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On the sintering of molybdenum with two liquid phases
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On the Sintering of Molybdenum with two Liquid Phases Andrea Huber , Lorenz S. Sigl PII: DOI: Reference:
S2589-1529(20)30017-X https://doi.org/10.1016/j.mtla.2020.100600 MTLA 100600
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Materialia
Received date: Accepted date:
17 December 2019 18 January 2020
Please cite this article as: Andrea Huber , Lorenz S. Sigl , On the Sintering of Molybdenum with two Liquid Phases, Materialia (2020), doi: https://doi.org/10.1016/j.mtla.2020.100600
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On the Sintering of Molybdenum with two Liquid Phases Andrea Huber1 and Lorenz S. Sigl2 1 2
Plansee SE, Reutte, Austria
Technical University Dresden, Dresden, Germany
Abstract The liquid phase sintering (LPS) of Molybdenum-base alloys containing 80 vol.% Mo and 20 vol.% Cu and Ni as sintering aids was investigated. Sintering experiments and subsequent microstructural analysis together with equilibrium and non-equilibrium thermodynamic calculations elucidate the densification behavior and the microstructure of Mo-Cu-Ni alloys. LPS is dominated by the liquid miscibility gap in the Mo-Cu-Ni system. Surrounding the Mograins during sintering, a Cu-rich and a Ni-rich liquid phase with moderate and high solubility for Mo respectively, generate fully dense and ductile composites via the solution and reprecipitation mechanism. The rates of densification and microstructural coarsening increase significantly with the amount of Ni-rich liquid, i.e. with the solubility of the binder for Mo. Upon cooling, the two immiscible liquids solidify into a dual-phase binder comprising a Cuand a Ni-rich fcc (Cu,Ni) solid solution respectively. Precipitation of the brittle MoNi phase (), which was reported in previous studies, can be suppressed provided that (i) sintering happens at temperatures where , which develops during heat-up of the elemental Mo-CuNi powder blend, dissolves in the liquid phases and (ii) cooling from sintering is sufficiently fast to suppress its reprecipitation.
Key Words Molybdenum alloy; liquid phase sintering; liquid immiscibility; densification; grain growth 1. Introduction Liquid phase sintering (LPS) is a well-established method to consolidate powder compacts into fully dense multi-phase components. The process is used to manufacture metals including tungsten heavy alloys [1] and some P/M steels [2, 3]. It is also applied to densify various ceramics and ceramic-metal composites, e.g. certain aluminas [4], the silicon nitride materials family [5], liquid phase sintered silicon carbide [6], and not least cemented carbides [7]. Sintering studies on various materials revealed that, despite the heterogeneities in the solid constituents, the melt during LPS is generally homogeneous, i.e. single phase [8]. Despite striking similarities between W and Mo, the assistance of a liquid phase analogous to the manufacturing of tungsten heavy alloys (WHAs) [1, 9-11] has, apart from “nonsintering” densification such as the infiltration of pre-sintered, porous Mo-skeletons with Cu [12,13], not been employed to densify Mo-base materials. This knowledge gap has triggered a thesis, which intended to assess the potential of various metallic binders for the densification of Mo by LPS [14]. Combining thermodynamic calculations and sintering experiments, page - 1 -
that study identified Cu and Ni as most promising alloy additions, because the Mo-Cu-Ni system can satisfy the essential criteria for liquid phase sintering: (i) wettability of Mo by the molten binder, (ii) sufficient solubility of Mo in the liquid binder and (iii) chemical compatibility between Mo and the binder to prevent the formation of brittle intermetallic phases. Only recently, Hwang and Huang investigated Mo-base alloys with a binder containing pure Ni and Ni/Cu alloys [15-17]. They demonstrated that the composites can be sintered to full density, yet found them to be brittle rather than ductile due to the formation of the intermetallic phase MoNi () in the binder [15,16]. Depending on the Ni/Cu composition, the authors identified activated or liquid phase sintering as operating densification mechanisms [16]. However, they did neither recognize the importance of the liquid miscibility gap (LMG) in the Mo-Cu-Ni system at temperatures above 1320 °C nor its impact on liquid phase sintering. It is the intention of this paper to elucidate the constitutional fundamentals for producing dense and ductile Mo-base alloys, by (i) investigating the constitution of the Mo-Cu-Ni system under equilibrium and non-equilibrium conditions, (ii) conducting sintering experiments and (iii) validating the computational results through microstructural and analytical observations. 2. Compositions and experimental procedures Commercial powders from industrial suppliers were used in this study. Their size and chemical composition are summarized in Table 1. To investigate liquid phase sintering in the MoCu-Ni system, a set of four compositions was chosen with the rationale to investigate Moalloys at a fixed binder fraction of 20 vol.%1 at ambient temperature. The choice of binder composition was guided by the isothermal section of the phase diagram at the intended sintering temperature of 1400 °C (Fig. 1). Characteristic for Mo-Cu-Ni is a miscibility gap in the liquid as shown in Fig. 1. Starting with the Cu-rich alloy 2N2, which is positioned in the two-phase field comprising solid Mo (bcc) and a single melt, LCu, an increase of the Ni-content in three equidistant steps of 4 wt.% Ni (alloys 6N, 10N, 14N) switches the phase configuration into the three-phase field which, besides Mo, contains two liquid phases, i.e. a Cu- and a Ni-rich melt, LCu and LNi, respectively (Figs. 1 and 2). The positions of the alloys 2N: 6N: 10N: 14N:
82 wt.% Mo – 16 wt.% Cu – 2 wt.% Ni 82 wt.% Mo – 12 wt.% Cu – 6 wt.% Ni 82 wt.% Mo – 8 wt.% Cu – 10 wt.% Ni 82 wt.% Mo – 4 wt.% Cu – 14 wt.% Ni
at the sintering temperature of 1400 °C are marked by the solid circles in Figs. 1-3. The alloys were prepared by dry mixing a total of 300 grams of elemental Mo/Cu/Ni powder blends with 100 grams zirconia balls for 1 hour in a Turbula mixer. After drying and subse-
1 2
20 vol.% binder corresponds approximately to 18 wt.% binder and 82 wt.% Mo the numeral denotes the Ni-content of the Mo-Cu-Ni alloy in wt.%
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quent compaction of cylindrical tablets (Ø 25 mm, height 5 mm) at 180 MPa in an axial press (Komage 100CNC), a green density of 65% of theoretical density (T.D.) was reached. Liquid phase sintering was carried out in a tube furnace (Xerion Xtube 1650) under flowing hydrogen at a flow rate of 0.5 l/min and holding times between 0h and 4h at temperatures of 1350 °C, 1400 °C and 1450 °C, applying a heating and a cooling rate of 10 K/min respectively. From WHAs it is well known, that large pores can develop during sintering under dry hydrogen, a phenomenon which has been attributed to the formation of water vapor [18]. Likewise, pore formation in Mo-Cu-Ni was successfully blocked by setting the hydrogen dew point to -55°C below 900 °C and to 20 °C above that temperature. The as-sintered densities were measured by Archimedes’ method. Metallographic samples were prepared by grinding on SiC papers up to P2400, followed by diamond polishing from 9 µm down to 1 µm and a final polishing step with 0.04 µm colloidal silica. Both optical (Leica Microsystems DMI 5000) and scanning electron microscopes (Carl Zeiss Microscopy Ultra 55 FEG) were used to investigate the microstructure on unetched specimens, as the different phases are easier visible on as-polished surfaces. The grain size, G, of the Mo particles was measured by the linear intercept method and their sphericity, S, was evaluated from the relationship S = 4 A/P2 where A denotes the area and P the perimeter of individual Mo crystallites. Quantitative energy dispersive X-ray analysis3 (acceleration voltage 20 kV, probe current 0.9 nA, standardless ZAF correction) was used to assess the chemical composition of the various phases. They were further analyzed by a combination of EDX and electron back scattered diffraction (EBSD3). Finally, differential scanning calorimetry (DSC) measurements (Netzsch STA 409 CD) were conducted in an Argon atmosphere by heating the samples up to 1400 °C with subsequent cooling at a cooling and heating rate of 10K/min respectively. 3. Thermodynamic calculations Thermodynamic calculations were performed with the commercial THERMOCALC software using the Ni-database TCNI8 [19], where the system Mo-Cu-Ni is covered in the full composition and temperature range based on the investigations of Yan et al. [20]. The Mo-Cu-Ni system was investigated by isothermal and isopleth sections as well as by evaluating the fraction of phases as a function of temperature for selected compositions. Notably, THERMOCALC permits non-equilibrium states to be explored by (i) calculating Scheil-Gulliversolidification simulations and (ii) by implementing those results into subsequent nonequilibrium phase diagram calculations. Both tools will prove to be useful for predicting the formation and spatial arrangement of phases under technical cooling conditions. 3.1. Equilibrium phase relationships in the Mo-Cu-Ni system As the topic of this paper is liquid phase sintering, it is convenient to start with a close look on the equilibrium phase diagrams at sintering temperature, e.g. on the sections of the MoCu-Ni system at 1400 °C (Fig. 1) and at 82 wt.% Mo (Fig. 2). The isothermal 1400 °C section reveals three one-phase rooms, i.e. bcc Mo, an fcc (Cu,Ni) solid solution, and a (Cu-Ni-Mo) liquid. That liquid is single-phase close to the Cu-Ni line, yet destabilized by the addition of
3
EDAX-AMETEK Materials Analysis Division
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Mo which causes a miscibility gap containing two liquids. While alloy 2N is located in the two-phase field with the Cu-rich melt, LCu, the Ni-rich alloys 6N-14N are located in a threephase field comprising the extra liquid, LNi. Hence, in the Ni-rich alloys, sintering above 1323 °C happens in the presence of two liquids, whereas only one liquid phase is active in alloy 2N. Notably, the Cu-rich liquid exhibits moderate solubility for Mo, whereas L Ni can dissolve 25 times more Mo (Table 2 and Fig. 1). 3.1.1. The liquid miscibility gap in the binder during sintering To our best knowledge, LPS in the presence of more than one liquid phase has as yet not been described in the literature. It is not the intension of this paper to discuss the reasons for the presence of an LMG in Mo-Cu-Ni, but to focus on the impact of the two liquid phases on the sinterability of Mo and on the microstructure developing from them. The ternary LMG in Mo-Cu-Ni originates from the liquid immiscibility in the Mo-Cu system, where it appears at temperatures above 2500 °C. Due to the high temperature level, this aspect has not been paid much attention, because the LMG is irrelevant for the processing of commercial Mo-Cu materials. In Cu-rich alloy 2N the liquid immiscibility becomes evident only at temperatures above 2300 °C, whereas a 2-phase field (Mo+LCu) obtains at sintering temperature (Fig. 3a). Adding more Ni drops the critical point of the ternary LMG significantly and expands the 3-phase field (Mo + LCu + LNi,) to temperatures as low as 1323 °C (Figs. 3b-d). Notably, two liquids are present at all sintering temperatures in alloys 6N, 10N and 14 N. The two liquids differ considerably with respect to their solubility for Mo: outside the LMG, the Mo-solubility of LCu increases only weakly with the Ni-concentration (c.f. Fig. 1), and dissolves just 0.45 at.% Mo in 2N at 1400 °C (Table 2). Within the three-phase field (Mo+ LCu+ LNi), where the compositions of LCu and LNi are fixed, the Mo-solubility at 1400 °C is still only 1.1 at.% for LCu, but 28.3 at.% in LNi (Table 2). Finally, it should be recalled that the volume fractions of LNi and LCu in- and decrease with the Ni-concentration as illustrated in Fig. 4. As will be shown below, the disproportionate solubility of Mo in the two liquids, LCu and LNi, affects densification via solution and reprecipitation deeply, and has a strong impact on the kinetics of grain growth. 3.1.2. Invariant reactions during equilibrium cooling from sintering temperature The isopleth sections at 82 wt.% Mo in Fig. 2 illustrate the phase relationships during cooling from sintering temperature. Under equilibrium conditions (Fig. 2a), THERMOCALC predicts three invariant reactions in the Ni-rich alloys 6N, 10N, and 14N: ↔
(U1)
↔
(U2)
↔
(U3)
where LCu and LNi are the liquid phases defined above, is an fcc (Cu,Ni) solid solution, Mo denotes the bcc Mo-phase, and (MoNi) is a brittle, intermetallic compound with orthorhombic lattice [16,21]. Cu-rich alloys with less than 4.5 wt.% Ni do not undergo the invari-
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ant reactions U1 - U3, but precipitate at elevated temperatures, e.g. for 2N below 1126°C, and also the -phase at low temperatures (Fig. 2a). Reaction U1 is calculated to take place at 1323 °C, i.e. slightly below the lowest sintering temperatures investigated here. THERMOCALC forecasts U1 to consume predominantly the Nirich liquid and to generate together with Mo precipitates. Below 1248 °C the intermetallic -phase is predicted to crystallize according to reaction U2. However, as U2 is a reaction between solid phases, mass transport is a lot slower than in reactions involving a liquid, suggesting that U2 is sluggish. Furthermore, can precipitate only in the temperature interval between U2 and U3. As this accounts for not more than 17 K, the experimental cooling rate of 10 K/min allows a time span of merely 100 seconds for the precipitation of . Finally, reaction U3 being inverse to U2 would partly reconsume the newly formed δ-phase during further cooling. These arguments suggest the precipitation of to be unlikely under practical cooling conditions. Consequently, the equilibrium prediction of -precipitation has to be checked against (i) thermodynamic non-equilibrium considerations as presented below and (ii) experimental observations as discussed in chapter 4. 3.2. Non-equilibrium solidification from sintering temperature With the Scheil-Gulliver module, THERMOCALC provides a computational tool that takes the kinetics of phase evolution during solidification into account, i.e. it computes the phases which develop during solidification and their volume fraction as a function of temperature assuming zero diffusivity in the solid, infinitely fast mass transport in the liquid and local equilibrium at the solid/liquid interface. In contradiction to the equilibrium results, ScheilGulliver predicts δ not to form during cooling, which supports the arguments discussed above. Whereas both, the liquid miscibility gap and reaction U1 remain virtually unaffected by the suppression of δ, i.e. we still have (U’1 U1)
↔ the invariant reactions U2 and U3 merge into reaction U’2 according to ↔
(U’2)
where Ni and Cu denote a Ni- and Cu-rich (Cu,Ni) fcc solid solution respectively. ScheilGulliver forecasts the following propositions, which will be compared with microstructural observations in section 4: i. Reaction U’1 generates Mo which is expected to co-precipitate with the Ni-rich (Cu,Ni) fcc solid solution Ni. This should generate a microstructural topology containing secondary molybdenum, Mosec., in intimate contact with Ni rather than the crystallization of excess Mo onto the primary Mo-crystals. THERMOCALC predicts the volume ratio of Ni /Mosec. to be 85%/15%. ii. During invariant reaction U’2 the residual Cu-rich melt, LCu, solidifies and eats up Ni to produce a Cu-rich (Cu,Ni) fcc solid solution Cu, and tertiary Mo precipitates , Motert., which should be, again, in close contact. iii. The intermetallic phase is unstable and does not form.
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Hence, a non-equilibrium phase diagram with being suppressed was calculated and is shown in Fig. 2b. It suggests that the Mo-Cu-Ni system meets the criteria for liquid phase sintering, i.e. i. LCu and LNi dissolve Mo in sufficient quantities to operate the solution-reprecipitation mechanism, despite their solubility for Mo differs by more than an order of magnitude (c.f. Table 2). ii. The brittle MoNi phase does not form under industrially relevant cooling conditions such that the criterion of chemical compatibility is met. iii. Solid Mo is wetted by the liquid phases [15]. iv. Finally, the densification temperature of Mo-Cu-Ni is markedly lower than for solid-state sintered molybdenum (> 2000 °C). 1. Results and Discussion 4.1. Densification Behavior The densification characteristics of Mo-Cu-Ni alloys are exemplified in Fig. 5. The as-sintered densities of alloys containing only LCu such as 2N (Fig. 5a) depend strongly on sintering time and temperature, i.e. they attain nearly full density only after soaking times > 2h. On the other hand, the Ni-rich compositions 6N-14N comprising the Ni-rich melt, LNi, achieve nearly full densities at all sintering temperatures and soaking times. In those alloys, nearly complete densification, as exemplified by alloy 6N in Fig. 5b, is obtained already at zero holding time. Analogous consolidation results were seen in alloys 10N and 14N [14]. Though at first surprising, this observation is readily understood as a consequence of the high Mo-solubility of LNi (Table 2). As the Ni-rich liquid becomes stable above 1323 °C (c.f. the isopleth in Fig. 2b) and increases its share with the Ni-content as shown in Fig. 4, it provides plenty of transportation capacity for Mo already during heat-up. Also, with the cooling and heating rate of 10 K/min, there is a sufficient time window of at least 5 minutes4 during which the solution-reprecipitation mechanism can redistribute Mo efficiently even before the onset of isothermal sintering. On the other hand, the as-sintered density of alloy 2N, containing just LCu with low Mosolubility, depends markedly on the soaking time ts (Fig. 5a), e.g. only 82 % T.D. is achieved at ts = 0 and 1350 °C. However, nearly full density is obtained after 4 hours of soaking, which can be understood as follows: Densification at short soaking times is to a good part due to solid state sintering during heat-up. Despite LCu becoming stable above 1126 °C, its low solubility for Mo, e.g. 0.36, 0.45 and 0.55 at.% at 1350, 1400 and 1450 °C respectively, limits the Mo-flux through the liquid. This view is supported by the moderate sphericity, S, of the Mo grains in 2N, which are less spherical than the Mo grains in 6N-14N (c.f. Figs 6a, 7a and 12b). Nearly theoretical density is nevertheless obtained after long soaking, indicating that LCu does transport Mo sufficiently over sintering time.
4
During heating from the invariant reaction U1 up to the sintering temperature, Ts, and cooling back to U1, there is a time interval tD for solution/reprecipitation given by tD = 2(Ts-1323)/R where Ts denotes the sintering temperature and R is the heating/cooling rate. With Ts = 1350 °C and R = 10 K/min, tD 5.5 min obtains.
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4.2. Microstructure and Microstructural Analysis Light microscopy (LM) images of alloys 2N-14N, which were sintered at 1400 °C for 4h, are displayed in Fig. 6 a-d. All compositions are fully dense showing virtually no porosity, in agreement with the density measurements from Fig. 5. The microstructure, which looks geometrically similar to tungsten heavy alloys, features equiaxed, rounded Mo particles (“primary” Mo) embedded into a continuous binder matrix. Corresponding results were obtained for the other sintering temperatures and soaking times, with the exception of (i) alloy 2N which does not sinter to full density at short soaking times as discussed above and (ii) alloys 6N-14N which can contain the -phase, an instance which will be considered in chapter 4.3. The binder of alloy 2N is homogeneous and comprises only the bright Cu-rich (Cu,Ni) fcc solid solution Cu (Fig. 6a). The primary Mo crystallites being not only much smaller but also less rounded than in the Ni-rich alloys, i.e. S 82%, support the hypothesis of limited Moflux during liquid phase sintering. On the other hand, the binder of the Ni-rich alloys 6N-14N features two fcc (Cu,Ni) solid solutions, i.e. the brightCu-phase and the dark Ni-phase (Fig. 6 b-d). As can be inferred from these micrographs, the amount of Ni increases with the Nicontent. The SEM micrographs in Fig. 7 confirm the light microscopy observations. Consistent with their mean atomic mass, the contrast of the phases in the secondary electron micrographs changes from bright Mo-crystals, to dark grey forCu, which contains virtually no Mo, to medium grey forNi owing to its high solubility for Mo. Moreover, the small dispersed particles, which precipitated within the Ni-binder phase as shown in Fig. 7c-d, were identified to be bcc Mo by EDX and by EBSD (Fig. 8, Table 3). The number of the secondary and tertiary Mo precipitates grows with increasing Ni-content. Following this qualitative assessment, the chemical composition and the crystal structure of the phases were quantitatively assessed by a combination of EDX and EBSD analysis. The numbers in Fig. 7 denote the positions of the EDX measurements which are documented in Table 3. Evidently, the binder of 2N comprises virtually no Mo in solid solution 5, whereas its Cu/Ni ratio of 89/11 at% matches the distribution of both elements in the initial powder blend. In alloys 6N-14N, the Mo-content ofCu assumes 1.2 at%, which is close to the predictions of THERMOCALC, as are the Cu- and Ni-concentrations of the two binder phases. Table 3 also documents the high solubility of Ni for Mo, and that the primary and secondary Mo grains contain almost no Cu and Ni6. Finally, EDX-mappings of Mo, Cu and Ni, and an EBSD map together with the corresponding SEM micrograph are presented in Fig. 8 for alloy 14N. Besides showing primary, secondary and tertiary Mo grains, they confirm the existence of two fcc binder phases with high and low Mo-concentrations respectively and validate the light microscopy results described above. Finally, the spatial arrangement and the domain size of the immiscible liquids L Cu and LNi can be estimated from Figs. 6 and 7. The size of the Cu- and Ni-regions, and consequently the 5 6
The EDX detection limit is 0.5 wt.% The apparent Cu- and Ni-contents of secondary Mo crystals are attributed to the small size of the particles with respect to the diameter ( 1µm) of the EDX probe.
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size of the immiscible liquid domains from which they originate, is smaller than the primary Mo grain size. This is remarkable, as the binder domain size in WC-Co, which features a single rather than two liquids during LPS, is up to 50 times greater than the mean size of the WC grains [22]. Apart from the major constituents, it is instructive to take a closer look on the details of the microstructure. The SEM micrograph in Fig. 9a displays small Mo-precipitates (Mosec.) which are entirely surrounded by the medium-grey Ni-phase. This spatial arrangement, i.e. small, secondary Mo grains in close proximity to Ni, is exactly what invariant reaction U1’ predicts. Originating from LCu + LNi, both constituents are expected to develop in immediate vicinity to each other. Specifically, Mo originating from reaction U1’, i.e. secondary Mo, is predicted to precipitate as a discrete precipitate rather than to crystallize onto primary Mo grains. Further peculiarities of the microstructure are tertiary Mo-precipitates (Motert.), e.g. the one in Fig. 9b. In contrast to the secondary Mo grain in Fig. 9a, this Mo-particle is surrounded by the dark Cu-phase which, in turn, has grown into the Ni-phase. Such an arrangement is readily explained in terms of reaction U’2 that predicts LCu to react with Ni thereby generating tertiary Mo-precipitates and Cu, again in close proximity to each other. 4.3. The δ phase The δ-phase was never observed at elevated sintering temperatures or long holding times. Only during sintering at 1350° and short to intermediate soaking times does the intermetallic phase emerge in the microstructures of the Ni-rich compositions 6N, 10N and 14N. Its volume fraction increases with the Ni-content and decreases with soaking time, i.e. MoNi is never present after soaking for more than 2 hours. An example of a -precipitate in alloy 6N is shown in Fig. 10. Generally, δ is distributed inhomogeneously and occasionally forms clusters in the microstructure. This suggests that it originates from the reaction between neighboring elemental Mo- and Ni-powder particles during heat-up [16] whereas the clusters are a consequence of incomplete mixing of the powder blend. To eliminate , it is most important, that sintering takes place at temperatures above invariant reaction U2 (1248°C), because only there can δ finally dissolve in the liquid binder owing to its thermodynamic instability above that temperature. Dissolution being a time-sensitive mechanism, explains why remnants of survive at short and intermediate soaking times (Fig. 10). As soon as δ has been fully dissolved, it does not reprecipitate during cooling at rates on the order of several K/min due to the slow kinetics of reaction U2. Thus sintering at sufficiently high temperatures and reasonably fast cooling are key to produce -free microstructures, i.e. LPS-Mo composites which contain only ductile constituents. In the light of these arguments, the presence of the δ-phase in Mo-4Ni-2Cu observed by Hwang and Huang [15-17] is readily understood as a result of the 1300 °C sintering, i.e. at a temperature where the dissolution rate of , despite being thermodynamically unstable, is very slow. 4.4. Invariant Reactions The solidification reactions and the temperatures at which they occur were further analyzed by DSC experiments. A feature common to all heating DSC charts (Fig. 11) is the peak at
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1084 °C, which we attribute to the melting of elemental Cu particles as a consequence of the heterogeneity of the elemental powder blend. It also indicates that interactions of Mo and Ni with Cu particles are negligible up to 1100 °C. The Cu-peak, being prominent in alloy 2N, is also present in the Ni-rich alloys, yet decreases, as expected, with increasing Nicontent, and is virtually absent in alloy 14N (Fig. 11d). After having attained equilibrium at 1400 °C, the DSC cooling-signal of alloy 2N (Fig. 11 a) reveals a solidification interval for LCu between 1143 °C and 1128 °C, which agrees well with THERMOCALC’s predictions of 1126 °C and 1147 °C respectively (c.f. Fig. 2b). Apart from the Cu-peak, the DSC graphs of alloys 6N-14N show two major transformations during cooling (Figs. 11 b-c), which we attribute to the invariant non-equilibrium reactions U1’ and U2’. Indications for reactions including the δ-phase, i.e. DSC signals corresponding to the equilibrium reactions U2 at 1248 °C and U3 at 1231 °C were not observed. Notably, we found the invariant reactions during heating, i.e. where the elemental powder blend is not yet in equilibrium, to occur at nearly identical temperatures as during cooling, i.e. after equilibrium had been attained during sintering. This indicates that invariant reactions U1’ and U2’ are already operative during heat-up. Furthermore, the amplitudes of the experimental DSC-signals are consistent with expectations: as reaction U1’ generates Ni and LNi respectively, its DSC signals should increase with growing Ni-content as can be inferred from Figs. 11 b-c, and vice-versa the intensity of the Cu and LCu DSC-signals corresponding to reaction U’2 decrease with the Ni-content. Taking the mean of the heating and cooling cycle as an indicator for the transformation temperature, the peak in the interval around 1210-1220 °C coincides well with THERMOCALC’s prediction of 1214 °C for the temperature of invariant reaction U’2. However, the DSC signals in the interval 1290-1300 °C (Figs. 11 b-d), which we attribute to invariant reaction U’1, deviate by 30 K from THERMOCALC’s prediction of 1323 °C. The reasons for that inconsistency are as yet not fully clear. We believe that Yan et al. [20], relying on optical microscopy only, misinterpreted the microstructure in their 1300 °C samples by identifying a 3-phase mixture containing (bcc + fcc + liquid). In contrast, our DSC results suggest a (bcc + liquid(1) + liquid(2)) phase composition at 1300 °C. Nevertheless, regardless of the real U’1 temperature, the impact of this ambiguity on the results and conclusions presented above is negligible. 4.5. Grain Size and Shape Coarsening of the Mo grains during sintering is significant. An example of the evolution of the size and shape of primary Mo grains during sintering at 1400 °C is shown in Fig. 12. Attaining a size of 9 μm after 4h of sintering, as compared to the initial size of the Mo powder (3 µm), grain growth in alloy 2N is pronounced. Yet, it is considerably slower than in alloys comprising LNi (c.f. also Figs. 6 and 7). While the primary Mo grains in 6N have grown to 25 µm, the final grain sizes in both alloys 10N and 14N reach 38 µm, i.e. grain growth in 10N and 14N is nearly identical. The size of the primary Mo grains, G, follows the relationship (1) where t is the isothermal sintering time and G0 denotes the size of primary Mo grains at the onset of isothermal sintering (Fig. 12a). Notably, G0 is definitely larger than the initial Mopowder size, suggesting that Mo grains coarsen considerably before the isothermal sintering
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temperature has been reached. However, G0 is significantly smaller in alloy 2N (5 µm) than in the Ni-rich alloys (10 µm). We attribute this difference to the coarsening that occurs in the temperature interval after the first liquid has formed and before isothermal sintering begins. The absence of LNi in alloy 2N retards grain growth and densification, contrary to the situation in alloys 6N-14N. Finally, the t1/3 time-dependence suggests that growth of the Mo grains is diffusion controlled [14, 23]. While the primary Mo crystals in the high Ni-alloys (6N-14N) are close to spherical (89% < S < 93%), they are less rounded in 2N (S 82%) (Fig. 12b, c.f. also Figs. 6 and 7). This indicates again that solution-reprecipitation in the high Ni-alloys is faster due to the high solubility of LNi for Mo. The Mo grains in “LCu-only” alloys being smaller, more facetted and less spherical, suggest that solid state sintering during heat-up plays a major role in densification. More details on the evolution of the microstructure in Mo-Cu-Ni alloys and on grain growth will be discussed in a companion paper [23]. 5. Concluding Remarks Two major conclusions can be drawn from this research. Firstly, the Mo-Cu-Ni system fulfills all criteria for liquid phase sintering, i.e. chemical compatibility, solubility and wetting by the liquid binder. Depending on the Ni-content, liquid phase sintering in Mo-Cu-Ni takes place in the presence of one or two liquid phases. Cu-rich alloys comprise a single-phase Cu-rich liquid, LCu, during sintering and, due to its low Mo-solubility, sinter slowly. In contrast, the Nirich alloys sinter in the presence of two melts, LCu and LNi, where the latter liquid promotes densification significantly due to its high solubility for Mo. In both cases, LPS generates a fully dense microstructure with equiaxed Mo particles embedded into a ductile fcc (Cu,Ni)-binder matrix, provided sintering conditions are selected which suppress the formation of the brittle MoNi phase. This requires (i) sintering above the U2’ temperature to dissolve that has possibly formed during heat-up and (ii) subsequent cooling at sufficiently fast rates to suppress its reprecipitation. Verified by sintering and DSC experiments as well as by microstructural and microanalytical observations, cooling rates on the order of 10 K/min ensure fully dense composites which comprise only ductile microstructural constituents. The non-equilibrium phase diagrams predict a miscibility gap not only in the liquid but also in the solid state of Mo-Cu-Ni, which explains the observed dual-phase microstructure in the binder of high Ni-alloys comprising a Cu- and a Ni-rich fcc (Cu,Ni) solid solution, Cu and Ni, respectively. They originate essentially from the two immiscible liquids, LCu and LNi, which solidify via reactions U1’ and U2’ into Cu, Ni and small Mo-precipitates. The existence of the two invariant reactions, U1’ and U2’, has been confirmed by DSC experiments. Densification and grain growth during LPS are largely governed by the Mo-flux during the solution-reprecipitation process. The Mo-transportation is controlled by the solubility of Mo in the liquid phases. As LCu and LNi comprise low and high solubility for Mo respectively, their actual volume fractions, which can be adjusted by variation of the Ni-concentration of the binder, dictate the overall Mo-transportation capacity during sintering. Thus the Ni-content is an efficient tool to control densification and grain growth in LPS-Mo.
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Finally, this investigation has generated insight into the densification and grain growth of solids which sinter in the presence of two liquid phases. On one hand those liquids generate a dual-phase binder, as compared to the single phase binders of conventional LPS systems. On the other hand they offer easy control of the solubility of the major phase during LPS. While the solubility in single liquid systems can be essentially affected by temperature only, sintering with two liquids permits the solubility of the major phase to be adjusted by selecting appropriate volume fractions of the two liquids.
Acknowledgements We would like to thank Plansee SE for the permission to publish this work. Valuable discussions with B. Kieback, H. Kestler and H.J. Seifert are gratefully acknowledged. Declaration of interests
☒ The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. ☐The authors declare the following financial interests/personal relationships which may be considered as potential competing interests:
page - 11 -
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M.V. Sundaram, K.B. Surreddi, E. Hryha, A. Veiga, S. Berg, F. Castro, L. Nyborg, Enhanced Densification of PM Steels by Liquid Phase Sintering with a Boron-Containing Master Alloy, Metall. and Mat. Trans. A 49A, (2018), 255-263
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O.-H. Kwon, G.L. Messing, Kinetic Analysis of Solution‐Precipitation during Liquid‐Phase Sintering of Alumina, J. Am. Cer. Soc. 73, (1990), 275-281,
(5)
S. Hampshire, Silicon nitride ceramics - review of structure, processing and properties, J. Achiev. Mater. Manuf. Eng. 24, (2007), 43-50
(6)
H.E. Exner, Physical and chemical nature of cemented carbides, Int. Met. Rev. 24, (1979), 149-173
(7)
L.S. Sigl, H.J. Kleebe, Core-Rim Structure of Liquid Phase Sintered SiC, J. Am. Cer. Soc. 76, (1993), 561-565
(8)
R.M. German, P. Suri, S.J. Park, Review: liquid phase sintering, J. Mater. Sci. 44, (2009), 1-39
(9)
D.N. Yoon, W.J. Huppmann, Grain Growth and Densification during Liquid Phase Sintering of W-Ni, Acta Met. 17, (1979), 693-698
(10) A. Bose, B.H. Rabin, R.M. German, Liquid Phase Sintering of Tungsten Heavy Alloys in Vacuum, Met. Powder Rep. 42, (1987), 834-839 (11) K. Hu, X. Li, S. Qu, Y. Li, Effect of Heating Rate on Densification and Grain Growth during Spark Plasma Sintering of 93W-5.6Ni-1.4Fe Heavy Alloys, Metall. Mater. Trans. A 44 (2013), 4323-4336 (12) E. Kny, Properties and Uses of the Pseudobinary Alloys of Cu with Refractory Metals, Proceedings of the 12th Plansee Seminar, Eds. H. Bildstein and H.M. Ortner, Metallwerk Plansee GmbH, Reutte (Austria), vol. 1, (1990), p. 763-772 (13) J.L. Johnson, R.M. German, Powder metallurgy processing of Mo-Cu for thermal management applications, Int. J. Powder Metall. 35, (1999), 39-48 (14) A. Huber, Flüssigphasensintern von Molybdän, Ph.D. thesis, Technical University Dresden, (2019) (15) K.S. Hwang, H.S. Huang, The liquid phase sintering of molybdenum with Ni and Cu additions, Materials Chemistry and Physics 67, (2001), 92-100
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(16) K.S. Hwang, H.S. Huang, Identification of the segregation layer and its effects on the activated sintering and ductility of Ni-doped molybdenum, Acta Mater. 51, (2003), 3915-26 (17) K.S. Hwang, H.S. Huang, Ductility improvement of Ni-added molybdenum compacts through the addition of Cu and Fe powders, Int. J. Ref. Met. Hard Mat. 22, (2004), 185191 (18) A. Bose, R.M. German, Sintering Atmosphere Effects on Tensile Properties of Tungsten Heavy Alloys, Met. Trans. A 19A, (1988), 2467-2476 (19) Thermo-Calc Software AB, Description of the TCNI.8 database of THERMOCALC, url: http://www.thermocalc.com/media/23650/tcni8_extended_info.pdf, last visited: 05.12.2018 (20) Y. Yan, W. Cuiping, L. Yifu, L. Xingjun, K. Ryosuke, I. Kiyohito, Experimental investigation and thermodynamic calculation of the phase equilibria in the Cu-Mo-Ni ternary system, Materials Chemistry and Physics 125, (2011), 37-45 (21) C.B. Shoemaker, D.P. Shoemaker, The crystal structure of the -phase MoNi, Acta Cryst. 16, (1963), 997-1009 (22) K.P. Mingard, B. Roebuck, J. Marshall, G. Sweetman, Some aspects of the structure of cobalt and nickel binder phases in hardmetals, Acta Mater. 59, (2011), 2277-2290 (23) A. Huber, L.S. Sigl, Grain Growth in Liquid Phase Sintered Molybdenum, to be published
Tables Table 1: Properties of the Mo, Ni, and Cu powders. Powder
Mo
Cu
Ni
Particle size D10, µm
1.2
3.4
1.5
Particle size D50, µm
2.0
4.7
5.2
Particle size D90, µm
3.4
6.5
12.5
Oxygen content, µg/g
345
1,451
2,865
Carbon content, µg/g
15
637
227
Nitrogen content, µg/g
19
10
17
Global Tungsten and Powders
Ecka Granules
Eurotungstene
Supplier
Table 2: Calculated compositions of the liquid phases LCu and LNi at 1400 °C.
page - 13 -
Composition in at.% at 1400 °C
Alloys 2N
6N, 10N, 14N
LCu
LCu
LNi
Mo
0.4
1.1
28.3
Cu
89.6
71.0
9.7
Ni
10.0
27.9
62.0
Table 3: Composition of the microstructural constituents in Fig. 7 (EDX measurements). Alloy 2N 6N 10N
14N
Position #
Microstructural Constituent
Cu at.%
Ni at.%
Mo at.%
1
Cu
89.1
10.9
0.0
2
primary Mo
0.0
0.4
99.6
3
Cu
69.2
29.6
1.2
4
Ni
16.7
65.9
17.4
5
Cu
66.3
32.6
1.1
6
Ni
19.5
65.2
15.3
7
Cu
59.9
38.8
1.3
8
Ni
18.4
66.5
15.1
9
secondary Mo
3.0
9.6
87.4
Figures
Fig. 1 Calculated isothermal section of Mo-Cu-Ni at 1400 °C with the four investigated alloys (solid red points) at 82 wt.% Mo (red solid line) and the vertical sections at constant Ni/Cu ratios (red dashed lines). The three-phase field with Mo and the two liquids LCu and LNi is shown in green. page - 14 -
2a
2b
Fig. 2: Calculated isopleth sections of the Mo-Cu-Ni system at 82 wt.% Mo showing the four investigated compositions 2N-14N (dashed vertical lines): (a) equilibrium diagram with three invariant reactions U1-U3 (b) non-equilibrium diagram with two invariant reactions U1’ and U2’.
3a
3b
3c
3d
Fig. 3: Calculated isopleth sections along the dashed red lines of Fig. 1 at constant Cu:Ni ratios of a) 16:2, b) 12:6, c) 8:10 and d) 4:14 showing the evolution of the liquid miscibility gap in Mo-Cu-Ni.
page - 15 -
Fig. 4: Calculated volume fractions of liquid phases at 1400 °C
5a
5b
Fig. 5: As-sintered densities of alloys 2N (a) and 6N (b) as a function of sintering temperature and soaking time.
6a)
6b)
page - 16 -
6c)
6d)
Fig. 6: Light microscopy images of alloys 2N (a), 6N (b), 10N (c) and 14N (d) sintered at 1400 °C for 4h with three phases: Mo (light grey), Cu (white),Ni (dark grey).
7a)
7b)
7c)
7d)
Fig. 7: Secondary electron images (SEM) of alloys 2N (a), 6N (b), 10N (c) and 14N (d) sintered at 1400 °C for 4h. Mo is white, Cu is dark grey, andNi is medium grey. The red circles point to the position of the EDX measurements documented in Table 3.
page - 17 -
8a
8b
8c
8d
8e Fig. 8 Composition and phase analysis of alloy 14N: (a-c) EDX-maps of Mo, Ni and Cu, (d) EBSD-map, (e) SE image.
9a
9b
page - 18 -
Fig. 9: Primary Mo grains and Mo precipitates (white) in the binder of alloy 10 N: (a) secondary Mo (Mosec.) surrounded by medium grey Ni binder-phase originating from invariant reaction U1’ (b) tertiary Mo surrounded by dark grey Cu binder-phase originating from invariant reaction U2’
Fig. 10: SEM micrograph (SE-image) of the δ-phase (point #1) in alloy 6NEDX composition of point #1 in at.%: Mo 47.7, Ni 49.5, Cu 2.9
11a)
11c)
11b)
11d)
page - 19 -
Fig. 11: DSC diagrams of Mo-Cu-Ni alloys a) 2N b) 6N, c) 10N and d) 14N
12a)
12b)
Fig. 12: Evolution of (a) grain size, G, and (b) grain shape, S, with isothermal soaking time, t, in alloys 2N-14N (sintering temperature 1400 °C). Grain growth follows a t1/3 law.
Graphical Abstract
page - 20 -
On the Sintering of Molybdenum with two Liquid Phases Andrea Huber , Lorenz S. Sigl PII: DOI: Reference:
S2589-1529(20)30017-X https://doi.org/10.1016/j.mtla.2020.100600 MTLA 100600
To appear in:
Materialia
Received date: Accepted date:
17 December 2019 18 January 2020
Please cite this article as: Andrea Huber , Lorenz S. Sigl , On the Sintering of Molybdenum with two Liquid Phases, Materialia (2020), doi: https://doi.org/10.1016/j.mtla.2020.100600
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On the Sintering of Molybdenum with two Liquid Phases Andrea Huber1 and Lorenz S. Sigl2 1 2
Plansee SE, Reutte, Austria
Technical University Dresden, Dresden, Germany
Abstract The liquid phase sintering (LPS) of Molybdenum-base alloys containing 80 vol.% Mo and 20 vol.% Cu and Ni as sintering aids was investigated. Sintering experiments and subsequent microstructural analysis together with equilibrium and non-equilibrium thermodynamic calculations elucidate the densification behavior and the microstructure of Mo-Cu-Ni alloys. LPS is dominated by the liquid miscibility gap in the Mo-Cu-Ni system. Surrounding the Mograins during sintering, a Cu-rich and a Ni-rich liquid phase with moderate and high solubility for Mo respectively, generate fully dense and ductile composites via the solution and reprecipitation mechanism. The rates of densification and microstructural coarsening increase significantly with the amount of Ni-rich liquid, i.e. with the solubility of the binder for Mo. Upon cooling, the two immiscible liquids solidify into a dual-phase binder comprising a Cuand a Ni-rich fcc (Cu,Ni) solid solution respectively. Precipitation of the brittle MoNi phase (), which was reported in previous studies, can be suppressed provided that (i) sintering happens at temperatures where , which develops during heat-up of the elemental Mo-CuNi powder blend, dissolves in the liquid phases and (ii) cooling from sintering is sufficiently fast to suppress its reprecipitation.
Key Words Molybdenum alloy; liquid phase sintering; liquid immiscibility; densification; grain growth 1. Introduction Liquid phase sintering (LPS) is a well-established method to consolidate powder compacts into fully dense multi-phase components. The process is used to manufacture metals including tungsten heavy alloys [1] and some P/M steels [2, 3]. It is also applied to densify various ceramics and ceramic-metal composites, e.g. certain aluminas [4], the silicon nitride materials family [5], liquid phase sintered silicon carbide [6], and not least cemented carbides [7]. Sintering studies on various materials revealed that, despite the heterogeneities in the solid constituents, the melt during LPS is generally homogeneous, i.e. single phase [8]. Despite striking similarities between W and Mo, the assistance of a liquid phase analogous to the manufacturing of tungsten heavy alloys (WHAs) [1, 9-11] has, apart from “nonsintering” densification such as the infiltration of pre-sintered, porous Mo-skeletons with Cu [12,13], not been employed to densify Mo-base materials. This knowledge gap has triggered a thesis, which intended to assess the potential of various metallic binders for the densification of Mo by LPS [14]. Combining thermodynamic calculations and sintering experiments, page - 1 -
that study identified Cu and Ni as most promising alloy additions, because the Mo-Cu-Ni system can satisfy the essential criteria for liquid phase sintering: (i) wettability of Mo by the molten binder, (ii) sufficient solubility of Mo in the liquid binder and (iii) chemical compatibility between Mo and the binder to prevent the formation of brittle intermetallic phases. Only recently, Hwang and Huang investigated Mo-base alloys with a binder containing pure Ni and Ni/Cu alloys [15-17]. They demonstrated that the composites can be sintered to full density, yet found them to be brittle rather than ductile due to the formation of the intermetallic phase MoNi () in the binder [15,16]. Depending on the Ni/Cu composition, the authors identified activated or liquid phase sintering as operating densification mechanisms [16]. However, they did neither recognize the importance of the liquid miscibility gap (LMG) in the Mo-Cu-Ni system at temperatures above 1320 °C nor its impact on liquid phase sintering. It is the intention of this paper to elucidate the constitutional fundamentals for producing dense and ductile Mo-base alloys, by (i) investigating the constitution of the Mo-Cu-Ni system under equilibrium and non-equilibrium conditions, (ii) conducting sintering experiments and (iii) validating the computational results through microstructural and analytical observations. 2. Compositions and experimental procedures Commercial powders from industrial suppliers were used in this study. Their size and chemical composition are summarized in Table 1. To investigate liquid phase sintering in the MoCu-Ni system, a set of four compositions was chosen with the rationale to investigate Moalloys at a fixed binder fraction of 20 vol.%1 at ambient temperature. The choice of binder composition was guided by the isothermal section of the phase diagram at the intended sintering temperature of 1400 °C (Fig. 1). Characteristic for Mo-Cu-Ni is a miscibility gap in the liquid as shown in Fig. 1. Starting with the Cu-rich alloy 2N2, which is positioned in the two-phase field comprising solid Mo (bcc) and a single melt, LCu, an increase of the Ni-content in three equidistant steps of 4 wt.% Ni (alloys 6N, 10N, 14N) switches the phase configuration into the three-phase field which, besides Mo, contains two liquid phases, i.e. a Cu- and a Ni-rich melt, LCu and LNi, respectively (Figs. 1 and 2). The positions of the alloys 2N: 6N: 10N: 14N:
82 wt.% Mo – 16 wt.% Cu – 2 wt.% Ni 82 wt.% Mo – 12 wt.% Cu – 6 wt.% Ni 82 wt.% Mo – 8 wt.% Cu – 10 wt.% Ni 82 wt.% Mo – 4 wt.% Cu – 14 wt.% Ni
at the sintering temperature of 1400 °C are marked by the solid circles in Figs. 1-3. The alloys were prepared by dry mixing a total of 300 grams of elemental Mo/Cu/Ni powder blends with 100 grams zirconia balls for 1 hour in a Turbula mixer. After drying and subse-
1 2
20 vol.% binder corresponds approximately to 18 wt.% binder and 82 wt.% Mo the numeral denotes the Ni-content of the Mo-Cu-Ni alloy in wt.%
page - 2 -
quent compaction of cylindrical tablets (Ø 25 mm, height 5 mm) at 180 MPa in an axial press (Komage 100CNC), a green density of 65% of theoretical density (T.D.) was reached. Liquid phase sintering was carried out in a tube furnace (Xerion Xtube 1650) under flowing hydrogen at a flow rate of 0.5 l/min and holding times between 0h and 4h at temperatures of 1350 °C, 1400 °C and 1450 °C, applying a heating and a cooling rate of 10 K/min respectively. From WHAs it is well known, that large pores can develop during sintering under dry hydrogen, a phenomenon which has been attributed to the formation of water vapor [18]. Likewise, pore formation in Mo-Cu-Ni was successfully blocked by setting the hydrogen dew point to -55°C below 900 °C and to 20 °C above that temperature. The as-sintered densities were measured by Archimedes’ method. Metallographic samples were prepared by grinding on SiC papers up to P2400, followed by diamond polishing from 9 µm down to 1 µm and a final polishing step with 0.04 µm colloidal silica. Both optical (Leica Microsystems DMI 5000) and scanning electron microscopes (Carl Zeiss Microscopy Ultra 55 FEG) were used to investigate the microstructure on unetched specimens, as the different phases are easier visible on as-polished surfaces. The grain size, G, of the Mo particles was measured by the linear intercept method and their sphericity, S, was evaluated from the relationship S = 4 A/P2 where A denotes the area and P the perimeter of individual Mo crystallites. Quantitative energy dispersive X-ray analysis3 (acceleration voltage 20 kV, probe current 0.9 nA, standardless ZAF correction) was used to assess the chemical composition of the various phases. They were further analyzed by a combination of EDX and electron back scattered diffraction (EBSD3). Finally, differential scanning calorimetry (DSC) measurements (Netzsch STA 409 CD) were conducted in an Argon atmosphere by heating the samples up to 1400 °C with subsequent cooling at a cooling and heating rate of 10K/min respectively. 3. Thermodynamic calculations Thermodynamic calculations were performed with the commercial THERMOCALC software using the Ni-database TCNI8 [19], where the system Mo-Cu-Ni is covered in the full composition and temperature range based on the investigations of Yan et al. [20]. The Mo-Cu-Ni system was investigated by isothermal and isopleth sections as well as by evaluating the fraction of phases as a function of temperature for selected compositions. Notably, THERMOCALC permits non-equilibrium states to be explored by (i) calculating Scheil-Gulliversolidification simulations and (ii) by implementing those results into subsequent nonequilibrium phase diagram calculations. Both tools will prove to be useful for predicting the formation and spatial arrangement of phases under technical cooling conditions. 3.1. Equilibrium phase relationships in the Mo-Cu-Ni system As the topic of this paper is liquid phase sintering, it is convenient to start with a close look on the equilibrium phase diagrams at sintering temperature, e.g. on the sections of the MoCu-Ni system at 1400 °C (Fig. 1) and at 82 wt.% Mo (Fig. 2). The isothermal 1400 °C section reveals three one-phase rooms, i.e. bcc Mo, an fcc (Cu,Ni) solid solution, and a (Cu-Ni-Mo) liquid. That liquid is single-phase close to the Cu-Ni line, yet destabilized by the addition of
3
EDAX-AMETEK Materials Analysis Division
page - 3 -
Mo which causes a miscibility gap containing two liquids. While alloy 2N is located in the two-phase field with the Cu-rich melt, LCu, the Ni-rich alloys 6N-14N are located in a threephase field comprising the extra liquid, LNi. Hence, in the Ni-rich alloys, sintering above 1323 °C happens in the presence of two liquids, whereas only one liquid phase is active in alloy 2N. Notably, the Cu-rich liquid exhibits moderate solubility for Mo, whereas L Ni can dissolve 25 times more Mo (Table 2 and Fig. 1). 3.1.1. The liquid miscibility gap in the binder during sintering To our best knowledge, LPS in the presence of more than one liquid phase has as yet not been described in the literature. It is not the intension of this paper to discuss the reasons for the presence of an LMG in Mo-Cu-Ni, but to focus on the impact of the two liquid phases on the sinterability of Mo and on the microstructure developing from them. The ternary LMG in Mo-Cu-Ni originates from the liquid immiscibility in the Mo-Cu system, where it appears at temperatures above 2500 °C. Due to the high temperature level, this aspect has not been paid much attention, because the LMG is irrelevant for the processing of commercial Mo-Cu materials. In Cu-rich alloy 2N the liquid immiscibility becomes evident only at temperatures above 2300 °C, whereas a 2-phase field (Mo+LCu) obtains at sintering temperature (Fig. 3a). Adding more Ni drops the critical point of the ternary LMG significantly and expands the 3-phase field (Mo + LCu + LNi,) to temperatures as low as 1323 °C (Figs. 3b-d). Notably, two liquids are present at all sintering temperatures in alloys 6N, 10N and 14 N. The two liquids differ considerably with respect to their solubility for Mo: outside the LMG, the Mo-solubility of LCu increases only weakly with the Ni-concentration (c.f. Fig. 1), and dissolves just 0.45 at.% Mo in 2N at 1400 °C (Table 2). Within the three-phase field (Mo+ LCu+ LNi), where the compositions of LCu and LNi are fixed, the Mo-solubility at 1400 °C is still only 1.1 at.% for LCu, but 28.3 at.% in LNi (Table 2). Finally, it should be recalled that the volume fractions of LNi and LCu in- and decrease with the Ni-concentration as illustrated in Fig. 4. As will be shown below, the disproportionate solubility of Mo in the two liquids, LCu and LNi, affects densification via solution and reprecipitation deeply, and has a strong impact on the kinetics of grain growth. 3.1.2. Invariant reactions during equilibrium cooling from sintering temperature The isopleth sections at 82 wt.% Mo in Fig. 2 illustrate the phase relationships during cooling from sintering temperature. Under equilibrium conditions (Fig. 2a), THERMOCALC predicts three invariant reactions in the Ni-rich alloys 6N, 10N, and 14N: ↔
(U1)
↔
(U2)
↔
(U3)
where LCu and LNi are the liquid phases defined above, is an fcc (Cu,Ni) solid solution, Mo denotes the bcc Mo-phase, and (MoNi) is a brittle, intermetallic compound with orthorhombic lattice [16,21]. Cu-rich alloys with less than 4.5 wt.% Ni do not undergo the invari-
page - 4 -
ant reactions U1 - U3, but precipitate at elevated temperatures, e.g. for 2N below 1126°C, and also the -phase at low temperatures (Fig. 2a). Reaction U1 is calculated to take place at 1323 °C, i.e. slightly below the lowest sintering temperatures investigated here. THERMOCALC forecasts U1 to consume predominantly the Nirich liquid and to generate together with Mo precipitates. Below 1248 °C the intermetallic -phase is predicted to crystallize according to reaction U2. However, as U2 is a reaction between solid phases, mass transport is a lot slower than in reactions involving a liquid, suggesting that U2 is sluggish. Furthermore, can precipitate only in the temperature interval between U2 and U3. As this accounts for not more than 17 K, the experimental cooling rate of 10 K/min allows a time span of merely 100 seconds for the precipitation of . Finally, reaction U3 being inverse to U2 would partly reconsume the newly formed δ-phase during further cooling. These arguments suggest the precipitation of to be unlikely under practical cooling conditions. Consequently, the equilibrium prediction of -precipitation has to be checked against (i) thermodynamic non-equilibrium considerations as presented below and (ii) experimental observations as discussed in chapter 4. 3.2. Non-equilibrium solidification from sintering temperature With the Scheil-Gulliver module, THERMOCALC provides a computational tool that takes the kinetics of phase evolution during solidification into account, i.e. it computes the phases which develop during solidification and their volume fraction as a function of temperature assuming zero diffusivity in the solid, infinitely fast mass transport in the liquid and local equilibrium at the solid/liquid interface. In contradiction to the equilibrium results, ScheilGulliver predicts δ not to form during cooling, which supports the arguments discussed above. Whereas both, the liquid miscibility gap and reaction U1 remain virtually unaffected by the suppression of δ, i.e. we still have (U’1 U1)
↔ the invariant reactions U2 and U3 merge into reaction U’2 according to ↔
(U’2)
where Ni and Cu denote a Ni- and Cu-rich (Cu,Ni) fcc solid solution respectively. ScheilGulliver forecasts the following propositions, which will be compared with microstructural observations in section 4: i. Reaction U’1 generates Mo which is expected to co-precipitate with the Ni-rich (Cu,Ni) fcc solid solution Ni. This should generate a microstructural topology containing secondary molybdenum, Mosec., in intimate contact with Ni rather than the crystallization of excess Mo onto the primary Mo-crystals. THERMOCALC predicts the volume ratio of Ni /Mosec. to be 85%/15%. ii. During invariant reaction U’2 the residual Cu-rich melt, LCu, solidifies and eats up Ni to produce a Cu-rich (Cu,Ni) fcc solid solution Cu, and tertiary Mo precipitates , Motert., which should be, again, in close contact. iii. The intermetallic phase is unstable and does not form.
page - 5 -
Hence, a non-equilibrium phase diagram with being suppressed was calculated and is shown in Fig. 2b. It suggests that the Mo-Cu-Ni system meets the criteria for liquid phase sintering, i.e. i. LCu and LNi dissolve Mo in sufficient quantities to operate the solution-reprecipitation mechanism, despite their solubility for Mo differs by more than an order of magnitude (c.f. Table 2). ii. The brittle MoNi phase does not form under industrially relevant cooling conditions such that the criterion of chemical compatibility is met. iii. Solid Mo is wetted by the liquid phases [15]. iv. Finally, the densification temperature of Mo-Cu-Ni is markedly lower than for solid-state sintered molybdenum (> 2000 °C). 1. Results and Discussion 4.1. Densification Behavior The densification characteristics of Mo-Cu-Ni alloys are exemplified in Fig. 5. The as-sintered densities of alloys containing only LCu such as 2N (Fig. 5a) depend strongly on sintering time and temperature, i.e. they attain nearly full density only after soaking times > 2h. On the other hand, the Ni-rich compositions 6N-14N comprising the Ni-rich melt, LNi, achieve nearly full densities at all sintering temperatures and soaking times. In those alloys, nearly complete densification, as exemplified by alloy 6N in Fig. 5b, is obtained already at zero holding time. Analogous consolidation results were seen in alloys 10N and 14N [14]. Though at first surprising, this observation is readily understood as a consequence of the high Mo-solubility of LNi (Table 2). As the Ni-rich liquid becomes stable above 1323 °C (c.f. the isopleth in Fig. 2b) and increases its share with the Ni-content as shown in Fig. 4, it provides plenty of transportation capacity for Mo already during heat-up. Also, with the cooling and heating rate of 10 K/min, there is a sufficient time window of at least 5 minutes4 during which the solution-reprecipitation mechanism can redistribute Mo efficiently even before the onset of isothermal sintering. On the other hand, the as-sintered density of alloy 2N, containing just LCu with low Mosolubility, depends markedly on the soaking time ts (Fig. 5a), e.g. only 82 % T.D. is achieved at ts = 0 and 1350 °C. However, nearly full density is obtained after 4 hours of soaking, which can be understood as follows: Densification at short soaking times is to a good part due to solid state sintering during heat-up. Despite LCu becoming stable above 1126 °C, its low solubility for Mo, e.g. 0.36, 0.45 and 0.55 at.% at 1350, 1400 and 1450 °C respectively, limits the Mo-flux through the liquid. This view is supported by the moderate sphericity, S, of the Mo grains in 2N, which are less spherical than the Mo grains in 6N-14N (c.f. Figs 6a, 7a and 12b). Nearly theoretical density is nevertheless obtained after long soaking, indicating that LCu does transport Mo sufficiently over sintering time.
4
During heating from the invariant reaction U1 up to the sintering temperature, Ts, and cooling back to U1, there is a time interval tD for solution/reprecipitation given by tD = 2(Ts-1323)/R where Ts denotes the sintering temperature and R is the heating/cooling rate. With Ts = 1350 °C and R = 10 K/min, tD 5.5 min obtains.
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4.2. Microstructure and Microstructural Analysis Light microscopy (LM) images of alloys 2N-14N, which were sintered at 1400 °C for 4h, are displayed in Fig. 6 a-d. All compositions are fully dense showing virtually no porosity, in agreement with the density measurements from Fig. 5. The microstructure, which looks geometrically similar to tungsten heavy alloys, features equiaxed, rounded Mo particles (“primary” Mo) embedded into a continuous binder matrix. Corresponding results were obtained for the other sintering temperatures and soaking times, with the exception of (i) alloy 2N which does not sinter to full density at short soaking times as discussed above and (ii) alloys 6N-14N which can contain the -phase, an instance which will be considered in chapter 4.3. The binder of alloy 2N is homogeneous and comprises only the bright Cu-rich (Cu,Ni) fcc solid solution Cu (Fig. 6a). The primary Mo crystallites being not only much smaller but also less rounded than in the Ni-rich alloys, i.e. S 82%, support the hypothesis of limited Moflux during liquid phase sintering. On the other hand, the binder of the Ni-rich alloys 6N-14N features two fcc (Cu,Ni) solid solutions, i.e. the brightCu-phase and the dark Ni-phase (Fig. 6 b-d). As can be inferred from these micrographs, the amount of Ni increases with the Nicontent. The SEM micrographs in Fig. 7 confirm the light microscopy observations. Consistent with their mean atomic mass, the contrast of the phases in the secondary electron micrographs changes from bright Mo-crystals, to dark grey forCu, which contains virtually no Mo, to medium grey forNi owing to its high solubility for Mo. Moreover, the small dispersed particles, which precipitated within the Ni-binder phase as shown in Fig. 7c-d, were identified to be bcc Mo by EDX and by EBSD (Fig. 8, Table 3). The number of the secondary and tertiary Mo precipitates grows with increasing Ni-content. Following this qualitative assessment, the chemical composition and the crystal structure of the phases were quantitatively assessed by a combination of EDX and EBSD analysis. The numbers in Fig. 7 denote the positions of the EDX measurements which are documented in Table 3. Evidently, the binder of 2N comprises virtually no Mo in solid solution 5, whereas its Cu/Ni ratio of 89/11 at% matches the distribution of both elements in the initial powder blend. In alloys 6N-14N, the Mo-content ofCu assumes 1.2 at%, which is close to the predictions of THERMOCALC, as are the Cu- and Ni-concentrations of the two binder phases. Table 3 also documents the high solubility of Ni for Mo, and that the primary and secondary Mo grains contain almost no Cu and Ni6. Finally, EDX-mappings of Mo, Cu and Ni, and an EBSD map together with the corresponding SEM micrograph are presented in Fig. 8 for alloy 14N. Besides showing primary, secondary and tertiary Mo grains, they confirm the existence of two fcc binder phases with high and low Mo-concentrations respectively and validate the light microscopy results described above. Finally, the spatial arrangement and the domain size of the immiscible liquids L Cu and LNi can be estimated from Figs. 6 and 7. The size of the Cu- and Ni-regions, and consequently the 5 6
The EDX detection limit is 0.5 wt.% The apparent Cu- and Ni-contents of secondary Mo crystals are attributed to the small size of the particles with respect to the diameter ( 1µm) of the EDX probe.
page - 7 -
size of the immiscible liquid domains from which they originate, is smaller than the primary Mo grain size. This is remarkable, as the binder domain size in WC-Co, which features a single rather than two liquids during LPS, is up to 50 times greater than the mean size of the WC grains [22]. Apart from the major constituents, it is instructive to take a closer look on the details of the microstructure. The SEM micrograph in Fig. 9a displays small Mo-precipitates (Mosec.) which are entirely surrounded by the medium-grey Ni-phase. This spatial arrangement, i.e. small, secondary Mo grains in close proximity to Ni, is exactly what invariant reaction U1’ predicts. Originating from LCu + LNi, both constituents are expected to develop in immediate vicinity to each other. Specifically, Mo originating from reaction U1’, i.e. secondary Mo, is predicted to precipitate as a discrete precipitate rather than to crystallize onto primary Mo grains. Further peculiarities of the microstructure are tertiary Mo-precipitates (Motert.), e.g. the one in Fig. 9b. In contrast to the secondary Mo grain in Fig. 9a, this Mo-particle is surrounded by the dark Cu-phase which, in turn, has grown into the Ni-phase. Such an arrangement is readily explained in terms of reaction U’2 that predicts LCu to react with Ni thereby generating tertiary Mo-precipitates and Cu, again in close proximity to each other. 4.3. The δ phase The δ-phase was never observed at elevated sintering temperatures or long holding times. Only during sintering at 1350° and short to intermediate soaking times does the intermetallic phase emerge in the microstructures of the Ni-rich compositions 6N, 10N and 14N. Its volume fraction increases with the Ni-content and decreases with soaking time, i.e. MoNi is never present after soaking for more than 2 hours. An example of a -precipitate in alloy 6N is shown in Fig. 10. Generally, δ is distributed inhomogeneously and occasionally forms clusters in the microstructure. This suggests that it originates from the reaction between neighboring elemental Mo- and Ni-powder particles during heat-up [16] whereas the clusters are a consequence of incomplete mixing of the powder blend. To eliminate , it is most important, that sintering takes place at temperatures above invariant reaction U2 (1248°C), because only there can δ finally dissolve in the liquid binder owing to its thermodynamic instability above that temperature. Dissolution being a time-sensitive mechanism, explains why remnants of survive at short and intermediate soaking times (Fig. 10). As soon as δ has been fully dissolved, it does not reprecipitate during cooling at rates on the order of several K/min due to the slow kinetics of reaction U2. Thus sintering at sufficiently high temperatures and reasonably fast cooling are key to produce -free microstructures, i.e. LPS-Mo composites which contain only ductile constituents. In the light of these arguments, the presence of the δ-phase in Mo-4Ni-2Cu observed by Hwang and Huang [15-17] is readily understood as a result of the 1300 °C sintering, i.e. at a temperature where the dissolution rate of , despite being thermodynamically unstable, is very slow. 4.4. Invariant Reactions The solidification reactions and the temperatures at which they occur were further analyzed by DSC experiments. A feature common to all heating DSC charts (Fig. 11) is the peak at
page - 8 -
1084 °C, which we attribute to the melting of elemental Cu particles as a consequence of the heterogeneity of the elemental powder blend. It also indicates that interactions of Mo and Ni with Cu particles are negligible up to 1100 °C. The Cu-peak, being prominent in alloy 2N, is also present in the Ni-rich alloys, yet decreases, as expected, with increasing Nicontent, and is virtually absent in alloy 14N (Fig. 11d). After having attained equilibrium at 1400 °C, the DSC cooling-signal of alloy 2N (Fig. 11 a) reveals a solidification interval for LCu between 1143 °C and 1128 °C, which agrees well with THERMOCALC’s predictions of 1126 °C and 1147 °C respectively (c.f. Fig. 2b). Apart from the Cu-peak, the DSC graphs of alloys 6N-14N show two major transformations during cooling (Figs. 11 b-c), which we attribute to the invariant non-equilibrium reactions U1’ and U2’. Indications for reactions including the δ-phase, i.e. DSC signals corresponding to the equilibrium reactions U2 at 1248 °C and U3 at 1231 °C were not observed. Notably, we found the invariant reactions during heating, i.e. where the elemental powder blend is not yet in equilibrium, to occur at nearly identical temperatures as during cooling, i.e. after equilibrium had been attained during sintering. This indicates that invariant reactions U1’ and U2’ are already operative during heat-up. Furthermore, the amplitudes of the experimental DSC-signals are consistent with expectations: as reaction U1’ generates Ni and LNi respectively, its DSC signals should increase with growing Ni-content as can be inferred from Figs. 11 b-c, and vice-versa the intensity of the Cu and LCu DSC-signals corresponding to reaction U’2 decrease with the Ni-content. Taking the mean of the heating and cooling cycle as an indicator for the transformation temperature, the peak in the interval around 1210-1220 °C coincides well with THERMOCALC’s prediction of 1214 °C for the temperature of invariant reaction U’2. However, the DSC signals in the interval 1290-1300 °C (Figs. 11 b-d), which we attribute to invariant reaction U’1, deviate by 30 K from THERMOCALC’s prediction of 1323 °C. The reasons for that inconsistency are as yet not fully clear. We believe that Yan et al. [20], relying on optical microscopy only, misinterpreted the microstructure in their 1300 °C samples by identifying a 3-phase mixture containing (bcc + fcc + liquid). In contrast, our DSC results suggest a (bcc + liquid(1) + liquid(2)) phase composition at 1300 °C. Nevertheless, regardless of the real U’1 temperature, the impact of this ambiguity on the results and conclusions presented above is negligible. 4.5. Grain Size and Shape Coarsening of the Mo grains during sintering is significant. An example of the evolution of the size and shape of primary Mo grains during sintering at 1400 °C is shown in Fig. 12. Attaining a size of 9 μm after 4h of sintering, as compared to the initial size of the Mo powder (3 µm), grain growth in alloy 2N is pronounced. Yet, it is considerably slower than in alloys comprising LNi (c.f. also Figs. 6 and 7). While the primary Mo grains in 6N have grown to 25 µm, the final grain sizes in both alloys 10N and 14N reach 38 µm, i.e. grain growth in 10N and 14N is nearly identical. The size of the primary Mo grains, G, follows the relationship (1) where t is the isothermal sintering time and G0 denotes the size of primary Mo grains at the onset of isothermal sintering (Fig. 12a). Notably, G0 is definitely larger than the initial Mopowder size, suggesting that Mo grains coarsen considerably before the isothermal sintering
page - 9 -
temperature has been reached. However, G0 is significantly smaller in alloy 2N (5 µm) than in the Ni-rich alloys (10 µm). We attribute this difference to the coarsening that occurs in the temperature interval after the first liquid has formed and before isothermal sintering begins. The absence of LNi in alloy 2N retards grain growth and densification, contrary to the situation in alloys 6N-14N. Finally, the t1/3 time-dependence suggests that growth of the Mo grains is diffusion controlled [14, 23]. While the primary Mo crystals in the high Ni-alloys (6N-14N) are close to spherical (89% < S < 93%), they are less rounded in 2N (S 82%) (Fig. 12b, c.f. also Figs. 6 and 7). This indicates again that solution-reprecipitation in the high Ni-alloys is faster due to the high solubility of LNi for Mo. The Mo grains in “LCu-only” alloys being smaller, more facetted and less spherical, suggest that solid state sintering during heat-up plays a major role in densification. More details on the evolution of the microstructure in Mo-Cu-Ni alloys and on grain growth will be discussed in a companion paper [23]. 5. Concluding Remarks Two major conclusions can be drawn from this research. Firstly, the Mo-Cu-Ni system fulfills all criteria for liquid phase sintering, i.e. chemical compatibility, solubility and wetting by the liquid binder. Depending on the Ni-content, liquid phase sintering in Mo-Cu-Ni takes place in the presence of one or two liquid phases. Cu-rich alloys comprise a single-phase Cu-rich liquid, LCu, during sintering and, due to its low Mo-solubility, sinter slowly. In contrast, the Nirich alloys sinter in the presence of two melts, LCu and LNi, where the latter liquid promotes densification significantly due to its high solubility for Mo. In both cases, LPS generates a fully dense microstructure with equiaxed Mo particles embedded into a ductile fcc (Cu,Ni)-binder matrix, provided sintering conditions are selected which suppress the formation of the brittle MoNi phase. This requires (i) sintering above the U2’ temperature to dissolve that has possibly formed during heat-up and (ii) subsequent cooling at sufficiently fast rates to suppress its reprecipitation. Verified by sintering and DSC experiments as well as by microstructural and microanalytical observations, cooling rates on the order of 10 K/min ensure fully dense composites which comprise only ductile microstructural constituents. The non-equilibrium phase diagrams predict a miscibility gap not only in the liquid but also in the solid state of Mo-Cu-Ni, which explains the observed dual-phase microstructure in the binder of high Ni-alloys comprising a Cu- and a Ni-rich fcc (Cu,Ni) solid solution, Cu and Ni, respectively. They originate essentially from the two immiscible liquids, LCu and LNi, which solidify via reactions U1’ and U2’ into Cu, Ni and small Mo-precipitates. The existence of the two invariant reactions, U1’ and U2’, has been confirmed by DSC experiments. Densification and grain growth during LPS are largely governed by the Mo-flux during the solution-reprecipitation process. The Mo-transportation is controlled by the solubility of Mo in the liquid phases. As LCu and LNi comprise low and high solubility for Mo respectively, their actual volume fractions, which can be adjusted by variation of the Ni-concentration of the binder, dictate the overall Mo-transportation capacity during sintering. Thus the Ni-content is an efficient tool to control densification and grain growth in LPS-Mo.
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Finally, this investigation has generated insight into the densification and grain growth of solids which sinter in the presence of two liquid phases. On one hand those liquids generate a dual-phase binder, as compared to the single phase binders of conventional LPS systems. On the other hand they offer easy control of the solubility of the major phase during LPS. While the solubility in single liquid systems can be essentially affected by temperature only, sintering with two liquids permits the solubility of the major phase to be adjusted by selecting appropriate volume fractions of the two liquids.
Acknowledgements We would like to thank Plansee SE for the permission to publish this work. Valuable discussions with B. Kieback, H. Kestler and H.J. Seifert are gratefully acknowledged. Declaration of interests
☒ The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. ☐The authors declare the following financial interests/personal relationships which may be considered as potential competing interests:
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M.V. Sundaram, K.B. Surreddi, E. Hryha, A. Veiga, S. Berg, F. Castro, L. Nyborg, Enhanced Densification of PM Steels by Liquid Phase Sintering with a Boron-Containing Master Alloy, Metall. and Mat. Trans. A 49A, (2018), 255-263
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(16) K.S. Hwang, H.S. Huang, Identification of the segregation layer and its effects on the activated sintering and ductility of Ni-doped molybdenum, Acta Mater. 51, (2003), 3915-26 (17) K.S. Hwang, H.S. Huang, Ductility improvement of Ni-added molybdenum compacts through the addition of Cu and Fe powders, Int. J. Ref. Met. Hard Mat. 22, (2004), 185191 (18) A. Bose, R.M. German, Sintering Atmosphere Effects on Tensile Properties of Tungsten Heavy Alloys, Met. Trans. A 19A, (1988), 2467-2476 (19) Thermo-Calc Software AB, Description of the TCNI.8 database of THERMOCALC, url: http://www.thermocalc.com/media/23650/tcni8_extended_info.pdf, last visited: 05.12.2018 (20) Y. Yan, W. Cuiping, L. Yifu, L. Xingjun, K. Ryosuke, I. Kiyohito, Experimental investigation and thermodynamic calculation of the phase equilibria in the Cu-Mo-Ni ternary system, Materials Chemistry and Physics 125, (2011), 37-45 (21) C.B. Shoemaker, D.P. Shoemaker, The crystal structure of the -phase MoNi, Acta Cryst. 16, (1963), 997-1009 (22) K.P. Mingard, B. Roebuck, J. Marshall, G. Sweetman, Some aspects of the structure of cobalt and nickel binder phases in hardmetals, Acta Mater. 59, (2011), 2277-2290 (23) A. Huber, L.S. Sigl, Grain Growth in Liquid Phase Sintered Molybdenum, to be published
Tables Table 1: Properties of the Mo, Ni, and Cu powders. Powder
Mo
Cu
Ni
Particle size D10, µm
1.2
3.4
1.5
Particle size D50, µm
2.0
4.7
5.2
Particle size D90, µm
3.4
6.5
12.5
Oxygen content, µg/g
345
1,451
2,865
Carbon content, µg/g
15
637
227
Nitrogen content, µg/g
19
10
17
Global Tungsten and Powders
Ecka Granules
Eurotungstene
Supplier
Table 2: Calculated compositions of the liquid phases LCu and LNi at 1400 °C.
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Composition in at.% at 1400 °C
Alloys 2N
6N, 10N, 14N
LCu
LCu
LNi
Mo
0.4
1.1
28.3
Cu
89.6
71.0
9.7
Ni
10.0
27.9
62.0
Table 3: Composition of the microstructural constituents in Fig. 7 (EDX measurements). Alloy 2N 6N 10N
14N
Position #
Microstructural Constituent
Cu at.%
Ni at.%
Mo at.%
1
Cu
89.1
10.9
0.0
2
primary Mo
0.0
0.4
99.6
3
Cu
69.2
29.6
1.2
4
Ni
16.7
65.9
17.4
5
Cu
66.3
32.6
1.1
6
Ni
19.5
65.2
15.3
7
Cu
59.9
38.8
1.3
8
Ni
18.4
66.5
15.1
9
secondary Mo
3.0
9.6
87.4
Figures
Fig. 1 Calculated isothermal section of Mo-Cu-Ni at 1400 °C with the four investigated alloys (solid red points) at 82 wt.% Mo (red solid line) and the vertical sections at constant Ni/Cu ratios (red dashed lines). The three-phase field with Mo and the two liquids LCu and LNi is shown in green. page - 14 -
2a
2b
Fig. 2: Calculated isopleth sections of the Mo-Cu-Ni system at 82 wt.% Mo showing the four investigated compositions 2N-14N (dashed vertical lines): (a) equilibrium diagram with three invariant reactions U1-U3 (b) non-equilibrium diagram with two invariant reactions U1’ and U2’.
3a
3b
3c
3d
Fig. 3: Calculated isopleth sections along the dashed red lines of Fig. 1 at constant Cu:Ni ratios of a) 16:2, b) 12:6, c) 8:10 and d) 4:14 showing the evolution of the liquid miscibility gap in Mo-Cu-Ni.
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Fig. 4: Calculated volume fractions of liquid phases at 1400 °C
5a
5b
Fig. 5: As-sintered densities of alloys 2N (a) and 6N (b) as a function of sintering temperature and soaking time.
6a)
6b)
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6c)
6d)
Fig. 6: Light microscopy images of alloys 2N (a), 6N (b), 10N (c) and 14N (d) sintered at 1400 °C for 4h with three phases: Mo (light grey), Cu (white),Ni (dark grey).
7a)
7b)
7c)
7d)
Fig. 7: Secondary electron images (SEM) of alloys 2N (a), 6N (b), 10N (c) and 14N (d) sintered at 1400 °C for 4h. Mo is white, Cu is dark grey, andNi is medium grey. The red circles point to the position of the EDX measurements documented in Table 3.
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8a
8b
8c
8d
8e Fig. 8 Composition and phase analysis of alloy 14N: (a-c) EDX-maps of Mo, Ni and Cu, (d) EBSD-map, (e) SE image.
9a
9b
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Fig. 9: Primary Mo grains and Mo precipitates (white) in the binder of alloy 10 N: (a) secondary Mo (Mosec.) surrounded by medium grey Ni binder-phase originating from invariant reaction U1’ (b) tertiary Mo surrounded by dark grey Cu binder-phase originating from invariant reaction U2’
Fig. 10: SEM micrograph (SE-image) of the δ-phase (point #1) in alloy 6NEDX composition of point #1 in at.%: Mo 47.7, Ni 49.5, Cu 2.9
11a)
11c)
11b)
11d)
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Fig. 11: DSC diagrams of Mo-Cu-Ni alloys a) 2N b) 6N, c) 10N and d) 14N
12a)
12b)
Fig. 12: Evolution of (a) grain size, G, and (b) grain shape, S, with isothermal soaking time, t, in alloys 2N-14N (sintering temperature 1400 °C). Grain growth follows a t1/3 law.
Graphical Abstract
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