RESEARCH PROJECT No. 32
IMPROVING THE PROPERTIES OF AUSTEMPERED DUCTILE IRON DIS Research Project No. 32
by R. B. Gundlach Climax Research Services
/ DUCTILE IRON \ SOCIETY
Issued by the Ductile Iron Society for the use of its Member Companies - Not for General Distribution
DUCTILE IRON SOCIETY 28938 Lorain Road North Olmsted, Ohio 44070 (440)734-8040
April 2001 RESEARCH PROJECT NUMBER 32
IMPROVING THE PROPERTIES OF AUSTEMPERED DUCTILE IRON DIS Research Project No. 32
R. B. Gundlach Climax Research Services
/ DUCTILE IRON \ SOCIETY
Issued by the Ductile Iron Society for the use of its Member Companies - Not for General Distribution
DUCTILE IRON SOCIETY 28938 Lorain Road North Olmsted, Ohio 44070 (440)734-8040
April 2001 IMPROVING THE PROPERTIES OF AUSTEMPERED DUCTILE IRON DIS Research Project No.32
R. B. Gundlach Climax Research Services March 19,2001
SUMMARY An investigation of the influence of austenite content on the properties of austempered ductile iron was undertaken. The "reacted" austenite content in the austempered structure was controlled by varying austenitizing temperature and, thus, the carbon content of the austenite during austenitizing. Ductile iron plates measuring 1.25 by 8 by 10 inches and containing 3.4C- 2.7%-0.26Mn-l.lNi-0.16Mowere austempered at 725F (385C). The properties evaluated included tensile properties, machinability and thermal expansion coefficient. Studies in a laboratory dilatometer were conducted to determine the transformation behavior of the ductile iron alloy with the intent to vary the carbon content during austenitizing. The lnvestigation included "step-austenitizing" in which the material was first heated above the upper critical temperature and subsequently cooled and held below the critical temperature. The experiments demonstrated that the carbon content of the austenite could be controlled by varying austenitizing temperature, even at temperatures below the upper critical temperature (a-transus). The stability of austenite, when cooled down into the intercritical region was found to be high, with no proeutectoid ferrite formation occurring; and specimens were successfully austempered from temperatures (1440Fl782C) well below the a-transus without pearlite or ferrite formation preceding the bainite reaction. Intercritical heat treatments were also performed with the intent to produce a mixed ferrite + ausferrite microstructure. Heat treatment in the intercritical (three-phase) region showed that an austempered structure of 50% proeutectoid ferrite + 50% ausferrite could be achieved when austempering from 1450F (788C). Austempering of the plate castings was performed at 725F (385C) from austenitizing temperatures above the critical temperature and by step-austenitizing to temperatures below the upper critical temperature. Austenite content in the austempered structure was found to decrease with decreasing austenitizing temperature. Tensile properties were found to generally increase with decreasing austenitizing temperature. Machinability in drilling was found to improve significantly as austenite content was reduced. The thermal expansion coefficient also decreased with decreasing austenite content. Austempering of the plate castings was also performed at 725F (385C) from an intercritical temperature [1450F (788C)l below the upper critical temperature. Tensile strength was reduced but tensile elongation was quite high. Machinability in drilling was found to improve significantly, comparing quite favorably with pearlitic (grade D5506) ductile iron. DIS Research Pro~ect#32
The benefits of lower final or reacted austenite content in AD1 structures were clearly demonstrated in this investigation. The results of austenitizing at temperatures near, and below, the a-transus produced lower reacted austenite contents and showed significant improvements in properties. The problem with formation of free ferrite when austempering from lower austenitizing temperatures should be addressed to better define the potential for step- austenitizing prior to austempering.
The properties obtained by austempering directly from the intercritical temperature range were also quite interesting. While the properties of this material did not meet the specifications of any of the AD1 grades, the good strength, high ductility and excellent machinability observed in this material are very encouraging and suggest that further investigations are warranted. The particular combination of properties observed for this material suggests that the material should be considered for high toughness applications. The material should be evaluated more thoroughly with a range of heat treatments (and ferritelausferrite ratios) to further develop the structure-property relationships. Fracture toughness and fatigue properties should also be evaluated to better define capabilities.
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INTRODUCTION Historical investigations in the development of austempered ductile iron (ADI) showed that, for a given austempering temperature, properties vary greatly with variation in composition. Likewise, several investigators have demonstrated that properties vary with austenitizing temperature. Variations in composition and austenitizing temperature clearly influence the austenite content of the austempered structure through their effects on carbon solubility and austenite stability. While specifications for AD1 grades, including ASTM 897 (97), specify minimum properties such as tensile strength, yield strength, elongation and hardness, AD1 properties can well exceed these minimums. Many have learned how to improve toughness and ductility through improving the quality of the base material, i.e., achieving higher nodule count and nodularity and reducing levels of microshrinkage, microsegregation, carbides and inclusions. Some corporate specifications go so far as to specify levels of some of these microstructural features for castings that are to be austempered. No one, to this investigator's knowledge, has attempted to specify (control) the final austenite content, that is, the "reacted" austenite content. Improving Strength in AD1 It seems clear that much of the variation in yield strength and tensile strength among various alloyed AD1 materials investigated is related to the volume fraction of austenite in the AD1 structure. The data in Table 1 illustrate the range of propertiqs observed in twelve ductile iron alloys austempered under the same heat treatment conditions . Note the wide range in strength; particularly yield strength, among the twelve alloys.
Table 1 Tensile Properties of Twelve AD1 Alloys with Identical Heat Treatments (Reference 1) [Austenitized 1600F (872C) + Austempered 700F (371C)]
Tensile Yield Strength Strength Percent. Alloy ksi ksi Elongation
Range 141-171 110-138 3.4-14.5 Mean 155 121 7.8 Std Dev 8.3 9.2 2.8
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Fig. I Influence of austenitizing temperature on the carbon content of the austenite matrix for 2.5%Si ductile iron alloys, after verhoeven3.
It is this investigator's contention that the variations in yield st;ength are due, in large part, to significant variations in austenite content. Work by Hayrynen showed that yield strength was governed by ferrite grain size, however, it is suspected that yielding is governed both by grain size and the amount of austenite. It would appear that less austenite leads to a higher flow stress. Preliminary examinations of data from the literature indicate that certain elements increase yield strength for a given heat treatment, while others decrease yield strength. Ductile iron alloys that contain Ni, which stabilizes austenite, have often displayed lower yield strengths than unalloyed ductile iron. One company recently inquired about obtaining AD1 with a combination of high yield strength (140 ksi) and high elongation (7%). A review of the literature indicated that five different investigators reported achieving this exceptional combination of properties in ADI. It occurred to us that yield strength could be raised significantly for a given austempering temperature, if compositio~andlor austenitizing conditions were adjusted to reduce the austenite content. As Verhoeven demonstrated in the attached Figure 1, modification of austenitizing conditions was effective in reducing dissolved carbon content. Preliminary experiments with variations in austenitizing temperature, using castings of AD1 grade 1.5, proved that yield strength could be raised by 15,000 psi with no apparent loss in ductility.
Tensile 0.2% Yield Strength Strength Elongation Red. in Condition ksi ksi % Area, % Std. Heat Trt 166.0 124.0 10.6 10.1 Modified Ht Trt 175.6 138.6 11.3 10.1
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The laboratory ex eriments listed above were quite encouraging. At the same austempering temperature, yiel8 strength continually pcreased with decreasing austenitizing temperature, yet there-was no significant change in ductility. The potential for further imp~ovementthrough additional modification of heat treatment a pears psssible. Even greater improvement might be achieved through optimization of chemica f'composition. Improving Machinability in AD1 The use of AD1 has been somewhat hampered by the difficulties associated with machining it. It has been well established by various investigators that the austenite in ausferrite structures both work-hardens and transforms to martensite during machining. Methods to overcome the machinability problems most often involve larger depths of cut in order to undercut the work- hardened layer at the machined surface of the casting. Depths of cut and feed rates can be readily increased in turning and milling operations, but quite difficult to accomplish in drilling. By reducing the amount of austenite in the austempered structure, it would seem possible that the machinability of AD1 can be improved. In this investigation, the austenite content will be widely varied and the influence of austenite content on machinability in drilling will be evaluated. Dimensional Compatibility Because the AD1 microstructure consists of a mixture of ferrite and austenite, the thermal expansion coefficient of AD1 is higher than conventional ferritic materials and generally falls between that of ferritic and austenitic steels. The question of dimensional compatibility with other ferrous materials in engine components that see temperature swings has at times created design problems that must be overcome and, as a result, there has been some resistance to the acceptance of ADI. With a reduction in austenite content, this investigation will show to what extent the thermal expansion coefficient can be reduced and thereby improve compatibility with other ferrous components. Controlling Reacted Austenite Content in AD1 Current understanding suggests that phase transformation in AD1 in the upper bainite region [temperatures above 650F (340C)], yields an acicular ferrite-austenite structure essentially free of carbide. It appears that transformation in ductile iron begins with the growth of acicular ferrite from the parent austenite without eutectoid cementite formation, but rather a stabilization of the remaining austenite. This acicular-ferrite-austenite transformation roduct has been defined as "ausferrite" in cast irons and the transformation reaction (the 18 stage reaction) has most often been described by the following equation.
where, y = parent austenite a = ferrite YHC = high-carbon or reacted austenite During transformation, the acicular ferrite platelets grow into the surrounding austenite, and the remaining austenite increases in carbon content up to 2.0 - 2.1%, coming in equilibrium with the ferrite phase. The fref energy diagram describing the reaction has been proposed by Rouns, Rundman and Moore and is depicted in Figure 2.
Page 5 of 40 0 0.5 1 .O 1.5 2.0
Weight Percent Carbon
Fig. 2 A section of the Fe-Si-Cphase diagram andpee energy diagram, at nominally 2% Si, illustrating the extension of the y + aphasejeld and its influence on the limit of carbon solubility in reacted austenite in austempering ductile iron, after Rouns, Rundman and Moore .
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Since the ferrite platelets that form at temperatures above 650F (340C) are essentially free of carbide and dissolved carbon, the carbon content of the austenite phase at the beginning of austempering directly affects the volume fraction of reacted austenite that remains at the end of transformation (1st stage). The following mass balance equations illustrate how carbon content controls the reacted austenite content and yla ratio.
where, Ci is the initial carbon content of austenite, in wt-% V, is volume fraction acicular ferrite formed, C, is the carbon content of ferrite, in wt-% V, is volume fraction reacted austenite formed, C, is the carbon content of reacted austenite, in wt-%.
In other words, an austenitizing heat treatment that produces austenite with a carbon content of 0.90% will yield an ausferrite structure with nominally 45% reacted austenite. Likewise, a heat treatment that produces a dissolved carbon content of 0.70% yields an AD1 structure with 35% austenite. Potential for Developing Low-Carbon Austenite As stated above, the carbon content in austenite is a finction of the austenitizing temperature employed in heat treatment. Solubility of carbon in austenite is governed by the phase boundary between the austenite and austenite + graphite phase fields, see the phase diagram in Figure 3. Austenite formed at a lower temperature in the y + Graphite phase field has a lower carbon content and, in turn, produces lower reacted austenite volume fractions in ADI, i.e. a lower yla ratio. It would appear that the lower limit of dissolved C in austenite occurs by heating just above the a-transus temperature, see Figure 3. Based on the phase diagram, however, the dissolved carbon content could be further reduced, under metastable conditions, by a further reduction in temperature below the a-transus temperature. Providing proeutectoid ferrite does not nucleate and grow, cooling into the 3-phase field (austenite + ferrite + graphite phase field), after first heating in the austenite + graphite phase field, should result in austenite of an even lower carbon content. It is proposed, therefore, that "step austenitizing" can produce low-carbon (<0.7%) austenite prior to austempering; that is, by first heating into the austenite + graphite region (above the a- transus) and then cooling down and holding in the three-phase region (y + a + G phase field). There is only a finite window of time that one can hold an alloy in the three-phase region before any significant amount of ferrite will form. IT and CCT diagrams for ductile iron indicate, however, that proeutectoid ferrite formation is quite sluggish above 1400F (8 16C), see Figures 4 and 5.
Page 7 of 40 LIQUID
a~ *I
m m m I; 0 1 2 3 4 % CARBON
Fig. 3 Schematic of Fe-2%Si-C phase diagram illustrating the limit of carbon solubility in austenite, when austenitizing above the upper critical (a-transus) temperature. Solubility follows the boundary line A-B.
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Element t Si Mn Me Ole 36 Z,C O,( 0,OI
5 10 15 20 30 40 60 90 120 sec 34 601015B30 lei! (Time) -+ min 10 loo Zeit (Time) -+ ain r-piTiJ Fig. 4 Isothermal transformation (TTT) Fig. 5 Continuous cooling transformation (CCT) diagram for a low-alloyed ductile diagram of a low-alloyed ductile iron, after iron, after Rohrig and ~airhurs?. Climax Molybdenum ~om~an~~.
The limit to how low ductile iron alloys can be cooled and held in the intercritical, three-phase region before ferrite forms is dependent on the stability of austenite and its resistance to proeutectoid ferrite formation. The kinetics of ferrite formation in this temperature range has not been well investigated. Furthermore, the effects of holding austenite in this temperature range on the solubility of carbon have, heretofore, not been considered, particularly in regard to its affect on ausferrite transformation and the yla ratio. To investigate the effects of austenitizing temperature and step austenitizing on (a) dissolved carbon content in austenite and (b) the yla ratio after austempering, transformation studies were conducted in a quenching dilatometer. The stability of austenite held in the intercritical region was also investigated. Subsequent to the transformation study, experimental austempering heat treatments were performed at varying austenitizing conditions (varying dissolved carbon contents) to determine the effect on mechanical properties and machinability. Ferritic-Ausferritic Ductile Iron Recent unpublished reports describe ductile iron parts that have been directly heat treated in the intercritical region and subsequently austempered. Intercritical heat treatment produces a duplex structure consisting of pro-eutectoid (polygonal) ferrite plus an ausferrite (ADI) structure, which is reportedly more ductile and has improved machinability. This approach to heat treatment was also investigated here. Dilatometer experiments were performed to determine the intercritical temperature that produces nominally 50% ferrite + 50% austenite. Subsequently, test plates were intercritically heat-treated and austempered to determine the mechanical properties and machinability of this modified form of ADI.
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EXPERIMENTAL PROCEDURES and RESULTS For this investigation, cast plates measuring 1.25 by 8 by 10 inches were produced from an alloy with moderate hardenability. The experimental procedure consisted of two phases of study -- Phases I and 11. These experiments were designed to confirm the validity of the hypotheses, and establish the relationships proposed between reacted austenite content, AD1 properties and machinability. Phase I consisted of phase transformation studies of a commercial AD1 alloy, using a quenching dilatometer, to determine the optimum austempering conditions for developing microstructures with widely varying yla ratios. Dilatometer studies were performed to determine the following: a. The stability of austenite (resistance to proeutectoid ferrite formation) in the intercritical region and the potential for step-austenitizing, b. Whether austenite carbon content can be reduced by step-austenitizing. c. The time needed to reach equilibrium carbon contents as a function of austenitizing temperature above and below the alpha-transus, d. The influence of austenitizing temperature on austenite carbon content through measurements of the Ms temperature and volume changes (% transformation) during austempering, e. The intercritical temperature that produces 50% ferrite + 50% austenite. Phase 11 consisted of determining the mechanical properties and machinability of the same commercial AD1 alloy for several austenitizing conditions that produce variations in austenite carbon contents and reacted austenite volume fractions. These experiments were designed to determine: a. The tensile properties and hardness of one alloy in various heat treatment conditions yielding various yla ratios, b. The austenite contents of the various heat treatment conditions of step a, c. The machinability (in drilling) of AD1 plates in three different heat treatment conditions (high, mid and low yla ratios), d. The machinability (in drilling) of one intercritically heat-treated AD1 plate having a mixed structure of proeutectoid ferrite and ausferrite, e. The microstructures of the various heat treatment conditions of step a, f. The correlation of mechanical properties with microstructure, g. The thermal expansion coefficient for each yla ratio.
EXPERIMENTAL ALLOY PLATES Ten plates measuring 1.25 by 8 by 10 inches were cast by a DIS member foundry in a moderately-alloyed ductile iron chemistry. The alloy included Ni, Cu and Mo to provide sufficient hardenability for the heat treatment of 1-114 inch thick plate sections. The chemical composition of the plates was determined by various methods. Carbon and sulfur were determined by combustometric methods, Mg and Cu were determined by atomic absorption, and all other elements were determined by optical emission spectroscopy using a re-melted and chill cast button. The results of chemical analysis are given in Table 2. 1 A metallographic sample was removed from one of the plates for examination. The plate was shown to have acceptable nodularity and nodule count, and to be largely ferritic with some pearlite. Representative micrographs of the graphite structure and metallic matrix are shown in Figure 6.
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Table 2 Chemical Composition of Experimental Test Plates
Element C
(a) as-polished (Mag. =I 003 (b) etched in 2% nital (Mag. =250X)
Fig. 6 As-cast microstructure in 1.25 by 8 by 10 inch plate castings, DIS Research Project #32
Cold .lunctjon C'ompensator
Fig. 7 Photograph of (a) Formastor quenching dilatometer and (b) schematic illustrating the specimen, specimen holder and L VDT transducer.
DILATOMETRY A Formastor quenching dilatometer was used in the phase transformation study. A photograph of the dilatometer is shown in Figure 7a. A schematic diagram illustrating the specimen, specimen holder and LVDT dilatation transducer are shown in Figure 7b. The dilatometer used a 118-inch diameter by 0.4-inch long test specimen, with a 0.080-inch diameter by 0.10-inch deep hole drilled in one end. Several dilatometer specimens were machined from the 1.25-inch plates. In the dilatometer, the test specimen was heated in vacuum by induction (1 MHz). A Type R thermocouple was spot-welded to the bottom of the hole in the specimen for monitoring and controlling temperature. The thermal cycle was controlled by a computer, and the change in length and temperature were monitored continuously throughout the thermal cycle. The dilatation curves were generated on a computer, plotting temperature versus change-in- length. From these curves any transformations that produced a volume change, such as graphitization, ausferrite transformation and martensitic transformation were readily detected; they are manifested by an inflection in the T vs. AL curve, i.e. the point where a deviation from linear thermal contraction occurs. Critical Temperatures The initial thermal cycle performed consisted of relatively rapid heating to 600C, followed by heating at a constant rate of 2 deg. Clminute to 900C to determine the critical temperatures, Al and a-transus, of the experimental alloy. The critical temperatures were detected as inflections in the change-in-length versus temperature curve, which is plotted in Figure 8 from the data generated by the dilatometer. The A1and a-transus for this alloy were determined to be 1370F (743C) and 1503F (8 17C), respectively.
Pase 12 of 40 DIS Research Project #32 Temperature, C
800 1000 1200 1400 1600 1800 Temperature, F
Fig. 8 Plot of the change-in-length versus temperature curve generated by the dilatometer. The critical temperatures are detected as injlections in the curve -- deviations in linear thermal expansion. Upper critical temperature = 1503F (81 7C), and lower critical temperature = 13 70F (743C).
I I I I I
-
1600F (872C) end -
- -
I
- -
I I I I I 0 2000 4000 6000 8000 10000 Time, seconds
Fig. 9 Plot of the change-in-length versus temperature during austenitizing. While being heated at 1600F (872C), the sample grew due to expansion of the austenite lattice as carbon (graphite) dissolved in the metallic matrix.
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Measure of Austenite Stability Several tests were performed using a variety of austenitizing and quenching cycles to determine the stability of austenite in the three-phase (intercritical) region, i.e. below the alpha-transus temperature and above the A1 temperature. Step-austenitizing was employed in which a sample was initially heated above the a-transus to 1600F (872C) and held for one hour, then cooled to a lower temperature. While being heated at 1600F (872C), the sample grew due to expansion of the austenite lattice as additional carbon from the graphite was dissolved, see Figure 9. The sample expanded at temperature for about 1 hour before reaching equilibrium. Subsequently, the samples were rapidly cooled to sequentially lower temperatures below the a- transus (into the intercritical region) and held at temperature for up to six hours to allow for transformation to ferrite. Samples were cooled to, and held at, temperatures as low as 1440F (782C). While the samples were cooled to the intermediate austenitizing temperature, they contracted due to thermal contraction. While the samples were held at the intermediate austenitizing temperature, they were observed to contract further, presumably as a result of the reduction in lattice spacing as carbon left the matrix and precipitated onto the existing graphite nodules. The contraction eventually stopped and, generally after one hour, the samples had contracted to within 90% of the total observed contraction, see Figure 10. Even after holding for six hours, no transformation to ferrite was detected at any intercritical temperature down to 1440F (782C).
Time, seconds Fig. 10 Plot of the change-in-length versus temperature during "step-austenitizing". While being heated at the intermediate austenitizing temperature, the samples contracted, presumably as a result of the reduction in lattice spacing as carbon left the matrix.
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Temperature, C Temperature, C
0 500 1000 1500 Temperature, F Temperature, F Fig. lla Plot of the change-in-length versus Fig. llb Plots of the martensite transformation temperature upon quenchingfrom austenitizing. for specimens quenchedj-om several austenitizing The martensite reaction and the martensite "start" temperatures. The plots illustrate the continuous temperature were recorded. rise in Ms Temperature with decreasing austenitizing temperature. Martensite Transformation Each sample was then quenched to ambient temperature to record the martensite "start" transformation temperature, see an example of the length versus temperature plot in Figure 1 1a. Several of the experimental thermal cycles employed are plotted in Figure 1 1b and illustrate the continuous rise in Ms temperature with decreasing hold (austenitizing) temperature down to 1440F (782C), which was well below the a-transus temperature 1503F (8 17C). The change in Ms temperature with austenitizing temperature reflected the decrease in austenite carbon content with austenitizing temperature and is plotted in Figure 12.
Austenitizing Temp., C
Austenitizing Temp., F Fig. 12 Correlation of martensite "start" temperature with austenitizing temperature. DIS Research Project #32
Ausferrite Transformation Additional samples were austempered following austenitization by single or step-austenitizing cycles. The samples were austenitized at 1700F, 1600F, 1500F and 1440F (927C, 872C, 8 16C and 782C), then austempered at 725F (385C) for up to two hours. The typical dilatation curves for the single and step-austenitizing cycles are given in Figure 13. The transformation at 725F (385C) was recorded for each austenitizing condition and plotted in Figure 14. From the plots it was seen that lower austenitizing temperatures resulted in greater volume changes at 725F (385C), indicating that more transformation took place and less reacted austenite remained. Note that as austenitizing temperature was reduced, the transformation times became shorter, also indicating lower carbon content in the austenite and less hardenability. Representative microstructures of selected samples are shown in Figure 15. Intercritical Austempering Experiments Additional dilatometer specimens were heated in the intercritical region to determine the temperature that would produce 50% austenite + 50% ferrite. The samples were heated directly to selected temperatures in the intercritical region between 1400 and 1475F (760 and 802C) without first exceeding the a-transus temperature. The samples were held at temperature for two hours and quenched to ambient temperature. Subsequently, the samples were metallographically prepared and analyzed to determine the amount of proeutectoid ferrite that had formed. A representative photomicrograph is shown in Figure 16. The amount of ferrite was determined by automated image analysis and plotted as a function of temperature in Figure 16b. It was determined that the desired structure of 50% proeutectoid ferrite was obtained when heating at 1450F (788C).
- Step-Austenitizing --- 1600F (872C) Austenitizing .
/------3 . / I .I I - . I I I I . I I . I I I I - I I _ I I I _____--___-___ 1, I / I I I,' I I/ I I I I
Time, seconds
Fig. 13 Change-in-length versus temperature curves for the single and step-austenitizing cycles. To better illustrate the differences, the curves were shifted in the "length" scale to avoid overlap.
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-0.2 1 10 100 1000 10000 100000 Time, seconds
Fig. 14 Change-in-length versus austempering time at 725F (385C). Lower austenitizing temperatures resulted in greater dimensional changes during the ausferrite reaction.
(a) Austemperedfrom 1700F (927C) (b) Austemperedfrom 1500F (816C) Fig. 15 Microstructures of specimens austempered at 725F (385C)from various austenitizing temperatures. (Mag. =5003
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Fig. 16a Microstructure of dilatometer specimen heated and quenchedfrom the intercritical temperature range (1460F/793C), illustrating proeutectoid ferrite in martensite. (Mag. =500JJ
lntercritical Temp., C
lntercritical Temp., F
Fig. 166 Proeutectoid ferrite content versus temperature in the intercritical region. Ferrite content represents the @action of ferrite in the metallic matrix. DIS Research Project #32
AUSTEMPERING TRIALS Machining blanks were cut from the some of the plate castings to obtain tensile test specimens. The test blanks were turned into round bars measuring about 718 inches diameter. Three samples each were then austempered according to the following eight schedules.
Austenitizing Austenitizing Austempering Cycle Temperature Process Temperature 1 1700Fl1 hr Single 725Fl1.5h 2 1600Fl1.5hr Single 725F11.5h 3 1500Fl2 hr Step 725Fl1.5h 4 1450Fl2 hr Step 725F11.5h 5 1700Fl1 hr Single 675Fl1.5h 6 1600Fl1.5 hr Single 675F11.5h 7 1500Fl2 hr Step 675F11.5h 8 1450Fl2 hr Step 675F11.5h
Step-austenitizing consisted of heating at 1600F (872C) for 1.5 hours prior to heating up to, or cooling down to, the designated austenitizing temperature. The test blanks were heated in an electric muffle furnace and subsequently quenched in a salt bath at the desired austempering temperature and held at the austempering temperature for 1.5 hours. Tensile Properties After removing the salt by washing in water, the test blanks were machined into tensile specimens with 0.50 inch diameter by 2.25 inch long reduced gauge sections. Tensile testing was conducted at room temperature according to ASTM E 8 recommended practice. The test results which include 0.2% offset yield strength, tensile strength and percent elongation are presented in Table 3. Austenite Measurements Metallographic samples were removed from the broken tensile specimens for examination of the microstructure and for the evaluation of austenite content. Austenite content for the series of bars austempered at 725F (385C) was determined by x-ray diffraction at an outside laboratory after an electrolytic polish. Austenite contents varied from 26% to 44% and the correlation of austenite with austenitizing temperature is shown in Table 4 and Figure 17. Representative micrographs of the microstructures in various heat treat conditions are shown in Figure 18a through 18d. Analysis for Dissolved Carbon Content Solid samples were removed from the broken tensile test specimens for determining combined carbon content, that is, carbon present as carbides or as dissolved carbon in the metallic matrix. The analysis for dissolved carbon consists of measuring total carbon and graphitic carbon. Both analyses are performed by a combustion method. Graphitic carbon is determined by dissolving a small (2 gram) solid sample in acid, filtering off the graphite, and analyzing the graphite content. Dissolved carbon is determined by difference. The dissolved carbon content of the various samples is presented in Table 5.
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Table 3 Tensile Properties of Austempered Test Bars at Different Austenitizing Cycles and Two Austempering Temperatures
Tensile 0.2% Yield Specimen Strength Strength Elongation Red.in ID ksi (MPa) ks i (MPa) % Area,% Austempered at 725F (385C)
Austempered at 675F (358C)
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Table 4 Final Austenite Content of Heat Treated Tensile Bars as Determined by X-ray Diffraction
Reacted Austenitizing Austenite Temperature Content
Austenitizing Temp., C
700 800 900 1000 50
Austenitizing Temp., F
Fig. 17 Correlation of reacted austenite in the austempered structure with austenitizing temperature.
Table 5 Dissolved Carbon Content of Heat Treated Plates
Austenitizing Total Graphitic Dissolved Temperature -Carbon -Carbon -Carbon 1700F 3.38% 2.75% 0.63% 1575F 3.38% 2.77% 0.61% 1450F 3.38% 2.91% 0.47% 1440F (IC) 3.38% 2.97% 0.41%
IC = Intercritical heat treatment DIS Research Project #32
(a) Austenitize 1700F -Amtemper 725F (b) Austenitize 1600F - Austemper 725F
(c) Austenitize 1500F -Amtemper 725F (d) Austenitize 1440F - Austemper 725F Fig. 18 Microstructures in bars, heat treated after removing from the plates. (Mag. ~5004
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MACHINABILITY STUDY Subsequent to the initial laboratory trials, whole plates were austempered at four conditions for conducting machinability trials. Six plates were shipped to a DIS member foundry for heat treatment. Two of the plates were split in half to produce sufficient material for the study. One whole plate and one half plate were heat treated by each of the following four conditions.
Austenitizing Austenitizing Austempering Cycle Temperature Process Temperature 1 1700Fl1 hr Single 725F11.5h 2 1575Fl2 hr Step 725F11.5h 3 1450Fl2 hr Step 725Fh 4 1450Fl2 hr Intercritical 725F11.5h
The whole plates were then shipped to Machining Research Inc., Florence, KY for machinability trials. Machinability was determined by drilling. Drilling Tests The evaluation of machinability was performed at Machining Research Inc. using a drilling test with high-speed steel (HSS) drills. The tests were conducted on a Tongil TNV-80 CNC three- axis machining center. The machine was equipped with a 24 hp drive motor with a spindle speed range of 80-8000 rpm. The programmable cutting feed rate ranged from 0.039 to 158 inches per minute. The feed rate for rapid traverse positioning was 600 inches per minute. This machine had a 30 position automatic tool changer. A Fanuc Model OM CNC controller provided accurate position and repeatability along with precise cutting speed and feed. The size of the machinability test plates was approximately 10" x 8" x 1.25". The plates were milled on the top and bottom faces to remove any heat treatment scale and provide a smooth, flat surface for drilling. They were also milled on the four sides to provide accurate positioning in the vise. Tool Materials -- The drilling tests on all four AD1 plates were initially performed using Cleveland list #2001 high speed steel, jobbers length drills. The drill diameter was 114 inch, with a point angle of 1 18 degrees, a helix angle of 29 degrees, relief angles of 12 degrees, and a plain point. Since the 1700F (927C) AD1 material was so difficult to machine, a premium class of drills, Cleveland list #2013/817 was required to obtain any tool life. These drills were made of cobalt (8%) HSS and were also coated with titanium nitride (TiN). The point angle was 135 degrees and the point style was split, or crankshaft. These drills were used only for the 1700F (927C) AD1 material. Cutting Fluid -- The cutting fluid used in these tests was a commercially available product, TRIM SOL, produced by the Master Chemical Co. The fluid was diluted in the ratio of one part concentrate to twenty parts tap water. The cutting fluid was delivered to the drill in a generous supply by two flexible pipes. The hole depth for all the tests was 0.5 inches, blind. The feed rate in all the drilling tests was 0.005 inch per revolution. The cutting speeds, as listed in Table 6 were varied from 20 feet per minute to 350 feet per minute. The drilling tests were terminated when the wear on the corners of the drills reached 0.01 5" or when the drill failed completely. At higher cutting speeds, it was
Pane 23 of 40 not unusual for a drill to fail with less than 0.01 5" on the corners, but in the process of failure the corners were completely rounded and worn off. A schematic illustrating the measurement of drill wear is shown in Figure 19. Drill life was recorded as number of holes per drill and is presented in Table 6 and plotted against cutting speed in Figure 20.
DRlLL WEAR
VEWA usom wm
VllEAR
UARW w ww VIEW A
Fig. 19 Schematic of drill bit illustrating measurement of tool wear.
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Table 6 Tool Life in Drilling versus Cutting Speed of Austempered Plates
Cutting TOOL LIFE, # of Holes Drilled Sveed. fpm 1700F 1575F 1450F 20 60 -- -- 25 50 -- -- 30 135 -- -- 35 100 -- -- 40 40 -- --
0 50 100 150 200 250 300 350 Cutting Speed - FPM
Fig.20 Drill life versus cutting speed in machinability study. Drill life represents the number of holes per drill to 0.015 inches wear at the corners of the drill.
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Further Tensile Testing The "half' plates were sectioned to produce additional tensile test specimens. Three tensile bars were machined from each half plate. The test blanks were machined into tensile specimens with 0.50 inch diameter by 2.25 inch long reduced gauge sections. Tensile testing was conducted at room temperature according to ASTM E 8 recommended practice. The test results, which include 0.2% offset yield strength, tensile strength and percent elongation, are presented in Table 7. Once again, metallographic samples were removed fiom the broken tensile specimens for examination of the structure. Representative micrographs of the microstructures in various heat treat conditions are shown in Figure 2 1.
Table 7 Tensile Properties of Test Bars from Plates Austempered at 725F (385C) from Different Austenitizing Cycles
Tensile 0.2% Yield Specimen Strength Strength Elongation Red. in ID ksi (MPa) ksi (MPa) % Area. % 725F (385C) Austempering:
AD1 Grades: ASTM 897 - 90 (Reapproved 1997) Grade 125/80/10 125 860 80 550 10 Grade 150/100/7 150 1030 100 690 7
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(a) Austenitized at1 700F
(b) Austenitized at 1575F (c) Austenitized at 1450F
Fig. 21 Microstructures in the broken tensile specimens removedfrom the heat treatedplates for the various heat treat conditions. (Mag. =SOOX)
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Thermal Expansion Coefficient The thermal expansion coefficient was determined for the experimental alloy in four austempered conditions and when austempered from four austenitizing temperatures (single and step- austenitizing). The test specimens were removed from the shoulders of the tensile bars and measured 0.25 inches diameter by 1.OO inch long. The samples were heated in an argon atmosphere to 500°C (930°F) at a rate of 3OC per minute. The change in length was recorded with temperature throughout the heating cycle, and from this record the average coefficient of thermal expansion was determined over the temperature range 40 to 200C. This testing was performed by an outside laboratory. The average thermal expansion coefficient for the four heat treated conditions are presented in Table 8.
Table 8 Thermal Expansion Coefficient after Austempering From Different Austenitizing Temperatures
Thermal Expansion Coefficient Specimen x 106pc x i O~PF ID J40 - 200C) J 104 - 392F')
Ferritic Ductile Pearlitic Ductile
DISCUSSION It has been hypothesized that there is a strong relationship between the yla ratio in AD1 and mechanical and physical properties. This project was initiated to investigate the controls needed to reduce the volume fraction of reacted austenite in austempered ductile iron and to determine its influence on properties. Austempered ductile iron with such a microstructure has been shown to develop higher yield and tensile strengths without loss in ductility. In addition to increased strength, other anticipated benefits of lower reacted austenite volume fraction are improved machinability and lower thermal expansion coefficient. It is believed that the carbon content of austenite during austenitizing is primary in determining the final reacted austenite content of the austempered structure. Therefore, the controls needed to reduce the dissolved carbon content at quenching and, thereby, reduce the volume fraction of reacted austenite in the austempered structure were examined. A heat treatment procedure (described as "step-austenitizing) was employed to extend austenitizing to temperatures below the upper critical (a-transus) temperature. Step-austenitizing consists of initially heating above the a-transus in the austenite+graphite phase field, then cooling down to, and holding at temperatures below the a-transus.
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Fig. 22 Microstructure in dilatometer specimen quenchedfiom temperature below the a-transus. A martensitic microstructure with no ferrite was observed in the quenched specimen. (Mag.=500JJ Austenite Stability The stability of austenite below the upper critical (a-transus) temperature was investigated by isothermal transformation studies in the dilatometer and by metallographic examinations. The a- transus was determined to be 1503F (8 17C). In dilatometer experiments where step- austenitizing was employed, no ferrite formation was detected during isothermal holding at temperatures down to 1440F (782C) for up to 6 hours. Subsequent to the 6-hour hold below the a-transus, the dilatometer specimens were quenched to promote martensitic transformation. No ferrite was observed in metallographic examination of the quenched specimens, see Figure 22. These findings clearly demonstrate that austenite is quite stable at temperatures well into the intercritical temperature range (in the 3-phase field). Note that the extended stability of the austenite in the 3-phase region would only be exhibited if the material was completely austenitic prior to heating in the intercritical region. Varying the Reacted Austenite Content (yla ratio) The principal goal of this investigation was to learn how-to reduce the yla ratio, that is, reduce the amount of reacted austenite in the ausferrite microstructure. Austenitizing heat treatments near and even below the a-transus were attempted. An approach used to reduce the austenite content was to develop heat treatment techniques (step-austenitizing) that allow austenitizing below the a-transus temperature. Direct proof that austenite (yla ratio) could be varied was found in the X-ray diffraction measurements, see Figure 17, where the austenite content was shown to decrease from 44% to 25% with a decrease in austenitizing temperature from 1700F to 1440F. Additional demonstration that the yla ratio was decreased was observed in the increase in the amount of acicular ferrite formed during transformation at the austempering temperature, see Figure 14. As the austenitizing temperature was decreased, the amount of growth that occurred during transformation at 725F (385C) increased. Transformation of austenite to ferrite produced
Page 29 of 40 DIS Research Project #32 a volume increase, due to the decrease in the packing factor of the unit cell, and thus a greater volume change indicated an increase in ferrite formation and a reduction in the amount of reacted austenite in the ausferrite structure. Controlling Austenite Carbon Content To reduce the reacted austenite content in the austempered structure, heat treatment procedures were employed that reduced the carbon content of the starting austenite. A strong indication that the carbon content of the austenite was reduced, by lowering the austenitizing temperature, was shown by the shift in Ms temperature to higher temperatures as austenitizing temperature was decreased, see Figure 12. The influence of carbon on the Ms temperature in ferrous alloys has been examined by numerous investigators. ~ldis~and ~ndrews~ are two among many who have published equations for calculating Ms temperature from chemical compositions of ferrous alloys. The equations that these investigators developed are as follows.
Based on these two equations the Ms temperature rises nominally 407 deg. C with a decrease of 1% carbon. Considering the range of Ms temperatures (190C to 275C) observed in the dilatometer experiments, one could predict that the carbon content in the austenite decreased 0.21% when decreasing the austenitizing temperature from 1700F to 1440F (927C to 782C). Furthermore, the Ms temperature (carbon content) was shown to continue to decrease as temperature was decreased below the a-transus temperature. The change in carbon content of the austenite with change in austenitizing temperature was also illustrated by the change in dimension of the dilatometer specimen in step-austenitizing, see Figure 23. When austenitizing was raised from 1600F to 1700F (872 to 927C), the specimen increased in length. When austenitizing temperature was decreased from 1600F to 1500F or 1450F (872 to 8 16 to 788C), the specimen decreased in length, see also Figure 10. The changes in dimension are attributed to the changes in lattice parameter as the solubility of carbon in austenite increases or decreases with temperature. Analysis of dissolved carbon content gives equally, if not more, compelling evidence that austenite carbon content was controllable. In the samples austempered by the DIS member foundry at 1700F, 1575F and 1450F (927, 857 and 788C), the dissolved carbon content decreased from 0.63% to 0.47% over a 250F (140C) temperature range, see Table 7. However, this change in carbon content (a reduction of 0.16%) was less than predicted by the Fe-Si-C phase diagram. In the diagram presented by Nieswaag and ~ijhofg,see Figure 24, austenite carbon content is shown to decrease 0.1% for each 72OF (40°C) decrease in austenitizing temperature. verhoeven3 also found that austenite carbon content decreased approximately 0.1 % for each 74°F (41OC) decrease in austenitizing temperature, see Figure 1. There appears to be a substantial difference in the alloys, however, because the dissolved carbon content in Verhoeven's alloy was nominally 1.O%C at 1700F (927C), whereas for the alloy of this investigation it was 0.63%C. The alloying elements are known to cause a significant shift in the phase boundary between the y and y + graphite phase fields, as shown for silicon in Figure 24. The alloy of this investigation also contains significant levels of Ni and Cu, both of which are expected to shift the boundary line to the left, i.e., to lower carbon contents.
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DIS Research Project #32
Influence of Austenite on Strength Yield strength and tensile strength increased with decreasing austenitizing temperature (austenite content) to 1500F (816C), see Tables 3 and 7. The data of Tables 3 and 7 are combined in Figure 25 for ease of comparison. This effect was seen with austempering at both 725F and 675F (385C and 357C). Yield strength and tensile strength reached 115 and 146 ksi, respectively, with austempering at 725F (385C), and 126 and 158 ksi with austempering at 675F (357C).
Tensile elongations were oftentimes quite acceptable for ASTM 897 grades 1 and 2, which are 10% and 7% respectively. However, at times elongations were unacceptably low, particularly at the lower austempering temperature. The occurrence of low elongations has been attributed to poor graphite structures that were sometimes encountered in the test bars. Several of the test bars displayed low nodularity, containing substantial degenerate graphite and vermicular graphite forms, see Figure 26. The deleterious effect of such graphite forms on tensile elongation in AD1 is substantial and quite well known and, therefore, the affect of various treatments on ductility could not be clearly determined by this study. Below 1500F (8 16C), both yield strength and tensile strength fall off. The loss in strength at the lower temperature [1450F (788C)l was attributed to a change in the microstructure to a mixed structure consisting of fiee ferrite and ausferrite, see Figure 27. The dilatometer experiments showed that the polygonal ferrite was not proeutectoid ferrite resulting fiom holding at the
Austenitizing Temp., C
TS in Bars o TS in Plates A YS in Bars aA YS in Plates
Austenitizing Temp., F Fig. 25 Yield strength and tensile strength as a function of austenitizing temperature in the heat treated plates and separately heat treated test bars.
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Fig. 26 Microstructures in tensile bars displaying low nodularity and containing substantial degenerate graphite and vermicular graphite forms. (Mag.=I 003 DIS Research Project #32
Fig. 27 Microstructure in plate casting austempered after step-austenitizing at 1450F (788C), displaying a mixed structure consisting ofpeeferrite and ausferrite. (Mag. =5OOY)
(a) Intercritical Austemper fiom 1450F (b) Intercritical Austemper @om 1450F (1OOX mag.) (500X mag.) Fig. 28 Microstructure in plate casting austempered directly porn the intercritical temperature 1450F (788C), illustrating ferrite fractions well above 50%
Page 34 of 40 DIS Research Project #32 intercritical temperature. Rather, it is believed that, with the reduction in carbon content at the low austenitizing temperatures, there was a significant decrease in hardenability, particularly with respect to ferrite formation. In other words, the ferrite nose had been shifted to the left as austenitizing temperature was decreased. Further experiments with alloys having greater hardenability and improved graphite structures should be conducted. Intercritical Heat Treatment In general, the strength of the intercritically heat treated ductile iron plate was substantially lower than the fully austenitized plates, see Table 6. The high ferrite content of the intercritically heat treated structure significantly reduced the strength and hardness. At 83 ksi and 57 ksi, the tensile and yield strengths meet the properties of grade D5506. However, the tensile elongations, at 19%, were quite high compared to conventional D5506. This high elongation appears extraordinary, especially considering the less than optimum graphite structure in these plates. Hardness was similarly low at 229 HB. The intercritical temperature that produced 50% ferrite - 50% austenite was determined to be 1450F (788C) for this alloy. Nevertheless, the austempered microstructure contained substantially more than 50% ferrite, see Figure 28. It is quite likely that additional ferrite forms on cooling to the austempering temperature due to the lack of hardenability when quenching from the intercritical region. Since proeutectoid ferrite already existed at the time of the quench, there was no need for ferrite nucleation to occur prior to ferrite growth. Consequently during quenching, without the delay associated with ferrite nucleation, additional ferrite readily formed on the existing ferrite grains and the kinetics of ferrite formation were accelerated. These findings suggest that a 50% ferrite - 50% ausferrite structure can only be achieved by austenitizing at higher temperatures in the intercritical region. Alternatively, the composition of the alloy could be more heavily alloyed to delay the formation of proeutectoid ferrite. Austenite and Machinability The results of the drilling studies were positive and demonstrated that reduction in austenite content and the yla ratio were beneficial. The drill life curves of Figure 20 show that drilling life increased significantly with decreasing austenitizing temperature. The machinability of the plate austempered from 1700F (927C) was so poor that the conventional high-speed steel drill could not be used. An alternate drill bit (high-speed steel with 8% Co) having significantly greater wear resistance had to be used to machine that plate. Even with the differences in drill bits, the plate austempered from 1575F (857C) gave at least a ten-fold increase in tool life at similar cutting speeds, see Table 6. With further reduction in austenitizing temperature to 1450F (788C), tool life increased multiple times. There was also a significant further improvement when machining the intercritically austempered material. However, the reader should be reminded that that material is largely ferritic and has significantly lower mechanical properties when compared with the plates austempered from a fully austenitic condition. For an equivalent tool life of 400 holes, the cutting speed increased progressively with decreasing austenitizing temperature (and austenite content), see Figure 20 and the table below. 80-55-06 65-45-12 Austenitizing Tem erature Pearlitic Ferritic ---1700F 1575F 1450% IC-1450F DI DI Hardness, HB 32 1 3 02 302 229 187-223 156- 163 Cutting <20 7 8 124 260 190-215 325-435 Speed, fpm DIS Research Project #32
The results of this study were compared with those of DIS Research Project 20 "Machinability of Ductile Iron Castings", in which the machinability of ferritic and pearlitic ductile irons were evaluated in drilling, turning and milling. Two materials each of two ductile iron grades were evaluated in that study - ASTM grades 80-55-06 and 65-45-12. Figure 29 combines the tool life curves from both studies. While the ferritic irons displayed good tool life at very high cutting speeds, the pearlitic irons showed poorer tool life compared with intercritically austempered ductile iron, even though they have comparable hardness, see the table above.
0 50 100 150 200 250 300 350 400 450 500 Cutting Speed - FPM
Fig. 29 Comparison of drill life versus cutting speed for austempered ductile iron of this investigation andferritic andpearlitic ductile irons of DIS Research Project 20. Drill life represents the number of holes per drill to 0.015 inches wear at the corners of the drill.
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Austenite and Coefficient of Thermal Expansion The results of the thermal expansion coefficient measurements clearly showed that as austenite content was reduced, so was the coefficient, see Table 8. The benefit of a lower thermal expansion rate is an improved compatibility with other ferritic components. The thermal expansion coefficient for specimens austempered at 725F (385C) appeared to have decreased 20% with a decrease in austenitizing temperature from 1700F to 1440F (927 to 782C). The correlation of thermal expansion coefficient with austenitizing temperature is shown in Figure 30.
Fig. 30 Correlation of thermal expansion coeflcient with austenitizing temperature for samples austempered at 725F (385C). DIS Research Project #32
CONCLUSIONS Methods of varying the austenite content or the yla ratio in austempered ductile iron were investigated using variations in austenitizing heat treatment to control austenite carbon content during austenitizing. Carbon content was varied by varying austenitizing temperature, both above and below the upper critical temperature (1503F18 17C). Step-austenitizing by heating above the upper critical and subsequently cooling into the intercritical region allowed austenitizing at lower temperatures. When austempering at 725F (385C), both mechanical properties and machinability improved with decreasing final or reacted austenite content in the austempered structure. The transformation behavior of this 3.4C-2.7%-1.1Ni-0.2Mo ductile iron was investigated, particularly with regard to the potential for step-austenitizing, and a number of characteristics were revealed. 1. The ability to reduce austenite carbon content during austenitizing was verified by increases in the measured Ms temperature and volume increases during austempering, i.e., increases in acicular ferrite formation. 2. In step-austenitizing, equilibrium carbon contents were achieved within about 1 hour after cooling down to and holding at the intercritical temperature. 3. The stability of the austenite (resistance to proeutectoid ferrite formation) in the intercritical region was high. No ferrite formation occurred, even after holding for 6 hours at temperatures as low as 1440F (782C). 4. In dilatometer experiments, no free ferrite was formed in quenching to the austempering temperature or during austempering. Based on the dilatometer study, austempering experiments were performed with the full plate castings. Tensile properties, machinability by drilling, austenite content, and thermal expansion coefficient were determined. 1. Reacted austenite content of the austempered structure decreased with decreasing austenitizing and step-austenitizing temperature. Reacted austenite content was reduced from 44% to 26%. 2. Tensile properties generally increased with decreasing austenite content and the yla ratio. 3. The machinability (in drilling) improved dramatically with decreasing reacted austenite content. AD1 plates in three different heat treatment conditions (above and below the upper critical) showed increases in cutting speeds. For a drill life of 400 holes, optimum cutting speeds increased from <20, 78 and 124 fpm, respectively with decreases in austenitizing temperature. 4. While no free ferrite was formed in step-austenitizing in dilatometer experiments, some amounts of proeutectoid ferrite formed in the plates austempered from step-austenitizing. 5. The tensile properties of the intercritically heat treated AD1 plate were extraordinary in that they displayed very high ductility for the observed strength, having 19% elongation with a yield strength of 57 ksi. The exceptional combination of properties was attributed to the mixed microstructure, which consists of proeutectoid ferrite and ausferrite. 6. The intercritically heat treated AD1 plate also displayed good machinability in drilling. For a drill life of 400 holes, cutting speeds of 260 fpm were achieved, comparing quite favorably with 190 to 215 Qm for grade 80-55-06 ductile iron.
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ACKNOWLEDGEMENTS
The research conducted for this project was funded by the Ductile Iron Society and the Nickel Development Institute. The generous contribution of materials and heat treatment activities of Intermet Corporation are gratefully acknowledged. A number of individuals also volunteered their time and resources to bring this project to successful completion, and their contributions are sincerely appreciated. They include the members of the DIS Research Committee, and the project monitoring committee. The constructive comments and encouragement of some are particularly appreciated, including A1 Alagarsamy, Bob Christ, A1 Druschitz, Kathy Hayrynen, Jim Mullins and Phil Seaton.
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REFERENCES 1. Report no.GRI-89/00 17, Gas Research Institute, (1 989). 2. Hayrynen, K.L. et al, "Austempered Ductile Iron Part 11: Microstructure and Tensile Properties of Low Alloy Small Section Size Castings", 1991 World Conference on Austempered Ductile Iron, Bloomingdale, IL, March 12-14, 1991, American Foundry Society. 3. Verhoeven, J.D., "A Study of Austempered Ductile Iron", The Physical Metallurgy of Cast Iron -- Proceedings of the Third International Symposium on the Physical Metallurgy of Cast Iron, Stockholm, Sweden, August 29-3 1, 1984. Edited H. Fredriksson and M. Hillert, Elsevier Science Publishing Co., Inc. 4. Rouns, T.N., Rundman, K.B. and Moore, D.M., "On the Structure and Properties of Austempered Ductile Cast Iron," AFS Trans. vol92, 1984. 5. Rohrig, K. and Fairhurst, W. Heat Treatment of Nodular Cast Iron, Climax Molybdenum Company, Dusseldorf, 1979. 6. McArdle, G. D., Continuous Cooling Transformation Diagrams for Ductile Iron, Climax Molybdenum Company, 1980. 7. Eldis, G.T., "A Critical Review of Data Sources for Isothermal Transformation and Continuous Cooling Transformation Diagrams," The Metallurgical Society of AIME, Warrendale, PA, 1978. 8. Andrews, K.W., "Empirical Formulae for the Calculation of Some Transformation Temperatures," JISI, vol203, 1965, p. 721-727. 9. Nieswaag, H. and Nijhof, J.W., "Influence of Silicon on Bainite Transformation in Ductile Iron; Relation to Mechanical Properties,", The Physical Metallurgy of Cast Iron -- Proceedings of the Third International Symposium on the Physical Metallurgy of Cast Iron, Stockholm, Sweden, August 29-3 1, 1984. Edited H. Fredriksson and M. Hillert, Elsevier Science Publishing Co., Inc.
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