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Brazing of titanium-based alloys with amorphous 25wt.%Ti-25wt.%Zr-50wt.%Cu filler metal
Materials Science and Engineering, A 188 (1994) 305-315
305
Brazing of titanium-based alloys with amorphous 25wt.%Ti-25wt.%Zr-50wt.%Cu filler metal O. Botstein Israel Institute of Metals, Technion, Haifa 32000 (Israel) A. Rabinkin Amorphous Metals, AlliedSignal Inc., Parsippany, NJ (USA) (Received January 18, 1994)
Abstract Ti-Pd (ASTM grade 7) and T i - 6 A I - 4 V alloys (where the composition is in approximate weight per cent) were brazed in a vacuum furnace, each to itself, by using a new rapidly solidified amorphous 25Ti-25Zr-50Cu brazing foil. The joint tensile strength, fatigue resistance and microstructure were determined, the latter by X-ray diffraction analysis, scanning electron microscopy-energy-dispersive spectroscopy and scanning microscopy. The joint tensile strength is close to that of each base metal. The fatigue properties of Ti-Pd (grade 7) joints do not differ from those of this base metal. The microstructure and mechanical properties of the brazed joints depend on the brazing cycle conditions: a fine lamellar eutectic joint microstructure consisting of a-Ti and y-[Ti(Zr)]zCu (tetragonal MoSie-type ) phase is observed after brazing of Ti-6A1-4V alloy at 900 °C for 10 min, followed by fast cooling. This brazing operation results is high strength joints. Brazing at temperatures higher than 900 °C and/or with a relatively low cooling rate results in a coarse dendritic microstructure consisting of y-[Ti(Zr)]eCu and hexagonal 2 Laves Cu2TiZr phases. Fine precipitations of y-tetragonal phase in the a-Ti matrix were also observed in the transition area between the base metal and joint in this case. Joints with such microstructures are brittle and have a low strength. It is shown that fast cooling suppresses formation of the 2 Laves brittle phase, thus resulting in high mechanical properties of the brazed joint.
1. Introduction
To date, the majority of titanium brazing has been carried out using silver-based and aluminum-based filler metals. Unfortunately, the service temperature of brazemements made with these filler metals is limited to about 300 °C [1-3]. The formation of brittle intermetallic compounds between titanium and filler metal elements at the base metal-filler metal interface results in rather brittle joints and creates galvanic corrosion problems [3, 4]. Therefore joints which have to sustain higher temperatures in service are mostly brazed with Ti-15Cu-15Ni clad strip (where the composition is in approximately weight per cent). In spite of its good wettability and high resultant joint strength, this clad strip material cannot be manufactured as homogeneous foil by conventional rolling owing to brittleness resulting from the alloy intermetallic phases forming upon crystallization of its conventional ingots. In addition, the clad strip melts in three separate stages [5] and upon crystallization forms a coarse joint micro0921-5093/94/$7.00 ,'~'SI)I 0921 ~5093(93)09545-8
structure which sometimes has a marked porosity [6, 7]. Because melting of the conventional clad strip is completed above 900 °C (1652 °F ), brazing operations should be carried out at temperatures higher than the a +/3 transformation temperature of titanium and close to the /3 transus temperature of titanium-based alloys. This results in a substantial grain coarsening and loss of ductility and strength of the base metal. Recently, amorphous titanium-zirconium-based filler metals were produced in ribbon form in a wide variety of compositions using rapid solidification technology [5]. To date, the data available on the properties of joints obtained with these more advanced filler metals are still limited. It was therefore logical to attempt brazing with the rapidly solidified amorphous ductile 25Ti-25Zr-50Cu alloy brazing foil [5]. This material's excellent melting characteristics (narrow T~-T~ temperature range and liquidus 100°C lower than the melting temperature of conventional clad strip material), ultimate uniformity of microstructure and the fact that it contains titanium and zirconium all © 1994 - Elsevier Sequoia. All rights reserved
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appear to make it compatible with titanium-based alloys and very attractive for application in titanium-zirconium brazing. To the best of our knowledge, the only reported results obtained with similar amorphous alloys are those of Onzawa and coworkers [8, 9] who describe the tensile properties of Ti-6AI-4V and commercial pure Ti joints. The best fatigue properties reported by them for Ti-6A1-4V joints, obtained after brazing for 10 min at 900 °C and using 35Ti-35Zr-15Cu-15Ni alloy are some 15-20% lower than those of the similarly heat-treated base metal. Good corrosion resistance in 5% NaCI solution was also reported. The joints were obtained by induction heating method which, on the one hand, permits high heating and cooling rates but, on the other hand, is applicable only to a very restricted range of cross-sections of the joined parts. In addition, high heating and cooling rates of induction heating are usually not representative of the actual temperature regimes in industrial vacuum furnaces. The joint microstructure obtained by these workers after a short brazing cycle using induction heating was described as a fine cellular or lamellar eutectoid. Coarse acicular Widmanst~itten structure distinct from the original base metal microstructure was observed at the interface when the brazing temperature was performed below the/3 transus. No phase analysis of the joint microstructure was performed by these workers. This study was undertaken in order first to determine brazing conditions which result in optimal mechanical properties of joints brazed with amorphous Ti-Zr-based filler metal under conditions which are close to those used in industrial vacuum furnace brazing, and second to perform a detailed phase analysis of the joint microstructure to establish optimal structure-mechanical properties relationships.
2. Materials and experimental procedure
Ti-Pd (grade 7) and Ti-6A1-4V alloys rods of 12 mm diameter were used as base metal parts for manufacturing tensile and fatigue specimens, whereas plates 3 mm thick were used for preparation of joints which were subjected to metallographic examination. Rapidly solidified amorphous 25Ti-25Zr-50Cu (Ti33Zr17.3Cu49.7) alloy brazing foil 6 mm wide and 60/~m thick was used as a filler metal. This foil has solidus T~= 842 °C and liquidus TI = 848 °C. Its vitrification starts at about 400 °C on heating. The foil density is about 6.67 g cm- 3. Brazing was carried out in a vacuum furnace for 5-10 min at 900-1000 °C in the case of Ti-6A1-4V and at 880 °C in the case of Ti-Pd (grade 7). Heating and cooling during the brazing cycles were carried out
at two different rates, i.e. a low rate of about 15 °C rain- l and a high rate of 35 °C min- i. The high cooling rate was achieved by introducing argon gas during the cooling part of the brazing cycle. Before brazing, the sample mating surfaces were polished and degreased. During brazing, the specimens were compressively loaded with a pressure on joints about 1 MPa. Cylindrical samples of 4.5 mm diameter were machined from the brazed rods. Three identical samples were subjected to tensile strength tests after each heat treatment. Samples for metallographic observations were prepared as single-lap shear test specimens with 5 mm × 20 mm cross-section. Fatigue tests of the brazed and the as-received Ti-Pd (grade 7) samples were carried out in the tension-tension mode with a stress ratio R =0.1. The microstructure of the brazed joints was determined by means of scanning electron microscopy (SEM)-energy-dispersive spectroscopy (EDS), X-ray diffraction analyses and scanning transmission electron microscopy (STEM). The STEM examination of a Ti-6A1-4V joint was carried out with a JEOL 2000FX electron microscope. STEM specimens were cut out perpendicular to the joint plane and thinned by ion milling to perforation.
3. Results and discussion
3.1. Mechanical properties The tensile properties of the as-received base metals and brazed joints are presented in Table 1. As can be seen, the tensile strength of Ti-Pd (grade 7) and Ti-6AI-4V joints is close to that of the as-received materials when the heating or cooling rate is high, about 35 °C rain- ~, but the plasticity of the Ti-6A1-4V joints is rather limited. When the heating and cooling rate is low, about 15 °C min-1, the fracture strength is lower than the yield strength o. of the joints heated and cooled with the 35 oC min- rate and the samples undergo brittle fracture within the joints. The plasticity achieved with the Ti-Pd (grade 7) joints heated and cooled at 35 °C min-1 is close to that of the base metal. Fracture occurs outside the joints, and the fracture faces have dimpled morphology. The morphology of the specimens fractured in the transition region or within the joints has features associated with pseudocleavage. The Vickers microhardness values measured across the Ti-6A1-4V joint are shown in Fig. 1, The microhardness of the joint obtained after brazing cycles is substantially higher than that of the base metal. The difference between the joint microhardness profiles after three different brazing cycles is a result of different microstructures, as will be shown below.
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Brazing of Ti-based metal with amorphous 25Ti-25Zr-50Cu
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TABLE 1. Tensile strength of Ti-6A1-4V and titanium (ASTM grade 7) brazed joints and corresponding as-received base metals Sample
Brazing conditions
Tensile properties o,y.2 (MPa)
Ti-6AI-4V joint
900 °C, 10 min, heating and cooling with V= 35 °C min800 °C, 15 min, heating and cooling with V= 35 °C min-
Location of fracture in the braze
(MPa)
e (%)
860
881
3
Transition region
982
1032
0.75
Transition region
au.r~
t
900 °C, 10 min, heating and cooling with V= 15 °C min-
--
667
--
Within the joint
900 °C, 10 min, heating and cooling with V= 35 °C minto + 840 °C, 4 h
--
938
--
Transition region
Base metal
As received
939
1027
15
--
Ti-Pd (ASTM grade 7) joint
880 °C, 10 min, heating and cooling with V= 35 °C min900 °C, 10 min, heating and cooling with V= 35 °C min-
367
483
27.5
Outside the joint
375
472
10
Transition region
414
501
28
--
Base metal
t
As received
550
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Fig. l. Vickers microhardness profile across the Ti-6AI-4V joint after brazing cycles: - - , 890 °C, 10 min, followed by cooling with Vc,~,~=35°C min-~; ---o---, 1000°C, 5 min, followed by cooling to 840 °C with Vcoo~= 35 °C min + annealing at 840°C for 4 h; - - z x - - , 900 °C, 10 min, Vcoo~=35 °C min -~ to 840 °C + annealing at 840 °C for 4 h. T h e fatigue properties of the T i - P d (grade 7) joints are s h o w n in Fig. 2. A s can be seen f r o m the S - N curve, the fatigue limit (greater than 106 cycles) of the b r a z e d s p e c i m e n s is at the same level as that of the base metal (about 3 0 0 MPa). S o m e s p e c i m e n s w e r e fractured in the base metal area. A n e x a m p l e of the fracture surface after fatigue failure of o n e of the samples (sample b l 1; ama× = 377 MPa) is s h o w n in Fig. 3.
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i
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Fig. 2. Fatigue properties of Ti-Pd (grade 7) joints brazed with 25Ti-25Zr-50Cu amorphous brazing filler metal at 880 °C for 10 min with a heating and cooling rate of 35 °C min-~: a, asreceived base metal; b, brazed samples.
T w o different zones can be distinguished o n this fracture surface (Fig. 3(a)): fatigue crack p r o p a g a t i o n z o n e F, and final fracture z o n e D. T h e fracture o c c u r r e d in the joint transition region. Z o n e F m o r p h o l o g y has fatigue striations a n d p s e u d o c l e a v a g e features (Figs. 3(b) and 3(c)), w h e r e a s the final fracture z o n e D has a d i m p l e d m o r p h o l o g y (Fig. 3(d)). In specimens u n d e r lower stresses (samples b6 and b 1 3 in Fig. 2) the whole fracture surface is c o v e r e d by fatigue striations. A similar fracture m o r p h o l o g y is o b s e r v e d in fatigue-fractured as-received base metal samples.
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Brazing of Ti-based metal with amorphous 25Ti-25Zr-50Cu
Fig. 3. Fatigue fracture face of the brazed specimen bl 1 (see Fig. 2). Fracture has occurred in the transition region of the braze: (a) general view showing the fatigue region F and the final fracture region D; (b), (c) fatigue striations combined with pseudocleavage in the zone region F; (c) magnified area shown by arrow in (d); (d) dimpled fracture in the region D.
4. Microstructure
4.1. The microstructure of the remelted and crystallized filler metal In order to clarify the genesis of the microstructure of the joints made with Ti-Pd (grade 7) and Ti-6A1-4V alloys, it was useful to consider first the microstructure of pure 2 5 T i - 2 5 Z r - 5 0 C u alloy melted and crystallized with the same temperature cycle. A comparison of the microstructure thus obtained with joint microstructure should reveal the nature of joint phases and to clarify what kind of changes are associated with the interaction of the liquid filler metal with the base metals used. Therefore a substantial amount of pure 2 5 T i - 2 5 Z r - 5 0 C u foil was placed on a titanium substrate and melted and cooled in a vacuum furnace. As can be seen in Fig. 4(a), pure filler metal crystallized in the form of faceted eutectic, the primary phase having the tetrahedral form. SEM X-ray digital mapp-
ing in the concentration form, shown in Figs. 4(b) and 4(d), indicates that the tetrahedra are enriched with titanium while the matrix is enriched with copper. Zirconium is dispersed almost homogeneously except for some restricted zirconium-depleted area, probably enriched in a-Ti (Fig. 4(d)). The compositions of these phases determined by EDS microanalysis are listed in Table 2. In fact, accurate estimates of the phase composition formulae yielded Til.6Zr0.sCu for the Ti-rich crystals and Cu2TiL2Zr0. s or Cu2Ti~.2Zr0. 6 for the matrix phase. These assumed phase formulae were confirmed by X-ray analysis. Indeed, the X-ray diffraction spectrum obtained from the above specimen (Fig. 5) contains lines of two main phases: tetragonal 7-[Ti(Zr)]2 Cu and hexagonal 2 Laves phase Cu2TiZr. Some amount of a-Ti may also be present in the spectrum. The lattice parameters that can be assigned to the tetragonal 14/ m m m MoSi2-type 7 phase are a = 3 . 0 6 A and
O. Botstein, A. Rabinkin
Brazingof Ti-based metal with amorphous 25Ti-25Zr-50Cu
309
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Fig. 4. (a) Scanning electron micrograph of the remelted and crystallized 25Ti-25Zr-50Cu filler metal ribbon and (b)-(d) corresponding X-ray maps ((b) titanium map; (c) copper map; (d) zirconium map).
TABLE 2. Chemical compositions of phases observed in the crystallized 25Ti-25Zr-50Cu alloy Phase
Tetrahedral crystals Matrix
Concentration (at.%)
Compound assumed
Ti
Zr
Cu
50.2
16.6
33.2
[Ti(Zr)]2Cu
31.4
18.2
50.4
Cu2TiZr
c = 10.86 A. These values correspond to lattice parameters of 17 at.% Zr pseudobinary alloy, which belongs to a group of isomorphous solid solution alloys located along Ti2Cu-Zr2Cu cross-section of the paternal ternary T i - Z r - C u phase diagram [10]. According to Woychic and Massalski [11] the phase relationships in this diagram are dominated by the existence of a ternary Cu2TiZr phase, which is in equilibrium with nearly all the various binary phases. An extended region of ternary y Laves phase is found in the central part of this system at subsolidus temperatures [12]. The
existence of an extended range of isomorphous solid solutions between Ti2Cu and Zr2Cu phases with the tetragonal I 4 / m m m MoSi2-type crystal structure at 700 °C was also proved in [10]. Thus the y phase zirconium concentration determined by X-ray analysis is similar to that found by SEM analysis (Table 2) and close to the zirconium concentration of the filler metal. The hexagonal 2 Laves phase, with a = 5 . 1 5 A and c = 8.24 A (c/a = 1.6) is similar to the ternary )t phase found in the T i - Z r - N i system [12]. Finally, if one accepts the construction of the T i - Z r - N i phase diagram as given by Woychic and Massalski [11], then it is proper to assign 2 5 T i - 2 5 Z r - 5 0 C u alloy as belonging to the narrow binary 7-[Ti(Zr)]2Cu+2-CuzTiZr region or ternary 7 + 2 + a region (Fig. 6). 4.2. The microstructure of Ti-6Al-4Vjoints The microstructure of a Ti-6A1-4V joint brazed at 900 °C for 10 min which is shown in Fig. 7 is strongly affected by the heating and cooling rates of the brazing cycle. A fine lamellar eutectic in the joint area develops when the heating and cooling rate is relatively high (35 °C min -1) (Figs. 7(a) and 7(b)). This eutectic
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Brazingof Ti-based metal with amorphous 25Ti-25Zr-50Cu
,z (
10.~
2t~ts
y
2W.
q9.970)
Lin~.ar
Fig. 5. X-ray diffractogram (Cu Ka radiation) obtained from the remelted and crystallized filler metal alloy. The identified phases are a-tetragonal [Ti(Zr)]2Cu, 2 hexagonal Cu2TiZr and a-Ti.
C~
Ti
10
20
30
40
50
O0
70
IlO
nO
Zr
ATOMIC ",¢,ZIRCONIUM
Fig. 6. Isothermal 703 °C section of the Cu-Ti-Zr system according to ref. 11. The arrow shows the location of 25Ti-25Zr-50Cu alloy.
occupies almost the whole braze area. Only traces of Widmanstfitten structure are found in the base-filler metal interface area. A substantially different microstructure is observed after brazing at the low heating and cooling rate (about 15 °C min-l). Here, a coarse eutectic microstructure containing large dendrites is
formed in the central region of the joint (Figs. 7(c) and 7(d)), displacing fine lamellar eutectic toward the base metal. This eutectic is similar to that obtained after high heating and cooling rate treatment. The third transition to the base metal zone with a coarse acicular Widmanst~itten microstructure is also present.
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Brazing of Ti-based metal with amorphous 25Ti-25Zr-50Cu
31 1
Fig. 7. The microstructure of Ti-6AI-4V joint brazed with 25Ti-25Zr-50Cu brazing metal at 900 °C for 10 min: (a), (b) high heating and cooling rate (35 °C min- ~), lamellar eutectic in the central part of the joint: (b) enlargement of the region in the square in (a); (c), (d) low heating and cooling rate, coarse dendritic eutectic in the central region of the joint; (d) enlargement of the region in the square in (c).
Obviously, the slow cooling rate causes the excessive development of unfavorable coarse brittle microstructural zones located in the center and the periphery of the braze which are only in an "embryonic" state after cooling at a high rate. The original microstructure of the base metal is still preserved for both these cases. There is a substantial difference between the joint thicknesses after brazing with high and low cooling rates: 90 ktm and 140 ~ m correspondingly (compare Figs. 7(a) and 7(c)). This different apparently is a consequence of a more extensive interaction and mutual dissolution of the liquid filler and the base metal in the case of slow cooling. The higher brazing temperature (5 min at about 1000 °C) combined with a low cooling rate also results in the coarse dendritic
structure within the brazed area and even partial a ~ fl transformation in the base metal regions adjoining the braze. The details of the slow-cooling brazement were also studied using high resolution STEM. The microstructure and a number of electron diffraction patterns thus obtained are shown in Figs. 8(a)-8(c) and 9(a)-(c). In addition, in situ EDS microanalysis permitted a more correct identification of the location of the observed braze zones. The appearance of a relatively coarse eutectic in the central part of the joint under high magnification and the corresponding electron diffraction pattern are shown in Figs. 8(a) and 8(b). Here, the phases were identified as v-[Ti(Zr)]2Cu and hexagonal 2-CuffiZr. Additional selected-area diffraction
312
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Brazing of Ti-based metal with amorphous 25Ti-25Zr-50Cu
Fig. 8. Scanning transmission electron micrographs and electron diffraction patterns, obtained in the central area of the Ti-6AI-4V joint after brazing with a low cooling rate: (a), (b) coarse eutectic and its corresponding diffraction pattern identified as y[Ti(Zr)]2Cu + 2-Cu2TiZr; (c), (d) electron diffraction patterns obtained from the coarse 2 phase in different orientations.
patterns taken from the coarse particles in the central part were also identified as 2 hexagonal phase (a = 5.15/k; c = 8.24 ,/k) (Figs. 8(c) and 8(d)). Interestingly, this phase has the same lattice parameters as the 2 phase formed in the "virgin" remelted filter metals under equilibrium conditions. This indicates that the duration of these brazing cycles is sufficient for formation equilibrium phases. Very fine precipitations, probably in the form of platelets 1 0 - 2 0 nm in diameter, are revealed in the transition region (Fig. 9(a)) and identified as [Ti(Zr)]2Cu particles in an a-Ti matrix. From the analysis of a number of diffraction patterns, shown in Figs. 9(b) and 9(c), the following orientation relationships between a-Ti and 7 phase were derived: [1120]a11[123] r (0i 13)a11(116)~
Fine lamellar eutectic zone 2, which is located between coarse dendritic zone 1 and Widmanst~itten zone 3 (Fig. 7(c)), built up from the a and 7 phases. Thus, summarizing the S T E M observations, one can conclude that several different regions can be distinguished in the joint microstructure obtained after the brazing cycle with the low heating and cooling rate: ( 1 ) region of coarse eutectic mixture of a and coarse 2 Laves hexagonal phases in the joint center; (2) region of a fine lamellar eutectic mixture of a-Ti + ~' tetragonal phases located between the central region and external braze zones; (3) region of the coarse acicular Widmanst~itten structure which also contains a small amount of fine a + ~' precipitations. Finally, it is worth considering how mechanical properties of the joints correlate with different joint microstructures created by different brazing condi-
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Brazing of N-based metal with amorphous 25Ti-25Zr-50Cu
313
NNS O
m
(d)
Fig. 9. (a) Scanning tunneling electron micrograph (bright field) showing fine precipitations of y-[Ti(Zr)]2Cu in the a-Ti matrix within the transition region of the slowly cooled Ti-6A1-4V joint and (b), (c) corresponding selected-area diffraction patterns taken in different orientations; (d) key diagram for the diffraction pattern shown in (c).
tions. As is observed, the fine lamellar eutectic is formed after brazing at 880 °C for 15 rain with the high heating and cooling rate. The interaction of the liquid filler metal with the base metal is rather short in this case, and therefore the joint composition is close to that of the "virgin" filler metal. Base metal regions adjoining the braze area are not substantially alloyed and their microstructure is close to the initial microstructure. Therefore only traces of Widmanst/itten structure are observed at the joint interface. As a consequence, joints having a moderately diluted filler metal area have a high strength and a high plasticity. The tensile fracture mostly occurs in the transition region (Fig. 10(a)) and the fracture face exhibits a dimpled morphology. Conversely, a low heating and cooling rate or a brazing temperature higher than 900°C results in a coarse dendritic structure having the highest microhardness, probably because of the formation of 2 Laves phase.
Under tensile stress this microstructure undergoes multiple brittle cracking within the joint (Fig. 10(b)) and has a low fracture strength. The joint fracture morphology in this case is pseudocleavage. A post-brazing anneal at 840 °C for 4 h results in a strong mutual base-filler metal interaction which draws the braze composition close to that of base metal. On the corresponding paternal phase diagram the alloy composition moves toward the titanium corner and the structures formed are determined now by the specifics of alloyed titanium-based alloys. Indeed, the braze expands up to 300 ~m, and its microhardness and chemical composition are only slightly different from those of the base metal. Still, the fl transus of alloyed braze material is lower than that of the as-received base metal. This leads upon cooling to the formation of Widmanst~itten structure throughout the joint (Fig. 11). The joint strength after such treat-
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Brazing of Ti-based metal with amorphous 25Ti-25Zr-50Cu
Fig. 11. The microstructure after post-brazing annealing treatment at 840 °C for 4 h which was preceded by brazing at 900 °C for 10 min and cooling of the 25Ti-25Zr-50Cu joint, showing Widmanst~itten structure formation throughout the joint (joint thickness, about 300/zm).
Fig. 10. The microstructure of the failure regions after tensile test of Ti-6AI-4V specimens brazed at 900°C for 10 min followed by cooling with (a) a high and (b) a low rate.
ment is substantially lower than that of the as-received base metal (Table 1). Therefore brazing procedures should be chosen in such a way that formation of Widmanstfitten structure will be avoided.
(2) The fatigue properties of the Ti-Pd (grade 7) joints are almost the same as those of the base metal. (3) The microstructure of the brazed joints is determined by the heating and cooling rate and the temperature of the brazing cycle. A fine lamellar eutectic is formed after the 900 °C, 10 min brazing cycle with a relatively high heating and cooling rate (35 °C min-1). A coarse dendritic structure is crystallized at the center of the joint after either cooling at the low ( 15 °C min- 1) rate or a higher brazing temperature. (4) STEM phase analysis of the brazed joint revealed that the fine lamellar eutectic is a mixture of ), tetragonal and a-Ti phases. The main constituent of a coarse 2 + 7 dendritic structure is hexagonal 2 Laves phase Cu2TiZr. The ), phase is identified as the [Ti(Zr)]2Cu phase having the tetragonal I 4 / m m m MoSi2-type structure. Fine precipitations of ~' phase in the a-Ti matrix are also observed in the transition region and their orientation relationships are proposed. (5) High joint mechanical properties are obtained when the formation of a brittle 2 Laves phase is suppressed by a fast heating and cooling regime of brazing operation.
5. Conclusions
(1) The use of 2 5 T i - 2 5 Z r - 5 0 C u amorphous filler metal permits brazing of Ti-Pd (grade 7) at a temperature lower than that of the a --" fl transformation, and T i - 6 A I - 4 V alloy at below the fl transus, thus preserving the original microstructure of these base metals. A very high joint strength, close to that of both base metals, is achieved.
References 1 M. Schwartz, in W. H. Culberly et al. (eds.), Brazing of reactive and refractory metals, in Metals Handbook, Vol. 6, American Society of Metals, Metals Park, OH, 1983, pp. 1049-1060. 2 M. Donachie, Titanium--A Technical Guide, American Society of Metals, Metals Park, OH, 1988, pp. 153-156. 3 T. Takemoto and I. Okamoto, J. Mater. Sci., 23 (1988) 1301-1308.
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Brazing of Ti-based metal with amorphous 25Ti-25Zr-50('u
4 B. Norris, in P. Lacombe et al. (eds.), Proc. 6th World Conf on Titanium, 1988, Les Editions de Physique, 1988, pp. 1209-1213. 5 A. Rabinkin, M. Liebermann, S. Pounds, T. Taylor, F. Reidinger and S.-C. Lui, Scr. MetalL, 25 ( 1991 ) 399-404. 6 S.W. Lan, Weld. J., 6(10)(1982)23-28. 7 D. G. Howden and R. W. Monroe, Weld. J., 51 (1) (1972) 31-36. 8 M. W. Ko, A. Suzumura and T. Onzawa, Proc. Conj. on Titanium: Products and Applications, Vol. 2, Lake Buena
9 10 11 12
315
Vista, I~L, September 30-October 3, 1990, Ti Dev. Ass., Dayton, OH, 1990, pp. 592-601. T. Onzawa, A. Suzumura and M. W. Ko, Weld. J., 69 (1990) 462~-467~. B. M. Chebotnikov and B. B. Molokanov, Neorg. Mater. (Russ.), 26 (1990) 96(/-964. C. Woychic and T. B. Massalski, Z. Metallkd., 7~) (3) (1988) 149-153. B. H. Eremenko, E. L. Semenova and L. A. Tretiachenko, Metals (Russ.) (6)(1990) 191-195.
305
Brazing of titanium-based alloys with amorphous 25wt.%Ti-25wt.%Zr-50wt.%Cu filler metal O. Botstein Israel Institute of Metals, Technion, Haifa 32000 (Israel) A. Rabinkin Amorphous Metals, AlliedSignal Inc., Parsippany, NJ (USA) (Received January 18, 1994)
Abstract Ti-Pd (ASTM grade 7) and T i - 6 A I - 4 V alloys (where the composition is in approximate weight per cent) were brazed in a vacuum furnace, each to itself, by using a new rapidly solidified amorphous 25Ti-25Zr-50Cu brazing foil. The joint tensile strength, fatigue resistance and microstructure were determined, the latter by X-ray diffraction analysis, scanning electron microscopy-energy-dispersive spectroscopy and scanning microscopy. The joint tensile strength is close to that of each base metal. The fatigue properties of Ti-Pd (grade 7) joints do not differ from those of this base metal. The microstructure and mechanical properties of the brazed joints depend on the brazing cycle conditions: a fine lamellar eutectic joint microstructure consisting of a-Ti and y-[Ti(Zr)]zCu (tetragonal MoSie-type ) phase is observed after brazing of Ti-6A1-4V alloy at 900 °C for 10 min, followed by fast cooling. This brazing operation results is high strength joints. Brazing at temperatures higher than 900 °C and/or with a relatively low cooling rate results in a coarse dendritic microstructure consisting of y-[Ti(Zr)]eCu and hexagonal 2 Laves Cu2TiZr phases. Fine precipitations of y-tetragonal phase in the a-Ti matrix were also observed in the transition area between the base metal and joint in this case. Joints with such microstructures are brittle and have a low strength. It is shown that fast cooling suppresses formation of the 2 Laves brittle phase, thus resulting in high mechanical properties of the brazed joint.
1. Introduction
To date, the majority of titanium brazing has been carried out using silver-based and aluminum-based filler metals. Unfortunately, the service temperature of brazemements made with these filler metals is limited to about 300 °C [1-3]. The formation of brittle intermetallic compounds between titanium and filler metal elements at the base metal-filler metal interface results in rather brittle joints and creates galvanic corrosion problems [3, 4]. Therefore joints which have to sustain higher temperatures in service are mostly brazed with Ti-15Cu-15Ni clad strip (where the composition is in approximately weight per cent). In spite of its good wettability and high resultant joint strength, this clad strip material cannot be manufactured as homogeneous foil by conventional rolling owing to brittleness resulting from the alloy intermetallic phases forming upon crystallization of its conventional ingots. In addition, the clad strip melts in three separate stages [5] and upon crystallization forms a coarse joint micro0921-5093/94/$7.00 ,'~'SI)I 0921 ~5093(93)09545-8
structure which sometimes has a marked porosity [6, 7]. Because melting of the conventional clad strip is completed above 900 °C (1652 °F ), brazing operations should be carried out at temperatures higher than the a +/3 transformation temperature of titanium and close to the /3 transus temperature of titanium-based alloys. This results in a substantial grain coarsening and loss of ductility and strength of the base metal. Recently, amorphous titanium-zirconium-based filler metals were produced in ribbon form in a wide variety of compositions using rapid solidification technology [5]. To date, the data available on the properties of joints obtained with these more advanced filler metals are still limited. It was therefore logical to attempt brazing with the rapidly solidified amorphous ductile 25Ti-25Zr-50Cu alloy brazing foil [5]. This material's excellent melting characteristics (narrow T~-T~ temperature range and liquidus 100°C lower than the melting temperature of conventional clad strip material), ultimate uniformity of microstructure and the fact that it contains titanium and zirconium all © 1994 - Elsevier Sequoia. All rights reserved
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appear to make it compatible with titanium-based alloys and very attractive for application in titanium-zirconium brazing. To the best of our knowledge, the only reported results obtained with similar amorphous alloys are those of Onzawa and coworkers [8, 9] who describe the tensile properties of Ti-6AI-4V and commercial pure Ti joints. The best fatigue properties reported by them for Ti-6A1-4V joints, obtained after brazing for 10 min at 900 °C and using 35Ti-35Zr-15Cu-15Ni alloy are some 15-20% lower than those of the similarly heat-treated base metal. Good corrosion resistance in 5% NaCI solution was also reported. The joints were obtained by induction heating method which, on the one hand, permits high heating and cooling rates but, on the other hand, is applicable only to a very restricted range of cross-sections of the joined parts. In addition, high heating and cooling rates of induction heating are usually not representative of the actual temperature regimes in industrial vacuum furnaces. The joint microstructure obtained by these workers after a short brazing cycle using induction heating was described as a fine cellular or lamellar eutectoid. Coarse acicular Widmanst~itten structure distinct from the original base metal microstructure was observed at the interface when the brazing temperature was performed below the/3 transus. No phase analysis of the joint microstructure was performed by these workers. This study was undertaken in order first to determine brazing conditions which result in optimal mechanical properties of joints brazed with amorphous Ti-Zr-based filler metal under conditions which are close to those used in industrial vacuum furnace brazing, and second to perform a detailed phase analysis of the joint microstructure to establish optimal structure-mechanical properties relationships.
2. Materials and experimental procedure
Ti-Pd (grade 7) and Ti-6A1-4V alloys rods of 12 mm diameter were used as base metal parts for manufacturing tensile and fatigue specimens, whereas plates 3 mm thick were used for preparation of joints which were subjected to metallographic examination. Rapidly solidified amorphous 25Ti-25Zr-50Cu (Ti33Zr17.3Cu49.7) alloy brazing foil 6 mm wide and 60/~m thick was used as a filler metal. This foil has solidus T~= 842 °C and liquidus TI = 848 °C. Its vitrification starts at about 400 °C on heating. The foil density is about 6.67 g cm- 3. Brazing was carried out in a vacuum furnace for 5-10 min at 900-1000 °C in the case of Ti-6A1-4V and at 880 °C in the case of Ti-Pd (grade 7). Heating and cooling during the brazing cycles were carried out
at two different rates, i.e. a low rate of about 15 °C rain- l and a high rate of 35 °C min- i. The high cooling rate was achieved by introducing argon gas during the cooling part of the brazing cycle. Before brazing, the sample mating surfaces were polished and degreased. During brazing, the specimens were compressively loaded with a pressure on joints about 1 MPa. Cylindrical samples of 4.5 mm diameter were machined from the brazed rods. Three identical samples were subjected to tensile strength tests after each heat treatment. Samples for metallographic observations were prepared as single-lap shear test specimens with 5 mm × 20 mm cross-section. Fatigue tests of the brazed and the as-received Ti-Pd (grade 7) samples were carried out in the tension-tension mode with a stress ratio R =0.1. The microstructure of the brazed joints was determined by means of scanning electron microscopy (SEM)-energy-dispersive spectroscopy (EDS), X-ray diffraction analyses and scanning transmission electron microscopy (STEM). The STEM examination of a Ti-6A1-4V joint was carried out with a JEOL 2000FX electron microscope. STEM specimens were cut out perpendicular to the joint plane and thinned by ion milling to perforation.
3. Results and discussion
3.1. Mechanical properties The tensile properties of the as-received base metals and brazed joints are presented in Table 1. As can be seen, the tensile strength of Ti-Pd (grade 7) and Ti-6AI-4V joints is close to that of the as-received materials when the heating or cooling rate is high, about 35 °C rain- ~, but the plasticity of the Ti-6A1-4V joints is rather limited. When the heating and cooling rate is low, about 15 °C min-1, the fracture strength is lower than the yield strength o. of the joints heated and cooled with the 35 oC min- rate and the samples undergo brittle fracture within the joints. The plasticity achieved with the Ti-Pd (grade 7) joints heated and cooled at 35 °C min-1 is close to that of the base metal. Fracture occurs outside the joints, and the fracture faces have dimpled morphology. The morphology of the specimens fractured in the transition region or within the joints has features associated with pseudocleavage. The Vickers microhardness values measured across the Ti-6A1-4V joint are shown in Fig. 1, The microhardness of the joint obtained after brazing cycles is substantially higher than that of the base metal. The difference between the joint microhardness profiles after three different brazing cycles is a result of different microstructures, as will be shown below.
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Brazing of Ti-based metal with amorphous 25Ti-25Zr-50Cu
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TABLE 1. Tensile strength of Ti-6A1-4V and titanium (ASTM grade 7) brazed joints and corresponding as-received base metals Sample
Brazing conditions
Tensile properties o,y.2 (MPa)
Ti-6AI-4V joint
900 °C, 10 min, heating and cooling with V= 35 °C min800 °C, 15 min, heating and cooling with V= 35 °C min-
Location of fracture in the braze
(MPa)
e (%)
860
881
3
Transition region
982
1032
0.75
Transition region
au.r~
t
900 °C, 10 min, heating and cooling with V= 15 °C min-
--
667
--
Within the joint
900 °C, 10 min, heating and cooling with V= 35 °C minto + 840 °C, 4 h
--
938
--
Transition region
Base metal
As received
939
1027
15
--
Ti-Pd (ASTM grade 7) joint
880 °C, 10 min, heating and cooling with V= 35 °C min900 °C, 10 min, heating and cooling with V= 35 °C min-
367
483
27.5
Outside the joint
375
472
10
Transition region
414
501
28
--
Base metal
t
As received
550
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Fig. l. Vickers microhardness profile across the Ti-6AI-4V joint after brazing cycles: - - , 890 °C, 10 min, followed by cooling with Vc,~,~=35°C min-~; ---o---, 1000°C, 5 min, followed by cooling to 840 °C with Vcoo~= 35 °C min + annealing at 840°C for 4 h; - - z x - - , 900 °C, 10 min, Vcoo~=35 °C min -~ to 840 °C + annealing at 840 °C for 4 h. T h e fatigue properties of the T i - P d (grade 7) joints are s h o w n in Fig. 2. A s can be seen f r o m the S - N curve, the fatigue limit (greater than 106 cycles) of the b r a z e d s p e c i m e n s is at the same level as that of the base metal (about 3 0 0 MPa). S o m e s p e c i m e n s w e r e fractured in the base metal area. A n e x a m p l e of the fracture surface after fatigue failure of o n e of the samples (sample b l 1; ama× = 377 MPa) is s h o w n in Fig. 3.
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-
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i
i
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IE+05 1E+Oe N U M B E R OF CYCLES (log)
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r
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Fig. 2. Fatigue properties of Ti-Pd (grade 7) joints brazed with 25Ti-25Zr-50Cu amorphous brazing filler metal at 880 °C for 10 min with a heating and cooling rate of 35 °C min-~: a, asreceived base metal; b, brazed samples.
T w o different zones can be distinguished o n this fracture surface (Fig. 3(a)): fatigue crack p r o p a g a t i o n z o n e F, and final fracture z o n e D. T h e fracture o c c u r r e d in the joint transition region. Z o n e F m o r p h o l o g y has fatigue striations a n d p s e u d o c l e a v a g e features (Figs. 3(b) and 3(c)), w h e r e a s the final fracture z o n e D has a d i m p l e d m o r p h o l o g y (Fig. 3(d)). In specimens u n d e r lower stresses (samples b6 and b 1 3 in Fig. 2) the whole fracture surface is c o v e r e d by fatigue striations. A similar fracture m o r p h o l o g y is o b s e r v e d in fatigue-fractured as-received base metal samples.
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Brazing of Ti-based metal with amorphous 25Ti-25Zr-50Cu
Fig. 3. Fatigue fracture face of the brazed specimen bl 1 (see Fig. 2). Fracture has occurred in the transition region of the braze: (a) general view showing the fatigue region F and the final fracture region D; (b), (c) fatigue striations combined with pseudocleavage in the zone region F; (c) magnified area shown by arrow in (d); (d) dimpled fracture in the region D.
4. Microstructure
4.1. The microstructure of the remelted and crystallized filler metal In order to clarify the genesis of the microstructure of the joints made with Ti-Pd (grade 7) and Ti-6A1-4V alloys, it was useful to consider first the microstructure of pure 2 5 T i - 2 5 Z r - 5 0 C u alloy melted and crystallized with the same temperature cycle. A comparison of the microstructure thus obtained with joint microstructure should reveal the nature of joint phases and to clarify what kind of changes are associated with the interaction of the liquid filler metal with the base metals used. Therefore a substantial amount of pure 2 5 T i - 2 5 Z r - 5 0 C u foil was placed on a titanium substrate and melted and cooled in a vacuum furnace. As can be seen in Fig. 4(a), pure filler metal crystallized in the form of faceted eutectic, the primary phase having the tetrahedral form. SEM X-ray digital mapp-
ing in the concentration form, shown in Figs. 4(b) and 4(d), indicates that the tetrahedra are enriched with titanium while the matrix is enriched with copper. Zirconium is dispersed almost homogeneously except for some restricted zirconium-depleted area, probably enriched in a-Ti (Fig. 4(d)). The compositions of these phases determined by EDS microanalysis are listed in Table 2. In fact, accurate estimates of the phase composition formulae yielded Til.6Zr0.sCu for the Ti-rich crystals and Cu2TiL2Zr0. s or Cu2Ti~.2Zr0. 6 for the matrix phase. These assumed phase formulae were confirmed by X-ray analysis. Indeed, the X-ray diffraction spectrum obtained from the above specimen (Fig. 5) contains lines of two main phases: tetragonal 7-[Ti(Zr)]2 Cu and hexagonal 2 Laves phase Cu2TiZr. Some amount of a-Ti may also be present in the spectrum. The lattice parameters that can be assigned to the tetragonal 14/ m m m MoSi2-type 7 phase are a = 3 . 0 6 A and
O. Botstein, A. Rabinkin
Brazingof Ti-based metal with amorphous 25Ti-25Zr-50Cu
309
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Fig. 4. (a) Scanning electron micrograph of the remelted and crystallized 25Ti-25Zr-50Cu filler metal ribbon and (b)-(d) corresponding X-ray maps ((b) titanium map; (c) copper map; (d) zirconium map).
TABLE 2. Chemical compositions of phases observed in the crystallized 25Ti-25Zr-50Cu alloy Phase
Tetrahedral crystals Matrix
Concentration (at.%)
Compound assumed
Ti
Zr
Cu
50.2
16.6
33.2
[Ti(Zr)]2Cu
31.4
18.2
50.4
Cu2TiZr
c = 10.86 A. These values correspond to lattice parameters of 17 at.% Zr pseudobinary alloy, which belongs to a group of isomorphous solid solution alloys located along Ti2Cu-Zr2Cu cross-section of the paternal ternary T i - Z r - C u phase diagram [10]. According to Woychic and Massalski [11] the phase relationships in this diagram are dominated by the existence of a ternary Cu2TiZr phase, which is in equilibrium with nearly all the various binary phases. An extended region of ternary y Laves phase is found in the central part of this system at subsolidus temperatures [12]. The
existence of an extended range of isomorphous solid solutions between Ti2Cu and Zr2Cu phases with the tetragonal I 4 / m m m MoSi2-type crystal structure at 700 °C was also proved in [10]. Thus the y phase zirconium concentration determined by X-ray analysis is similar to that found by SEM analysis (Table 2) and close to the zirconium concentration of the filler metal. The hexagonal 2 Laves phase, with a = 5 . 1 5 A and c = 8.24 A (c/a = 1.6) is similar to the ternary )t phase found in the T i - Z r - N i system [12]. Finally, if one accepts the construction of the T i - Z r - N i phase diagram as given by Woychic and Massalski [11], then it is proper to assign 2 5 T i - 2 5 Z r - 5 0 C u alloy as belonging to the narrow binary 7-[Ti(Zr)]2Cu+2-CuzTiZr region or ternary 7 + 2 + a region (Fig. 6). 4.2. The microstructure of Ti-6Al-4Vjoints The microstructure of a Ti-6A1-4V joint brazed at 900 °C for 10 min which is shown in Fig. 7 is strongly affected by the heating and cooling rates of the brazing cycle. A fine lamellar eutectic in the joint area develops when the heating and cooling rate is relatively high (35 °C min -1) (Figs. 7(a) and 7(b)). This eutectic
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Brazingof Ti-based metal with amorphous 25Ti-25Zr-50Cu
,z (
10.~
2t~ts
y
2W.
q9.970)
Lin~.ar
Fig. 5. X-ray diffractogram (Cu Ka radiation) obtained from the remelted and crystallized filler metal alloy. The identified phases are a-tetragonal [Ti(Zr)]2Cu, 2 hexagonal Cu2TiZr and a-Ti.
C~
Ti
10
20
30
40
50
O0
70
IlO
nO
Zr
ATOMIC ",¢,ZIRCONIUM
Fig. 6. Isothermal 703 °C section of the Cu-Ti-Zr system according to ref. 11. The arrow shows the location of 25Ti-25Zr-50Cu alloy.
occupies almost the whole braze area. Only traces of Widmanstfitten structure are found in the base-filler metal interface area. A substantially different microstructure is observed after brazing at the low heating and cooling rate (about 15 °C min-l). Here, a coarse eutectic microstructure containing large dendrites is
formed in the central region of the joint (Figs. 7(c) and 7(d)), displacing fine lamellar eutectic toward the base metal. This eutectic is similar to that obtained after high heating and cooling rate treatment. The third transition to the base metal zone with a coarse acicular Widmanst~itten microstructure is also present.
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Brazing of Ti-based metal with amorphous 25Ti-25Zr-50Cu
31 1
Fig. 7. The microstructure of Ti-6AI-4V joint brazed with 25Ti-25Zr-50Cu brazing metal at 900 °C for 10 min: (a), (b) high heating and cooling rate (35 °C min- ~), lamellar eutectic in the central part of the joint: (b) enlargement of the region in the square in (a); (c), (d) low heating and cooling rate, coarse dendritic eutectic in the central region of the joint; (d) enlargement of the region in the square in (c).
Obviously, the slow cooling rate causes the excessive development of unfavorable coarse brittle microstructural zones located in the center and the periphery of the braze which are only in an "embryonic" state after cooling at a high rate. The original microstructure of the base metal is still preserved for both these cases. There is a substantial difference between the joint thicknesses after brazing with high and low cooling rates: 90 ktm and 140 ~ m correspondingly (compare Figs. 7(a) and 7(c)). This different apparently is a consequence of a more extensive interaction and mutual dissolution of the liquid filler and the base metal in the case of slow cooling. The higher brazing temperature (5 min at about 1000 °C) combined with a low cooling rate also results in the coarse dendritic
structure within the brazed area and even partial a ~ fl transformation in the base metal regions adjoining the braze. The details of the slow-cooling brazement were also studied using high resolution STEM. The microstructure and a number of electron diffraction patterns thus obtained are shown in Figs. 8(a)-8(c) and 9(a)-(c). In addition, in situ EDS microanalysis permitted a more correct identification of the location of the observed braze zones. The appearance of a relatively coarse eutectic in the central part of the joint under high magnification and the corresponding electron diffraction pattern are shown in Figs. 8(a) and 8(b). Here, the phases were identified as v-[Ti(Zr)]2Cu and hexagonal 2-CuffiZr. Additional selected-area diffraction
312
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Brazing of Ti-based metal with amorphous 25Ti-25Zr-50Cu
Fig. 8. Scanning transmission electron micrographs and electron diffraction patterns, obtained in the central area of the Ti-6AI-4V joint after brazing with a low cooling rate: (a), (b) coarse eutectic and its corresponding diffraction pattern identified as y[Ti(Zr)]2Cu + 2-Cu2TiZr; (c), (d) electron diffraction patterns obtained from the coarse 2 phase in different orientations.
patterns taken from the coarse particles in the central part were also identified as 2 hexagonal phase (a = 5.15/k; c = 8.24 ,/k) (Figs. 8(c) and 8(d)). Interestingly, this phase has the same lattice parameters as the 2 phase formed in the "virgin" remelted filter metals under equilibrium conditions. This indicates that the duration of these brazing cycles is sufficient for formation equilibrium phases. Very fine precipitations, probably in the form of platelets 1 0 - 2 0 nm in diameter, are revealed in the transition region (Fig. 9(a)) and identified as [Ti(Zr)]2Cu particles in an a-Ti matrix. From the analysis of a number of diffraction patterns, shown in Figs. 9(b) and 9(c), the following orientation relationships between a-Ti and 7 phase were derived: [1120]a11[123] r (0i 13)a11(116)~
Fine lamellar eutectic zone 2, which is located between coarse dendritic zone 1 and Widmanst~itten zone 3 (Fig. 7(c)), built up from the a and 7 phases. Thus, summarizing the S T E M observations, one can conclude that several different regions can be distinguished in the joint microstructure obtained after the brazing cycle with the low heating and cooling rate: ( 1 ) region of coarse eutectic mixture of a and coarse 2 Laves hexagonal phases in the joint center; (2) region of a fine lamellar eutectic mixture of a-Ti + ~' tetragonal phases located between the central region and external braze zones; (3) region of the coarse acicular Widmanst~itten structure which also contains a small amount of fine a + ~' precipitations. Finally, it is worth considering how mechanical properties of the joints correlate with different joint microstructures created by different brazing condi-
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Brazing of N-based metal with amorphous 25Ti-25Zr-50Cu
313
NNS O
m
(d)
Fig. 9. (a) Scanning tunneling electron micrograph (bright field) showing fine precipitations of y-[Ti(Zr)]2Cu in the a-Ti matrix within the transition region of the slowly cooled Ti-6A1-4V joint and (b), (c) corresponding selected-area diffraction patterns taken in different orientations; (d) key diagram for the diffraction pattern shown in (c).
tions. As is observed, the fine lamellar eutectic is formed after brazing at 880 °C for 15 rain with the high heating and cooling rate. The interaction of the liquid filler metal with the base metal is rather short in this case, and therefore the joint composition is close to that of the "virgin" filler metal. Base metal regions adjoining the braze area are not substantially alloyed and their microstructure is close to the initial microstructure. Therefore only traces of Widmanst/itten structure are observed at the joint interface. As a consequence, joints having a moderately diluted filler metal area have a high strength and a high plasticity. The tensile fracture mostly occurs in the transition region (Fig. 10(a)) and the fracture face exhibits a dimpled morphology. Conversely, a low heating and cooling rate or a brazing temperature higher than 900°C results in a coarse dendritic structure having the highest microhardness, probably because of the formation of 2 Laves phase.
Under tensile stress this microstructure undergoes multiple brittle cracking within the joint (Fig. 10(b)) and has a low fracture strength. The joint fracture morphology in this case is pseudocleavage. A post-brazing anneal at 840 °C for 4 h results in a strong mutual base-filler metal interaction which draws the braze composition close to that of base metal. On the corresponding paternal phase diagram the alloy composition moves toward the titanium corner and the structures formed are determined now by the specifics of alloyed titanium-based alloys. Indeed, the braze expands up to 300 ~m, and its microhardness and chemical composition are only slightly different from those of the base metal. Still, the fl transus of alloyed braze material is lower than that of the as-received base metal. This leads upon cooling to the formation of Widmanst~itten structure throughout the joint (Fig. 11). The joint strength after such treat-
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Brazing of Ti-based metal with amorphous 25Ti-25Zr-50Cu
Fig. 11. The microstructure after post-brazing annealing treatment at 840 °C for 4 h which was preceded by brazing at 900 °C for 10 min and cooling of the 25Ti-25Zr-50Cu joint, showing Widmanst~itten structure formation throughout the joint (joint thickness, about 300/zm).
Fig. 10. The microstructure of the failure regions after tensile test of Ti-6AI-4V specimens brazed at 900°C for 10 min followed by cooling with (a) a high and (b) a low rate.
ment is substantially lower than that of the as-received base metal (Table 1). Therefore brazing procedures should be chosen in such a way that formation of Widmanstfitten structure will be avoided.
(2) The fatigue properties of the Ti-Pd (grade 7) joints are almost the same as those of the base metal. (3) The microstructure of the brazed joints is determined by the heating and cooling rate and the temperature of the brazing cycle. A fine lamellar eutectic is formed after the 900 °C, 10 min brazing cycle with a relatively high heating and cooling rate (35 °C min-1). A coarse dendritic structure is crystallized at the center of the joint after either cooling at the low ( 15 °C min- 1) rate or a higher brazing temperature. (4) STEM phase analysis of the brazed joint revealed that the fine lamellar eutectic is a mixture of ), tetragonal and a-Ti phases. The main constituent of a coarse 2 + 7 dendritic structure is hexagonal 2 Laves phase Cu2TiZr. The ), phase is identified as the [Ti(Zr)]2Cu phase having the tetragonal I 4 / m m m MoSi2-type structure. Fine precipitations of ~' phase in the a-Ti matrix are also observed in the transition region and their orientation relationships are proposed. (5) High joint mechanical properties are obtained when the formation of a brittle 2 Laves phase is suppressed by a fast heating and cooling regime of brazing operation.
5. Conclusions
(1) The use of 2 5 T i - 2 5 Z r - 5 0 C u amorphous filler metal permits brazing of Ti-Pd (grade 7) at a temperature lower than that of the a --" fl transformation, and T i - 6 A I - 4 V alloy at below the fl transus, thus preserving the original microstructure of these base metals. A very high joint strength, close to that of both base metals, is achieved.
References 1 M. Schwartz, in W. H. Culberly et al. (eds.), Brazing of reactive and refractory metals, in Metals Handbook, Vol. 6, American Society of Metals, Metals Park, OH, 1983, pp. 1049-1060. 2 M. Donachie, Titanium--A Technical Guide, American Society of Metals, Metals Park, OH, 1988, pp. 153-156. 3 T. Takemoto and I. Okamoto, J. Mater. Sci., 23 (1988) 1301-1308.
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Brazing of Ti-based metal with amorphous 25Ti-25Zr-50('u
4 B. Norris, in P. Lacombe et al. (eds.), Proc. 6th World Conf on Titanium, 1988, Les Editions de Physique, 1988, pp. 1209-1213. 5 A. Rabinkin, M. Liebermann, S. Pounds, T. Taylor, F. Reidinger and S.-C. Lui, Scr. MetalL, 25 ( 1991 ) 399-404. 6 S.W. Lan, Weld. J., 6(10)(1982)23-28. 7 D. G. Howden and R. W. Monroe, Weld. J., 51 (1) (1972) 31-36. 8 M. W. Ko, A. Suzumura and T. Onzawa, Proc. Conj. on Titanium: Products and Applications, Vol. 2, Lake Buena
9 10 11 12
315
Vista, I~L, September 30-October 3, 1990, Ti Dev. Ass., Dayton, OH, 1990, pp. 592-601. T. Onzawa, A. Suzumura and M. W. Ko, Weld. J., 69 (1990) 462~-467~. B. M. Chebotnikov and B. B. Molokanov, Neorg. Mater. (Russ.), 26 (1990) 96(/-964. C. Woychic and T. B. Massalski, Z. Metallkd., 7~) (3) (1988) 149-153. B. H. Eremenko, E. L. Semenova and L. A. Tretiachenko, Metals (Russ.) (6)(1990) 191-195.