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Precipitation in 9Ni–12Cr–2Cu maraging steels
PII:
Acta mater. Vol. 46, No. 17, pp. 6063±6073, 1998 # 1998 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved Printed in Great Britain S1359-6454(98)00267-5 1359-6454/98 $19.00 + 0.00
PRECIPITATION IN 9Ni±12Cr±2Cu MARAGING STEELS K. STILLER1{, M. HAÈTTESTRAND1 and F. DANOIX2 Department of Physics, Chalmers University of Technology, S-412 96 GoÈteborg, Sweden and Laboratoire de Microscopie Ionique, URA CNRS 808, Universite de Rouen, F-76821 Mont Saint Aignan, France 1
2
(Received 9 February 1998; accepted 28 July 1998) AbstractÐTwo maraging steels with the compositions 9Ni±12Cr±2Cu±4Mo (wt%) and 9Ni±12Cr±2Cu and with small additions of Al and Ti were investigated using atom probe ®eld ion microscopy. Tomographic atom probe investigations were performed to clarify the spatial distribution of elements in and close to the precipitates. Materials heat treated at 4758C for 5, 25 min, 1, 2, 4 and 400 h were analyzed. Precipitates in the Mo-rich material were observed already after 5 min of aging, while in the material without Mo, precipitation started later. In both materials precipitation begins with the formation of Cu-rich particles which work as nucleation sites for a Ni-rich phase of type Ni3(Ti,Al). A Mo-rich phase was detected in the Mo-rich steel after 2 h of aging. The distribution of alloying elements in the precipitates, their role in the precipitation process, and the mechanism of hardening in the two materials are discussed. # 1998 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved.
1. INTRODUCTION
Maraging steels, developed some 30 years ago, were created for applications requiring ultrahigh strength in combination with good fracture toughness. The materials achieve their good mechanical properties because of small, densely distributed intermetallic precipitates formed during heat treatment, a process called maraging. The type and structure of the precipitates depend on the composition of the material, aging temperature and time. However, dierent results have been reported even for the same material and after the same heat treatment. These discrepancies in observations can be explained by the similarity of diraction patterns from dierent possible precipitate phases making their identi®cation using transmission electron microscopy (TEM) very hard. The situation is specially dicult, when analyzing small precipitates that cannot be extracted from the bulk, but must be studied in the presence of a highly magnetic matrix that impedes the investigation. Atom probe ®eld ion microscopy (APFIM) proved to be very powerful in the investigation of small precipitates when other microscopy techniques are unsatisfactory [1]. Its unique capability of measuring compositional variations on a nanometer scale together with equal detection eciency for all elements makes it particularly suitable for investigation of early stages of precipitation. Recent developments of the technique allow not only a registration of compositional changes along the direction of analysis (one-dimensional APFIM), but {To whom all correspondence should be addressed.
also a three-dimensional reconstruction of the distribution of dierent elements in the analyzed volume. In this investigation both conventional, i.e. onedimensional and three-dimensional APFIM techniques were applied for studies of precipitation in two maraging steels with the compositions 9Ni± 12Cr±2Cu±4Mo and 9Ni±12Cr±2Cu (wt%). Despite the similarity in composition the hardening behavior of the materials was very dierent. Molybdenum-containing material proved to be stable against overaging at temperatures below 5408C, while Mo-poor material showed continuous hardening only for temperatures lower than 4258C [2, 3]. TEM investigations of the precipitation in the materials were successful only after prolonged heat treatments where the precipitates were large enough to allow their identi®cation by analysis of electron diraction patterns [4]. The present study was undertaken to investigate early stages of precipitation and in this way understand the mechanisms of hardening in the two materials after commercial, short time tempering.
2. EXPERIMENTAL
Two materials; one containing 4 wt% of Mo (denoted 1RK91) and one with a very small amount of this element (denoted C455), were studied after aging at 4758C. This is the temperature used for the commercial heat treatment of Mo-rich steel. The detailed composition of the materials and the applied aging times are presented in Tables 1 and 2, respectively.
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STILLER et al.: PRECIPITATION IN MARAGING STEELS Table 1. Nominal composition (at.%) of investigated materials
Material
1RK91 C455
Element C
Al
Si
Ti
Cr
Fe
Mn
Ni
Cu
Mo
0.05 Ð
0.69 0.14
0.26 0.18
1.16 1.3
12.82 12.3
72.04 76.1
0.28 Ð
8.63 7.9
1.74 1.8
2.32 <0.3
The 1RK91 material was produced as a 0.9 mm wire from full scale melt by Sandvik AB. It was annealed at 1050±11508C, air cooled and cold worked (90% of cold reduction) before the ®nal heat treatment at 4758C. C455 material was cold drawn to 1 mm wire and heat treated in a similar manner to 1RK91. Cold reduction proceeding aging was about 20% in this material. Both materials were investigated by APFIM in four dierent conditions (Table 2). 1RK91 was additionally investigated after 0, 4 and 400 h of aging. Tomographic atom probe (TAP) three-dimensional investigations [5] were performed on all variants of 1RK91, while C455 material was analyzed by TAP only after 5 min of aging. APFIM specimens were prepared using a standard electropolishing method described elsewhere [6]. Analysis was performed at 65 K using 20% pulse ratio. Field ion microscopy (FIM) micrographs were obtained using Ne as the imaging gas. Two dierent methods of compositional APFIM analysis of the precipitates were used; random analysis, in which the measurement starts at a random position on the specimen surface; and selected area analysis, in which the aperture is positioned on the chosen feature on the specimen surface. The latter method generally gives a more precise composition of the analyzed features, since the contribution from the matrix is minimized by optimizing the position and the size of the aperture. However, for a successful selected area analysis the contrast on FIM micrographs from the precipitates must be good enough. Small dark appearing precipitates (such as Cu precipitates in the Fe matrix) are dicult to analyze using this procedure. After random area investigations, ladder diagrams (the number of collected ions of some speci®c element as Table 2. Heat treatments and investigation methods of studied materials
1RK91
C455
Aging time
AP
TAP
Unaged 5 min 25 min 1h 2h 4h 400 h Aged 5 min 25 min 1h 2h
X X X X X X X X X X X
X X X X X X X X
a function of the total number of ions) were used to estimate the position of the interface between the precipitate and the matrix, and in this way to minimize the contribution from the matrix to the composition of the precipitates. TAP analyses were applied to follow compositional evolution of the precipitates in 1RK91. The average compositions of the precipitates after each heat treatment were obtained using small cubic volumes (1 nm 1 nm 1 nm) positioned in the center of each precipitate. Besides, concentration pro®les in dierent directions through the precipitates were obtained moving step by step the small volumes in the chosen direction and calculating the concentration of elements in each position. Mass resolution in TAP analysis is not good enough to allow separation of Mo3+ isotopes (at mass to charge ratio of 31.33, 31.66 and 32.33) from both Cu2+ isotopes (at 31.5 and 32.5). Therefore, when calculating the Mo content in the precipitates, only the contribution from the nonoverlapping Mo ions was considered. This resulted in the Mo concentration measured by TAP, being slightly lower and the Cu content slightly higher than their real values. Thus, to calculate an overall composition of the precipitates the results from APFIM analysis (where separation of Mo3+ and Cu2+ is accessible) were used, while TAP results were used to estimate compositions in dierent parts of the precipitates. In 1RK91, the size of the precipitates and their number density were measured using three-dimensional reconstruction of the analyzed volumes in TAP. In the C455 material, precipitates were observed only using APFIM. Therefore, their sizes were estimated using a geometrical model proposed by Blavette and Chambreland [7]. The two methods give slightly dierent results. Thus, to enable comparison of the precipitation in the two materials, the corresponding values from APFIM analysis of 1RK91 are also presented. 3. RESULTS
3.1. Unaged material The matrix composition of unaged 1RK91 material was measured using APFIM. The result is summarized in Table 3. The composition is in good agreement with the nominal composition of the material. The only exception is the Cu content, being slightly lower. This is most probably due to some
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Table 3. Matrix composition of 1RK91 and C455 by APFIM (at.%) Aging
Element C
(a) 1RK91 non 0.052 0.02 5 min 0.062 0.03 25 min 0.092 0.03 1h 0.002 0.00 2h 0.022 0.01 4h 0.002 0.00 400 h 0.002 0.00 (b) C455 5 min 0.022 0.01 25 min 0.032 0.01 1h 0.052 0.02 2h 0.022 0.02
Al
Si
Ti
Cr
Fe
Mn
Ni
Cu
Mo
0.532 0.08 0.762 0.11 0.502 0.08 0.412 0.05 0.152 0.03 0.232 0.09 0.132 0.06
0.442 0.07 0.402 0.08 0.572 0.09 0.392 0.04 0.242 0.03 0.232 0.09 0.102 0.05
1.052 0.11 1.032 0.12 0.902 0.11 0.512 0.03 0.132 0.02 0.102 0.06 0.032 0.03
12.47 2 0.39 12.41 2 0.42 12.51 2 0.40 16.00 2 0.19 14.90 2 0.25 15.30 2 0.72 13.47 2 0.60
72.65 20.74 73.05 20.80 73.48 20.73 74.19 20.28 77.43 20.37 77.76 21.08 82.58 21.02
0.33 20.06 0.21 20.06 0.18 20.05 0.26 20.02 0.17 20.03 0.33 20.10 0.33 20.09
8.79 20.33 9.04 20.37 8.36 20.32 5.26 20.11 4.51 20.14 3.99 20.35 2.77 20.26
1.272 0.12 0.942 0.12 0.712 0.09 0.422 0.03 0.262 0.03 0.202 0.08 0.052 0.04
2.432 0.17 2.062 0.17 2.692 0.18 2.172 0.07 2.152 0.10 1.832 0.24 0.542 0.12
0.162 0.04 0.132 0.06 0.142 0.04 0.072 0.04
0.122 0.03 0.092 0.02 0.112 0.06 0.072 0.04
1.912 0.13 1.042 0.08 1.112 0.11 0.652 0.12
14.37 2 0.35 13.15 2 0.26 14.32 2 0.41 12.94 2 0.52
74.67 20.59 78.29 20.50 77.17 20.74 79.99 20.74
0.06 20.02 0.05 20.02 0.02 20.02 0.15 20.06
7.42 20.25 6.27 20.19 5.80 20.27 5.76 20.37
1.152 0.11 0.792 0.07 1.112 0.11 0.222 0.07
0.122 0.03 0.122 0.03 0.172 0.04 0.132 0.05
minor preferential ®eld evaporation of Cu ions (i.e. evaporation of ions between high voltage pulses), which prevents them from being detected. No evidence of the existence of precipitates was found on FIM micrographs. This result was also con®rmed by TAP investigation. Statistical treatment of the randomness of ion arrivals to the detector, using the mean separation method, denied clustering of any element. 3.2. Materials aged for 5 min The matrix compositions, obtained using APFIM, are presented in Table 3. No indication of clustering of any element in C455, using statistics, was found. TAP analysis revealed a homogeneous distribution of all elements [see Fig. 1(a)] in this material. The situation for 1RK91 was very dierent. The matrix composition clearly showed a de®cit of the Cu content compared to the nominal level of this element. The composition pro®les indicated the existence of very small clusters rich in Cu, Ni, Ti and Al. This observation was also con®rmed by TAP analysis, where a dense distribution (average number density of precipitates was 4.0 1024/m3) of small precipitates (1.5 nm in diameter) was observed [Fig. 1(b)]. TAP analysis revealed that the distribution of elements in the precipitates was inhomogeneous. They consisted of a Cu-rich nucleus surrounded by an Al-, Ti- and Ni-rich region (Fig. 2). Concentration pro®les through precipitates, obtained using TAP analysis, revealed that the regions rich in Ni, Al and Ti were more pronounced on one side of the Cu nuclei. No preferential direction for this enrichment was, however, observed. Moreover, Al- and Ni-enriched zones were not exactly coincident with each other. Aluminum-enriched regions seemed to be situated slightly closer to the Cu-rich nuclei than Nienriched (Fig. 3). The composition of the precipitates, measured by APFIM using selected area analysis, was incorrect because they showed a tendency to be ripped from
the specimen surface during one single pulse. This phenomenon contributed to the relatively low measured Cu content [Table 4(a)]. The average compositions in the center of Cu- and Ni-rich parts obtained by TAP are summarized in Table 5(a) and (b). Only the most important elements are listed in these tables. 3.3. Materials aged for 25 min Depletion of Cu in the bulk together with a small decrease in Al and Ti contents were observed in both materials (Table 3), indicating the presence of precipitation. According to TAP analysis, the spherical precipitates in 1RK91 were slightly larger (2.0 nm average diameter obtained by TAP and 1.4 nm estimated from APFIM) than those observed after 5 min aging. Some Mo and Si were incorporated into the precipitates but concentration pro®les through the precipitates showed that these elements were mostly situated outside the Ti- and Al-rich zones. The compositions of the precipitates obtained by APFIM are presented in Table 4(a). The average density of precipitates was 2.9 1024/ m3. The composition of precipitates in C455 material was similar to those in 1RK91, i.e. rich in Cu, Ni and Ti. However, the amount of Al in the precipitates was much smaller, re¯ecting a smaller nominal amount of this element in the material. The precipitates were also smaller (1.1 nm in diameter) than in 1RK91 material. The composition of precipitates obtained by APFIM is presented in Table 4(b). 3.4. Materials aged for 1 h APFIM analysis of the materials revealed further increase of the Ni and Ti content in the precipitates (Table 4) with the concomitant decrease of Ni, Ti, Cu and Al content in the matrix (Table 3). The low level of Ti, compared to that after 25 min of aging, in the precipitates in C455, is most probably caused by the fact that the most of the analyzed precipitates lay at the peripheries of the analyzed volume. This statement is also con®rmed by the low estimated size of the precipitates in this material, being
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Fig. 1. Distribution of elements by TAP after 5 min of aging: (a) C455; (b) 1RK91.
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Fig. 3. Concentration pro®le through one precipitate in 1RK91 after 5 min of aging.
Fig. 2. Distribution of Cu and Al in precipitates in 1RK91 after 5 min of aging.
3.5. Materials aged for 2 h
about the same as after 25 min of aging (1.2 nm). In both materials the amount of Cu in the Ni-rich zones has decreased with aging time. Instead volumes rich in Cu were found in connection to the Ni-rich precipitates. One precipitate much larger than others (2.2 nm in diameter) was observed in C455 material. The composition of this precipitate could therefore be more exactly determined, i.e. without large contribution from the matrix. It consisted of 70.5Ni±22.9Ti±2.3Al±0.6Cu±3.3Fe (at.%) which agrees well with the composition of the ZNi3Ti phase. TAP analysis of 1RK91 showed that precipitates rich in Ni, Ti and Al (denoted as Ni-rich in the following) grow at the surfaces of Cu-rich particles (Fig. 4). No preferential direction of the growth of the Ni-rich precipitates was observed. One typical concentration pro®le through the Ni-rich precipitates is shown in Fig. 5. The average size of precipitates obtained by TAP was 2.5 nm and their number density was 2.6 1024/m3. The compositions measured in the dierent parts of precipitates are summarized in Table 5(a) and (b). No sign of a clustering of Mo in the matrix of 1RK91 was observed (Fig. 6).
Further depletion of Ni, Ti, Al and Cu in the matrix was observed (Table 3). The separation of the Ni- and Cu-rich part of precipitates became clearer (Fig. 4) and the concentrations of Ni and Ti in Ni-rich precipitates had increased. The precipitate number density in 1RK91 by TAP was 2.2 1024/m3 and their size was 2.8 nm. The precipitates in C455 material were smaller than in 1RK91 (1.3 nm average size in C455 as compared to 2.3 nm estimated from APFIM in 1RK91) and contained less Ni. Selected area analysis of one large Ni-rich particle in 1RK91 revealed that it contained 72.3Ni±14.9Ti±6.0Al±2.1Cu±2.8Fe±1.0Cr±0.4Mo (at.%). Enrichment of Mo was observed at the interface between this precipitate and the matrix. The amount of Mo at the interface was higher than 10 at.% (Fig. 7). The Mo-rich zone was not coincident with the Cu-rich region. At this stage of aging, the mean separation method indicated also clustering of Mo in the bulk. One martensitic lath boundary was observed in the FIM image of the 1RK91 specimen (Fig. 8). No FIM contrast revealing the presence of precipitates at this boundary was observed.
Table 4. Composition of Ni-rich precipitates in 1RK91 and C455 by APFIM (at.%) Aging
(a) 1RK91 5 min 25 min 1h 2h 4h 400 h (b) C455 25 min 1h 2h
Element Al
Si
Ti
Cr
Fe
Mn
Ni
Cu
Mo
3.2 20.4 4.0 20.4 3.2 20.5 5.2 20.5 5.8 20.2 9.2 20.6
1.4 20.3 1.3 20.2 1.7 20.4 1.4 20.3 1.2 20.1 0.9 20.2
3.0 20.4 4.1 20.4 6.3 20.7 10.1 20.7 10.7 20.3 11.8 20.7
11.62 0.8 8.92 0.5 7.72 0.7 6.22 0.5 5.52 0.2 1.82 0.3
59.9 21.2 52.8 20.9 50.4 21.4 38.0 21.1 31.0 20.5 7.9 2 0.6
0.6 20.2 0.2 20.1 0.6 20.2 0.5 20.2 0.7 20.1 1.8 20.3
15.9 20.9 18.2 20.7 26.7 21.2 34.9 21.1 40.4 20.5 62.7 21.1
2.6 20.4 8.5 20.5 1.6 20.3 2.3 20.3 2.4 20.2 1.4 20.3
1.8 20.3 1.9 20.3 1.7 20.3 1.3 20.3 2.4 20.2 2.5 20.4
0.8 20.3 0.1 20.1 1.4 20.7
0.4 20.2 0.6 20.3 0.4 20.4
5.6 20.7 3.4 20.6 14.0 22.5
8.62 0.9 10.62 1.1 7.92 1.6
56.0 21.9 67.6 22.2 48.2 23.4
0.1 20.1 0.0 20.0 0.0 20.0
20.4 21.3 17.1 21.5 27.3 22.7
7.8 20.9 0.5 20.2 0.7 20.5
0.5 20.3 0.1 20.1 0.0 20.0
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Table 5. Composition in the middle of Ni-rich and Cu-rich parts of precipitates in 1RK91 by TAP (at.%) Aging
(a) Ni-rich 5 min 1h 2h
(b) Cu-rich 5 min 1h 2h 400 h
Element Al
Ti
Ni
Cu
Mo
12.3 9.6 10.6
3.8 12.8 14.0
16.0 42.5 49.5
7.8 7.8 4.1
Ð 0.6 Ð
Cu
Fe
Ni
Al
Ti
53.6 63.9 70.6 90.0
31.9 14.2 14.9 6.4
5.0 14.2 11.0 5.7
2.8 2.8 5.3 Ð
0.7 2.1 2.4 Ð
Fig. 5. Concentration pro®le through one precipitate in 1RK91 after 1 h of aging.
3.6. Material aged for 4 h Only 1RK91 was investigated in this condition. The precipitates were clearly visible as bright and dark areas in FIM micrographs (Fig. 9) and selected area investigations could be performed. By correlating the contrast on FIM micrographs with the composition of particles it was shown that Curich parts appeared as dark areas in FIM images while Ni-rich parts appeared as spherical bright areas. The precipitates were 4.7 nm in size and their number density was 1.6 1024/m3. Further decrease of Cu content in the Ni-rich precipitates [Table 4(a)] was observed. The Mo content in the matrix was lower than the nominal, indicating the existence of Mo-rich precipitates. This type of particle was indeed found in the material. Molybdenum-rich volumes were observed at the interfaces between the Ni-rich precipitates and the matrix. The measured Mo content in these regions could be as high as 14 at.%. The average composition of the Mo-rich phase is presented in Table 6. In addition to Fe, Cr and Ni the precipitates contained some Si and small amounts of P and C. An almost continuous line of disk-shaped Morich precipitates was found in one specimen, indicating precipitation of this phase at martensite lath boundaries or at prior austenite grain boundaries (Fig. 10). The composition of only one of such precipitates was obtained. It was similar to that obtained for precipitates connected to the Ni-rich particles. 3.7. Material aged for 400 h
Fig. 4. Distribution of elements in precipitates in 1RK91 after 1 and 2 h of aging.
The matrix of 1RK91 alloy in this condition was extremely depleted of Cu. Large amounts of Al, Ti and Mo were also missing in the matrix (Table 3). The Ni- and Mo-rich precipitates were observed both at lath boundaries and in the matrix. The Ni-rich particles became larger, 7±8 nm in size, but preserved their spheroidal shape. Their number density decreased to about 8 1023/m3. The
STILLER et al.: PRECIPITATION IN MARAGING STEELS
Fig. 6. Distribution of elements by TAP in 1RK91 after 1 and 2 h of aging.
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Fig. 7. Concentration pro®le through the interface of one precipitate in 1RK91 after 2 h of aging.
Ni content in the precipitates increased to almost 63 at.% [Table 4(a)]. TAP investigation revealed existence of plates of almost pure Cu [Table 5(b)] on the sides of Ni-rich particles (Fig. 11). The Mo-rich precipitates were observed both in the matrix and at lath boundaries, often in direct contact to the Ni-rich precipitates but not in connection to the Cu-rich plates. They became larger than in the material aged for 4 h and their average Mo content increased (Table 6). Figure 12 shows a three-dimensional reconstruction of one of the precipitates. The observed amount of Mo-rich precipitates in the matrix was much lower than the amount of Ni-rich precipitates. 4. DISCUSSION
The investigation shows that the nucleation and growth of the precipitates are very complex. No doubt, Cu plays a very important role in the
Fig. 8. FIM micrograph of one 1RK91 specimen after 2 h of aging. The position of an observed lath boundary is marked with arrows.
Fig. 9. FIM micrograph of one 1RK91 specimen after 4 h of aging: (a) bright appearing precipitates with two morphologies, spherical and disk-shaped, are marked with arrows; (b) dark appearing precipitates.
nucleation of the Ni-rich phase in both materials. The observation of Cu-rich nuclei, with a Cu content as high as 54 at.% and more widely dispersed Ni-, Al- and Ti-rich regions, at the beginning of the precipitation sequence, indicates that Cu must precipitate ®rst. In 1RK91 the nucleation of the Cu-rich phase is very fast and must occur during a period shorter than 5 min. This early precipitation of Cu should be expected since the solubility of Cu in Fe is known to be low [8]. The solubility of copper in ferrite is approximately 2 wt% at 8508C and much lower for the aging temperature used in the present investigation. The dense and homogeneous distribution of precipitates in the matrix also indicates a high driving force for the precipitation of this phase.
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Table 6. Composition of Mo-rich precipitates by APFIM (at.%) Material
aged 4 h aged 400 h
Element Si
Cr
Fe
Ni
Mo
C
P
1.5420.4 4.0 21.5
14.92 1.4 12.82 2.4
60.4 22.3 38.6 24.2
8.5 21.0 6.8 21.9
13.9 21.2 36.9 23.6
0.3 20.2 1.1 20.8
0.3 20.2 Ð
The Cu-rich particles seem to act as preferential nucleation sites for the precipitation of a phase rich in Ni, Ti and Al. The nucleation of a Ni-rich phase in 1RK91 is also very fast since an enrichment of Ni, Al and Ti was observed in connection to all Cu-rich nuclei after 5 min of aging. The distribution of elements in this phase is, however, inhomogeneous after this aging time. The maximum Ni content in the concentration pro®les by TAP (Fig. 3) is not coincident with the maximum of Al content. The observed shift in the distribution of Al towards Cu-rich nuclei compared to that of Ni, indicates faster diusion of Al in the steel. Another indi-
Fig. 10. FIM micrograph of one 1RK91 specimen after 4 h of aging. An almost continuous line of precipitates indicates the presence of a lath or prior austenite boundary.
Fig. 11. Three-dimensional reconstruction of Cu and Al distribution in a Ni-rich precipitate in a specimen aged for 400 h.
cation for this faster diusion of Al is its high content in the Ni-rich phase at the beginning of precipitation. The ratio between Al and Ni decreases with aging time from 0.8 after 5 min to 0.2 after 2 h [Table 5(a)]. This is to be expected since the ratio between the diusion coecient for Al in a-Fe and the coecient for self-diusion of Fe is about 30, while it is only 2 for Ni. Aluminum is also incorporated into the Cu-rich nuclei to much larger extent than Ti and Ni. After 5 min of aging the Al content in the Cu-rich volumes is about four times that of the bulk. On the other hand the amount of Ni in the center of Cu-rich nuclei is extremely low. It is therefore reasonable to assume that the Ni found in the Nirich parts, after this heat treatment, results from its rejection from the Cu-rich volumes and not from its diusion from the matrix. After 1 and 2 h of aging, the maxima of Ti, Al and Ni are coincident and the observed depletion of elements (especially Al and Ti) close to the precipitate surface (Fig. 5), indicates that the growth process is diusion limited. The calculated number density of precipitates in 1RK91 is highest after 5 min of aging but does not change much before 4 h of aging. The observed growth of the precipitates is also very slow and does not follow the classical size±time0.5 behavior
Fig. 12. Three-dimensional reconstruction of Mo and Si distribution in a Mo-rich precipitate in a specimen aged for 400 h.
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STILLER et al.: PRECIPITATION IN MARAGING STEELS
expected from the theory of diusion controlled growth of precipitates. It seems therefore that, because of the very high driving force for precipitation, most of the nucleation and growth occurs before 5 min of aging. After this time the rearrangement of elements in the precipitates, rather than further growth and coarsening, takes place. The precipitates split into two parts of dierent chemistry, and the contributions from Fe and Cr to both the Ni- and the Cu-rich parts decrease. Recent investigation of Cu precipitation in ferritic steels and alloys [9] reveals that the precipitation sequence is initiated by coherent, spherical b.c.c. precipitates which transform to a stable f.c.c. structure (e-phase) via intermediate 9R and 3R phases. It has been also shown that the critical diameter for transformation of the metastable b.c.c. precipitates to 9R structure is approximately 5 nm. This is much larger than the diameter of the observed Cu-rich parts of precipitates after 2 h of aging in 1RK91. Thus, the small size of the Cu- and Ni-rich precipitates together with the apparent lack of any preferential direction for growth of the Ni-rich phase indicate that the precipitates are coherent with the martensitic matrix at this stage. The Ni-rich precipitates preserve their spheroidal shape up to 400 h of aging, while the Cu precipitates form thin shields on the surface of the Ni-rich particles. The change of the shape of the Cu precipitates must be connected to the minimization of the surface energy between the Ni- and Cu-rich precipitates. The observed close ®tting of Cu shields on the Ni precipitates indicates some coherency between the two types of precipitates. If the Ni-rich phase is g' (see Section 5) then its 0.209 nm spacing between {111} planes ®ts best with the 0.207 nm spacing between {111} planes in the e-Cu. The next best candidates are 9R and 3R phases with 0.204 nm spacing between the close packed planes. e-Copper forms in ferritic steels only after prolonged aging and therefore, it seems most probable that Cu shields have 9R or 3R structure. It is also tempting to conclude that the Cu shields work as strain reducers between Ni precipitates and the martensite. This would however, implicate some relationship between the direction of the shields and the matrix which was not observed in the present investigation and therefore more detailed investigation is needed to con®rm this statement. It is interesting that Mo is not incorporated into the Ni-rich precipitates in 1RK91. TAP analysis of precipitates shows that the amount of this element in precipitates is very low for all aging conditions [see Table 5(a)]. Instead Mo enriches the peripheries of Ni-rich precipitates already after 25 min of aging, which shows that this element is rejected from precipitates. This observation is in contradiction to investigations of Ni-rich phases in other maraging steels [10]. According to these studies a large amount of Mo (replacing Ni) was observed in
Z-Ni3Ti precipitates. One possible explanation of the low amount of Mo in our precipitates would be that their structure diers from that of the Z phase. A possible candidate is g'-Ni3Al since the observed Ni-rich phase in 1RK91 includes large amounts of Al. It has been shown that Ti prefers to substitute for the Al sites in g' [11] and that at 11508C more than half of the Al atoms can be replaced by Ti [12]. On the other hand this phase can incorporate Mo only when the Al content is higher than 20 at.% [12]. In our study the highest measured Al content is 12 at.% which shows that Mo should not be included in the precipitates. Taking into consideration all these arguments it seems probable that the observed Ni-rich phase in 1RK91 material is g'. Precipitation in C455 material starts later than in 1RK91, de®nitely after 5 min. It follows, however, the general trend observed in 1RK91, i.e. precipitation of a Ni-rich phase nucleated on the Cu-rich particles. The largest dierence in the chemistry of the investigated alloys is their Mo content, being higher in 1RK91. It can therefore be tempting to explain the eect in terms of Mo speeding up precipitation of Cu, which is assumed to initiate precipitation of the Ni-rich phases. However, according to Thermo-Calc calculations the actual amount of Mo present in 1RK91 should not have any eect on Cu precipitation in the material. Another possible explanation of faster precipitation in 1RK91 alloy is its higher degree of cold deformation (90% cold reduction in 1RK91 as compared to 20% in C455) proceeding aging. The high cold deformation rate used is necessary in 1RK91 because it induces the martensite formation, while in C455 martensite forms during cooling [3]. Thus, the apparent diusivity in 1RK91, which contains a higher dislocation density, should be lower as this line defect acts as a diusion short circuit. There is no doubt that Mo precipitates last in the precipitation sequence. The ®rst signs of Mo clustering were observed after 2 h of aging. However, Mo precipitation at this stage was only initiated, which is also con®rmed by the observed lath boundary free from precipitates. After 4 h of aging, the Mo-rich phase was observed both in connection to the Ni-rich precipitates and at the lath boundaries (Figs 7 and 10). The amount of Mo in the precipitates is close to that in Fe2Mo Laves phase. This intermetallic compound with a hexagonal structure was observed in other maraging steels with low Co content [13, 14]. The hardening eect of 1RK91 after 4 h of aging, which corresponds to the commercial heat treatment, is mainly due to the observed dense distribution of Ni- and Cu-rich precipitates. The contribution of the Mo-rich phase must be less important at this stage. After 400 h of aging the contribution from Mo-rich precipitates to the material hardening becomes more important. However, the
STILLER et al.: PRECIPITATION IN MARAGING STEELS
Ni-rich precipitates are still small enough to be active in the hardening process. The in¯uence of Mo on Ni-rich precipitates could not be elucidated using Thermo-Calc since the accessible data base does not contain information about Ni-rich phases. However, since these precipitates proved a slow coarsening rate in 1RK91 material, while Mo-poor C455 material showed softening only after a few hours at 4758C, due to the coarsening of precipitates [3], it seems reasonable to assume that the slow growth of the precipitates in 1RK91 is caused by the presence of Mo. One possible explanation is that Mo enrichment at the Ni-rich precipitate interfaces (leading to its subsequent precipitation) hampers growth of the Ni-rich phase. Therefore, the superior properties of 1RK91 compared to C455 are due to its high Mo content, which both provides an additional precipitation hardening mechanism (by precipitation of the Morich phase) and impedes coarsening of the Ni-rich phase. 5. CONCLUSIONS
. The precipitation in both investigated materials starts with the formation of Cu-rich particles only after a few minutes at 4758C. These particles act as nucleation sites for a Ni-rich phase of type Ni3(Ti,Al). . After 2 h of aging Cu- and Ni-rich phases are well separated. . Nucleation of a Mo-rich phase on the Ni-rich precipitates in 1RK91 starts after 2 h of aging. . The observed slow growth of Ni-rich precipitates in 1RK91 up to 400 h of aging is most likely due to a hampering eect of Mo that enriches (and precipitates) at their surfaces. . Hardening of 1RK91 after 4 h of aging is mainly due to the Ni- and Cu-rich precipitates, while the
6073
contribution from precipitation of the Mo-rich phase must be less important at this stage. . The Mo-rich phase provides an additional hardening mechanism after prolonged aging. . The presence of Mo is responsible for the superior properties of 1RK91 compared to C455 after long heat treatment.
AcknowledgementsÐAB Sandvik Steel is acknowledged for supply of material and fruitful discussions.
REFERENCES 1. Stiller, K., Svensson, L.-E., Howell, P. R., Wang Rong, AndreÂn, H.-O. and Dunlop, G. L., Acta metall., 1984, 33, 1457. 2. Nilsson, J. O., Stigenberg, A. M. and Liu, P., Metall. Mater. Trans. A, 1994, 25A, 2225. 3. Stigenberg, A. M., Nilsson, J. O., Liu, P. and Wilson, A., Wire, 1994, 44, 375. 4. Liu, P., Stigenberg, A. M. and Nilsson, J. O., Acta metall., 1995, 43, 2881. 5. Bostel, A., Blavette, D., Menand, A. and Sarrau, J. M., J. Physique (Paris), 1989, C8-50, 501. 6. Stiller, K., Danoix, F. and Bostel, A., Appl. Surf. Sci., 1996, 94/95, 326. 7. Blavette, D. and Chambreland, S., J. Physique (Paris), 1986, C47, 503. 8. Fourlaris, G., Baker, A. J. and Papadimitrious, G. D., Acta metall. mater., 1995, 43, 2589. 9. Othen, P. J., Jenkins, M. L. and Smith, G. D. W., Phil. Mag. A, 1994, 70, 1. 10. Sha, W., Cerezo, A. and Smith, G. D. W., Metall. Trans., 1993, 24A, 1221. 11. Mukarani, H., Harada, H. and Bhadeshia, H. K. D. H., Appl. Surf. Sci., 1994, 76/77, 177. 12. Kornilov, I. I., in Intermetallic Compounds, ed. J. H. Westbrook. John Wiley and Sons, 1967, p. 368. 13. Servant, C. and Bouzid, N., Phil. Mag. B, 1989, 60, 659. 14. Vasudevan, V. K., Kim, S. J. and Wayman, C. M., Metall. Trans. A, 1990, 21A, 2655.
Acta mater. Vol. 46, No. 17, pp. 6063±6073, 1998 # 1998 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved Printed in Great Britain S1359-6454(98)00267-5 1359-6454/98 $19.00 + 0.00
PRECIPITATION IN 9Ni±12Cr±2Cu MARAGING STEELS K. STILLER1{, M. HAÈTTESTRAND1 and F. DANOIX2 Department of Physics, Chalmers University of Technology, S-412 96 GoÈteborg, Sweden and Laboratoire de Microscopie Ionique, URA CNRS 808, Universite de Rouen, F-76821 Mont Saint Aignan, France 1
2
(Received 9 February 1998; accepted 28 July 1998) AbstractÐTwo maraging steels with the compositions 9Ni±12Cr±2Cu±4Mo (wt%) and 9Ni±12Cr±2Cu and with small additions of Al and Ti were investigated using atom probe ®eld ion microscopy. Tomographic atom probe investigations were performed to clarify the spatial distribution of elements in and close to the precipitates. Materials heat treated at 4758C for 5, 25 min, 1, 2, 4 and 400 h were analyzed. Precipitates in the Mo-rich material were observed already after 5 min of aging, while in the material without Mo, precipitation started later. In both materials precipitation begins with the formation of Cu-rich particles which work as nucleation sites for a Ni-rich phase of type Ni3(Ti,Al). A Mo-rich phase was detected in the Mo-rich steel after 2 h of aging. The distribution of alloying elements in the precipitates, their role in the precipitation process, and the mechanism of hardening in the two materials are discussed. # 1998 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved.
1. INTRODUCTION
Maraging steels, developed some 30 years ago, were created for applications requiring ultrahigh strength in combination with good fracture toughness. The materials achieve their good mechanical properties because of small, densely distributed intermetallic precipitates formed during heat treatment, a process called maraging. The type and structure of the precipitates depend on the composition of the material, aging temperature and time. However, dierent results have been reported even for the same material and after the same heat treatment. These discrepancies in observations can be explained by the similarity of diraction patterns from dierent possible precipitate phases making their identi®cation using transmission electron microscopy (TEM) very hard. The situation is specially dicult, when analyzing small precipitates that cannot be extracted from the bulk, but must be studied in the presence of a highly magnetic matrix that impedes the investigation. Atom probe ®eld ion microscopy (APFIM) proved to be very powerful in the investigation of small precipitates when other microscopy techniques are unsatisfactory [1]. Its unique capability of measuring compositional variations on a nanometer scale together with equal detection eciency for all elements makes it particularly suitable for investigation of early stages of precipitation. Recent developments of the technique allow not only a registration of compositional changes along the direction of analysis (one-dimensional APFIM), but {To whom all correspondence should be addressed.
also a three-dimensional reconstruction of the distribution of dierent elements in the analyzed volume. In this investigation both conventional, i.e. onedimensional and three-dimensional APFIM techniques were applied for studies of precipitation in two maraging steels with the compositions 9Ni± 12Cr±2Cu±4Mo and 9Ni±12Cr±2Cu (wt%). Despite the similarity in composition the hardening behavior of the materials was very dierent. Molybdenum-containing material proved to be stable against overaging at temperatures below 5408C, while Mo-poor material showed continuous hardening only for temperatures lower than 4258C [2, 3]. TEM investigations of the precipitation in the materials were successful only after prolonged heat treatments where the precipitates were large enough to allow their identi®cation by analysis of electron diraction patterns [4]. The present study was undertaken to investigate early stages of precipitation and in this way understand the mechanisms of hardening in the two materials after commercial, short time tempering.
2. EXPERIMENTAL
Two materials; one containing 4 wt% of Mo (denoted 1RK91) and one with a very small amount of this element (denoted C455), were studied after aging at 4758C. This is the temperature used for the commercial heat treatment of Mo-rich steel. The detailed composition of the materials and the applied aging times are presented in Tables 1 and 2, respectively.
6063
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STILLER et al.: PRECIPITATION IN MARAGING STEELS Table 1. Nominal composition (at.%) of investigated materials
Material
1RK91 C455
Element C
Al
Si
Ti
Cr
Fe
Mn
Ni
Cu
Mo
0.05 Ð
0.69 0.14
0.26 0.18
1.16 1.3
12.82 12.3
72.04 76.1
0.28 Ð
8.63 7.9
1.74 1.8
2.32 <0.3
The 1RK91 material was produced as a 0.9 mm wire from full scale melt by Sandvik AB. It was annealed at 1050±11508C, air cooled and cold worked (90% of cold reduction) before the ®nal heat treatment at 4758C. C455 material was cold drawn to 1 mm wire and heat treated in a similar manner to 1RK91. Cold reduction proceeding aging was about 20% in this material. Both materials were investigated by APFIM in four dierent conditions (Table 2). 1RK91 was additionally investigated after 0, 4 and 400 h of aging. Tomographic atom probe (TAP) three-dimensional investigations [5] were performed on all variants of 1RK91, while C455 material was analyzed by TAP only after 5 min of aging. APFIM specimens were prepared using a standard electropolishing method described elsewhere [6]. Analysis was performed at 65 K using 20% pulse ratio. Field ion microscopy (FIM) micrographs were obtained using Ne as the imaging gas. Two dierent methods of compositional APFIM analysis of the precipitates were used; random analysis, in which the measurement starts at a random position on the specimen surface; and selected area analysis, in which the aperture is positioned on the chosen feature on the specimen surface. The latter method generally gives a more precise composition of the analyzed features, since the contribution from the matrix is minimized by optimizing the position and the size of the aperture. However, for a successful selected area analysis the contrast on FIM micrographs from the precipitates must be good enough. Small dark appearing precipitates (such as Cu precipitates in the Fe matrix) are dicult to analyze using this procedure. After random area investigations, ladder diagrams (the number of collected ions of some speci®c element as Table 2. Heat treatments and investigation methods of studied materials
1RK91
C455
Aging time
AP
TAP
Unaged 5 min 25 min 1h 2h 4h 400 h Aged 5 min 25 min 1h 2h
X X X X X X X X X X X
X X X X X X X X
a function of the total number of ions) were used to estimate the position of the interface between the precipitate and the matrix, and in this way to minimize the contribution from the matrix to the composition of the precipitates. TAP analyses were applied to follow compositional evolution of the precipitates in 1RK91. The average compositions of the precipitates after each heat treatment were obtained using small cubic volumes (1 nm 1 nm 1 nm) positioned in the center of each precipitate. Besides, concentration pro®les in dierent directions through the precipitates were obtained moving step by step the small volumes in the chosen direction and calculating the concentration of elements in each position. Mass resolution in TAP analysis is not good enough to allow separation of Mo3+ isotopes (at mass to charge ratio of 31.33, 31.66 and 32.33) from both Cu2+ isotopes (at 31.5 and 32.5). Therefore, when calculating the Mo content in the precipitates, only the contribution from the nonoverlapping Mo ions was considered. This resulted in the Mo concentration measured by TAP, being slightly lower and the Cu content slightly higher than their real values. Thus, to calculate an overall composition of the precipitates the results from APFIM analysis (where separation of Mo3+ and Cu2+ is accessible) were used, while TAP results were used to estimate compositions in dierent parts of the precipitates. In 1RK91, the size of the precipitates and their number density were measured using three-dimensional reconstruction of the analyzed volumes in TAP. In the C455 material, precipitates were observed only using APFIM. Therefore, their sizes were estimated using a geometrical model proposed by Blavette and Chambreland [7]. The two methods give slightly dierent results. Thus, to enable comparison of the precipitation in the two materials, the corresponding values from APFIM analysis of 1RK91 are also presented. 3. RESULTS
3.1. Unaged material The matrix composition of unaged 1RK91 material was measured using APFIM. The result is summarized in Table 3. The composition is in good agreement with the nominal composition of the material. The only exception is the Cu content, being slightly lower. This is most probably due to some
STILLER et al.: PRECIPITATION IN MARAGING STEELS
6065
Table 3. Matrix composition of 1RK91 and C455 by APFIM (at.%) Aging
Element C
(a) 1RK91 non 0.052 0.02 5 min 0.062 0.03 25 min 0.092 0.03 1h 0.002 0.00 2h 0.022 0.01 4h 0.002 0.00 400 h 0.002 0.00 (b) C455 5 min 0.022 0.01 25 min 0.032 0.01 1h 0.052 0.02 2h 0.022 0.02
Al
Si
Ti
Cr
Fe
Mn
Ni
Cu
Mo
0.532 0.08 0.762 0.11 0.502 0.08 0.412 0.05 0.152 0.03 0.232 0.09 0.132 0.06
0.442 0.07 0.402 0.08 0.572 0.09 0.392 0.04 0.242 0.03 0.232 0.09 0.102 0.05
1.052 0.11 1.032 0.12 0.902 0.11 0.512 0.03 0.132 0.02 0.102 0.06 0.032 0.03
12.47 2 0.39 12.41 2 0.42 12.51 2 0.40 16.00 2 0.19 14.90 2 0.25 15.30 2 0.72 13.47 2 0.60
72.65 20.74 73.05 20.80 73.48 20.73 74.19 20.28 77.43 20.37 77.76 21.08 82.58 21.02
0.33 20.06 0.21 20.06 0.18 20.05 0.26 20.02 0.17 20.03 0.33 20.10 0.33 20.09
8.79 20.33 9.04 20.37 8.36 20.32 5.26 20.11 4.51 20.14 3.99 20.35 2.77 20.26
1.272 0.12 0.942 0.12 0.712 0.09 0.422 0.03 0.262 0.03 0.202 0.08 0.052 0.04
2.432 0.17 2.062 0.17 2.692 0.18 2.172 0.07 2.152 0.10 1.832 0.24 0.542 0.12
0.162 0.04 0.132 0.06 0.142 0.04 0.072 0.04
0.122 0.03 0.092 0.02 0.112 0.06 0.072 0.04
1.912 0.13 1.042 0.08 1.112 0.11 0.652 0.12
14.37 2 0.35 13.15 2 0.26 14.32 2 0.41 12.94 2 0.52
74.67 20.59 78.29 20.50 77.17 20.74 79.99 20.74
0.06 20.02 0.05 20.02 0.02 20.02 0.15 20.06
7.42 20.25 6.27 20.19 5.80 20.27 5.76 20.37
1.152 0.11 0.792 0.07 1.112 0.11 0.222 0.07
0.122 0.03 0.122 0.03 0.172 0.04 0.132 0.05
minor preferential ®eld evaporation of Cu ions (i.e. evaporation of ions between high voltage pulses), which prevents them from being detected. No evidence of the existence of precipitates was found on FIM micrographs. This result was also con®rmed by TAP investigation. Statistical treatment of the randomness of ion arrivals to the detector, using the mean separation method, denied clustering of any element. 3.2. Materials aged for 5 min The matrix compositions, obtained using APFIM, are presented in Table 3. No indication of clustering of any element in C455, using statistics, was found. TAP analysis revealed a homogeneous distribution of all elements [see Fig. 1(a)] in this material. The situation for 1RK91 was very dierent. The matrix composition clearly showed a de®cit of the Cu content compared to the nominal level of this element. The composition pro®les indicated the existence of very small clusters rich in Cu, Ni, Ti and Al. This observation was also con®rmed by TAP analysis, where a dense distribution (average number density of precipitates was 4.0 1024/m3) of small precipitates (1.5 nm in diameter) was observed [Fig. 1(b)]. TAP analysis revealed that the distribution of elements in the precipitates was inhomogeneous. They consisted of a Cu-rich nucleus surrounded by an Al-, Ti- and Ni-rich region (Fig. 2). Concentration pro®les through precipitates, obtained using TAP analysis, revealed that the regions rich in Ni, Al and Ti were more pronounced on one side of the Cu nuclei. No preferential direction for this enrichment was, however, observed. Moreover, Al- and Ni-enriched zones were not exactly coincident with each other. Aluminum-enriched regions seemed to be situated slightly closer to the Cu-rich nuclei than Nienriched (Fig. 3). The composition of the precipitates, measured by APFIM using selected area analysis, was incorrect because they showed a tendency to be ripped from
the specimen surface during one single pulse. This phenomenon contributed to the relatively low measured Cu content [Table 4(a)]. The average compositions in the center of Cu- and Ni-rich parts obtained by TAP are summarized in Table 5(a) and (b). Only the most important elements are listed in these tables. 3.3. Materials aged for 25 min Depletion of Cu in the bulk together with a small decrease in Al and Ti contents were observed in both materials (Table 3), indicating the presence of precipitation. According to TAP analysis, the spherical precipitates in 1RK91 were slightly larger (2.0 nm average diameter obtained by TAP and 1.4 nm estimated from APFIM) than those observed after 5 min aging. Some Mo and Si were incorporated into the precipitates but concentration pro®les through the precipitates showed that these elements were mostly situated outside the Ti- and Al-rich zones. The compositions of the precipitates obtained by APFIM are presented in Table 4(a). The average density of precipitates was 2.9 1024/ m3. The composition of precipitates in C455 material was similar to those in 1RK91, i.e. rich in Cu, Ni and Ti. However, the amount of Al in the precipitates was much smaller, re¯ecting a smaller nominal amount of this element in the material. The precipitates were also smaller (1.1 nm in diameter) than in 1RK91 material. The composition of precipitates obtained by APFIM is presented in Table 4(b). 3.4. Materials aged for 1 h APFIM analysis of the materials revealed further increase of the Ni and Ti content in the precipitates (Table 4) with the concomitant decrease of Ni, Ti, Cu and Al content in the matrix (Table 3). The low level of Ti, compared to that after 25 min of aging, in the precipitates in C455, is most probably caused by the fact that the most of the analyzed precipitates lay at the peripheries of the analyzed volume. This statement is also con®rmed by the low estimated size of the precipitates in this material, being
6066
STILLER et al.: PRECIPITATION IN MARAGING STEELS
Fig. 1. Distribution of elements by TAP after 5 min of aging: (a) C455; (b) 1RK91.
STILLER et al.: PRECIPITATION IN MARAGING STEELS
6067
Fig. 3. Concentration pro®le through one precipitate in 1RK91 after 5 min of aging.
Fig. 2. Distribution of Cu and Al in precipitates in 1RK91 after 5 min of aging.
3.5. Materials aged for 2 h
about the same as after 25 min of aging (1.2 nm). In both materials the amount of Cu in the Ni-rich zones has decreased with aging time. Instead volumes rich in Cu were found in connection to the Ni-rich precipitates. One precipitate much larger than others (2.2 nm in diameter) was observed in C455 material. The composition of this precipitate could therefore be more exactly determined, i.e. without large contribution from the matrix. It consisted of 70.5Ni±22.9Ti±2.3Al±0.6Cu±3.3Fe (at.%) which agrees well with the composition of the ZNi3Ti phase. TAP analysis of 1RK91 showed that precipitates rich in Ni, Ti and Al (denoted as Ni-rich in the following) grow at the surfaces of Cu-rich particles (Fig. 4). No preferential direction of the growth of the Ni-rich precipitates was observed. One typical concentration pro®le through the Ni-rich precipitates is shown in Fig. 5. The average size of precipitates obtained by TAP was 2.5 nm and their number density was 2.6 1024/m3. The compositions measured in the dierent parts of precipitates are summarized in Table 5(a) and (b). No sign of a clustering of Mo in the matrix of 1RK91 was observed (Fig. 6).
Further depletion of Ni, Ti, Al and Cu in the matrix was observed (Table 3). The separation of the Ni- and Cu-rich part of precipitates became clearer (Fig. 4) and the concentrations of Ni and Ti in Ni-rich precipitates had increased. The precipitate number density in 1RK91 by TAP was 2.2 1024/m3 and their size was 2.8 nm. The precipitates in C455 material were smaller than in 1RK91 (1.3 nm average size in C455 as compared to 2.3 nm estimated from APFIM in 1RK91) and contained less Ni. Selected area analysis of one large Ni-rich particle in 1RK91 revealed that it contained 72.3Ni±14.9Ti±6.0Al±2.1Cu±2.8Fe±1.0Cr±0.4Mo (at.%). Enrichment of Mo was observed at the interface between this precipitate and the matrix. The amount of Mo at the interface was higher than 10 at.% (Fig. 7). The Mo-rich zone was not coincident with the Cu-rich region. At this stage of aging, the mean separation method indicated also clustering of Mo in the bulk. One martensitic lath boundary was observed in the FIM image of the 1RK91 specimen (Fig. 8). No FIM contrast revealing the presence of precipitates at this boundary was observed.
Table 4. Composition of Ni-rich precipitates in 1RK91 and C455 by APFIM (at.%) Aging
(a) 1RK91 5 min 25 min 1h 2h 4h 400 h (b) C455 25 min 1h 2h
Element Al
Si
Ti
Cr
Fe
Mn
Ni
Cu
Mo
3.2 20.4 4.0 20.4 3.2 20.5 5.2 20.5 5.8 20.2 9.2 20.6
1.4 20.3 1.3 20.2 1.7 20.4 1.4 20.3 1.2 20.1 0.9 20.2
3.0 20.4 4.1 20.4 6.3 20.7 10.1 20.7 10.7 20.3 11.8 20.7
11.62 0.8 8.92 0.5 7.72 0.7 6.22 0.5 5.52 0.2 1.82 0.3
59.9 21.2 52.8 20.9 50.4 21.4 38.0 21.1 31.0 20.5 7.9 2 0.6
0.6 20.2 0.2 20.1 0.6 20.2 0.5 20.2 0.7 20.1 1.8 20.3
15.9 20.9 18.2 20.7 26.7 21.2 34.9 21.1 40.4 20.5 62.7 21.1
2.6 20.4 8.5 20.5 1.6 20.3 2.3 20.3 2.4 20.2 1.4 20.3
1.8 20.3 1.9 20.3 1.7 20.3 1.3 20.3 2.4 20.2 2.5 20.4
0.8 20.3 0.1 20.1 1.4 20.7
0.4 20.2 0.6 20.3 0.4 20.4
5.6 20.7 3.4 20.6 14.0 22.5
8.62 0.9 10.62 1.1 7.92 1.6
56.0 21.9 67.6 22.2 48.2 23.4
0.1 20.1 0.0 20.0 0.0 20.0
20.4 21.3 17.1 21.5 27.3 22.7
7.8 20.9 0.5 20.2 0.7 20.5
0.5 20.3 0.1 20.1 0.0 20.0
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STILLER et al.: PRECIPITATION IN MARAGING STEELS
Table 5. Composition in the middle of Ni-rich and Cu-rich parts of precipitates in 1RK91 by TAP (at.%) Aging
(a) Ni-rich 5 min 1h 2h
(b) Cu-rich 5 min 1h 2h 400 h
Element Al
Ti
Ni
Cu
Mo
12.3 9.6 10.6
3.8 12.8 14.0
16.0 42.5 49.5
7.8 7.8 4.1
Ð 0.6 Ð
Cu
Fe
Ni
Al
Ti
53.6 63.9 70.6 90.0
31.9 14.2 14.9 6.4
5.0 14.2 11.0 5.7
2.8 2.8 5.3 Ð
0.7 2.1 2.4 Ð
Fig. 5. Concentration pro®le through one precipitate in 1RK91 after 1 h of aging.
3.6. Material aged for 4 h Only 1RK91 was investigated in this condition. The precipitates were clearly visible as bright and dark areas in FIM micrographs (Fig. 9) and selected area investigations could be performed. By correlating the contrast on FIM micrographs with the composition of particles it was shown that Curich parts appeared as dark areas in FIM images while Ni-rich parts appeared as spherical bright areas. The precipitates were 4.7 nm in size and their number density was 1.6 1024/m3. Further decrease of Cu content in the Ni-rich precipitates [Table 4(a)] was observed. The Mo content in the matrix was lower than the nominal, indicating the existence of Mo-rich precipitates. This type of particle was indeed found in the material. Molybdenum-rich volumes were observed at the interfaces between the Ni-rich precipitates and the matrix. The measured Mo content in these regions could be as high as 14 at.%. The average composition of the Mo-rich phase is presented in Table 6. In addition to Fe, Cr and Ni the precipitates contained some Si and small amounts of P and C. An almost continuous line of disk-shaped Morich precipitates was found in one specimen, indicating precipitation of this phase at martensite lath boundaries or at prior austenite grain boundaries (Fig. 10). The composition of only one of such precipitates was obtained. It was similar to that obtained for precipitates connected to the Ni-rich particles. 3.7. Material aged for 400 h
Fig. 4. Distribution of elements in precipitates in 1RK91 after 1 and 2 h of aging.
The matrix of 1RK91 alloy in this condition was extremely depleted of Cu. Large amounts of Al, Ti and Mo were also missing in the matrix (Table 3). The Ni- and Mo-rich precipitates were observed both at lath boundaries and in the matrix. The Ni-rich particles became larger, 7±8 nm in size, but preserved their spheroidal shape. Their number density decreased to about 8 1023/m3. The
STILLER et al.: PRECIPITATION IN MARAGING STEELS
Fig. 6. Distribution of elements by TAP in 1RK91 after 1 and 2 h of aging.
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Fig. 7. Concentration pro®le through the interface of one precipitate in 1RK91 after 2 h of aging.
Ni content in the precipitates increased to almost 63 at.% [Table 4(a)]. TAP investigation revealed existence of plates of almost pure Cu [Table 5(b)] on the sides of Ni-rich particles (Fig. 11). The Mo-rich precipitates were observed both in the matrix and at lath boundaries, often in direct contact to the Ni-rich precipitates but not in connection to the Cu-rich plates. They became larger than in the material aged for 4 h and their average Mo content increased (Table 6). Figure 12 shows a three-dimensional reconstruction of one of the precipitates. The observed amount of Mo-rich precipitates in the matrix was much lower than the amount of Ni-rich precipitates. 4. DISCUSSION
The investigation shows that the nucleation and growth of the precipitates are very complex. No doubt, Cu plays a very important role in the
Fig. 8. FIM micrograph of one 1RK91 specimen after 2 h of aging. The position of an observed lath boundary is marked with arrows.
Fig. 9. FIM micrograph of one 1RK91 specimen after 4 h of aging: (a) bright appearing precipitates with two morphologies, spherical and disk-shaped, are marked with arrows; (b) dark appearing precipitates.
nucleation of the Ni-rich phase in both materials. The observation of Cu-rich nuclei, with a Cu content as high as 54 at.% and more widely dispersed Ni-, Al- and Ti-rich regions, at the beginning of the precipitation sequence, indicates that Cu must precipitate ®rst. In 1RK91 the nucleation of the Cu-rich phase is very fast and must occur during a period shorter than 5 min. This early precipitation of Cu should be expected since the solubility of Cu in Fe is known to be low [8]. The solubility of copper in ferrite is approximately 2 wt% at 8508C and much lower for the aging temperature used in the present investigation. The dense and homogeneous distribution of precipitates in the matrix also indicates a high driving force for the precipitation of this phase.
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Table 6. Composition of Mo-rich precipitates by APFIM (at.%) Material
aged 4 h aged 400 h
Element Si
Cr
Fe
Ni
Mo
C
P
1.5420.4 4.0 21.5
14.92 1.4 12.82 2.4
60.4 22.3 38.6 24.2
8.5 21.0 6.8 21.9
13.9 21.2 36.9 23.6
0.3 20.2 1.1 20.8
0.3 20.2 Ð
The Cu-rich particles seem to act as preferential nucleation sites for the precipitation of a phase rich in Ni, Ti and Al. The nucleation of a Ni-rich phase in 1RK91 is also very fast since an enrichment of Ni, Al and Ti was observed in connection to all Cu-rich nuclei after 5 min of aging. The distribution of elements in this phase is, however, inhomogeneous after this aging time. The maximum Ni content in the concentration pro®les by TAP (Fig. 3) is not coincident with the maximum of Al content. The observed shift in the distribution of Al towards Cu-rich nuclei compared to that of Ni, indicates faster diusion of Al in the steel. Another indi-
Fig. 10. FIM micrograph of one 1RK91 specimen after 4 h of aging. An almost continuous line of precipitates indicates the presence of a lath or prior austenite boundary.
Fig. 11. Three-dimensional reconstruction of Cu and Al distribution in a Ni-rich precipitate in a specimen aged for 400 h.
cation for this faster diusion of Al is its high content in the Ni-rich phase at the beginning of precipitation. The ratio between Al and Ni decreases with aging time from 0.8 after 5 min to 0.2 after 2 h [Table 5(a)]. This is to be expected since the ratio between the diusion coecient for Al in a-Fe and the coecient for self-diusion of Fe is about 30, while it is only 2 for Ni. Aluminum is also incorporated into the Cu-rich nuclei to much larger extent than Ti and Ni. After 5 min of aging the Al content in the Cu-rich volumes is about four times that of the bulk. On the other hand the amount of Ni in the center of Cu-rich nuclei is extremely low. It is therefore reasonable to assume that the Ni found in the Nirich parts, after this heat treatment, results from its rejection from the Cu-rich volumes and not from its diusion from the matrix. After 1 and 2 h of aging, the maxima of Ti, Al and Ni are coincident and the observed depletion of elements (especially Al and Ti) close to the precipitate surface (Fig. 5), indicates that the growth process is diusion limited. The calculated number density of precipitates in 1RK91 is highest after 5 min of aging but does not change much before 4 h of aging. The observed growth of the precipitates is also very slow and does not follow the classical size±time0.5 behavior
Fig. 12. Three-dimensional reconstruction of Mo and Si distribution in a Mo-rich precipitate in a specimen aged for 400 h.
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expected from the theory of diusion controlled growth of precipitates. It seems therefore that, because of the very high driving force for precipitation, most of the nucleation and growth occurs before 5 min of aging. After this time the rearrangement of elements in the precipitates, rather than further growth and coarsening, takes place. The precipitates split into two parts of dierent chemistry, and the contributions from Fe and Cr to both the Ni- and the Cu-rich parts decrease. Recent investigation of Cu precipitation in ferritic steels and alloys [9] reveals that the precipitation sequence is initiated by coherent, spherical b.c.c. precipitates which transform to a stable f.c.c. structure (e-phase) via intermediate 9R and 3R phases. It has been also shown that the critical diameter for transformation of the metastable b.c.c. precipitates to 9R structure is approximately 5 nm. This is much larger than the diameter of the observed Cu-rich parts of precipitates after 2 h of aging in 1RK91. Thus, the small size of the Cu- and Ni-rich precipitates together with the apparent lack of any preferential direction for growth of the Ni-rich phase indicate that the precipitates are coherent with the martensitic matrix at this stage. The Ni-rich precipitates preserve their spheroidal shape up to 400 h of aging, while the Cu precipitates form thin shields on the surface of the Ni-rich particles. The change of the shape of the Cu precipitates must be connected to the minimization of the surface energy between the Ni- and Cu-rich precipitates. The observed close ®tting of Cu shields on the Ni precipitates indicates some coherency between the two types of precipitates. If the Ni-rich phase is g' (see Section 5) then its 0.209 nm spacing between {111} planes ®ts best with the 0.207 nm spacing between {111} planes in the e-Cu. The next best candidates are 9R and 3R phases with 0.204 nm spacing between the close packed planes. e-Copper forms in ferritic steels only after prolonged aging and therefore, it seems most probable that Cu shields have 9R or 3R structure. It is also tempting to conclude that the Cu shields work as strain reducers between Ni precipitates and the martensite. This would however, implicate some relationship between the direction of the shields and the matrix which was not observed in the present investigation and therefore more detailed investigation is needed to con®rm this statement. It is interesting that Mo is not incorporated into the Ni-rich precipitates in 1RK91. TAP analysis of precipitates shows that the amount of this element in precipitates is very low for all aging conditions [see Table 5(a)]. Instead Mo enriches the peripheries of Ni-rich precipitates already after 25 min of aging, which shows that this element is rejected from precipitates. This observation is in contradiction to investigations of Ni-rich phases in other maraging steels [10]. According to these studies a large amount of Mo (replacing Ni) was observed in
Z-Ni3Ti precipitates. One possible explanation of the low amount of Mo in our precipitates would be that their structure diers from that of the Z phase. A possible candidate is g'-Ni3Al since the observed Ni-rich phase in 1RK91 includes large amounts of Al. It has been shown that Ti prefers to substitute for the Al sites in g' [11] and that at 11508C more than half of the Al atoms can be replaced by Ti [12]. On the other hand this phase can incorporate Mo only when the Al content is higher than 20 at.% [12]. In our study the highest measured Al content is 12 at.% which shows that Mo should not be included in the precipitates. Taking into consideration all these arguments it seems probable that the observed Ni-rich phase in 1RK91 material is g'. Precipitation in C455 material starts later than in 1RK91, de®nitely after 5 min. It follows, however, the general trend observed in 1RK91, i.e. precipitation of a Ni-rich phase nucleated on the Cu-rich particles. The largest dierence in the chemistry of the investigated alloys is their Mo content, being higher in 1RK91. It can therefore be tempting to explain the eect in terms of Mo speeding up precipitation of Cu, which is assumed to initiate precipitation of the Ni-rich phases. However, according to Thermo-Calc calculations the actual amount of Mo present in 1RK91 should not have any eect on Cu precipitation in the material. Another possible explanation of faster precipitation in 1RK91 alloy is its higher degree of cold deformation (90% cold reduction in 1RK91 as compared to 20% in C455) proceeding aging. The high cold deformation rate used is necessary in 1RK91 because it induces the martensite formation, while in C455 martensite forms during cooling [3]. Thus, the apparent diusivity in 1RK91, which contains a higher dislocation density, should be lower as this line defect acts as a diusion short circuit. There is no doubt that Mo precipitates last in the precipitation sequence. The ®rst signs of Mo clustering were observed after 2 h of aging. However, Mo precipitation at this stage was only initiated, which is also con®rmed by the observed lath boundary free from precipitates. After 4 h of aging, the Mo-rich phase was observed both in connection to the Ni-rich precipitates and at the lath boundaries (Figs 7 and 10). The amount of Mo in the precipitates is close to that in Fe2Mo Laves phase. This intermetallic compound with a hexagonal structure was observed in other maraging steels with low Co content [13, 14]. The hardening eect of 1RK91 after 4 h of aging, which corresponds to the commercial heat treatment, is mainly due to the observed dense distribution of Ni- and Cu-rich precipitates. The contribution of the Mo-rich phase must be less important at this stage. After 400 h of aging the contribution from Mo-rich precipitates to the material hardening becomes more important. However, the
STILLER et al.: PRECIPITATION IN MARAGING STEELS
Ni-rich precipitates are still small enough to be active in the hardening process. The in¯uence of Mo on Ni-rich precipitates could not be elucidated using Thermo-Calc since the accessible data base does not contain information about Ni-rich phases. However, since these precipitates proved a slow coarsening rate in 1RK91 material, while Mo-poor C455 material showed softening only after a few hours at 4758C, due to the coarsening of precipitates [3], it seems reasonable to assume that the slow growth of the precipitates in 1RK91 is caused by the presence of Mo. One possible explanation is that Mo enrichment at the Ni-rich precipitate interfaces (leading to its subsequent precipitation) hampers growth of the Ni-rich phase. Therefore, the superior properties of 1RK91 compared to C455 are due to its high Mo content, which both provides an additional precipitation hardening mechanism (by precipitation of the Morich phase) and impedes coarsening of the Ni-rich phase. 5. CONCLUSIONS
. The precipitation in both investigated materials starts with the formation of Cu-rich particles only after a few minutes at 4758C. These particles act as nucleation sites for a Ni-rich phase of type Ni3(Ti,Al). . After 2 h of aging Cu- and Ni-rich phases are well separated. . Nucleation of a Mo-rich phase on the Ni-rich precipitates in 1RK91 starts after 2 h of aging. . The observed slow growth of Ni-rich precipitates in 1RK91 up to 400 h of aging is most likely due to a hampering eect of Mo that enriches (and precipitates) at their surfaces. . Hardening of 1RK91 after 4 h of aging is mainly due to the Ni- and Cu-rich precipitates, while the
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contribution from precipitation of the Mo-rich phase must be less important at this stage. . The Mo-rich phase provides an additional hardening mechanism after prolonged aging. . The presence of Mo is responsible for the superior properties of 1RK91 compared to C455 after long heat treatment.
AcknowledgementsÐAB Sandvik Steel is acknowledged for supply of material and fruitful discussions.
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