SINTERING OF POWDER PREMIXES A BRIEF OVERVIEW

Transcription

1 SINTERING OF POWDER PREMIXES A BRIEF OVERVIEW Kalathur S.Narasimhan and Frederick J.Semel Hoeganaes Corporation 1001 Taylors lane Cinnaminson,NJ Abstract Advances in the understanding of the sintering of powder premixes have contributed significantly to the growth of the ferrous powder metallurgy industry. This includes sintering both in the solid state and in the presence of a liquid phase. In this article, the sintering of iron powder premixes containing: 1) graphite; 2) nickel and graphite; 3) copper and graphite; 4) Phosphorus as ferrophosphorus; and, 5) boron as ferroboron are discussed. The evolution of microstructure and mechanical properties are discussed as well.

2 INTRODUCTION The growth of ferrous powder metallurgy (P/M) over the past five decades has been outstanding as this technology is proving itself as an alternative lower cost process to machining, casting, stamping, forging, and other similar metal working technologies. The growth has been spurred by continuing technological advancements in powder making and processing, alloy development, and parts production methods. The markets for the powder metallurgical parts are varied. Automotive applications have so far dominated the industry growth. However, non-automotive applications are becoming increasingly important. The latter applications mainly include hand tools, household and lawn appliances and industrial motor controls and hydraulics The growth of P/M is dependent on the ability of technological advancements to deliver products with higher densities than are currently available. Figure 1 shows how parts distribution by density has progressed and the anticipated use of higher density parts in the future. The indicated density increases to date are mainly due to improvements in powder making and compaction processing. The future advances are likely to come mostly from our understanding of the sintering process. P/M parts are characterized by density, composition, and microstructure. For a given part, these three parameters are optimized. Sintering and additional heat treatments of powder premixes generate the microstructures that are needed to meet the required performance. Generally, there are four types of powder premixes that are used in the P/M industry, (2). These are: 1) iron powder mixed with alloying ingredients such as C, Cu, Ni and P; 2) iron pre-alloyed with Mo or with Mo and Ni and mixed with C; 3) partially alloyed iron with Mo, Ni and Cu and mixed with carbon; and, 4) hybrid systems that use various combinations of 1, 2, and 3. All of these premixes contain a lubricant and in many cases also contain a binder to prevent compositional variations due to demixing and dusting. Prior to sintering and any additional heat treatments, the premixes are compacted to near net shape. Typically, the compaction process determines the final density to within one or two percent and the sintering process provide the balance.

3 P/M Part Distribution by Density Figure 1: History projection of density advances with time. In this article, we will discuss the sintering, microstructure, and mechanical properties of the four most common alloy systems including: 1) iron-carbon; 2) iron-carbon-nickel; 3) iron-carbon-copper; and, 4) iron-phosphorus. The iron-boron system as potential future alloy will also be discussed. In general, when powder premixes are sintered, sintering is by solid-state processes or a combination thereof with liquid phase processes. In both cases, atom movements plays a key role. The basic mechanisms include vapor transport and surface, volume and grain boundary diffusion. The vapor transport and surface diffusion mechanisms effect metallurgical bonding and pore rounding but not densification. The volume and grain boundary diffusion mechanisms effect all three: bonding, pore rounding and densification. Typically, in ferrous PM, all four processes occur to some extent. Since vapor transport and diffusion are thermally activated processes and iron is subject to oxidation, sintering necessarily involves high temperatures and protective atmospheres. In the case of iron base premixes, this usually means temperatures in the range of 1100 to 1300 o C, (2000 to 2350 o F) and nitrogen/hydrogen atmospheres ranging anywhere from 5% to 100% hydrogen. Hold times are typically from 15 to 30 minutes but on occasion may be as long as an hour or more. Excellent reviews of the sintering process and of the various sintering mechanisms and their effects are available in the open literature, (3, 4). Fe-C and Fe-C-Ni Systems: The response of Fe-C and Fe-C-Ni premixes to sintering in terms of the microstructures and mechanical properties that are obtained are very similar. A typical Fe-C

4 microstructure based on an iron powder premix containing 0.5% graphite as the carbon source is shown below in Figure 2. Such a composition is normally designated as F-0005, (F = Iron, 5 = 0.5% C). F Figure 2: Microstructures of sintered iron with 0.5% carbon. Sinter bonding of the powder particles and two phases, pearlite and ferrite, are visible. The micrographs each show two phases. The light-etching phase is ferrite, essentially, pure iron or if nickel is present, a solid solution of nickel in iron. The dark etching phase is pearlite, a lamellar structure of the ferrite and iron carbide, Fe 3 C, known as cementite. The relative amounts of the two phases are largely dependent on the carbon content and to a lesser extent on the process conditions. Their morphologies, on the other hand, are almost entirely dependent on the process. Typically, higher carbons and faster cooling rates subsequent to sintering lead to higher pearlite contents, decreased lamellae spacing and finer grains of both phases. The micrographs also clearly show a good degree of sinter in terms of pore rounding and sinter bonding of the iron particles. The porosity seen here is typical of most ferrous P/M materials. More generally, the numbers and shapes of the pores are a function of the as-compacted density and the sintering conditions, especially the temperature. The microstructures of the sinterings are important, of course, because they largely determine the resultant mechanical properties. In general, increases in the pearlite content and decreases in the lamellae spacing and grain size increase strength but decrease ductility while increases in the ferrite content increase ductility but decrease strength. Increases in the density and degree of sinter increase both strength and ductility. The mechanical properties of an iron powder premix with 0.8% graphite, (i.e. F- 0008), are shown in Table 1, (5).

5 Table 1: Mechanical Properties of F-0008 carbon steel sintered at 1120 C Density (g/cm³) UTS (MPa) YS (MPa) Elongation In 25.4 mm Impact Energy (J) Hardness (HRB) < to The microstructure of an iron powder premix containing 0.5% graphite and 2% elemental nickel, (i.e. FN-0205), and otherwise submitted to the same processing as the earlier F-0005 premix is shown below in the micrographs of Figure 3. Notice, in particular, the effect of the nickel in increasing the pearlite and decreasing the ferrite contents relative to those of the F-0005 microstructure. The nickel does this in part by decreasing the carbon content of the pearlite and in part by altering the response of the phases to the process conditions. FN-0205 Figure 3: Microstructure of sintered iron with 0.5% graphite and 2% nickel. Nickel also increases the diffusivity of the iron and consequently has noticeable effects on the densification that occurs during the process as well as on the pore morphology and the degree of sinter of the resultant microstructure. Although the latter effects are not as qualitatively obvious as its effect on the pearlite content, they are easily demonstrated by quantitative means.

6 FN-0208 Figure 4: Microstructure of sintered iron with 0.8% graphite and 2% nickel. Note the nickel-rich regions. Another interesting effect of the nickel is illustrated by the micrographs of Figure 4. These show the microstructure of an iron powder premix containing 0.8% graphite and 2% nickel, (i.e. FN-0208). As the carbon content of steel increases to 0.8%, the pearlite content should increase to 100% and correspondingly, the ferrite content should decrease to 0%. However, the micrographs in this case clearly show the existence of light etching colonies, which are superficially indistinguishable from the ferrite of the earlier F-0008 microstructure. As it turns out, these are so called nickel-rich areas, i.e., the remnants of the elemental nickel due to its incomplete solution during the sintering process. These areas are also present at lower carbon contents as well but, of course, are not so obvious as here. In any case, since the evident effect of the nickel in these lower carbon steels is to increase the pearlite content, the mechanical property expectation is that it will correspondingly increase strength and decease ductility. However, as it happens, the anticipated strength increases are observed but the ductility decreases are not. Instead, the ductility either remains the same or increases slightly. This effect is generally attributed in part to the improved degree of sinter due to the nickel and in part to the presence of these nickel rich areas. The mechanical properties of the most common of the Fe-C-Ni and Fe-C-Cu premix compositions are shown in Figures 5 through 9. The sintering and microstructure of the Fe-C-Cu alloys will be discussed shortly. The data in the figures indicate the effects at two carbon contents, 0.5% and 0.8%, and of the 2% alloy level in each case. The FN and FN-0208 designations correspond to the nickel bearing compositions and the FC-0205 and FC-0208 designations likewise correspond to the copper compositions.

7 The figures are all based on the Standard 35 data compilations of the Metal Powder Industry Federation and are presented to reflect the impact of sintered density on properties, (6). Properties of the Nickel and Copper Steels For Optimized Microstructures Figure 5: Variation of the ultimate tensile strength versus sintered density. Figure 6: Variation of the compressive yield strength with sintered density.

8 Properties of the Nickel and Copper Steels Continued Figure 7: Variation of apparent hardness with sintered density. Figure 8: Variation of impact energy with sintered density -unnotched specimens. Figure 9: Rotating bending fatigue strength versus sintered density.

9 Heat Treated Properties As will be evident, these nickel and copper premix compositions generate excellent mechanical properties, which typically increase with increasing density. If the application requires even higher properties, then the parts can be submitted to a hardening treatment. This typically involves reheating to a temperature in the range of 850 to 950 o C followed by rapid cooling as by quenching in oil and a subsequent low temperature stress relief or so-called tempering treatment. The tensile properties of the nickel and copper steels resulting from this heat treatment are shown below in Figures 10 and 11. A typical quenched and tempered microstructure is shown in Figure 12. The dark etching phase in this micrograph is so-called tempered martensite and the light-etching phase is again an instance of the nickel rich areas that were previously discussed. Properties of the Nickel and Copper Steels in the Heat Treated Condition Figure 10: Variation of the ultimate tensile strength versus sintered density. Figure 11: Variation of the compressive yield strength with sintered density.

10 Quenched and Tempered FN0208 Figure 12: Microstructure of heat-treated nickel steel showing tempered martensite And nickel-rich areas. Fe-C-Cu System The Fe-Cu-C system is widely used in the P/M industry whenever hardness, higher strength, and moderate wear resistance is required. As a review of the earlier mechanical property compilations in Figures 8 through 13 will confirm, the copper steels exhibit higher strengths and hardnesses than the nickel steels at similar copper addition levels and otherwise the same carbon contents. Premixes of the Fe-C-Cu system are typically made with elemental copper additions. Since copper melts at 1084 o C and these premixes are normally sintered at temperatures in the range of 1100 to 1300 o C, their sintering behavior differs from those of the Fe-C and Fe-C-Ni systems by virtue of the presence of a liquid phase. As a general matter, sintering in the presence of a liquid phase can have profound effects. However, the effects may be positive or negative and depend on several factors. In general, the three main factors concern both the liquid and the solid phases and include: 1) their surface tension properties; 2) their relative solubilities; and 3) their volume fractions. In the case of the Fe-C-Cu premixes, neither the solubilities nor the surface tension properties that characterize them are very favorable. In addition, there are certain

11 adverse effects having to do with the appearance of the liquid and the behavior of the copper subsequently. Consequently, the contribution of the copper to the overall sinterability of these premixes is essentially negative. However, as was indicated in the earlier mechanical property compilations, the nominal copper addition is typically 2%. Thus, the volume fraction of the liquid that forms is fairly low and its effects, although generally negative, are also reasonably moderate. The details are as follows. The effects of a liquid phase during sintering depend primarily on its ability to penetrate the interparticle boundaries and secondarily on its ability to dissolve and transport the solid phase. If the liquid is able to penetrate the boundaries and is present in sufficient quantity to spread throughout the compact, then significant densification will occur in a relatively short period of time after the liquid first appears due to the effects of its capillary action in rearranging the particles in a more compact form. Thereafter, provided the solid phase is reasonably soluble in the liquid and the liquid is conversely not too soluble in the solid, the densification will continue by a process that is similar to but generally faster than the grain boundary diffusion mechanism in the solid state. In order that the liquid penetrate the grain boundaries sufficiently to effect significant densification, the dihedral angle that it makes with the solid at the point of the boundaries must be low. Dihedral angles of 5 o or less are ideal but higher angles up to about 20 o are known to be suitable. In the case of the binary Fe-Cu system, the dihedral angle is in the neighborhood of 10 o or less. However, it is also very sensitive to the presence of carbon and increases sharply as the carbon content increases. Thus, for the premix compositions under discussion, (i.e. 0.5 to 0.8%C), the angle is in the neighborhood of 35 to 45 o. As a result, the liquid penetrates the boundaries but only at their intersections, (e.g. at so-called triple points), and consequently, not sufficiently to promote much densification, (7, 8). Moreover, even if this were not the case, there is another difficulty in this system. As discussed below, the relative solubilities of the solid and liquid phases are also unfavorable. At the carbon and copper contents of interest, the phase relations indicated by the binary Fe-Cu diagram provide a reasonable approximation of the relations in the ternary Fe-C-Cu system. This diagram is shown overleaf in Figure 13, (9). After the rearrangement step or in lieu of it, the liquid phase promotes sintering by providing a medium of rapid transport for the solid. To do so, the solid must dissolve in the liquid and the greater its solubility, the better. On the other hand, as already indicated, the liquid should not be too soluble in the solid. The reason is that its dissolution will diminish its presence and hence, its sintering effect. Unfortunately, in the case of the Fe-Cu and Fe-C-Cu systems, the solubilities of the two phases are precisely the opposite of what they should be. As indicated by the phase relations in Figure 13, the copper only dissolves up to about 4.5% iron at all but the highest temperatures under discussion whereas the iron dissolves upwards of 9% copper. As a result, the

12 Figure 13: The Fe-Cu binary diagram. copper's capacity as a medium of transport for the iron is very limited and its existence as a liquid is likewise limited. In fact, as to the latter, the copper starts to dissolve in the iron even before it melts and then dissolves even faster once it does melt so that the actual liquid presence during the sintering of these compositions is short lived at best. Nevertheless, if the foregoing were the only effects of the copper, then the overall effect of the presence of the liquid phase would be innocuous or, perhaps, even somewhat beneficial. However, both the melting of the copper and its subsequent dissolution are accompanied by growth effects, which more than offset any beneficial effects that might otherwise accrue. The growth associated with the melting of the copper is thought to arise as follows. During the compaction step, some of the copper particles are completely encapsulated by the iron and some end up forming part of the internal surfaces of the pores. When the copper melts, it undergoes a substantial volume expansion. The copper particles that are part of the pore surfaces simply expand into the pores. However, the particles that were initially encapsulated force the adjacent material to expand as well and so produce the indicated growth effect. The growth associated with the dissolution of the copper is generally attributed to a difference in the diffusion characteristics of the copper and the iron. In general, copper diffuses faster in iron than visa versa. Thus, as it goes into solution, it causes growth by adding its volume to the iron while simultaneously leaving an almost equivalent volume of porosity in its wake. The latter effect is frequently manifest as small isolated pores in the iron matrix. For example, a number of such pores are evident in the high magnification micrograph of Figure 14 which otherwise shows the typical microstructure of a sintered FC-0208 composition.

13 Figure 14: Microstructure of a Fe-Cu-C sintered part. Notice the small isolated pores in the micrograph on the right which are possibly due to copper solution during sintering. Fe-P System Iron-phosphorus compositions find application in parts that require a combination of high yield strength and better than average ductility and also in parts that require improved magnetic properties. The increased strengths and superior magnetic performance of these compositions are directly traceable to the solution hardening effects of the phosphorus. The associated distortion of the iron lattice both strengthens it and increases it s resistivity. The ductility improvements are likewise traceable to the phosphorus but in this case, primarily to its effects in improving the sintering response of the iron. The yield strength and elongation properties of a series of iron-phosphorus compositions as sintered to different densities are shown overleaf in Figure 15. As a review of this figure will confirm, the yield strength is a monotonic increasing function of the phosphorus content but the elongation first decreases then increases and finally decreases again with increasing phosphorus. The yield strength variation is thought

14 Figure 15: Yield strength and elongation values of Fe-P premixes as a function of sintered density and phosphorus content. to be mostly attributable to the comparatively large solution hardening effect of the phosphorus. For example, according to findings based on microalloying theory, its effect in increasing both the yield and tensile strength of iron is upwards of 20 times that of copper, (10). In contrast, the more complex variation of the elongation property with increasing phosphorus is apparently due to a combination of three effects including: 1) enhanced sintering; 2) solution hardening; and, 3) at high phosphorus contents, the precipitation of Fe 3 P as an embrittleling grain boundary phase. Thus, as shown below in Figure 16, the Fe-P binary diagram is relevant to explain the elongation variation of these mixes as well as their sintering response, (11). Figure 16: Fe-P binary diagram. The diagram shows that the phosphorus closes the γ loop and stabilizes the α phase. Since the iron diffusivity of the α is about an order of magnitude greater than that of the

15 γ, the general effect of stabilizing this phase is to increase the sinterability of the resulting alloys. The insert on the right hand side of the diagram indicates that the knee of the loop occurs at about 1100 o C and that the γ phase is stable at this temperature for contents up to about 0.3% P. Further increases, however, destabilize it in favor of the α which eventually becomes the dominant phase at contents in excess of about 0.65% P. Thus, the greatest sintering response is typically observed at this content with higher phosphorus contents having little additional effect. The diagram also indicates the existence of a binary eutectic at 1050 o C and 10.5 % P. Since the solidus limit on the iron rich side here occurs a 2.8% P, the existence of this eutectic would ordinarily have little bearing on the behavior of the low alloy compositions that are typical of most P/M applications. However, phosphorus-containing premixes are normally made with the intermetallic compound Fe 3 P as the phosphorus source. As will be evident from the diagram, the Fe 3 P determines the solidus limit on the phosphorus rich side of the eutectic. Thus, in spite of their relatively low alloy contents, the premixes benefit from the presence of a transient liquid phase, which appears during the early stages of the sintering process. Since its appearance is, at best, very brief, the evident effect of the phase is to even out the phosphorus distribution along the interparticle boundaries of the compact and thereby create a more continuous intergranular network of the rapid sintering α phase than would otherwise be the case. In other words, its unlikely that the enhanced sintering response that is observed with these alloys is due to actual liquid phase sintering per se but rather to an effect of the transient presence of the liquid in promoting improved conditions for subsequent rapid sintering in the solid state. Lastly, the phase diagram also indicates Fe 3 P precipitation during cooling subsequent to sintering for alloys with phosphorus contents in excess of about 0.3%. The precipitates are dispersed in the α matrix and at grain boundaries. At phosphorus contents in excess of about 0.6%, the latter are sufficient to form essentially continuous films and embrittle the grain boundaries. Thus, as indicated in the earlier Figure 15, the ductility of these alloys tends to peak with increasing phosphorus at about 0.6% and decrease thereafter. Since parts of the phosphorus alloys may undergo substantial shrinkage during sintering, they are sometimes alloyed with Cu and/or C to return them to near die size. The Cu does this by virtue of the growth that accompanies its dissolution during the sintering process whereas the C does it by stabilizing the γ phase and reducing the amount of α phase sintering that occurs. Since the Cu and C additions also improve the strength of the resulting parts, several material systems containing one or both of these alloys in addition to phosphorus have been developed accordingly. Fe-B System Boron is capable of producing extraordinary effects in ferrous-based alloys. However, in P/M, these effects largely remain as untapped potentials since they are accompanied by equally extraordinary process control and/ or metallurgical problems. For example, in as-cast and wrought steels, as little as % B has been shown to double and, in

16 many cases, more than double the hardenability of virtually any ferrous base alloy, (12). However, boron s affinity for oxygen and, to a lesser extent, nitrogen and the associated inconsistencies which attend their perennial presence in iron and iron base powders have so far completely prevented its use as a practical hardenability agent in P/M applications. Similarly, Jandeska has developed an extremely effective abrasion resistant alloy for automotive applications, which takes advantage of boron s ability to combine with iron and carbon to form massive hard carbides in an otherwise relatively ductile pearlitic matrix, (13). But, here again, the necessity to sinter in pure hydrogen and to use expensive oxygen and nitrogen getters such as titanium or aluminum as well as the unfortunate further complication of having to sinter the alloy in a narrow temperature window to form the carbides and control dimensions has largely mitigated against its widespread usage. In P/M terms, probably the most spectacular attribute of boron is the potential it offers to achieve high sintered densities, (14). For instance, an addition of as little as 0.15% is sufficient to attain near full density in most compositions. Of course, once again, these increases are accompanied by the aforementioned process problems relating to oxygen, nitrogen and dimensional control. However, given the fact that the indicated density potential has been known for twenty years or more, if these were the only problems, they would have almost certainly been solved by now. Unfortunately, they are not the only problems. In particular, as explained below, it turns out that in spite of the density improvements, there are microstructual changes attending the use of the boron that prevent all but marginal improvements in the resulting mechanical properties. The present findings are based on an investigation that included studies of three alloy systems: Fe-B; Fe-B-C; and Fe-B-X where X was any one or a combination of carbon and one or two of the most common of the first and second series transition elements. Most of the latter work was with iron based prealloyed Mo as in Ast 85 HP or prealloyed iron with Mo and Cr as in various experimental compositions that were made for Högänas AB in Sweden during their development of Astaloy 3CrM. The phase relations in the Fe-B system are shown below in Figure 17, (15). According to the indications of this figure, the principal phases at all temperatures on the iron rich side of the diagram are Fe and the intermetallic compound, Fe 2 B or so-called iron boride. At high temperatures, the phase relations in the Fe-B-C and Fe-B- X systems containing C are similar. They include the iron base phase with whatever carbon is present as a dissolved solute and Fe 2 B. At room temperature, the phases are marginally different. They include the iron base phase in the essentially carbon free condition and either the so-called borocarbide phase, Fe 23 (C,B) 6 which, of course, is also an intermetallic compound or for carbon lean compositions, a mixture of the borocarbide with Fe 2 B, (16). The borocarbide phase is the product of a ternary reaction that starts at 965 o C or somewhat higher and is essentially complete at 800 o C or slightly below. In the case of the Fe-B-X system in general, the indications of observations made in this laboratory are that the phase relations are qualitatively similar to those of the Fe-B system. Accordingly, on the iron rich side of the diagram, the principal phases at all temperatures appear to be Fe with most of the X as a dissolved

17 solute and a simple substitution intermetallic that appears at high temperature as the alloy boride, (Fe,X) 2 B and at lower temperatures as nearly pure X 2 B. Figure 17: Fe-B phase diagram. The binary Fe-B alloy and, owing to the similarity in phase relations, each of the ternary alloys mentioned are classic liquid phase sintering systems. Contrary to the indications of the foregoing figure, the eutectic liquid in the binary system appears to form at about 1160 o C and anywhere from 5 to 15 o C lower in the ternary systems. In all three systems, it s dihedral angle is of the order of 10 o or less, so it both wets the iron matrix and easily penetrates the original interparticle boundaries and apparently with more difficulty or, at least, more slowly, the grain boundaries. In accordance with the indications of the figure, the boron is virtually insoluble in the solid phase, so unlike the transient nature of the liquid in the case of the Fe-Cu-C system, the liquid presence in this case is permanent or so-called, persistent. Since, as also indicated in the diagram, the liquid s iron content is upwards of 80% atomic, (i.e. approaching 97% by weight), its sintering potential in terms of ability to transfer iron is substantial. Of course, the actual amount of sintering that takes place is also dependent on how much liquid is actually present, (i.e. on its volume fraction). This, in turn, is determined mainly by the boron content and the temperature. As a general matter, increasing the boron content at a constant temperature or increasing the temperature at a constant boron content will increase the amount of liquid and thus the amount of sintering that occurs. For example, the effect of increasing the temperature at a constant boron content is shown below in Figure 18 for sintering in the Fe-B-(C, Cr, Mo) system.

18 Figure 18: Variation of density as a function of sintering temperature. The mechanical properties of these same sinterings are shown below in Figure 19. As a review of this figure will confirm, the indicated properties are unexpectedly poor in view of the high densities that were achieved. In fact, as shown by the data presented in the earlier Figure 10, equivalent or better properties are obtainable by conventional sintering of a simple FC-0208 composition. Unfortunately, the present poor response in terms of properties is typical of all three of the subject alloy systems and is in every case a consequence of the disastrous effects of the room temperature remnants of the liquid phase. More specifically, the boron latent liquid that initially penetrated the interparticle and grain boundaries and was responsible for the observed densification is likewise responsible for the poor properties because during cooling it subsequently remains in the boundaries and precipitates embrittling intermetallic borides and/or borocarbides. The only possibility to improve matters with these alloys appears to be by a spheroidizing anneal. However, preliminary trials with this objective suggested that the process is likely to be costly and the potential improvements too small to justify the effort.

19 Figure 19: Tensile properties and ductility properties of the sinterings of 1.5Cr 0.5 Mo Pre-alloy with 0.5%C, 0.4%B. To conclude, the use of boron to improve hardenability and abrasion resistance appears to have potential for future applications. Although processing difficulties tend to mitigate against its present widespread usage in these regards, its only a matter of time before the general advance of technology provides the process control remedies that are needed to overcome the problems. On the other hand, it seems unlikely that boron will play a significant role as a liquid phase sintering additive in the foreseeable future. Unless a practical expedient can be found to alter the morphology of the boron containing phase from what it presently is in the as-sintered condition, the liquid phase compositions will continue to exhibit properties that simply do not justify the effort and expense that are required to produce them. Unfortunately, based on present theory, the possibility that such an expedient will turn up any time soon appears minimal. All the systems discussed thus far use pure iron powder as a base. However prealloyed materials and diffusion-alloyed materials are also used replacing pure iron. The driver for these is to increase tensile strength, hardenability and fatigue performance. Number of alloy systems have been developed and reported at various MPIF conferences in the past years. Of particular interest are two alloy systems namely, a nickel-molybdenum- manganese prealloyed material(ancorsteel737sh) and a chromium-silicon-molybdenum (Ancorsteel4300) material. Both of these systems were developed to reduce cost of processing powder metallurgy parts. Table (2) below shows the effect of copper and graphite additions to Ancorsteel 737SH on the mechanical properties:

20 Table 2: Ancorsteel 737SH Mixes with 0.75w/oAcrawax C Sintered at 2080F with full VariCool, Tempered at 400F for 1 hour. Pressure Pressure Green Density Green Exp. Sintered Density Dimensional Change Hardness YS YS TS TS Elong. (tsi) (MPa) (g/cm³) (%) (g/cm³) (%) (HRC) (10 3 psi) (MPa) (10 3 psi) (MPa) (%) 737SH w/o Gr SH w/o Cu w/o Gr SH w/o Cu w/o Gr Microstructure developed in a 1Cu,0.7% Gr containing material is shown below. It is fully martensitic in as sintered and rapidly cooled condition. Figure 20: Microstructure of a Sample Produced from Premix #1-5 (1 w/o Copper 0.7 w/o Graphite). Etched with 1% Nital / 4% Picral. Original Magnification 500X. In the case of Ancorsteel 4300,chromium is introduced as an alloying element along with silicon without adversely affecting the compressibility.further this material can be sintered at conventional sintering temperature of 1120C to 1150C.Sintering at higher temperature than this can result in superior properties. A comparison is made of this material with a diffusion-alloyed material and a hybrid material of 0.85%molybdenum prealloyed material premixed with 2%nickel and 0.6% graphite. Ancorsteel 4300 contains lower alloying elements than FD0405 and does not require heat treatment as is necessary with FLN

21 Table 3: Summary of mechanical properties for the three alloys with conventional cooling. Table 4: Nominal compositions of the alloys. APPLICATIONS P/M parts manufactured today utilize several of the material systems described above. Listed below are some of the applications. Any specific application is dependent on design and material performance to meet the design needs. For the same application, there could be more than one material system and process to match the performance characteristics. The exact material selection is made by discussions with the powder producers, parts producers, and the end user. The list below shows typical applications, and is by no means exhaustive.

22 Table 5: Application of sintered powder mixtures. Material System Fe Fe-Cu-C Fe-Cu-C (double pressed-double sintered) Fe-Ni-C Fe-Cu-C (heat-treated) Fe-P Fe-Cu-P-C Application Magnetic clutch housings Lock-up clutch collar Magnetic applications Powder forged connecting rod Crankshaft belt drive Camshaft pulley Valve seat inserts Clutch adjustment ring Converter turbine sleeve Oil pump gears Synchronizer hubs Crankshaft belt drive Camshaft sprocket assembly Engine balance shaft components Water pump pulley Output shaft hub Parking gears Crankshaft sprocket Synchronizer hubs Synchronizer interlock sleeve Magnetic applications Bearing cap CONCLUSIONS Iron powder premixes containing carbon and either or both of nickel and copper sinter mainly by surface diffusion which produces metallurgical bonding but has no effect on dimensions. Thus, they provide the opportunity to manufacture parts to net shape with a wide range of mechanical properties. The use of phosphorus in the form of Fe 3 P provides strength and/or magnetic properties via solution hardening as well as good ductility as a result of densification due to the combined effects of a transient liquid phase and alpha phase sintering. Boron systems have potential for future applications requiring either optimum hardenability or abrasion resistance but probably not densification by liquid phase sintering. Growth of P/M will continue over the next several years. Sintering, a key technique, will be needed to reach the densities and properties required to optimize this growth. However, persistent liquid phase sintering as exemplified by the present boron compositions is not likely to be one of these techniques. Nevertheless, there are three other liquid phase processes that have yet to be thoroughly explored including in order

23 of greatest probability of near term success: 1) infiltration; 2) alloy design for transient liquid phase sintering; and, 3) supersolidus liquid phase sintering. In addition, there are also the possibilities of using either persistent or transient alpha phase sintering as well. Unfortunately, none of these processes is as easy or straightforward as simple compaction. However, compaction is limited by the pore free density of the premix, which in some cases may be quite low. Moreover, although its possible to increase the pore free density by reducing the lubricant content and/or otherwise re-designing the premix, its easily shown that even under the best of circumstances, compaction is incapable of achieving the density values that are physically possible by sintering. Thus, to realize its maximum potential, the P/M industry will inevitably need to take advantage of this key technique. ACKNOWLEDGEMENTS Tom Murphy and his group provided microstructures provided in this report. REFERENCES 1. Source Automotive Magazines Data Compiled by Eric Boreczky (Private Communication) 2. K. S. Narasimhan, Recent Advances in Ferrous Powder Metallurgy, Advanced Performance Materials 3, 7-27 (1996) 3. R. M. German, Power Metallurgy Science, Metal Powder Industry Federation, Princeton, NJ, R. M. German and K.A. D Angelo, Enhanced Sintering Treatments for Ferrous Powders, International Metals Reviews, 1984, Vol. 29, No. 4, MPIF Standard 35, Published by Metal powder Industries Federation, 105 College Road East, Princeton, NJ W.B. James, Ferrous Powder Metallurgy, Hoeganaes Corporation Training Manual. 7. R. L. Lawcock and T. J. Davies, Effect of Carbon on Dimensional and Microstructural Characteristics of Fe-Cu Compacts During Sintering, Powder Metallurgy, 1990, Vol. 33, No. 2, pp R. M. German, Liquid Phase Sintering, Plenum Press, New York and London, 1985, pp M. Hanson, Constitution of Binary Alloys, Mcgraw-Hill, New York and Toronto, 1958, pp 581.

24 10. K. J. Irvine and F.B. Pickering, Low-Carbon Steels With Ferrite-Pearlite Structures, JISI, Vol. 201, Nov. 1963, pp Metals Handbook, Volume 8, American Society for Metals, Metals Park, Ohio, USA, pp Atlas of Isothermal Transformation Diagrams, United States Steel Co. Pittsburgh, Pa. 1951, pp W. F. Jandeska Jr., Wear Resistant Iron Powder Article, U. S. Patent No. 4,678,510, Jul B. Loy and R. J. Dower, Effect of Boron on Some Properties of Sintered Iron- Carbon Alloys, International Powder Metal Conference, Associazione Italiana di Metallurgia, Florence, Italy, pp M. Hanson, Constitution of Binary Alloys, ibid., pp Metals Handbook, Volume 8, bid, pp M.C. Baran, A. H. Graham, A.B. Davala, R.J. Causton, C. Schade, A. Superior Sinter-Hardenable Material, PM 2 TEC 99, Vancouver, BC Canada, Jun 1999