Abstract
Using a combination of mixed elemental powders and TiB2, a series of Ti–Ni and Ti–Ni–B alloys were optimised for sintering by varying the nickel and boron contents, the particle size of the elemental powders and the compaction pressure. The sintering temperature was maintained at 1200°C to limit the costs of a potential commercial sintering operation. For Ti–Ni alloys, a density of 99% was attained in Ti–7Ni made using fine Ti and Ni powders sintered in the solid state, and from liquid phase sintering of Ti–8Ni made using coarser powders. Porosity was almost eliminated from Ti–7Ni–xB alloys made by adding 1–3%TiB2 to the coarser Ti and Ni powders. The action of TiB2 as a sintering aid is possibly owing to a combination of the formation of a small amount of liquid at the sintering temperature and the restriction of grain growth owing to the presence of TiB particles.
Introduction
Near net shape processing using powder metallurgy techniques is a low cost method of manufacturing components, especially in the case of titanium alloys. To obtain good mechanical properties, particularly dynamic properties including fracture toughness and fatigue strength, it is important to reduce porosity to a minimum. This paper presents some methods of minimising porosity in Ti–Ni alloys by manipulation of the composition of the alloy, the particle size and the compaction pressure and other processing variables. In keeping with the secondary goal of minimising cost, the basic press and sinter technique has been applied to mixed elemental powders; techniques such as hot isostatic pressing and extrusion were avoided. Ti–Ni alloys were chosen because of their high rate of densification during sintering.1– 3
Previous work has shown that it is possible to achieve almost full density using the press and sinter approach and via metal injection moulding. For commercial purity Ti, sintered density in excess of 95% has usually been achieved by employing powders with particle size of less than 40–45 μm4– 10 or a size distribution augmented with fine particles.11 In other cases, density above 95% has been reached by pressing at pressures of 800 MPa or higher. 4 12 4,12,13 Sintered density of above about 99% has been reported for particle size of less than about 10 μm 7 14 7,14,15 or compaction pressure >1000 MPa. 4 4,13 Softer and higher purity powder is advantageous for sintering to high density,4 possibly because of its higher green density.
Making alloying additions to commercially pure Ti has been observed to increase the sintered density, but the Ti and alloy powders are usually required to be fine if the sintered density over 95% is to be achieved.16 High density has been reached for additions of up to 10% of the fast diffusing elements Fe, Co and Ni. Majima et al. 8 obtained ∼95% sintered density for commercially pure Ti (<45 μm) and >99% by adding up to 10%Co (1·4 μm). The density was about 99% for Ti–Fe and Ti–Ni binary alloys (∼5 μm) and 95–98% for several other binary alloys.10 There have been other similar results for Ti–Fe. 15 17 15,17,18 A Ti–13Nb–13Zr alloy made from mixed elemental powders with particle sizes of 9–10 μm was sintered to density of 93–95% using sintering temperatures as high as 1600°C.19 The addition of Al powder to a fine Ti powder (both <40 μm) resulted in a sintered density of 97%, possibly because Al increased the green density.20 However, Liu et al. 15 found that additions of Al (∼15 μm) and Mo (∼5 μm) to an 8 μm Ti powder reduced the sintered density from 99 to ∼95%.
For Ti–6Al–4V and some similar alloys, the sintered density >95% has been reached by means of fine Ti particle size (33–45 and ∼25 μm), 21 21,22 high compaction pressure (1515–1760 MPa)23 or Ti particle size somewhat less than the usual size (<150 μm) in combination with high uniaxial compaction pressure (926 MPa),24 cold isostatic pressing (480 MPa)25 or high pressure, high sintering temperature and long sintering time (640–960 MPa, 1350°C and 4 h).26 The use of TiH2 powder is beneficial.27 The sintered density of above about 99% has been achieved in the following ways:
pressing fine Al–V master alloy powder (<20 μm) mixed with Ti to relatively high green density (>80%).28– 31 The best combination of particle sizes was reported as 2–15 μm for the master alloy and 44–105 μm for the Ti.29 A similar procedure was followed for the alloy SP-70032
using titanium hydride powder, instead of Ti, in combination with high pressure, high sintering temperature and long sintering time (640–960 MPa, 1350°C and 4 h),26 and avoiding coarse alloying powder.
Saito et al. 33 33,34 attributed sintered density of >99% in TiB/Ti–6·8Mo–4·2Fe–1·4Al–1·4V and other Mo bearing composites to the formation of a transient liquid owing to cosegregation of Mo and B at the grain boundaries. However, similar results were reported for TiB/Ti–6Al–4V without Mo.35 The particle size of the alloying powders was fine, with an average size of 10 μm or less, and the powder was compacted by cold isostatic pressing at 392 MPa. The presence of TiB particles in Ti alloys results in the refinement of the microstructure, improved mechanical properties and superplastic behaviour at elevated temperatures.36– 39
The objective of the work described below was to sinter to almost full density without requiring especially fine Ti particles or high compaction pressure, at a relatively low sintering temperature of 1200°C. After an initial examination of the effects of Ti and Ni particle sizes, the effects of compaction pressure and additions of TiB2 were explored.
Calculation of theoretical density
The theoretical density was initially estimated by40
use the phase diagram to calculate the mass fractions of the phases present (including the TiNi phase even though it is not stable at room temperature)
use the results of (i) and the published densities of intermetallic compounds listed in Table 1 to obtain ρth* as in equation (1). The phases are effectively the pure elements or line compounds for Ti–Ni at room temperature, so it is not necessary to calculate the density of off stoichiometric compositions or solid solutions. Similar calculations are also presented for Ti–Si, for comparative purposes only.
Density of equilibrium Ti–Ni and Ti–Si phases
*Calculated from lattice parameters.
Results are shown in Fig. 1. The difference between the theoretical density of elemental mixtures and that of a mixture of the equilibrium phases is greater for Ti–Si than for Ti–Ni. To maintain continuity with results presented in previous publications, the relative density of sintered alloys reported below is with respect to the theoretical density of the corresponding elemental mixture ρth rather than ρth*. The values of ρth* relative to ρth are 100·34% for Ti–5Ni, 100·41% for Ti–6Ni, 100·48% for Ti–7Ni, 100·55% for Ti–8Ni and 100·62% for Ti–9Ni.
Comparison of room temperature theoretical density of Ti–Ni and Ti–Si alloys calculated for mixture of elemental particles and mixture of phases according to equilibrium phase diagram
Experimental
Experiments were carried out with the aim of minimising the porosity of alloys prepared using the mixed elemental approach. A moderate sintering temperature of 1200°C was selected, and other process variables including alloy composition, particle size and compaction pressure were varied to determine their effect on the final porosity. The three main experiments were as follows:
Ti powders with three different particle sizes mixed with Ni powders of two different sizes to form binary alloys with 7, 8 or 9%Ni, compacted at 400 MPa and sintered
Ti, Ni and TiB2 powders mixed to produce alloys with up to 7%Ni and 1%TiB2, compacted at 400 MPa and sintered (TiB2 was initially added as a potential grain refiner33)
Ti, Ni and TiB2 powders mixed to produce alloys with 7%Ni and up to 3%TiB2, compacted at 200–800 MPa and sintered.
The experiments were carried out using three types of Ti powders: −100 mesh (<150 μm) hydride-mill-dehydride Ti powders supplied by CERAC Incorporated (Milwaukee, WI, USA) and Sumitomo Titanium Corporation (TSP-100 grade, Tokyo, Japan) and <20 μm powder supplied by Atlantic Equipment Engineers (AEE, Bergenfield, NJ, USA). Particle size distributions were determined using a Malvern Mastersizer 2000 laser scattering instrument (Malvern Instruments Ltd, Malvern, UK). The chemical composition was determined using inductively coupled plasma atomic emission spectroscopy, atomic absorption spectroscopy and LECO combustion techniques (LECO Corporation, St. Joseph, MI, USA). The Ti powders were mixed with two types of Ni powders supplied by CERAC with particle sizes of −325 mesh and 5 μm average (both 99·9% pure), and 99·9% pure −325 mesh (<45 μm) TiB2 powder with Fisher median particle size of 7·2 μm supplied by Alfa Aesar (Ward Hill, MA, USA). The nominal particle sizes of the Ni powders are 30 and 5 μm. The methods of manufacturing Ti (<20 μm) and Ni powders are not known but based on the morphology as observed in the SEM, the Ti powder appears to be hydride-mill-dehydride powder produced from Ti sponge and the Ni powders appear to have been produced by the carbonyl process.
After mixing in a Turbula mixer for 30 min, the powders were cold pressed in a floating cylindrical die, at pressures of 200–800 MPa, to produce 10 mm diameter compacts ∼10 mm in length. Acrawax C lubricant was applied lightly to the die wall before it was filled with the required mass of powder. Compositions are reported as mass percentages. The green compacts were sintered in a vacuum tube furnace by heating at a rate of 4 K min−1 to a temperature of 1200°C, holding at this temperature for 2 h and cooling at a rate of 4 K min−1. The vacuum was maintained at about 1–10 MPa by an oil diffusion pump backed by a rotary vane pump. The compacts were placed on lightly sintered yttria supports in the furnace.
The characterisation of the sintered compacts was conducted using optical microscopy and SEM and density measurements. Cross-sections for metallography were polished using colloidal silica in 3% hydrogen peroxide and examined unetched or after etching in Kroll's etchant. The SEM was a Philips XL30 model (FEI Company, Hillsboro, OR, USA) equipped with EDX analysis capability. Green density was calculated from the mass and dimensions of the cylindrical compact. Sintered density was measured using the Archimedes method according to ASTM Standard B 328 (MPIF Standard 42). The pore filling liquid was Mobil DTE25 oil (Exxon Mobil Corporation, Irving, TX, USA) with a specific gravity of 0·87 and the test liquid was a hydrofluoropolyether heat transfer fluid with a specific gravity of 1·68 (H-Galden ZT-180, Solvay S.A., Brussels, Belgium). Total porosity was obtained from the density of the alloy relative to the theoretical density ρth. Open porosity was determined in accordance with ASTM B 328, and the amount of closed porosity determined by difference.
Equilibrium phase diagram data were estimated from the diagrams attributed to Murray51 and calculations carried out using Thermo-Calc for Windows (TCW4) software from Thermo-Calc Software AB (Stockholm, Sweden) (2006) and the Ti alloy database from Thermo Tech Ltd (Thermotech Ltd, Guildford, United Kingdom).
Results and discussion
Characterisation of Ti powders
The particle size distributions of the Ti powders are summarised in Table 2. More complete descriptions of the same Sumitomo and AEE powders and a different batch of the CERAC powder are available elsewhere.52 The Sumitomo powder is finer than the CERAC powder and its size distribution has a fat tail at fine particle sizes (∼10 μm). The particle size of the AEE powder is slightly greater than the nominal 20 μm maximum. The chemical analysis results are listed in Table 3. The Sumitomo powder appears to have slightly fewer impurities than the CERAC powder. As might be expected from the fine particle size, the AEE powder has higher impurity content.
Particle size of Ti powders, μm
Composition of Ti powders, mass ppm
Experiment 1: Effect of Ti and Ni particle sizes
Relative to the CERAC Ti powder, the green density of Sumitomo Ti was slightly higher and that of AEE Ti somewhat less. This is attributable to the slightly higher purity (lower hardness) and broader size distribution of the Sumitomo powder, and the lower purity and finer particle size of the AEE powder. For all three Ti powders the green density was slightly higher when mixed with 5 μm Ni than with 30 μm Ni. The green density and the density after sintering at 1200°C for 2 h are shown in Fig. 2.
Geometric relative green density and Archimedes relative sintered density of Ti–7Ni, Ti–8Ni and Ti–9Ni prepared from CERAC −100 mesh Ti (C), Sumitomo −100 mesh Ti (S), AEE <20 μm Ti (A), CERAC −325 mesh Ni (30) and CERAC 5 μm Ni (5) powders
A maximum sintered density of ∼99% was achieved. Reaching this density for the Ti–7Ni composition required the use of fine Ti and fine Ni powders. However, for Ti–8Ni, and especially Ti–9Ni, 99% density was achieved using coarser powders. One of the main reasons for the increase in density with Ni content is that liquid phase sintering occurs for the higher Ni contents. As shown in Fig. 3, a small amount of material that had clearly been liquid during sintering was found at the base of all compacts with 9%Ni. A lesser amount was observed for 8%Ni in the Sumitomo and AEE Ti powders (but not CERAC Ti). No liquid appeared to have moved to the base of the compact for any of the compacts with 7%Ni. The Ti–Ni equilibrium phase diagram indicates that at 1200°C the solidus is at ∼8%Ni.
Macrographs of sintered compacts made from Sumitomo Ti mixed with 30 μm Ni powder showing material that had been liquid during sintering at base of compacts
Figure 4 shows that there is some distortion of the compact and interaction with the support material when there is a significant amount of liquid (Ti–9Ni at 1200°C). It is therefore preferable to operate at the lowest Ni content consistent with achieving the required density. This of course also minimises the amount of material lost from the compact as liquid if the liquid wets the material used to support the compact in the sintering furnace.
Macrographs of sintered compacts produced from Sumitomo Ti powder mixed with 30 μm Ni powder: compacts have similar appearance for other Ti powders and 5 μm Ni but spangles are not as prominent for 5 μm Ni
Determinations of the amounts of open and closed porosity in the compacts of all three experiments (sections on ‘Experiment 1: effect of Ti and Ni particle sizes’, ‘Experiment 2: TiB2 additions’ and ‘Experiment 3: compaction pressure and TiB2’) are plotted in Fig. 5. The results suggest that there is little or no open porosity in the Ti–Ni compacts with sintered density above about 92%. There appears to be a very sharp transition from open to closed porosity as the density increases from 86 to 92%.
Amounts of open and closed porosity for CERAC, Sumitomo or AEE Ti powders mixed with 7, 8 or 9% −325 mesh or 5 μm Ni powder pressed at 400 MPa and sintered at 1200°C for 2 h: additional data for Ti–Ni–TiB2 compositions added as open and filled circles
Sufficient data are available to estimate the statistical variation in green density, sintered density and open porosity measured for the three unalloyed Ti powders and CERAC Ti with 7% −325 mesh Ni. The results of the statistical analysis of compacts pressed at 400 MPa and sintered at 1200°C for 2 h in several different furnace runs are listed in Table 4. For other compositions the variation appears to be similar.
Estimated mean and standard deviation of green density, sintered density and open porosity of compacts pressed at 400 MPa and sintered at 1200°C for 2 h
Experiment 2: TiB2 additions
In this experiment, the relatively coarse CERAC Ti powder was mixed with −325 mesh (30 μm) Ni powder and Alfa Aesar TiB2 powder. The Ni content of the powders ranged from 0 to 7%, and the TiB2 content from 0 to 1%. The Ti–B equilibrium phase diagram indicates that only a very small amount of B dissolves in Ti, and that the intermetallic compound TiB is present during and after sintering. The motivation for the TiB2 addition was the expected reduction in grain growth during sintering owing to the presence of TiB particles in the microstructure.36– 39 Compaction and sintering conditions were the same as for the previous experiment (400 MPa, 1200°C for 2 h).
The results of measurements of green and sintered density are shown in Fig. 6. Adding a small amount of TiB2 to Ti–6Ni had little or no effect on the density. However, there was a noticeable increase in sintered density when up to 1%TiB2 was added to Ti–7Ni, possibly because a small amount of liquid phase formed in Ti–7Ni–TiB2 compacts. The liquid pooled at the base of the compacts in the same way as for the compacts shown in Fig. 3.
Effect of Ni and TiB2 on the green and Archimedes sintered density of CERAC −100 mesh Ti powder: the horizontal line is the average of the sintered density for the two Ti–7Ni compacts
Calculations performed by the Thermo-Calc TCW4 software indicate that at 1200°C, the Ti–7Ni composition is within the β Ti single phase field. Adding TiB2 to Ti–7Ni initially moves the alloy towards the (β Ti+liquid) field and then almost along the boundary between the (β+TiB) and the (β+TiB+liquid) fields as shown in Fig. 7. The junction between the fields shown in Fig. 7 is at 7·1%Ni and 0·026%B (0·08%TiB2). These results suggest that the liquid observed experimentally is due to the TiB2 itself rather than any impurities in the TiB2 powder.
Calculated Ti rich corner of 1200°C isothermal section of Ti–Ni–B ternary phase diagram: symbol ‘×’ indicates Ti–7Ni with 1, 2 and 3%TIB2
Open and closed porosity are plotted in Fig. 5 (almost no open porosity for density >92%, almost all open at 86% dense) and are individually identified in Fig. 8. The addition of TiB2 to Ti–7Ni had an effect on the sintered density, which is similar to the increase in Ni content to Ti–8Ni (cf. Figure 2 Figs. 2 and 6). The maximum sintered density was only ∼93%. Therefore, a finer Ti powder was used in the third experiment. The typical pore structure is shown in Fig. 9 and the distribution of TiB particles in Fig. 10.
Effect of Ni and TiB2 on the open and closed porosity of sintered CERAC −100 mesh Ti powder: results of two furnace runs with unalloyed Ti and Ti–7Ni compacts included in both runs; the horizontal line is the average of the total porosity for the two Ti–7Ni compacts
Optical micrograph of Ti–7Ni–1TiB2 pressed at 400 MPa and sintered at 1200°C for 2 h
Backscattered electron image of pores (black), TiB particles (dark grey), Ti solid solution (midgrey) and eutectoidal structure (light grey) in Ti–6Ni–1TiB2 pressed at 400 MPa and sintered at 1200°C for 2 h
Experiment 3: Compaction pressure and TiB2
In the third experiment, the Ti powder was changed to the finer Sumitomo −100 mesh powder, both Ni powders (30 and 5 μm) were used, the Ni content was fixed at 7%, the TiB2 content ranged from 0 to 3%, and the compaction pressure was changed to 200 or 800 rather than 400 MPa. Significantly higher sintered density was achieved compared with experiment 2. Density results are shown in Fig. 11 (results for Ti–7Ni with 0–3%TiB2 are compared with unalloyed Ti). Sintered density reached almost 100% relative to ρth for Ti–7Ni with 1–3%TiB2 after pressing at 800 MPa (99·7% for 1%TiB2 with both Ni powders, 99·8% for 3%TiB2 with the coarser Ni powder and 100·2% for 3%TiB2 with the finer Ni powder). The true density of Ti–7Ni was calculated as 100·5% of ρth in the section on ‘Calculation of theoretical density’. Compaction at 200 instead of 800 MPa resulted in only slightly lower sintered density. Owing to the increase in TiB2 content from 1 to 3% (0·2 to 0·4 percentage point), the increase in compaction pressure from 200 to 800 MPa (0·1 to 0·4 percentage point) or the reduction in Ni particle size from about 30 to about 5 μm (0·1 to 0·4 percentage point), the increases in sintered density are small. However, dynamic mechanical properties are very sensitive to small amounts of porosity.31,53– 55
Geometric green density and Archimedes sintered density of Sumitomo −100 mesh Ti powder (plotted at slightly negative TiB2 content) and Sumitomo −100 mesh Ti with 7%Ni powder and 0–3%TiB2 powder, pressed at 200 or 800 MPa and sintered at 1200°C for 2 h: also data for 400 MPa from experiment 1
The beneficial effect of TiB2 may be due to its grain refining effect but is possibly due to the presence of the optimal amount of liquid at the sintering temperature. As shown in Fig. 7, the amount of liquid is not expected to change as much with an additional 1% of TiB2 as with an increase in Ni content from 7 to 8%. As for CERAC Ti powder in experiment 2, some liquid pooled at the base of the compacts containing TiB2 but not the Ti–7Ni compacts. Macrographs are shown in Fig. 12 to illustrate the minimal distortion of sintered compacts and the pooling of liquid at the base. Despite the results shown in Fig. 7, there appeared to be less liquid for compacts with 3%TiB2 than for the compacts with 1%TiB2 shown in Fig. 12.
Macrographs of sintered Ti–7Ni and Ti–7Ni–TiB2 compacts made using Sumitomo Ti and ∼30 μm Ni powders pressed at 200 or 800 MPa
Optical micrographs of polished sections revealed uniformly dispersed porosity in the sintered Ti–7Ni–xTiB2 compacts, consistent with the density measurements. The size of the pores decreases with increasing compaction pressure and decreasing Ni particle size (i.e. with increasing density). Micrographs of unetched Ti–7Ni–3TiB2 compacts are presented in Fig. 13. In addition to the dark well rounded pores, the grey plate-like TiB phase is shown to be well dispersed.
Optical micrographs of Sumitomo −100 mesh Ti with 7%Ni and 3%TiB2, sintered at 1200°C for 2 h using two Ni particle sizes and two compaction pressures
There are at least three possible reasons for the beneficial effect conferred by TiB2 additions:
restriction of grain coarsening owing to the pinning effect of the well dispersed TiB particles
formation of a small amount of liquid
modification of grain boundary structure and behaviour owing to segregation of boron to the grain boundaries.39
Conclusion
Close to full density was achieved after sintering mixed elemental powders at the relatively low sintering temperature of 1200°C for 2 h, after compaction at 400 MPa. For binary Ti–Ni alloys, high density was achieved by solid state sintering of Ti–7Ni made using fine Ti (∼20 μm) and fine Ni (∼5 μm) powders. Liquid phase sintering of Ti–8Ni allowed almost full density to be reached using coarser powders: <150 μm Ti and ∼30 μm Ni. Higher Ni content resulted in more liquid, more distortion and more reaction with the yttria used to support compacts in the sintering furnace.
Close to full density was also achieved for Ti–Ni–B alloys prepared by sintering <150 μm Ti powder mixed with 7%Ni powder and 2%TiB2 powder. Smaller pores and slightly higher density were obtained for finer Ni powder (5 versus 30 μm), higher compaction pressure (800 versus 200 MPa) and higher TiB2 content (3 versus 1%). It is anticipated that even such small differences in porosity may have a significant effect on mechanical properties such as ductility55 and fatigue strength.28
Footnotes
Acknowledgements
This research was funded by the Australian Research Council through the Centre of Excellence for Design in Light Metals. The authors thank Cheryl Berquist for some of the experimental data.