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On spinodal decomposition in the Co–W system
Scripta Materialia 54 (2006) 595–598 www.actamat-journals.com
On spinodal decomposition in the Co–W system ¨ stberg Gustaf O a
a,*
, Bo Jansson b, Hans-Olof Andre´n
a
Department of Applied Physics, Chalmers University of Technology, SE-41296 Go¨teborg, Sweden b Seco Tools AB, SE-73782 Fagersta, Sweden Received 23 June 2005; received in revised form 13 October 2005; accepted 21 October 2005 Available online 18 November 2005
Abstract Spinodal decomposition in a Co-rich binder phase of a Ti(C, N)–WC–Co cermet has been studied. Thermodynamic calculations of the Co–W system predict a miscibility gap between a paramagnetic and a ferromagnetic face-centred cubic-phase at the Curie temperature. For the first time these predictions were confirmed with atom probe measurements of the binder phase. 2005 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Cermets; CALPHAD; Atom-probe field-ion microscopy; Binder phase; Spinodal decomposition
1. Introduction Cermets are composite materials produced by a powder metallurgical compaction and liquid phase sintering process. Their microstructure consists of a hard carbonitride skeleton embedded in a tough binder metal, which is often Co based. During sintering, carbide and carbonitride powders are dissolved in the binder phase and when the binder is saturated the carbonitride phase will coarsen, small grains dissolve and reprecipitation occurs on large undissolved grains. In this way grains with a core/rim structure form where the core is a remnant of undissolved powder and the rim consists of reprecipitated material. The rim forms in equilibrium with the binder phase, the composition of which depends on the solubility of the different elements at different temperatures. For cermets containing W, the carbon activity during sintering determines the solid solution of W in the binder phase; for a higher carbon content the W content of the binder phase will be lower. During cool-down after sintering, the solubility decreases and reprecipitation will con-
*
Corresponding author. Tel.: +46 31772 3325; fax: +46 31772 3224. ¨ stberg). E-mail address: [email protected] (G. O
tinue, although the diffusion paths will be gradually shorter and the composition is said to be frozen in at temperatures around 1000 C [1]. Cermets are often used as tool materials in metal cutting applications due to their high hardness and wear resistance. The W content of the binder phase plays an important role as a solid solution hardener of the material and formation of a W rich phase is therefore of high interest as a hardening mechanism for cermets. A similar process has been seen in cermet materials with Ni binder phase, although in this case small particles of the intermetallic phase Ni3Ti nucleated in the binder [2]. Recently, magnetically induced phase separations in the Co–W system have been evaluated with thermodynamic calculations and have also been experimentally confirmed, both in bulk materials [3] and in sputtered films [4]. The phase separations were directly observed by electron probe microanalysis and energy dispersive X-ray analysis, respectively. In this work, thermodynamic calculations of the Co–W system have been performed to study the miscibility gap occurring at the Curie temperature. The calculations have also been compared with atom-probe compositional measurements of a Co binder phase of a cermet in order to find indications of spinodal decomposition.
1359-6462/$ - see front matter 2005 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.scriptamat.2005.10.040
596
¨ stberg et al. / Scripta Materialia 54 (2006) 595–598 G. O
2. Material and experimental methods The material studied was a Ti(C, N)–WC–Co cermet made from two 1.5 lm Ti(C, N) powders, 19 wt.% of a 0.8 lm WC powder and 14.4 wt.% of a 1.3–1.7 lm Co powder, which were mixed by milling and then sintered at 1480 C. In the final material the volume fraction of binder phase was 10.2% and the average grain size 2 lm. To study the microstructure of the material a Jeol 2000FX transmission electron microscope (TEM) was used. Specimens were produced by cutting B3 mm discs which were dimple ground with diamond paste and finally thinned to electron transparency with ion polishing using a Gatan precision ion polishing station. Needle-shaped specimens for atom-probe analysis were prepared by first cutting rods with a 250 · 250 lm2 crosssection from the material. The rods were then thinned in one end by dimple grinding and finally sharpened by focused ion beam milling with a procedure described elsewhere [5]. Atom-probe analysis was performed with an instrument described in detail earlier [6]. During analysis the specimen temperature was held at 30 K and the residual gas pressure in the analysis chamber was kept below 5 · 108 Pa. The evaporation pulse height was 20% of the positive direct current voltage applied to the specimen. 3. Results and discussion The microstructure of the studied material is shown in the TEM micrograph in Fig. 1. A number of planar faults can be seen as straight lines in the Co binder phase but there are no signs of precipitation. Since the solubility of Ti in the Co binder phase is very low, the behaviour of the binder can be understood by con-
Fig. 1. TEM micrograph of the microstructure of the studied material. Areas of fcc binder phase have been marked with B. No signs of precipitation in the binder phase can be seen but planar faults can be seen as straight lines. In addition, some radiation damage introduced during the specimen preparation can be seen as dark spots.
sidering the binary system Co–W. The Co–W phase diagram shown in Fig. 2 was calculated utilising the thermodynamic CALPHAD (computer coupling of phase diagrams and thermochemistry) assessment due to Guillermet [7]. Addition of W to a Co-rich face-centred cubic (fcc) phase results in a drastic decrease in the Curie temperature. Alloy systems with a rapid decrease in Curie temperature often have miscibility gaps with an unusual form. This type of miscibility gap, discussed by Nishizikawa et al. [8], has the shape of a sharp horn protruding along the Curie temperature line. In the Co–W system there is a miscibility gap in the fcc phase, a Nishizikawa horn, with a top temperature of approximately 1240 K (967 C). The miscibility gap is not stable below the invariant three-phase equilibrium fcc + fcc + Co3W at 1171 K (898 C) but the fcc phase will separate into two phases with different composition and magnetic character, a Co-rich ferromagnetic and a more W-rich paramagnetic phase. On cooling of a homogeneous fcc phase, with a composition of e.g. xW = 0.1, nucleation and growth of the Co3W phase can be very sluggish. However, a spinodal decomposition that only requires atomic movement over very short distances may occur. It is therefore interesting to consider the metastable miscibility gap at lower temperatures. Interesting parts of the Co–W phase diagram where also the metastable miscibility gap has been included can be found in Fig. 3. An fcc phase that on cooling falls within the spinodal will decompose without the need of nucleation. Note that the lower limit of the spinodal is about xW = 0.04 and does not change much with temperature. The total W content of the binder phase was 6.2 at.% and the carbon content was 0.53 at.% as measured by atom-probe analysis. Since the C content is very low and does not affect the Curie temperature, according to the thermodynamic model used, the carbon is not considered to have any effect on the spinodal decomposition.
Fig. 2. Phase diagram of the Co–W system.
¨ stberg et al. / Scripta Materialia 54 (2006) 595–598 G. O
597
dent on pffiffiffiffiffithe diffusion constant, D, and time, t, as r ¼ 2:4 Dt where D is dependent on temperature which in turn is dependent on time [9]. The diffusion constant used for W in Co was: DW ¼ 7 105 e
Fig. 3. Magnification of the region where the miscibility gap occurs. The extension of the miscibility gap at lower temperatures is in full lines and the variation of the Curie temperature with W content is drawn with a dotted line. The spinodal lines are given by the dashed lines.
Fig. 4. Composition profile of the binder phase of a Ti(C, N)–WC–Co cermet as measured by atom-probe analysis. The error bars correspond to one standard deviation of the counting statistics and the curve is balanced by Co.
It can be seen in the compositional profile shown in Fig. 4 that the W content varies between roughly 4 and 10 at.% and the wavelength of the fluctuations is between 5 and 10 nm. The compositional variations between Wdepleted and more W-rich regions agree well with the compositions around 700 C (1000 K) as can be seen in Fig. 3. It is possible to calculate the expected amplitude and wave length of the periodic variations if also the coherent elastic energy is considered. A complete such calculation has not been performed within this study but an estimation of the plausibility of the results was made by making some assumptions. The temperature during cooling after sintering was supposed to vary exponentially with time as: 4
T ¼ T 0 e1:2910
t
where the constant was calculated from a cooling time t = 4 h between the temperatures T0 = 1275 C and T = 200 C [1]. The range of diffusion, r, is roughly depen-
282;100 RT
where R is the gas constant and T is the temperature in K [10]. To estimate the range of diffusion at a certain temperature the temperature was divided into a number of intervals. Within such an interval D was considered to be constant during the time, Dt, it took to cool the material from the upper to the lower temperature. By calculating D for the temperature in the middle of the interval the range of diffusion could then be calculated for this temperature and the time Dt. With the composition measured, the spinodal in Fig. 2 should be intersected somewhere between 1100 and 1150 K (850 C) during cooling and the distribution of W corresponded to the compositions around 700 C. Hence, the intervals used were 875–825, 725–675 and 625–575 C resulting in diffusion lengths of 120, 15 and 2 nm, respectively. The diffusion path is, thus, rather long when the spinodal curve is reached at 850 C and a separation to the equilibrium wavelength should occur quite easily, although the driving force for phase separation is very low at this point. Moreover, the range of diffusion at 700 C is still longer than the wavelength measured for the corresponding W distribution and diffusion lengths of several nm are possible down to about 600 C. Despite this, the W distribution corresponding to 700 C appears to be frozen in. An explanation of this may be that even though the driving force for phase separation increases with lower temperatures, the elastic energy between W-rich and Wdepleted regions increases as the separation is getting more pronounced. The importance of the elastic energy as a phase separation suppressor has also been pointed out earlier by Oikawa et al. [4]. In addition, the wavelength of the decomposition gets shorter and the amplitude is increased with lower temperatures [9]. Hence, W has to be continuously redistributed between the W-rich and the W-depleted regions but the diffusion paths are getting gradually shorter so the redistribution from longer to shorter wavelengths will be increasingly difficult. The short wavelength of the compositional variation should present a considerable resistance to dislocation movement and therefore make a large contribution to the strength of the binder, probably much larger than the solid solution hardening. In fact, a similar cermet with a lower carbon content and thus a higher W solid solution of the binder exhibited a lower hardness, in spite of the expected higher solution hardening [11]. The reason for this is most likely a lower hardening by spinodal composition due to the much shorter time available for decomposition at higher W contents.
598
¨ stberg et al. / Scripta Materialia 54 (2006) 595–598 G. O
4. Conclusions • The W–Co system is predicted to exhibit spinodal decomposition into a paramagnetic and a ferromagnetic fcc phase. • For the first time, the spinodal decomposition was confirmed by atom-probe analysis in a cermet with a Co based binder phase containing 6.2 at.% W. • The W distribution corresponded to the compositions of the separated phases around 700 C. • The wavelength of the decomposition was measured to between 5 and 10 nm. • Spinodal decomposition of the binder gives an important contribution to the hardness of the material.
Acknowledgement AB Sandvik Coromant is gratefully thanked for financial support and for providing materials.
References [1] Hellsing M. Mater Sci Technol 1988;4:824. [2] Zackrisson J, Larsson A, Andre´n H-O. Micron 2001;32:707. [3] Sato J, Katsunari O, Kainuma R, Ishidi K. Mater Trans 2005;46:1199. [4] Oikawa K, Qin GW, Sato M, Okamoto S, Kitakami O, Shimada Y, et al. Appl Phys Lett 2004;85:2559. ¨ stberg G. Submitted for publication. [5] O [6] Andre´n H-O, Norde´n H. Scand J Metall 1979;8:147. [7] Guillermet AF. Metall Trans A 1989;20A:935. [8] Nishizikawa T, Hasabe M, Ko M. Acta Metall 1979;27:817. [9] Porter DA, Easterling KE. Phase transformations in metals and alloys. Second ed. Cheltenham: Nelson Thornes; 2001. [10] Haglund S. Sintering of cemented carbides—experiments and modelling. PhD Thesis. Department of Materials Science and Engineering: Stockholm; 1998. ¨ stberg G, Buss K, Christensen M, Norgren S, Andre´n H-O, Mari D, [11] O Reineck I. Int J Refract Met Hard Mater 2006;24:145.
On spinodal decomposition in the Co–W system ¨ stberg Gustaf O a
a,*
, Bo Jansson b, Hans-Olof Andre´n
a
Department of Applied Physics, Chalmers University of Technology, SE-41296 Go¨teborg, Sweden b Seco Tools AB, SE-73782 Fagersta, Sweden Received 23 June 2005; received in revised form 13 October 2005; accepted 21 October 2005 Available online 18 November 2005
Abstract Spinodal decomposition in a Co-rich binder phase of a Ti(C, N)–WC–Co cermet has been studied. Thermodynamic calculations of the Co–W system predict a miscibility gap between a paramagnetic and a ferromagnetic face-centred cubic-phase at the Curie temperature. For the first time these predictions were confirmed with atom probe measurements of the binder phase. 2005 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Cermets; CALPHAD; Atom-probe field-ion microscopy; Binder phase; Spinodal decomposition
1. Introduction Cermets are composite materials produced by a powder metallurgical compaction and liquid phase sintering process. Their microstructure consists of a hard carbonitride skeleton embedded in a tough binder metal, which is often Co based. During sintering, carbide and carbonitride powders are dissolved in the binder phase and when the binder is saturated the carbonitride phase will coarsen, small grains dissolve and reprecipitation occurs on large undissolved grains. In this way grains with a core/rim structure form where the core is a remnant of undissolved powder and the rim consists of reprecipitated material. The rim forms in equilibrium with the binder phase, the composition of which depends on the solubility of the different elements at different temperatures. For cermets containing W, the carbon activity during sintering determines the solid solution of W in the binder phase; for a higher carbon content the W content of the binder phase will be lower. During cool-down after sintering, the solubility decreases and reprecipitation will con-
*
Corresponding author. Tel.: +46 31772 3325; fax: +46 31772 3224. ¨ stberg). E-mail address: [email protected] (G. O
tinue, although the diffusion paths will be gradually shorter and the composition is said to be frozen in at temperatures around 1000 C [1]. Cermets are often used as tool materials in metal cutting applications due to their high hardness and wear resistance. The W content of the binder phase plays an important role as a solid solution hardener of the material and formation of a W rich phase is therefore of high interest as a hardening mechanism for cermets. A similar process has been seen in cermet materials with Ni binder phase, although in this case small particles of the intermetallic phase Ni3Ti nucleated in the binder [2]. Recently, magnetically induced phase separations in the Co–W system have been evaluated with thermodynamic calculations and have also been experimentally confirmed, both in bulk materials [3] and in sputtered films [4]. The phase separations were directly observed by electron probe microanalysis and energy dispersive X-ray analysis, respectively. In this work, thermodynamic calculations of the Co–W system have been performed to study the miscibility gap occurring at the Curie temperature. The calculations have also been compared with atom-probe compositional measurements of a Co binder phase of a cermet in order to find indications of spinodal decomposition.
1359-6462/$ - see front matter 2005 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.scriptamat.2005.10.040
596
¨ stberg et al. / Scripta Materialia 54 (2006) 595–598 G. O
2. Material and experimental methods The material studied was a Ti(C, N)–WC–Co cermet made from two 1.5 lm Ti(C, N) powders, 19 wt.% of a 0.8 lm WC powder and 14.4 wt.% of a 1.3–1.7 lm Co powder, which were mixed by milling and then sintered at 1480 C. In the final material the volume fraction of binder phase was 10.2% and the average grain size 2 lm. To study the microstructure of the material a Jeol 2000FX transmission electron microscope (TEM) was used. Specimens were produced by cutting B3 mm discs which were dimple ground with diamond paste and finally thinned to electron transparency with ion polishing using a Gatan precision ion polishing station. Needle-shaped specimens for atom-probe analysis were prepared by first cutting rods with a 250 · 250 lm2 crosssection from the material. The rods were then thinned in one end by dimple grinding and finally sharpened by focused ion beam milling with a procedure described elsewhere [5]. Atom-probe analysis was performed with an instrument described in detail earlier [6]. During analysis the specimen temperature was held at 30 K and the residual gas pressure in the analysis chamber was kept below 5 · 108 Pa. The evaporation pulse height was 20% of the positive direct current voltage applied to the specimen. 3. Results and discussion The microstructure of the studied material is shown in the TEM micrograph in Fig. 1. A number of planar faults can be seen as straight lines in the Co binder phase but there are no signs of precipitation. Since the solubility of Ti in the Co binder phase is very low, the behaviour of the binder can be understood by con-
Fig. 1. TEM micrograph of the microstructure of the studied material. Areas of fcc binder phase have been marked with B. No signs of precipitation in the binder phase can be seen but planar faults can be seen as straight lines. In addition, some radiation damage introduced during the specimen preparation can be seen as dark spots.
sidering the binary system Co–W. The Co–W phase diagram shown in Fig. 2 was calculated utilising the thermodynamic CALPHAD (computer coupling of phase diagrams and thermochemistry) assessment due to Guillermet [7]. Addition of W to a Co-rich face-centred cubic (fcc) phase results in a drastic decrease in the Curie temperature. Alloy systems with a rapid decrease in Curie temperature often have miscibility gaps with an unusual form. This type of miscibility gap, discussed by Nishizikawa et al. [8], has the shape of a sharp horn protruding along the Curie temperature line. In the Co–W system there is a miscibility gap in the fcc phase, a Nishizikawa horn, with a top temperature of approximately 1240 K (967 C). The miscibility gap is not stable below the invariant three-phase equilibrium fcc + fcc + Co3W at 1171 K (898 C) but the fcc phase will separate into two phases with different composition and magnetic character, a Co-rich ferromagnetic and a more W-rich paramagnetic phase. On cooling of a homogeneous fcc phase, with a composition of e.g. xW = 0.1, nucleation and growth of the Co3W phase can be very sluggish. However, a spinodal decomposition that only requires atomic movement over very short distances may occur. It is therefore interesting to consider the metastable miscibility gap at lower temperatures. Interesting parts of the Co–W phase diagram where also the metastable miscibility gap has been included can be found in Fig. 3. An fcc phase that on cooling falls within the spinodal will decompose without the need of nucleation. Note that the lower limit of the spinodal is about xW = 0.04 and does not change much with temperature. The total W content of the binder phase was 6.2 at.% and the carbon content was 0.53 at.% as measured by atom-probe analysis. Since the C content is very low and does not affect the Curie temperature, according to the thermodynamic model used, the carbon is not considered to have any effect on the spinodal decomposition.
Fig. 2. Phase diagram of the Co–W system.
¨ stberg et al. / Scripta Materialia 54 (2006) 595–598 G. O
597
dent on pffiffiffiffiffithe diffusion constant, D, and time, t, as r ¼ 2:4 Dt where D is dependent on temperature which in turn is dependent on time [9]. The diffusion constant used for W in Co was: DW ¼ 7 105 e
Fig. 3. Magnification of the region where the miscibility gap occurs. The extension of the miscibility gap at lower temperatures is in full lines and the variation of the Curie temperature with W content is drawn with a dotted line. The spinodal lines are given by the dashed lines.
Fig. 4. Composition profile of the binder phase of a Ti(C, N)–WC–Co cermet as measured by atom-probe analysis. The error bars correspond to one standard deviation of the counting statistics and the curve is balanced by Co.
It can be seen in the compositional profile shown in Fig. 4 that the W content varies between roughly 4 and 10 at.% and the wavelength of the fluctuations is between 5 and 10 nm. The compositional variations between Wdepleted and more W-rich regions agree well with the compositions around 700 C (1000 K) as can be seen in Fig. 3. It is possible to calculate the expected amplitude and wave length of the periodic variations if also the coherent elastic energy is considered. A complete such calculation has not been performed within this study but an estimation of the plausibility of the results was made by making some assumptions. The temperature during cooling after sintering was supposed to vary exponentially with time as: 4
T ¼ T 0 e1:2910
t
where the constant was calculated from a cooling time t = 4 h between the temperatures T0 = 1275 C and T = 200 C [1]. The range of diffusion, r, is roughly depen-
282;100 RT
where R is the gas constant and T is the temperature in K [10]. To estimate the range of diffusion at a certain temperature the temperature was divided into a number of intervals. Within such an interval D was considered to be constant during the time, Dt, it took to cool the material from the upper to the lower temperature. By calculating D for the temperature in the middle of the interval the range of diffusion could then be calculated for this temperature and the time Dt. With the composition measured, the spinodal in Fig. 2 should be intersected somewhere between 1100 and 1150 K (850 C) during cooling and the distribution of W corresponded to the compositions around 700 C. Hence, the intervals used were 875–825, 725–675 and 625–575 C resulting in diffusion lengths of 120, 15 and 2 nm, respectively. The diffusion path is, thus, rather long when the spinodal curve is reached at 850 C and a separation to the equilibrium wavelength should occur quite easily, although the driving force for phase separation is very low at this point. Moreover, the range of diffusion at 700 C is still longer than the wavelength measured for the corresponding W distribution and diffusion lengths of several nm are possible down to about 600 C. Despite this, the W distribution corresponding to 700 C appears to be frozen in. An explanation of this may be that even though the driving force for phase separation increases with lower temperatures, the elastic energy between W-rich and Wdepleted regions increases as the separation is getting more pronounced. The importance of the elastic energy as a phase separation suppressor has also been pointed out earlier by Oikawa et al. [4]. In addition, the wavelength of the decomposition gets shorter and the amplitude is increased with lower temperatures [9]. Hence, W has to be continuously redistributed between the W-rich and the W-depleted regions but the diffusion paths are getting gradually shorter so the redistribution from longer to shorter wavelengths will be increasingly difficult. The short wavelength of the compositional variation should present a considerable resistance to dislocation movement and therefore make a large contribution to the strength of the binder, probably much larger than the solid solution hardening. In fact, a similar cermet with a lower carbon content and thus a higher W solid solution of the binder exhibited a lower hardness, in spite of the expected higher solution hardening [11]. The reason for this is most likely a lower hardening by spinodal composition due to the much shorter time available for decomposition at higher W contents.
598
¨ stberg et al. / Scripta Materialia 54 (2006) 595–598 G. O
4. Conclusions • The W–Co system is predicted to exhibit spinodal decomposition into a paramagnetic and a ferromagnetic fcc phase. • For the first time, the spinodal decomposition was confirmed by atom-probe analysis in a cermet with a Co based binder phase containing 6.2 at.% W. • The W distribution corresponded to the compositions of the separated phases around 700 C. • The wavelength of the decomposition was measured to between 5 and 10 nm. • Spinodal decomposition of the binder gives an important contribution to the hardness of the material.
Acknowledgement AB Sandvik Coromant is gratefully thanked for financial support and for providing materials.
References [1] Hellsing M. Mater Sci Technol 1988;4:824. [2] Zackrisson J, Larsson A, Andre´n H-O. Micron 2001;32:707. [3] Sato J, Katsunari O, Kainuma R, Ishidi K. Mater Trans 2005;46:1199. [4] Oikawa K, Qin GW, Sato M, Okamoto S, Kitakami O, Shimada Y, et al. Appl Phys Lett 2004;85:2559. ¨ stberg G. Submitted for publication. [5] O [6] Andre´n H-O, Norde´n H. Scand J Metall 1979;8:147. [7] Guillermet AF. Metall Trans A 1989;20A:935. [8] Nishizikawa T, Hasabe M, Ko M. Acta Metall 1979;27:817. [9] Porter DA, Easterling KE. Phase transformations in metals and alloys. Second ed. Cheltenham: Nelson Thornes; 2001. [10] Haglund S. Sintering of cemented carbides—experiments and modelling. PhD Thesis. Department of Materials Science and Engineering: Stockholm; 1998. ¨ stberg G, Buss K, Christensen M, Norgren S, Andre´n H-O, Mari D, [11] O Reineck I. Int J Refract Met Hard Mater 2006;24:145.