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Materials Science and Engineering A 480 (2008) 456–463

Formation process of the bonding jointin Ti/Al diffusion bonding

Yao Wei ∗, Wu Aiping, Zou Guisheng, Ren JialieKey Laboratory for Advanced Materials Processing Technology of the Ministry of Education,

Department of Mechanical Engineering, Tsinghua University, Beijing 100084, PR China

Received 6 December 2006; received in revised form 10 July 2007; accepted 13 July 2007

bstract

The process of the formation of Ti/Al diffusion bonding joints was studied by means of scanning electron microscopy (SEM), X-ray diffractometryXRD) and shear strength measurement. Pure titanium and pure aluminum were used as bonding couples. The results show that the process of jointormation can be separated into four stages, and the product of the diffusion reaction is only TiAl under a particular range of holding time. There

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s a delay time tD before TiAl3 is generated, which is mainly affected by temperature. The joint strength depends on the metallurgical combinationercentage and the interface structure in the diffusion zone, and it can reach or even exceed the strength of pure aluminum after TiAl3 forms aayer. The position where shear fracture occurs depends on interface structure in the diffusion zone. 2007 Elsevier B.V. All rights reserved.

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eywords: Diffusion bonding; Diffusion reaction; Interface structure

. Introduction

Ti and Al are widely used in engineering practice, and theirpplication will become more extensive with the increasingemand for lightweight components. In some special locations,he complementary characteristics of Ti and Al are required, suchs lower weight, increased strength and lower cost. Therefore,t is necessary to obtain Ti/Al compound structures. However,here are great differences in the performance of Ti and Al, andhey do not have the same metallurgical features. For these dis-imilar metal couples, diffusion bonding is a suitable method1].

The process of bonding pure Ti and pure Al is fundamentaln studies of Ti-alloy/Al-alloy diffusion bonding, but there haveeen few reports about this bonding process. Some researchersave reported that the performance of Ti-alloy/Al-alloy jointsould be improved by inserting a piece of pure Al or dip plating

l on Ti alloy [2–4], but these techniques only provide a way tobtain better joints; they do not indicate the process of elementiffusion and reaction in diffusion bonding. When pure Ti and

∗ Corresponding author. Tel.: +86 10 6277 3859x5; fax: +86 10 6277 3859.E-mail addresses: icerss@gmail.com, yaow01@mails.tsinghua.edu.cn

Y. Wei).

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921-5093/$ – see front matter © 2007 Elsevier B.V. All rights reserved.oi:10.1016/j.msea.2007.07.027

ure Al, including pure Al interlayers or Al cladding, are bonded,he intermetallic compound of Ti and Al will appear in the jointsnd influence the performance [5–10]. Thus, it is very importanto study the process of intermetallic compound growth in bond-ng joints, which will be helpful for designing Ti/Al diffusiononding techniques and providing a better understanding of dif-usion bonding. According to the above, we studied the structurend formation process of Ti/Al diffusion bonding joints.

. Experimental procedure

Commercially pure Ti TA2 and commercially pure Al L4ere employed in this study. The chemical compositions andhysical performance of TA2 and L4 are shown in Table 1.he sizes of the samples were TA2 – 5 mm × 10 mm × 5 mm,4 – 5 mm × 10 mm × 10 mm, with a bonding face measuringmm × 10 mm.

Oxide films of Ti and Al could form easily on eachurface. The surfaces of Ti and Al had to be processedefore bonding or they would prevent the diffusion ofi and Al. The Ti surfaces were processed as follows:

round by metallographic sandpaper → etched by 3% HF + 30%NO3 solution → water flushed → ultrasonically cleaned in

cetone → dried. The Al surfaces were processed as fol-ows: ground by metallographic sandpaper → etched by 6%

Y. Wei et al. / Materials Science and Engineering A 480 (2008) 456–463 457

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the maximum vacuum degree was 1.0 × 10 Pa; the highestbonding temperature was 1200 ◦C; the heating curves were setand controlled by an intelligent temperature controller. The shearstrength of the Ti/Al joints was measured by a Gleeble 1500D

Fig. 1. (a–d) Microstructures of the fracture surf

aOH solution → water flushed → etched by 40% HNO3olution → water flushed → ultrasonically cleaned in ace-one → dried. After these treatments, a quite thin and compactayer of Al2O3 formed on each Al surface. These Al2O3 filmsould protect Al surfaces from being oxidized continually, andhey were so thin that they could be crushed easily wheniffusion bonding started. Diffusion bonding was conducted

mmediately after the sample surfaces were processed.

Vacuum diffusion bonding equipment was used in the exper-ments. The main parameters of the equipment were as follows:

ig. 2. XRD comparison between �Ti and pure Ti. Line a, pure Ti; line b,racture surface on Ti side at 625 ◦C 10 min.

n Ti side with holding different times at 625 ◦C.

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Fig. 3. EDS analysis of the grain phases on Ti side fracture surface.

4 and Engineering A 480 (2008) 456–463

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Table 1Chemical composition and physical performance

Material Ti Al Fe Si Cr Mg C O

Chemical composition (wt.%)TA2 99.25 – 0.30 0.15 – – 0.10 0.20L4 – 99.58 0.20 0.17 0.03 0.02 – –

Material Density(g cm−3)

Meltingpoint (◦C)

Coefficient of linearexpansion (10−6 K−1)

Specific heat(J (kg K)−1)

Physical performanceTA2 4.5 1677 8.2 539.1

58 Y. Wei et al. / Materials Science

hermal–mechanical simulation machine. The push-shear modeas used, and the pushing speed was 1 mm/min. The microstruc-

ures of the fracture surfaces and the backscattered electronmages of the bonded joints were observed using a JSM-6301Fcanning electron microscope (SEM). The phase structures inhe joints were analyzed using a D/max-IIIA X-ray diffractionnstrument (XRD).

The technical parameters used in the diffusion bondingere: static pressure, 5 MPa; bonding temperature, 500 ◦C,50 ◦C, 600 ◦C, 625 ◦C, 650 ◦C; holding time, 10–600 min; heat-ng rate, 60 ◦C/min; vacuum degree, 2.5–3.0 × 10−3 Pa. The

ooling process was conducted until the temperature droppedo 100 ◦C in the vacuum chamber. The sample deformationatios were controlled at about 10% in all diffusion BONDINGxperiments.

Fig. 4. (a–e) Microstructures of Ti side fracture surface

L4 2.7 660 23.8 934.8

s at holding 60 min with different temperatures.

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. Results

.1. Fracture site of joints

The XRD results of fracture surfaces on the Al side showedhat there were only Al diffraction peaks under all bondingemperature and holding time conditions. But on the Ti side,here were not only Al peaks but also Ti and TiAl3 peaks underertain conditions. By this token, the shear fracture occurs onhe interface of Al and Ti where elements are not sufficientlynterdiffused or in Al where elements are sufficiently interdif-used; shear fracture may also occur on the interface of reactionroducts and Al.

.2. Structures of diffusion zone

Samples bonded at 625 ◦C with different holding times weretudied in order to understand the transformation of the struc-ures in the diffusion zone. Fig. 1 shows the microstructures ofhe fracture surfaces on the Ti side with holding times of 10, 60nd 600 min at 625 ◦C. For the convenience of comparison, theicrostructure of the Ti sample surface after etching treatment

s also shown (Fig. 1(a)). These pictures indicate that the areahere elements are not sufficiently interdiffused retains the orig-

nal pattern, just as the vertical arrow denotes in Fig. 1(b); in therea where interdiffusion is sufficient, metallurgical bonding isormed and ductile fracturing is displayed, just as the horizontalrrows denote in Fig. 1(b)–(d). Tiny grain phases, as denoted

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Fig. 5. (a–d) Microstructures of Ti side fracture surface

ineering A 480 (2008) 456–463 459

y the slanted arrows in Fig. 1(c) and (d), appear in the ductileracture area after a holding time of 60 min; the grain phasesrow with increasing holding time.

The results from XRD analysis, Fig. 2, show that there isome deviation in the diffraction angle of Ti peaks between purei and the fracture surfaces on the Ti side when ductile fracturesppear. This indicates that solid solution �Ti is generated in theiffusion zone. Low TiAl3 peaks appear on the XRD patternsith the appearance of grain phases, and they also grow with

he grain phases. Fig. 3 shows the results of energy dispersivepectrometry (EDS) analysis for the grain phases. They weredentified as TiAl3, so the moment when the grain phases appears the moment when the diffusion reaction occurs.

.3. Influence of bonding parameters on joints

The main parameters of diffusion bonding are temperature,olding time and pressure. In this test, the same pressure andositive stop were used to restrict the deformation ratio of theamples.

The microstructures of the Ti side fracture surfaces at a hold-ng time of 60 min with different temperatures are shown inig. 4. The figure shows that the original surfaces, as the verti-al arrows denote in Fig. 4(a) and (b), decrease with the increase

n temperature from 500 to 600 ◦C. When the bonding tempera-ure reaches 600 ◦C, the original surfaces disappear. This meanshe elements become increasingly more acutely diffused. Grainhases appear at 625 and 650 ◦C, as the slanted arrows denote

s at 650 ◦C bonding with different holding times.

4 and Engineering A 480 (2008) 456–463

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n Fig. 4(d) and (e). Fig. 5 shows the microstructures of thei side fracture surfaces at 650 ◦C bonding with different hold-

ng times. No original surfaces appear, and grain phases, as thelanted arrows denote, were formed after 10 min. From then on,hey grew gradually. There was a time delay before the diffu-

ion reaction, which was close to 60 min at 625 ◦C but less than0 min at 650 ◦C, so the influence of the bonding temperatureith a constant holding time seems much more evident than the

nfluence of the holding time at each bonding temperature.

ig. 6. (a–c) Backscattered electron images of bonded joint under 650 ◦C andifferent times.

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ig. 7. Plot of ln �x vs. ln(t − tD) by linear regression analysis using Eq. (3).

Fig. 6 shows the backscattered electron images of a jointonded at 650 ◦C with different holding times. The white section

s Ti, the black section is Al, and the gray section in the middle ishe diffusion reaction layer, which was detected to be TiAl3 byDS analysis. The maximum thickness of the diffusion reaction

ayer in each sample was measured, because the diffusion reac-

ig. 8. XRD patterns of fracture surfaces on Ti side under different temperaturesnd different times.

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ion occurred early in the maximum thickness spot. Accordingo the diffusion rules, at a given temperature, the thickness ofhe diffusion reaction layer should match the empirical formula11]:

x = ktn (1)

here �x is the thickness of the diffusion reaction layer, k theate constant, t the holding time, and n is the kinetic exponent.f the delay time is considered, the formula should be rewrittens

x = k(t − tD)n (2)

n �x = n ln(t − tD) + ln k (3)

here tD is the diffusion reaction delay time. The results of mea-urement (Fig. 6) are as follows: holding time 60 min, maximumhickness 1.2 �m; holding time 240 min, maximum thickness.5 �m; holding time 600 min, maximum thickness 4 �m. At50 ◦C, the tD is less than 10 min. The results were substitutednto Eq. (3), and a plot of ln �x versus ln(t − tD) was obtained,s shown in Fig. 7. Linear regression analysis gave the best-t straight line, and the value of the kinetic exponent n waseckoned as 0.49, approaching 1/2. Therefore, the growth of theiffusion reaction layer obeys the parabolic law:

x = k(t − tD)1/2 (4)

This is the standard diffusion growth model. Xu and co-orkers suggested that the parabolic growth behavior could be

ttributed to the limits imposed by lattice diffusion [12].

ig. 9. Free energy of formation of different Ti–Al intermetallic compounds asfunction of temperature.

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.4. Product of diffusion reaction

Fig. 8 shows the XRD patterns of fracture surfaces on the Tiide. Only TiAl3 peaks appear in addition to Ti and Al peaks. Ofhe Ti–Al intermetallic compounds, three phases, Ti3Al, TiAlnd TiAl3, can exist stably. Their free energies of formationave been calculated by Kattner [13], and the results obtainedre shown in Fig. 9 with the temperature range of 0–1000 ◦C.he results show that TiAl3 has the lowest free energy.onsequently, the first diffusion reaction product must beiAl3.

Fig. 10 shows the XRD patterns of the Ti side surfacesbtained by layer-by-layer grinding from the fracture sur-ace on the Ti side to the Ti base-metal. This sample wasroduced at 650 ◦C with a holding time of 600 min. Thereere Ti, Al and TiAl3 peaks in the XRD patterns, so onlyiAl3 was the product of diffusion reaction within a long

ime.

.5. Shear strength

Fig. 11 shows the shear strength of the bonding sam-les using different bonding parameters. According to theicrostructures of the fracture surfaces, it can be concluded

hat the metallurgical combination percentage and inter-ace structure of TiAl3/Al significantly influence the sheartrength.

Before the new TiAl3 phase appears, the joint strengthepends on the size of the interdiffusion solid solution area; theyoth increase together. In the early stage of TiAl3 formation, therain phase area replaces the interdiffusion solid solution area,nd the combination of grain phases and Al is not as firm as theolid solution, so the rate at which the joint strength increaseslows down, as shown in Fig. 11(a). After the grain phases con-ect to form a layer, the combination of bedded TiAl3 and Al

s firmer again, and the joint strength increases to reach or evenxceed the strength of pure Al after the thermal bonding cycle,s shown in Fig. 11(b). Finally, fracture occurs in Al.

ig. 10. XRD patterns of the Ti side surfaces obtained by layer-by-layer grindingrom fracture surface on Ti side to Ti base-metal at 650 ◦C, 600 min.

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by plastic deformation, and Al2O3 films on the Al surface arecracked because of the great differences in plasticity and thermalexpansion coefficient between Al2O3 and Al. After that, Ti/Alinterdiffusion occurs, and solid solution �Ti is generated. The

ig. 11. (a and b) Shear strength of the joints at 625 and 650 ◦C with holdingifferent times.

. Discussion

.1. Cause of appearance of grain patterns

At the initial stages of the diffusion reaction, new phase TiAl3s nucleated. These new phase particles cannot connect with eachther and are separate in the diffusion zone. When the joints areheared at this moment, plastic deformation occurs and disloca-ion pile-ups will form at the particles. These piled-up loops areepelled by the particles through the action of their image forces14]. On the other hand, the leading loop will be pushed towardshe particles by stresses set up by the pile-up and the appliedhear stress. As the bond strength between TiAl3 particles andi was higher than that between TiAl3 particles and Al, the inter-

ace of TiAl3 particles and Al will ultimately fail when one or twooops are pushed toward it. If this occurs, a void is formed. Theonsequence is that the repelling forces on subsequent loops arerastically reduced and the greater part of the pile-up can empty

tself into the newly formed void. The void growth eventuallyeads to fractures. Fig. 12 shows a schematic diagram of therocess above. Therefore, TiAl3 particles remain at the Ti side

ngineering A 480 (2008) 456–463

racture surface and exhibit grain patterns (Fig. 13(a)). At thel side fracture surface, however, there are only shear dimples

Fig. 13(b)).

.2. Complete process of diffusion bonding

As described above, the process of diffusion bonding of Ti/Alan be summarized as follows.

In the first stage, the real contact area of Ti/Al is augmented

Fig. 12. Schematic pattern of the grain phases appearing process.

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ig. 13. (a and b) Microstructures of the fracture surfaces with holding 240 mint 625 ◦C.

oint strength increases with the expansion of the solid solutionrea. If the bonding is stopped in this stage, fractures will appearn Al near �Ti where the elements are sufficiently interdiffused,nd they will also be found on the interface of Al and Ti wherelements are insufficiently interdiffused.

In the second stage, the diffusion reaction begins when theoncentration rate of Ti and Al reaches that of TiAl3. The newhase TiAl3 is nucleated and displays grain patterns. The jointtrength increases slowly. If the bonding is stopped in this stage,

ractures will appear on the interface of grain phases and Al asell as in Al near �Ti.In the third stage, grain phases grow and join together to

orm a layer. The direct interdiffusion of Ti/Al becomes indirect

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ineering A 480 (2008) 456–463 463

nterdiffusion through the TiAl3 layer. The joints become firmer.f the bonding is stopped in this stage, fractures will appear onhe interface of the reaction layer and Al.

In the last stage, a TiAl3 layer expands to the regions of bothi and Al. This growth obeys the parabolic law. If the bonding

s stopped in this stage, fractures will appear on the interface ofhe reaction layer and Al or in Al.

Actually, the whole joint cannot be in the same stage at theame time. That is to say, some area may precede to the lasttage while another area is still in the third, the second, or evenhe first stage. Therefore, the joint strength is a synthetic resultf all the stages at the same time.

. Conclusions

1) In Ti/Al diffusion bonding, the forming process of joint for-mation can be separated into four stages: Ti and Al elementsinterdiffuse to become a solid solution; the new phase TiAl3is generated as grain patterns; grain phases join together toform a layer; the TiAl3 layer grows in accordance with theparabolic law.

2) The product of the diffusion reaction is only TiAl3 under aparticular range of holding time.

3) The joint strength depends on the metallurgical combinationpercentage and interface structure in the diffusion zone, andit can reach or even exceed the strength of pure Al afterTiAl3 forms a layer.

4) The position where shear fracture occurs depends on inter-face structure in the diffusion zone.

5) There is a delay time tD before TiAl3 is generated, which isclose to 60 min at 625 ◦C and less than 10 min at 650 ◦C.

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1244–1246.[8] Z. Junshan, W. Tao, Z. Meili, et al., Acta Metall. Sin. 38 (10) (2002)

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638–647.13] U.R. Kattner, J.-C. Lin, Y.A. Chang, Metall. Trans. A 23A (8) (1992)

2081–2090.14] D. Broek, Elementary Engineering Fracture Mechanics, 3rd ed., Hague,

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