High purity Al2O3 ceramic:Metallizing strategy, microstructure and sealing properties

ABSTRACT A high purity Al2O3 ceramic (HPAOC) metallizing strategy was developed via gradient coating process of metallizing pastes with different ratios of Mo to manganese glass (MnG) contents, to improve the wettability and reactivity of metallized layer (ML) to the ceramic substrate and the secondary metallizing layer or sealed metals. Self-made HPAOC samples firstly coated by a layer of metallizing paste with a lower proportion of Mo:MnG and superposed a layer of metallizing paste with a higher proportion of Mo:MnG were fired at 1450°C in hydrogen atmosphere. The crystal phase structure, microstructure and element distribution of the metallized samples wer characterized by XRD, SEM and EDS. The results of sealing properties show that the tensile sealing strength of as high as 121MPa and the He leakage rate of as low as 4.2 × 10−11 Pa.m3/s can be obtained of the sealed joints of the as-metallized HPAOC and Kovar.


Introduction
Alumina ceramics with 92 ~ 99%wt Al 2 O 3 (92 ~ 99AOC)are widely used in vacuum electronic devices due to its high temperature resistance, corrosion resistance, excellent dielectric properties, excellent mechanical properties, low cost, suitable thermal properties, available metallization and proper bonding properties with Kovar or high purity copper [1], [2][3][4]. High purity Al 2 O 3 ceramics (HPAOC)with more than 99%wt Al 2 O 3 content have superior insulation and dielectric properties, which are expected to meet the higher requirements of high frequency, high power and long life for high vacuum electronic devices, such as radio frequency windows, microwave tube and traveling wave tube [3,5]. Nevertheless, HPAOC with more than 99%wt Al 2 O 3 content is hardly metallized well using the conventional formula and method for 92 ~ 99AOC.
Many methods of ceramic metallization have been developed up to now, including various modified sintered metal powder process (SMPP) [6], direct copper coating (DBC) [ 7], magnetron sputtering [8,9], chemical vapor deposition (CVD) [10], vacuum evaporation, etc. SMPP was earlier invented by Siemens in Germany in 1935. Over years of research and development, the SMPP metallization layer of 92 ~ 99AOC with dense, excellent air tightness and stable performance has been obtained by a stable technology [11,12].
The metallization mechanisms of Mo-Mn method [13,14] for 92 ~ 99AOC consist mainly of glass phase migration, chemical reaction and molybdenum sintering [15][16][17][18]. The quality of the metallized layer (ML) has the most important impact on the performance of ceramic metal sealing devices. In ML for the activated Mo-Mn method, Mo forms a skeleton, Mn and other oxides form a glassy phase under as high as~1450°C temperature, which evenly fills into the pores of Mo skeleton and form a dense ML [2,[19][20][21].
The metallization mechanisms for 92 ~ 99AOC with superior performance can be explained well by the glass phase migration mechanism [22]. However, the high-purity 99 ~ 99.9AOC has no or little glass phase, the activated constituent forming the bonding phase in the metallization process [12,16]. In our early study on the HPAOC metallization, an idea occurs to us naturally that the glass phase in ML composition was enhanced, aiming at forming a reverse penetrating glass bonding phase. However, the more glass phase in ML leads to very bad sealing properties, a lower sealing strength and a bad air tightness. There are two reasons for that: one belongs to the brittle behavior of the more glass phase in ML; another reason is the less metallic phase and more glass phase in ML surface result into a poor Nickel layer plated, presenting a bad bonding behavior with Kovar metals.
In order to acquire the good bonding behavior of ML with both the ceramic matrix and brazed metals simultaneously, a high purity 998AOC metallization strategy was developed via gradient coating process of metallizing pastes with different proportions of Mo to manganese glass(Mo:MnG), i.e. the samples of 998AOC were coated firstly a layer of ML paste with a lower proportion of Mo:MnG and subsequently coated the second layer paste with a higher proportion of Mo:MnG, dried and fired in hydrogen atmosphere. As a result, it is expected that a dense Mo top layer in ML and a good wettability and bonding of ML to the ceramic matrix both are obtained.
The microstructure and sealing properties of 998AOC with as-prepared ML were studied in detail. It is hopeful that the study can provide a method of improving HPAOC sealing properties with kovar or other metals, and will lay a solid foundation for the application of HPAOC, especially in high frequency field above 100 MHz, as a promising candidate for wave traveling vacuum window materials due to the expected lower energy dissipation.

Raw materials
The starting materials of Al 2 O 3 powder (0.5um, purity 99.9%) and MgO powder (0.5um, purity≥99%) as sintering aid for 998AOC are produced by Guangzhou Sinoma Xinyan Co., Ltd. The powders of molybdenum powder (1-3um, purity≥99%) and manganese powder (10-40um, purity≥99%) used in the metallized slurry come from Hebei Zhuyan Alloy Material Co., Ltd; Al 2 O 3 , SiO 2 and CaO (all analytical purity) used with Mn powder as raw materisls of MnG in ML are produced by Tianjin Zhiyuan Chemical Reagent Co., Ltd. The reagent grade organic vehicles of Terpene alcohol as solvent, castor oil plasticizer and soybean lecithin as dispersant are produced by Guangdongxilong Science Co., Ltd, and ethyl cellulose as binder comes from by Tianjin Fuchen chemical reagent factory.

Preparation process
High purity alumina ceramics 998AOC were prepared by dry press method with the pressure of 100MPa. The ringlike green bodies with external diameter of 20 mm, inner diameter of 12 mm and height of 3 mm were sintered at 1700°C for 2 hr, and the necessary end surface of the sintered samples was polished to a ~0.8 μm roughness. The six series of ML powder batches (Mo+MnG) were designed according to Table 1. RMetallized samples with different mass ratios of Molybdenum to MnG different mass ratios of Mo:MnG of 4/6, 5/5, 6/4, 7/3, 8/2, and 9/1, and the corresponding metallized samples were designated as M1~M6 respectively ( Table 1). The nominal mass composition of manganese glass (MnG) is 35%MnO, 39%Al 2 O 3 , 23%SiO 2 and 3%CaO. The series of metallizing pastes of M1~M6 was prepared by fully mixing the starting solid powders and organic carriers.
In order to investigate the effect of MnG on 998AOC matrix, the samples of M1~M6 with once paste coating were prepared firstly and studied. As-prepared M1~M6 metallizing pastes were evenly coated on the polished end surfaces of the ring-shaped 998AOCs by the screen printing, respectively, and hereafter, dried at 80°C for 2 hours. The dried paste coated samples were put into the hydrogen sintering furnace, and sintered at 1450°C for 50 minutes with the volume proportion of N 2 to H 2 (dew point 30) of 2:1. According to the metallized results of M1~M6, for the gradient coating samples, the paste of M2 (Mo:MnG = 5/5) was selected as the first coating layer, and the pastes of M4 (Mo:MnG = 7/3) and M5(Mo:MnG = 8/2) as the second coating layers. The corresponding gradient paste coating samples of M4@M2 and M5@M2 are named by C1 and C2, respectively. The superposed second layer coating method and the sintering program of C1 and C2 are the same as those of M1~M6.
All the metallized samples and Kovar alloy samples were plated with a 3.5-5 μm layer of nickel [23]. The brazing of the nickel-electroplated ringlike metallized samples and the matched Kovar alloy (see Figure 1) was made in a vacuum furnace at 790°C with silver copper solder AgCu28.
The micromorphology of the samples was observed by scanning electron microscopy (SEM, Gemini SEM500, ZEISS, Germany). The distribution of molybdenum and glass phase in ML was analyzed by energy dispersive spectrometer (EDS, Gemini SEM500, ZEISS, Germany). The X-ray diffraction (XRD, D8, Advance, Bruker, Germany)spectra of the samples were recorded with Cu Ka radiation at a scanning speed of 4°/min in the 2θ range of 10-80°. He mass spectrometer leak detector (ASM 340, Pfeiffer,Germany) was used to detect the vacuum tightness of the sealed samples, and the universal testing machine (RG2000-20A)was used to detect the tensile strength of the sealed samples.

The wettability and phase structure of MnG
During the metallization of ceramics, the glass phase in the metallized layer is crucial to the quality of ML and the properties of the sealed device. In order to explore the inverse infiltration and wettability of MnG on HPAOC, the MnG raw material was dry pressed into a small cylinder of diameter 5 mm and height 5 mm and placed on 998AOC (see Fig.2, a), which were fired under the same heating scheme as metallized samples. The result shows that the MnG has a good bonding phenomenon with the substrate of 998AOC. The contact angle between MnG and high purity alumina ceramic matrix was~27° (see Figure 2,b), indicating good wetting at the temperature of 1450°C.
The phase structure for as-prepared MnG was characterized by XRD and its patterns are shown in Figure 3. It shows that the main phases in the MnG are α-Al 2 O 3 , and manganese spinel phase, manganese peridotite phase, and calcium feldspar phase, in reference with the PDF cards of No:78-2427, No:29-0880, No:35-0748, and No:89-1472. A proper amount of Manganese spinel phase is considered as taking a role of combining the AOC substrate and ML [24]. Calcium feldspar phase (CaAl 2 Si 2 O 8 ) can also improve the compatibility of glass phases with ceramics [25].

The cross sectional microstructures of MLs and interfaces of M1~M6
The interface bonding state of AOC and ML is vital for the sealing properties. Figure 4 presents the cross-sectional SEM microstructures of M1~M6, which shows clearly that the bonding states of 998AOC and ML are gradually changed from tight to loose, from blur interfaces of M1~M2 to clear interfaces of M3~M6, with the decrease of MnG weight percent content of from M1: 60% to M6:10%.
M1 and M2 both have a better bonding state with a blur and dense boundary, which is attributed to the good wettability and interpenetration of higher content MnG to HP-998AOC. Molybdenum particles are evenly distributed in ML of M1 and M2, and the MLs are dense and have no pores. However, the SEM of M3~M6 (see Figure 4 c~f) shows that there occur the pores in the ML. The more appeared pores with Mo increasing can be attributed to the local agglomerating effect of the less amount of MnG melt, which results into the unevenly stress domain, and leads to form a hollow Mo skeleton forming by side during sintering [26]. The kind of pores in ML, especially in the boudary, will seriously degenerate the combining strength and tightness of ceramic matrix and ML. In combination with the nickel plating properties (see Figure 5), the MnG for M2 was chosen as the first coating layer for the gradient coating process strategy, and MLs of M5 and M6 were chosen as the the second coating layer for the gradient coating process strategy.     Figure 5 (a1)~(f1) shows the surface microstructure images of M1~M6. It can be seen clearly that the Mo particles are isolated and distributed evenly in the MnG matrix in M1 and M2, and no pores has been found. Mo particles gradually become a continuous phase with the Mo content increase from M3 to M6, and the more pores appear.

The surface microstructures of M1~M6 and the samples nickel plated
Correspondingly, for all the samples after nickel plating, the morphologies of nickel plating layers have the same trend as surface microstructures of the once metallized samples of M1~M6 ( Figure 5 a2~f2), i.e. the nickel layer is gradually become dense and has a smaller average grain with Mo content increase from M1 to M6, accordingly.

SEM images of cross sectional, external surfaces and nickel plating surfaces of C1 and C2
Figure 6 (a) and (b) presents the cross-sectional, external surfaces and nickel plating surfaces SEM of as twice coated samples C1 and C2, which uses the ML pastes of M4 and M5 as the second layer coating pastes, respectively, and both use the ML paste of M2 as the first layer of coating paste.
The MLs of C1 and C2 have a dense microstructure without observable pores. The microstructure of as twice coated MLs of C2 shows that Mo particles are distributed evenly in the whole ML, which is different from the initially expected image of Mo particles being   C1(a,a1,a2) and C2 (b,b1,b2). gradient distribution form external surface of paste layer to the interface of ceramic and paste layer. This can be interpreted as the precipitation behavior of Mo particles in high-temperature MnG melt due to the higher density of Mo. Figure 6 (a1) and (b1) represents that the surface microstructure images of C1 and C2. The surface microstructure images of show that C1 and C2 have a dense surfaces, and the Mo particles on the surface of C2 are more dense. Figure 6 (a2) and (b2) shows that C1 and C2 have nearly the same nickel morphologies as those of M4~M5, since C1 and C2 use them as the top coating layer.

The element maps of C2
Figure 7(a) shows that there is a distinct composition transition interface. The uniform and continuous distribution of Mo element is distributed in ML and obstructively changed at the interface, due to Mo dissolution within 998AOC, while the distribution of Mn, Al, Si and Ca in the glass phase renders interpenetrating phenomena, especially for Mn. In the matrix of 998AOC, Al is predominant, and also has an evident change at the interface.The distribution of oxygen element is uniform in the entire section for being the constituent element for both 998AOC and ML.
During the metallization process, the molybdenum particles form a continuous and homogeneous molybdenum skeleton. There is no apparent Mo distribution layer emerging in ML as expected, which can be explained as the viscosity of the manganese-containing glass phase at high temperatures decreases and the metal Mo has a larger density, which makes Mo particles migrate toward the bottom of ML. The migration process of Mo particles are considered as a promotion effect for MnG flowing and filling the interstice of Mo particles, and eventually, a dense ML with a continuous and uniformly distributed molybdenum skeleton was formed.

The microstructures and element maps of cross-section of sealed joints
The cross-section morphology in Figure 8(a) and element maps of the sealing samples of 998AOC and Kovar alloy are presented in Figure 8(b~j). The image of Figure 8(a) shows that there are six layers of 998AOC-ML-Ni-AgCu28-Ni-KJ33 well bonded with each other. The MnG of ML has a good infiltration into 998AOC and partly connects into a whole, and there are no observable holes in the whole sealing belt.
The distribution maps of Fe and Co in Fig 8(b,c) originate from Kovar alloy. That of Ni in Figure 8(d) partly originates from Kovar and partly comes from the two electroplated nickel layers of Kovar and ML. The Ni map shows that it diffuses into silver copper solder layer forming NiCu alloy and by a minor amount enters into the ML. Figure 8 (e) and (f) show the profile of silver and copper elements. Ag atoms keep its original sites without noticeable diffusion. Cu atoms have strong interdiffusion with Ni, that demonstrates that Ni layer can promote the fluidity and sealing properties of silver copper solder.
It is observed again that the distribution of Mo in Figure 8(g) is evenly distributed in the ML without noticeable gradient; Si is mainly distributed in the glass phase of ML in Figure 8(h), and a small amount of it migrates into alumina ceramics. The presence of Cu and Ni elements in Figure 8(d,f) in the ML is considered as the introduction during sample grinding. Al and O elements in Figure 8

Metallization principle of gradient coating process
In order to demonstrate the effect of MnG on the surface of ceramic substrate after metallization firing, the MLs of the metallized sample by pure Mo paste (Figure 9(b1), M2 (Figure 9(c1), M5 (Figure 9(d1) and C2 (Figure 9,e1) were eroded off by the mixture of concentrated nitric acid and concentrated hydrofluoric acid by the volume ratio of 1:1, and then the surface microstruture images of correspondingly treated samples were observed by SEM.
Compared to the natural surface of 998AOC in Figure 9 shows that the grains of the alumina ceramic surface keep intact and grain boundaries are natural and clear, indicating that the pure Mo metallized paste does not react with the ceramic substrate and form chemical bonding of high strength. Figure 9(c1) shows the surface microstructure SEM images of ML of M2 (Mo:MnG = 5:5) after acid corrosion. It is evident that the surface grains of alumina ceramics are corroded off in a large area, the surface of which appears pitted and porous, no complete grains and grain boundaries can be seen. Figure 9 (d1) shows that of ML of M5 (Mo: MnG = 8:2) after acid corrosion. Compared with that of ML of M2 in Figure 9 (c1), M5 has the less MnG phase, and the acid corrosion degree is relatively weaker, indicating the more MnG phase of the activated  constituent, the severe reaction with alumina ceramic and the denser binding intermediate layer. Figure 9 (e1) shows that of the composite gradient metallization paste coating sample C2, which was made by M2 metallization paste as the base layer and M5 metallization paste as the top layer, so that in the metallization process, the base layer with higher manganese glass phase content can fully react with the ceramic substrate to form a solid fixed layer. The degree of corrosion on the surface of the ceramic substrate is similar to that of ML of M5.

Mechanical properties and air tightness
The He leakage rates and tensile strength of the sealed joints of M1-C2 all the metallized samples and Kovar KJ33 are shown in Figure 10 (a&b). It can be seen that there exists an apparent extremity effect of mechanical properties and air tightness for M1~M6 with the decrease of MnG, in which M3 has the minimum leakage of 1.5 × 10 −10 Pa.m 3 /s, and M5 has the maximum tensile strength of 91MPa. This indicates that the proper proportion of Mo and glass phase is vital for the sealing properties, in which the glass phase is considered as the bonding constituent for ceramic and ML, and the metal Mo phase is regarded as the bonding constituent to metals, such as electroplated nickel layer or copper.
The lower He leakage rates and higher tensile strength of as twice coated samples C1 and C2 indicate that the proposed strategy of twice coated ML is feasible. C2 has maximum tensile strength up to 121MPa, and the He leakage rate gets as low as 4.2 × 10 −11 Pa. m 3 /s due to the combining effect of the good weldability of ML of M2 to ceramic substrate and the good wettability of ML of M5 to metals.
Comparing with the relatively reported research results (see Table 2), it can be found that the sealing strength of 121MP in this study is lower than that of 150MPa for 95AOC [27] by activated Mo-Mn method, and is higher than that of 103MPa of for 99.5AOC [19].

Conclusions
A high purity Al 2 O 3 ceramic metallization with high performance has been successfully developed via gradient coating process strategy, and the better wettability and reactivity of ML to the ceramic substrate and the secondary metallizing layer have been obtained. The tensile sealing strength of as high as 121 MPa and the He leakage rate of as low as 4.2 × 10 −11 Pa.m 3 /s have been obtained of the sealed joints of the as-metallized ceramic and Kovar KJ33. It is hopeful that the metallization strategy can provide a reference of high-performance sealing for the high purity ceramic with metals.