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Microstructural and mechanical properties of rapidly solidified CuNiSn alloys
Materials Science and Engineering, A124 (1990) 223-231
223
Microstructural and Mechanical Properties of Rapidly Solidified Cu-Ni-Sn Alloys L. DEYONG*, R. TREMBLAY and R. ANGERS
Department of Mining and Metallurgy, Laval University, Ste. Foy, Quebec G I K 7P4 (Canada) (Received July 5, 1989; in revised form September 5, 1989)
Abstract
Rapid solidification (RS) by melt spinning of Cu-Ni-Sn alloys rendered more chemically homogeneous materials beneficial to mechanical strengthening. The structure was microcrystalline and manifested none of the tin segregation evident in conventionally cast alloys. Although there was microsegregation of tin in the form of y precipitates at the grain boundaries of rapidly solidified ribbons, their degree of segregation was significantly reduced. Age hardening increased with tin (by spinodal decomposition) and nickel (by alloy hardening) content of the alloy. Solid solubility of tin in a high tin content alloy otherwise unattainable by conventional methods was achieved by RS. I. Introduction Cu-Ni-Sn alloys have been developed for use as high performance connector materials in electronics applications [1, 2]. The spinodal age hardening of these alloys leads to outstanding mechanical properties. However, beneficial exploitation of these alloys has been hindered by severe segregation of tin during solidification. In order to overcome these segregation problems, they were made by rapid solidification via melt spinning and the resultant properties thus developed were compared to those properties obtained by conventional practice. Alloys assessment was also made by microstructural and mechanical analyses. The alloys studied were Cu-15Ni-8Sn and a series of Cu-10Ni-xSn alloys with tin concentrations of 6, 8, 10 and 12 wt.%. The rationale for examining this particular range of Cu-Ni-Sn alloys follows. *Present address: Alcan Rolled Products Company, Kingston Works, PO Box 2000, Kingston, Ontario K7L 4Z5, Canada. 0921-5093/90/$3.50
(1) Study the effect of tin content on the mechanical properties of Cu-Ni-Sn alloys keeping nickel fixed at 10 wt.%. (2) Likewise study the effect of nickel content keeping tin fixed at 8 wt.%. (3) Compare the mechanical performance of a range of Cu-Ni-Sn alloys, including the present most popular composition, Cu- 15Ni-8Sn. (4) Study the mechanical performance of an alloy composition, Cu-10Ni-12Sn, that cannot be obtained in solid solution at any temperature by conventional methods.
2. Experimental procedure All alloys were fabricated by melting the constituent elements in an induction furnace under argon. The resultant cast ingots (1.2 cm diameter by 8 cm length) were homogenized for 50 h at 825 °C under argon and then water quenched. Disks (0.5 mm thick) were cut from the homogenized ingots and used as samples of conventionally processed materials. The rapidly solidified materials were produced by melt spinning of pieces (30 g) from the homogenized ingots. These pieces were heated to 40 °C above the liquidus temperature for melt spinning. The liquidus and solidus temperatures of the alloys studied are given in Table 1. Details of this chill block melt spin apparatus are reported in refs. 3 and 4. The resultant chemical compositions for these alloys in ingot and ribbon form are shown in Table 2. These alloys age harden by spinodal decomposition. Both ribbon and ingot samples were thus age hardened at 400 °C in a resistance furnace under argon for all alloy compositions. Chemical homogeneity was examined by electron microprobe analysis for ribbons and by energy-dispersive spectroscopy (EDS) for ingots. Phase identification was made by X-ray © Elsevier Sequoia/Printed in The Netherlands
224 TABLE 1
Liquidus and solidus temperatures for Cu-Ni-Sn alloys
Tliq (°C) T~ol(°C)
TABLE 2
Cu- t ONi-6Sn
Cu- l ONi-8Sn
Cu- l ONi- I OSn
Cu- l ONi- 12Sn
Cu- 15Ni-8Sn
11 I0 953
1090 923
1075 895
1060 875
1115 950
Chemical composition of conventionally prepared and melt-spun alloys Element
Composition (wt.%) Cu-lONi-6Sn
Cu-lONi-8Sn
Cu-lONi-lOSn
Cu-lONi-12Sn
Cu-15Ni-8Sn
Conventionally cast ingots
Sn Ni Cu
6.03 9.58 Bal.
7.92 9.63 Bal.
9.47 9.48 Bal.
12.02 9.89 Bal.
7.81 14.95 Bal.
Rapidly solidified ribbons
Sn Ni Cu
5.29 10.29 Bal.
8.05 10.13 Bal.
10.22 9.57 Bal.
11.83 9.99 Bal.
8.18 15.37 Bal.
diffraction analysis. Ribbon samples were mounted on edge, polished and etched with alcoholic ferric chloride solution (FeCl3, 5 g; HCI, 2.5 ml; alcohol, 93.5 ml) for metallographic examination. Knoop microhardness measurements were taken on all samples using a 20 g load. 3. Results
3.1. Microstructure 3.1.1. Ingots The solution heat treatment of the as-cast ingots served to eliminate the severe segregation of tin manifested by cored dendrites formed during casting. Since the cooling rate of the water quench (around 103 °C S - 1) was n,~t quick enough to avoid a secondary phase precipitation, the resulting alloys consisted of an a matrix containing tiny y phase particles (~, has the formula ( C u x N i l _ x ) 3 s n with approximately 40% Ni, 40% Sn and D O 3 structure)[5, 6]. Figure 1 shows the microstructure of such conventionally processed alloys representative of alloys containing up to 10wt.%Sn. All ingot alloys had similar microstructures except Cu-10Ni-12Sn, a two-phase material where the homogenization heat treatment nodularized the gamma phase (Fig. 2). 3.1.2. Ribbons The microstructures tor all ribbon alloys, including Cu-10Ni-12Sn, were similar with grain sizes of the order of 1 pm. These ribbons were on the average 40 ktm thick and 2.5 mm wide. A representative micrograph is shown in Fig. 3. Such rapidly solidified microcrystalline materials are
Fig. 1. Typical microstructure of conventionally cast and homogenized Cu-10Ni-xSn ingot (representative of alloys containing 6 - 10 wt.% Sn).
typically produced with a cooling rate of the order of 105 °C s-~. Two regions were evident in the microstructure of the rapidly solidified material: a columnar crystal structure growing from the substrate surface and a fine equiaxed microstructure on the free side of the ribbon. At higher magnification, precipitates can be seen along the grain boundaries (Fig. 4). These precipitates are fewer and finer near the substrate surface. Towards the ribbon free surface, the precipitates become elongated. This is due to a greater microsegregation of tin to the cell boundaries. This fact was substantiated by microprobe
225
Q
•
.
~.
200 I~m
Fig. 2. Microstructure of conventionally cast and homogenized Cu- 10Ni- 12Sn ingot.
Fig. 4. Higher magnification of a ribbon microstructure' showing V precipitates at grain boundaries (substrate side; Cu- 10Ni- 10Sn alloy).
cipitation increases with tin content. Moreover, a quantitative electron microprobe analysis showed no definitive trend in the degree of segregation with tin concentration. It was noted though that the ribbon surface that was in direct contact with the substrate had the greatest solidification rate and consequently the least microsegregation of tin. Also, the compositional variation at any depth in the ribbon was far less than that in the homogenized ingots [8, 9].
3.2. Energy-dispersive spectroscopy of ingots
Fig. 3. Microstructure of as-melt-spun Cu-10Ni-8Sn ribbon typical of all alloy compositions (SEM; transverse sectioh; substrate surface on bottom).
analysis, which showed that tin segregation increases with depth from the substrate surface. It has been reported that these precipitates are rich in tin and nickel compared to the matrix [7]. Clearly, for these alloys, such precipitates are the 7 phase (CuxNi I x)3Sn found in other studies [5, 6,8,9]. From the micrographs it is difficult to determine qualitatively whether grain boundary pre-
On a microscopic scale, EDS confirmed the chemical analyses of the ingot slices containing 6-10wt.%Sn. As for the 12wt.%Sn alloy, chemical analysis gave the nominal Cu-10Ni-12Sn composition but EDS revealed that two distinct phases were present. The difference in composition between precipitates and matrix in Cu-10Ni-12Sn is illustrated in Fig. 5. This micrograph indicates three different zones of composition. The nodular precipitates (B) were found to be very rich in tin and nickel whereas the matrix (A) was tin poor with nickel at the nominal composition. Interestingly, the matrix region (C) between intranocular 7 lamellae was nickel poor and tin rich. That the nickel remained at the nominal composition in the bulk matrix indicates that nickel diffused locally in the tin-rich regions to form the 7 pre-
226
cipitates which created within the precipitate nodules a spheroidized pearlitic structure with nickel-rich (),) and nickel-poor (a) regions.
3.3. Electron microprobe analysis of ribbons Electron microprobe analysis of tin at various depths within the as-melt-spun ribbons showed similar characteristics for all alloys. Figure 6 shows a microprobe trace for Cu-10Ni-8Sn. Traces were taken at the substrate surface, the mid-thickness and the free surface. From these traces the chemical variations were quantified by a method that permitted the determination of the range in tin content, Rsn, as the difference between the average of the five highest values of tin content and the average of the five lowest
values observed over a scanning length of 32/~m [7]. Table 3 shows the Rsn values for all alloys at the three different depths. There was very little variation in tin content at the substrate surface where it was generally within a range of about 0.6wt.%. Away from the substrate surface, tin segregation at the cell boundaries was consonant with the observed precipitates found in this region. Tin segregation was greatest at the free surface where the Rsn value attained 2.3 wt.%. Table 3 does not manifest a conspicuous trend towards greater segregation with increasing tin content. Although the 6 wt.% Sn alloy had undisputedly the least segregation, the rest of the alloys seem to have Rsn values independent of alloy composition for a given depth. Segregation, especially at the substrate surface, will depend on factors that govern the solidification rate, among them ribbon thickness. As for the variation of
9 8
7 A
~ml0 9 t/) --i/:
ii¸
Fig. 5. Regions of compositional variation in Cu-10Ni-12Sn ingot: A, matrix (10wt.%Ni, 9wt.%Sn); B, nodular )' phase (15wt.%Ni, 35wt.%Sn); C, intranodular a phase (7wt.%Ni, 13wt.%Sn); D, fine quench-induced y precipitate.
TABLE 3 ribbon
7 6 LENGTH = 25z,~m
Fig. 6. Microprobe trace showing tin segregation at various depths within Cu-10Ni-8Sn ribbon (typical of all melt-spun alloys).
Range in chemical variation measured by electron microprobe analysis at three different depths in the melt-spun
Ribbon depth
Substrate surface
Range
Cu-lONi-6Sn
Cu-lONi-SSn
Cu-lONi-lOSn
Cu-lONi-12Sn
Cu-15Ni-8Sn
All alloys (average)
0.32 1.20
0.30 1.41
0.63 1.43
0.48 1.13
0.36 1.70
0.42 1.37
RNi
0.62 1.20
1.35 1.35
0.78 1.37
1.31 1.52
1.62 1.16
1.14 1.32
Rs, RNi
1.60 1.40
2.00 1.62
1.71 1.52
2.03 1.36
2.32 1.70
1.93 1.52
Rsn
RNi Mid-thickness Outer surface
Chemical variation range (wt.%)
Rs.
227
nickel content, _+0.7wt.%, this was consistent throughout for all alloys at all ribbon depths.
3.4. Age hardening of conventionallyprepared ingots The Knoop microhardness curves with respect to aging time for all conventionally prepared alloys (except for Cu-10Ni-12Sn) are shown in Fig. 7. The curves interpolate the average value of ten microhardness measurements taken at eleven heat treatment durations for each alloy. The standard deviation of the Knoop hardness for each heat treatment duration was on the average 18 HKN (Hardness Knoop Numbers) or about 7%. Alloys with higher tin content have a greater maximum hardness, which confirms that it is primarily the tin which contributes to the hardening by spinodal decomposition. These curves exhibit the same characteristics as the aging response for other mechanical properties such as yield strength. In fact, the aging time to maximum hardness corroborates with the aging time to maximum strength for similar alloys studied elsewhere. For Cu-15Ni-8Sn the time to maximum yield strength was found to be 1 h [6]; for Cu-9Ni-6Sn it was 2.5 h [10]. Overaging severely reduces hardness from the maximum value. No microhardness measurements for the 12 wt.% Sn alloy are reported because the resolution of those measurements was far too great to make any meaningful assessment of this alloy hardness as an ensemble of phases.
.160
--
I I 10 Ni-6Sn 10 Ni-8Sn 10 Ni- 10Sn 15 Nt-8Sn
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3.5. Age hardening of rapidly solidified ribbons
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220 0.0
The 12wt.%Sn alloy presents an anomaly to the interpretation of microhardness measurements owing to its dual-phase nature. The y precipitates are very hard (around 465 HKN) in comparison to the matrix (around 280 HKN before heat treatment). Notice in Fig. 2 how polishing has removed the a phase, leaving the harder precipitate nodules to protrude from the matrix relief. In addition, the matrix is not of the nominal composition and so Knoop hardness measurements in the matrix resemble those determined in an alloy lower in tin concentration. Nickel apparently improves mechanical properties as the maximum hardness increases with the change in nickel content from 10 to 15 wt.%, but does not alter the aging time to maximum hardness. Nickel serves to strengthen the matrix by alloy hardening and not by any phase transformation. As inferred from the age-hardening curves (Fig. 7), the Cu-15Ni-8Sn alloy potentially offers the greatest strength. Figure 8 indicates the microstructural aging response of Cu-10Ni-8Sn, which is representative of the aging response for ingot alloys containing 6-10wt.%Sn. The figure shows specimens before heat treatment, at maximum hardness, slightly overaged, and overaged for 2040 rain. At maximum hardness (Fig. 8(b)) the onset of the discontinuous transformation appears to commence at the grain boundaries. This discontinuous transformation continues into the matrix (Fig. 8(c)) until the material has entirely transformed, rendering the overaged condition (Fig. 8(d)). The same behaviour is also observed within the matrix of the 12 wt.% Sn alloy. Apart from the large y precipitate nodules, the matrix of this alloy undergoes spinodal decomposition and discontinuous precipitation like the other alloys.
I
I
I
I
I
I
I
I
0.4
0.8
1.2
1.6
2.0
2.4
2.8
3.2
LOG ( a g i n g t i m e ) , m i n u t e s
Fig. 7. Hardness as a function of aging time at 400 °C for conventionally cast C u - N i - S n ingots.
Age hardening of ribbons follows the same tendencies as observed for ingots except that here Cu-10Ni-12Sn may be included in the comparison (Fig. 9). The homogeneity offered by rapid solidification gave greater distinction between the mechanical properties for each alloy. The standard deviation of the Knoop hardness measurements for each heat treatment duration was at the most 12 HKN (5%) for ribbons whereas for ingots the deviation varied between 11 and 23 HKN (4% and 8%). Changing the nickel content from 10 to 15 wt.% did not alter the aging time to maximum
228
Fig. 8. Effectof aging time at 400 °C on the microstructure of Cu-10Ni-8Sn ingot:(a) 0 min, as-cast and homogenized;(b) 67 rain, maximumhardness; (c) 90 min, slightlyoveraged;(d) 2040 min, overaged microstructure.
hardness but increased the maximum hardness and raised the overall hardness curve. C u - 1 5 N i - 8 S n has a maximum hardness comparable to that of Cu-10Ni-12Sn. All ribbon alloys evolved a similar microstructural development with aging. The microstructural aging response for C u - 1 0 N i - 1 2 S n is shown in Fig. 10. Again the microstructure at maximum hardness manifests a partial discontinuous transformation product (Fig. 10(b)). The overaged condition (Fig. 10(c)) shows the resultant discontinuous transformation with the formation of duplex cells composed of the a and ~ phases.
Overaging brought about by cellular discontinuous precipitation produced similar dualphase structures in all alloys and both nickel and tin contents have little effect on differentiating mechanical performance. 4. Discussion
4.1. Microstructural comparison between conventionally processed ingots and rapidly solidified ribbons A certain inhomogeneity of both ribbons and ingots has been observed owing to a practically
229
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Fig. 9. Hardness as a function of aging time at 400 °C for rapidly solidified C u - N i - S n ribbons.
unavoidable V precipitation no matter what the casting process. However, this precipitation is on a much finer scale in ribbons than in ingots. For the rapidly solidified material, as can be seen in the micrographs and the microprobe trace data, segregation spacing was of the order of 5/~m. By comparison, it was reported that a continuous-cast Cu-10Ni-2Sn plate (15 mm thick) had a segregation spacing (secondary dendrite arm spacing) of the order of 20 /~m where tin varied from 0.6 to 3.6 wt.% across the dendritic structure [7]. The pseudobinary equilibrium phase diagram (for Cu-9Ni-xSn) indicates that for these alloys a melt must be cooled from liquidus to solidus by at least 160 °C to obtain a single phase, a [9]. This undercooling was ostensibly obtained at the substrate surface of the ribbons. However, the rate of heat extraction diminishes with distance from the substrate surface and the required cooling was not obtained in the ribbon towards the free surface. A segregated microstructure with secondary phase particles at cell boundaries then resulted in these regions. The homogenization heat treatment served to decrease segregation in dendritic structures and at grain boundaries in the ingots. However, even though these ingots were relatively small (1.2 cm diameter by 8 cm long cylinder) compared to commercially available products, V precipitation was unavoidable despite a water quench after the solution heat treatment, since a and Y are stable at lower temperatures and the cooling rate was
Fig. 10. Microstructural response to aging at 400°C in Cu-10Ni-12Sn rapidly solidified ribbons: (a) 0 min, as melt spun; (b) 45 min, maximum hardness; (c) 24 h, overaged.
230
not fast enough to keep tin in supersaturated solution. Although the heat extraction in the melt-spinning process utilized in this study was not rapid enough to completely avoid precipitation, the rapidly solidified materials had significantly reduced segregation spacing and greatly diminished chemical inhomogeneity in comparison to the as-cast alloys. Both of these effects are beneficial in curtailing the requirements of subsequent homogenization treatments. An evident case where rapid solidification retains homogeneity when conventional casting processes cannot is the Cu-10Ni-12Sn composition. This alloy in ribbon form has a uniform microstructure no different from such alloys with lower tin concentration, whereas in conventionally prepared materials a strongly heterogeneous microstructure is apparent.
4.2. Comparison of age hardening between conventionallyprocessed ingots and rapidly solidified ribbons The general age-hardening behaviours for both ingots and ribbons are similar. Equivalent maximum hardness values were obtained for the same alloy composition in both ingot and ribbon forms. However, the heat treatment duration to maximum hardness for ribbons was slightly longer than for ingots. This was due to the fact that the cooling rate of the ingot by water quench (from the homogenizing temperature of 825 °C) was less than that of the ribbon by melt spinning. Thus spinodal decomposition had already advanced in the ingot ahead of the ribbon as confirmed by hardness values measured at zero aging time. Rapid sofidification leads to mechanical characteristics (as implied by microhardness testing) equivalent to homogenized ingots in addition to greater homogeneity of both chemical and mechanical properties within the material.
5. Conclusions
The effect of alloy composition and solidification processing on mechanical properties has been studied for a range of Cu-Ni-Sn alloys. Clearly chemical homogeneity was the key factor in all aspects of material performance. Melt spinning was found to be an excellent means for reducing tin segregation which could otherwise
be eliminated only by complex processing methods. Age hardening by spinodal composition was studied in both rapidly solidified and conventionally processed materials, and mechanical properties were found to be comparable in both materials as inferred from microhardness tests. Where rapidly solidified alloys manifest superiority over their conventional counterparts is in the uniformity of metallurgical properties throughout the material on a microscopic scale. This has been confirmed by Knoop microhardness measurements, microprobe and X-ray diffraction analyses and metallography. The results obtained from mechanical and microstructural investigations lead to the following conclusions. (1) Segregation and segregation spacing in Cu-Ni-Sn alloys were significantly reduced by rapid solidification. Microsegregation persisted in the form of 7 precipitates at grain boundaries. (2) Elemental tin segregation in the melt-spun material was independent of the alloy content of tin ranging from 6 to 12 wt.%. (3) Rapid solidification by melt spinning allowed the solid solubility of tin at a higher concentration otherwise unattainable by conventional alloy preparation techniques. (4) A new homogeneous Cu-10Ni-12Sn rapidly solidified alloy, unavailable by conventional means, potentially offers mechanical strength greater than currently available alloys with nickel content less than or equal to 10 wt.%. (5) The age-hardening response of rapidly solidified alloys was similar to that of their conventional counterpart. (6) Increasing the tin content served to augment the maximum hardness of the alloy by spinodal decomposition. This maximum hardness was similar irrespective of the preparation method (except for the C u - 10Ni- 12Sn alloy). (7) Hardness increases with nickel content by alloy hardening of the matrix. (8) As indicated by the variation in microhardness values, the rapidly solidified alloys had mechanical properties more uniform throughout the material. Rapid solidification technology improves the material performance of Cu-Ni-Sn alloys, including greater resistance to corrosion [11]. It remains for further research to investigate the integrity of working pieces consolidated from such rapidly solidified materials.
231
References 1 J. M. DeVanssay, Matdriaux et Technique, 75 (1-2) (1987) 37. 2 C.R. Scorey, S. Chin, M. J. White and R. J. Livak, J. Met., 36 (11)(1984)52. 3 L. Deyong, Rapidly solidified Cu-Ni-Sn alloys: their metallurgical and electrochemical properties, M.Sc. Thesis, Laval University, Quebec, 1989. 4 R. Angers and C. Gelinas, Fabrication of shaped charge liners from rapidly solidified powders. Rep. DSS File lOSD-97701-6-O768-A, January 1988 (Defense Research Establishment, Valcartier, Quebec). 5 M. Miki and Y. Ogino, Trans. Jpn. Inst. Met., 25 (9)
(1984) 593. 6 B. G. Lefevre, A. T. D'Annessa and D. Kalish, Metall. Trans. A, 9(1978) 577. 7 L. E. Collins, J. R. Barry and J. Masounave, Reduced segregation in rapidly solidified Cu-Ni-Sn alloys, Rep. PRML 87-43(OP-J), June 1987 (CANMET, Physical Metallurgy Research Laboratories, Ottawa, Ontario). 8 Y. Watanabe, Y, Murakami and S. Kachi, J. Jpn. Inst. Met., 45(1981) 551. 9 J.T. Plewes, Metall. Trans. A, 6 (1975), 537. 10 P. Kratochvil, J. Mencl and J. Pesicka, Freib. Forschungsh. B, Metall. Werkstofftech., 256(1986) 87. 11 L. Deyong, M. Elboujdai'ni, R. Tremblay and E. Ghali, J. Appl. Electrochemistry, in the press.
223
Microstructural and Mechanical Properties of Rapidly Solidified Cu-Ni-Sn Alloys L. DEYONG*, R. TREMBLAY and R. ANGERS
Department of Mining and Metallurgy, Laval University, Ste. Foy, Quebec G I K 7P4 (Canada) (Received July 5, 1989; in revised form September 5, 1989)
Abstract
Rapid solidification (RS) by melt spinning of Cu-Ni-Sn alloys rendered more chemically homogeneous materials beneficial to mechanical strengthening. The structure was microcrystalline and manifested none of the tin segregation evident in conventionally cast alloys. Although there was microsegregation of tin in the form of y precipitates at the grain boundaries of rapidly solidified ribbons, their degree of segregation was significantly reduced. Age hardening increased with tin (by spinodal decomposition) and nickel (by alloy hardening) content of the alloy. Solid solubility of tin in a high tin content alloy otherwise unattainable by conventional methods was achieved by RS. I. Introduction Cu-Ni-Sn alloys have been developed for use as high performance connector materials in electronics applications [1, 2]. The spinodal age hardening of these alloys leads to outstanding mechanical properties. However, beneficial exploitation of these alloys has been hindered by severe segregation of tin during solidification. In order to overcome these segregation problems, they were made by rapid solidification via melt spinning and the resultant properties thus developed were compared to those properties obtained by conventional practice. Alloys assessment was also made by microstructural and mechanical analyses. The alloys studied were Cu-15Ni-8Sn and a series of Cu-10Ni-xSn alloys with tin concentrations of 6, 8, 10 and 12 wt.%. The rationale for examining this particular range of Cu-Ni-Sn alloys follows. *Present address: Alcan Rolled Products Company, Kingston Works, PO Box 2000, Kingston, Ontario K7L 4Z5, Canada. 0921-5093/90/$3.50
(1) Study the effect of tin content on the mechanical properties of Cu-Ni-Sn alloys keeping nickel fixed at 10 wt.%. (2) Likewise study the effect of nickel content keeping tin fixed at 8 wt.%. (3) Compare the mechanical performance of a range of Cu-Ni-Sn alloys, including the present most popular composition, Cu- 15Ni-8Sn. (4) Study the mechanical performance of an alloy composition, Cu-10Ni-12Sn, that cannot be obtained in solid solution at any temperature by conventional methods.
2. Experimental procedure All alloys were fabricated by melting the constituent elements in an induction furnace under argon. The resultant cast ingots (1.2 cm diameter by 8 cm length) were homogenized for 50 h at 825 °C under argon and then water quenched. Disks (0.5 mm thick) were cut from the homogenized ingots and used as samples of conventionally processed materials. The rapidly solidified materials were produced by melt spinning of pieces (30 g) from the homogenized ingots. These pieces were heated to 40 °C above the liquidus temperature for melt spinning. The liquidus and solidus temperatures of the alloys studied are given in Table 1. Details of this chill block melt spin apparatus are reported in refs. 3 and 4. The resultant chemical compositions for these alloys in ingot and ribbon form are shown in Table 2. These alloys age harden by spinodal decomposition. Both ribbon and ingot samples were thus age hardened at 400 °C in a resistance furnace under argon for all alloy compositions. Chemical homogeneity was examined by electron microprobe analysis for ribbons and by energy-dispersive spectroscopy (EDS) for ingots. Phase identification was made by X-ray © Elsevier Sequoia/Printed in The Netherlands
224 TABLE 1
Liquidus and solidus temperatures for Cu-Ni-Sn alloys
Tliq (°C) T~ol(°C)
TABLE 2
Cu- t ONi-6Sn
Cu- l ONi-8Sn
Cu- l ONi- I OSn
Cu- l ONi- 12Sn
Cu- 15Ni-8Sn
11 I0 953
1090 923
1075 895
1060 875
1115 950
Chemical composition of conventionally prepared and melt-spun alloys Element
Composition (wt.%) Cu-lONi-6Sn
Cu-lONi-8Sn
Cu-lONi-lOSn
Cu-lONi-12Sn
Cu-15Ni-8Sn
Conventionally cast ingots
Sn Ni Cu
6.03 9.58 Bal.
7.92 9.63 Bal.
9.47 9.48 Bal.
12.02 9.89 Bal.
7.81 14.95 Bal.
Rapidly solidified ribbons
Sn Ni Cu
5.29 10.29 Bal.
8.05 10.13 Bal.
10.22 9.57 Bal.
11.83 9.99 Bal.
8.18 15.37 Bal.
diffraction analysis. Ribbon samples were mounted on edge, polished and etched with alcoholic ferric chloride solution (FeCl3, 5 g; HCI, 2.5 ml; alcohol, 93.5 ml) for metallographic examination. Knoop microhardness measurements were taken on all samples using a 20 g load. 3. Results
3.1. Microstructure 3.1.1. Ingots The solution heat treatment of the as-cast ingots served to eliminate the severe segregation of tin manifested by cored dendrites formed during casting. Since the cooling rate of the water quench (around 103 °C S - 1) was n,~t quick enough to avoid a secondary phase precipitation, the resulting alloys consisted of an a matrix containing tiny y phase particles (~, has the formula ( C u x N i l _ x ) 3 s n with approximately 40% Ni, 40% Sn and D O 3 structure)[5, 6]. Figure 1 shows the microstructure of such conventionally processed alloys representative of alloys containing up to 10wt.%Sn. All ingot alloys had similar microstructures except Cu-10Ni-12Sn, a two-phase material where the homogenization heat treatment nodularized the gamma phase (Fig. 2). 3.1.2. Ribbons The microstructures tor all ribbon alloys, including Cu-10Ni-12Sn, were similar with grain sizes of the order of 1 pm. These ribbons were on the average 40 ktm thick and 2.5 mm wide. A representative micrograph is shown in Fig. 3. Such rapidly solidified microcrystalline materials are
Fig. 1. Typical microstructure of conventionally cast and homogenized Cu-10Ni-xSn ingot (representative of alloys containing 6 - 10 wt.% Sn).
typically produced with a cooling rate of the order of 105 °C s-~. Two regions were evident in the microstructure of the rapidly solidified material: a columnar crystal structure growing from the substrate surface and a fine equiaxed microstructure on the free side of the ribbon. At higher magnification, precipitates can be seen along the grain boundaries (Fig. 4). These precipitates are fewer and finer near the substrate surface. Towards the ribbon free surface, the precipitates become elongated. This is due to a greater microsegregation of tin to the cell boundaries. This fact was substantiated by microprobe
225
Q
•
.
~.
200 I~m
Fig. 2. Microstructure of conventionally cast and homogenized Cu- 10Ni- 12Sn ingot.
Fig. 4. Higher magnification of a ribbon microstructure' showing V precipitates at grain boundaries (substrate side; Cu- 10Ni- 10Sn alloy).
cipitation increases with tin content. Moreover, a quantitative electron microprobe analysis showed no definitive trend in the degree of segregation with tin concentration. It was noted though that the ribbon surface that was in direct contact with the substrate had the greatest solidification rate and consequently the least microsegregation of tin. Also, the compositional variation at any depth in the ribbon was far less than that in the homogenized ingots [8, 9].
3.2. Energy-dispersive spectroscopy of ingots
Fig. 3. Microstructure of as-melt-spun Cu-10Ni-8Sn ribbon typical of all alloy compositions (SEM; transverse sectioh; substrate surface on bottom).
analysis, which showed that tin segregation increases with depth from the substrate surface. It has been reported that these precipitates are rich in tin and nickel compared to the matrix [7]. Clearly, for these alloys, such precipitates are the 7 phase (CuxNi I x)3Sn found in other studies [5, 6,8,9]. From the micrographs it is difficult to determine qualitatively whether grain boundary pre-
On a microscopic scale, EDS confirmed the chemical analyses of the ingot slices containing 6-10wt.%Sn. As for the 12wt.%Sn alloy, chemical analysis gave the nominal Cu-10Ni-12Sn composition but EDS revealed that two distinct phases were present. The difference in composition between precipitates and matrix in Cu-10Ni-12Sn is illustrated in Fig. 5. This micrograph indicates three different zones of composition. The nodular precipitates (B) were found to be very rich in tin and nickel whereas the matrix (A) was tin poor with nickel at the nominal composition. Interestingly, the matrix region (C) between intranocular 7 lamellae was nickel poor and tin rich. That the nickel remained at the nominal composition in the bulk matrix indicates that nickel diffused locally in the tin-rich regions to form the 7 pre-
226
cipitates which created within the precipitate nodules a spheroidized pearlitic structure with nickel-rich (),) and nickel-poor (a) regions.
3.3. Electron microprobe analysis of ribbons Electron microprobe analysis of tin at various depths within the as-melt-spun ribbons showed similar characteristics for all alloys. Figure 6 shows a microprobe trace for Cu-10Ni-8Sn. Traces were taken at the substrate surface, the mid-thickness and the free surface. From these traces the chemical variations were quantified by a method that permitted the determination of the range in tin content, Rsn, as the difference between the average of the five highest values of tin content and the average of the five lowest
values observed over a scanning length of 32/~m [7]. Table 3 shows the Rsn values for all alloys at the three different depths. There was very little variation in tin content at the substrate surface where it was generally within a range of about 0.6wt.%. Away from the substrate surface, tin segregation at the cell boundaries was consonant with the observed precipitates found in this region. Tin segregation was greatest at the free surface where the Rsn value attained 2.3 wt.%. Table 3 does not manifest a conspicuous trend towards greater segregation with increasing tin content. Although the 6 wt.% Sn alloy had undisputedly the least segregation, the rest of the alloys seem to have Rsn values independent of alloy composition for a given depth. Segregation, especially at the substrate surface, will depend on factors that govern the solidification rate, among them ribbon thickness. As for the variation of
9 8
7 A
~ml0 9 t/) --i/:
ii¸
Fig. 5. Regions of compositional variation in Cu-10Ni-12Sn ingot: A, matrix (10wt.%Ni, 9wt.%Sn); B, nodular )' phase (15wt.%Ni, 35wt.%Sn); C, intranodular a phase (7wt.%Ni, 13wt.%Sn); D, fine quench-induced y precipitate.
TABLE 3 ribbon
7 6 LENGTH = 25z,~m
Fig. 6. Microprobe trace showing tin segregation at various depths within Cu-10Ni-8Sn ribbon (typical of all melt-spun alloys).
Range in chemical variation measured by electron microprobe analysis at three different depths in the melt-spun
Ribbon depth
Substrate surface
Range
Cu-lONi-6Sn
Cu-lONi-SSn
Cu-lONi-lOSn
Cu-lONi-12Sn
Cu-15Ni-8Sn
All alloys (average)
0.32 1.20
0.30 1.41
0.63 1.43
0.48 1.13
0.36 1.70
0.42 1.37
RNi
0.62 1.20
1.35 1.35
0.78 1.37
1.31 1.52
1.62 1.16
1.14 1.32
Rs, RNi
1.60 1.40
2.00 1.62
1.71 1.52
2.03 1.36
2.32 1.70
1.93 1.52
Rsn
RNi Mid-thickness Outer surface
Chemical variation range (wt.%)
Rs.
227
nickel content, _+0.7wt.%, this was consistent throughout for all alloys at all ribbon depths.
3.4. Age hardening of conventionallyprepared ingots The Knoop microhardness curves with respect to aging time for all conventionally prepared alloys (except for Cu-10Ni-12Sn) are shown in Fig. 7. The curves interpolate the average value of ten microhardness measurements taken at eleven heat treatment durations for each alloy. The standard deviation of the Knoop hardness for each heat treatment duration was on the average 18 HKN (Hardness Knoop Numbers) or about 7%. Alloys with higher tin content have a greater maximum hardness, which confirms that it is primarily the tin which contributes to the hardening by spinodal decomposition. These curves exhibit the same characteristics as the aging response for other mechanical properties such as yield strength. In fact, the aging time to maximum hardness corroborates with the aging time to maximum strength for similar alloys studied elsewhere. For Cu-15Ni-8Sn the time to maximum yield strength was found to be 1 h [6]; for Cu-9Ni-6Sn it was 2.5 h [10]. Overaging severely reduces hardness from the maximum value. No microhardness measurements for the 12 wt.% Sn alloy are reported because the resolution of those measurements was far too great to make any meaningful assessment of this alloy hardness as an ensemble of phases.
.160
--
I I 10 Ni-6Sn 10 Ni-8Sn 10 Ni- 10Sn 15 Nt-8Sn
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3.5. Age hardening of rapidly solidified ribbons
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o
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220 0.0
The 12wt.%Sn alloy presents an anomaly to the interpretation of microhardness measurements owing to its dual-phase nature. The y precipitates are very hard (around 465 HKN) in comparison to the matrix (around 280 HKN before heat treatment). Notice in Fig. 2 how polishing has removed the a phase, leaving the harder precipitate nodules to protrude from the matrix relief. In addition, the matrix is not of the nominal composition and so Knoop hardness measurements in the matrix resemble those determined in an alloy lower in tin concentration. Nickel apparently improves mechanical properties as the maximum hardness increases with the change in nickel content from 10 to 15 wt.%, but does not alter the aging time to maximum hardness. Nickel serves to strengthen the matrix by alloy hardening and not by any phase transformation. As inferred from the age-hardening curves (Fig. 7), the Cu-15Ni-8Sn alloy potentially offers the greatest strength. Figure 8 indicates the microstructural aging response of Cu-10Ni-8Sn, which is representative of the aging response for ingot alloys containing 6-10wt.%Sn. The figure shows specimens before heat treatment, at maximum hardness, slightly overaged, and overaged for 2040 rain. At maximum hardness (Fig. 8(b)) the onset of the discontinuous transformation appears to commence at the grain boundaries. This discontinuous transformation continues into the matrix (Fig. 8(c)) until the material has entirely transformed, rendering the overaged condition (Fig. 8(d)). The same behaviour is also observed within the matrix of the 12 wt.% Sn alloy. Apart from the large y precipitate nodules, the matrix of this alloy undergoes spinodal decomposition and discontinuous precipitation like the other alloys.
I
I
I
I
I
I
I
I
0.4
0.8
1.2
1.6
2.0
2.4
2.8
3.2
LOG ( a g i n g t i m e ) , m i n u t e s
Fig. 7. Hardness as a function of aging time at 400 °C for conventionally cast C u - N i - S n ingots.
Age hardening of ribbons follows the same tendencies as observed for ingots except that here Cu-10Ni-12Sn may be included in the comparison (Fig. 9). The homogeneity offered by rapid solidification gave greater distinction between the mechanical properties for each alloy. The standard deviation of the Knoop hardness measurements for each heat treatment duration was at the most 12 HKN (5%) for ribbons whereas for ingots the deviation varied between 11 and 23 HKN (4% and 8%). Changing the nickel content from 10 to 15 wt.% did not alter the aging time to maximum
228
Fig. 8. Effectof aging time at 400 °C on the microstructure of Cu-10Ni-8Sn ingot:(a) 0 min, as-cast and homogenized;(b) 67 rain, maximumhardness; (c) 90 min, slightlyoveraged;(d) 2040 min, overaged microstructure.
hardness but increased the maximum hardness and raised the overall hardness curve. C u - 1 5 N i - 8 S n has a maximum hardness comparable to that of Cu-10Ni-12Sn. All ribbon alloys evolved a similar microstructural development with aging. The microstructural aging response for C u - 1 0 N i - 1 2 S n is shown in Fig. 10. Again the microstructure at maximum hardness manifests a partial discontinuous transformation product (Fig. 10(b)). The overaged condition (Fig. 10(c)) shows the resultant discontinuous transformation with the formation of duplex cells composed of the a and ~ phases.
Overaging brought about by cellular discontinuous precipitation produced similar dualphase structures in all alloys and both nickel and tin contents have little effect on differentiating mechanical performance. 4. Discussion
4.1. Microstructural comparison between conventionally processed ingots and rapidly solidified ribbons A certain inhomogeneity of both ribbons and ingots has been observed owing to a practically
229
460
I I 10 Ni-6$n 10 NMISn 10 Ni-10$n 10 Ni-12Sn 15 Ni.8Sn
I
.
.
.
.
.
I
I
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.
~
.,.,
.-" S " , %1 % ,,./ ..-/'j,,,,.-",.~,,, '~,~ ..- /'~ "...~-,,. ,~
-
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.=
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.
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/
,-" X ' /
340 ,# /
/
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0.4
.o
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0.8
I
1.2
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2.0
2.4
2.8
32
LOG (aglng lime), e i s m s
Fig. 9. Hardness as a function of aging time at 400 °C for rapidly solidified C u - N i - S n ribbons.
unavoidable V precipitation no matter what the casting process. However, this precipitation is on a much finer scale in ribbons than in ingots. For the rapidly solidified material, as can be seen in the micrographs and the microprobe trace data, segregation spacing was of the order of 5/~m. By comparison, it was reported that a continuous-cast Cu-10Ni-2Sn plate (15 mm thick) had a segregation spacing (secondary dendrite arm spacing) of the order of 20 /~m where tin varied from 0.6 to 3.6 wt.% across the dendritic structure [7]. The pseudobinary equilibrium phase diagram (for Cu-9Ni-xSn) indicates that for these alloys a melt must be cooled from liquidus to solidus by at least 160 °C to obtain a single phase, a [9]. This undercooling was ostensibly obtained at the substrate surface of the ribbons. However, the rate of heat extraction diminishes with distance from the substrate surface and the required cooling was not obtained in the ribbon towards the free surface. A segregated microstructure with secondary phase particles at cell boundaries then resulted in these regions. The homogenization heat treatment served to decrease segregation in dendritic structures and at grain boundaries in the ingots. However, even though these ingots were relatively small (1.2 cm diameter by 8 cm long cylinder) compared to commercially available products, V precipitation was unavoidable despite a water quench after the solution heat treatment, since a and Y are stable at lower temperatures and the cooling rate was
Fig. 10. Microstructural response to aging at 400°C in Cu-10Ni-12Sn rapidly solidified ribbons: (a) 0 min, as melt spun; (b) 45 min, maximum hardness; (c) 24 h, overaged.
230
not fast enough to keep tin in supersaturated solution. Although the heat extraction in the melt-spinning process utilized in this study was not rapid enough to completely avoid precipitation, the rapidly solidified materials had significantly reduced segregation spacing and greatly diminished chemical inhomogeneity in comparison to the as-cast alloys. Both of these effects are beneficial in curtailing the requirements of subsequent homogenization treatments. An evident case where rapid solidification retains homogeneity when conventional casting processes cannot is the Cu-10Ni-12Sn composition. This alloy in ribbon form has a uniform microstructure no different from such alloys with lower tin concentration, whereas in conventionally prepared materials a strongly heterogeneous microstructure is apparent.
4.2. Comparison of age hardening between conventionallyprocessed ingots and rapidly solidified ribbons The general age-hardening behaviours for both ingots and ribbons are similar. Equivalent maximum hardness values were obtained for the same alloy composition in both ingot and ribbon forms. However, the heat treatment duration to maximum hardness for ribbons was slightly longer than for ingots. This was due to the fact that the cooling rate of the ingot by water quench (from the homogenizing temperature of 825 °C) was less than that of the ribbon by melt spinning. Thus spinodal decomposition had already advanced in the ingot ahead of the ribbon as confirmed by hardness values measured at zero aging time. Rapid sofidification leads to mechanical characteristics (as implied by microhardness testing) equivalent to homogenized ingots in addition to greater homogeneity of both chemical and mechanical properties within the material.
5. Conclusions
The effect of alloy composition and solidification processing on mechanical properties has been studied for a range of Cu-Ni-Sn alloys. Clearly chemical homogeneity was the key factor in all aspects of material performance. Melt spinning was found to be an excellent means for reducing tin segregation which could otherwise
be eliminated only by complex processing methods. Age hardening by spinodal composition was studied in both rapidly solidified and conventionally processed materials, and mechanical properties were found to be comparable in both materials as inferred from microhardness tests. Where rapidly solidified alloys manifest superiority over their conventional counterparts is in the uniformity of metallurgical properties throughout the material on a microscopic scale. This has been confirmed by Knoop microhardness measurements, microprobe and X-ray diffraction analyses and metallography. The results obtained from mechanical and microstructural investigations lead to the following conclusions. (1) Segregation and segregation spacing in Cu-Ni-Sn alloys were significantly reduced by rapid solidification. Microsegregation persisted in the form of 7 precipitates at grain boundaries. (2) Elemental tin segregation in the melt-spun material was independent of the alloy content of tin ranging from 6 to 12 wt.%. (3) Rapid solidification by melt spinning allowed the solid solubility of tin at a higher concentration otherwise unattainable by conventional alloy preparation techniques. (4) A new homogeneous Cu-10Ni-12Sn rapidly solidified alloy, unavailable by conventional means, potentially offers mechanical strength greater than currently available alloys with nickel content less than or equal to 10 wt.%. (5) The age-hardening response of rapidly solidified alloys was similar to that of their conventional counterpart. (6) Increasing the tin content served to augment the maximum hardness of the alloy by spinodal decomposition. This maximum hardness was similar irrespective of the preparation method (except for the C u - 10Ni- 12Sn alloy). (7) Hardness increases with nickel content by alloy hardening of the matrix. (8) As indicated by the variation in microhardness values, the rapidly solidified alloys had mechanical properties more uniform throughout the material. Rapid solidification technology improves the material performance of Cu-Ni-Sn alloys, including greater resistance to corrosion [11]. It remains for further research to investigate the integrity of working pieces consolidated from such rapidly solidified materials.
231
References 1 J. M. DeVanssay, Matdriaux et Technique, 75 (1-2) (1987) 37. 2 C.R. Scorey, S. Chin, M. J. White and R. J. Livak, J. Met., 36 (11)(1984)52. 3 L. Deyong, Rapidly solidified Cu-Ni-Sn alloys: their metallurgical and electrochemical properties, M.Sc. Thesis, Laval University, Quebec, 1989. 4 R. Angers and C. Gelinas, Fabrication of shaped charge liners from rapidly solidified powders. Rep. DSS File lOSD-97701-6-O768-A, January 1988 (Defense Research Establishment, Valcartier, Quebec). 5 M. Miki and Y. Ogino, Trans. Jpn. Inst. Met., 25 (9)
(1984) 593. 6 B. G. Lefevre, A. T. D'Annessa and D. Kalish, Metall. Trans. A, 9(1978) 577. 7 L. E. Collins, J. R. Barry and J. Masounave, Reduced segregation in rapidly solidified Cu-Ni-Sn alloys, Rep. PRML 87-43(OP-J), June 1987 (CANMET, Physical Metallurgy Research Laboratories, Ottawa, Ontario). 8 Y. Watanabe, Y, Murakami and S. Kachi, J. Jpn. Inst. Met., 45(1981) 551. 9 J.T. Plewes, Metall. Trans. A, 6 (1975), 537. 10 P. Kratochvil, J. Mencl and J. Pesicka, Freib. Forschungsh. B, Metall. Werkstofftech., 256(1986) 87. 11 L. Deyong, M. Elboujdai'ni, R. Tremblay and E. Ghali, J. Appl. Electrochemistry, in the press.