129 (1966). 25. D. P. WHITTLE, D. J. EVANS, D. B. SCULLY and G. C. WOOD, Acta. Met. 15, 1747 (1967). 26. G. C. WOOD and T. HODGKIESS, J. Electrochem. Soc. 113, 319 (1966). 27. W. A. FISCHER and G. LORENZ, Arch. Eissenhutten ICes. 28, 497 (1957). 28. R. K. DI CERBO and A. V. SEYBOLT,J. Am. Ceram. Soc. 42, 430 (1959). 29. A. MUAN and S. SOMIYA,J. Am. Ceram. Soc. 43, 207 (1960). 30. D. MORTIMER and W. B. A. SHARP, unpublished results (1967). 31. C. COATI, D. MORTIMER and W. B. A. SHARP, 1967 A.E.R.E. Meeting on Structural Processes in High Temp. Materials. Harwell (June 1967). 32. D. CAPLAN and M. COHEN, Nature, Lond. 205, 690 (1965). 33. D. CAPLAN,P. E. BEAUBIEN and M. COHEN, Tratm. Am. Inst. Min. Engrs 233, 766 (1965). 34. J. M. FRANCIS and W. H. WHITLOW, J. Iron Steel Inst. 203, 468 (1965). 35. V. R. HOWLS, Nature, Lond. 216, 362 (1967). 36. E. C. Po'vrER and G. M. W. MANN, Proc. 1st International Congress on Metallic Corrosion, London, p. 49 (1961). 37. P. W. TEARE, P. L. HARRISON and D. R. HOLMES, VGB-Speisewassertagung, p. 2 (1965). 38. E. J. FELTON, J. electrochem. Soc. 108, 490 (1961). 39. W. C. HAOLE, Trans. Am. Soc. Metals 56, 583 (1963). 40. J. M. FRANCIS and W. H. WHITLOW, Corros. Sci. 5, 701 (1965). 41. A. V. SEYBOLT, CottoN. Sci. 6, 263 (1966). 42. W. D. KINGERY, Introduction to Ceramics, p. 232. Wiley, London (1963). 43. K. HAUFFE and B. ILSCHNER, Z. Elektrochem. 58, 467 (1954). 44. W. C. HAGEL, Corrosion 21, 316 (1965). 45. G. HORN, A. W. SHARP and W. R. HRYNISZAK, Phil. Trans. Roy. Soc. 261A, 514 (1967). 46. M. HANSEN, Constitution of Binary Alloys. McGraw-Hill, New York (1958). 47. S. G. EPSTEIN, A. A. BAUER and R. F. DICKENSON, Battelle Memorial Institute, B.M.I. 1376 (1959). 48. W. CHUBB, S. ALEANT,A. A. BAUER, J. JABLINOWSKI,F. R. SHOBER and R. F. DICKENSON, Battelle Memorial Institute Report No. 1298 (1958). 49. E. N. BUNTING, J. Res. Nat. Bur. Stand. 6, 948 (1931). 50. J. R. HENSLER and E. C. HENRY, J. Am. Ceram. Soc. 36, 76 (1953). 51. D. MORTIMER and W. B. A. SHARP, unpublished results (1966). 1. 2. 3. 4. 5. 6. 7.

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D . R . HOLMES

52. V. R. HowEs, Paper at the C.A.P.A. Meeting at Leatherhead, England, on MechanicalProperties and Stresses in Oxide Films (1967). 53. O. D. SHERBY, Acta. Met. 10, 135 (1962). 54. O. D. SHERBYand J. L. LYTTON, Trans. Am. Inst. Min. Engrs 206, 928 (1956). 55. C. E. BmCHENALL and R. F. MEHL, Trans. Am. Inst. Min. Engrs 188, 144 (1950). 56. G. S. ANSELL and J. WEERTMAN, Trans. Am. Inst. Min. Engrs 215, 838 (1959). 57. C. L. MEYERS and O. D. SHERBY, J. Inst. Metal 90, 380 (1961-2). 58. H. LEWIS, J. lnst. o f Fuel39, 8 (1966). 59. K. WICKERT, Ni Berichte 24, 9 (1966). 60. F. E. ASBURY, private communication (1967). 61. J. HUMXFaK(Ed.) High Temperature Inorganic Coatings. Reinhold Publishing Corp., New York (1963). 62. J. F. BURST and J. A. SPIECKERMAN, Chem. Eng. 85 (1967). 63. J. HUMINIK (Ed.) High Temperature Inorganic Coatings, p. 77. 64. E. D. TEAGUE, Foundry Trade J. 102, 2112 (1957). 65. E. D. TEAGUE, Information from Norton Abrasives Ltd. (1963). 66. D. J. A. DEAR, unpublished results (1964). 67. R. J. BISHOP and C. J. HOLLAND, B.C.U.R.A., Information Circular No. 280 (1964). 68. W. N. HARRISON and D. G. MOORE, Nat. Bur. Stand. Tech. News Bull. 35, 89 (1951). 69. J. C. RICHMOND, H. G. LEFORT, C. N. WILLIAMS and W. N. HARRISON, J. Am. Ceram. Soc. 38, 72 (1955). 70. J. R. CUTHILL,J. C. RICHMOND and N. J. TIGHE, Am. Ceram. Soc. Bull. 38, 4 (1959). 71. B.P., 981,794. 72. M. SCHULZ, Energie 11, 2 (1959). 73. J. JONAKIN and D. H. BARNHARTin ref. 6, p. 648. 74. A. BOLTAXand J. R. HARDWERK (Eds.), Proceedings o f the Conference on Nuclear Applications o f Nonfissionable Ceramics, American Nuclear Society (1966). 75. E. J. D. SMITH, Thesis, Battersea College of Advanced Technology (1966).