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Transmission electron microscopy study of austempered nodular iron: Influence of silicon content, austenitizing time and austempering temperature
Materials Science and Engineering, 96 (1987) 231-245
231
Transmission Electron Microscopy Study of Austempered Nodular Iron: Influence of Silicon Content, Austenitizing Time and Austempering Temperature VJEKOSLAV FRANETOVIC, MICHAEL M. SHEA and EDWARD F. RYNTZ Metallurgy Department, General Motors Research Laboratories, Warren, MI 48090 (U.S.A.) (Received January 16, 1987; in revised form May 20, 1987)
ABSTRACT
The structures produced by austempering nodular iron were evaluated by transmission electron microscopy to explain the influence o f austempering temperature, austenitizing time and silicon content on the impact properties. No carbide precipitation was evident in a 2.88 wt. % Si nodular iron austenitized for 15 min at 950 °C and austempered for 60 min at 370°C. This iron had the highest impact strength. Austempering at 340 and 400 °C resulted in the formation o f transition carbides in bainitic ferrite laths and a lower impact strength. In addition, both a lower silicon content and a longer austenitizing time promoted carbide formation in the bainitic ferrite during austempering at 3 70 °C.
1. INTRODUCTION
Recent investigations [1, 2] have shown the outstanding mechanical properties produced by austempering nodular iron castings. However, the properties achievable by austempering can be divided into two groups: one consisting of high tensile and fatigue strength obtained with lower bainitic matrix structures and the other consisting of high ductility and impact strength obtained with upper bainitic matrix structures containing a large amount of untransformed austenite. Recent work [3] showed that a combination of high hardness and high impact strength can be produced in nodular iron by a unique austempering heat treatment. This heat treatment, consisting of a short austenitizing time, generally 15 min at 950 °C, followed by austempering at 370 °C for 1 h, produced a Brinell 0025-5416/87/$3.50
hardness of 310 HB and a room temperature impact strength of 170 J in a 2.88 wt.% Si nodular iron. The austempered microstructure consisted of upper bainite and retained austenite with a small amount of lower bainite. When the austenitizing time was increased to 75 rain or the silicon content was reduced to 1.57 wt.%, the impact strength decreased significantly at all austempering temperatures (Fig. 1). However, the maximum impact strength was achieved by austempering at 370 °C for each of the austenitizing conditions. Metallographic examination revealed that the superior properties produced by the 370 °C austemper were not the result of a significantly higher retained austenite content since a similar amount of austenite was present after austempering at 400 °C. Rather, the results indicated that austempering at 370 °C may produce a unique bainite structure and/or enhanced austenite stability. In nodular iron, similar to high silicon steels [4-9], the bainite reaction is believed to involve the nucleation and growth of bainitic ferrite unaccompanied by carbide formation. Consequently, during growth of the ferrite, the remaining austenite becomes increasingly enriched in carbon. Recent X-ray diffraction [10] and electron microprobe [3] results and, as will be shown later, the results of this study have confirmed the increase in austenite carbon content during austempering at 370 °C to form upper bainite. In addition, the electron microprobe study showed that the carbon content of the austenite decreased at extended austempering times (120 min), supporting the belief that decomposition of the carbon~nriched austenite ultimately occurs, similar to the behavior of high silicon steels [ 11 ]. Furthermore, the results showed that at 370 °C a bay region © Elsevier Sequoia/Printed in The Netherlands
232 180
stabilizing the austenite. However, the a m o u n t of retained austenite in lower bainite is known [8] to be less than in upper bainite, consistent with less enrichment in carbon of the austenite. In the present investigation the bainitic structures formed by austenitizing a 2.88 wt.% Si iron for 15 min and 75 min at 950 °C followed by austempering at 340 °C (lower bainite), 370 °C and 400 °C (upper bainite) for 60 min were evaluated by TEM in an attempt to explain the difference in impact properties. Of particular interest was the presence of carbides in the bainitic ferrite, in the retained austenite or at the ferrite-austenite interface. Also, the influence of silicon content on the structures produced during austempering at 370 °C for 60 min was evaluated to provide a better understanding of the effect of this factor on the structure and properties of austempered nodular iron.
160
140
120 w
L. @
100
m
ebu
a
Q.
E
80
60
m
m
2. PROCEDURE 40
m
20
I
340
I
370
I
400
Austempering Temperature (°C)
Fig. 1. Influence of austempering temperature, austenitizing time at 950 °C, and silicon content on the impact strength of nodular cast iron: e, 15 rain, 2.88 wt.% Si; m, 75 rain, 2.88 wt.% Si; A, 15 min, 1.57 wt.% Si.
occurs in the isothermal transformation diagram indicating higher stability of the austenite to bainite transformation. The occurrence of carbide-free bainite in nodular iron may only be applicable to transformation to upper bainite. The results of a recent transmission electron microscopy (TEM) investigation [12] on nodular iron {2.33 wt.% Si), austempered at 240 °C for 2 h to produce a lower bainite, indicated that cementite (FeaC) formed in the bainitic ferrite although no carbides were observed in the retained austenite. Although cementite precipitation occurred, an increase in the austenite carbon content presumably still t o o k place,
2.1. Casting process and heat treatment Nodular iron was prepared by treating 270 kg of iron in a ladle with 0.15 wt.% Mg added as a magnesium ferrosilicon alloy containing 5.5 wt.% Mg. The treated iron was inoculated with 0.50 wt.% Si using ferrosilicon (50 wt.% Si) during transfer to a pouring ladle. Keel block castings were poured into oilbonded dry sand molds. The composition of the irons is shown in Table 1. The b o t t o m section of the keel blocks was completely annealed prior to machining unnotched Charpy impact specimens. The annealing treatment, consisting of a 12 h hold at 930 °C followed by a slow cool of 55 °C h -1 in the furnace to 650 °C and air cooling to room temperature, produced a completely ferritic matrix structure although some intercellular carbides were present. The bars were copper plated to prevent decarburization during the austenitizing heat treatment. The copper plate was removed prior to impact testing. Austenitizing was performed in a resistance-heated vertical retort furnace with an atmosphere of nitrogen (1.7 m 3 h -1 ) and methane (0.4 m 3 h-l). The heat treatment consisted of austenitizing for 15 or 75 min at 950 0C followed by quenching in a nitrate-nitrite salt bath to austempering temperatures of 340, 370 or 400 °C. The
233 TABLE 1 Nodular iron compositions
Heat
A m o u n t ( wt.% )
C
equivalent
CN 2 9 6 5 CN 2 9 6 6
C
Si
Mn
Mg
S
P
(wt.%)
3.97 3.51
1.57 2.88
0.49 0.54
0.059 0.053
0.010 0.012
0.03 --
4.49 4.46
specimens were held in the salt for 60 min and air cooled to room temperature. After impact testing, TEM specimens were prepared according to the procedures below.
2.2. X-ray diffraction The retained austenite content of all specimens was determined by analysis of the X-ray diffraction spectra produced by Mo K s radiation using the procedure developed b y Miller [ 13 ]. The specimen holder provided b o t h rotating and rocking m o t i o n to minimize the effect of preferred orientation of the austenite. The volume fraction o f austenite was determined using the integrated intensities of the austenite 220 and 311 peaks and the ferrite 211 peak.
2.3. Transmission electron microscopy sample preparation The procedure developed for preparation of nodular iron foils suitable for viewing with the TEM consisted of three major steps: (a) mechanical polishing of nodular iron wafers to approximately 0.13 mm thickness or less; (b) electrochemical polishing to approximately 0.03 mm thickness or less; (c) ion milling to produce perforation on the thin edges of the specimen.
2.4. Mechanical polishing The starting point for the foil preparation was a 10 mm × 10 mm square slice 0.18 mm thick from an impact bar o f the desired material. The slice was obtained b y cutting the bar using a low speed saw with a 0.305 mm diamond wafering blade and a 152 gf load. Further mechanical thinning to provide a uniform sample 0.13 mm thick was accomplished b y mounting the slice on a 16 mm × 50 mm × 50 mm brass holding block that had a slot 0.13 mm deep and 12 mm wide ground on the center on one face. The slice was
m o u n t e d in this slot using an electrical conducting mixture of Duco cement, powdered graphite and enough acetone to form a thin even layer. A vinyl separator sheet, approximately 13 mm thick, was placed on t o p of the slice and another metal block was placed on top of the vinyl. The entire assembly was clamped together and allowed to dry. Following drying and disassembly, the slice was mechanically polished flush with the holding block using 3 2 0 , 4 0 0 and 600 grit silicon carbide paper. The specimens 3 mm in diameter were obtained from the slice by electrodischarge machining (EDM) using a brass electrode. The adhesive used to m o u n t the foil to the holding block allowed for electrical continuity for the EDM and prevented the cut specimen from being washed away b y the EDM cooling oil. The cut samples were separated from the holding block b y soaking in acetone.
2.5. Electropolishing Electropolishing was performed in an electrolyte consisting of 75 g of CrO3, 400 ml of glacial acetic acid and 21 ml of distilled H20, at r o o m temperature using a Fischione twin-jet electropolishing unit. The o p t i m u m electropolishing condition occurred at a current density of 18 mA c m - 2 . Although the time of polishing varied with specimen thickness, approximately 15 min was sufficient to produce minimal perforation. After the specimens had been po lished, they were rinsed successively in acetic acid and methanol.
2.6. Ion milling This technique is widely used to thin nonmetallic and metallic materials which cannot be uniformly etched electrochemically because of large differences in the electrochemical properties of the microconstituents. With ion milling, it was possible to produce perforations
234
having edges sufficiently thin to allow TEM examination. Thinning was accomplished with an argon ion energy of 1.5-2.0 kV. The ion thinning was extremely slow, requiring approximately 20 h at an initial angle of 20 ° between the ion beam and specimen and an additional 6 h at an angle of 7°-10 ° . The shallow finishing angle was necessary to obtain satisfactory TEM foils.
3. AUSTEMPERED MICROSTRUCTURES
3.1. Optical microscopy Optical micrographs o f the specimens evaluated are shown in Fig. 2. For the 2.88 wt.% Si iron austenitized at 950 °C for 15 min, the bainite formed by austempering at 340 °C (Fig. 2(a)) appeared to be lower bainite while upper bainite was primarily formed at 370 °C (Fig. 2(b)) and 400 °C (Fig. 2(c)). Increasing the austenitizing time to 75 min or decreasing the silicon content to 1.57 wt.% resulted in no significant change in the microstructure produced b y austempering at 370 °C. X-ray diffraction results indicated that the retained austenite content o f the 370 °C austempered samples was in the range 4 0 - 4 4 wt.%.
4. TRANSMISSION ELECTRON MICROSCOPY
The TEM results are presented in the following order: (1) 2.88 wt.% Si nodular iron austenitized at 950 °C for 15 min (sample 1, austempered at 370 °C (highest impact strength); sample 2, austempered at 400 °C; sample 3, austempered at 340 °C); (2) 2.88 wt.% Si nodular iron austenitized at 950 °C for 75 min (sample 4, austempered at 370 °C); (3) 1.57 wt.% Si nodular iron austenitized at 950 °C for 15 rain (sample 5, austempered at 370 °C). Samples 1-3 were selected to show the influence of austempering temperature on the microstructure of the 2.88 wt.% Si nodular iron austenitized at 950 °C for 15 min. Sample 4 was selected to reveal the effect of a longer austenitizing time (75 min) on the structure of the 2.88 wt.% Si iron. Finally, sample 5 was selected to reveal the influence of silicon content (1.57 wt.% Si compared with 2.88 wt.% Si).
Fig. 2. Microstructures of 2.88 wt.% Si nodular iron austenitized for 15 rain at 950 °C and austempered for 60 rain at (a) 340 °C, (b) 370 °C and (c) 400 °C (Nital etch). (Magnification, 400×.)
4.1. Sample 1 To characterize the austempered structure by TEM and to determine whether carbides had formed, it was essential initially to identify the ferrite and austenite phases. For that purpose, numerous micrographs and selected area diffraction (SAD) patterns corresponding to the ferrite and austenite phases with different zone axis were taken and analyzed. Figure 3 represents a typical TEM micrograph of sample 1 in which the ferrite and austenite are labeled b y F and A respectively.
235
Fig. 3. Transmission electron micrograph of sample 1.
Both phases show contrast due to a high density of dislocations. No contrast due to precipitates was revealed during careful tilting of the specimen and no additional reflections or other diffraction effects due to precipitates were f o u n d on the corresponding SAD patterns. Therefore, for this stage o f austempering, both phases were free of carbides. In addition, no carbide formation was observed at interphase boundaries between the two phases. The measured interplanar spacing dhk~ obtained for ferrite matched the published data [14] calculated on the basis of the lattice parameter a = 2.866 A well. An austenite lattice parameter a of 3.621 A was calculated using the measured interplanar spacings for retained austenite. Using the relationship between the austenite lattice parameter and carbon content given in ref. 15, the carbon content of the austenite was calculated to be 1.5 wt.%. The calculated value is in agreement with the results of a microprobe carbon determination for the austenite in the same sample [3]. The very high carbon content found in sample 1 can be correlated with the absence of carbide precipitation for this austempering temperature. In addition, the high austenite carbon content promotes stabilization of the austenite and prevents significant transformation to martensite during cooling to room temperature. The orientation relationship between ferrite and austenite was determined from SAD patterns taken at the interphase b o u n d a r y between these ~hases. Figure 4(a) is a bright
Fig. 4. (a) Bright field micrograph showing the ferriteaustenite interphase boundary; (b), (c) SAD patterns taken from areas A and F showing the N-W orientation relationship.
field micrograph showing ferrite at F and austenite at A and the associated interphase boundary. Several partly overlapping areas containing the interphase boundary were chosen and tilting experiments performed to obtain satisfactory SAD patterns for determination of the crystallographic relationship between the phases. Figure 4(b) is an SAD
236 pattern obtained mainly from austenite (primary diffraction pattern) but also contains ferrite reflections. The austenite orientations is close to [211]fcc. The superimposed diffraction pattern of ferrite indicates that the ferrite orientation is close to [110]bc¢. This zone axis, [110]b¢~ of ferrite, is clearly recognizable and indexed on Fig. 4(c) which was obtained b y slightly moving the specimen to orient the ferrite in a strong diffraction condition. Additional spots in Fig. 4(b) are due to double diffraction which appeared because of symmetrical orientation between the two phases. Two reflections due to double diffraction are indicated by arrows to distinguish them from the overlapping austenite (i 11) and ferrite ( i 1 0 ) reflections. The SAD patterns are correctly rotated to show the exact crystallographic orientation between the phases in the corresponding bright field micrograph. The correspondence of reflections to the observed phases was checked successively b y the dark field technique. The orientation relationship between ferrite and austenite was determined to be the Nishiyama-Wasserman orientation, which for the particular example is [211]~c~// [110]b¢c, (022)~¢~//(000.)bc¢ and (i11)~¢// (ilO)b~. The appearance of a small amount of an additional phase was evident in the sample. The structure of this phase was f.c.c, with a measured lattice parameter of 4.300 A. This phase belongs to the group of MC-type f.c.c. carbides where, generally, M -= Ti, Ta, Nb, Hf, Th or Zr [16]. According to measured interplanar spacings, structure determination and possible impurities in the investigated alloy, it was concluded that the phase was TiC. The measured interplanar spacings dhm of the identified phase are in very good agreement with that published for TiC [16]. This comparison is shown in Table 2. TiC adjacent to austenite is shown in Figs. 5(a) and 5(b) which are bright field and dark field micrographs respectively. The dark field micrograph was obtained b y using the (111)~c reflection of TiC. Also shown in Fig. 5(b) is the corresponding SAD pattern of the area shown in Fig. 5(a) which is in the foil orientation [011]fee. The presence of TiC in the iron was expected since titanium is a normal residual element in the Sorel metal used for the charge material. However, the TiC is highly stable in austenite and, consequently, would not dissolve during
TABLE 2 Comparison between the mean values of measured dhk ! with that of TiC [16] dhkl, measured (A) Sample 1
Sample 2
2.47 2.16 1.509 1.301 1.236 1.085
2.50 2.14 1.523 1.295 1.236 1.070 0.987
~.980
dhkl, TiC (A)
hkl
2.48 2.15 1.520 1.297 1.241 1.075 0.986
111 200 022 311 222 400 331
Fig. 5. TiC phase in sample 1 shown in (a) bright field and (b) dark field micrographs. The inset in (b) is the SAD pattern taken from the area shown in (a). austenitization and re-form during austempering. All specimens used in the previous mechanical property evaluation and this TEM study should contain TiC.
237
4.2. Sample 2 Ferrite and austenite were identified using the procedure described for sample 1. Analysis of the micrographs and SAD patterns indicated that precipitation had occurred in both phases. Unfortunately, neither precipitate gave satisfactory diffraction patterns to enable unambiguous identification. Fine precipitates revealed by the fact that the bend contour (A) passes over the austenite grain can be seen in Fig. 6(a), while Fig. 6(b) is the corresponding SAD pattern with the zone axis [011 ]~c¢. Additional strong reflections are from an adjacent austenite grain (top of Fig. 6(a)) and the faint reflections are probably due to the precipitates. The fine precipitates in bainitic ferrite are shown in Fig. 7(a) which is a dark field micrograph obtained from a (ll0)b~¢ diffuse reflection. The observed diffraction effect due to the fine precipitates was diffuse spikes, which can be seen around fundamental spots in Fig. 7(b) with the spike directions close to (102).
In Fig. 7(b) the foil orientation is [001]bcc. Such images can be attributed to interstitial carbon atom clusters. Similar diffraction effects have been observed during the preliminary stage of tempering of carbon steel [17]. Clusters appear in martensitic carbon steel at low tempering temperatures, but no tetragonality which would indicate tetragonal bainitic ferrite (or martensite) was detected in these SAD patterns. Although 400 °C appears to be too high a temperature for the interstitial carbon clusters to form (they formed at up to 100 °C in carbon steel), it can be assumed that the high silicon content moved the normal sequence of carbide formation to higher tem-
Fig. 6. (a) Transmission e l e c t r o n micrograph of austenite in sample 2; ( b ) t h e corresponding SAD pattern.
Fig. 7. (a) Dark field micrograph of ferrite in sample 2; (b) the corresponding SAD pattern.
238
peratures. In addition, no interphase carbide was observed in this specimen. The interplanar spacings obtained for ferrite were in very good agreement with calculated values using a lattice parameter a of 2.866 A [14]. However, the calculated austenite lattice parameter a of 3.586 A was significantly smaller than for sample 1 and represents a carbon content in the austenite of approximately 0.68 wt.%. On the basis of the lower austenite carbon content in sample 2 than in sample 1, the precipitates observed in sample 2 are believed to be carbides. The carbide precipitation probably occurred because of the higher austempering temperature (400 °C) compared with that for sample 1 (370 °C). Evidence for precipitation, which was not observed in sample 1, was clear in sample 2 and was confirmed b y additional precipitate reflections (Fig. 6(b)). This finding can be correlated with the lower impact strength of sample 2 than that of sample 1.
Since the carbide precipitation in austenite was observed on a very fine scale, it is believed that the precipitation reaction did not occur to an extent which would destabilize the austenite completely and result in decomposition to ferrite during further austempering or to martensite during cooling.
4.3. Sample 3 In sample 3, no contrast or diffraction effects due to precipitates could be found in the austenite. Austenite (labeled A) free of precipitation and with a high density of dislocations is shown at the t o p of the micrograph in Fig. 8(a). Calculation of the austenite lattice parameter indicated a carbon content similar to that of sample 2. However, there was indication that the start of precipitation occurred in the ferrite as shown in Figs. 8(a) and 8(b). Fine precipitates were observed in the ferrite (labeled F) in both the bright field micrograph (Fig. 8(a)) and the dark field micrograph (Fig. 8(b)) using a (211)bc¢ matrix reflection. No
Fig. 8. (a) Electron micrograph of sample 3 showing austenite at A, TiC at C and ferrite at F; (b) TiC illuminated by its (200) reflection; (c) ferrite illuminated by its (211) reflection.
239
satisfactory diffraction pattern could be obtained from these microstructural features. TiC with f.c.c, structure was also found in this sample. TiC (marked C) is oriented obliquely in the bright field micrograph (Fig. 8(a)) and dark field micrograph (Fig. 8(c)) where it is illuminated by its (9.00)fcc reflection. Additional micrographs of sample 3 (Fig. 9(a)) showed the presence of ferrite at F, a B2 structure phase, austenite at A and a carbide phase at C. Although the structure determination of the carbide phase was not completed, the interplanar spacings of the phase, obtained from different foil orientations, matched very well with that of x-carbide [18]. This comparison is shown in Table 3. The same type of carbide but with a different morphology was detected in an Fe-l.22wt.%C alloy [19] ' tempered at temperatures similar to the austempering temperatures of this study. This carbide phase is shown in the bright and dark field micrographs and the corresponding SAD
TABLE 3 Comparison between the mean values of measured
dhk ! with that of X-carbide [18]
dhkz, measured
(A) 2.06
2.01
1.81
1.508 1.364 1.339 1.315
1.277 1.266
dhkl, FesC2
hkl
2.06 2.03 1.98 1.91 1.80 1.76 1.72 1.67 1.62 1.57 1.50 1.37 1.34 1.32 1.27 1.25
510,021 315,402 511,211 221 312,511 402 421 512 602 113 422 331 331 802 531 712
(A)
Fig. 9. (a) Transmission electron micrograph of sample 3 showing austenite at A, ferrite at F and x-carbide at C; (b) dark field micrograph of x-carbide obtained from spot C; (c), (d) SAD patterns corresponding to area C in (a).
240
pattern in Figs. 9(a)-9(c). The dark field micrograph was made using the spot marked C in Fig. 9(c). Also, superlattice spots belonging to a B2 structure can be seen (indicated by an arrow) on the SAD pattern in Fig. 9(d). This diffraction pattern was obtained by tilting the foil from the position in Fig. 9(c). Hence, spot 2 in Fig. 9(c) corresponds to spot 2 in Fig. 9(d), but the zone axis of the carbide is different in Figs. 9(c) and 9(d). Indexing of the diffraction pattern (Fig. 9(d)) is shown in Fig. 10, where three separate diffraction patterns are identified: ferrite with the zone axis []33], a B2 structure with the zone axis [001] and an unidentified zone axis of monoclinic x-carbide. The observed superlattice reflections in the diffraction patterns of sample 3 cannot currently be explained. Superlattice spots observed for tempered martensite were attributed by Izotov and Utevsky [20] to the precipitation of tetragonal carbide, Fe4C. However, on the basis of the identification of X-carbide in sample 3, the presence of Fe4C seems unlikely.
4.4. Sample 4 After identification of the ferrite and austenite phases by SAD, each phase was examined for the presence of carbide precipitation. Figure 11 is a typical bright field micrograph showing highly dislocated retained austenite free of precipitation in the foil orientation [013]tcc. In addition, no diffraction effects due to the presence of precipitates were observed during subsequent examination. Careful evaluation of the ferrite phase indicated two modifications of the structure: ferrite free of precipitates and ferrite containing coherent carbide precipitates with two different morphologies, which both had symmetrical orientations with the ferrite. The bright field micrograph and dark field micrograph in Figs. 12(a) and 12(b) respectively show isolated ferrite with a high density o f dislocations but with no evidence of precipitation. The dark field micrograph was obtained by ( 0 1 i ) ferrite reflection. In the corresponding SAD pattern (Fig. 12(c)) the zone axis for the ferrite was [i11]bc¢. No tetragonal distortion of ferrite could be detected within the error of measuring the diffraction pattern. The measured interplanar spacings agreed well
J
5 - ~- ___-__ _
!
II0
I I
-
X I..,=X,.~_.
o~o
t tTo
301
310
Fig. 10. Indexed diffraction pattern shown in Fig. 9(d): , [133]bcc;- - - , [ O 0 1 ] o r d e r e d b c c with superlattice spots (×); -----, zone axis not identified, probably monoclinic X-carbide.
Fig. 11. Transmission electron micrograph of sample 4, showing precipitate-free austenite with a high density of dislocations.
with those of b.c.c, ferrite calculated using a = 2.865 A. Carbide precipitation was found in the ferrite in the region of the specimen containing alternate laths of ferrite and austenite (Fig. 13(a)). These carbides were similar to the e-carbide observed b y Sandvik [9] in an Fe-Si-C alloy austempered in the temperature range 2 9 0 - 3 8 0 °C. However, in the present
241
O
Fig. 13. (a) Transmission electron micrograph of sample 4, showing ferrite laths containing ~7-carbide; (b) SAD pattern showing the orientation relationship between 7?-carbide, ferrite and austenite.
Fig. 12. Transmission electron micrographs of sample 4: (a) precipitate-free ferrite; (b) ferrite in a dark field taken using the ferrite (011) reflection; (c) SAD pattern corresponding to the area shown in (a).
study the carbide was identified as ~?-carbide. This t y p e o f unstable transition carbide is usually formed during the early stages o f tempering of martensite. The structure of the carbide was determined to be orthorhombic with the lattice parameters a -- 4.704 A, b -4.318 A and c = 2.830 A [21, 22]. This structure belongs to the space group P n n m and the carbon atoms regularly occupy half the octahedral interstices [17, 22]. Since 7?-carbide
can be present simultaneously with e-carbide during tempering of martensite [ 21] and all first-order reflections from e-carbide are covered by ~?-carbide reflections, identification o f the ~?-carbide was difficult. However, carefully measured angles between corresponding g vectors, which were indexed according to orthorhombic unit cells, agreed very well with theoretical values calculated using a = 4.704/~, b = 4.318 A and c = 2.830 A [22]. The interplanar spacings of the carbide were measured and matched with the calculated values for 1?-carbide (Table 4). The mutual orientation relationships between ferrite, austenite and ~?-carbide were determined using an SAD pattern taken from the ferrite-austenite interface (Fig. 13(b)). A symmetrical orientation relationship was found between the phases. In Fig. 13(b), ferrite was
242 TABLE 4 Comparison between the measured dhM with that of ~?-carbide [4]
dhM, measured
dhM, ~7-carb id e
hkl
4.69 3.15 2.86 2.44 2.39 1.81 1.65
4.70 3.18 2.83 2.42 2.37 1.809 1.668
100 110 001 101 011 201 211
(h)
(h)
indexed as [012]bcc and this zone axis is parallel to the [i11]fcc zone axis of austenite on the superimposed austenite diffraction pattern. Faint reflections due to 7?-carbide were indexed according to the orthorhombic structure with the zone axis [010]. This zone axis is parallel to the zone axis of both ferrite and austenite. Additional reflections indicated by the arrow appeared because of double diffraction due to symmetrical orientation between the phases. The above results indicate that (001) of ~?-carbide is parallel to (001) of ferrite and ( 0 i l ) of austenite. The relationship between ferrite and austenite is consistent with the Nishiyama-Wasserman orientation relationship between b.c.c, and f.c.c, phases. The simultaneous relationship between 7?-carbide, ferrite and austenite has not been previously reported. Similar 7?-carbide has been observed in carbon steel [17, 22], nickel steel [23], manganese steel [24] and Si-Mn steel [21]. According to Jacks [25], the loss of tetragonality of martensite can be attributed to the formation o f a transitional carbide which he identified as e-carbide, although more recent studies [19, 22, 23, 26] indicated that the carbide was orthorhombic 7?-carbide. This is consistent with the present findings of no tetragonal distortion in the ferrite phase. Another type of carbide was observed in the ferrite as elongated images with the axis close to [001]be c. Figures 14(a) and 14(b) respectively are a bright field micrograph of ferrite and a dark field micrograph of an ordered carbide phase taken with the ½ ~ ½ superlattice reflection. According to the SAD pattern (Fig. 14(c)) which represents the
Fig. 14. Transmission electron micrographs of sample 4: (a) ferrite; (b) ferrite in a dark field taken using the 31 231 3 superlattice reflection; (c) SAD pattern corresponding to the area shown in (a).
[ 2 i 0 ] zone of ferrite, the carbide can be described as a long-range-ordered phase. The variation in diffracted intensity due to the ordered structure is clearly shown by the superlattice reflections marked S in Fig. 14(c). 4.5. Sample 5 A typical electron micrograph of this sample (Fig. 15) indicated that the retained austenite (A) was free of carbide precipitation. How-
243
Fig. 15. Transmission electron micrograph of sample 5 showing carbide-free austenite at A and two types of carbide in ferrite at B and C. The inset dark field micrograph is from the carbides at B.
ever, two types of precipitate were observed in ferrite at B and C. The precipitates at B are shown in the dark field inset taken by (121) orthorhombic reflection in Fig. 15. The diffraction pattern o f the precipitates was indexed as an orthorhombic structure with a = 4.524 A, b = 5.088 • and c = 6.741 A, which corresponds to cementite [14]. The measured angles between the g vectors corresponded to those calculated for orthorhombic cementite. In addition, the measured interplanar spacings matched those for cementite [14]. The morphology of the precipitates at C is similar to that of the carbides observed by Sandvik [9] in austempered high silicon steel. Analysis of diffraction patterns indicated t h a t the precipitates were ~7-carbide. A typical diffraction pattern parallel to the [ 3 i 0 ] zone axis of ferrite contained faint q-carbide reflec-
tions and double
5. DISCUSSION On the basis of these TEM results, the higher impact strength produced in the 2.88 wt.% Si iron by austenitizing for 15 min at 950 °C
244
followed by austempering at 370 °C for 60 min compared with 340 or 400 °C, appears to be related to the absence of carbide formation and the resulting high stability of the retained austenite due to enrichment with carbon. At the present time, it is unclear w h y carbide formation was not observed in either ferrite or retained austenite for austempering at 370 °C. However, the results of other investigations on high silicon steel are similar to those of this study. Carbide-free upper bainite and retained austenite have been reported [8] for an Fe0.43wt.%C-3.00wt.%Mn-2.12wt.%Si alloy transformed at 350 °C and held for 74 h. However, when transformed at 300 °C to lower bainite, cementite precipitation occurred in the bainitic ferrite although no precipitation was observed in the retained austenite. Another investigation [9] also reported carbide-free upper bainite in Fe-0.74wt.%C-2.40wt.%Si0.52wt.%Cr-0.51wt.%Mn transformed at 380 °C. However, e-carbide was observed to form on the bainitic ferrite-displaced austenire twin boundaries after 30 min. No such interphase precipitation was found in the present study. From the TEM results of this study and others [ 8, 9], carbide-free transformation appears to occur only for upper bainite formation in both high silicon steels and nodular cast iron. Although evidence of carbide precipitation in upper bainitic ferrite and retained austenite was observed for a 400 °C austemper, it is believed that the precipitation t o o k place after the bainite reaction because of the aging of bainitic ferrite supersaturated with carbon and decomposition of carbon-enriched retained austenite. The results of previous studies [8, 11] on high silicon steel support this hypothesis. The lack of carbide formation in bainitic ferrite and retained austenite for a 370 °C austemper appears to be related to the austempering time. Previous work [3] indicated that a bay region exists in the isothermal transformation diagram for bainitic formation at about 370 °C (Fig. 16). Consequently, for austempering times less than B f, the extent of the bainite reaction is less at 370 °C than at other temperatures. For the 60 min austemper used in this study, specimens austempered at 400 °C would be fully transformed and precipitation processes in the bainitic ferrite and
1000 8 S 10%
50~
50%
90% Bf
900 450
o~
800 4oo
TOO ~" 35o
'I
300
600
\
250
200
i 10
I
i it 102
I
~
~ll ?03
500
I
J III
I
104
I
400
I 105
Time (see)
Fig. 16. I s o t h e r m a l t r a n s f o r m a t i o n diagram for b a i n i t e f o r m a t i o n in a 2.3 wt.% Si n o d u l a r cast iron a u s t e n i t i z e d at 9 0 0 °C.
retained austenite would have started. In contrast, the transformation to bainite would be incomplete in specimens austempered at 370 °C, and precipitation would be less advanced or not started. This interpretation of the effect of austempering time appears to explain the TEM observations. Consequently, the austempering time and not the temperature may be the unique factor determining carbide formation in upper bainitic nodular cast iron. The 2.88 wt.% Si nodular iron austenitized at 950 °C for 75 min and austempered at 370 °C for 60 min exhibited a transition ~carbide in the bainitic ferrite laths. Also, the presence of an elongated carbide precipitate, which was described as a long-range-ordered phase, was detected in the ferrite laths. The formation of these carbides after the long austenitizing time appeared to result from precipitation in bainitic ferrite laths supersaturated with carbon, after holding at the austempering temperatures. The long austenitizing time produced a high carbon content austenite which may have exceeded a critical level necessary for carbide formation. A more complete austenitizing-time-dependent study would be required to establish the level of critical carbon content. The presence of orthorhombic cementite and H-carbide in the bainitic ferrite of 1.57 wt.% Si nodular iron austenitized at 950 °C for 15 min and austempered at 370 °C for 60 min can be attributed to the low silicon
245
content. This same heat treatment produced carbide-free structures in the 2.88 wt.% Si nodular iron because carbide precipitation was suppressed as a result of the high silicon content. In addition, the presence of martensite in this iron indicates that the extent of carbon precipitation during austempering was sufficient to prevent complete stabilization of the austenite b y enrichment with carbon. The poor impact strength of austempered 1.57 wt.% Si nodular iron can be explained by these observations.
6. CONCLUSIONS
(1) A silicon content of 2.88 wt.% and austenitizing treatment of 15 min at 950 °C are necessary to avoid carbide precipitation in nodular iron when austempering at 370 °C for 60 min. These conditions result in the o p t i m u m combination of strength and impact toughness. (2) A lower silicon content and a longer austenitizing time both tend to p r o m o t e carbide precipitation in the bainitic ferrite laths. This is attributed to enhanced carbide stability at a lower silicon content (lower carbon activity) and a higher carbon content in the austenite for a longer austenitizing time (increased dissolution of graphite).
ACKNOWLEDGMENTS
The authors thank W. M. Gruszczynski for the preparation of the thin foils and A. K. Sachdev and M. D. Hanna for reviewing this report.
REFERENCES 1 P. A. Blackmore and R. A. Harding, The effects of metallurgical process variables on the properties of austempered ductile irons, Proc. 1st Int. Conf. on Austempered Ductile Iron, American Society for Metals, Metals Park, OH, 1984, p. 117. 2 Y. Ueda and M. Takita, Impact strength and isothermal transformation diagram of austempered ductile iron, Proc. 2nd Int. Conf. on Austempered
Ductile Iron, American Society for Metals, Metals Park, OH, 1986, p. 141. 3 M. M. Shea and E. F. Ryntz, Austempering nodular iron for optimum toughness, Trans. Am. Foundrymen's Soc., 94 (1986) 683. 4 W. S. Owen, Trans. Am. Soc. Met., 46 (1954) 812. 5 S. J. Matas and R. Hehemann, Trans. Metall. Soc. AIME, 221 (1961) 179. 6 J. Gordine and I. Codd, J. Iron Steel Inst., London, 207 (1969) 461. 7 R.M. Hobbs, G. W. Lorimer and N. Ridley, J. Iron Steel Inst., London, 210 (1972) 757. 8 H . K . D . H . Bhadeshia and D. V. Edmonds, MetaU. Trans. A, 10 (1979) 895. 9 B. P. J. Sandvik, Metall. Trans. A, 13 (1982) 777. 10 K. B. R u n d m a n and R. C. Klug, Trans. Am. Foundryrnen's Soc., 90 (1982) 499. 11 B. P. J. Sandvik, MetaU. Trans. A, 13 (1982) 789. 12 V. Franetovic, A. K. Sachdev and E. F. Ryntz, A transmission electron microscopy study of austempered lower bainitic nodular cast iron, Metallography, to be published. 13 R. L. Miller, Trans. Am. Soc. Met., 61 (1968) 592. 14 K. W. Andrews, D. J. Dyson and S. R. Keown, Interpretation o f Electron Diffraction Patterns, Plenum, New York, 2nd edn., 1971. 15 C. S. Roberts, J. Met., 5 (1953) 203. 16 M. J. Donachie, Jr., and O. H. Kriege, Phase extraction and analysis in superalloys -- summary of investigations B ASTM Committee E-4 Task Group 1, J. Mater., 7 (1972) 269-278. 17 S. Nagakura, Y. Hirotsu, M. Kusunoki, T. Suzuki and Y. Nakamura, Crystallographic study of the tempering of martensitic carbon steel by electron microscopy and diffraction, MetaU. Trans. A, 14 (1983) 1025-1031. 18 Powder Diffraction File, Joint Committee on Powder Diffraction Standards, International Center for Diffraction Data, Swarthmore, PA, 1982. 19 C. B. Ma, T. Ando, D. L. Williamson and G. Krauss, Chi-carbide in tempered high carbon martensite, Metall. Trans. A, 14 (1983) 10331045. 20 V. I. Izotov and L. M. Utevsky, Fiz. Met. Metalloved., 25 (1967) 98-110. 21 M. Dirand and L. Afqir, Acta MetaU., 31 (7) (1983) 1089. 22 ~Y. Hirotsu and S. Nagakura, Acta Metall., 20 (1972) 665. 23 Y. Hirotsu, K. C. Su and S. Nagakura, Trans. Jpn. Inst. Met., 17 (1976) 503. 24 Y. Tanaka and K. Shimizu, Trans. Jpn. Inst. Met., 22 (11) (1981) 779. 25 K. H. Jacks, J. Iron Steel Inst., London, 169 (1951) 26. 26 K. Shimizu and H. Okamoto, Trans. Jpn. Inst. Met., 15 (1974) 193.
231
Transmission Electron Microscopy Study of Austempered Nodular Iron: Influence of Silicon Content, Austenitizing Time and Austempering Temperature VJEKOSLAV FRANETOVIC, MICHAEL M. SHEA and EDWARD F. RYNTZ Metallurgy Department, General Motors Research Laboratories, Warren, MI 48090 (U.S.A.) (Received January 16, 1987; in revised form May 20, 1987)
ABSTRACT
The structures produced by austempering nodular iron were evaluated by transmission electron microscopy to explain the influence o f austempering temperature, austenitizing time and silicon content on the impact properties. No carbide precipitation was evident in a 2.88 wt. % Si nodular iron austenitized for 15 min at 950 °C and austempered for 60 min at 370°C. This iron had the highest impact strength. Austempering at 340 and 400 °C resulted in the formation o f transition carbides in bainitic ferrite laths and a lower impact strength. In addition, both a lower silicon content and a longer austenitizing time promoted carbide formation in the bainitic ferrite during austempering at 3 70 °C.
1. INTRODUCTION
Recent investigations [1, 2] have shown the outstanding mechanical properties produced by austempering nodular iron castings. However, the properties achievable by austempering can be divided into two groups: one consisting of high tensile and fatigue strength obtained with lower bainitic matrix structures and the other consisting of high ductility and impact strength obtained with upper bainitic matrix structures containing a large amount of untransformed austenite. Recent work [3] showed that a combination of high hardness and high impact strength can be produced in nodular iron by a unique austempering heat treatment. This heat treatment, consisting of a short austenitizing time, generally 15 min at 950 °C, followed by austempering at 370 °C for 1 h, produced a Brinell 0025-5416/87/$3.50
hardness of 310 HB and a room temperature impact strength of 170 J in a 2.88 wt.% Si nodular iron. The austempered microstructure consisted of upper bainite and retained austenite with a small amount of lower bainite. When the austenitizing time was increased to 75 rain or the silicon content was reduced to 1.57 wt.%, the impact strength decreased significantly at all austempering temperatures (Fig. 1). However, the maximum impact strength was achieved by austempering at 370 °C for each of the austenitizing conditions. Metallographic examination revealed that the superior properties produced by the 370 °C austemper were not the result of a significantly higher retained austenite content since a similar amount of austenite was present after austempering at 400 °C. Rather, the results indicated that austempering at 370 °C may produce a unique bainite structure and/or enhanced austenite stability. In nodular iron, similar to high silicon steels [4-9], the bainite reaction is believed to involve the nucleation and growth of bainitic ferrite unaccompanied by carbide formation. Consequently, during growth of the ferrite, the remaining austenite becomes increasingly enriched in carbon. Recent X-ray diffraction [10] and electron microprobe [3] results and, as will be shown later, the results of this study have confirmed the increase in austenite carbon content during austempering at 370 °C to form upper bainite. In addition, the electron microprobe study showed that the carbon content of the austenite decreased at extended austempering times (120 min), supporting the belief that decomposition of the carbon~nriched austenite ultimately occurs, similar to the behavior of high silicon steels [ 11 ]. Furthermore, the results showed that at 370 °C a bay region © Elsevier Sequoia/Printed in The Netherlands
232 180
stabilizing the austenite. However, the a m o u n t of retained austenite in lower bainite is known [8] to be less than in upper bainite, consistent with less enrichment in carbon of the austenite. In the present investigation the bainitic structures formed by austenitizing a 2.88 wt.% Si iron for 15 min and 75 min at 950 °C followed by austempering at 340 °C (lower bainite), 370 °C and 400 °C (upper bainite) for 60 min were evaluated by TEM in an attempt to explain the difference in impact properties. Of particular interest was the presence of carbides in the bainitic ferrite, in the retained austenite or at the ferrite-austenite interface. Also, the influence of silicon content on the structures produced during austempering at 370 °C for 60 min was evaluated to provide a better understanding of the effect of this factor on the structure and properties of austempered nodular iron.
160
140
120 w
L. @
100
m
ebu
a
Q.
E
80
60
m
m
2. PROCEDURE 40
m
20
I
340
I
370
I
400
Austempering Temperature (°C)
Fig. 1. Influence of austempering temperature, austenitizing time at 950 °C, and silicon content on the impact strength of nodular cast iron: e, 15 rain, 2.88 wt.% Si; m, 75 rain, 2.88 wt.% Si; A, 15 min, 1.57 wt.% Si.
occurs in the isothermal transformation diagram indicating higher stability of the austenite to bainite transformation. The occurrence of carbide-free bainite in nodular iron may only be applicable to transformation to upper bainite. The results of a recent transmission electron microscopy (TEM) investigation [12] on nodular iron {2.33 wt.% Si), austempered at 240 °C for 2 h to produce a lower bainite, indicated that cementite (FeaC) formed in the bainitic ferrite although no carbides were observed in the retained austenite. Although cementite precipitation occurred, an increase in the austenite carbon content presumably still t o o k place,
2.1. Casting process and heat treatment Nodular iron was prepared by treating 270 kg of iron in a ladle with 0.15 wt.% Mg added as a magnesium ferrosilicon alloy containing 5.5 wt.% Mg. The treated iron was inoculated with 0.50 wt.% Si using ferrosilicon (50 wt.% Si) during transfer to a pouring ladle. Keel block castings were poured into oilbonded dry sand molds. The composition of the irons is shown in Table 1. The b o t t o m section of the keel blocks was completely annealed prior to machining unnotched Charpy impact specimens. The annealing treatment, consisting of a 12 h hold at 930 °C followed by a slow cool of 55 °C h -1 in the furnace to 650 °C and air cooling to room temperature, produced a completely ferritic matrix structure although some intercellular carbides were present. The bars were copper plated to prevent decarburization during the austenitizing heat treatment. The copper plate was removed prior to impact testing. Austenitizing was performed in a resistance-heated vertical retort furnace with an atmosphere of nitrogen (1.7 m 3 h -1 ) and methane (0.4 m 3 h-l). The heat treatment consisted of austenitizing for 15 or 75 min at 950 0C followed by quenching in a nitrate-nitrite salt bath to austempering temperatures of 340, 370 or 400 °C. The
233 TABLE 1 Nodular iron compositions
Heat
A m o u n t ( wt.% )
C
equivalent
CN 2 9 6 5 CN 2 9 6 6
C
Si
Mn
Mg
S
P
(wt.%)
3.97 3.51
1.57 2.88
0.49 0.54
0.059 0.053
0.010 0.012
0.03 --
4.49 4.46
specimens were held in the salt for 60 min and air cooled to room temperature. After impact testing, TEM specimens were prepared according to the procedures below.
2.2. X-ray diffraction The retained austenite content of all specimens was determined by analysis of the X-ray diffraction spectra produced by Mo K s radiation using the procedure developed b y Miller [ 13 ]. The specimen holder provided b o t h rotating and rocking m o t i o n to minimize the effect of preferred orientation of the austenite. The volume fraction o f austenite was determined using the integrated intensities of the austenite 220 and 311 peaks and the ferrite 211 peak.
2.3. Transmission electron microscopy sample preparation The procedure developed for preparation of nodular iron foils suitable for viewing with the TEM consisted of three major steps: (a) mechanical polishing of nodular iron wafers to approximately 0.13 mm thickness or less; (b) electrochemical polishing to approximately 0.03 mm thickness or less; (c) ion milling to produce perforation on the thin edges of the specimen.
2.4. Mechanical polishing The starting point for the foil preparation was a 10 mm × 10 mm square slice 0.18 mm thick from an impact bar o f the desired material. The slice was obtained b y cutting the bar using a low speed saw with a 0.305 mm diamond wafering blade and a 152 gf load. Further mechanical thinning to provide a uniform sample 0.13 mm thick was accomplished b y mounting the slice on a 16 mm × 50 mm × 50 mm brass holding block that had a slot 0.13 mm deep and 12 mm wide ground on the center on one face. The slice was
m o u n t e d in this slot using an electrical conducting mixture of Duco cement, powdered graphite and enough acetone to form a thin even layer. A vinyl separator sheet, approximately 13 mm thick, was placed on t o p of the slice and another metal block was placed on top of the vinyl. The entire assembly was clamped together and allowed to dry. Following drying and disassembly, the slice was mechanically polished flush with the holding block using 3 2 0 , 4 0 0 and 600 grit silicon carbide paper. The specimens 3 mm in diameter were obtained from the slice by electrodischarge machining (EDM) using a brass electrode. The adhesive used to m o u n t the foil to the holding block allowed for electrical continuity for the EDM and prevented the cut specimen from being washed away b y the EDM cooling oil. The cut samples were separated from the holding block b y soaking in acetone.
2.5. Electropolishing Electropolishing was performed in an electrolyte consisting of 75 g of CrO3, 400 ml of glacial acetic acid and 21 ml of distilled H20, at r o o m temperature using a Fischione twin-jet electropolishing unit. The o p t i m u m electropolishing condition occurred at a current density of 18 mA c m - 2 . Although the time of polishing varied with specimen thickness, approximately 15 min was sufficient to produce minimal perforation. After the specimens had been po lished, they were rinsed successively in acetic acid and methanol.
2.6. Ion milling This technique is widely used to thin nonmetallic and metallic materials which cannot be uniformly etched electrochemically because of large differences in the electrochemical properties of the microconstituents. With ion milling, it was possible to produce perforations
234
having edges sufficiently thin to allow TEM examination. Thinning was accomplished with an argon ion energy of 1.5-2.0 kV. The ion thinning was extremely slow, requiring approximately 20 h at an initial angle of 20 ° between the ion beam and specimen and an additional 6 h at an angle of 7°-10 ° . The shallow finishing angle was necessary to obtain satisfactory TEM foils.
3. AUSTEMPERED MICROSTRUCTURES
3.1. Optical microscopy Optical micrographs o f the specimens evaluated are shown in Fig. 2. For the 2.88 wt.% Si iron austenitized at 950 °C for 15 min, the bainite formed by austempering at 340 °C (Fig. 2(a)) appeared to be lower bainite while upper bainite was primarily formed at 370 °C (Fig. 2(b)) and 400 °C (Fig. 2(c)). Increasing the austenitizing time to 75 min or decreasing the silicon content to 1.57 wt.% resulted in no significant change in the microstructure produced b y austempering at 370 °C. X-ray diffraction results indicated that the retained austenite content o f the 370 °C austempered samples was in the range 4 0 - 4 4 wt.%.
4. TRANSMISSION ELECTRON MICROSCOPY
The TEM results are presented in the following order: (1) 2.88 wt.% Si nodular iron austenitized at 950 °C for 15 min (sample 1, austempered at 370 °C (highest impact strength); sample 2, austempered at 400 °C; sample 3, austempered at 340 °C); (2) 2.88 wt.% Si nodular iron austenitized at 950 °C for 75 min (sample 4, austempered at 370 °C); (3) 1.57 wt.% Si nodular iron austenitized at 950 °C for 15 rain (sample 5, austempered at 370 °C). Samples 1-3 were selected to show the influence of austempering temperature on the microstructure of the 2.88 wt.% Si nodular iron austenitized at 950 °C for 15 min. Sample 4 was selected to reveal the effect of a longer austenitizing time (75 min) on the structure of the 2.88 wt.% Si iron. Finally, sample 5 was selected to reveal the influence of silicon content (1.57 wt.% Si compared with 2.88 wt.% Si).
Fig. 2. Microstructures of 2.88 wt.% Si nodular iron austenitized for 15 rain at 950 °C and austempered for 60 rain at (a) 340 °C, (b) 370 °C and (c) 400 °C (Nital etch). (Magnification, 400×.)
4.1. Sample 1 To characterize the austempered structure by TEM and to determine whether carbides had formed, it was essential initially to identify the ferrite and austenite phases. For that purpose, numerous micrographs and selected area diffraction (SAD) patterns corresponding to the ferrite and austenite phases with different zone axis were taken and analyzed. Figure 3 represents a typical TEM micrograph of sample 1 in which the ferrite and austenite are labeled b y F and A respectively.
235
Fig. 3. Transmission electron micrograph of sample 1.
Both phases show contrast due to a high density of dislocations. No contrast due to precipitates was revealed during careful tilting of the specimen and no additional reflections or other diffraction effects due to precipitates were f o u n d on the corresponding SAD patterns. Therefore, for this stage o f austempering, both phases were free of carbides. In addition, no carbide formation was observed at interphase boundaries between the two phases. The measured interplanar spacing dhk~ obtained for ferrite matched the published data [14] calculated on the basis of the lattice parameter a = 2.866 A well. An austenite lattice parameter a of 3.621 A was calculated using the measured interplanar spacings for retained austenite. Using the relationship between the austenite lattice parameter and carbon content given in ref. 15, the carbon content of the austenite was calculated to be 1.5 wt.%. The calculated value is in agreement with the results of a microprobe carbon determination for the austenite in the same sample [3]. The very high carbon content found in sample 1 can be correlated with the absence of carbide precipitation for this austempering temperature. In addition, the high austenite carbon content promotes stabilization of the austenite and prevents significant transformation to martensite during cooling to room temperature. The orientation relationship between ferrite and austenite was determined from SAD patterns taken at the interphase b o u n d a r y between these ~hases. Figure 4(a) is a bright
Fig. 4. (a) Bright field micrograph showing the ferriteaustenite interphase boundary; (b), (c) SAD patterns taken from areas A and F showing the N-W orientation relationship.
field micrograph showing ferrite at F and austenite at A and the associated interphase boundary. Several partly overlapping areas containing the interphase boundary were chosen and tilting experiments performed to obtain satisfactory SAD patterns for determination of the crystallographic relationship between the phases. Figure 4(b) is an SAD
236 pattern obtained mainly from austenite (primary diffraction pattern) but also contains ferrite reflections. The austenite orientations is close to [211]fcc. The superimposed diffraction pattern of ferrite indicates that the ferrite orientation is close to [110]bc¢. This zone axis, [110]b¢~ of ferrite, is clearly recognizable and indexed on Fig. 4(c) which was obtained b y slightly moving the specimen to orient the ferrite in a strong diffraction condition. Additional spots in Fig. 4(b) are due to double diffraction which appeared because of symmetrical orientation between the two phases. Two reflections due to double diffraction are indicated by arrows to distinguish them from the overlapping austenite (i 11) and ferrite ( i 1 0 ) reflections. The SAD patterns are correctly rotated to show the exact crystallographic orientation between the phases in the corresponding bright field micrograph. The correspondence of reflections to the observed phases was checked successively b y the dark field technique. The orientation relationship between ferrite and austenite was determined to be the Nishiyama-Wasserman orientation, which for the particular example is [211]~c~// [110]b¢c, (022)~¢~//(000.)bc¢ and (i11)~¢// (ilO)b~. The appearance of a small amount of an additional phase was evident in the sample. The structure of this phase was f.c.c, with a measured lattice parameter of 4.300 A. This phase belongs to the group of MC-type f.c.c. carbides where, generally, M -= Ti, Ta, Nb, Hf, Th or Zr [16]. According to measured interplanar spacings, structure determination and possible impurities in the investigated alloy, it was concluded that the phase was TiC. The measured interplanar spacings dhm of the identified phase are in very good agreement with that published for TiC [16]. This comparison is shown in Table 2. TiC adjacent to austenite is shown in Figs. 5(a) and 5(b) which are bright field and dark field micrographs respectively. The dark field micrograph was obtained b y using the (111)~c reflection of TiC. Also shown in Fig. 5(b) is the corresponding SAD pattern of the area shown in Fig. 5(a) which is in the foil orientation [011]fee. The presence of TiC in the iron was expected since titanium is a normal residual element in the Sorel metal used for the charge material. However, the TiC is highly stable in austenite and, consequently, would not dissolve during
TABLE 2 Comparison between the mean values of measured dhk ! with that of TiC [16] dhkl, measured (A) Sample 1
Sample 2
2.47 2.16 1.509 1.301 1.236 1.085
2.50 2.14 1.523 1.295 1.236 1.070 0.987
~.980
dhkl, TiC (A)
hkl
2.48 2.15 1.520 1.297 1.241 1.075 0.986
111 200 022 311 222 400 331
Fig. 5. TiC phase in sample 1 shown in (a) bright field and (b) dark field micrographs. The inset in (b) is the SAD pattern taken from the area shown in (a). austenitization and re-form during austempering. All specimens used in the previous mechanical property evaluation and this TEM study should contain TiC.
237
4.2. Sample 2 Ferrite and austenite were identified using the procedure described for sample 1. Analysis of the micrographs and SAD patterns indicated that precipitation had occurred in both phases. Unfortunately, neither precipitate gave satisfactory diffraction patterns to enable unambiguous identification. Fine precipitates revealed by the fact that the bend contour (A) passes over the austenite grain can be seen in Fig. 6(a), while Fig. 6(b) is the corresponding SAD pattern with the zone axis [011 ]~c¢. Additional strong reflections are from an adjacent austenite grain (top of Fig. 6(a)) and the faint reflections are probably due to the precipitates. The fine precipitates in bainitic ferrite are shown in Fig. 7(a) which is a dark field micrograph obtained from a (ll0)b~¢ diffuse reflection. The observed diffraction effect due to the fine precipitates was diffuse spikes, which can be seen around fundamental spots in Fig. 7(b) with the spike directions close to (102).
In Fig. 7(b) the foil orientation is [001]bcc. Such images can be attributed to interstitial carbon atom clusters. Similar diffraction effects have been observed during the preliminary stage of tempering of carbon steel [17]. Clusters appear in martensitic carbon steel at low tempering temperatures, but no tetragonality which would indicate tetragonal bainitic ferrite (or martensite) was detected in these SAD patterns. Although 400 °C appears to be too high a temperature for the interstitial carbon clusters to form (they formed at up to 100 °C in carbon steel), it can be assumed that the high silicon content moved the normal sequence of carbide formation to higher tem-
Fig. 6. (a) Transmission e l e c t r o n micrograph of austenite in sample 2; ( b ) t h e corresponding SAD pattern.
Fig. 7. (a) Dark field micrograph of ferrite in sample 2; (b) the corresponding SAD pattern.
238
peratures. In addition, no interphase carbide was observed in this specimen. The interplanar spacings obtained for ferrite were in very good agreement with calculated values using a lattice parameter a of 2.866 A [14]. However, the calculated austenite lattice parameter a of 3.586 A was significantly smaller than for sample 1 and represents a carbon content in the austenite of approximately 0.68 wt.%. On the basis of the lower austenite carbon content in sample 2 than in sample 1, the precipitates observed in sample 2 are believed to be carbides. The carbide precipitation probably occurred because of the higher austempering temperature (400 °C) compared with that for sample 1 (370 °C). Evidence for precipitation, which was not observed in sample 1, was clear in sample 2 and was confirmed b y additional precipitate reflections (Fig. 6(b)). This finding can be correlated with the lower impact strength of sample 2 than that of sample 1.
Since the carbide precipitation in austenite was observed on a very fine scale, it is believed that the precipitation reaction did not occur to an extent which would destabilize the austenite completely and result in decomposition to ferrite during further austempering or to martensite during cooling.
4.3. Sample 3 In sample 3, no contrast or diffraction effects due to precipitates could be found in the austenite. Austenite (labeled A) free of precipitation and with a high density of dislocations is shown at the t o p of the micrograph in Fig. 8(a). Calculation of the austenite lattice parameter indicated a carbon content similar to that of sample 2. However, there was indication that the start of precipitation occurred in the ferrite as shown in Figs. 8(a) and 8(b). Fine precipitates were observed in the ferrite (labeled F) in both the bright field micrograph (Fig. 8(a)) and the dark field micrograph (Fig. 8(b)) using a (211)bc¢ matrix reflection. No
Fig. 8. (a) Electron micrograph of sample 3 showing austenite at A, TiC at C and ferrite at F; (b) TiC illuminated by its (200) reflection; (c) ferrite illuminated by its (211) reflection.
239
satisfactory diffraction pattern could be obtained from these microstructural features. TiC with f.c.c, structure was also found in this sample. TiC (marked C) is oriented obliquely in the bright field micrograph (Fig. 8(a)) and dark field micrograph (Fig. 8(c)) where it is illuminated by its (9.00)fcc reflection. Additional micrographs of sample 3 (Fig. 9(a)) showed the presence of ferrite at F, a B2 structure phase, austenite at A and a carbide phase at C. Although the structure determination of the carbide phase was not completed, the interplanar spacings of the phase, obtained from different foil orientations, matched very well with that of x-carbide [18]. This comparison is shown in Table 3. The same type of carbide but with a different morphology was detected in an Fe-l.22wt.%C alloy [19] ' tempered at temperatures similar to the austempering temperatures of this study. This carbide phase is shown in the bright and dark field micrographs and the corresponding SAD
TABLE 3 Comparison between the mean values of measured
dhk ! with that of X-carbide [18]
dhkz, measured
(A) 2.06
2.01
1.81
1.508 1.364 1.339 1.315
1.277 1.266
dhkl, FesC2
hkl
2.06 2.03 1.98 1.91 1.80 1.76 1.72 1.67 1.62 1.57 1.50 1.37 1.34 1.32 1.27 1.25
510,021 315,402 511,211 221 312,511 402 421 512 602 113 422 331 331 802 531 712
(A)
Fig. 9. (a) Transmission electron micrograph of sample 3 showing austenite at A, ferrite at F and x-carbide at C; (b) dark field micrograph of x-carbide obtained from spot C; (c), (d) SAD patterns corresponding to area C in (a).
240
pattern in Figs. 9(a)-9(c). The dark field micrograph was made using the spot marked C in Fig. 9(c). Also, superlattice spots belonging to a B2 structure can be seen (indicated by an arrow) on the SAD pattern in Fig. 9(d). This diffraction pattern was obtained by tilting the foil from the position in Fig. 9(c). Hence, spot 2 in Fig. 9(c) corresponds to spot 2 in Fig. 9(d), but the zone axis of the carbide is different in Figs. 9(c) and 9(d). Indexing of the diffraction pattern (Fig. 9(d)) is shown in Fig. 10, where three separate diffraction patterns are identified: ferrite with the zone axis []33], a B2 structure with the zone axis [001] and an unidentified zone axis of monoclinic x-carbide. The observed superlattice reflections in the diffraction patterns of sample 3 cannot currently be explained. Superlattice spots observed for tempered martensite were attributed by Izotov and Utevsky [20] to the precipitation of tetragonal carbide, Fe4C. However, on the basis of the identification of X-carbide in sample 3, the presence of Fe4C seems unlikely.
4.4. Sample 4 After identification of the ferrite and austenite phases by SAD, each phase was examined for the presence of carbide precipitation. Figure 11 is a typical bright field micrograph showing highly dislocated retained austenite free of precipitation in the foil orientation [013]tcc. In addition, no diffraction effects due to the presence of precipitates were observed during subsequent examination. Careful evaluation of the ferrite phase indicated two modifications of the structure: ferrite free of precipitates and ferrite containing coherent carbide precipitates with two different morphologies, which both had symmetrical orientations with the ferrite. The bright field micrograph and dark field micrograph in Figs. 12(a) and 12(b) respectively show isolated ferrite with a high density o f dislocations but with no evidence of precipitation. The dark field micrograph was obtained by ( 0 1 i ) ferrite reflection. In the corresponding SAD pattern (Fig. 12(c)) the zone axis for the ferrite was [i11]bc¢. No tetragonal distortion of ferrite could be detected within the error of measuring the diffraction pattern. The measured interplanar spacings agreed well
J
5 - ~- ___-__ _
!
II0
I I
-
X I..,=X,.~_.
o~o
t tTo
301
310
Fig. 10. Indexed diffraction pattern shown in Fig. 9(d): , [133]bcc;- - - , [ O 0 1 ] o r d e r e d b c c with superlattice spots (×); -----, zone axis not identified, probably monoclinic X-carbide.
Fig. 11. Transmission electron micrograph of sample 4, showing precipitate-free austenite with a high density of dislocations.
with those of b.c.c, ferrite calculated using a = 2.865 A. Carbide precipitation was found in the ferrite in the region of the specimen containing alternate laths of ferrite and austenite (Fig. 13(a)). These carbides were similar to the e-carbide observed b y Sandvik [9] in an Fe-Si-C alloy austempered in the temperature range 2 9 0 - 3 8 0 °C. However, in the present
241
O
Fig. 13. (a) Transmission electron micrograph of sample 4, showing ferrite laths containing ~7-carbide; (b) SAD pattern showing the orientation relationship between 7?-carbide, ferrite and austenite.
Fig. 12. Transmission electron micrographs of sample 4: (a) precipitate-free ferrite; (b) ferrite in a dark field taken using the ferrite (011) reflection; (c) SAD pattern corresponding to the area shown in (a).
study the carbide was identified as ~?-carbide. This t y p e o f unstable transition carbide is usually formed during the early stages o f tempering of martensite. The structure of the carbide was determined to be orthorhombic with the lattice parameters a -- 4.704 A, b -4.318 A and c = 2.830 A [21, 22]. This structure belongs to the space group P n n m and the carbon atoms regularly occupy half the octahedral interstices [17, 22]. Since 7?-carbide
can be present simultaneously with e-carbide during tempering of martensite [ 21] and all first-order reflections from e-carbide are covered by ~?-carbide reflections, identification o f the ~?-carbide was difficult. However, carefully measured angles between corresponding g vectors, which were indexed according to orthorhombic unit cells, agreed very well with theoretical values calculated using a = 4.704/~, b = 4.318 A and c = 2.830 A [22]. The interplanar spacings of the carbide were measured and matched with the calculated values for 1?-carbide (Table 4). The mutual orientation relationships between ferrite, austenite and ~?-carbide were determined using an SAD pattern taken from the ferrite-austenite interface (Fig. 13(b)). A symmetrical orientation relationship was found between the phases. In Fig. 13(b), ferrite was
242 TABLE 4 Comparison between the measured dhM with that of ~?-carbide [4]
dhM, measured
dhM, ~7-carb id e
hkl
4.69 3.15 2.86 2.44 2.39 1.81 1.65
4.70 3.18 2.83 2.42 2.37 1.809 1.668
100 110 001 101 011 201 211
(h)
(h)
indexed as [012]bcc and this zone axis is parallel to the [i11]fcc zone axis of austenite on the superimposed austenite diffraction pattern. Faint reflections due to 7?-carbide were indexed according to the orthorhombic structure with the zone axis [010]. This zone axis is parallel to the zone axis of both ferrite and austenite. Additional reflections indicated by the arrow appeared because of double diffraction due to symmetrical orientation between the phases. The above results indicate that (001) of ~?-carbide is parallel to (001) of ferrite and ( 0 i l ) of austenite. The relationship between ferrite and austenite is consistent with the Nishiyama-Wasserman orientation relationship between b.c.c, and f.c.c, phases. The simultaneous relationship between 7?-carbide, ferrite and austenite has not been previously reported. Similar 7?-carbide has been observed in carbon steel [17, 22], nickel steel [23], manganese steel [24] and Si-Mn steel [21]. According to Jacks [25], the loss of tetragonality of martensite can be attributed to the formation o f a transitional carbide which he identified as e-carbide, although more recent studies [19, 22, 23, 26] indicated that the carbide was orthorhombic 7?-carbide. This is consistent with the present findings of no tetragonal distortion in the ferrite phase. Another type of carbide was observed in the ferrite as elongated images with the axis close to [001]be c. Figures 14(a) and 14(b) respectively are a bright field micrograph of ferrite and a dark field micrograph of an ordered carbide phase taken with the ½ ~ ½ superlattice reflection. According to the SAD pattern (Fig. 14(c)) which represents the
Fig. 14. Transmission electron micrographs of sample 4: (a) ferrite; (b) ferrite in a dark field taken using the 31 231 3 superlattice reflection; (c) SAD pattern corresponding to the area shown in (a).
[ 2 i 0 ] zone of ferrite, the carbide can be described as a long-range-ordered phase. The variation in diffracted intensity due to the ordered structure is clearly shown by the superlattice reflections marked S in Fig. 14(c). 4.5. Sample 5 A typical electron micrograph of this sample (Fig. 15) indicated that the retained austenite (A) was free of carbide precipitation. How-
243
Fig. 15. Transmission electron micrograph of sample 5 showing carbide-free austenite at A and two types of carbide in ferrite at B and C. The inset dark field micrograph is from the carbides at B.
ever, two types of precipitate were observed in ferrite at B and C. The precipitates at B are shown in the dark field inset taken by (121) orthorhombic reflection in Fig. 15. The diffraction pattern o f the precipitates was indexed as an orthorhombic structure with a = 4.524 A, b = 5.088 • and c = 6.741 A, which corresponds to cementite [14]. The measured angles between the g vectors corresponded to those calculated for orthorhombic cementite. In addition, the measured interplanar spacings matched those for cementite [14]. The morphology of the precipitates at C is similar to that of the carbides observed by Sandvik [9] in austempered high silicon steel. Analysis of diffraction patterns indicated t h a t the precipitates were ~7-carbide. A typical diffraction pattern parallel to the [ 3 i 0 ] zone axis of ferrite contained faint q-carbide reflec-
tions and double
5. DISCUSSION On the basis of these TEM results, the higher impact strength produced in the 2.88 wt.% Si iron by austenitizing for 15 min at 950 °C
244
followed by austempering at 370 °C for 60 min compared with 340 or 400 °C, appears to be related to the absence of carbide formation and the resulting high stability of the retained austenite due to enrichment with carbon. At the present time, it is unclear w h y carbide formation was not observed in either ferrite or retained austenite for austempering at 370 °C. However, the results of other investigations on high silicon steel are similar to those of this study. Carbide-free upper bainite and retained austenite have been reported [8] for an Fe0.43wt.%C-3.00wt.%Mn-2.12wt.%Si alloy transformed at 350 °C and held for 74 h. However, when transformed at 300 °C to lower bainite, cementite precipitation occurred in the bainitic ferrite although no precipitation was observed in the retained austenite. Another investigation [9] also reported carbide-free upper bainite in Fe-0.74wt.%C-2.40wt.%Si0.52wt.%Cr-0.51wt.%Mn transformed at 380 °C. However, e-carbide was observed to form on the bainitic ferrite-displaced austenire twin boundaries after 30 min. No such interphase precipitation was found in the present study. From the TEM results of this study and others [ 8, 9], carbide-free transformation appears to occur only for upper bainite formation in both high silicon steels and nodular cast iron. Although evidence of carbide precipitation in upper bainitic ferrite and retained austenite was observed for a 400 °C austemper, it is believed that the precipitation t o o k place after the bainite reaction because of the aging of bainitic ferrite supersaturated with carbon and decomposition of carbon-enriched retained austenite. The results of previous studies [8, 11] on high silicon steel support this hypothesis. The lack of carbide formation in bainitic ferrite and retained austenite for a 370 °C austemper appears to be related to the austempering time. Previous work [3] indicated that a bay region exists in the isothermal transformation diagram for bainitic formation at about 370 °C (Fig. 16). Consequently, for austempering times less than B f, the extent of the bainite reaction is less at 370 °C than at other temperatures. For the 60 min austemper used in this study, specimens austempered at 400 °C would be fully transformed and precipitation processes in the bainitic ferrite and
1000 8 S 10%
50~
50%
90% Bf
900 450
o~
800 4oo
TOO ~" 35o
'I
300
600
\
250
200
i 10
I
i it 102
I
~
~ll ?03
500
I
J III
I
104
I
400
I 105
Time (see)
Fig. 16. I s o t h e r m a l t r a n s f o r m a t i o n diagram for b a i n i t e f o r m a t i o n in a 2.3 wt.% Si n o d u l a r cast iron a u s t e n i t i z e d at 9 0 0 °C.
retained austenite would have started. In contrast, the transformation to bainite would be incomplete in specimens austempered at 370 °C, and precipitation would be less advanced or not started. This interpretation of the effect of austempering time appears to explain the TEM observations. Consequently, the austempering time and not the temperature may be the unique factor determining carbide formation in upper bainitic nodular cast iron. The 2.88 wt.% Si nodular iron austenitized at 950 °C for 75 min and austempered at 370 °C for 60 min exhibited a transition ~carbide in the bainitic ferrite laths. Also, the presence of an elongated carbide precipitate, which was described as a long-range-ordered phase, was detected in the ferrite laths. The formation of these carbides after the long austenitizing time appeared to result from precipitation in bainitic ferrite laths supersaturated with carbon, after holding at the austempering temperatures. The long austenitizing time produced a high carbon content austenite which may have exceeded a critical level necessary for carbide formation. A more complete austenitizing-time-dependent study would be required to establish the level of critical carbon content. The presence of orthorhombic cementite and H-carbide in the bainitic ferrite of 1.57 wt.% Si nodular iron austenitized at 950 °C for 15 min and austempered at 370 °C for 60 min can be attributed to the low silicon
245
content. This same heat treatment produced carbide-free structures in the 2.88 wt.% Si nodular iron because carbide precipitation was suppressed as a result of the high silicon content. In addition, the presence of martensite in this iron indicates that the extent of carbon precipitation during austempering was sufficient to prevent complete stabilization of the austenite b y enrichment with carbon. The poor impact strength of austempered 1.57 wt.% Si nodular iron can be explained by these observations.
6. CONCLUSIONS
(1) A silicon content of 2.88 wt.% and austenitizing treatment of 15 min at 950 °C are necessary to avoid carbide precipitation in nodular iron when austempering at 370 °C for 60 min. These conditions result in the o p t i m u m combination of strength and impact toughness. (2) A lower silicon content and a longer austenitizing time both tend to p r o m o t e carbide precipitation in the bainitic ferrite laths. This is attributed to enhanced carbide stability at a lower silicon content (lower carbon activity) and a higher carbon content in the austenite for a longer austenitizing time (increased dissolution of graphite).
ACKNOWLEDGMENTS
The authors thank W. M. Gruszczynski for the preparation of the thin foils and A. K. Sachdev and M. D. Hanna for reviewing this report.
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Ductile Iron, American Society for Metals, Metals Park, OH, 1986, p. 141. 3 M. M. Shea and E. F. Ryntz, Austempering nodular iron for optimum toughness, Trans. Am. Foundrymen's Soc., 94 (1986) 683. 4 W. S. Owen, Trans. Am. Soc. Met., 46 (1954) 812. 5 S. J. Matas and R. Hehemann, Trans. Metall. Soc. AIME, 221 (1961) 179. 6 J. Gordine and I. Codd, J. Iron Steel Inst., London, 207 (1969) 461. 7 R.M. Hobbs, G. W. Lorimer and N. Ridley, J. Iron Steel Inst., London, 210 (1972) 757. 8 H . K . D . H . Bhadeshia and D. V. Edmonds, MetaU. Trans. A, 10 (1979) 895. 9 B. P. J. Sandvik, Metall. Trans. A, 13 (1982) 777. 10 K. B. R u n d m a n and R. C. Klug, Trans. Am. Foundryrnen's Soc., 90 (1982) 499. 11 B. P. J. Sandvik, MetaU. Trans. A, 13 (1982) 789. 12 V. Franetovic, A. K. Sachdev and E. F. Ryntz, A transmission electron microscopy study of austempered lower bainitic nodular cast iron, Metallography, to be published. 13 R. L. Miller, Trans. Am. Soc. Met., 61 (1968) 592. 14 K. W. Andrews, D. J. Dyson and S. R. Keown, Interpretation o f Electron Diffraction Patterns, Plenum, New York, 2nd edn., 1971. 15 C. S. Roberts, J. Met., 5 (1953) 203. 16 M. J. Donachie, Jr., and O. H. Kriege, Phase extraction and analysis in superalloys -- summary of investigations B ASTM Committee E-4 Task Group 1, J. Mater., 7 (1972) 269-278. 17 S. Nagakura, Y. Hirotsu, M. Kusunoki, T. Suzuki and Y. Nakamura, Crystallographic study of the tempering of martensitic carbon steel by electron microscopy and diffraction, MetaU. Trans. A, 14 (1983) 1025-1031. 18 Powder Diffraction File, Joint Committee on Powder Diffraction Standards, International Center for Diffraction Data, Swarthmore, PA, 1982. 19 C. B. Ma, T. Ando, D. L. Williamson and G. Krauss, Chi-carbide in tempered high carbon martensite, Metall. Trans. A, 14 (1983) 10331045. 20 V. I. Izotov and L. M. Utevsky, Fiz. Met. Metalloved., 25 (1967) 98-110. 21 M. Dirand and L. Afqir, Acta MetaU., 31 (7) (1983) 1089. 22 ~Y. Hirotsu and S. Nagakura, Acta Metall., 20 (1972) 665. 23 Y. Hirotsu, K. C. Su and S. Nagakura, Trans. Jpn. Inst. Met., 17 (1976) 503. 24 Y. Tanaka and K. Shimizu, Trans. Jpn. Inst. Met., 22 (11) (1981) 779. 25 K. H. Jacks, J. Iron Steel Inst., London, 169 (1951) 26. 26 K. Shimizu and H. Okamoto, Trans. Jpn. Inst. Met., 15 (1974) 193.