alumina system

Acta Materialia 54 (2006) 2205–2214 www.actamat-journals.com

The relation between wetting and interfacial chemistry in the CuAgTi/alumina system R. Voytovych a, F. Robaut b, N. Eustathopoulos a

a,*

Laboratoire de Thermodynamique et de Physico-Chimie Metallurgiques, URA-CNRS 29, Institut National Polytechnique de Grenoble, ENSEEG, BP 75-Domaine Universitaire 1130, F-38402, Saint Martin-d’He`res Cedex, France b CMTC, Institut National Polytechnique de Grenoble, F-38402, Saint Martin-d’He`res Cedex, France Received 24 May 2005; received in revised form 15 November 2005; accepted 30 November 2005 Available online 15 March 2006

Abstract From results obtained over the last 10 years from simple reactive metal/ceramic model systems, a concept of reactive wetting has been formulated which suggests a direct relation between wetting and the physicochemical properties of the interfacial reaction products. The aim of this study is to examine whether this concept can help to understand and rationalise the sometimes conflicting results for the CuAg–Ti/alumina system, which is more complex but important in brazing. This was done using new wetting data generated by sessile drop experiments performed by varying the Ti content in the alloy and the furnace atmosphere (neutral gas and two levels of high vacuum). Experiments were carried out on monocrystalline alumina and also on polycrystalline alumina substrates of varying purity and roughness.  2006 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Wetting; Interfaces; Ceramics; Brazing; Capillary phenomena

1. Introduction Brazing of alumina to itself or to metallic substrates is often performed using Cu or CuAg alloys containing a few per cent of Ti, an element that promotes wetting and adhesion. Although it is well established that Ti dissolved in these alloys reacts with alumina forming at the metal– ceramic interface layers of Ti oxides or Cu–Ti–O compounds [1–6], the relation between the interfacial chemistry and wetting in these types of systems has not been established clearly. A major reason for this is the particularly wide scattering of contact angle values (several tens of degrees) of Ti-containing alloys (Cu–Ti, CuAg eutectic– Ti, Sn–Ti) on alumina measured by different researchers. Saiz et al. [7,8] suggested that this high variability of wetting data is an intrinsic feature of this type of metal–

*

Corresponding author. Tel.: +33 476826504; fax: +33 476826767. E-mail address: [email protected] (N. Eustathopoulos).

ceramic system, caused by metastable states generated at the interface by ridges growing at the solid–liquid–vapour triple line. The possible effect on wetting of these types of ridges in the system under study will be discussed further. As for the variability of contact angle data for Ti-containing alloys, this is an indisputable fact. This is also seen by comparing the results of Naidich et al. [1] and Loehman and Tomsia [3] (Table 1) obtained at nearly the same temperature for two CuAg–Ti alloys with very similar compositions. It should be noted that neither the furnace atmosphere nor the type of alumina substrate was the same in these two studies. The results of Loehman and Tomsia and those of Shiue et al. [4] show greater agreement for h values. However, the interfacial reaction products found by these authors are different: TiO0.5 and the M6X-type Ti3Cu3O compound, respectively. In view of the results of Table 1, a relation between interfacial chemistry and wetting cannot be arrived at conclusively. This seems to disagree with conclusions drawn in previous studies [9–11]. It is true that these conclusions were drawn using

1359-6454/$30.00  2006 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2005.11.048

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Table 1 Wetting data from the literature for the CuAg–Ti/alumina system Author Naidich et al. [1] Loehman and Tomsia [3] Shiue et al. [4] Janicˇkovicˇ et al. [12] a

T (C) 980 1000 900 950

Type of aluminaa Mono Poly Poly Poly

Atmosphere 3

10 Pa Gettered Ar 103 Pa 103 Pa

Ti (at.%)

Contact angle ()

Reaction products

2.0 1.9 8.0 3.0

85 20 10 35

TiO TiO0.5 M6O Ti-rich layer

Mono, monocrystalline alumina; poly, polycrystalline alumina.

results obtained for very simple model systems (mainly metal/carbon systems). The aim of this work is to study the relationship, if any, between Ti activity in the alloy, interfacial chemistry and wetting characteristics in the alumina/CuAgTi couple. It will thus be possible to verify whether the concept developed during the last 10 years concerning reactive wetting using simple model metal/ceramic couples can be applied to more complex systems that are interesting for applications. Wetting was studied using the sessile drop technique in various furnace atmospheres for three different Ti contents. The effects on the results of alumina type (monoor polycrystalline), purity and surface roughness as well as the influence of the configuration of the system (sessile drop configuration corresponding to a low interface/liquid alloy volume ratio or brazing configuration corresponding to a high value of this ratio) are also discussed. 2. Experimental Wetting experiments were performed by the sessile drop technique at 900 C in a metallic furnace under a vacuum of 105 Pa or in a static atmosphere of He. The gas, initially containing 1 ppm of O2 and H2O, was further purified by being passed through a bed of Zr–Al getter before being introduced into the furnace. For comparison, a limited number of experiments were also done in a vacuum of 103 Pa. The drop images produced with an optical system fitted with a zoom (magnification 30·) were recorded with a video camera connected to a computer, permitting automatic image analysis. This device enables the linear dimensions (drop base radius R and height) and contact angle h of the drop to be obtained with an accuracy of ±1% and ±2, respectively. Fully isothermal spreading of a homogeneous alloy droplet can be performed using the transferred drop (TD) [13] and dispensed drop (DD) [14] versions of the sessile drop technique. It should be noted, however, that both techniques require auxiliary substrates (TD) or crucibles (DD) that do not react with and are not wetted by the alloy. This is hard to achieve in the case of Ti-containing alloys. Wetting experiments were therefore performed using the classic version of the sessile drop technique in which a piece of the alloy is placed on the substrate under study and the system is heated to the experimental temperature. Experiments were carried out using either CuAg–Ti prepared in situ or commercial brazing alloys. In the first type of experiment, a CuAg eutectic alloy droplet was first

prepared by melting Cu (99.999 wt.%) with Ag (99.999 wt.%) on an alumina substrate in a high vacuum. The CuAg–Ti alloy, weighing a total of 100–300 mg, was processed in situ during the sessile drop experiment by directly melting a piece of Ti (purity 99.7 wt.%) over the CuAg on the substrate. This procedure was used to avoid any direct contact between Ti and the ceramic before melting the alloy. Selected drop profile images from one of these experiments are presented in Fig. 1. The first image corresponds to the configuration before melting. Three alloy compositions with Ti contents of 0.7, 2.9 and 8.0 at.% were studied. Using this technique, it is difficult to study alloys with less Ti. Moreover, as will be seen later, interfacial reactions in the case of very dilute alloys can lead to a notable decrease in Ti concentration during interface layer formation. Several sessile drop experiments were also performed using commercial active fillers with a nominal composition of Ag–48.1Cu–3.1Ti (at.%) and Ag–36.7Cu–8.0Ti (at.%). The braze alloys, supplied by Wesgo Metal, were in the form of a foil 0.1 mm thick. The active filler foil containing 3.1 at.% Ti had a nearly uniform chemical composition along its entire length. As for the active filler containing 8.0 at.% Ti, scanning electron microscopy (SEM) inspection of a cross-section showed that it had a AgCu/Ti/AgCu sandwich structure and the thickness of the Ti layer was not uniform along the entire length of the sample. Metal samples weighing about 0.1 g and of approximately cylindrical shape were prepared by folding the initial alloy foil. A limited number of experiments were performed by placing a foil of the active filler between two alumina plates. Most of the experiments were performed using high-purity (99.993 wt.%) a-alumina single crystals with a random orientated surface. Their surface was polished to an average roughness Ra value of a few nanometres. For compar-

Fig. 1. Selected drop profiles of a CuAg–2.9 at.% Ti droplet prepared in situ on a sapphire substrate.

R. Voytovych et al. / Acta Materialia 54 (2006) 2205–2214

ison purposes, polycrystalline a-alumina substrates of 99.5 wt.% purity and two roughness values (Ra = 80 and 1300 nm) as well as 96 wt.% purity a-alumina with Ra = 150 nm were also used. The chemistry, morphology and microstructure of the interfacial reaction products were determined using microprobe analysis (EPMA), SEM with energy dispersive X-ray spectroscopy (EDXS) facility and X-ray diffraction (XRD). XRD was performed on the sessile drop samples using a D500 Siemens diffractometer in h/2h mode with Cu Ka radiation, after dissolution of the non-reacted alloy in 5% nitric aqueous solution at 100 C. 3. Results and discussion 3.1. Spreading kinetics

temperature, ˚C

Fig. 2 shows the variation in contact angle h, drop base diameter and temperature as a function of time for the wetting of sapphire by a CuAg–2.9 at.% Ti alloy prepared in situ in purified He. The time origin was taken to be the melting of the CuAg alloy while Ti on the alloy drop was still undissolved. During the first few minutes, no spreading of the alloy was observed while the temperature was raised at a rate of 15 C/min. The contact angle of about 140 observed at this stage, which is typical of non-reactive metal/ionocovalent oxide systems [15], is that of CuAg eutectic on sapphire. Afterwards, the contact angle decreases, first rapidly and then more slowly to attain

920 880

3.2. Interfacial chemistry

800

The reaction zone in the case of the CuAg–Ti (2.9 at.%)/ alumina couple consists of two layers (Fig. 3(b)) with one layer bordering the substrate, referred to subsequently as layer I, and the other, layer II, being in contact with the alloy. Table 2 presents the chemical composition of these layers obtained by EPMA. The machine was operated at an accelerated voltage of 15 kV. For this voltage, a Monte-Carlo simulation performed assuming TiO as the reaction product leads to a depth of analysis of 1–1.3 lm. For this reason, the layer adjacent to the alumina (layer I) was analysed at spots where the thickness of the new phase was at least 1.5 lm. The fact that the volume analysed was in layer I was further confirmed by the absence of any significant Cu signal (Cu is a major constituent of layer II). For each layer, an analysis was performed at 10 spots at least and, as can be seen from the data, the reproducibility is fairly good. The value of the ratio xTi/xO = 1.75 ± 0.10 in layer I (referred to as Ti1.75O) is between 1.5 and 2, corresponding to Ti3O2 and Ti2O oxides. Both these oxides appear in the Ti–O phase diagram [19] but Ti2O seems to be unstable above 600 C, which

0

500

1000

1500

2000

2500

time, s 10

160 140

8 theta d

100

6

80 4

60 40

Drop base diameter, mm

120 contact angle, deg

a steady contact angle hf = 10 ± 3. The use of a smooth polycrystalline a-alumina substrate leads to nearly the same wetting curves. The sudden spreading appears to be triggered by Ti attaining the triple line by diffusion from the top of the drop. Indeed, a simple evaluation of the distance l that Ti could cover by diffusion in t = 500 s (the approximate experimental time needed to initiate spreading) using a typical coefficient of diffusion in the molten alloy, D  5 · 109 m2/s, gives l = (Dt)1/2  1–2 mm, which is of the order of the initial drop height. It should be noted that during the spreading process the Ti concentration in the drop is not constant, so its kinetics cannot be properly analysed using existing reactive spreading models [14,16]. The only parameter that can be considered intrinsic to the system at the study temperature is the final or steady contact angle. The values of this angle were found to be equal to within a few degrees when measured under purified He or in high vacuum. It is worth noting that contact angles as low as 10–20 are typical of liquid metals on solid metal systems, not of metals on ceramic ones [15]. In terms of work of adhesion W = r(1 + cos hf) (where r is the surface tension of the liquid alloy), if one neglects the effect on r of Ti additions, the change in the interfacial chemistry caused by Ti corresponds to an increase in W by a factor of more than 5. It should be noted that the assumption of a constant r is reasonable as the surface tension of Ti at 900 C, extrapolated from the melting point of Ti, r = 1740 mJ/m2 [15], is much higher than the surface tension of CuAg eutectic alloy (r  1000 mJ/ m2) [17]. The absence of any significant effect on r of small additions (a few per cent) of Ti to CuAg eutectic also agrees with the model calculations performed by Novacovic et al. [18].

840

a

2 20 0 0

b

2207

500

1000

1500

2000

0 2500

time, s

Fig. 2. Contact angle (filled circles), drop base diameter (open squares) and temperature as a function of time for an in situ processed alloy containing 2.9 at.% Ti on sapphire (He atmosphere).

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*

50

Lin (Cps)

40

30

20

* 10

*

*

* *

* * *

0 37

50

40

*

**

* 60

80

70

2-Theta - Scale Fig. 4. XRD pattern from CuAg–Ti (2.9 at.%)/sapphire sample (T = 900 C, 30 min) etched in nitric acid solution. The set of peaks of Ti3Cu3O is marked with asterisks. The unmarked peaks belong to layer I and to alloy retentions.

3.3. Effect of time on interfacial chemistry

Fig. 3. SEM images of drop/substrate interface for three CuAg–Ti alloys on sapphire (in black) with 0.7 at.% Ti (a), 2.9 at.% Ti (b) and 8.0 at.% Ti (c). T = 900 C, t = 30 min, He atmosphere.

suggests that it may precipitate during cooling, as proposed by Stephens et al. [20]. It should be noted, however, that the part of the Ti–O phase diagram close to Ti2O is not yet well established. Layer II has a very similar composition to the Ti3Cu3O compound first characterised by Carim [6]. Further characterisation of the reaction products performed using XRD after dissolution of the alloy in a nitric solution made it possible to identify the Ti3Cu3O compound unambiguously (diamond cubic, Fd3m, with a lattice parameter of 1.137 nm) (Fig. 4). Both layers I and II contain 1.5 ± 0.5 at.% Al released from reduction of Al2O3.

Are layers I and II really formed at the wetting temperature or, as suggested by the results given in Ref. [21], during cooling? This is an important question in determining whether the formation of three-dimensional compounds at interfaces plays a key role in reactive wetting. In the case of the CuAg–Ti/alumina system, this question has been discussed in detail in Ref. [22] on the basis of experiments carried out with an alloy of the same composition, at the same temperature T = 900 C by varying the holding time between 3 and 600 min. Briefly, it was found that both layers I and II grow with time and that the total thickness increases parabolically with time. These results confirm that both layers are formed at the experimental temperature and not during solidification. The reasons for the discrepancy between these results and those reported in Ref. [21] are discussed in Ref. [22]. 3.4. Wetting versus interfacial chemistry While the results presented above demonstrate that wetting in the CuAg–Ti/alumina system correlates with the three-dimensional interfacial chemistry, it has not yet been shown which of the two compounds formed at the inter-

Table 2 Microprobe analysis of the areas indicated as 1–6 in Figs. 3(a)–(c) Alloy (at.% Ti)

Area

Ti (at.%)

Cu (at.%)

Ag (at.%)

0.7

1

64.8 ± 1.5

0.2 ± 0.5

0.2 ± 0.1

O (at.%) 34.8 ± 0.5

Al (at.%) 0

2.9

2 3

61.0 ± 1.5 43.0 ± 1.5

2.6 ± 0.5 39.5 ± 1.0

0.33 ± 0.1 0.33 ± 0.1

34.4 ± 2.0 15.5 ± 1.0

1.4 ± 0.4 1.6 ± 0.5

8.0

4 40 5 6

45.0 ± 1.5 44.3 ± 1.5 22.0 ± 0.5 <0.1

34.5 ± 1.0 34.1 ± 1.0 74 ± 1.0 12.0 ± 1.0

0.25 ± 0.1 0.20 ± 0.1 2.0 ± 0.4 87.0 ± 1.5

17.1 ± 1.0 17.8 ± 1.0 1.5 ± 0.1 0.4 ± 0.1

3.0 ± 0.5 3.5 ± 0.5 <0.1 0.6 ± 0.1

R. Voytovych et al. / Acta Materialia 54 (2006) 2205–2214

face, Ti1.75O or Ti3Cu3O, is responsible for the improvement in wetting due to the addition of Ti. For this purpose a CuAg–Ti/alumina sample was cooled after being held for 3 min at 900 C, before the contact angle attained its steady value. SEM observation of the triple line from above (Fig. 5(a)) showed that the triple line receded slightly during solidification from its initial position, as evidenced by the presence of small metallic particles on the liberated area. Receding is very interesting in this case because it allows the reaction product layer to be analysed directly. EDXS analysis of this layer gives a Ti/Cu atomic ratio close to unity, as in interfacial layer II, consisting of Ti3Cu3O compound. EDXS analysis also gives large amounts of O and Al, indicating that the thickness of the reaction layer at this point is less than the vertical resolution of the EDXS analysis, i.e., less than 1 lm. This is consistent with the surface profile presented in Fig. 5(b). It should be noted that the Ti3Cu3O layer has also been shown to form at the interface in the vicinity of the triple line at the end of the spreading process, after the contact angle attains its steady value [22]. The excellent wetting of alumina by CuAg–2.9 at.% Ti alloy can thus be explained by the formation of this M6X-type compound (Tm = 1112 C [23]), which is known to have a metallic character, as shown by the value of the electrical resistivity q [23]. Indeed, Kelkar et al. found q = 5 · 106 X m, a value very close to that of Ti alloys (1.5 · 106 X m) [24] and lower by 20 orders of magnitude than that of alumina.

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3.5. Effect of Ti content Fig. 6 shows the wetting curves for three alloys with different Ti contents processed in situ on monocrystalline alumina substrates in a purified He atmosphere. In the case of the alloys with a higher Ti content (2.9 and 8.0 at.%), the steady contact angle values did not differ significantly (10 and 7, respectively) but they were essentially lower than in the case of the alloy with 0.7 at.% Ti (66). This different wetting was confirmed by a second experiment performed in He with the alloy with the lowest Ti content, which led to hf = 60, while a third experiment performed by replacing the He atmosphere with a vacuum of 105 Pa resulted in hf = 62. Fig. 3 shows SEM images of the interfacial zones of the samples after the wetting experiments. In the case of the melt containing 0.7 at.% Ti, one reaction layer about 1 lm thick was found at the interface (Fig. 3(a)). The thickness and morphology of this layer did not allow its composition to be analysed on a cross-section sample. For this reason EPMA was performed directly on the layer after dissolution of the remaining alloy in nitric aqueous solution. The accelerated voltage (15 kV) was adjusted so that the depth of sampling was shallow enough to avoid the electron beam penetrating into the alumina substrate. This was attested to by the fact that no Al was detected in the reaction layer. The reaction product consisted essentially of Ti and O (area 1, Table 2). The ratio xTi/xO obtained by averaging the analyses performed at five different spots of the layer was found to be 1.85 ± 0.10, which is very close to the value of 1.75 ± 0.10 determined for this ratio in layer I in the case of the 2.9 at.% Ti alloy. The experiment performed with the alloy containing 8.0 at.% Ti resulted, at first sight, in a more complicated picture. The reason for this is a break in CuAg–Ti (8 at.%) miscibility in the liquid phase at 900 C that causes the melt to separate into two phases, an Ag-rich liquid (L1) with a low Ti content and a Cu-rich liquid (L2) with a high

160

900 880

120 860

100

840

80 60

820

40 800

20

780

0 0

Fig. 5. (a) SEM image of a triple line region of CuAg–2.9 at.% Ti droplet on sapphire after a wetting experiment of 3 min at 900 C. A and B denote the positions of the triple line before and after receding, respectively. (b) Profile taken in a direction normal to the triple line. The position of points A, B, C are given in (a).

temperature, ˚C

contact angle, deg

140

500

1000

1500

2000

time,s Fig. 6. Contact angle versus time for in situ performed experiments with the alloys containing 0.7 at.% Ti (open circles), 2.9 at.% Ti (filled squares) and 8.0 at.% Ti (filled triangles), He atmosphere.

R. Voytovych et al. / Acta Materialia 54 (2006) 2205–2214

Ti content [25]. In view of their chemical compositions (see Table 2), areas 5 and 6 in Fig. 3(c) would appear to correspond to liquids L2 and L1, respectively. The interface consists of a continuous layer about 5 lm thick bordering the sapphire substrate and in contact with the Cu-rich liquid. Grains can also be observed inside liquid L2; their shape suggests that they have been detached from the interfacial layer. This shows that liquid L2 wets the grain boundaries of the reaction products. Intergranular liquid films can lead to fast diffusion, thus increasing the layer’s growth rate. The reaction layer (areas 4 and 4 0 ) has a similar composition to that of layer II in Fig. 3(b) but with a significantly higher xTi/xCu + Al ratio (1.20 versus 1.05). Further characterisation of the interface using XRD clearly revealed that the layer consists of Ti3Cu3O (as for a 2.9 at.% Ti alloy) and also of Ti4Cu2O, both compounds containing some dissolved Al (about 1.5 at.%). Like Ti3Cu3O, Ti4Cu2O is a metal-like compound with a melting point Tm = 1127 C and an electrical resistivity q = 5 · 106 X m equal to that of Ti3Cu3O [23]. In view of the very similar chemical, crystallographic and electrical properties of these compounds, it is expected that the presence of Ti4Cu2O in the layer will not significantly modify the wetting and adhesion properties of the system. The data on interfacial chemistry and final contact angles obtained from the experiments performed with alloys with different Ti contents are summarised in Table 3. It is clear that the addition of an increasing amount of Ti to CuAg resulted in a change in the reaction compounds formed at the interface. The very similar final contact angle values obtained with the 2.9 and 8.0 at.% Ti alloys can easily be explained by the fact that, in both cases, wetting is caused by the same M6X-type compound. In the case of the alloy with 0.7 at.% Ti, such a compound is not formed and as a result a significantly higher contact angle is observed. An even further increase in hf is to be expected by decreasing the thermodynamic activity aTi of Ti in the alloy. Indeed it has been shown that such a decrease in aTi leads to the formation, at the interface, of Ti oxides with a high degree of oxidation, such as Ti3O5 or Ti5O9, resulting in hf values close to and even higher than 90. This has been observed by Kritsalis et al. [26] for a NiPd–Ti/ alumina system (Fig. 7) in which aTi is very small due to extremely strong Ni–Ti and Pd–Ti interactions in the alloy.

120 110

Ti5O9

100

contact angle, deg

2210

90 80

Ti3O5

70 60

Ti2O3 50 40 0.00

0.05

0.10

0.15

0.20

0.25

X Ti Fig. 7. Variation of equilibrium contact angle of NiPd–Ti alloys on monocrystalline alumina at T  1573 K as a function of the molar fraction of Ti. The three plateaus correspond to different Ti oxides [26].

3.6. Mechanical interface As shown by the high value of work of adhesion (2000 mJ/m2), the alloy/alumina interfaces are strong from the point of view of thermodynamics. This favours the formation of a mechanically strong interface but does not necessarily lead to it [27–30]. A qualitative indication of the mechanical strength of the interface is given by the behaviour of a solidified droplet/alumina sample during cooling. The coefficient of thermal expansion of alumina (8.2 · 106 C1) is known to be about half that of Cu (17.7 · 106 C1) [24]. During cooling from the experimental temperature, this difference generates internal stress in the samples, leading to cohesive failure of the alumina (Fig. 8) whatever its type (mono- or polycrystalline alumina) or the composition of the alloy. This behaviour indicates that the mechanical interface in the alumina/alloy system is strong, stronger than the alumina itself. It should be noted, however, that when Ti oxides with a high degree of oxidation are formed at the interface, failure is interfacial, indicating a mechanically weak interface [26,31].

Table 3 Data for interfacial chemistry and wetting for three different Ti content alloys on sapphire Ti content (at.%)

Interfacial compounds

Final contact angle ()

0.7

Ti1.85O

63 ± 3

2.9

Layer I: Ti1.75O Layer II: Ti3(CuAl)3O

10 ± 3

8.0

Ti3(CuAl)3O and Ti4(CuAl)2O

7±3

Fig. 8. Cohesive failure of monocrystalline alumina occurring during cooling to room temperature after a wetting experiment with a CuAg– 2.9 at.% Ti alloy.

R. Voytovych et al. / Acta Materialia 54 (2006) 2205–2214

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3.7. Effect of furnace atmosphere Whereas the final contact angles obtained in purified He and in a vacuum of 105 Pa are equal to within a few degrees, very different values are observed when the experiments are performed in a vacuum of 103 Pa. This is shown by the results of two experiments performed using commercial braze containing 3.1 at.% Ti (Fig. 9). It should be noted that when the alloy was prepared from folded foil, the sample collapsed while melting and immediately formed a drop with a contact angle of 50–80, depending on the experiment, with a complete absence of initial spreading. In purified He, a drop with a contact angle of 60 formed after melting and subsequently spread towards the final contact angle. In contrast, the sample melted in a vacuum of 103 Pa led to a contact angle of 75 that barely changed over time, even after thousands of seconds at the experimental temperature. Inspection of the drops after the experiments showed a loss of metallic appearance in the case of the experiment performed in a poor vacuum, thereby indicating oxidation of the drop. The strong tendency of Ti-containing alloys to oxidise and the consequences of this on wetting were analysed by Kritsalis et al. [32], who showed that oxidation is more pronounced in melts with low Ti concentrations. This is because when xTi increases, the solubility of oxygen in the alloy also increases, leading to droplet deoxidation through dissolution of the oxide skin into the molten alloy. A comparison of Figs. 10(a) and Fig. 3(b) shows that, in contrast to wetting, interfacial chemistry is not significantly affected by furnace atmosphere, 103 Pa vacuum or gettered He. In both cases, the interface consists of two layers, and the total thickness of the interfacial zone is very similar. Thus, the interfacial chemistry in this system is determined by the oxido-reduction reactions occurring between Ti and alumina, while the furnace oxygen has little or no effect. More generally, the type of reaction products for a given oxide substrate and temperature depends on the

Fig. 10. SEM images of commercial braze containing 3.1 at.% Ti on monocrystalline (a) and polycrystalline (b) alumina after wetting experiments in vacuum of 103 Pa, T = 900 C, t = 30 min.

thermodynamic activity of Ti in the alloy, which, in turn, is determined by its concentration and the strength of its interactions with the metal matrix. 3.8. Effect of substrate quality and purity Fig. 11 and Table 4 show that the spreading curves and the steady contact angles on mono- and polycrystalline 160

920

T

880

60 840 50 800

40 30

He atm. -3 10 Pa

temperature, ˚C

contact angle, deg

70

a b c d

140

760

contact angle, deg

80

120 100 80 60 40 20

20 0

100

200

300

400

500

720 600

time, s Fig. 9. Contact angle versus time on alumina single crystals for two experiments carried out with a 3.1 at.% Ti commercial braze in vacuum of 103 Pa (filled triangles) and in He (filled circles).

0 0

500

1000 1500 time, s

2000

2500

Fig. 11. Wetting curves of in situ processed alloys containing 2.9 at.% Ti on different types of alumina presented in Table 4. T = 900 C, He atmosphere.

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Table 4 Contact angles of CuAg–Ti (2.9 at.%) on different types of alumina, 900 C, 30 min

unlikely that ridges of nanometric size will have a significant effect on wetting.

Substrate

Type of alumina (purity, %)

Ra (lm)

Final contact angle ()

3.9. Wetting and interfacial chemistry: comparison with the literature

a b c d

Sapphire Poly (99.5) Poly (96.0) Poly (99.5)

0.004 0.075 0.15 1.3

10 ± 3 17 ± 3 19 ± 3 13.5 ± 3

substrates with different purity and surface roughness values are quite similar. No significant differences in interfacial chemistry and thickness of the reaction layers were found for any of these types of alumina. It should be noted that EPMA of 96% alumina far from the interface revealed the presence of Mg and Si in a second phase located at the grain boundaries. No such elements were detected in reaction layer I or in the bulk of the solidified drop. However, a very small but significant concentration of Si (0.3 at.%) was found in layer II. The hf values obtained with alumina substrates differing in average roughness by a factor of up to 300 (Table 4) lie within ±5, which is close to the experimental error of ±3. The results do not reveal any systematic variation in hf with Ra. It should be noted that in reactive systems the roughness relevant for wetting is not that of the initial substrate but that of the reaction product layer in contact with the liquid during the spreading process. This is equal to the initial roughness, possibly modified by the reaction. However, it was shown in Ref. [22] that during the wetting of a CuAg–2.9 at.% Ti alloy on alumina, the maximum thickness of the reaction layer at the vicinity of the triple line is only a few tens of nanometres, meaning that the effective roughness is close to the initial roughness. The weak effect of roughness on wetting observed in this work in the case of a reactive system is very similar to the results obtained in Ref. [33] for non-reactive systems that also exhibit very good wetting (hf = 20). Indeed, the variation in Ra by two orders of magnitude in these non-reactive systems led to a change in the contact angle of only a few degrees. These results, which contrast with the strong effect of roughness in the case of systems with intrinsic hf of more than 90 (changes of several tens of degrees in hf have been observed with a variation in Ra from a few nanometres to about 1 lm [33]) underline the fact that the energy barrier of metastable states induced by surface defects is greatly reduced at very low hf [15]. The interfacial chemistry on mono- and polycrystalline alumina is very similar (Fig. 10). A final remark concerns the possible effect on spreading and wettability of ridges formed at the triple line as suggested by Saiz et al. [7,8] (see Section 1). As is seen from the surface profile taken at the triple line junction after receding of the alloy (Fig. 5(b)), no ridge can be distinguished in the vicinity of point B implying that the ridge would be of nanometric size, as alumina asperities. As the increase of Ra from a few nanometers to 1300 nm barely modifies the wetting curves (Fig. 11), it is very

Although the average roughness of the substrates used in the experiments reported in Table 1 is not known, the results given in Table 4 and discussed in the previous section suggest that the dispersion of hf values found in the literature is not due to roughness. High sensitivity to oxidation may explain the large differences in contact angles (more than 60) measured by Naidich et al. [1] and by Loehman and Tomsia [3] obtained at nearly the same temperature in the case of two CuAg–Ti alloys with very similar compositions (Table 1). The relatively poor vacuum used in Ref. [1] may be the reason for the large contact angle obtained in this study. Indeed, as indicated in Section 3.7, experiments carried out in a vacuum of 103 mbar led to non-reproducible hf values lying between 50 and 80 instead of the expected value of 10. On the other hand, the results of Loehman and Tomsia [3] for the 1.9 at.% Ti alloy are consistent with ours in terms of wetting but not in terms of interfacial chemistry. These authors found a 3–5 lm thick layer of TixO (2.0 6 x 6 2.8) and a much larger zone rich in Cu and Ti, which they attributed to a Cu–Ti intermetallic formed during solidification rather than to the Ti3Cu3O compound formed at the experimental temperature, as in our experiments. In the case of the alloy containing 8.0 at.% Ti, Shiue et al. [4] (see Table 1) obtained contact angles ranging between 10 and 25. Their lowest value of 10 agrees well with the value observed in our study for the same alloy. It should be noted that Shiue et al. also detected a Ti3Cu3O compound at the interface even after very short periods of contact between the liquid and solid phases. 3.10. Wetting versus brazing When alumina is brazed to alumina using CuAg–Ti alloys, the joint is completely filled, without any pores or voids. An example is given in Fig. 12 for a CuAg–8 at.% Ti braze. A continuous 3–5 lm thick reaction layer is formed at the interface, consisting of the Ti3Cu3O and Ti4Cu2O compounds, as in the wetting experiments. However, as a general rule, the interfacial chemistry in the brazing (‘‘sandwich’’) configuration is less well defined than in sessile drop experiments. Indeed, in brazing, the thickness of the metallic foil (a few tens of micrometres) is two orders of magnitude lower than the average height of the droplets used in a typical sessile drop experiment (a few millimetres). Hence, when brazing is performed with an alloy containing a few per cent of a reactive element, the formation of micrometre-thick reaction layers may significantly modify the average concentration in the reactive element, which in turn may affect the interfacial chemistry and adhesion. In other words, the type of reaction products formed at

R. Voytovych et al. / Acta Materialia 54 (2006) 2205–2214

Fig. 12. SEM image of a braze/alumina interface obtained by brazing alumina (in black) to alumina with a CuAg–8.0 at.% Ti alloy at 900 C for 30 min.

metal/ceramic interfaces during brazing depends not only on the brazing temperature and the reactive element content in the braze but also on the brazing time and thickness of the braze. For this reason it is very difficult to compare reactivity data in the literature obtained using thin foils [4,5,20] or metallic layers deposited on alumina substrates [34]. When alumina is brazed to metallic solids using reactive brazes, a crucial factor is the interaction between the reactive element and the elements contained in the metallic solid [31,35]. For example, when alumina is brazed to a CuNi plate using a CuAg–3.1 at.% Ti alloy, the strong Ni–Ti interactions cause a significant decrease in the thermodynamic activity of Ti, which in turn leads to a dramatic decrease in the thickness of the Ti reaction layer at the metal/alumina interface and to a shift in the composition of this layer towards Ti oxides with a higher degree of oxidation of Ti. This change was found to be responsible for mechanical weakening of the interface and of the joint itself [31]. 4. Conclusions The results obtained in this study demonstrate that a direct relation exists for the CuAg–Ti/alumina system between interfacial chemistry and wetting. As expected, droplets of CuAg alloys without Ti do not wet alumina. They form contact angles much higher than 90, which are typical of noble metal/ionocovalent oxide systems. Additions of 3 and 8 at.% Ti lead to a dramatic decrease in the steady contact angle to about 10, a value which is typical of metal/metal systems, not of metal/ceramic ones. This is attributed to the formation of continuous layers of metal-like M6X-type Ti3Cu3O and Ti4Cu2O compounds at the interface, by reaction between the Ti-containing alloy and alumina. When the Ti content in the alloy is reduced to 0.7 at.%, M6X-type compounds are no longer formed and wetting is controlled by a Ti1.75O layer. This change in the interface chemistry is responsible for a significant increase in the contact angle from about 10 to 60–65. Even higher hf

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contact angles, close to 90, are expected to form with a further decrease in the thermodynamic activity of Ti, resulting in the formation at the interface of TiOn oxides with n P 1.5, such as Ti3O5 or Ti5O9. The wetting of CuAg–Ti alloys on alumina does not vary significantly with the roughness of the solid. Indeed, changes of ±5 were observed in hf when the average roughness varied by three orders of magnitude. This weak effect of roughness on steady contact angles observed in the case of a reactive alloy with hf  10 is also typical of nonreactive liquid–solid systems exhibiting a high degree of wetting. In the case of the experiments carried out in a high vacuum, the quality of the vacuum has little or no effect on interfacial chemistry, which is determined mainly by Ti activity in the alloy. In contrast, the furnace atmosphere strongly affects wetting. Indeed, a medium vacuum can lead to oxidation of the free surface of droplet and to high apparent contact angles. The variability in the contact angles of Ti-containing alloys observed in the literature is very probably the result of furnace atmosphere effects, and not of metastable states generated by ridges growing at the triple line as suggested by Saiz et al. [7,8]. In the CuAg–Ti/alumina system, relatively limited variations in Ti concentration were shown to change the interfacial chemistry, which, in turn, affects wetting and adhesion. This is particularly critical in brazing, where, because of the thinness of the braze foils used in practice and the low concentration of the reactive elements, consumption of the reactive element by the interfacial reactions can significantly modify its average concentration in the braze. Controlling interfacial chemistry in brazing therefore calls for strict control of the process parameters, especially of the temperature cycle. When alumina is brazed to metallic solids, further complications may arise from a change in the thermodynamic activity of the reactive element caused by its interactions with the components of the metallic solid. Acknowledgement R.V. and N.E. acknowledge the support for this study provided by the NEDO International Research Grant. References [1] Naidich Y, Zhuravlev V, Chuprina V, Strashinskaya L. Sov Powder Metall Met Ceram 1973;12:895. [2] Naidich Y. Prog Surf Membr Sci 1981;14:353. [3] Loehman RE, Tomsia AP. Acta Mater 1994;40:547. [4] Shiue R, Wu S, O JM, Wang JY. J Metall Mater Trans A 2000;31:2527. [5] Paulasto M, Kivilahti J. J Mater Res 1998;13:343. [6] Carim A. Scripta Metall Mater 1991;25:51. [7] Saiz E, Cannon R, Tomsia A. Acta Mater 2000;48:4449. [8] Saiz E, Cannon R, Tomsia A. Acta Mater 1998;46:2349. [9] Labrousse B, PhD. thesis, Polytechnical Institute of Grenoble, France; 2000.

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