Material formulation for galvanizing equipment submerged in molten aluminum and aluminum/zinc melts

An article of equipment intended to be submerged in molten zinc, molten aluminum and mixtures thereof, said article containing an alloy material comprised of: Percent Carbon 1.6-2.6 Chromium 15.0-30.0 Tungsten 10.0-30.0 Molybdenum 2.0-8.0 Iron 10.0-50.0 and including, vanadium, niobium, cobalt, boron, and manganese and being substantially free of silicon.

BACKGROUND OF THE INVENTION 
 Galvanized steel utilized in the automotive, construction and appliance 
 industries is formed in very thin strips (0.015 to 0.060 inch thick), 
 which is passed through a molten bath of either zinc (galvanizing), 
 aluminum (aluminizing), or aluminum/zinc (galvanneal, galfan, galvalume, 
 etc.), in which the levels of aluminum vary from a fraction of a percent 
 to as much as 100 percent. 
 The "hot dip" metalizing coating process requires equipment that runs 
 submerged in the molten metal. The molten metal temperature usually varies
 from as low as 820.degree. F. to as high as 1300.degree. F. 
 A heated metal pot contains a bath of molten zinc/aluminum. A continuous 
 moving strip of low carbon steel is introduced into the bath from a 
 furnace in the conventional manner. The strip passes around a sink roll 
 and tensor rolls while submerged in the bath, so the surface of the strip 
 picks up a zinc/aluminum coating. The strip is delivered to the bath 
 through a conventional tubular snout. The interior of the snout housing 
 contains an inert gas such as nitrogen or a mix of nitrogen and hydrogen. 
 This procedure, as is well known to those skilled in the art, is useful in
 preventing the steel strip from oxidizing. 
 Because of the extremely large dimensions of the equipment and in spite of 
 efforts to prevent all possible air leaks into the furnace, small leaks do
 occur, generating ferrous oxides (Fe.sub.2 O.sub.3 FeO, etc.) When the 
 steel strip enters the bath, a chemical process occurs in which the melt 
 in the bath reacts with the iron in the steel strip (inducing the coating)
 but also reacts with the oxides to form dross that contains ZnFe, ZnAlFe, 
 ZnFe+Al.sub.2 O.sub.3, etc. The free iron settles to the bottom of the 
 molten metal pot. Because of the slightly or nearly identical density to 
 the molten metal, the oxides (Al.sub.2 O.sub.3, ZnO) and the 
 intermetallics formed (ZnFe, ZnAlFe, etc.) remain in suspension or float 
 to the surface in the form of dross. The dross increases its concentration
 by being nearly entrapped in the zone comprised by the snout, the strip, 
 the sink roll and the tensor rolls, where it gradually forms deposits on 
 top of the sink roll and the strip being processed. 
 Outline of Requirements for Galvanizing Equipment 
 Standard rolls and equipment used in the hot dip metalizing process, when 
 the alloy melt is zinc or zinc/aluminum with aluminum concentrations of 
 less than 60%, are 316-L stainless steel. The rolls and bearings, in 
 particular, require continuous maintenance of their surfaces. The rolls 
 are removed weekly from the pot and their surfaces machined to remove 
 accumulated dross, to smooth the roll surfaces as well as to return them 
 to a round and straight condition. The main reason for this continuous 
 maintenance is because 316-L stainless steel is not a material formulated 
 specifically for this application and, consequently, it lacks the 
 properties to meet the operational needs. 
 In order of importance, although all requirement must be met to a minimum 
 degree, the properties required for a proper roll material that meets the 
 operational needs are as follows: 
 1. Very low solubility in molten zinc or zinc/aluminum alloys. In other 
 words, 0=S&lt;4.times.10.sup.-5 in/hr Where S=the amount of roll radial loss 
 due to molten metal dissolution. 
 2. Low adhesion (non-wettable) to zinc/iron and zinc/iron/aluminum dross. 
 Wetting plays the main role in the bonding of solid-liquid state metals. 
 3. High surface hardness (R.sub.c larger than 40). Abrasive wear 
 contributes nearly half of the loss of roll life in metalizing 
 applications. 
 4. Dimensional stability at operating temperatures up to 1300.degree. F., 
 for straightness and roundness. This property is necessary because of the 
 difficulties encountered when the lines operate at over 100 RPM, 
 generating excessive vibration and damage to the holding equipment. 
 5. Thermal shock resistance. The roll should be capable of withstanding a 
 thermal shock of no less than 500.degree. F. when going from air to the 
 molten metal, and 1300.degree. F. when going from the molten metal to air.
 6. Good impact and notch resistance strength. This is importance due to the
 severity of the application. 
 7. Centrifugally castable and machinable by standard procedures in order to
 provide simple and available maintenance. 
 8. Economic viability. 
 The following expands, in corresponding order, each of the material 
 properties required to meet the listed operational needs. 
 Evaluation of Specification Requirements 
 In order to obtain a material formulation that is capable of having a 
 dissolution rate of 
EQU O=S&lt;4.times.10.sup.-5 inches/hour 
 It is important to understand the interaction of dissimilar metals in 
 solid-liquid states. The joining of dissimilar metals in a solid-liquid 
 state is governed by their physico-chemical properties and by the 
 interaction between them; or, in the case of more complex systems, such as
 super alloys, by their interaction with all other alloying elements and 
 impurities. When the melting point of the corrosive metal (the coating 
 alloy in our case) is much lower than that of the metal being attacked 
 (the roll material), the roll material may remain in a solid state 
 throughout the process. In this case, a strong metallic bond between the 
 atoms of the coating metal and the roll material occurs in the wetting 
 process. It is true, however, that other associated processes can 
 significantly influence the attack rate and kinetics of solubility, i.e., 
 dissolution, interdiffusion and formation of intermetallics that have a 
 significant effect on the bonding properties of the intermetallic layers 
 being formed. 
 Experimental as well as theoretical findings have shown that the attack on 
 a solid metal by zinc and zinc/aluminum alloys is a topochemical reaction 
 in which a two-stage formation of strong bonds between atoms of the two 
 materials is a characteristic feature. 
 In the first stage, a physical contact is established by the close 
 proximity of the two metals allowing interaction between the atoms. The 
 electrostatic interaction between the surface atoms is of great importance
 at this stage. 
 In the second stage, the chemical interaction takes place and the formation
 of a strong bond is completed. At this stage, quantum processes between 
 the electrons prevail. Thus, the occurrence of electron interaction of 
 different types of materials requires a definite quantity of energy for 
 surface activation. This energy, in the case of "hot dip" metalizing is 
 imparted in the form of heat retained in the molten metal that is 
 maintained at temperatures well above their melting temperature in order 
 to improve the coating capability of the melt alloy in accelerated 
 production. In other words, the lower the temperature of the melt in the 
 pot, the slower the two basic stages of alloying formation. 
 During galvanizing and aluminizing, both stages as well as the subsequent 
 diffusion take place so fast that it is difficult to join zinc/aluminum to
 steel without the formation of brittle intermetallic layers at the contact
 zone. Zinc/aluminum alloys are so active that adhesion and diffusion into 
 steel is achieved even in the presence of a passive film of iron oxides, 
 as long as the oxide layer is no thicker than 100 .ANG. (see FIG. 1). 
 In order to improve the resistance of ferrous alloys to molten aluminum, it
 is necessary to study the dissolution process that follows wetting in 
 detail. The dissolution of solid ferrous alloys into molten aluminum has 
 been studied by M. Kosaka and S. Minowa (Transactions Iron & Steel 
 Institute of Japan, Vol. 50 and 52, 1964.) It is also theoretically 
 described by Nernst-Shchukarev's equation 
EQU dc/dt=K.sub.s A/V.(C.sub.s -C.sub.i) (1) 
 where C.sub.i =the instantaneous concentration of the dissolved metal in 
 the melt (weight percent) 
 C.sub.s =the saturation concentration at operational temperature (weight 
 percent) 
 K.sub.s =the dissolution rate constant (or mass transfer coefficient) 
 A=the surface exposed to the Zn/Al melt 
 V=the volume of the melt 
 From this equation and assuming the dissolution of the solid metal is 
 controlled by mass transfer into the molten metal, the rate in weight loss
 of a roll submerged in the molten alloy is expressed by the equation 
EQU -dw/dt=K.sub.s A(C.sub.s -C)-(2) 
 Where t=time 
 K.sub.s =dissolution rate constant 
 A=roll surface area exposed to the melt 
 C.sub.s =saturation concentration of the melt at the operating temperature 
 in weight percent 
 C=instantaneous concentration of the weight dissolved metal in the melt in 
 weight percent 
 w=the weight of the roll 
 Since in a coating line the melt alloy is being added continuously, it can 
 then be accepted that 
EQU C=O (or approximately zero) 
 Then, 
EQU -dw/dt=K.sub.s A C.sub.s (3) 
 or 
EQU .delta.: Density of the Material (4) 
 In other words, the dissolution of the roll material into the melt depends 
 ##EQU1## 
 on two elements: 
 a. C.sub.s -A coefficient, independent of time, whose value can be obtained
 from the concentration of the liquidus curve at the operating temperature 
 in the phase diagram for each of the components of the roll material (see 
 FIGS. 3 and 4). 
 ##EQU2## 
 b. K.sub.s -A coefficient, dependent on time, that establishes the kinetics
 of dissolution of each one of the component elements of the super alloy 
 roll material (see FIGS. 6a and 6b). 
 Utilizing metals or transition metals with a C.sub.s equal to 0 at the 
 operating temperature of the melt, obtains a non-wetting, zero-solubility 
 alloy for operation in the alloy melt. Unfortunately, only two such 
 materials exist for zinc, namely, tungsten and carbon. Only one exists for
 aluminum, carbon. 
 It has then been necessary to study the dissolution coefficient, K.sub.s, 
 for metals and transition metals, and its change with changes in operating
 temperatures, velocity, etc. (N. Tunca, G. W. Delamore and R. W. Smith) 
 (Metallurgical Transactions Association, Vol. 21A, Nov. 1990) and to 
 establish its variation in value for binary and ternary alloys (V. R. 
 Ryabov) (Aluminizing of Steel, Oxonian Press Pvt Ltd N.D.) 
 Due to the large quantity of experimental data needed and the scarcity of 
 it, an attempt has been made to establish a correlation between the 
 theoretical values of the dissolution coefficient, K.sub.s, with the 
 existing experimental values (Mitsuo Niinomi and Masamichi Sano). 
 (Dissolution of Ferrous Alloys into Molten Aluminum, Transactions of the 
 Japan Institute of Metals, Vol. 23, No. 12) Unfortunately, it was 
 established that the kinetics of dissolution of metals and transition 
 metal alloys in zinc/aluminum melts do not follow the Nernst/Shchukarev 
 equation. The differences of the coefficients K.sub.s obtained should be 
 attributed to: 
 a. The mechanism of dissolution (static, natural convection dynamic, etc.).
 b. The relationship to the appearance and growth peculiarities of the 
 intermetallic phases formed at the interface of the solid and liquid 
 metals. The growth of these intermetallic phases in zinc/aluminum alloy 
 melts, as discussed earlier, is extremely fast. Their growth decreases the
 dissolution rate, and with C.sub.s and A values constant, the value of 
 K.sub.s must decrease with time below the theoretical value (see Equation 
 3). Finally, the dissolution process changes to an intermetallic 
 layer/alloy melt diffusion controlled process. This occurs when the 
 critical thickness of the intermetallic layer is reached and dissolution 
 reaches equilibrium (see FIGS. 1a-1f, 6a and 6b). 
 Additional studies were conducted using V. G. Levich's equation in order to
 enter into consideration the rotational velocity of the roll (See FIGS. 6a
 and 6b). 
EQU dc/dt=0.62 S/V D.sup.2/3 v.sup.-1/8 w.sup.1/2 (C.sub.s -C.sub.i) (5) 
 where C.sub.i =the concentration of additive 
 C.sub.s =the saturation concentration 
 t=the time 
 S=the surface area of the specimen 
 V=the volume of the aluminum/zinc melt 
 D=the diffusion coefficient 
 V=the kinematic viscosity 
 w=the angular speed of rotation 
 Equation 5 at t=C=0 and w=1 may be used to theoretically evaluate the 
 diffusion of metals in molten metals. This effect was partially initiated 
 and conducted with some degree of correlation by T. Heumann and S. 
 Dittrich. (Z. Metallkunde, Vol. 50, 1959, p. 47-617) 
 A fourth approach undertaken to analyze the behavioral change of the 
 diffusion coefficient utilizes the Stokes-Einstein formula. The difference
 between the coefficient obtained by the theoretical calculations with that
 reduced from experimental tests utilizing the theoretical formulae differ 
 in value by 10 to 20 times, even though the techniques used in determining
 the coefficients in most cases were almost the same. 
 Perhaps the most valuable information is that derived from the following 
 facts: 
 a. Aluminum and zinc do not attack or wet most oxides, carbides or 
 nitrides. 
 b. At steady-state equilibrium, KS is no longer a variable function of time
 (K.sub.s =f(t)) but a constant. 
 c. The investigations (V. R. Ryabov, Alitirovanie Stali, Chapter IV, 
 Metalurgiya Publishers Moscow) on how the addition and concentration of 
 other elements to iron affected the diffusion zone, and formation of 
 intermetallics and change in the dissolution rate, lead to the following 
 conclusions: 
 Carbon 
 The structure of iron-carbon alloys formed by slow cooling from the 
 .gamma.-solid solution region is well known. Only the nature and the 
 properties of the diffusion zone developed when immersing the prepared 
 alloys in aluminum as a function of temperature, duration and the 
 percentage of carbon content were studied. 
 With an increase in carbon content from 0.2 to 0.56% the thickness of the 
 intermediate layer varies (FIG. 8a) insignificantly at both 750 and 
 850.degree. C. (from 110 to 125 and from 90 to 110 .mu.m respectively). 
 The layer has greater solubility at 850.degree. C. and is, therefore, 
 thinner, increasing the mass transfer rate. 
 The microstructures showed the presence of an irregular, tongue-shaped 
 intermediate layer (see FIGS. 7a-7f). 
 Since aluminum decreases the solubility of carbon in liquid and solid iron,
 carbon is forced out from the solid solution of iron during formation of 
 the intermediate layer and an area rich in carbon develops immediately in 
 front of the diffusion zone (see area labeled 3 in FIG. 7f. 
 This happens, because carbon, unlike iron and aluminum, cannot penetrate 
 through the intermetallic layer. Carbon is deprived of its solvent and 
 aluminum travels deeper into the specimen and pushes carbon ahead of 
 itself; nonetheless, carbon retards the dissolution of iron by acting like
 a barrier to the aluminum diffusion progress. 
 Aluminum may be partially combined with carbon forming either Al.sub.4 
 C.sub.3 AlC.sub.3 carbides or Fe.sub.3 AlC.sub.x carbide. The 
 microhardness of the intermediate layer showed little dependance on the 
 carbon content in the alloy. It should be noted that the structure of the 
 layer changed with an increase in carbon content in the alloy. This can be
 explained by the change in the structure of steel from purely ferritic to 
 pearlitic. 
 In the interaction of aluminum with iron, when the latter has a 
 body-centered cubic lattice .alpha.-Fe, the diffusion layer is always 
 jagged towards the iron side (FIG. 7f). In the case of interaction of 
 aluminum with iron or an alloy containing iron in a face-centered cubic, 
 .gamma.-Fe, the diffusion layer has smooth boundaries. 
 Nickel 
 The maximum nickel (and chromium) content in the alloys were the same as in
 stainless steel to examine the effect of each percentage element 
 individually. 
 Nickel belongs to the group of those elements forming a continuous series 
 of solid solutions with iron. Introduction of nickel into iron widens the 
 .gamma.-Fe region. Nickel has a low C.sub.s in zinc, but it is very high 
 in aluminum; and its addition is equivalent to an increase in temperature 
 of the alloy melt (see FIGS. 17a and 17b). 
 The thickness of the intermediate layer significantly changes with an 
 increase in the nickel content in the steel substrate from 1.92 to 12% at 
 750.degree. as well as at 850.degree. C. The thickness of the intermediate
 layer decreases rapidly from 70-100 .mu.m (1.92% Ni) to 10-14 .mu.m (8.5% 
 Ni), afterward increasing as the nickel is increased up to 12% (see FIG. 
 8b). The intermediate layer is smooth along the thickness without tongues.
 X-ray analysis of the phase composition of joints of aluminum with steel 
 alloyed with nickel, established that the quantity of the intermetallic 
 phases Fe.sub.2 Al.sub.5 decreases with an increase in nickel 
 concentration in steel. Thus, in the specimens with 1.92% Ni, a wide 
 intermediate layer of Fe.sub.2 Al.sub.5 phases was detected, whereas, in 
 the specimens with 8.5% Ni contents, this phase was negligible. 
 Chromium 
 Chromium belongs to the group of alloying elements, which narrow the 
 .gamma.-region. The chosen chromium content and the aluminizing 
 temperature do not alter the region of phase changes, as seen in the 
 iron-chromium phase diagram (FIG. 15). 
 Thus, the diffusion zone, formed in aluminizing, varied only as a function 
 of the chromium content in the substrate alloy, temperature and duration 
 of aluminizing. The thickness of the intermetallic zone is almost 
 independent of the temperature but increases rapidly (80 .mu.m to 140 
 .mu.m) with an increase in chromium content from 2.20 to 7.2% and 
 decreases when chromium reaches 12.4% to 75 .mu.m (FIG. 9a). 
 The thickness of the intermetallic layer (Fe.sub.2 Al.sub.5 phase) 
 decreased when alloyed with chromium also, but to a much lesser extent 
 than when alloyed with nickel. Thus, the Fe.sub.2 Al.sub.5 phase with 7.2%
 Cr is about the same as in unalloyed steels. The thickness of the 
 intermetallic layer (Fe.sub.2 Al.sub.5 phase) decreased only in the alloys
 containing more than 10% Cr. 
 Manganese 
 Manganese is one of the alloying elements which widens the .gamma.-region, 
 behaving very much like nickel. A continuous series of solid solutions 
 does not appear in a solid state in the iron manganese system. Diffusion 
 of manganese in .alpha.- and .gamma.-iron is more difficult than the 
 diffusion of carbon. The alloys prepared cross only through .alpha.- and 
 .alpha.+.gamma.-region during aluminizing. 
 As seen in FIG. 9b, the thickness and hardness of the intermetallic zone, 
 formed in aluminizing, decrease with an increase in manganese content in 
 the steel substrate thus increasing the dissolution rate and the 
 dissolution coefficient, K.sub.s. 
 Silicon 
 Although silicon belongs to those elements which narrow the .gamma.-region,
 it behaves in a more detrimental manner than nickel because of its high 
 C.sub.s in aluminum and the reduction of aluminum's melting temperature as
 silicon percentage increases to 12.6% at the eutectic ratio (FIGS. 18a and
 18b). The silicon content was chosen so that only the .gamma.-region was 
 covered at the aluminizing temperatures. No other phase changes occurred 
 in the alloys during aluminum immersion. Thus, the pattern of the 
 diffusion zone formed during aluminum immersion depended upon the silicon 
 content in the alloy, aluminum melt temperature, and time the specimens 
 were held in liquid aluminum. 
 An additional problem with silicon is that it does not generate carbides at
 the standard processing temperatures, as do vanadium, tungsten and 
 molybdenum; so, its use in resistant alloy formulations should be 
 excluded. 
 Silicon, if added to the aluminum bath, has a greater effect than 
 increasing its content in steel. An increase of silicon in either the bath
 or steel reduces the hardness and thickness of the diffusion layer to the 
 same extent. With an increase of silicon in steel, the silicon content in 
 the diffusion layer increases. 
 Boron, Titanium, Vanadium and Molybdenum 
 Boron very strongly narrows the .gamma.-region. There are only two 
 mechanisms by which a crystal can dissolve atoms of a different element: 
 interstitial and substitutional. Boron and carbon are the only elements 
 with atoms small enough to fit into the interstices of iron crystals. The 
 other small-diameter-atom elements, such as oxygen, hydrogen and nitrogen,
 tend to form compounds with metals instead of dissolving in them. The 
 addition of these atoms (boron and carbon) creates a strong increase in 
 the crystal's internal energy, strengthening the alloy and reducing its 
 solubility in zinc/aluminum melts. 
 Phase changes in iron-titanium alloys set in only above 900.degree. C. An 
 introduction of titanium in iron strongly narrows the .gamma.-region. 
 Vanadium and molybdenum drastically limit the .gamma.-region. No phase 
 changes took place at aluminum immersion temperatures tested with the 
 selected molybdenum and vanadium content in the alloys (FIG. 9c). 
 Component Elements Selection Criteria 
 Thus, all the elements which increase the thickness of the diffusion layer 
 and reduce the mass transfer rate narrow the field of the 
 .gamma.-modification in the iron alloying element phase diagram. The 
 elements acting in the opposite manner widen the .gamma.-region. This 
 occurs because the diffusion rate of different elements in the 
 .alpha.-modification of iron with a bcc (Body-centered cubic) structure is
 greater than in the .gamma.-modification, with an fcc (Face-centered 
 cubic) structure (V. R. Ryabov and V. D. Duplyak, 1968). (Protective 
 Coatings on Metals, Naukova-Domka Kiev No. 5, p. 89-94). 
 Alloying elements which narrow the .gamma.-region at the process 
 temperature favor the .gamma.-.alpha.transformation of iron at small 
 concentrations of aluminum, thus increasing the overall depth of 
 diffusion. The effect of this phase change, the variation of the lower 
 limit of dissolution of aluminum in .alpha.-iron by the alloying elements,
 is overlapped by the effect of the added alloying elements on the rate of 
 diffusion of aluminum in each of the modifications of iron and also by the
 change of the upper limit of dissolution of aluminum in iron. The combined
 effect of these facts determines the relationship of the depth of aluminum
 penetration to the content of the alloying elements. This is clearly seen 
 when alloying iron with manganese, which reduces the thickness of the 
 diffusion layer only on the basis of widening of the .gamma.-region in the
 Fe--Mn system, accelerating the corrosion attack. 
 From the kinetics formation of the diffusion layer and the growth thickness
 and properties of the intermetallic layers between solid-liquid phases, it
 can be concluded that if a metallic alloy must be formulated to resist 
 zinc/aluminum melts, it should meet the following requirements: 
 a. The components of the alloy should have the lowest saturation 
 concentration possible, i.e., 1%&gt;C.sub.s =0 at the metalizing operating 
 temperature (see FIGS. 2-5). 
 b. The alloying elements must narrow the .gamma.-Fe region, and their 
 percentage content should be such that only the .gamma.-region is covered 
 at the coating operational temperature. 
 c. The intermetallic layer thickness formed during molten zinc/aluminum 
 immersion at steady-state conditions should not be less than 
 1.2.times.10.sup.2 .mu.m. 
 d. Elements that reduce the melting temperature of either zinc or aluminum 
 should not be used as components of the melt resistant alloy (FIGS. 18a 
 and 18b). In other words, 
EQU 1.times.10.sup.-2 1/.degree. C.&gt;dc.sub.s /dT.ltoreq.0 
 e. The formation of strong covalent bonded molecules of the type M.sub.x 
 C.sub.y should be promoted to generate a microstructure rich in hard and 
 steady carbides, resistant to molten zinc/aluminum, having tough complex 
 matrix structures. 
 f. Maximization of the carbides to matrix ratio should be secured by proper
 selection of the carbon ration to carbide forming elements, thus, also 
 assuring a reduction of the exposed effective area (Equation 3). 
 Based on the preceding studies and conclusions, the optimum components for 
 materials to be used in a zinc/aluminum alloy melt should be: boron, 
 carbon, cobalt, chromium, molybdenum, niobium (columbium), titanium, 
 vanadium, tungsten and zirconium (FIGS. 2, 3, 4, 5, 11a, 11b, 12a, 12b, 
 13a, 13b, 14a and 14b). Despite the fact that chromium does not show 
 optimum C.sub.s values with aluminum in its phase diagram, it has been 
 included because of the excellent percentage of chromium carbides that can
 be generated during processing (Cr.sub.23 C.sub.6). Nickel, iron and 
 cobalt necessary to form the solid solution matrix and increase thermal 
 shock and impact resistance, on the other hand, must be used with great 
 discretion. In FIGS. 17a and 17b, the nickel/zinc diagram shows a low 
 C.sub.s. On the other hand, it's use is extremely detrimental in the 
 presence of even moderate quantities of aluminum, since it not only has a 
 high C.sub.s, but reduces the melting temperature of aluminum. 
 This is similar to silicon that acts equivalently to an increase in the 
 aluminum temperature, accelerating the corrosion attack. For this reason, 
 the percentage of nickel used in a zinc/aluminum melt resistant alloys 
 should decrease with increases of percentage of aluminum in the melt, and 
 its replacement by cobalt or iron will show a beneficial reduction in the 
 mass transfer rate by the formation of F.sub.3 C and CO.sub.3 C during 
 processing. The reason for avoiding the use of iron with low percentages 
 of aluminum is because of its high C.sub.s in zinc (FIG. 16b)--for cobalt,
 its extremely high price. 
 The second requirement for a roll material requires it to be non-wettable 
 to a zinc/iron or zinc/iron/aluminum dross. Surface phenomena as stated 
 earlier plays a decisive role in the formation of strong bonds (J. A. 
 Morando, U.S. Pat. No. 5,338,280, Columns 1 and 2). Published data on the 
 subject is scarce or non-existent. Although some efforts have been made, 
 especially by V. R. Ryabov, ("Aluminizing of Steel Alloys Oxonian Press, 
 1985, pages 1-7, 61-83) The latter has devoted considerable effort toward 
 understanding the kinetic relationship of steels and iron wetting with 
 aluminum and zinc, establishing the fact that the spreading rate increased
 with increase in temperature as well as the activation energy required. 
 Data and analyses performed by J. A. Morando strongly suggest that the work
 of adhesion of metals and transition metal alloys decreases with increases
 in the surface hardness and a reduction of surface energy of the adhesion 
 resistant alloy. This is due perhaps to the fact that the surface hardness
 of the resistant alloy is a consequence of the concentrations of low 
 surface energy carbides (WC, MoC, VC, etc.) present on the roll surface. 
 By formulating a material based on the restraints of the selected criteria,
 the mass transfer rate is reduced with the increase in complexity of the 
 intermetallic layer and with a decrease in the bonding strength of the 
 diffusion layer, as a consequence of the minimization of matrix exposure 
 and reduction of exposed effective area. The carbides' surface coverage 
 makes the dross adhesion by mechanical action less likely, since the strip
 running through the roll surface can easily remove it before it can build 
 up to a detrimental thickness that affects the quality of the coating 
 finish. 
 In testing performed with very high percentages of aluminum (50% to 60%), 
 this principle of dross attachment to the roll was corroborated. The roll 
 whose material was formulated as suggested in this patent application was 
 allowed to run for a period of 45 hours. In that period, the thickness of 
 the dross accumulated in the zone in contact with the steel strip built up
 to nearly 0.125 inch thick. The portion of the roll not in contact with 
 the steel strip showed no dross build-up whatsoever. The test was repeated
 while the submerged roll surface was being periodically scrapped through a
 mechanical device. The roll continued to operate in excess of 120 hours 
 before the dross build-up accumulated to 0.125 inch thick. In other words,
 the bond strength of the diffusion layer of the zinc/aluminum alloy to the
 roll material was so weak that it could be removed by simple mechanical 
 means and with it the zinc/iron/aluminum dross that had mechanically 
 locked itself to it. 
 Surface Hardness 
 A high surface hardness (R.sub.c larger than 40) is necessary. This is a 
 mechanical requirement imposed by the fact that the roll surface is acting
 as the bearing surface for the steel strip being processed, and sliding 
 friction between the two will occur during operation. The wear caused by 
 this sliding friction can be greatly reduced if the material hardness is 
 above R.sub.c 40. One of the many reasons for the poor performance of 316L
 stainless steel is the fact that it cannot be scraped to remove melt dross
 due to its very soft surface. On the other hand, if it could have been 
 removed, it would fail even faster due to abrasive wear. 316L surface 
 hardness is approximately R.sub.c 10 and highly inadequate. 
 The materials formulated in accord with the selection criteria, because of 
 the high carbide densification and distribution as well as the toughness 
 of the solid solution matrix that contains them, have shown excellent 
 dimensional stability at temperatures up to 1600.degree. F. Roundness of 
 rolls removed from the melt at operating temperatures up to 1380.degree. 
 F. was within 1.times.10.sup.-3 inches total indicated reading. On the 
 other hand, 316L stainless steel, as well as any other stainless steel 
 materials, because of the interaction of nickel/chrome/iron and the change
 in phases when subject to temperatures above 600.degree. F. show poor 
 geometrical stability and tend to deform. 
 Thermal shock resistance and impact resistance, can be achieved by proper 
 utilization of nickel, iron and cobalt necessary to form the solid 
 solution matrix that will contain the carbides as outlined in the 
 discussion of the selection criteria. 
 Because of the discussions presented regarding nickel's performance in 
 molten aluminum and molten zinc, the following rules should apply: 
 a) If the aluminum percentage in the melt is: 
EQU 0.ltoreq.Al&gt;5% 
 then, the nickel percentage in the melt resistant super alloy should be: 
EQU 10%&lt;Ni&lt;30% 
 b) If aluminum percentage in the melt is: 
EQU 50%&lt;Al&lt;100% 
 then 
EQU Ni=0 
 c) If zinc percentage in the melt is: 
EQU Zn&gt;50% 
 then, 
EQU 10%&lt;Fe&lt;30% 
 d) The percentage of cobalt should be based on the balance of the 
 composition needed to maximize the carbide-to-matrix ratio while 
 maintaining the material hardness. 
EQU 40&lt;Rc&lt;50 
 Following the preceding general rules will provide a material having good 
 thermal shock resistance, good impact and notch resistance strength and 
 which is centrifugally castable and machinable. 
 Molten Zinc/Aluminum Resistant Advanced Material Formulation General 
 Chemical Composition 
 Taking into consideration all the previous discussions included in this 
 disclosure, we are now in a position to formulate the chemical composition
 limits of a super alloy material capable of resisting molten zinc/aluminum
 alloy melts and their drosses. 
 
 % Component % 
 1.6 &lt; C &lt; 2.6 
 15.0 &lt; Cr &lt; 30.0 
 0.0 .ltoreq. Ni &lt; 30.0 
 10.0 &lt; W &lt; 30.0 
 2.0 &lt; Mo &lt; 8.0 
 0.0 .ltoreq. V &lt; 6.0 
 0.0 .ltoreq. Nb &lt; 6.0 
 0.0 .ltoreq. Co &lt; 20.0 
 0.0 .ltoreq. B &lt; 5.0 
 10.0 &lt; Fe &lt; 50.0 
 0.0 .ltoreq. Zr .ltoreq. 6.0 
 Super Alloy Chemical Composition for Zinc/Aluminum Alloy Melts Containing 
 Less than 5% Aluminum 
 The general chemical composition converts to: 
 
 % Component % 
 1.9 &lt; C &lt; 2.3 
 24.0 &lt; Cr &lt; 30.0 
 18.0 &lt; Ni &lt; 26.0 
 15.0 &lt; W &lt; 25.0 
 4.0 &lt; Mo &lt; 8.0 
 4.0 &lt; V &lt; 6.0 
 0.0 .ltoreq. Nb &lt; 2.0 
 0.0 .ltoreq. Co &lt; 6.0 
 0.0 .ltoreq. B &lt; 2.0 
 18.0 &lt; Fe &lt; 24.0 
 0.0 .ltoreq. Zr &lt; 6.0 
 The range of modification to this chemical composition should be based on 
 the melt's operating temperature which imposes additional restrictions, 
 especially with respect to thermal shock resistance of the super alloy 
 material under consideration. 
 Super Alloy Chemical Composition for Zinc/Aluminum Alloy Melts Containing 
 More than 50% Aluminum 
 From all previous discussions, the following chemical composition for the 
 melt resistant super alloy can be derived: 
 
 % Component % 
 1.9 &lt; C &lt; 2.3 
 16.0 &lt; Cr &lt; 24.0 
 0.0 .ltoreq. Ni &lt; 2.0 
 15.0 &lt; W &lt; 25.0 
 4.0 &lt; Mo &lt; 8.0 
 4.0 &lt; V &lt; 6.0 
 0.0 .ltoreq. Nb &lt; 2.0 
 5.0 &lt; Co &lt; 15.0 
 0.0 .ltoreq. B &lt; 2.0 
 35.0 &lt; Fe &lt; 45.0 
 0.0 .ltoreq. Zr &lt; 6.0 
 The range of modification to this chemical composition should be based on 
 the melt's operating temperature which imposes additional restrictions, 
 especially with respect to thermal shock resistance of the super alloy 
 material under consideration. 
 Super Alloy Chemical Composition for Zinc/Aluminum Alloy Melts Containing 
 Different Percentages of Aluminum 
 From the general discussions, it will not be extremely difficult to 
 formulate zinc/aluminum melt resistant alloys if a proportional 
 interpolation is conducted. Final optimization of the alloy selected must 
 be performed in a test of actual application, since other variables may 
 exist such as additives used in the alloy melt (silicon, strontium, etc.) 
 may require slight modifications in the composition, although the 
 principles for the formulation remain. 
 Economic Viability 
 The material formulations arrived at in spite of their superior performance
 to standard materials cannot escape the need for economic viability. The 
 number of special components required in the formulation of the materials 
 under consideration, will translate into a price increase. This price 
 increase can be overcome if the roll designs follow the methods outlined 
 in my patent disclosures: 
 a) "Composite Centrifugally Cast Furnace Roll Rings for Furnace Rolls and 
 Method of Making Same", Ser. No. 08/383,578, filed Feb. 3, 1995; 
 b) "Multi-Cast Furnace Roll and Method for Making Same" (pending). 
 In other words, the roll comprises a centrifugally cast outside layer made 
 of the melt resistant super alloy (as outlined in this disclosure) and an 
 inner liner of a material having different solubility and hardness 
 characteristics from the outside layer. The object of the utilization of 
 the multicast process, in this case, is that the high hardness and 
 solubility resistance of the outer layer is not necessary throughout the 
 thickness of the roll, that is mostly dictated as a function of mechanical
 requirements and stress levels. 
 Experimental rolls made in this fashion, with an outer layer thickness of 
 0.750 to 1.0 inch and an inner layer thickness of 1.0 to 1.5 inches 
 operating in molten zinc/aluminum melts have shown continuous operation 
 without maintenance three to five times longer than standard materials and
 a total roll life of six to ten times that of standard materials. Similar 
 performance has been obtained from roll bearings manufactured of materials
 formulated in accordance with the present disclosure, yet their cost was 
 only 50% to 80% higher than those made of standard roll material.

DESCRIPTION OF THE PREFERRED EMBODIMENTS 
 According to one preferred embodiment of the invention, rolls made in 
 accordance with the invention and generally as shown in FIG. 19, were made
 of the following alloys: 
 
 B C Co Cr Mn Mo Nb Ni Si Ti V W Zr Fe 
 Composition of AT101 (For AL .gtoreq. 50%) in percent by weight: 
 -- 2.2 15 18 1.0 4 -- -- -- -- 3 10 -- 45 
 Rc = 42 
 Composition of AT103 (For Zn .gtoreq. 90.degree.) in percent by weight: 
 -- 2.0 4 27 6 4 -- 20 -- -- 4 12 -- 21 
 Rc = 42 
 Composition of AT101G (For AL .gtoreq. 95%) 
 2.0 2.0 5 16 1.0 4.0 2.0 -- -- -- 3.0 16.0 1.0 47 
 Rc = 50 
 Referring to FIG. 19, a typical roll 10 is illustrated in a galvanizing 
 bath 10 below metal line 14. A metal strip 16 is trained around roll 10 
 and tensor rolls 18 and 20. Mechanical scraper 22 removes dross build-up 
 on the roll. 
 These rolls were prepared by centrifugal casting methods, tested and the 
 results, measured in terms of alloy loss onto the bath were as follows: 
 (.DELTA.D) DIAMETER LOSS PER DAY vs ROLL MATERIAL AND MELT COMPOSITION 
 
 BATH MELT COMPOSITION 
 Roll 45% Zn + 95% AL + 
 Material of Zn @ 55% AL 5% Si 100% AL 
 Test Samples 975.degree. F. 1120.degree. F. 1280.degree. F. 1360.degree. 
 F. 
 Cast Iron -- -- .080 IN/ .120 IN/ 
 (2.0 C) DAY DAY 
 52100 -- .046 IN/ .100 IN/ -- 
 DAY DAY 
 316L .003 IN/ .030 IN/ .130 IN/ .200 IN/ 
 DAY DAY DAY DAY 
 AT101 .0003 IN/ .0035 IN/ .012 IN/ .020 IN/ 
 DAY DAY DAY DAY 
 AT103 .0000 IN/ -- -- -- 
 DAY 
 112 Day 30 Day 10 Day 6 Day 
 Test Test Test Test 
 CORROSION ATTACK RATIOS 
 
 ##EQU3## 
 The following is a comparison of tests of a sink roll made with the 
 inventive alloy and a conventional 52100 low carbon steel roll, when 
 operating in molten 95% aluminum and 5% silicon at 1280.degree. F. 
 
 ROLL MATERIAL 
 AT-101 52100 
 Roll thickness .650 in. 3.00 in. 
 Roll hardness R.sub.c 42 R.sub.c 20 
 Roll diameter spooling .012/day .080/day 
 loss 
 Roll diameter corrosion .012/day .080/day 
 loss 
 Roll dross adhesion None Medium to high 
 Roll surface damage Minor (.50 in. dia.) Major (2.0 in dia.) 
 Roll stability Excellent (.001 in.) Good (.010 in.) 
 Strip quality Excellent Good 
 Strip edge Straight Waving 
 Strip surface None to minor Minor to medium 
 imperfection 
 Strip marks spandrel Excellent Good 
 Campaign length 5 days* 3 days*** 
 Number of campaigns 4** 5 
 Total life 20 days 15 days 
 *Line never stopped because of sink roll failure or dross build-up. 
 **Discontinued test because line maintenance shut down. 
 ***Line stopped four times because of sink roll dross build-up. 
 In another test, stabilizer roll half bushings made of the AT103 low 
 solubility alloy was run in a 99% zone galvanizing line. Previous half 
 bushings and sleeves made of Stellite material would wear out in four to 
 five weeks. 
 The AT103 half bushings were placed into service with Stellite roll 
 sleeves. The Stellite sleeves wore away within a one month period. The 
 AT103 half bushings showed no "scoring" or "wear" on the I.D. The same 
 bushings were reused with a new set of Stellite sleeves which were also 
 worn through in a three week period. Visual inspection of the bushings 
 after the second run showed no "scoring" and only slight wear on one side 
 (probably due to misalignment).