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Effect of alloying time and composition on the mechanical properties of Ti alloy
Materials Science and Engineering A 394 (2005) 220–228
Effect of alloying time and composition on the mechanical properties of Ti alloy Margam Chandrasekaran∗ , Zhang Su Xia Singapore Institute of Manufacturing Technology, 71 Nanyang Drive, 638075 Singapore Received 19 November 2003; received in revised form 8 November 2004; accepted 12 November 2004
Abstract Titanium and its alloys find wide applicability in aerospace and medicine. The key restricting factor in using the Ti-based alloys is the difficulties associated with the processing of the alloy as they are forged at temperatures >800 ◦ C. In the current work, Ti was mechanically alloyed in a tumbler mixer with Al, Fe and Zr with a charge to ball ratio of 1:2. The mechanical properties such as tensile and compression strengths and the hardness profile were evaluated as per the ASTM/MPIF standards. SEM analysis of the fracture surfaces indicated that samples sintered in vacuum exhibited brittle failure while those sintered in argon exhibited ductile–brittle failure. The tensile strengths ranged from 300 to 850 MPa for the samples tested. The alloying time affects the fracture mode and the strength of the alloy. Strain in excess of 20% was obtainable during upsetting samples alloyed for 48 h which is comparable to the conventional Ti–6Al–4V alloy. © 2004 Elsevier B.V. All rights reserved. Keywords: Titanium alloy; Mechanical alloying; Tensile strength; Sintering atmosphere; Sintering time
1. Introduction The use of titanium and its alloys in the aerospace and biomedical industries is well known due to their high strength to weight ratio and excellent biocompatibility. The market for titanium alloys is fast expanding with applications ranging from aerospace, automobiles, and biomedical to even electronic applications such as hard disk bases or hand phone covers. The most common alloy currently being used is the Ti–6Al–4V because of its performance. However, despite the advantages of Ti alloys and their application potential, the growth in the application of Ti alloys has been rather slow. This is due to the difficulty associated with the processing of Ti alloys at room/warm temperature ranges. Another significant disadvantage associated with titanium alloys is that they become highly reactive at temperatures higher than 600 ◦ C and undergo rapid oxidation. Thus, Ti components are either ∗
Corresponding author. Tel.: +65 67932730; fax: +65 67925362. E-mail address: [email protected] (M. Chandrasekaran). 0921-5093/$ – see front matter © 2004 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2004.11.056
fabricated by investment casting or by forging to near net shape at elevated temperatures in a controlled environment. This contributes to the increase in processing cost and most of the research works on titanium are aimed at development of microstructure useful for easier processing of the Ti alloys. Forging and mechanical working Ti alloys occupy a principal position in processing of titanium alloy components owing to the properties achievable of alloys [1–5]. Due to high flow stress and limited elongation of titanium alloys at room temperature, titanium and its alloys are normally processed at elevated temperatures close to 900 ◦ C to enable forming the desired shape of the component. Moreover, Ti alloys are highly sensitive to strain rate during processing and require careful control of parameters to minimize hardening due to strain rate sensitivity. Both the requirements of controlled environment and strain rates during forming contribute to higher processing costs of Ti alloys. Titanium exhibits polymorphism at different temperatures and its room temperature phase (alpha) has a hexagonal closely packed (HCP) structure while the high temperature phase (beta) has body centered cubic (BCC) crystal structure [6]. The hexagonal closely packed
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structure has only three slip systems active at room temperature which affects the plastic deformation while the high temperature phase of BCC has about 48 slip systems thus aiding plastic deformation upon application of stress. Addition of certain alloying elements to titanium results in stabilization of the phase and the elements can be broadly classified in to alpha stabilizers and beta stabilizers. One of the ways to improve the processability is to use beta phase stabilizers to facilitate forging at lower temperatures. The elements used for stabilizing the particular phase can be further classified into ‘isomorphous’ and ‘eutectoid’ stabilizers depending on the crystal structure of the alloying elements and the type of resulting compounds upon cooling from an elevated temperature. Isomorphous stabilizers do not form brittle intermetallics as on slow cooling, the beta phase does not transform in to alpha Ti + alloying elements, but are expensive. On the other hand, the eutectoid stabilizers have limitation on the percentage that can be used. Typical isomorphous stabilizers include, molybdenum, vanadium, tungsten, tantalum while iron, chromium, nickel, cobalt, manganese are termed as eutectoid stabilizers. The critical minimum percentage for the eutectoid stabilizers required goes as low as 3.5% in iron to as high as 13% in case of copper [4]. Moreover, the addition of vanadium in titanium has been found to retard the bone mineralization and formation in patients with implants and the bone growth is found to be lower with the presence of V and Al. Vanadium has higher toxicity compared with Al on the biocompatibility chart [14]. The alloying elements were selected based on the earlier works on cold forgeability of Ti alloys [7–14]. In the current work, mechanical alloying of the powders was performed for different lengths of times to improve the processability of the Ti alloys at lower temperatures. The present work studies the effect of alloying time and composition on the properties of Ti–Al–Fe–Zr alloy.
2. Experimental methodology and materials used Pure elemental powders of Ti, Al, Fe, Zr were obtained from different sources and their size distribution and morphology are provided in Table 1. The alloying addition of Fe and Zr were kept lower to minimize their adverse effects on the creep resistance and oxidation resistance, respectively [6]. Fe has a higher interdiffusion coefficient in Ti and has a retained beta phase on cooling. Moreover, alloying of Fe contributes to increase in the proof strength of the alloy, but on the other hand, decreases the cryogenic toughness of the alloy [4,12,13].
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The flowability tests were conducted under normal atmosphere. The resistance to free flow in the Carney flow meter is due, possibly, to the low particle size and high relative humidity. The powders were weighed in their respective proportions and filled in a plastic container, which was filled with Al2 O3 balls with a charge to ball ratio of 1:2 (v/v) and were mechanically alloyed from 8 to 48 h using a tumbler mixer. After alloying, the containers were carefully opened to avoid any oxidation of the powders and the contents were sieved to remove the balls and collect the powder blend. About 4–7 g of powders was used to prepare tensile bars and pellets of 11 mm diameter × 10 mm height specimens. In order to ascertain the effect of sintering temperature and sintering atmosphere, samples alloyed for 8 h were sintered in vacuum for 2 h at 1150 and 1250 ◦ C, respectively. The optimum sintering conditions were identified based on the tensile strength and the strain at fracture besides the phases formed during sintering. The samples alloyed for 8 h were also sintered in argon at 1250 ◦ C. The optimum conditions were used for sintering of the compositions alloyed for different time durations and the properties such as tensile strength, strain at failure, fractograph and hardness were evaluated. The alloyed samples were compacted to form tensile specimens (MPIF 10) and pellets of 12 mm diameter × 9 mm height. The green pellets were used to countercheck the repeatability of results obtained from powder X-ray diffraction. The pellets and the tensile specimens were then sintered at respective conditions as described in the experimental methodology. The density of the alloyed powder, green pellets and the sintered compacts were identified using a pycnometer. The X-ray phase analysis was done using a Philips X-ray Diffractometer with 2θ values ranging from 20 to 120◦ . The reason for the wider scan angle is basically to identify any stray phases that might form in the new formulation. The tensile tests were performed on the sintered samples as per ASTM E8-03/MPIF 10 at room temperature while the compressive strengths were evaluated as per procedures in ASTM E9-89a. The hardness of the samples was evaluated using the Vickers micro hardness tester at a load of 300 g. A minimum of three samples were used for each of the tests to confirm the repeatability of the results within ±5% and the average values were plotted. One pellet sample in each condition was polished and etched with Keller’s reagent for approximately 30 s to reveal the grains and phases present. Scanning electron microscopy was performed on the fractured samples and the etched samples for identifying the fracture mode. EDX was performed on individual grains on samples to identify the elements present to be used as indicative evidence of spe-
Table 1 Morphology/size of the powders used for the study and the compositions studied Powders
Titanium
Aluminium
Iron
Zirconium
Size/morphology Flowability Alloy composition (wt.%)
−270 Mesh/irregular Non-flowing Bal.
−325 Mesh/rounded Non-flowing 2
−270 Mesh/irregular Non-flowing 1
−325 Mesh/rounded Non-flowing 1
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Fig. 1. (a and b) XRD spectrum of the composition mechanically alloyed for 8 h and sintered in vacuum and argon for 2 h at 1250 ◦ C, respectively. ( Ti and () elemental Fe.
cific allotropic form of the alloy besides having the evidence from X-ray spectrum.
3. Results and discussions 3.1. X-ray phase analysis X-ray phase analysis was carried out on the alloyed powder sample and the sintered samples to know the phase formations with the alloying time and sintering, respectively. Since some of the peaks were overlapping, a clear distinction of the element or alloy could not be made with three strong peaks. However, efforts were made to identify the elements with two strong peaks wherever such overlapping of such strong peaks occurred with other prominent element or compound. Besides this, the samples alloyed for 8 h and sintered in vacuum at 1150 and 1250 ◦ C was also evaluated to know the effect of sintering temperature and time on the phase formations. X-ray phase analysis was also done on alloyed powder samples, which indicated presence of elemental peaks of alloying elements for those samples alloyed for less than 32 h. However, samples alloyed for 32 h and above did not show any evidence of elemental peak presence in the samples. The XRD analysis was repeated on random samples to confirm this observation. Both the powder samples and those sintered in vacuum revealed presence of some of the elemental peaks
) Alpha
along with suppression of peak intensities. Samples sintered in argon did not show any evidence of trace elemental peak of alloying elements (Fig. 1). This suggested that the alloying was uniform due, possibly, to the better distribution of heat during sintering. The XRD spectrum of the samples mechanically alloyed for different time durations and sintered at 1250 ◦ C in argon were taken to know the effect of alloying duration on the phase formations. Two representative XRD spectrums are presented above for the effect of alloying duration on the phase formations. The higher duration of alloying possibly resulted in texturing in the alloy causing suppression of few peaks and increasing the intensity along specific lattice directions of the crystal structure. This can be clearly seen from the relative peak heights recorded for 8 h of alloying and that of 40 h of alloying. Alloying for 16 and 24 h and sintering in argon for 1250 ◦ C resulted in increase in peak intensities of certain peaks such as alpha Ti and formation of intermetallics of TiAl3 as observed from the XRD spectrums for the respective alloying/sintering conditions. The samples alloyed for 32 h and sintered in argon indicated the presence of both alpha Ti and intermetallics of TiAl3 and FeAl3 (Fig. 2). The peaks consists mainly the alpha Ti peaks and a beta peak in the case of samples alloyed for 40 h and sintered in argon along with peaks of ZrAl3.3 Fe1.3 . The Zr–Al–Fe peaks were insignificant compared to the base matrix peaks. Samples al-
Fig. 2. (a and b) XRD spectrum of the composition mechanically alloyed for 32 h and 40 h and sintered for 2 h in argon at 1250 ◦ C. ( beta Ti.
) Alpha Ti and (
)
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Fig. 3. Volumetric dimensional change with alloying time. A simple regression analysis was performed based on the data obtained for different alloying time which fitted in to a polynomial curve given by dimensional shrinkage = 35.76 − 2.119X + 0.1255X2 − 0.003140X3 + 0.00002725X4 , where X represents time (h), R2 = 0.89.
loyed for 48 h and sintered in argon showed alpha Ti peaks with intermetallics of Al–Ti and Al3 Ti0.7 Fe0.25 in the XRD spectrum. The XRD spectrums of the samples alloyed for 48 h also showed some unidentified peak in addition to the above. 3.2. Volumetric shrinkage The volumetric shrinkage of the samples alloyed for different durations and sintered was calculated from the change in dimensions. The volumetric dimensional change was the highest for samples alloyed for 8 h at 22–25% which reduced with the alloying time. The reduction in volumetric dimensional change was not significant for samples alloyed for 40 and 48 h (Fig. 3). The volumetric dimensional change observed for samples sintered in vacuum was lower (17–19%) compared to those samples sintered in argon suggesting that the densification and bonding was not significant. This is due, possibly, to the thermal diffusion giving rise to non-uniform heat distribution. The powders after mechanical alloying were analyzed for morphology under the SEM. Fig. 4 shows a representative image of the particle distribution and the morphology after 8 and 40 h of alloying. After 8 h of alloying, there were some isolated agglomerated particles in addition to elemental particles whose size had reduced from the initial value. The SEM image of the sample alloyed for 40 h showed irregular particles with evidence of alloying on the surface. The particle size distribution was found to be more or less uniform compared to the blend alloyed for 8 h. The powder samples
from each of the alloy blend prepared were then compacted using a cylindrical pelletizing die and a tensile die set. The dimensions of the green compacts were measured and the blend alloyed for 8 h were sintered in vacuum at a sintering temperature of 1150 and 1250 ◦ C, respectively. Sintering temperature of 1250 ◦ C was selected based on the previous study on Tibased alloys [1]. The samples so sintered were analyzed for phases that form and the tensile strength values. The XRD of the sample sintered at 1250 ◦ C for 2 h in vacuum indicated the formation of alpha Ti mainly with intermetallics of ZrAl3 and Ti Al. The tensile strengths recorded were low for the samples sintered at 1150 ◦ C and that of 1250 ◦ C at 280 ± 25 and 340 ± 20 MPa, respectively. The hardness of the samples on the Vickers scale ranged from 300 to 340 VHN for both sintering temperatures. The strain at fracture was low at 2.5–3% for the samples sintered at 1150 ◦ C and 3.5–4.5% for samples sintered at 1250 ◦ C in vacuum. Despite the higher hardness, the lower strength obtained for the samples sintered in vacuum was due, possibly, to the inherent porosity and brittle intermetallic phases. Since the sintering was done in vacuum, the convection component of heat transfer across the porous network is absent. Thus, the bonding between the adjacent particulates was not uniform throughout the sample core. This possibly resulted in a poorly bonded network of matrix materials with the intermetallic phases thereby weakening the samples. In the case of hardness testing, since the compressive load is applied on the surface instead of the bulk specimen, which is restricted to phases present on the surface, the hardness is measure over small area compared to the bulk samples in tensile testing.
Fig. 4. Powder blend after 8 and 40 h of alloying, respectively.
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Table 2 Effect of Sintering Temperature and atmosphere on the properties of the sample alloyed for 48 h Sintering temperature (◦ C)
Sintering time (min)
Strength (MPa)
Strain at fracture (%)
Atmosphere
1150 1250 1250
120 120 120
280 ± 25 340 ± 20 590 ± 12
3 ± 0.25 3.75 ± 0.2 7.5 ± 1
Vacuum Vacuum Argon
This possibly explains the discrepancy between the hardness values and the strength measured. The lower strength and strain values for samples sintered at 1150 ◦ C is due, possibly, to the temperature not being high enough to promote diffusion and bonding of particulates to initiate recrystallization and grain growth compared to those sintered at 1250 ◦ C. To confirm this behavior, XRD was performed on both samples sintered at 1150 and 1250 ◦ C, respectively. The powder blend was then compacted and sintered in argon atmosphere of 1250 ◦ C to know the effect of the sintering atmosphere. The sample had strength of approximately 590 ± 12 MPa. Minimum three samples were tested in each condition to verify the repeatability of the results. Based on the above, the sintering conditions were fixed at 1250 ◦ C in argon for 2 h. The powders were mixed in the respective proportions and were mechanically alloyed with a charge to ball ratio of 1:2 by volume from times varying from 16 to 48 h in steps of 8 h to know the effect of alloying time on the phase formation, tensile strength and hardness. The average particle size ranged from 30 ± 10 m. The distribution was ascertained by optical microscopic examination of the alloy blends. The density of the alloyed powders was measured using a pycnometer. The density was found to decrease with increase in alloying time. The initial density of the elemental powders was also evaluated using the pycnometer and the obtained values were used to calculate the alloy density by the rule of mixtures. The values of the densities are given in the graph below, which shows a decreasing trend with increase in alloying time. The initial increase in the density for alloying time of 16 and 24 h is due, possibly, to non-uniform distribution of elemental powders in the blend and the lot selected for density measurement. However, with increase in alloying time, the density decreases which is due, possibly, to the uniformity in the alloying and most of the elemental powders contributing to the alloy/intermetallic formation (Fig. 5). This is confirmed by the absence of elemental peaks/reduction in intensity of
Fig. 5. Density variation with alloying time. Density variation with the alloying time was fitted to a non-linear regression curve fitting and the equation obtained was compared with the experimental values. Typical equation used for the prediction of density was: density = 4.480 + 005486X − 0002414X2 + 0.000002727X3 , where X represents the alloying time (h), R2 = 0.973.
elemental peaks for those powder samples alloyed for higher time. Table 2 shows the strength values obtained for different temperatures of sintering and that of atmospheres. It can be clearly seen that both the density and the volumetric shrinkage follows a third/fourth-order polynomial instead of a linear fit. The length of mixing/alloying has a strong influence on the porosity of the compact which possibly decides the shrinkage by controlling the pore morphology and distribution. It can be clearly seen that the sintering temperature and the atmosphere has profound influence both on the strength and the strain at failure. Fig. 6 shows the microstructure of the sample sintered in vacuum at 1250 ◦ C and the corresponding fracture surface. It can be seen that the microstructure shows predominantly an alpha grain boundaries with porosities. The poorly bonded secondary phases of ZrAl3 weaken the matrix in addition to the inherent porosity which possibly causes poorer distribution of heat across the samples. This is due, possibly, to the
Fig. 6. (a and b) Microstructure and fracture surface of the sample sintered in vacuum.
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Fig. 7. (a and b) Model of heat flow path through pore with different sintering atmosphere.
fact that only conductive heat transfer across the sample is active in vacuum as compared to convection and conduction heat transfers in argon atmosphere. The fracture surface shows clearly a cleavage fracture with sharp features. The facets that are found on the fracture surface are possibly due to separation of slip bands under the action of tensile stress normal to the slip plane. This is due, possibly, to the inhomogeneous distribution of slip systems and microstructure resulting in weaker regions of off-loading stress [15]. The sample sintered in vacuum had a lower strength compared to the samples sintered in argon. An A simple model described below can depict this effect (Fig. 7). The heat flow is affected by the convection coefficient of the gas in addition to the conductivity of the Ti alloy thus assisting more uniform distribution of the heat across the cross section leading to uniform phase distribution. Figs. 8 and 9 show the strain and the strengths recorded at room temperature for the samples alloyed between 8 and 48 h. The average strength of the sintered compacts alloyed for different time durations varied from 600 MPa for 8 h alloying to about 280 MPa for 32 h of alloying which increased to close to 600 MPa on further increase in alloying time. In order to confirm the lowering of strength for samples alloyed for 32 h and sintered tests were repeated keeping all conditions identical and the variation obtained was within ±20 MPa indicating that 32 h indeed exhibited different behaviour. The XRD spectrum of the samples alloyed for 32 h and sintered were analyzed to check for the anomalous behaviour of the samples. It was found that the elemental powers upon me-
chanical alloying an alpha phase matrix with widmanstratten colonies. This formation is typical of beta Ti alloys especially with metastable intermetallic precipitates in the alpha grain before stable alloys and compounds form during the transition to a stable alloy/intermetallic. Moreover, the orientation of widmanstratten colonies also contribute to the difference in behavior and lower strengths experienced in the specimens alloyed for 32 h and sintered [16–18]. The higher strength for lower time periods of alloying was due, possibly, to the minimal alloying effect confirmed by the presence of elemental peaks in the X-ray analysis of the composition after respective alloying time durations. The strain at failure for the samples alloyed for 32 h also showed a decreasing trend compared to other alloying durations. This behavior was confirmed by repeating the alloying at least for four times and analyzing the specimen. Though a variation of ±20 to 40 MPa was recorded, but the behavior was not significantly different. The average strain at failure varied from as high as 10.8% for 40 h alloying condition to 6.7% for 32 h of alloying. The higher strain values and the strength after 40 h of alloying could be attributed to the microstructural development which has alpha grains with mostly grain boundary beta grains assisting is deformation. This possibly results in correspondence of alpha slip systems with beta grain slip systems [19]. Both the strength and the strain at failure showed a decreasing trend beyond 40 h alloying, i.e., 48 h of alloying. Fig. 10 shows the average hardness values of the samples with the alloying addition. The lowest hardness was recorded for samples alloyed for 8 h. The hardness also showed a typical trend, which was in conformity with the results obtained
Fig. 8. Effect of alloying time on the strength.
Fig. 9. Effect of alloying time on the strain.
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Fig. 10. Hardness changes with the alloying duration for the Ti alloy.
Fig. 11. Fracture surface of the sample alloyed for 8 h.
Fig. 14. Fracture surface of the sample alloyed for 32 h.
Fig. 15. Fracture surface of the sample alloyed for 40 h.
from the phase analysis and the tensile tests. The variation in hardness was minimum with alloying time up to 24 h of alloying. Further increase in alloying time increases the hardness by about 15–25 VHN. The increase in hardness with alloying time above 24 h is due, possibly, to the increased formation of intermetallics causing a solid solution strengthening effect. The fracture surfaces of the samples alloyed for different time durations and tested in the instron were examined using a scanning electron microscope. The fractographs are presented in Figs. 11–16.
The figures indicate a smooth transition from a ductile–brittle failure to a quasi cleavage fracture for 8 h of alloying to 32 h of alloying. This can be seen in the reduction in the dimple formation with increase in alloying time from 8 to 32 h (Figs. 11–14) and the predominance of sharp edges indicative of brittle failure. Fracture surface of the sample alloyed for 8 h, Fig. 11 shows equiaxed dimples with some trans-granular cleavage. While the fracture surface of specimens alloyed for 16 h, Fig. 12 indicates mixed cleavage fracture and shear rupture. Fracture surfaces of specimens of sample alloyed for 24 h, Fig. 13 has scattered distribution of dimples in addition to cleavage fracture and further increase to 32 h of alloying, Fig. 14 results in predominant cleavage fracture with no or little evidence of dimples. The fracture surface shows a completely different morphology as compared to the rest of the samples. Moreover, XRD of the samples alloyed for 32 h and sintered were repeated on multiple samples (at least three samples were prepared by alloying the powders individually at 32 h to confirm the anomalous behavior of the samples) from different alloying lots. This anomalous behavior of the samples at 32 h is due, possibly, to the non-
Fig. 13. Fracture surface of the sample alloyed for 24 h.
Fig. 16. Fracture surface of the sample alloyed for 48 h.
Fig. 12. Fracture surface of the sample alloyed for 16 h.
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Fig. 17. Microstructure of the sample alloyed for 8 h. Fig. 19. Microstructure of the sample alloyed for 40 h and sintered at 1250 ◦ C.
Table 3 Compressive strength and strain values during upsetting tests
Fig. 18. Microstructure of the sample alloyed for 32 h, sintered at 1250◦ .
uniform distribution of elemental powders in the blend resulting in a super saturated solution after 32 h of alloying to form intermetallics that are detrimental to the properties of the alloy. Further increase of alloying time to 40 h, Fig. 15 results in a combinational fracture mode (ductile–brittle) with deformation dimples seen across the fracture surface. Increasing to time 48 h of alloying, Fig. 16 results in reduction in dimple formation but with mixed mode failure mechanism similar to 40 h alloying. Further increase in alloying time possibly results in the breaking up of the intermetallics to facilitate stabilization of beta phase at the surface of the particulates yielding a alpha core with beta surface which on sintering produces grain boundary beta phase to aid deformation. This is confirmed by the X-ray phase analysis after sintering of the sample, which indicates absence of Fe–Al intermetallics for the samples alloyed for durations higher than 32 h. Figs. 17–19 show the representative microstructures of the samples alloyed for different time durations. It can be seen from the above microstructure that the surface has predominantly an alpha phase grains with few grain boundary phases. The microstructure of the sample alloyed for 32 h and sintered showed mainly alpha case with porosity along the grain boundary and intermetallics within and near the grain boundary. The weak intermetallics that form during alloying in combination with the porosity possibly contributed to low strains and strength values recorded for the samples. Fig. 19 shows the microstructure of the sample alloyed for 40 h duration and sintered. The microstructure showed a well-distributed alpha case with beta phase along alpha grain
Alloying time (h)
Strength recorded (MPa)
Strain values (%)
40 48
1039 ± 5 1070 ± 7
23 ± 0.5 18 ± 0.7
boundaries. Though porosity and intermetallics were seen on the grain boundaries but the deformation was facilitated possibly by the beta phase distribution resulting in higher strains before failure. Higher strengths compared to 32 h of alloying was also recorded due, possibly, to the beta phase arresting the crack propagation that was causing failure and expending it on increasing the strain. After ascertaining the strength values and the alloying conditions under which the elemental peaks disappear and the phases start to form, upsetting tests were done on those sample alloyed for 40 and 48 h to evaluate the workability of the alloy. Pellets of dimension of 12 mm diameter × 9 mm length were pressed, sintered and used for the above tests. Strength at the maximum load of 95 tonnes was recorded for the samples alloyed for 48 and 40 h and sintered at 1250 ◦ C, respectively. The results obtained from the upsetting tests are presented below in Table 3. The samples did not show a clear failure (the samples did not fail during compression tests as the limiting load of the equipment was possibly lower than the strength of the alloy tested at room temperature) when subjected to compression at the maximum load of the equipment (95 kN). The compressive tests were performed at a speed of 2.5 mm/min and were continued until failure or maximum load is reached depending on whichever occurs first. The results had a similar trend of that of the tensile results where the samples alloyed for 40 h and sintered had a strain value of 9.6 ± 1.2% at failure while the samples alloyed for 48 h had a strain value of 9.4 ± 0.6% at failure.
4. Conclusions From the above results, the following conclusions were derived:
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1. Alloying time has significant effect on the formation of intermetallics and hence the mechanical properties of the alloy blend. 2. Sintering atmosphere influences the phase formations in the alloy and therefore affects the tensile properties of the alloy. 3. The phase formations, microstructure and ratio of the phases formed affect the mechanical strength of the alloy. 4. Significant alloying could not be observed in samples alloyed for less than 24 h. 5. The fracture surfaces indicated a combination of ductile and brittle failure whose relative proportions varied with the alloying time. 6. Alloying time of 40 h resulted in optimum combination of properties with respect to strength and strain achievable.
Acknowledgments The authors are thankful to Singapore Institute of Manufacturing Technology for their support of this work by the way of in-house project (C01-P-150 AR). The help of Mr. Ho Meng Kwong on the compaction and sintering is acknowledged with thanks.
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Effect of alloying time and composition on the mechanical properties of Ti alloy Margam Chandrasekaran∗ , Zhang Su Xia Singapore Institute of Manufacturing Technology, 71 Nanyang Drive, 638075 Singapore Received 19 November 2003; received in revised form 8 November 2004; accepted 12 November 2004
Abstract Titanium and its alloys find wide applicability in aerospace and medicine. The key restricting factor in using the Ti-based alloys is the difficulties associated with the processing of the alloy as they are forged at temperatures >800 ◦ C. In the current work, Ti was mechanically alloyed in a tumbler mixer with Al, Fe and Zr with a charge to ball ratio of 1:2. The mechanical properties such as tensile and compression strengths and the hardness profile were evaluated as per the ASTM/MPIF standards. SEM analysis of the fracture surfaces indicated that samples sintered in vacuum exhibited brittle failure while those sintered in argon exhibited ductile–brittle failure. The tensile strengths ranged from 300 to 850 MPa for the samples tested. The alloying time affects the fracture mode and the strength of the alloy. Strain in excess of 20% was obtainable during upsetting samples alloyed for 48 h which is comparable to the conventional Ti–6Al–4V alloy. © 2004 Elsevier B.V. All rights reserved. Keywords: Titanium alloy; Mechanical alloying; Tensile strength; Sintering atmosphere; Sintering time
1. Introduction The use of titanium and its alloys in the aerospace and biomedical industries is well known due to their high strength to weight ratio and excellent biocompatibility. The market for titanium alloys is fast expanding with applications ranging from aerospace, automobiles, and biomedical to even electronic applications such as hard disk bases or hand phone covers. The most common alloy currently being used is the Ti–6Al–4V because of its performance. However, despite the advantages of Ti alloys and their application potential, the growth in the application of Ti alloys has been rather slow. This is due to the difficulty associated with the processing of Ti alloys at room/warm temperature ranges. Another significant disadvantage associated with titanium alloys is that they become highly reactive at temperatures higher than 600 ◦ C and undergo rapid oxidation. Thus, Ti components are either ∗
Corresponding author. Tel.: +65 67932730; fax: +65 67925362. E-mail address: [email protected] (M. Chandrasekaran). 0921-5093/$ – see front matter © 2004 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2004.11.056
fabricated by investment casting or by forging to near net shape at elevated temperatures in a controlled environment. This contributes to the increase in processing cost and most of the research works on titanium are aimed at development of microstructure useful for easier processing of the Ti alloys. Forging and mechanical working Ti alloys occupy a principal position in processing of titanium alloy components owing to the properties achievable of alloys [1–5]. Due to high flow stress and limited elongation of titanium alloys at room temperature, titanium and its alloys are normally processed at elevated temperatures close to 900 ◦ C to enable forming the desired shape of the component. Moreover, Ti alloys are highly sensitive to strain rate during processing and require careful control of parameters to minimize hardening due to strain rate sensitivity. Both the requirements of controlled environment and strain rates during forming contribute to higher processing costs of Ti alloys. Titanium exhibits polymorphism at different temperatures and its room temperature phase (alpha) has a hexagonal closely packed (HCP) structure while the high temperature phase (beta) has body centered cubic (BCC) crystal structure [6]. The hexagonal closely packed
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structure has only three slip systems active at room temperature which affects the plastic deformation while the high temperature phase of BCC has about 48 slip systems thus aiding plastic deformation upon application of stress. Addition of certain alloying elements to titanium results in stabilization of the phase and the elements can be broadly classified in to alpha stabilizers and beta stabilizers. One of the ways to improve the processability is to use beta phase stabilizers to facilitate forging at lower temperatures. The elements used for stabilizing the particular phase can be further classified into ‘isomorphous’ and ‘eutectoid’ stabilizers depending on the crystal structure of the alloying elements and the type of resulting compounds upon cooling from an elevated temperature. Isomorphous stabilizers do not form brittle intermetallics as on slow cooling, the beta phase does not transform in to alpha Ti + alloying elements, but are expensive. On the other hand, the eutectoid stabilizers have limitation on the percentage that can be used. Typical isomorphous stabilizers include, molybdenum, vanadium, tungsten, tantalum while iron, chromium, nickel, cobalt, manganese are termed as eutectoid stabilizers. The critical minimum percentage for the eutectoid stabilizers required goes as low as 3.5% in iron to as high as 13% in case of copper [4]. Moreover, the addition of vanadium in titanium has been found to retard the bone mineralization and formation in patients with implants and the bone growth is found to be lower with the presence of V and Al. Vanadium has higher toxicity compared with Al on the biocompatibility chart [14]. The alloying elements were selected based on the earlier works on cold forgeability of Ti alloys [7–14]. In the current work, mechanical alloying of the powders was performed for different lengths of times to improve the processability of the Ti alloys at lower temperatures. The present work studies the effect of alloying time and composition on the properties of Ti–Al–Fe–Zr alloy.
2. Experimental methodology and materials used Pure elemental powders of Ti, Al, Fe, Zr were obtained from different sources and their size distribution and morphology are provided in Table 1. The alloying addition of Fe and Zr were kept lower to minimize their adverse effects on the creep resistance and oxidation resistance, respectively [6]. Fe has a higher interdiffusion coefficient in Ti and has a retained beta phase on cooling. Moreover, alloying of Fe contributes to increase in the proof strength of the alloy, but on the other hand, decreases the cryogenic toughness of the alloy [4,12,13].
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The flowability tests were conducted under normal atmosphere. The resistance to free flow in the Carney flow meter is due, possibly, to the low particle size and high relative humidity. The powders were weighed in their respective proportions and filled in a plastic container, which was filled with Al2 O3 balls with a charge to ball ratio of 1:2 (v/v) and were mechanically alloyed from 8 to 48 h using a tumbler mixer. After alloying, the containers were carefully opened to avoid any oxidation of the powders and the contents were sieved to remove the balls and collect the powder blend. About 4–7 g of powders was used to prepare tensile bars and pellets of 11 mm diameter × 10 mm height specimens. In order to ascertain the effect of sintering temperature and sintering atmosphere, samples alloyed for 8 h were sintered in vacuum for 2 h at 1150 and 1250 ◦ C, respectively. The optimum sintering conditions were identified based on the tensile strength and the strain at fracture besides the phases formed during sintering. The samples alloyed for 8 h were also sintered in argon at 1250 ◦ C. The optimum conditions were used for sintering of the compositions alloyed for different time durations and the properties such as tensile strength, strain at failure, fractograph and hardness were evaluated. The alloyed samples were compacted to form tensile specimens (MPIF 10) and pellets of 12 mm diameter × 9 mm height. The green pellets were used to countercheck the repeatability of results obtained from powder X-ray diffraction. The pellets and the tensile specimens were then sintered at respective conditions as described in the experimental methodology. The density of the alloyed powder, green pellets and the sintered compacts were identified using a pycnometer. The X-ray phase analysis was done using a Philips X-ray Diffractometer with 2θ values ranging from 20 to 120◦ . The reason for the wider scan angle is basically to identify any stray phases that might form in the new formulation. The tensile tests were performed on the sintered samples as per ASTM E8-03/MPIF 10 at room temperature while the compressive strengths were evaluated as per procedures in ASTM E9-89a. The hardness of the samples was evaluated using the Vickers micro hardness tester at a load of 300 g. A minimum of three samples were used for each of the tests to confirm the repeatability of the results within ±5% and the average values were plotted. One pellet sample in each condition was polished and etched with Keller’s reagent for approximately 30 s to reveal the grains and phases present. Scanning electron microscopy was performed on the fractured samples and the etched samples for identifying the fracture mode. EDX was performed on individual grains on samples to identify the elements present to be used as indicative evidence of spe-
Table 1 Morphology/size of the powders used for the study and the compositions studied Powders
Titanium
Aluminium
Iron
Zirconium
Size/morphology Flowability Alloy composition (wt.%)
−270 Mesh/irregular Non-flowing Bal.
−325 Mesh/rounded Non-flowing 2
−270 Mesh/irregular Non-flowing 1
−325 Mesh/rounded Non-flowing 1
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Fig. 1. (a and b) XRD spectrum of the composition mechanically alloyed for 8 h and sintered in vacuum and argon for 2 h at 1250 ◦ C, respectively. ( Ti and () elemental Fe.
cific allotropic form of the alloy besides having the evidence from X-ray spectrum.
3. Results and discussions 3.1. X-ray phase analysis X-ray phase analysis was carried out on the alloyed powder sample and the sintered samples to know the phase formations with the alloying time and sintering, respectively. Since some of the peaks were overlapping, a clear distinction of the element or alloy could not be made with three strong peaks. However, efforts were made to identify the elements with two strong peaks wherever such overlapping of such strong peaks occurred with other prominent element or compound. Besides this, the samples alloyed for 8 h and sintered in vacuum at 1150 and 1250 ◦ C was also evaluated to know the effect of sintering temperature and time on the phase formations. X-ray phase analysis was also done on alloyed powder samples, which indicated presence of elemental peaks of alloying elements for those samples alloyed for less than 32 h. However, samples alloyed for 32 h and above did not show any evidence of elemental peak presence in the samples. The XRD analysis was repeated on random samples to confirm this observation. Both the powder samples and those sintered in vacuum revealed presence of some of the elemental peaks
) Alpha
along with suppression of peak intensities. Samples sintered in argon did not show any evidence of trace elemental peak of alloying elements (Fig. 1). This suggested that the alloying was uniform due, possibly, to the better distribution of heat during sintering. The XRD spectrum of the samples mechanically alloyed for different time durations and sintered at 1250 ◦ C in argon were taken to know the effect of alloying duration on the phase formations. Two representative XRD spectrums are presented above for the effect of alloying duration on the phase formations. The higher duration of alloying possibly resulted in texturing in the alloy causing suppression of few peaks and increasing the intensity along specific lattice directions of the crystal structure. This can be clearly seen from the relative peak heights recorded for 8 h of alloying and that of 40 h of alloying. Alloying for 16 and 24 h and sintering in argon for 1250 ◦ C resulted in increase in peak intensities of certain peaks such as alpha Ti and formation of intermetallics of TiAl3 as observed from the XRD spectrums for the respective alloying/sintering conditions. The samples alloyed for 32 h and sintered in argon indicated the presence of both alpha Ti and intermetallics of TiAl3 and FeAl3 (Fig. 2). The peaks consists mainly the alpha Ti peaks and a beta peak in the case of samples alloyed for 40 h and sintered in argon along with peaks of ZrAl3.3 Fe1.3 . The Zr–Al–Fe peaks were insignificant compared to the base matrix peaks. Samples al-
Fig. 2. (a and b) XRD spectrum of the composition mechanically alloyed for 32 h and 40 h and sintered for 2 h in argon at 1250 ◦ C. ( beta Ti.
) Alpha Ti and (
)
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Fig. 3. Volumetric dimensional change with alloying time. A simple regression analysis was performed based on the data obtained for different alloying time which fitted in to a polynomial curve given by dimensional shrinkage = 35.76 − 2.119X + 0.1255X2 − 0.003140X3 + 0.00002725X4 , where X represents time (h), R2 = 0.89.
loyed for 48 h and sintered in argon showed alpha Ti peaks with intermetallics of Al–Ti and Al3 Ti0.7 Fe0.25 in the XRD spectrum. The XRD spectrums of the samples alloyed for 48 h also showed some unidentified peak in addition to the above. 3.2. Volumetric shrinkage The volumetric shrinkage of the samples alloyed for different durations and sintered was calculated from the change in dimensions. The volumetric dimensional change was the highest for samples alloyed for 8 h at 22–25% which reduced with the alloying time. The reduction in volumetric dimensional change was not significant for samples alloyed for 40 and 48 h (Fig. 3). The volumetric dimensional change observed for samples sintered in vacuum was lower (17–19%) compared to those samples sintered in argon suggesting that the densification and bonding was not significant. This is due, possibly, to the thermal diffusion giving rise to non-uniform heat distribution. The powders after mechanical alloying were analyzed for morphology under the SEM. Fig. 4 shows a representative image of the particle distribution and the morphology after 8 and 40 h of alloying. After 8 h of alloying, there were some isolated agglomerated particles in addition to elemental particles whose size had reduced from the initial value. The SEM image of the sample alloyed for 40 h showed irregular particles with evidence of alloying on the surface. The particle size distribution was found to be more or less uniform compared to the blend alloyed for 8 h. The powder samples
from each of the alloy blend prepared were then compacted using a cylindrical pelletizing die and a tensile die set. The dimensions of the green compacts were measured and the blend alloyed for 8 h were sintered in vacuum at a sintering temperature of 1150 and 1250 ◦ C, respectively. Sintering temperature of 1250 ◦ C was selected based on the previous study on Tibased alloys [1]. The samples so sintered were analyzed for phases that form and the tensile strength values. The XRD of the sample sintered at 1250 ◦ C for 2 h in vacuum indicated the formation of alpha Ti mainly with intermetallics of ZrAl3 and Ti Al. The tensile strengths recorded were low for the samples sintered at 1150 ◦ C and that of 1250 ◦ C at 280 ± 25 and 340 ± 20 MPa, respectively. The hardness of the samples on the Vickers scale ranged from 300 to 340 VHN for both sintering temperatures. The strain at fracture was low at 2.5–3% for the samples sintered at 1150 ◦ C and 3.5–4.5% for samples sintered at 1250 ◦ C in vacuum. Despite the higher hardness, the lower strength obtained for the samples sintered in vacuum was due, possibly, to the inherent porosity and brittle intermetallic phases. Since the sintering was done in vacuum, the convection component of heat transfer across the porous network is absent. Thus, the bonding between the adjacent particulates was not uniform throughout the sample core. This possibly resulted in a poorly bonded network of matrix materials with the intermetallic phases thereby weakening the samples. In the case of hardness testing, since the compressive load is applied on the surface instead of the bulk specimen, which is restricted to phases present on the surface, the hardness is measure over small area compared to the bulk samples in tensile testing.
Fig. 4. Powder blend after 8 and 40 h of alloying, respectively.
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Table 2 Effect of Sintering Temperature and atmosphere on the properties of the sample alloyed for 48 h Sintering temperature (◦ C)
Sintering time (min)
Strength (MPa)
Strain at fracture (%)
Atmosphere
1150 1250 1250
120 120 120
280 ± 25 340 ± 20 590 ± 12
3 ± 0.25 3.75 ± 0.2 7.5 ± 1
Vacuum Vacuum Argon
This possibly explains the discrepancy between the hardness values and the strength measured. The lower strength and strain values for samples sintered at 1150 ◦ C is due, possibly, to the temperature not being high enough to promote diffusion and bonding of particulates to initiate recrystallization and grain growth compared to those sintered at 1250 ◦ C. To confirm this behavior, XRD was performed on both samples sintered at 1150 and 1250 ◦ C, respectively. The powder blend was then compacted and sintered in argon atmosphere of 1250 ◦ C to know the effect of the sintering atmosphere. The sample had strength of approximately 590 ± 12 MPa. Minimum three samples were tested in each condition to verify the repeatability of the results. Based on the above, the sintering conditions were fixed at 1250 ◦ C in argon for 2 h. The powders were mixed in the respective proportions and were mechanically alloyed with a charge to ball ratio of 1:2 by volume from times varying from 16 to 48 h in steps of 8 h to know the effect of alloying time on the phase formation, tensile strength and hardness. The average particle size ranged from 30 ± 10 m. The distribution was ascertained by optical microscopic examination of the alloy blends. The density of the alloyed powders was measured using a pycnometer. The density was found to decrease with increase in alloying time. The initial density of the elemental powders was also evaluated using the pycnometer and the obtained values were used to calculate the alloy density by the rule of mixtures. The values of the densities are given in the graph below, which shows a decreasing trend with increase in alloying time. The initial increase in the density for alloying time of 16 and 24 h is due, possibly, to non-uniform distribution of elemental powders in the blend and the lot selected for density measurement. However, with increase in alloying time, the density decreases which is due, possibly, to the uniformity in the alloying and most of the elemental powders contributing to the alloy/intermetallic formation (Fig. 5). This is confirmed by the absence of elemental peaks/reduction in intensity of
Fig. 5. Density variation with alloying time. Density variation with the alloying time was fitted to a non-linear regression curve fitting and the equation obtained was compared with the experimental values. Typical equation used for the prediction of density was: density = 4.480 + 005486X − 0002414X2 + 0.000002727X3 , where X represents the alloying time (h), R2 = 0.973.
elemental peaks for those powder samples alloyed for higher time. Table 2 shows the strength values obtained for different temperatures of sintering and that of atmospheres. It can be clearly seen that both the density and the volumetric shrinkage follows a third/fourth-order polynomial instead of a linear fit. The length of mixing/alloying has a strong influence on the porosity of the compact which possibly decides the shrinkage by controlling the pore morphology and distribution. It can be clearly seen that the sintering temperature and the atmosphere has profound influence both on the strength and the strain at failure. Fig. 6 shows the microstructure of the sample sintered in vacuum at 1250 ◦ C and the corresponding fracture surface. It can be seen that the microstructure shows predominantly an alpha grain boundaries with porosities. The poorly bonded secondary phases of ZrAl3 weaken the matrix in addition to the inherent porosity which possibly causes poorer distribution of heat across the samples. This is due, possibly, to the
Fig. 6. (a and b) Microstructure and fracture surface of the sample sintered in vacuum.
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Fig. 7. (a and b) Model of heat flow path through pore with different sintering atmosphere.
fact that only conductive heat transfer across the sample is active in vacuum as compared to convection and conduction heat transfers in argon atmosphere. The fracture surface shows clearly a cleavage fracture with sharp features. The facets that are found on the fracture surface are possibly due to separation of slip bands under the action of tensile stress normal to the slip plane. This is due, possibly, to the inhomogeneous distribution of slip systems and microstructure resulting in weaker regions of off-loading stress [15]. The sample sintered in vacuum had a lower strength compared to the samples sintered in argon. An A simple model described below can depict this effect (Fig. 7). The heat flow is affected by the convection coefficient of the gas in addition to the conductivity of the Ti alloy thus assisting more uniform distribution of the heat across the cross section leading to uniform phase distribution. Figs. 8 and 9 show the strain and the strengths recorded at room temperature for the samples alloyed between 8 and 48 h. The average strength of the sintered compacts alloyed for different time durations varied from 600 MPa for 8 h alloying to about 280 MPa for 32 h of alloying which increased to close to 600 MPa on further increase in alloying time. In order to confirm the lowering of strength for samples alloyed for 32 h and sintered tests were repeated keeping all conditions identical and the variation obtained was within ±20 MPa indicating that 32 h indeed exhibited different behaviour. The XRD spectrum of the samples alloyed for 32 h and sintered were analyzed to check for the anomalous behaviour of the samples. It was found that the elemental powers upon me-
chanical alloying an alpha phase matrix with widmanstratten colonies. This formation is typical of beta Ti alloys especially with metastable intermetallic precipitates in the alpha grain before stable alloys and compounds form during the transition to a stable alloy/intermetallic. Moreover, the orientation of widmanstratten colonies also contribute to the difference in behavior and lower strengths experienced in the specimens alloyed for 32 h and sintered [16–18]. The higher strength for lower time periods of alloying was due, possibly, to the minimal alloying effect confirmed by the presence of elemental peaks in the X-ray analysis of the composition after respective alloying time durations. The strain at failure for the samples alloyed for 32 h also showed a decreasing trend compared to other alloying durations. This behavior was confirmed by repeating the alloying at least for four times and analyzing the specimen. Though a variation of ±20 to 40 MPa was recorded, but the behavior was not significantly different. The average strain at failure varied from as high as 10.8% for 40 h alloying condition to 6.7% for 32 h of alloying. The higher strain values and the strength after 40 h of alloying could be attributed to the microstructural development which has alpha grains with mostly grain boundary beta grains assisting is deformation. This possibly results in correspondence of alpha slip systems with beta grain slip systems [19]. Both the strength and the strain at failure showed a decreasing trend beyond 40 h alloying, i.e., 48 h of alloying. Fig. 10 shows the average hardness values of the samples with the alloying addition. The lowest hardness was recorded for samples alloyed for 8 h. The hardness also showed a typical trend, which was in conformity with the results obtained
Fig. 8. Effect of alloying time on the strength.
Fig. 9. Effect of alloying time on the strain.
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Fig. 10. Hardness changes with the alloying duration for the Ti alloy.
Fig. 11. Fracture surface of the sample alloyed for 8 h.
Fig. 14. Fracture surface of the sample alloyed for 32 h.
Fig. 15. Fracture surface of the sample alloyed for 40 h.
from the phase analysis and the tensile tests. The variation in hardness was minimum with alloying time up to 24 h of alloying. Further increase in alloying time increases the hardness by about 15–25 VHN. The increase in hardness with alloying time above 24 h is due, possibly, to the increased formation of intermetallics causing a solid solution strengthening effect. The fracture surfaces of the samples alloyed for different time durations and tested in the instron were examined using a scanning electron microscope. The fractographs are presented in Figs. 11–16.
The figures indicate a smooth transition from a ductile–brittle failure to a quasi cleavage fracture for 8 h of alloying to 32 h of alloying. This can be seen in the reduction in the dimple formation with increase in alloying time from 8 to 32 h (Figs. 11–14) and the predominance of sharp edges indicative of brittle failure. Fracture surface of the sample alloyed for 8 h, Fig. 11 shows equiaxed dimples with some trans-granular cleavage. While the fracture surface of specimens alloyed for 16 h, Fig. 12 indicates mixed cleavage fracture and shear rupture. Fracture surfaces of specimens of sample alloyed for 24 h, Fig. 13 has scattered distribution of dimples in addition to cleavage fracture and further increase to 32 h of alloying, Fig. 14 results in predominant cleavage fracture with no or little evidence of dimples. The fracture surface shows a completely different morphology as compared to the rest of the samples. Moreover, XRD of the samples alloyed for 32 h and sintered were repeated on multiple samples (at least three samples were prepared by alloying the powders individually at 32 h to confirm the anomalous behavior of the samples) from different alloying lots. This anomalous behavior of the samples at 32 h is due, possibly, to the non-
Fig. 13. Fracture surface of the sample alloyed for 24 h.
Fig. 16. Fracture surface of the sample alloyed for 48 h.
Fig. 12. Fracture surface of the sample alloyed for 16 h.
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Fig. 17. Microstructure of the sample alloyed for 8 h. Fig. 19. Microstructure of the sample alloyed for 40 h and sintered at 1250 ◦ C.
Table 3 Compressive strength and strain values during upsetting tests
Fig. 18. Microstructure of the sample alloyed for 32 h, sintered at 1250◦ .
uniform distribution of elemental powders in the blend resulting in a super saturated solution after 32 h of alloying to form intermetallics that are detrimental to the properties of the alloy. Further increase of alloying time to 40 h, Fig. 15 results in a combinational fracture mode (ductile–brittle) with deformation dimples seen across the fracture surface. Increasing to time 48 h of alloying, Fig. 16 results in reduction in dimple formation but with mixed mode failure mechanism similar to 40 h alloying. Further increase in alloying time possibly results in the breaking up of the intermetallics to facilitate stabilization of beta phase at the surface of the particulates yielding a alpha core with beta surface which on sintering produces grain boundary beta phase to aid deformation. This is confirmed by the X-ray phase analysis after sintering of the sample, which indicates absence of Fe–Al intermetallics for the samples alloyed for durations higher than 32 h. Figs. 17–19 show the representative microstructures of the samples alloyed for different time durations. It can be seen from the above microstructure that the surface has predominantly an alpha phase grains with few grain boundary phases. The microstructure of the sample alloyed for 32 h and sintered showed mainly alpha case with porosity along the grain boundary and intermetallics within and near the grain boundary. The weak intermetallics that form during alloying in combination with the porosity possibly contributed to low strains and strength values recorded for the samples. Fig. 19 shows the microstructure of the sample alloyed for 40 h duration and sintered. The microstructure showed a well-distributed alpha case with beta phase along alpha grain
Alloying time (h)
Strength recorded (MPa)
Strain values (%)
40 48
1039 ± 5 1070 ± 7
23 ± 0.5 18 ± 0.7
boundaries. Though porosity and intermetallics were seen on the grain boundaries but the deformation was facilitated possibly by the beta phase distribution resulting in higher strains before failure. Higher strengths compared to 32 h of alloying was also recorded due, possibly, to the beta phase arresting the crack propagation that was causing failure and expending it on increasing the strain. After ascertaining the strength values and the alloying conditions under which the elemental peaks disappear and the phases start to form, upsetting tests were done on those sample alloyed for 40 and 48 h to evaluate the workability of the alloy. Pellets of dimension of 12 mm diameter × 9 mm length were pressed, sintered and used for the above tests. Strength at the maximum load of 95 tonnes was recorded for the samples alloyed for 48 and 40 h and sintered at 1250 ◦ C, respectively. The results obtained from the upsetting tests are presented below in Table 3. The samples did not show a clear failure (the samples did not fail during compression tests as the limiting load of the equipment was possibly lower than the strength of the alloy tested at room temperature) when subjected to compression at the maximum load of the equipment (95 kN). The compressive tests were performed at a speed of 2.5 mm/min and were continued until failure or maximum load is reached depending on whichever occurs first. The results had a similar trend of that of the tensile results where the samples alloyed for 40 h and sintered had a strain value of 9.6 ± 1.2% at failure while the samples alloyed for 48 h had a strain value of 9.4 ± 0.6% at failure.
4. Conclusions From the above results, the following conclusions were derived:
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1. Alloying time has significant effect on the formation of intermetallics and hence the mechanical properties of the alloy blend. 2. Sintering atmosphere influences the phase formations in the alloy and therefore affects the tensile properties of the alloy. 3. The phase formations, microstructure and ratio of the phases formed affect the mechanical strength of the alloy. 4. Significant alloying could not be observed in samples alloyed for less than 24 h. 5. The fracture surfaces indicated a combination of ductile and brittle failure whose relative proportions varied with the alloying time. 6. Alloying time of 40 h resulted in optimum combination of properties with respect to strength and strain achievable.
Acknowledgments The authors are thankful to Singapore Institute of Manufacturing Technology for their support of this work by the way of in-house project (C01-P-150 AR). The help of Mr. Ho Meng Kwong on the compaction and sintering is acknowledged with thanks.
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