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Quantitative microstructural analysis of M2 grade high speed steel during high temperature treatment
PII:
Acta mater. Vol. 46, No. 14, pp. 5207±5220, 1998 # 1998 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved Printed in Great Britain S1359-6454(98)00110-4 1359-6454/98 $19.00 + 0.00
QUANTITATIVE MICROSTRUCTURAL ANALYSIS OF M2 GRADE HIGH SPEED STEEL DURING HIGH TEMPERATURE TREATMENT M. R. GHOMASHCHI School of Engineering, University of South Australia, The Levels, SA 5095, Australia (Received 17 September 1997; accepted 13 February 1998) AbstractÐThe microstructural changes of M2 grade high speed steel were examined quantitatively to identify the mechanisms responsible for changes during high temperature treatment. Reheating at high temperature leads to major microstructural changes involving phase transformation, spheroidization and coarsening of carbides formed in the as-cast M2 grade high speed steel. The M2C carbide decomposes rapidly to form MC and M6C in less than 1 h after attaining the reheating temperatures of 1150 or 12008C. Further reheating brings about the eutectic M6C carbide spheroidization and general coarsening of both M6C and MC. Carbide particles situated at the original austenite grain boundaries grow preferentially until impingement of carbides along these boundaries takes place. It is the diusion of tungsten in g-Fe which controls spheroidization rate of M6C carbides, while diusion of vanadium and tungsten in g-Fe is the rate controlling mechanism for coarsening of MC and M6C carbides, respectively. The process of spheroidization of alloy carbides such as M6C, is somewhat dierent from that of cementite where impingement of the already spheroidized carbide particles make the analysis more dicult. # 1998 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved.
1. INTRODUCTION
It has been reported by the present author [1] that the solidi®cation of M2 grade high speed steel leads to the formation of coarse M2C particles at the centre of dendrites as a result of a peritectic decomposition of d-ferrite. Subsequent enrichment of solute atoms around the growing austenite dendrites leads to the formation of coarse interdendritic eutectic carbides of MC, M2C and M6C. Further cooling below the solidus results in general precipitation of small carbides of M2C and MC. The eutectic and small carbides were then reported to have dierent morphologies and chemistry [2, 3]. In addition some ultra-®ne carbides were also detected [1]. Reheating leads to rapid dissolution of the ultra®ne carbides. The M2C particles decompose rapidly to form MC and M6C in less than 1 h after attaining the reheating temperature of 1150 or 12008C. Further reheating brings about the eutectic M6C carbide spheroidization and general coarsening of both M6C and MC. Carbide particles situated at the original austenite grain boundaries grow preferentially until impingement of carbides along these boundaries takes place. The majority of small carbides dissolve as coarsening occurs, but occasional M6C particles grow to sizes equivalent to those of {The same author reported the eect of hot forging on carbides in M2 grade high speed steel. Interested readers are referred to Metall. Trans., 1993, 24A, 2171.
the original central large or eutectic carbides. These newly-formed carbides are called intragranular grown carbides. Figure 1 illustrates the original cast structures and the resultant structures due to reheating at high temperatures. The schematic diagram included in Fig. 1 speci®es the nomenclature used in this investigation. In all previous publications [1±3] a precise qualitative picture was given for carbide formation during solidi®cation and their subsequent changes due to reheating. The present article is designed to characterize the microstructural changes quantitatively for better understanding of such a complex system and to provide insights for all high speed steels.{ 2. MATERIALS AND EXPERIMENTAL PROCEDURES
The alloy used in this study was prepared in the laboratory by air melting using ARMCO iron as the base metal with the addition of white cast-iron, ferrotungsten, ferrovanadium, pure chromium and molybdenum as the alloy elements. Silicon was added for deoxidization and casting was carried out at a temperature of 15608C in two sand and chill moulds. The size of the sand and chill cast ingots were 292 75 24 and 420 78 38 mm3, respectively. These steels will be referred to as sand and chill cast and their chemical analysis is as follows (wt%): 0.86 C, 6.85 W, 5.3 Mo, 4.22 Cr, 1.96 V, 0.25 Si and 0.07 Mn, (Al, Nb, Ti, Co R 0.02).
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Table 1. The list of specimens with their soaking time and temperature Sand
Chill
Time (h)
12008C
11508C
12008C
11508C
H51 H50 H53 H55 H57 H41 H24 H26 H22 H32 H29 H34 H30
H60 H62 H64 H76 H74 H43 H45 H47 H48 Ð H78 Ð H80
H52 Ð H54 H56 H58 H40 H23 H25 H27 H31 H28 H33 H35
H61 H63 H65 H77 H75 H42 H44 H46 H49 Ð H79 Ð H81
0.5 1 2 4 10 25 50 75 100 125 150 175 200
Specimen preparation, thermal treatments and various etching techniques together with dierent etchants and microscopes used for dierent purposes are similar to those reported before [1]. Table 1 summarizes reheating times and temperatures employed during the course of the present investigation. In order to characterize microstructural features in the as-cast and heat treated conditions, the following quantitative metallographic techniques were employed. (i) Point counting. Systematic point counting was carried out on parallel lines on the plane of polish with a light microscope with a travelling stage connected to a digital counter. The stage moved in 0.05 mm increments. A total of 1200 points was counted for each specimen to reduce the statistical error to less than 1%. (ii) Size distribution. A semi-automatic Carl Ziess particle size analyser was used in conjunction with light micrographs taken from specimens etched electrolytically in chromic acid solution for size distribution measurements. The size evaluation is based on equivalent circle area diameter, ECAD; that is the diameter of a circle having the same area as the object to be measured. A total of 250 particles was measured for each specimen and the arithmetic mean particle size d was calculated. In addition, both light and scanning electron microscopes were used to measure the number of particles per unit area. For light microscopy, two magni®cations of 1600 and 200 times were used for counting the number of small and intragranular grown carbides, respectively. The size criterion for small carbides was based on the ability of the light microscope to distinguish between two individual particles, i.e. resolution. The microscope resolving power was improved by application of immersion oil, d 1= 2m sin a, where the resolving power, d, was improved to 0.2 mm. The intragranular grown carbides, however, were considered dierently, where the 200 times magni®cation made them easy
to recognize in a background of ®ne carbides uniformly distributed within the grains. Any particles with dimensions of about 1 mm at 200 times were counted as intragranular grown carbides. A total number of ten randomly selected ®elds of view were examined. In order to measure the number of grain boundary eutectic carbides, SEM was used at 5000 times magni®cation. A total number of 30 ®elds of view was selected randomly from each specimen. 3. RESULTS
The microstructures of both sand and chill cast specimens were similar with chill cast being slightly ®ner. Furthermore, both responded equally to high temperature treatments although the kinetics of reactions were slightly faster for chill cast. Therefore, the following results are presented mainly for sand cast specimens. 3.1. Volume fraction As has been reported before [1], the carbides may be divided into the following categories. 3.1.1. Grain boundary eutectic carbides. The ascast structures contain M2C, MC and M6C eutectic carbides where the M2C portion transforms to M6C and MC within 30 min of reheating at 1150± 12008C. Therefore, the volume fractions of M6C and MC carbides were measured and recorded in Table 2. As is evident, there is a general trend towards increasing the percentage of grain boundary eutectic carbides with soaking times at 1150 and 12008C for both sand and chill cast structures. However, after a certain time it appears to have reached an equilibrium stage, see Fig. 2. 3.1.2. Spheroidization. Carbide plates are not stable at high temperatures and break down into smaller segments, i.e. spheroidization. The grain boundary eutectic carbides were examined very closely and as reported earlier [1], spheroidization occurs only of M6C carbides, which is then followed by coarsening and impingement processes active within the already spheroidized M6C particles. Therefore, to study spheroidization, the following measurements were carried out on M6C carbides. (a) The amount of spheroidized M6C carbide was obtained for only isolated carbide particles Table 2. Volume fraction of grain boundary eutectic carbides for sand cast specimens reheated at 12008C followed by water quench Specimen H50 H53 H41 H24 H26 H29 H30
Vv% MC
2DVv%
Vv% M6C
2DVv%
1.6 1.7 2.1 2.1 2.4 1.9 2.0
0.3 0.3 0.3 0.4 0.4 0.3 0.3
9.1 9.5 8.4 10.3 10.3 10.2 9.2
0.6 0.7 0.6 0.7 0.7 0.7 0.7
GHOMASHCHI: MICROSTRUCTURAL ANALYSIS OF HIGH SPEED STEEL
Fig. 1. Optical micrographs of the as-cast and reheated M2 grade high speed steel: (a) as-cast sand structure; (b) reheated sand cast structure, 150 h at 11508C, water quenched; (c) schematic diagram to show the nomenclature for carbides.
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having an aspect ratio R5. If there are N1 plates, N2 necked plates and N3 isolated M6C particles, the percentage spheroidized M6C carbide is N3 100: 1 Sph%M6 C N1 N2 N3 Table 3(a) shows the results for sand cast structure at 12008C which are plotted in Fig. 3. As can be seen, the resultant graph is similar to conventional spheroidization of cementite, for instance, where percentage spheroidized cementite increases with time [4]. However, the process of spheroidization appears to have stopped at about 80% spheroidized M6C. (b) It was further decided to count the necked plates as spheroids due to impingement eect, provided that the aspect ratio of the impinged particle is within the set criterion, i.e. AR R 5. The percentage spheroidization is then calculated as N2 N3 Sph%M6 C 100: 2 N1 N2 N3 Table 3(b) shows the new results which are superimposed on Fig. 3. The percentage of impinged or necked particles was then calculated from equations (1) and (2) and plotted in Fig. 3. 3.1.3. Central large carbides. The central large carbides form due to a complex peritectic reaction at around 13308C where part of the d-ferrite within the peritectic reaction transforms to a low carbon austenite and carbide{ 0 1 L d ÿ4g @ d ÿ4g M2 C A L d ÿ4g M2 C and are mainly M2C in the as-cast state, but as for grain boundary eutectic carbides they too transform to M6C and MC within the ®rst 30 min of reheating. The amount of central large carbides increases with increasing reheating time, see Table 4 and Fig. 2. 3.1.4. Intragranular grown carbides. A few M6C carbide particles grow preferentially within the initial austenite dendrites where their percentage increases initially with time but remains steady after a speci®c time, depending on the reheating temperature and cast structure [1]. The results are shown in Table 4 and Fig. 2.
Table 3. To show the eect of soaking time on the percentage spheroidization in M6C grain boundary eutectic carbide (a) Volume fraction of isolated M6C grain boundary eutectic carbide particles with AR R5 N3 2DVv Specimen Vv sph% 100 N1 N2 N3 H53 H57 H41 H24 H26 H29 H30
58.9 76.2 70.9 75.5 78.8 383 84.2
5.7 4.4 5 4.2 3.9 3.6 3.6
(b) Volume fraction of isolated and impinged M6C carbide particles with AR R 5 N2 N3 2DVv Specimen Vv sph% 100 N1 N2 N3 H57 H41 H24 H26 H32 H29 H30
78.3 89.2 91.9 96.1 96.7 99 100
4 3.2 2.5 1.7 1.6 0.9 0
during early stages of heat treatment which is concurrent with general dissolution of MC and growth of occasional M6C. The NA is a useful means to study both MC dissolution and M6C growth. It was found that the small carbides disappear after a certain period of time depending upon the cast structure and soaking temperature, where ®ner cast structure and higher soaking temperatures are recipes for greater dissolution and coarsening rates. Figure 4 shows the results for sand cast structure. It is noticeable, in Fig. 4, that the ®nal stages of small carbides dissolution is concurrent with a steady state for the number of intergranular grown carbides. The results are shown in Table 5. 3.2.2. Grain boundary eutectic carbides. The number of grain boundary eutectic carbides was measured with the view that impinged carbides
3.2. Number of particles per unit area (NA) 3.2.1. Small carbides. It has been shown that the small carbides consist of M2C needles and MC cubes [2]. The M2C transforms to M6C and MC
{For more information, see Ref. [1] and J. McLaughlin et al., Metall. Trans., 1977, 8A, 1787.
Fig. 2. Eect of reheating time on the volume fraction of carbides.
GHOMASHCHI: MICROSTRUCTURAL ANALYSIS OF HIGH SPEED STEEL
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Table 4. Volume fraction of central large and intragranular grown carbides for sand cast specimens at 12008C Specimen
H53 H57 H41 H24 H26 H29 H30
Intragranular grown carbides
Central large carbides
Vv%
2DVv%
Vv% M6C
2DVv%
0.4 0.7 1.2 1.5 1.9 1.7 1.0
0.2 0.2 0.3 0.4 0.4 0.4 0.3
0.3 0.9 1.2 1.4 1.3 2 1.5
0.2 0.3 0.3 0.4 0.4 0.4 0.4
were counted as isolated ones after impingement. The results are given in Table 6 and plotted in Fig. 5. After an initial rise, the number of eutectic carbides decreases to an almost constant value after about 75 h at temperature. 3.3. Size distribution Size distribution analysis was carried out on MC and M6C grain boundary eutectic carbides to characterize both spheroidization and coarsening phenomena. However, since MC carbide particles were almost fully spheroidized both in the as-cast and reheated structures, coarsening appears to be the only active process within MC carbide. 3.3.1. MC. The size values for MC carbide were arranged into a number of equal valued class intervals and plotted against the relative frequency of each respective class interval. The resulting distribution graphs are presented as histograms from which the arithmetic mean particle size was calculated for each specimen. The distribution histograms are almost the same for both sand and chill cast steels at 1150 and 12008C, and thus the sand cast structure at 12008C is presented here, Fig. 6. The X-axis in Fig. 6 is both the ratio of each inter i.e. val's size (di) over the arithmetic mean size (d), and the actual size of each interval. As r di =d, soaking time increases, ®rst, the number of smaller particles decreases, and second, some larger particles are introduced into the distribution graphs. The cut-o value is generally in the range of 1.6 < r < 1.8, although according to the theoretical treatments [5±7], it should be about 1.5. The mechanism responsible for coarsening may be deduced from a logarithmic plot of mean particle size vs time according to the theory of coarsening, i.e. dn ÿ dn0 Kt where d and d0 are the mean particle sizes at time t and zero, respectively. K is the rate constant and n the mode of coarsening, where for n = 2, 3, 4, 5, coarsening is dominated by interface, bulk or volume, grain boundary and pipe (dislocation) diusion, respectively. In order to estimate the activation energy for coarsening of the MC grain boundary eutectic carbide, the cubes of mean {See Appendix A. {See Section 4.3 for more details.
Fig. 3. Eect of reheating time on the percentage spheroidization of M6C grain boundary eutectic carbides.
sizes{ obtained from the histograms were plotted against time for both the sand and chill cast structures at 1200 and 11508C. Figure 7 shows the result for sand cast only. The activation energy was determined by calculating the slope of d3±t curves, i.e. a @ d3 =@ t, at t>60 h, in which a is equivalent to the rate constant, K. Since K is both temperature and diusion dependent [5±7],{ a may be presented in terms of an Arrhenius equation, i.e.
Table 5. The number of small, and intragranular grown carbides per 1 mm2 for sand cast structure at: (a) 12008C; (b) 11508C Small carbides$ Specimen
NA
H51 H50 H53 H55 H57 H41 H24 H26 H22 H32 H29 H34 H30
438 400 453 120 368 000 342 400 292 480 228 480 245 760 53 120 0 0 0 0 0
H60 H62 H64 H76 H74 H43 H45 H47 H48 H78 H80
766 080 691 840 604 160 560 640 538 240 533 760 477 440 416 000 366 720 167 680 151 040
2S.E. (a) 20 480 35 840 18 560 18 560 26 880 19 200 25 920 20 480 0 0 0 0 0 (b) 29 120 49 600 39 040 33 920 36 480 46 080 18 880 20 800 21 120 19 840 16 320
Intragranular grown carbides$ NA
2S.E.
50 160 160 220 300 440 566 630 390 520 1080 380 460
48 72 92 120 88 124 152 224 88 128 232 128 148
30 0 170 90 80 420 410 570 576 830 710
30 0 72 48 64 152 92 104 68 100 96
$All measurements are based on 95% con®dence limit (NA22S.E.) as plotted in Fig. 4 (S.E. = standard error).
Fig. 4. Eect of reheating time on the dissolution of small carbides and precipitation of intragranular grown carbides for sand cast structure: (a) 11508C; (b) 12008C.
5212 GHOMASHCHI: MICROSTRUCTURAL ANALYSIS OF HIGH SPEED STEEL
GHOMASHCHI: MICROSTRUCTURAL ANALYSIS OF HIGH SPEED STEEL
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Table 6. The number of grain boundary eutectic carbides per 1 mm2 for sand cast structure at: (a) 12008C; (b) 11508C Specimen Mean (NA) Std. 2S.E. Specimen Mean (NA) Std. 2S.E.
H51
H50
H53
H55
H57
30 000 18 500 10 000
40 500 34 500 18 000
62 500 33 500 17 000
60 500 27 000 14 000
46 500 18 500 10 000
H60
H62
H64
H76
H74
30 000 20 000 10 000
24 000 12 000 6000
49 500 32 000 17 000
70 500 38 500 23 000
74 000 45 000 23 000
(a)
(b)
H41
H24
H26
H32
H34
H30
45 000 15 000 8000
55 500 22 000 12 000
36 500 15 500 8000
34 500 10 000 5000
36 500 12 500 7000
38 500 19 500 10 000
H43
H45
H47
H78
H80
49 000 23 500 12 000
51 000 33 000 17 000
39 000 22 500 12 000
49 000 17 500 9000
58 000 19 500 10 000
$All measurements are based on 95% con®dence limit (NA22S.E.) as plotted in Fig. 5 (S.E. = standard error).
a b eÿQ=RT , from which the activation energy may be calculated. The activation energy of MC carbide together with the arithmetic mean particle sizes and their cubes are given in Table 7. The activation energy for both sand and chill cast structures is very close, i.e. 250±300 kJ/mol. This is close to the activation energy for vanadium diusion in g-iron, i.e. QgV =293 kJ/mol [8]. 3.3.2. M6C. The M6C carbide, unlike MC, is more complicated since it is exposed to dierent processes of spheroidization, coarsening and impingement. The sizing of the M6C particles, Table 8, was carried out the same as for MC, and the impinged particles were counted as separate particles. Figure 8 shows the histogram distribution of M6C carbide particles at 12008C for sand cast structure. Close examination of the histograms at various times reveals the eectiveness of spheroidization and coarsening processes as for instance H55 and H57 for spheroidization and H41 and H24 for coarsening. The cut-o size is between 1.6 < r < 1.8 and an activation energy of 240± 280 kJ/mol was calculated for both sand and chill cast structures, see Fig. 9. In order to study the eect of coarsening separately, the top 25% largest M6C particles were plotted against time in Fig. 10. Coarsening is
Fig. 5. Eect of reheating time on the number of grain boundary eutectic carbides.
expected as the dominant phenomenon within these carbide particles, and the rate of coarsening is faster in early stages of treatment.
4. DISCUSSION
4.1. Volume fraction of carbides Figure 2 shows that the volume fraction of grain boundary eutectic carbides increases with time. There are also the occasionally grown intragranular grown carbides and a slight increase in the central large carbide percentage. Therefore, there is an increase of carbide content by as much as 03% in total. This is due to dissolution of ultra-®ne and small carbides and their subsequent precipitation onto the central large and grain boundary carbides as reported before [1, 2]. In the cast materials the coarse austenite grains have boundaries which connect and run through the eutectic carbide pools. Frequently, the carbides on these boundaries are larger than others in the pools with the result that they tend to grow preferentially during coarsening, leading to the gradual disappearance of the eutectic pools and the development of strings of coarse carbides at sites provided by the original austenite grain boundaries. As the carbides in these strings grow, they tend to impinge on each other, leading to a polycrystalline network of coarse M6C carbides with MC. If such a mechanism is responsible for grain boundary eutectic carbide evolution, why should its percentage increase with time? However, it was further pointed out that there are denuded zones around grain boundary eutectic carbides which may be due to dissolution of small carbides at the vicinity of these carbides [2]. Therefore, they may be the feeding source for volume fraction change of these carbides during reheating. The occasionally coarse M6C carbides appeared to have grown from selected small carbides within the original grains that are probably those which are favourably situated at the grain corners in the ®ne grained austenite structure produced by transformation on reheating. This conclusion is supported by the observation that such particles are frequently polygonal in shape. Therefore, it is true to say that the carbide increase
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GHOMASHCHI: MICROSTRUCTURAL ANALYSIS OF HIGH SPEED STEEL
Fig. 6. Histogram distribution of MC grain boundary eutectic carbide particles for sand cast structure at 12008C.
GHOMASHCHI: MICROSTRUCTURAL ANALYSIS OF HIGH SPEED STEEL
5215
Fig. 7. Eect of reheating time on the MC grain boundary eutectic carbide particle size: (a) sand cast; (b) chill cast.
is due to dissolution of small carbides in the as-cast condition. Furthermore, such a conclusion may be drawn from Fig. 4 where as the number of small carbides drops to a minimum, the number of intragranular grown carbide remains steady with time. In other words, the small carbides are also the feeding source for the intragranular grown carbides. The volume fraction of small carbides was reported [2] to be about 3±4% which is compatible with the overall increase in volume fraction of the carbides, Fig. 2. 4.2. Spheroidization As mentioned before, the M6C grain boundary eutectic carbides spheroidize with time, and the resulting curve, Fig. 3, is similar to the previous studies of carbide spheroidization [4, 9]. The reason why spheroidization has appeared to stop after
about 100 h is attributed to the impingement of the already spheroidized carbide particles as reported before [1]. However, Fig. 3 may be used to estimate the beginning of the impingement mechanism which is, for instance, about 10 h at 12008C for the sand cast structure. This is in con®rmation of the qualitative analysis of microstructure. The mechanism of spheroidization is based on the edge-wise growth of channels in the M6C plates as was suggested qualitatively [10]. The activation energy for spheroidization could not be calculated since such a result, Fig. 3, is based on both quantitative and qualitative examination of microstructure. However, diusion of tungsten atoms in g-iron may be the rate controlling mechanism since diusion of the slowest species should control diusion needed for spheroidization [11, 12]. The large radius of tungsten atoms makes it dicult to diuse in g-Fe.
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Table 7. The mean arithmetic size of MC grain boundary eutectic carbide for sand cast structure (Q = 250±300 kJ/mol): (a) 12008C; (b) 11508C dMC (mm)
Specimen
Std. (mm)
H51 H50 H53 H55 H57 H41 H24 H26 H29 H30
0.61 0.63 0.65 0.77 0.81 0.99 1.27 1.45 1.39 1.74
0.19 0.16 0.35 0.32 0.31 0.41 0.42 0.54 0.60 0.61
H60 H62 H64 H76 H74 H43 H45 H47 H78 H80
0.36 0.43 0.44 0.43 0.53 0.65 0.79 0.84 0.94 0.93
0.06 0.07 0.10 0.09 0.15 0.20 0.25 0.31 0.37 0.32
(dMC)3 (mm3)
S.E. (mm) (a)
(b)
2DX (mm)
0.01 0.01 0.02 0.02 0.03 0.04 0.03 0.04 0.05 0.06
0.23 0.25 0.27 0.46 0.53 0.97 2.05 3.05 2.69 5.27
0.03 0.03 0.06 0.08 0.10 0.23 0.33 0.48 0.53 1.07
0.004 0.01 0.01 0.01 0.01 0.01 0.02 0.02 0.03 0.02
0.05 0.08 0.08 0.08 0.15 0.28 0.49 0.59 0.83 0.81
0.003 0.01 0.01 0.01 0.02 0.03 0.06 0.08 0.15 0.12
DX = error bar values in Fig. 7, Std. = standard deviation and S.E. = standard error.
4.3. Coarsening
where subscript M de®nes the element in question, OM is the atomic volume, VM =u, DM the diusion coecient, s the interfacial energy, SM the solubility, and u the stoichiometric factor depending on the composition of the precipitate. However, in 1981 Yong Wey et al. [13] modi®ed the above-mentioned equation for carbides in the Fe±C±M system where a carbide, (FeM)aCb, is formed:
4.3.1. MC. The MC carbide formed during the M2 C ÿ4M6 C MC transformation is almost fully spheroidized and thus, is mainly dominated by coarsening. Such a conclusion may easily be extracted from the histogram distribution graphs in Fig. 6, where after about 24 h, there are not many small particles, d00.30 mm, left. The coarsening of second phase particles as a result of reduction in interfacial energy was theoretically treated by Lifshitz and Slyozov, and Wagner [5±7] where the result was shown as r3 ÿ r30
8sDM SM OM t 9RT
r3 ÿ r30
8 a bsV y DgM U gM t 9aRT U yM ÿ U gM
4
where y represents the carbide and g is austenite. U is the concentration parameter de®ned by Hillert et
3
Table 8. The mean arithmetic size of M6C grain boundary eutectic carbide for sand cast structure (Q = 240±280 kJ/mol): (a) 12008C; (b) 11508C Specimen
dM6 C (mm)
Std. (mm)
H51 H50 H53 H55 H57 H41 H24 H26 H29 H30
1.02 0.87 1.01 1.53 1.33 1.60 1.77 2.07 2.09 2.24
0.39 0.21 0.28 0.64 0.42 0.55 0.49 0.67 0.76 0.75
H60 H62 H64 H76 H74 H43 H45 H47 H78 H80
0.60 0.67 0.67 0.67 0.85 1.18 1.41 1.41 1.60 1.57
0.16 0.21 0.20 0.18 0.24 0.39 0.43 0.46 0.55 0.59
S.E. (mm) (a)
(b)
dM6 C 3 (mm)3
2DX (mm3)
0.03 0.02 0.02 0.03 0.02 0.03 0.03 0.04 0.05 0.05
1.06 0.66 1.03 3.58 2.35 4.10 5.55 8.87 9.13 11.24
0.18 0.08 0.10 0.46 0.21 0.52 0.60 1.03 1.36 1.54
0.01 0.02 0.01 0.01 0.01 0.02 0.03 0.03 0.04 0.04
0.22 0.29 0.30 0.30 0.61 1.64 2.80 2.80 4.10 3.87
0.02 0.04 0.03 0.03 0.06 0.16 0.30 0.31 0.55 0.55
DX = error bar values in Fig. 9, Std. = standard deviation and S.E. = standard error.
GHOMASHCHI: MICROSTRUCTURAL ANALYSIS OF HIGH SPEED STEEL
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Fig. 8. Histogram distribution of M6C grain boundary eutectic carbide particles for sand cast structure at 12008C.
al. [14] as Ui Si = 1 ÿ Sc , where i is the alloy element with iron and carbon in the ternary system. The observed growth rate of the MC alloy carbide
is compared with the rate calculated from equation (4). Since the activation energy for MC carbide, Fig. 7, is close to the activation energy of
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GHOMASHCHI: MICROSTRUCTURAL ANALYSIS OF HIGH SPEED STEEL
Fig. 9. Eect of reheating time on the M6C grain boundary eutectic carbide particle size: (a) sand cast; (b) chill cast.
vanadium diusion in g-Fe, QgV =293 kJ/mol, the calculation of the rate constant in equation (4) was based on vanadium diusion as the rate controlling mechanism. The calculated K-value, Kcal=1.996 10ÿ6 mm3/s, see Appendix A, is close to the experimental results, K Sand exp: 12008C = =3.95 10ÿ6 2.26 10ÿ6 mm3/s and K Chill exp: 12008C 3 mm /s. Therefore, it may be concluded that, ®rst, the coarsening of MC carbide is volume-diusion controlled and, second, it is the diusion of vanadium in g-Fe which controls the growth rate. Finally, according to the theory of second-phase particle coarsening, the distribution curve cut-o size should be around 1.5 times the average size, r di =d 1:5. Although the experimental cut-o value (r), 1.6 < r < 1.8, is somewhat greater than the theoretically predicted one, it appears to be reasonable within the experimental scatter. 4.3.2. M6C. In contrast to MC, M6C carbide is initially dominated by morphological changes of spheroidization followed by coarsening. The histogram distribution graphs, Fig. 8, which are for the sand cast structure at 12008C may be utilized to describe the occurrence of spheroidization and coarsening and their subsequent eect on the microstructure. The introduction of smaller sizes in H50, 1 h at 12008C, and reduction of larger particles in comparison to H51, i.e. 0.5 h at 12008C, shows the dominance of spheroidization in H50, while formation of larger sizes in H53, i.e. 2 h at 12008C, may be due to coarsening. However, in some cases
the overlapping between spheroidization and coarsening makes it dicult to identify the domination of either process. As the soaking time increases, the state of the microstructure changes in such a way that coarsening will be the dominant process although spheroidization is still active. This is recognizable from the distribution histograms in which, some smaller particles formed even after 75 h at 12008C, H26. In order to show that coarsening is an active phenomenon within the large particles even at the early hours of treatment, the average size of the top 25% largest particles was plotted against time in Fig. 10. The higher coarsening rate at the early stages of thermal treatment may be further substantiated if Fig. 5 is considered more closely where the slope of the curve decreases continuously with time after about 4 h at 12008C. The activation energy obtained for the coarsening of M6C carbide from Fig. 9, i.e. 240±280 kJ/mol for sand and chill cast structures, is close to activation energies of self diusion of iron in austenite, i.e. QgFe =284 kJ/mol, and tungsten diusion in g-iron, i.e. QgW =267 kJ/mol. On the other hand, since M6C in the present investigation contains high percentages of both elements [1], i.e. Fe and W, it seems that diusion of both the iron and tungsten is necessary for the coarsening mechanism. According to Wagner [12], the diusion of the slowest moving species should be the rate controlling, and therefore it is the diusion of tungsten in g-iron, due to its larger atomic sizeÐa barrier for tungsten move-
GHOMASHCHI: MICROSTRUCTURAL ANALYSIS OF HIGH SPEED STEEL
5219
Fig. 10. Eect of reheating time on the size variation of the top 25% largest M6C grain boundary eutectic carbide for sand cast structure at 12008C. (Size intervals refer to the interval limits of the Carl Ziess Particle Size Analyser used in this study. They vary between 1 and 48 with equivalent sizes of 1.4±27.7 mm, respectively.)
ment, which controls the coarsening rate of M6C carbide. The above conclusion may be reached if the rate constant of M6C carbide during coarsening is considered. According to equation (4), it was 8C =1.36 10ÿ3 mm3/s while the found that K 1200 Fe rate constant for tungsten diusion was calculated 8C =3.36x10ÿ5 mm3/s, see Appendix A. as K 1200 W The experimental K-values obtained from d3±t graphs for sand and chill cast structures, i.e. ÿ5 8C 8C =1.11 K 1200 mm3/s and K 1200 sand =0.76 10 chill 10ÿ5 mm3/s, showed reasonable agreement with the rate constant obtained for tungsten diusion. Therefore, it might be concluded that the diusion rate of tungsten within g-iron is the rate controlling mechanism during coarsening. 5. CONCLUSIONS
It was shown that reheating of M2 grade high speed steel to temperatures of 1150±12008C brings about dissolution, spheroidization and coarsening of carbides. It is the diusion of tungsten in g-Fe which controls the spheroidization rate of M6C carbides, while diusion of vanadium and tungsten in g-Fe is the rate controlling mechanism for coarsening of MC and M6C carbides, respectively. The process of spheroidization of alloy carbides such as M6C, is somewhat dierent from that of cementite where impingement of the already spheroidized carbide particles make the analysis more dicult. Therefore, the resultant graph could be misleading where it appears that complete spheroidization is unachievable even after 200 h at 12008C. However, if the impinged carbide particles are considered as spheroids, complete spheroidization of M6C carbide plates may be reached after about 150±175 h at 12008C.
AcknowledgementsÐThe author would like to thank C. M. Sellars of Sheeld University for many useful suggestions and comments. REFERENCES 1. Ghomashchi, M. R. and Sellars, C. M., Metals Sci., 1984, 18, 44. 2. Ghomashchi, M. R., Z. Metallk., 1985, 76(10), 701. 3. Ghomashchi, M. R., Metall. Trans., 1985, 16A, 2341. 4. Chattopadhyay, S. and Sellars, C. M., Acta metall., 1982, 30, 157. 5. Lifschitz, I. M. and Slyosov, V. V., J. Phys. Chem. Solids, 1961, 19(1/2), 35. 6. Lifschitz, I. M. and Slyosov, V. V., Zh. eÂksp. teor. Fiz., 1958, 35, 218. 7. Wagner, C., Z. Elektrochem., 1961, 65, 581. 8. Haworth, C. W., Private communication, Sheeld University, 1983. 9. Chattopadhyay, S., Ph.D. thesis, University of Sheeld, 1978. 10. Nichols, F. A. and Mullins, W. W., Trans. Am. Inst. Min. Engrs, 1965, 233, 180. 11. Ghomashchi, M. R., Diusion in MaterialsÐAn Introduction (to be published). 12. Wagner, C., Unpublished results, 1969. 13. Yong Wey, M., Sakuma, T. and Nishizawa, T., Trans. Japan Inst. Metals, 1981, 22, 733. 14. Hillert, M., Nilsson, K. and Torndahl, L., J. Iron Steel Inst., 1971, 20, 949. 15. Ghomashchi, M. R., Ph.D. thesis, University of Sheeld, Jan. 1983. 16. Kramer, J. J., Pound, G. M. and Mehl, R. F., Acta metall., 1958, 6, 763. 17. Uhrenius, B. and Harvig, H., Metals Sci., 1975, 9, 67.
APPENDIX A Initially the particle sizes were plotted against time on a ln±ln scale and the slope was calculated. It was about 0.28 for sand cast at 12008C. The following equation was used for the purpose: d Ct1=n
A1
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GHOMASHCHI: MICROSTRUCTURAL ANALYSIS OF HIGH SPEED STEEL
where C is a constant and d, t and n are particle size, time and diusion exponent, respectively. For n = 2, 3, 4, 5, the diusion is interface, volume, grain boundary and dislocation controlled, respectively. Equation (A1) was then rearranged as 1 A2 ln d ln C ln t: n The above check has supported the bulk or volume diusion during coarsening of both MC and M6C carbides. The K-values were then calculated using the following equation as explained in the text: 8 a bsV y DgV U gV r3 ÿ r30 t: 9aRT U yV ÿ U gV 2
APPENDIX B Nomenclature MC
rMC Vy
Vy a=b=1
density of MC carbide = 7.84 Mg/m3 [15] molar volume of MC carbide = (weight of 1 mol)/(density) 343.007/7.84 = 43.75 10ÿ6 m3/mol for MC carbide
s DgV DgV
interfacial energy for ferrite/cementite interface = 0.7 J/m2 [16] diusion coecient of vanadium in austenite D0 exp ÿQ=RT
D0 gV
365 mm2/s [8]
QgV
activation energy of vanadium diusion in g-Fe = 293 kJ/mol [8] 1.0 (0.016 + 0.016/20 16) = 2.88 10ÿ2 (8(1 + 1) 0.7 J/m2 43.75 10ÿ6 m3/mol 1.47 10ÿ14 m2/s)/(9 1 8.31 J/(mol K) 1473 K(1 ÿ 2.88 10ÿ2)2 1.996 10ÿ6 mm3/s.
UyV UgV K K
A3 M6C
rM6 C Vy a=6 b=1 g D0 QgW
UgW 10008C UyW 10008C K K
10.29 [15] 8424.65/10.29 = 818.72 10ÿ6 m3/mol ferrite/cementite = 0.7 J/m2 [16] 13 mm2/s [8] activation energy of tungsten diusion in g-Fe = 267 kJ/mol [8] 0.0040 [17] 0.163 [17] (8(6 + 1)0.7 818.72 10ÿ6 4.37 10ÿ15 0.0040)/(9 6 8.31 1473(0.163 ÿ 0.0040)2 3.36 10ÿ23 m3/s = 3.36 10ÿ5 mm3/s.
Acta mater. Vol. 46, No. 14, pp. 5207±5220, 1998 # 1998 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved Printed in Great Britain S1359-6454(98)00110-4 1359-6454/98 $19.00 + 0.00
QUANTITATIVE MICROSTRUCTURAL ANALYSIS OF M2 GRADE HIGH SPEED STEEL DURING HIGH TEMPERATURE TREATMENT M. R. GHOMASHCHI School of Engineering, University of South Australia, The Levels, SA 5095, Australia (Received 17 September 1997; accepted 13 February 1998) AbstractÐThe microstructural changes of M2 grade high speed steel were examined quantitatively to identify the mechanisms responsible for changes during high temperature treatment. Reheating at high temperature leads to major microstructural changes involving phase transformation, spheroidization and coarsening of carbides formed in the as-cast M2 grade high speed steel. The M2C carbide decomposes rapidly to form MC and M6C in less than 1 h after attaining the reheating temperatures of 1150 or 12008C. Further reheating brings about the eutectic M6C carbide spheroidization and general coarsening of both M6C and MC. Carbide particles situated at the original austenite grain boundaries grow preferentially until impingement of carbides along these boundaries takes place. It is the diusion of tungsten in g-Fe which controls spheroidization rate of M6C carbides, while diusion of vanadium and tungsten in g-Fe is the rate controlling mechanism for coarsening of MC and M6C carbides, respectively. The process of spheroidization of alloy carbides such as M6C, is somewhat dierent from that of cementite where impingement of the already spheroidized carbide particles make the analysis more dicult. # 1998 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved.
1. INTRODUCTION
It has been reported by the present author [1] that the solidi®cation of M2 grade high speed steel leads to the formation of coarse M2C particles at the centre of dendrites as a result of a peritectic decomposition of d-ferrite. Subsequent enrichment of solute atoms around the growing austenite dendrites leads to the formation of coarse interdendritic eutectic carbides of MC, M2C and M6C. Further cooling below the solidus results in general precipitation of small carbides of M2C and MC. The eutectic and small carbides were then reported to have dierent morphologies and chemistry [2, 3]. In addition some ultra-®ne carbides were also detected [1]. Reheating leads to rapid dissolution of the ultra®ne carbides. The M2C particles decompose rapidly to form MC and M6C in less than 1 h after attaining the reheating temperature of 1150 or 12008C. Further reheating brings about the eutectic M6C carbide spheroidization and general coarsening of both M6C and MC. Carbide particles situated at the original austenite grain boundaries grow preferentially until impingement of carbides along these boundaries takes place. The majority of small carbides dissolve as coarsening occurs, but occasional M6C particles grow to sizes equivalent to those of {The same author reported the eect of hot forging on carbides in M2 grade high speed steel. Interested readers are referred to Metall. Trans., 1993, 24A, 2171.
the original central large or eutectic carbides. These newly-formed carbides are called intragranular grown carbides. Figure 1 illustrates the original cast structures and the resultant structures due to reheating at high temperatures. The schematic diagram included in Fig. 1 speci®es the nomenclature used in this investigation. In all previous publications [1±3] a precise qualitative picture was given for carbide formation during solidi®cation and their subsequent changes due to reheating. The present article is designed to characterize the microstructural changes quantitatively for better understanding of such a complex system and to provide insights for all high speed steels.{ 2. MATERIALS AND EXPERIMENTAL PROCEDURES
The alloy used in this study was prepared in the laboratory by air melting using ARMCO iron as the base metal with the addition of white cast-iron, ferrotungsten, ferrovanadium, pure chromium and molybdenum as the alloy elements. Silicon was added for deoxidization and casting was carried out at a temperature of 15608C in two sand and chill moulds. The size of the sand and chill cast ingots were 292 75 24 and 420 78 38 mm3, respectively. These steels will be referred to as sand and chill cast and their chemical analysis is as follows (wt%): 0.86 C, 6.85 W, 5.3 Mo, 4.22 Cr, 1.96 V, 0.25 Si and 0.07 Mn, (Al, Nb, Ti, Co R 0.02).
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GHOMASHCHI: MICROSTRUCTURAL ANALYSIS OF HIGH SPEED STEEL
Table 1. The list of specimens with their soaking time and temperature Sand
Chill
Time (h)
12008C
11508C
12008C
11508C
H51 H50 H53 H55 H57 H41 H24 H26 H22 H32 H29 H34 H30
H60 H62 H64 H76 H74 H43 H45 H47 H48 Ð H78 Ð H80
H52 Ð H54 H56 H58 H40 H23 H25 H27 H31 H28 H33 H35
H61 H63 H65 H77 H75 H42 H44 H46 H49 Ð H79 Ð H81
0.5 1 2 4 10 25 50 75 100 125 150 175 200
Specimen preparation, thermal treatments and various etching techniques together with dierent etchants and microscopes used for dierent purposes are similar to those reported before [1]. Table 1 summarizes reheating times and temperatures employed during the course of the present investigation. In order to characterize microstructural features in the as-cast and heat treated conditions, the following quantitative metallographic techniques were employed. (i) Point counting. Systematic point counting was carried out on parallel lines on the plane of polish with a light microscope with a travelling stage connected to a digital counter. The stage moved in 0.05 mm increments. A total of 1200 points was counted for each specimen to reduce the statistical error to less than 1%. (ii) Size distribution. A semi-automatic Carl Ziess particle size analyser was used in conjunction with light micrographs taken from specimens etched electrolytically in chromic acid solution for size distribution measurements. The size evaluation is based on equivalent circle area diameter, ECAD; that is the diameter of a circle having the same area as the object to be measured. A total of 250 particles was measured for each specimen and the arithmetic mean particle size d was calculated. In addition, both light and scanning electron microscopes were used to measure the number of particles per unit area. For light microscopy, two magni®cations of 1600 and 200 times were used for counting the number of small and intragranular grown carbides, respectively. The size criterion for small carbides was based on the ability of the light microscope to distinguish between two individual particles, i.e. resolution. The microscope resolving power was improved by application of immersion oil, d 1= 2m sin a, where the resolving power, d, was improved to 0.2 mm. The intragranular grown carbides, however, were considered dierently, where the 200 times magni®cation made them easy
to recognize in a background of ®ne carbides uniformly distributed within the grains. Any particles with dimensions of about 1 mm at 200 times were counted as intragranular grown carbides. A total number of ten randomly selected ®elds of view were examined. In order to measure the number of grain boundary eutectic carbides, SEM was used at 5000 times magni®cation. A total number of 30 ®elds of view was selected randomly from each specimen. 3. RESULTS
The microstructures of both sand and chill cast specimens were similar with chill cast being slightly ®ner. Furthermore, both responded equally to high temperature treatments although the kinetics of reactions were slightly faster for chill cast. Therefore, the following results are presented mainly for sand cast specimens. 3.1. Volume fraction As has been reported before [1], the carbides may be divided into the following categories. 3.1.1. Grain boundary eutectic carbides. The ascast structures contain M2C, MC and M6C eutectic carbides where the M2C portion transforms to M6C and MC within 30 min of reheating at 1150± 12008C. Therefore, the volume fractions of M6C and MC carbides were measured and recorded in Table 2. As is evident, there is a general trend towards increasing the percentage of grain boundary eutectic carbides with soaking times at 1150 and 12008C for both sand and chill cast structures. However, after a certain time it appears to have reached an equilibrium stage, see Fig. 2. 3.1.2. Spheroidization. Carbide plates are not stable at high temperatures and break down into smaller segments, i.e. spheroidization. The grain boundary eutectic carbides were examined very closely and as reported earlier [1], spheroidization occurs only of M6C carbides, which is then followed by coarsening and impingement processes active within the already spheroidized M6C particles. Therefore, to study spheroidization, the following measurements were carried out on M6C carbides. (a) The amount of spheroidized M6C carbide was obtained for only isolated carbide particles Table 2. Volume fraction of grain boundary eutectic carbides for sand cast specimens reheated at 12008C followed by water quench Specimen H50 H53 H41 H24 H26 H29 H30
Vv% MC
2DVv%
Vv% M6C
2DVv%
1.6 1.7 2.1 2.1 2.4 1.9 2.0
0.3 0.3 0.3 0.4 0.4 0.3 0.3
9.1 9.5 8.4 10.3 10.3 10.2 9.2
0.6 0.7 0.6 0.7 0.7 0.7 0.7
GHOMASHCHI: MICROSTRUCTURAL ANALYSIS OF HIGH SPEED STEEL
Fig. 1. Optical micrographs of the as-cast and reheated M2 grade high speed steel: (a) as-cast sand structure; (b) reheated sand cast structure, 150 h at 11508C, water quenched; (c) schematic diagram to show the nomenclature for carbides.
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GHOMASHCHI: MICROSTRUCTURAL ANALYSIS OF HIGH SPEED STEEL
having an aspect ratio R5. If there are N1 plates, N2 necked plates and N3 isolated M6C particles, the percentage spheroidized M6C carbide is N3 100: 1 Sph%M6 C N1 N2 N3 Table 3(a) shows the results for sand cast structure at 12008C which are plotted in Fig. 3. As can be seen, the resultant graph is similar to conventional spheroidization of cementite, for instance, where percentage spheroidized cementite increases with time [4]. However, the process of spheroidization appears to have stopped at about 80% spheroidized M6C. (b) It was further decided to count the necked plates as spheroids due to impingement eect, provided that the aspect ratio of the impinged particle is within the set criterion, i.e. AR R 5. The percentage spheroidization is then calculated as N2 N3 Sph%M6 C 100: 2 N1 N2 N3 Table 3(b) shows the new results which are superimposed on Fig. 3. The percentage of impinged or necked particles was then calculated from equations (1) and (2) and plotted in Fig. 3. 3.1.3. Central large carbides. The central large carbides form due to a complex peritectic reaction at around 13308C where part of the d-ferrite within the peritectic reaction transforms to a low carbon austenite and carbide{ 0 1 L d ÿ4g @ d ÿ4g M2 C A L d ÿ4g M2 C and are mainly M2C in the as-cast state, but as for grain boundary eutectic carbides they too transform to M6C and MC within the ®rst 30 min of reheating. The amount of central large carbides increases with increasing reheating time, see Table 4 and Fig. 2. 3.1.4. Intragranular grown carbides. A few M6C carbide particles grow preferentially within the initial austenite dendrites where their percentage increases initially with time but remains steady after a speci®c time, depending on the reheating temperature and cast structure [1]. The results are shown in Table 4 and Fig. 2.
Table 3. To show the eect of soaking time on the percentage spheroidization in M6C grain boundary eutectic carbide (a) Volume fraction of isolated M6C grain boundary eutectic carbide particles with AR R5 N3 2DVv Specimen Vv sph% 100 N1 N2 N3 H53 H57 H41 H24 H26 H29 H30
58.9 76.2 70.9 75.5 78.8 383 84.2
5.7 4.4 5 4.2 3.9 3.6 3.6
(b) Volume fraction of isolated and impinged M6C carbide particles with AR R 5 N2 N3 2DVv Specimen Vv sph% 100 N1 N2 N3 H57 H41 H24 H26 H32 H29 H30
78.3 89.2 91.9 96.1 96.7 99 100
4 3.2 2.5 1.7 1.6 0.9 0
during early stages of heat treatment which is concurrent with general dissolution of MC and growth of occasional M6C. The NA is a useful means to study both MC dissolution and M6C growth. It was found that the small carbides disappear after a certain period of time depending upon the cast structure and soaking temperature, where ®ner cast structure and higher soaking temperatures are recipes for greater dissolution and coarsening rates. Figure 4 shows the results for sand cast structure. It is noticeable, in Fig. 4, that the ®nal stages of small carbides dissolution is concurrent with a steady state for the number of intergranular grown carbides. The results are shown in Table 5. 3.2.2. Grain boundary eutectic carbides. The number of grain boundary eutectic carbides was measured with the view that impinged carbides
3.2. Number of particles per unit area (NA) 3.2.1. Small carbides. It has been shown that the small carbides consist of M2C needles and MC cubes [2]. The M2C transforms to M6C and MC
{For more information, see Ref. [1] and J. McLaughlin et al., Metall. Trans., 1977, 8A, 1787.
Fig. 2. Eect of reheating time on the volume fraction of carbides.
GHOMASHCHI: MICROSTRUCTURAL ANALYSIS OF HIGH SPEED STEEL
5211
Table 4. Volume fraction of central large and intragranular grown carbides for sand cast specimens at 12008C Specimen
H53 H57 H41 H24 H26 H29 H30
Intragranular grown carbides
Central large carbides
Vv%
2DVv%
Vv% M6C
2DVv%
0.4 0.7 1.2 1.5 1.9 1.7 1.0
0.2 0.2 0.3 0.4 0.4 0.4 0.3
0.3 0.9 1.2 1.4 1.3 2 1.5
0.2 0.3 0.3 0.4 0.4 0.4 0.4
were counted as isolated ones after impingement. The results are given in Table 6 and plotted in Fig. 5. After an initial rise, the number of eutectic carbides decreases to an almost constant value after about 75 h at temperature. 3.3. Size distribution Size distribution analysis was carried out on MC and M6C grain boundary eutectic carbides to characterize both spheroidization and coarsening phenomena. However, since MC carbide particles were almost fully spheroidized both in the as-cast and reheated structures, coarsening appears to be the only active process within MC carbide. 3.3.1. MC. The size values for MC carbide were arranged into a number of equal valued class intervals and plotted against the relative frequency of each respective class interval. The resulting distribution graphs are presented as histograms from which the arithmetic mean particle size was calculated for each specimen. The distribution histograms are almost the same for both sand and chill cast steels at 1150 and 12008C, and thus the sand cast structure at 12008C is presented here, Fig. 6. The X-axis in Fig. 6 is both the ratio of each inter i.e. val's size (di) over the arithmetic mean size (d), and the actual size of each interval. As r di =d, soaking time increases, ®rst, the number of smaller particles decreases, and second, some larger particles are introduced into the distribution graphs. The cut-o value is generally in the range of 1.6 < r < 1.8, although according to the theoretical treatments [5±7], it should be about 1.5. The mechanism responsible for coarsening may be deduced from a logarithmic plot of mean particle size vs time according to the theory of coarsening, i.e. dn ÿ dn0 Kt where d and d0 are the mean particle sizes at time t and zero, respectively. K is the rate constant and n the mode of coarsening, where for n = 2, 3, 4, 5, coarsening is dominated by interface, bulk or volume, grain boundary and pipe (dislocation) diusion, respectively. In order to estimate the activation energy for coarsening of the MC grain boundary eutectic carbide, the cubes of mean {See Appendix A. {See Section 4.3 for more details.
Fig. 3. Eect of reheating time on the percentage spheroidization of M6C grain boundary eutectic carbides.
sizes{ obtained from the histograms were plotted against time for both the sand and chill cast structures at 1200 and 11508C. Figure 7 shows the result for sand cast only. The activation energy was determined by calculating the slope of d3±t curves, i.e. a @ d3 =@ t, at t>60 h, in which a is equivalent to the rate constant, K. Since K is both temperature and diusion dependent [5±7],{ a may be presented in terms of an Arrhenius equation, i.e.
Table 5. The number of small, and intragranular grown carbides per 1 mm2 for sand cast structure at: (a) 12008C; (b) 11508C Small carbides$ Specimen
NA
H51 H50 H53 H55 H57 H41 H24 H26 H22 H32 H29 H34 H30
438 400 453 120 368 000 342 400 292 480 228 480 245 760 53 120 0 0 0 0 0
H60 H62 H64 H76 H74 H43 H45 H47 H48 H78 H80
766 080 691 840 604 160 560 640 538 240 533 760 477 440 416 000 366 720 167 680 151 040
2S.E. (a) 20 480 35 840 18 560 18 560 26 880 19 200 25 920 20 480 0 0 0 0 0 (b) 29 120 49 600 39 040 33 920 36 480 46 080 18 880 20 800 21 120 19 840 16 320
Intragranular grown carbides$ NA
2S.E.
50 160 160 220 300 440 566 630 390 520 1080 380 460
48 72 92 120 88 124 152 224 88 128 232 128 148
30 0 170 90 80 420 410 570 576 830 710
30 0 72 48 64 152 92 104 68 100 96
$All measurements are based on 95% con®dence limit (NA22S.E.) as plotted in Fig. 4 (S.E. = standard error).
Fig. 4. Eect of reheating time on the dissolution of small carbides and precipitation of intragranular grown carbides for sand cast structure: (a) 11508C; (b) 12008C.
5212 GHOMASHCHI: MICROSTRUCTURAL ANALYSIS OF HIGH SPEED STEEL
GHOMASHCHI: MICROSTRUCTURAL ANALYSIS OF HIGH SPEED STEEL
5213
Table 6. The number of grain boundary eutectic carbides per 1 mm2 for sand cast structure at: (a) 12008C; (b) 11508C Specimen Mean (NA) Std. 2S.E. Specimen Mean (NA) Std. 2S.E.
H51
H50
H53
H55
H57
30 000 18 500 10 000
40 500 34 500 18 000
62 500 33 500 17 000
60 500 27 000 14 000
46 500 18 500 10 000
H60
H62
H64
H76
H74
30 000 20 000 10 000
24 000 12 000 6000
49 500 32 000 17 000
70 500 38 500 23 000
74 000 45 000 23 000
(a)
(b)
H41
H24
H26
H32
H34
H30
45 000 15 000 8000
55 500 22 000 12 000
36 500 15 500 8000
34 500 10 000 5000
36 500 12 500 7000
38 500 19 500 10 000
H43
H45
H47
H78
H80
49 000 23 500 12 000
51 000 33 000 17 000
39 000 22 500 12 000
49 000 17 500 9000
58 000 19 500 10 000
$All measurements are based on 95% con®dence limit (NA22S.E.) as plotted in Fig. 5 (S.E. = standard error).
a b eÿQ=RT , from which the activation energy may be calculated. The activation energy of MC carbide together with the arithmetic mean particle sizes and their cubes are given in Table 7. The activation energy for both sand and chill cast structures is very close, i.e. 250±300 kJ/mol. This is close to the activation energy for vanadium diusion in g-iron, i.e. QgV =293 kJ/mol [8]. 3.3.2. M6C. The M6C carbide, unlike MC, is more complicated since it is exposed to dierent processes of spheroidization, coarsening and impingement. The sizing of the M6C particles, Table 8, was carried out the same as for MC, and the impinged particles were counted as separate particles. Figure 8 shows the histogram distribution of M6C carbide particles at 12008C for sand cast structure. Close examination of the histograms at various times reveals the eectiveness of spheroidization and coarsening processes as for instance H55 and H57 for spheroidization and H41 and H24 for coarsening. The cut-o size is between 1.6 < r < 1.8 and an activation energy of 240± 280 kJ/mol was calculated for both sand and chill cast structures, see Fig. 9. In order to study the eect of coarsening separately, the top 25% largest M6C particles were plotted against time in Fig. 10. Coarsening is
Fig. 5. Eect of reheating time on the number of grain boundary eutectic carbides.
expected as the dominant phenomenon within these carbide particles, and the rate of coarsening is faster in early stages of treatment.
4. DISCUSSION
4.1. Volume fraction of carbides Figure 2 shows that the volume fraction of grain boundary eutectic carbides increases with time. There are also the occasionally grown intragranular grown carbides and a slight increase in the central large carbide percentage. Therefore, there is an increase of carbide content by as much as 03% in total. This is due to dissolution of ultra-®ne and small carbides and their subsequent precipitation onto the central large and grain boundary carbides as reported before [1, 2]. In the cast materials the coarse austenite grains have boundaries which connect and run through the eutectic carbide pools. Frequently, the carbides on these boundaries are larger than others in the pools with the result that they tend to grow preferentially during coarsening, leading to the gradual disappearance of the eutectic pools and the development of strings of coarse carbides at sites provided by the original austenite grain boundaries. As the carbides in these strings grow, they tend to impinge on each other, leading to a polycrystalline network of coarse M6C carbides with MC. If such a mechanism is responsible for grain boundary eutectic carbide evolution, why should its percentage increase with time? However, it was further pointed out that there are denuded zones around grain boundary eutectic carbides which may be due to dissolution of small carbides at the vicinity of these carbides [2]. Therefore, they may be the feeding source for volume fraction change of these carbides during reheating. The occasionally coarse M6C carbides appeared to have grown from selected small carbides within the original grains that are probably those which are favourably situated at the grain corners in the ®ne grained austenite structure produced by transformation on reheating. This conclusion is supported by the observation that such particles are frequently polygonal in shape. Therefore, it is true to say that the carbide increase
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GHOMASHCHI: MICROSTRUCTURAL ANALYSIS OF HIGH SPEED STEEL
Fig. 6. Histogram distribution of MC grain boundary eutectic carbide particles for sand cast structure at 12008C.
GHOMASHCHI: MICROSTRUCTURAL ANALYSIS OF HIGH SPEED STEEL
5215
Fig. 7. Eect of reheating time on the MC grain boundary eutectic carbide particle size: (a) sand cast; (b) chill cast.
is due to dissolution of small carbides in the as-cast condition. Furthermore, such a conclusion may be drawn from Fig. 4 where as the number of small carbides drops to a minimum, the number of intragranular grown carbide remains steady with time. In other words, the small carbides are also the feeding source for the intragranular grown carbides. The volume fraction of small carbides was reported [2] to be about 3±4% which is compatible with the overall increase in volume fraction of the carbides, Fig. 2. 4.2. Spheroidization As mentioned before, the M6C grain boundary eutectic carbides spheroidize with time, and the resulting curve, Fig. 3, is similar to the previous studies of carbide spheroidization [4, 9]. The reason why spheroidization has appeared to stop after
about 100 h is attributed to the impingement of the already spheroidized carbide particles as reported before [1]. However, Fig. 3 may be used to estimate the beginning of the impingement mechanism which is, for instance, about 10 h at 12008C for the sand cast structure. This is in con®rmation of the qualitative analysis of microstructure. The mechanism of spheroidization is based on the edge-wise growth of channels in the M6C plates as was suggested qualitatively [10]. The activation energy for spheroidization could not be calculated since such a result, Fig. 3, is based on both quantitative and qualitative examination of microstructure. However, diusion of tungsten atoms in g-iron may be the rate controlling mechanism since diusion of the slowest species should control diusion needed for spheroidization [11, 12]. The large radius of tungsten atoms makes it dicult to diuse in g-Fe.
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GHOMASHCHI: MICROSTRUCTURAL ANALYSIS OF HIGH SPEED STEEL
Table 7. The mean arithmetic size of MC grain boundary eutectic carbide for sand cast structure (Q = 250±300 kJ/mol): (a) 12008C; (b) 11508C dMC (mm)
Specimen
Std. (mm)
H51 H50 H53 H55 H57 H41 H24 H26 H29 H30
0.61 0.63 0.65 0.77 0.81 0.99 1.27 1.45 1.39 1.74
0.19 0.16 0.35 0.32 0.31 0.41 0.42 0.54 0.60 0.61
H60 H62 H64 H76 H74 H43 H45 H47 H78 H80
0.36 0.43 0.44 0.43 0.53 0.65 0.79 0.84 0.94 0.93
0.06 0.07 0.10 0.09 0.15 0.20 0.25 0.31 0.37 0.32
(dMC)3 (mm3)
S.E. (mm) (a)
(b)
2DX (mm)
0.01 0.01 0.02 0.02 0.03 0.04 0.03 0.04 0.05 0.06
0.23 0.25 0.27 0.46 0.53 0.97 2.05 3.05 2.69 5.27
0.03 0.03 0.06 0.08 0.10 0.23 0.33 0.48 0.53 1.07
0.004 0.01 0.01 0.01 0.01 0.01 0.02 0.02 0.03 0.02
0.05 0.08 0.08 0.08 0.15 0.28 0.49 0.59 0.83 0.81
0.003 0.01 0.01 0.01 0.02 0.03 0.06 0.08 0.15 0.12
DX = error bar values in Fig. 7, Std. = standard deviation and S.E. = standard error.
4.3. Coarsening
where subscript M de®nes the element in question, OM is the atomic volume, VM =u, DM the diusion coecient, s the interfacial energy, SM the solubility, and u the stoichiometric factor depending on the composition of the precipitate. However, in 1981 Yong Wey et al. [13] modi®ed the above-mentioned equation for carbides in the Fe±C±M system where a carbide, (FeM)aCb, is formed:
4.3.1. MC. The MC carbide formed during the M2 C ÿ4M6 C MC transformation is almost fully spheroidized and thus, is mainly dominated by coarsening. Such a conclusion may easily be extracted from the histogram distribution graphs in Fig. 6, where after about 24 h, there are not many small particles, d00.30 mm, left. The coarsening of second phase particles as a result of reduction in interfacial energy was theoretically treated by Lifshitz and Slyozov, and Wagner [5±7] where the result was shown as r3 ÿ r30
8sDM SM OM t 9RT
r3 ÿ r30
8 a bsV y DgM U gM t 9aRT U yM ÿ U gM
4
where y represents the carbide and g is austenite. U is the concentration parameter de®ned by Hillert et
3
Table 8. The mean arithmetic size of M6C grain boundary eutectic carbide for sand cast structure (Q = 240±280 kJ/mol): (a) 12008C; (b) 11508C Specimen
dM6 C (mm)
Std. (mm)
H51 H50 H53 H55 H57 H41 H24 H26 H29 H30
1.02 0.87 1.01 1.53 1.33 1.60 1.77 2.07 2.09 2.24
0.39 0.21 0.28 0.64 0.42 0.55 0.49 0.67 0.76 0.75
H60 H62 H64 H76 H74 H43 H45 H47 H78 H80
0.60 0.67 0.67 0.67 0.85 1.18 1.41 1.41 1.60 1.57
0.16 0.21 0.20 0.18 0.24 0.39 0.43 0.46 0.55 0.59
S.E. (mm) (a)
(b)
dM6 C 3 (mm)3
2DX (mm3)
0.03 0.02 0.02 0.03 0.02 0.03 0.03 0.04 0.05 0.05
1.06 0.66 1.03 3.58 2.35 4.10 5.55 8.87 9.13 11.24
0.18 0.08 0.10 0.46 0.21 0.52 0.60 1.03 1.36 1.54
0.01 0.02 0.01 0.01 0.01 0.02 0.03 0.03 0.04 0.04
0.22 0.29 0.30 0.30 0.61 1.64 2.80 2.80 4.10 3.87
0.02 0.04 0.03 0.03 0.06 0.16 0.30 0.31 0.55 0.55
DX = error bar values in Fig. 9, Std. = standard deviation and S.E. = standard error.
GHOMASHCHI: MICROSTRUCTURAL ANALYSIS OF HIGH SPEED STEEL
5217
Fig. 8. Histogram distribution of M6C grain boundary eutectic carbide particles for sand cast structure at 12008C.
al. [14] as Ui Si = 1 ÿ Sc , where i is the alloy element with iron and carbon in the ternary system. The observed growth rate of the MC alloy carbide
is compared with the rate calculated from equation (4). Since the activation energy for MC carbide, Fig. 7, is close to the activation energy of
5218
GHOMASHCHI: MICROSTRUCTURAL ANALYSIS OF HIGH SPEED STEEL
Fig. 9. Eect of reheating time on the M6C grain boundary eutectic carbide particle size: (a) sand cast; (b) chill cast.
vanadium diusion in g-Fe, QgV =293 kJ/mol, the calculation of the rate constant in equation (4) was based on vanadium diusion as the rate controlling mechanism. The calculated K-value, Kcal=1.996 10ÿ6 mm3/s, see Appendix A, is close to the experimental results, K Sand exp: 12008C = =3.95 10ÿ6 2.26 10ÿ6 mm3/s and K Chill exp: 12008C 3 mm /s. Therefore, it may be concluded that, ®rst, the coarsening of MC carbide is volume-diusion controlled and, second, it is the diusion of vanadium in g-Fe which controls the growth rate. Finally, according to the theory of second-phase particle coarsening, the distribution curve cut-o size should be around 1.5 times the average size, r di =d 1:5. Although the experimental cut-o value (r), 1.6 < r < 1.8, is somewhat greater than the theoretically predicted one, it appears to be reasonable within the experimental scatter. 4.3.2. M6C. In contrast to MC, M6C carbide is initially dominated by morphological changes of spheroidization followed by coarsening. The histogram distribution graphs, Fig. 8, which are for the sand cast structure at 12008C may be utilized to describe the occurrence of spheroidization and coarsening and their subsequent eect on the microstructure. The introduction of smaller sizes in H50, 1 h at 12008C, and reduction of larger particles in comparison to H51, i.e. 0.5 h at 12008C, shows the dominance of spheroidization in H50, while formation of larger sizes in H53, i.e. 2 h at 12008C, may be due to coarsening. However, in some cases
the overlapping between spheroidization and coarsening makes it dicult to identify the domination of either process. As the soaking time increases, the state of the microstructure changes in such a way that coarsening will be the dominant process although spheroidization is still active. This is recognizable from the distribution histograms in which, some smaller particles formed even after 75 h at 12008C, H26. In order to show that coarsening is an active phenomenon within the large particles even at the early hours of treatment, the average size of the top 25% largest particles was plotted against time in Fig. 10. The higher coarsening rate at the early stages of thermal treatment may be further substantiated if Fig. 5 is considered more closely where the slope of the curve decreases continuously with time after about 4 h at 12008C. The activation energy obtained for the coarsening of M6C carbide from Fig. 9, i.e. 240±280 kJ/mol for sand and chill cast structures, is close to activation energies of self diusion of iron in austenite, i.e. QgFe =284 kJ/mol, and tungsten diusion in g-iron, i.e. QgW =267 kJ/mol. On the other hand, since M6C in the present investigation contains high percentages of both elements [1], i.e. Fe and W, it seems that diusion of both the iron and tungsten is necessary for the coarsening mechanism. According to Wagner [12], the diusion of the slowest moving species should be the rate controlling, and therefore it is the diusion of tungsten in g-iron, due to its larger atomic sizeÐa barrier for tungsten move-
GHOMASHCHI: MICROSTRUCTURAL ANALYSIS OF HIGH SPEED STEEL
5219
Fig. 10. Eect of reheating time on the size variation of the top 25% largest M6C grain boundary eutectic carbide for sand cast structure at 12008C. (Size intervals refer to the interval limits of the Carl Ziess Particle Size Analyser used in this study. They vary between 1 and 48 with equivalent sizes of 1.4±27.7 mm, respectively.)
ment, which controls the coarsening rate of M6C carbide. The above conclusion may be reached if the rate constant of M6C carbide during coarsening is considered. According to equation (4), it was 8C =1.36 10ÿ3 mm3/s while the found that K 1200 Fe rate constant for tungsten diusion was calculated 8C =3.36x10ÿ5 mm3/s, see Appendix A. as K 1200 W The experimental K-values obtained from d3±t graphs for sand and chill cast structures, i.e. ÿ5 8C 8C =1.11 K 1200 mm3/s and K 1200 sand =0.76 10 chill 10ÿ5 mm3/s, showed reasonable agreement with the rate constant obtained for tungsten diusion. Therefore, it might be concluded that the diusion rate of tungsten within g-iron is the rate controlling mechanism during coarsening. 5. CONCLUSIONS
It was shown that reheating of M2 grade high speed steel to temperatures of 1150±12008C brings about dissolution, spheroidization and coarsening of carbides. It is the diusion of tungsten in g-Fe which controls the spheroidization rate of M6C carbides, while diusion of vanadium and tungsten in g-Fe is the rate controlling mechanism for coarsening of MC and M6C carbides, respectively. The process of spheroidization of alloy carbides such as M6C, is somewhat dierent from that of cementite where impingement of the already spheroidized carbide particles make the analysis more dicult. Therefore, the resultant graph could be misleading where it appears that complete spheroidization is unachievable even after 200 h at 12008C. However, if the impinged carbide particles are considered as spheroids, complete spheroidization of M6C carbide plates may be reached after about 150±175 h at 12008C.
AcknowledgementsÐThe author would like to thank C. M. Sellars of Sheeld University for many useful suggestions and comments. REFERENCES 1. Ghomashchi, M. R. and Sellars, C. M., Metals Sci., 1984, 18, 44. 2. Ghomashchi, M. R., Z. Metallk., 1985, 76(10), 701. 3. Ghomashchi, M. R., Metall. Trans., 1985, 16A, 2341. 4. Chattopadhyay, S. and Sellars, C. M., Acta metall., 1982, 30, 157. 5. Lifschitz, I. M. and Slyosov, V. V., J. Phys. Chem. Solids, 1961, 19(1/2), 35. 6. Lifschitz, I. M. and Slyosov, V. V., Zh. eÂksp. teor. Fiz., 1958, 35, 218. 7. Wagner, C., Z. Elektrochem., 1961, 65, 581. 8. Haworth, C. W., Private communication, Sheeld University, 1983. 9. Chattopadhyay, S., Ph.D. thesis, University of Sheeld, 1978. 10. Nichols, F. A. and Mullins, W. W., Trans. Am. Inst. Min. Engrs, 1965, 233, 180. 11. Ghomashchi, M. R., Diusion in MaterialsÐAn Introduction (to be published). 12. Wagner, C., Unpublished results, 1969. 13. Yong Wey, M., Sakuma, T. and Nishizawa, T., Trans. Japan Inst. Metals, 1981, 22, 733. 14. Hillert, M., Nilsson, K. and Torndahl, L., J. Iron Steel Inst., 1971, 20, 949. 15. Ghomashchi, M. R., Ph.D. thesis, University of Sheeld, Jan. 1983. 16. Kramer, J. J., Pound, G. M. and Mehl, R. F., Acta metall., 1958, 6, 763. 17. Uhrenius, B. and Harvig, H., Metals Sci., 1975, 9, 67.
APPENDIX A Initially the particle sizes were plotted against time on a ln±ln scale and the slope was calculated. It was about 0.28 for sand cast at 12008C. The following equation was used for the purpose: d Ct1=n
A1
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GHOMASHCHI: MICROSTRUCTURAL ANALYSIS OF HIGH SPEED STEEL
where C is a constant and d, t and n are particle size, time and diusion exponent, respectively. For n = 2, 3, 4, 5, the diusion is interface, volume, grain boundary and dislocation controlled, respectively. Equation (A1) was then rearranged as 1 A2 ln d ln C ln t: n The above check has supported the bulk or volume diusion during coarsening of both MC and M6C carbides. The K-values were then calculated using the following equation as explained in the text: 8 a bsV y DgV U gV r3 ÿ r30 t: 9aRT U yV ÿ U gV 2
APPENDIX B Nomenclature MC
rMC Vy
Vy a=b=1
density of MC carbide = 7.84 Mg/m3 [15] molar volume of MC carbide = (weight of 1 mol)/(density) 343.007/7.84 = 43.75 10ÿ6 m3/mol for MC carbide
s DgV DgV
interfacial energy for ferrite/cementite interface = 0.7 J/m2 [16] diusion coecient of vanadium in austenite D0 exp ÿQ=RT
D0 gV
365 mm2/s [8]
QgV
activation energy of vanadium diusion in g-Fe = 293 kJ/mol [8] 1.0 (0.016 + 0.016/20 16) = 2.88 10ÿ2 (8(1 + 1) 0.7 J/m2 43.75 10ÿ6 m3/mol 1.47 10ÿ14 m2/s)/(9 1 8.31 J/(mol K) 1473 K(1 ÿ 2.88 10ÿ2)2 1.996 10ÿ6 mm3/s.
UyV UgV K K
A3 M6C
rM6 C Vy a=6 b=1 g D0 QgW
UgW 10008C UyW 10008C K K
10.29 [15] 8424.65/10.29 = 818.72 10ÿ6 m3/mol ferrite/cementite = 0.7 J/m2 [16] 13 mm2/s [8] activation energy of tungsten diusion in g-Fe = 267 kJ/mol [8] 0.0040 [17] 0.163 [17] (8(6 + 1)0.7 818.72 10ÿ6 4.37 10ÿ15 0.0040)/(9 6 8.31 1473(0.163 ÿ 0.0040)2 3.36 10ÿ23 m3/s = 3.36 10ÿ5 mm3/s.