Chapter 8 Ion-mixing

Chapter 8 Ion-Mixing B. X. LIU

8.1. I N T R O D U C T I O N

The earliest use of an energetic ion beam generated from low-energy accelerators (several 100-keV or higher) to modify material properties was initiated in early 1960s, i.e. the doping of semiconductors by implanting the desired ion species into silicon wafers (Carter and Grant 1976). Since then, ion implantation has significantly developed and has become a supporting technique in the semiconductor and VLSI technology (Dearnaley et al. 1973). Since early 1970s, ion-implantation has also been employed to deal with metallic materials, e.g., for improving the wear and corrosion resistance, catalytic properties, superconductivity, etc. (Mayer and Lau 1988). As a versatile metal ion source has not been available for some time, it has not been possible to implant various metal species to achieve a high ion concentration frequently required for modifying the surface properties of the metallic materials. An alternative scheme, namely ion-mixing of multiple metal layers, was therefore introduced and a new research area of synthesizing metastable alloys by ion-mixing has boomed since the early 1980s (Liu et al. 1983). Figure 8.1 shows schematically how ion-mixing is realized in multilayered films consisting of metal layers A and B. The A-B multilayered films are prepared by depositing alternately the metals A and B on a substrate. The as-deposited films are then irradiated by a scanning ion beam, e.g., 300 keV xenon ions, to a certain dose to achieve a uniform mixing between metal layers A and B, resulting in forming an Al-xBx alloy. The substrate can be SiO2 or NaC1 acting as a supporting substance or it can be metal A as a matrix for forming a surface Al-xBx alloy coating on metal A. It is known that the first amorphous alloy, also called metallic glass, was obtained by liquid melt quenching (LMQ) in 1960 (Klement et al. 1960). Because amorphous alloys have a non-crystalline structure (with no vulnerable grain boundaries), they are expected to possess unique properties. During the past 30 years, various techniques have been developed to synthesize metallic glasses. From a physical point of view, high cooling rates of 106 K/s or higher are needed to avoid crystallization against amorphization (Davies 1983, Hafner 1981, Giessen and Whang 1980) and cooling rates of up to 107 K/s have been achieved so far. Such developments lead to expand the number and types of metallic glasses, including a variety of metal-metalloid and metal-metal amorphous alloys (Takayama 1976). Laser-quenching has also been applied for producing amorphous alloys (von Allmen 1983, Lin and Spaepen 1984), as many intense laser sources have 197

198

B. X. Liu

amorph~s SiOz

... ..

: NaCI 3,.keVXeions

:~~7

[

cn'statline Figure 8.1. Ion-mixingof multilayered films consisting of metal layers A and B.

become available and they can achieve a cooling rate of 1012 K]s in the material's surface layer (See the chapter on "Laser Processing" by K.E Kobayashi in this book). In comparison, effective cooling rates can reach as high as 1014 K]s (Thompson 1969) in ion-mixing, and therefore this technique can also produce a variety of amorphous alloys. In fact, the cooling rates achieved during ion-mixing are probably the highest among the currently available glass producing techniques. Up-to-now, numerous metastable alloys with either an amorphous or crystalline structure have been obtained by ion-mixing in some 100 binary metal systems (Liu and Jin 1997), and in many alloy systems that are not obtainable by other non-equilibrium processing techniques. The glass-forming ability (GFA) of a system can be considerably improved by ion-mixing because of its high cooling rate. Furthermore, ion-mixing offers a unique opportunity for studying the amorphization mechanism, since the ion dose can be varied at small intervals and so the process of ion-mixing can be studied step-by-step. Based on experimental data, several empirical models have been proposed to predict the GFA and the glass-forming composition range (abbreviated hereafter as GFR) of the binary metal systems. Liu et al. (1983) proposed the first rule, namely the structural difference rule, which predicted a sufficient condition and demonstrated the predominant role of the crystalline structures of the component metals in forming metallic glasses by ionmixing. Further studies led Liu to formulate a two-parameter model. The first parameter is the maximum possible amorphization range (MPAR), which was defined as the total width of the two-phase region, equaling the whole composition range (100%) minus the maximum solid solubilities, observed from the corresponding equilibrium phase diagram (Liu et al. 1987). The second parameter is the heat of formation A H f calculated by Miedema's model (Miedema 1976). Based on a quantitative combination of the parameters, the binary metal systems are classified into readily, possibly, and hardly glass-forming ones (Liu et al. 1987). Alonso and Simozar (1983) constructed a two-dimensional map with

199

Ion-Mixing

the ratio of atomic radii, rA/rB, and the A H f and by summarizing the published data, they separated the binary metal systems into glass-forming and non-glass-forming ones and predicted that no amorphous alloy could be formed in a system with a A H f exceeding + 10 kJ/mol. Ossi (1990) proposed an atomistic model based on collision cascades and the related bombardment-induced surface compositional changes. Though the prediction by the two-parameter model matched best with the experimental data, lack of a relevant thermodynamic interpretation was the main drawback until the early 1990s. It was also noted that metallic glasses were unexpectedly obtained in some alloy systems with a positive A Hu, which have long been recognized as non-glass-forming ones (Liu 1987, Jin et al. 1995, Zhang et al. 1995). The thermodynamic stability of the different alloy phases was evaluated by Miedema, who developed a semi-quantitative model in the mid 1970s to calculate the free energy curves of the alloy phases versus alloy compositions by using the available characteristic parameters of the constituent metals (de Boer et al. 1989), rather than the parameters of the alloy phases, as required in CALPHAD calculation. This chapter attempts to present a brief review of the up-to-date progress on the synthesis of metastable alloys by ion-mixing of multilayered films. It is divided into the following sections: a brief description of the underlying physics in ion-mixing, thermodynamic calculations based on Miedema's model, experimentation and sample design, formation of metallic glasses in systems with a positive or a negative heat of formation, metastable crystalline alloys synthesized by ion mixing, and applicability of the interface concept to solid-state amorphization.

8.2. B R I E F D E S C R I P T I O N O F U N D E R L Y I N G P H Y S I C S IN I O N M I X I N G

This section briefly describes the underlying physics in ion-mixing and the experimental details of ion-mixing will be discussed in Section 8.4. The ion-mixing process is commonly divided into two steps, viz., an atomic collision cascade and subsequent relaxation. After relaxation, one can consider having a delayed step and it has no significant effect on alloy phase formation. When an ion with an energy E1 impacts on the material surface and penetrates inside, it collides with the atoms in the material. If an atom in the material receives an energy E2 exceeding a threshold energy Ed, the atom will leave its own lattice site and a vacancy is created. This displaced atom is referred to as knock-out atom and it will, if E2 is greater than Ed, trigger secondary collisions, then tertiary ones, and so on, and this is named an atomic collision cascade. To describe atomic displacements and defect creation caused by ion irradiation, a parameter "dpa", (abbreviation for displacement p e r atom) is defined, which can be calculated through atomic collision theory, as

do-

E2 (max)

d p a -- Nd = ~ . Crc

No

rJed

v(E2)

9dE2 ~

(8 1) "

whereNd and No are the numbers of displaced and total atoms per volume, respectively; r is the ion flux, ere is the collision cross-section; Ed and E2 are the displacement threshold energy and the energy received by the knock-out atom, respectively; v(E2) is the number

200

B. X. Liu

of displaced atoms created by the knock-out atom with an energy of E2, and dot/dE2 is the differential collision cross-section at E2. The atomic collision cascades trigger intermixing of the metal layers A and B in the multilayered films. As the energy of the ions is on the order of several 100 keV and the binding energy of the solids is typically 5-10 eV, atomic collision is a highly dynamic process of far-from equilibrium. By receiving an adequate irradiation dose, the layered structure of the multilayered films is smeared out and metals A and B mix into a uniform mixture, which is in a highly energetic state with a disordered atomic configuration. At the termination of the cascade, the highly energetic mixture should somehow relax towards equilibrium, and at this stage, the equilibrium thermodynamics comes into play in governing the direction of relaxation. Whether the mixture can reach the equilibrium state or not, however, depends on the temperature and time conditions available during this period. During the atomic collision cascades, all the atoms are involved in violent motion, it is hard to imagine the formation of any kind of alloy phase. Consequently, it is commonly considered that the structure of the alloy phase is to be formed during the relaxation period. As the relaxation period is very short (10-1~ -9 seconds, as we will discuss soon) and the mixture cannot straightforwardly go to the equilibrium state in most cases, it frequently resides in some intermediate states, corresponding to the metastable states of either an amorphous or a simple-structured crystalline phase. From an energy point of view, a thermal spike concept was introduced to discuss the ion-mixing process. Suppose a target is irradiated by ions of 300 keV energy and a current density of 1-4/zA/cm 2 for a certain time period, a huge amount of energy would be deposited into the material surface within a very thin layer. Estimation showed that in a local region, e.g., using a cylinder geometry, with a diameter of 5 nm and a length of 100 nm, the temperature could be as high as 103-104 K, while the surrounding material remains close to the ambient temperature of 300 K. According to thermal conduction calculations, the temperature drops to 300 K within approximately 10-1~ -9 second, which is the duration of relaxation. Apparently, during relaxation the effective cooling rate can be as high as 1014 K/s (Thompson 1969).

8.3. T H E R M O D Y N A M I C S

OF A L L O Y P H A S E F O R M A T I O N

As described above, the process of ion-mixing, as a whole, takes place under conditions that are far from equilibrium. However, during the relaxation period, the highly energetic state relaxes towards equilibrium. At this stage, equilibrium thermodynamics plays a role in governing the possible evolution path, thus determining the nature (amorphous or crystalline) of the alloy phase formed. It is therefore necessary to calculate the free energy of the competing phases in the alloy system to understand and predict alloy phase formation upon ion-mixing.

8.3.1 Miedema's theory and Alonso's method In this section, the discussion is focused on the amorphous phase formation and the thermodynamics for metastable crystalline phase formation will be discussed later.

201

Ion-Mixing

Generally, the Gibbs free energy A G of a phase is calculated by A G = A H -- T A S, where A H and AS are the enthalpy and entropy of formation, respectively. The entropy usually plays only a secondary role in solid phase formation, and its expression can be taken, as a first approximation, as that of an ideal solid solution, i.e., AS = - - R [ x A ln(xa) + XB ln(XB)], where R is the gas constant, and XA and XB are the atomic concentrations of metals A and B, respectively. For a substitutional solid solution of transition metals, the enthalpy of formation, according to Miedema's model and Alonso's method (Alonso et al. 1990, Niessen and Miedema 1983, Lopez and Alonso 1985), can be considered as a sum of three contributions A H = A H c + A H e + A H s,

(8.2)

where A H c, A H e and A H s correspond to the chemical, elastic, and structural contributions, respectively, and they can be calculated according to the well documented theory and methods. For an amorphous phase, the enthalpy of formation due to both elastic and structural terms is absent. Following the suggestion of Niessen et al. (1988), the enthalpy of the amorphous phase is then given by

(8.3)

AHamorphous-- A H c @OI(XATm,A @ XBTm,B),

where ot is an empirical constant equal to 3.5 J/(mol 9 K) and Tm,i is the melting point of the component i. 8.3.2 I n t e r f a c i a l f r e e energy in the multilayers

Since ion-mixing (IM) begins with multilayered films, which definitely consist of a number of interfaces, the effect of interfacial free energy on the alloying behavior should be taken into account. To include the effect of interfaces, the Gibbs free energy of the initial state of the A-B multilayered films was calculated by adding the interfacial free energy to the ground state of a mixture of A and B in bulk form. As pointed out by Turnbull (1955), the interfacial free energy between two solid phases can be split into geometrical and chemical contributions and can be written as ysAB m

.geo .chem r s s + Yss 9

(8.4) geo For two solids with quite different lattice structures, the geometrical term Yss stands for the energy of a large-angle grain boundary. It is estimated that the grain boundary enthalpy is about 30 percent of the surface energy at 0 K (Murr 1975) and the entropy contribution is neglected. Thus .

.chem , y~sAsB - - 0 . 1 5 x (ys,o + YBB,0) + YSS

(8.5)

where yjS,0, ys,0 are the surface energies of metals A and B, respectively, and Yss'chem is given by Miedema (1978)

B. X. Liu

202 ----d

---'B

--dA'-~ ~ d B - ~ d

A

I I I I I I

A

B

iI I I I I I I

A

Figure 8.2. Schematicconfiguration of A-B multilayered films.

ychem

SS

i12/3

-- AnOinB/(CO, a ),

(8.6)

where Co is a constant being 4.5 • 108, A H ~ in B is the heat of solution of metal A in metal B at infinite dilution (Miedema 1978) and VA is the atomic volume of metal A. According to Gerkema and Miedema (1983), the interfacial free energy of the multilayers, AGmul was calculated as follows"

A Gmul - S f • ysAsB,

(8.7)

where Sf is the surface area occupied by one mole of interfacial atoms. Figure 8.2 shows schematically the configuration of the A-B multilayers. Let us assume that NA and N8 are the numbers of layers of metals A and B, dA,i and dB,i are the thicknesses of the ith layers of metals A and B, respectively; nA and n8 are numbers of interfaces (A on B) and (B on A); AdA,i and AdB,i are the thicknesses of the ith interfaces (A on B) and (B on A), respectively. Obviously, the relationship of n A+ n 8 + 1 = NA +N8 always holds for these four parameters. The interfacial free energy of the multilayers can then be calculated as

(8.8)

AGmul -- OtASf AYBA + OIBSfBYAB,

where SfA and SfB are the surface areas occupied by one mole of A and B atoms, respectively, lYa and OrB are the ratios of interfacial atoms A and B to the total number of atoms in the multilayered films, respectively, and are given by (Zhang and Liu 1995)

S nA AdA,i OlA -- XA Z NA dA,i

Z j B mdB,j ,

or8 - x 8

~NB NB,j

,

(8.9)

Z ; A AriA, i / ~-~NA dA,i and )--~.8 A d s , j / ~--~8 ds,j are the ratios of interfacial A atoms to the total number of A atoms and the ratio of interfacial B atoms to the total number of B atoms, respectively. These terms, multiplied by XA and x s, respectively, give the fractions of A and B interfacial atoms versus the total number of (A + B) atoms where

Ion-Mixing

203

in the multilayers. According to experimental studies on thin films (Clemens 1986) and for simplicity of calculation, the thickness of each interface is assumed to have the same value of 0.5 nm. It should be pointed out that in experimental studies one can always add a few extra interfaces in a sample design to ensure a higher energetic state of the multilayers than that of the phase of interest, so that the assumption of the effect of this thickness can be ignored. Figure 8.3 shows two free energy diagrams calculated by the above methods for two representative systems with negative (Ni-Nb) and positive (Y-Nb) AHf to show a significant difference due to a different sign of the heat of formation. One sees that the free energy curve of the amorphous phase for the Ni-Nb system is concave, i.e., it is always lower than that of an equilibrium mixture of two crystalline metals, suggesting a favored condition for forming glass. The free energy curve of the amorphous phase in the Y-Nb system is convex and is always higher than that of the equilibrium state without considering the interfacial free energy. It is noted that taking into account the interfacial free energy, the energetic sequence of the respective states in the Y-Nb system is changed significantly, while the effect of interface is negligible for the Ni-Nb system. The details of this will be further discussed together with the experimental results later.

8.4. E X P E R I M E N T A T I O N

OF I O N - M I X I N G

8.4.1 Sample design In designing the multilayered samples, the total thickness of the multilayers should match with the range of the irradiating ions. Inert gas ions, such as Xe, Kr, Ne, or Ar, are frequently chosen to be the irradiating ions. The advantages of employing inert gas ions are that the chemical impurity effect is negligible, if the ion dose is below 1016 ion/cm 2, and that the gas ion sources are frequently available easily for the common implanters. As heavier ions induce more efficient mixing than the lighter ones when the other conditions are kept the same, Xe + ions are therefore mostly used. The typical energy of the irradiating ions is several 100 keV. It was found experimentally that if the total thickness of the multilayered film was designed to be the projected range Rp plus the projected range straggling ARp of the irradiating ions, uniform mixing between the layers could be achieved well at an ion dose range of 1014-1016 ion/cm 2. The Rp and ARp for elemental metals can be calculated easily by the LSS theory (Dearnaley et al. 1973). For an alloy target of AxB 100-x, where x refers to the atomic percentage of metal A, the projected range of the alloy can be calculated by the following equation (Townsend et al. 1976):

Rp (Ax B 100-x ) --

Rp (A) Rp (B) yRp(B) + ( 1 0 0 - x)Rp(A)'

(8.10)

where y refers to the weight percentage of metal B in the alloy and can be obtained from x easily. The calculation of the projected range straggling ARp for the alloys is rather complicated. However, considering the errors involved in the calculation itself, as well as in controlling the thickness during deposition, the above equation is also good enough to estimate ARp of the alloy, by substituting ARp and Rp. The total thickness is divided into the thicknesses of metals A and B, according to the desired atomic ratio of the alloy being

204

B. X. Liu

Nb

Y ,

,

.

,

(a)

Muitilayers(17 interfaces)

+60

Amorphous +40

+20

0

&

+20

&

I

9

Multilayers(7 interfaces) (17 interfaces) --"-'--'"'"'""-7~.~ --

-20

----7,

Amorphous

(b) I

Ni

20

9

I

1

40 60 80 Nb concentration (at. %)

J

Nb

Figure 8.3. Calculated free energy diagrams of the Y-Nb (a) and Ni-Nb (b) systems showing the different effects of interfacial free energy on the energetic sequence in the positive and negative heat of formation systems.

studied. Typically, the total thickness of the films of transition metal alloys is about 60100 nm under 300 keV xenon ion-mixing. Another parameter is the fraction of interfacial atoms, which determines the number of metal layers. Each individual layer is always designed to be 5-8 nm or even thinner. The multilayered films are prepared in a vacuum system to deposit alternately the metals A and B onto an inert substrate supporting the film or on a metal matrix. The supporting materials can be SiO2, A1203 or NaC1 single crystals. The system is often a multi-crucible electron gun evaporation system with a vacuum level better than 10 -5 Pa.

Ion-Mixing

205

8.4.2 lon-mixing parameters It was found experimentally that the ion dose required to homogenize the discrete layered films is in the range of 1014 to 1016 ions/cm 2 and varies from system to system. The ionbeam current density should be controlled generally around 1/~A/cm 2, or at most up to 4/zA/cm 2 to avoid overheating of the multilayers during ion irradiation. In operating the implanter, the ion dose, which is measured by the electrical charge, is easy to adjust and control. Ion-mixing thus offers a unique method for investigating the phase transformation process step-by-step. Another important parameter in ion-mixing is the temperature at which the experiment is conducted. As ion-mixing often forms metastable alloy phases, ion-irradiation is generally performed at low temperatures, e.g., at room temperature or down to liquid nitrogen temperature. Nevertheless, the temperature mentioned for the experiments is often the nominal temperature measured at the sample holder instead of the effective temperature of the irradiated films. The effective temperature is an important parameter in the process induced by ion irradiation, and it should be considered when discussing the alloy phase formation mechanism. The stable temperature of the expecting alloy should be considered, i.e., if one expects to form an amorphous alloy with a recrystallization temperature Ta-c (from amorphous-to-crystalline), the effective temperature should of course be kept at a temperature far below Ta-c.

8.4.3 Characterization methods 1-2 MeV He + ion Rutherford backscattering spectrometry (RBS) (Chu et al. 1978) is frequently used to analyze the amount of mixing in the multilayered films. Figure 8.4 shows a typical example for Fe-Mo multilayered films before and after ion-mixing and one can clearly see the information given by the RBS spectra. Besides, XPS, SIMS and Auger spectroscopy are employed to analyze the composition depth profile of the ion mixed films. The structure of the phase is identified mainly by X-ray diffraction and TEM. Fourpoint probe resistivity measurement is a supplemental method to give evidence that the amorphous phase is formed by a crystal-to-amorphous transformation, because the resistivity of the amorphous alloys is often considerably greater than that of their crystalline counterparts.

8.5. A M O R P H O U S

PHASE FORMATION

The maximum possible amorphization range (MPAR) is a decisive parameter in determining the GFA or GFR. When a system has an MPAR exceeding 20%, the amorphous phase can probably be formed in the system. If the MPAR is less than 20%, the related two-phase region is too narrow and can be further decreased by the extension of the solid solubilities from two metal ends under ion-mixing. In other words, it is difficult to form an amorphous alloy (Liu 1987). For an extreme value of MPAR = 0, the phase diagram of the system is of continuous solid solution type with no two-phase region, and GFA and GFR are zero. The following discussion will show that the empirical prediction agrees

206

B. X. Liu

~.~

+

Mo

2.0MeV

SiO2 ~ ~~70.

He

Fe

Mo

.,4 ~a 4~

Ix 1016Xe,./ c m

2

\

as-deposited'

|

. ~ _

500

i_ =.........

600

Channel. Number Figure 8.4. Backscattering spectra of Fe-Mo multilayered films before (solid line) and after room temperature ion-mixing to a dose of 1 • 1016 Xe+/cm 2.

reasonably well with the experimental results for the systems featuring either negative or positive A H f 8.5.1 Glass-forming ability of systems with a negative heat of formation A large body of data confirmed that an amorphous alloy is most likely to form by IM in the two-phase region (Liu et al. 1983), which can be attributed to the competition of crystallization between two crystalline phases, resulting in a frustration of nucleation and growth of the crystalline phase. It is also easy to form amorphous alloys at or in the vicinity of intermetallic compound compositions. This is even more so when the intermetallics have complex crystal structures and therefore cannot form due to a very restricted kinetic condition available in IM (Liu and Jin 1997). Consider the Ni-Nb system (A H i = - 4 2 kJ/mol) as a representative example, whose calculated free energy diagram is shown in Figure 8.3. One sees clearly that the energetic sequence of the phases involved is thermodynamically favorable for the formation of a glassy phase. As the formation of intermetallic compounds is usually hindered kinetically, they need not be considered when discussing the GFA. The GFR was therefore regarded as the composition range in which the free energy of the amorphous phase is lower than

Ion-Mixing

207

that of the solid solutions. The points of intersection between the free energy curve of the amorphous phase with those of the two solid solutions are around 20 and 85 at. % Ni. The GFR in Ni-Nb is therefore from 20 to 85 at.% Ni, which is in excellent agreement with the IM results, i.e. 20-85 at.% Ni, observed by Zhang et al. (1993), as well as very close to the defined MPAR of the system.

8.5.2 Glass-forming ability of systems with a positive heat of formation For the Y-Nb system with a positive A H i, shown in Figure 8.3, the free energy curve of the amorphous phase is usually higher than that of an equilibrium mixture of crystalline metals A and B in bulk form, leading to a situation unfavorable for amorphization. However, as the interfaces in multilayered films are generally regarded as transient layers of mixed A and B atoms in a highly disordered configuration, they contribute extra free energy to the films. With an increasing fraction of the interfacial atoms, it is possible to have the free energy curve of the multilayers intersect with that of the amorphous phase in some specific composition ranges, making it possible to form a glass. These composition ranges depend on the designed fraction of the interfacial atoms of the films (Zhang et al. 1995). Some examples of experimental results are given in the following paragraphs.

8.5.2.1 Amorphous alloys formed within restricted compositions Figure 8.5 shows the calculated free energy diagram of the Ag-Mo system with a A Hf of +56 kJ/mol, the second largest A H f for transition metal binary systems. One sees that the free energy curve of the multilayered films containing 12 layers (or 11 interfaces) intersects with that of the amorphous phase, dividing the diagram into three regions along the composition axis. Obviously, in regions I and III, the amorphous phase is possible to be formed, while in region II, amorphization is hardly to be achieved. Ion-mixing by 200 keV xenon ions at liquid nitrogen temperature did result in the formation of an amorphous alloy at near 80 at.% Mo, and the formation of another amorphous phase co-existed with a new Mo-based f.c.c, phase around 40 at.% Mo (Jin et al. 1995). Such compositional discontinuity in glass formation has also been observed in other systems with a positive A H f (Jin and Liu 1994).

8.5.2.2 Amorphous alloys formed in a broad composition range With a further increase in the fraction of interfacial atoms versus the total number of atoms in the multilayered films, the two separated composition ranges favoring amorphization may become broader and meet each other to become a continuous one. The calculated free energy diagram of the Y-Mo system with A H f = +35 kJ/mol, shown in Figure 8.6, includes three free energy curves for the multilayered films with 9, 11 and 19 layers, corresponding to the fractions of interfacial atoms of 8.1%, 10.0% and 18.1%, respectively. Table 8.1 lists the predicted composition ranges favoring amorphization for these three cases. The Y-Mo multilayered films were prepared accordingly for ion-mixing investigation and the experimental results are in excellent agreement with the prediction (Zhang et al. 1995). These results suggest that through appropriately designing the multilayers, it is possible to synthesize metallic glasses in systems with positive A Hf even with a composition near the equiatomic stoichiometry, where the amorphous phase frequently has the

208

B. X. Liu

_~-

O

6O

>'

40

"

i

~

I

'

I

'

I

E

,

K,.,

ii

~

2o

~

10

L_

I

~

v

30

'

f.c

12 L a y e r s

)

-

rphous

(.9 0

1

l

Mo 10

1

,

1

30

50

9

1

70

j.

~'~

I

90

Ag

Ag C o n c e n t r a t i o n Figure 8.5. Calculated free energy diagram of the A g - M o system.

Table 8.1 Predicted and experimental results of amorphous phase formation in the Y-Mo multilayered films upon 200 keV xenon ion mixing at room temperature, n" number of interfaces in the multilayers; o~" fraction of interfacial atoms in the multilayers. Composition

Y80 Mo20

Y60 Mo40

Y50 Mo50

Y40 Mo60

Y20 Mo80 Yes

n = 18

Predicted

Yes

Yes

Yes

Yes

(a = 18.1%)

Ion mixing

Yes

Yes

Yes

Yes

*

n = 10

Predicted

Yes

Yes

No

No

Yes

(or = 10%)

Ion mixing

Yes

Yes

No

No

Yes

n =8

Predicted

Yes

No

No

No

Yes

ot = 8.1%)

Ion mixing

Yes

No

No

No

Yes

* Experiment not conducted.

highest free energy and is most difficult to form. It is worth mentioning that the IM results listed in Table 8.1 include both "yes" and "no" and that they strongly depend on the correlation between interfacial fraction and alloy composition. These results also support the argument that IM induced amorphization can really be attributed to the interfacial free energy, but not to the irradiation energy. The irradiation energy has exactly the same effect on all the multilayered films and does not depend on alloy composition. In the past 15 years, nearly 100 binary metal-metal systems, according to published results, have been studied by ion-mixing of multilayers. Table 8.2 lists these systems in alphabetical order of the chemical symbols of the constituent metals for convenience of searching.

Ion-Mixing

209

Table 8.2 Binary metal-metal systems that have so far been studied by ion mixing AHf

Any

<0

System

Ref.

Remarks

A1-Au A1-Co A1-Fe A1-Mn A1-Mo A1-Nb A1-Ni A1-Pd A1-Pt A1-Ta A1-Ti A1-Y A1-Yb Au-Ta Au-Ti Au-V Co--Gd Co-Hf Co-Mo Co-Nb Co--Ta Co-Tb Co-Ti Co-Y Co-Yb Co-Zr Cu-Y Cu-Zr Er-Ni Fe-Gd Fe-Mo Fe-Nb Fe-Ta Fe-Tb Fe-Ti Fe-Y Fe-Zr Hf-Ni Mo-Ni Mo-Ru Mo-Zr Nb-Ni Ni-Ta Ni-Ti Ni-Y Ni-Yb Ni-Zr Ru-Ti Ru-Zr

Schmid and Ziemann (1985) Hung and Mayer (1985) Rauschenbach and Hohmuth (1987) Knapp and Follstaedt (1987) Meissner et aL (1987) Liu et al. (1983) Hung et al. (1983) Hung et al. (1983) Hung et al. (1983) Meissner et al. (1987) Cheng (1986) Liu et al. (1991) Ding et al. (1992) Pan et aL (1996) Liu et al. (1982) Tsaur et aL (1991) Yan et aL (1986) Anderson et al. (1990) Liu et al. (1983) Zhang and Liu (1994) Zhang and Liu (1994) Yan et al. (1987) Hung and Mayer (1985) Liu et al. (1991) Liu et aL (1991) BCttiger et al. (1987) Liu et aL (1991) Liu et al. (1991) Liu et al. (1983) Yan et al. (1987) Liu et aL (1987) Zhang and Liu (1994) Zhang (1995) Yan et al. (1987) Brenier et al. (1987) Liu et al. (1991) BCttiger et al. (1987) van Rossum et al. (1983) Liu et al. (1983) Liu et al. (1983) Jin and Liu (1995) Zhang et al. (1993) Liu and Zhang (1994) Cheng et aL (1984) Liu et al. (1991) Liu et al. (1991) BCttiger et al. (1989) Cheng et al. (1984) Liu (1985)

Several glasses have been formed in many systems. Or, at least one glass has been formed in some systems.

B. X. Liu

210 Table 8.2 Continued AHf

System

Ref.

Remarks

Any < 0

Au-Cu Cu-Rh Mo-Nb Fe-W

Rauschenbach (1986) Peiner and Kopitzki (1988) Liu (1982) Goltz (1982)

No glass has so far been formed.

Cu-Ir Ti-Zr

Peiner and Kopitzki (1988) Liu (1982)

Ag-Cr Ag--Cu Ag-Mo Ag-Nb Ag-Ni Ag-Ta Ag-V Au-Co Au-Ir Au-Mo Au-Os Au-Ru Au-W Co-Cu Cu-Fe Cu-Mo Cu-Nb Cu-Ta Fe-Nd Hf-Nb Hf-Ta Mo-Y Nb-Ti Nb-Zr Ta-Ti Ta-Y Ta-Zr Ti-Y Y-Zr

Liu et al. (1987) Liu et al. (1986) Jin et al. (1985) Jin and Liu (1996) Liu et al. (1986) Andersen et al. (1990) Andersen et al. (1990) Tsaus et al. (1981 ) Peiner and Kopitzki (1988) Pan and Liu (1995) Peiner and Kopitzki (1988) Meissner et al. (1987) Meissner et al. (1987) Liu et al. (1987) Huang and Liu (1987) Chen and Liu (1997) Andersen et al. (1990) Tsaur (1980) Yan et al. (1987) Jin et al. (1995) Liu et al. (1996) Zhang et al. (1995) Chen and Liu (1988) Jin and Liu (1994) Chen et al. ( 1991 ) Zhang and Liu (1995) Jin and Liu (1997) Liu et al. (1991 ) Liu et al. (1998)

Several glasses have been formed in many systems. Or, at least one glass has been formed in some systems.

Ag-Co Ag-Fe Ag-Ir Ag-Os Ag-Rh Ag-Ru Ag-W Au-Fe Au-Ni Au-Rh Cr-Cu Cu-Os Fe-Yb

Liu et al. (1998) Andersen et al. (1990) Peiner and Kopitzki (1988) Peiner and Kopitzki (1988) Peiner and Kopitzki (1988) Peiner and Kopitzki (1988) Hiller et al. (1989) Tsaur et al. (1981 ) Tsaur et al. (1981) Peiner and Kopitzki (1988) Shin et al. (1983) Meissner et al. (1987) Liu et al. (1991)

No glass has has so far been formed.

AHf = 0

AHf = 0

AHf

> 0

AHf>O

Several glasses have been formed. No glass has so far been formed.

Ion-Mixing

211

18.1% +40 10.0% 8 1%

*20

1

0 Bulk L..

,I,,

Y

20 Mo

State 9

40

J

,

60

Concentration

,,

&

80

Mo

(at%)

Figure 8.6.

Calculated free energy diagram of the Y-Mo system. The numbers of 8.1%, 10.0% and 18.1% are for the fractions of interfacial atoms versus the total atoms in the multilayers, respectively.

8.5.2.3 Nominal and intrinsic glass-forming ability From both practical and theoretical points of view, the GFA and its prediction are of vital importance. The originally defined GFA was only to classify the systems into two categories of glass-forming and non-glass-forming ones in a qualitative way and the classification mostly depended on glass preparation technique employed. However, by employing powerful techniques, e.g., IM, many new metallic glasses have been obtained in those systems considered as non-glass-forming ones by LMQ. Apparently, the definition or the physical meaning of GFA needs to be extended. Firstly, GFA of a system should reflect qualitatively whether a metallic glass can be formed or not in the system. As mentioned above, the GFA determined by more powerful techniques frequently possessing higher cooling rates is expected to be closer to the real ability of the system. In other words, from a physical point of view, there exists an intrinsic GFA reflecting the capability of the system itself to form a metallic glass and it does not depend on the technique employed to produce the glass. Furthermore, it is necessary to have a quantitative measure of GFA to compare the capabilities of glass formation among the glass-forming systems. The experimentally determined GFR, in the author's view, can be taken as a quantitative measure of GFA. Apparently, the larger the GFR, the greater the GFA. It is true that if one employs a technique with a higher cooling rate, one may obtain the glassy phase in a broader composition range than that observed by a technique with a lower cooling rate. In this sense, the experimentally observed GFR can only be considered as a nominal GFA. Obviously, the greater the GFR observed by a specific technique, the closer the GFR is to the intrinsic GFA. Consequently, to approach the intrinsic GFA of a system, one should

B. X. Liu

212

Table 8.3 Some examples of metallic glasses formed in binary metal systems with a negative heat of formation by ion-mixing No.

System

AHf (kJ/mol)

MPAR (%)

Glass-forming range

1 2 3 4 5

Zr-Ru Ni-Nb Co-Nb Nb-Fe Co-Mo

-86 -44 - 38 -23 -7

90 86 86 85 76

25-75 at.% Ru 20-85 at.% Nb 20-77 at.% Nb 20-78 at.% Fe 35-77 at.% Mo

Table 8.4 Some examples of metallic glasses formed in binary metal systems with a positive heat of formation by ion-mixing. (Total thickness of the multilayered films is 40-50 nm.) No.

System

AHf (kJ/mol)

MPAR (%)

Number of interfaces

Glass-forming range

1 2 3 4 5

Zr-Nb Ag-Mo Ti-Ta Y-Nb Y-Mo

+6 +56 +3 +44 +35

70 100 100 100 100

7 11 15 17 18

at 12 or 60 at. % Nb at 30 or 80 at.% Mo 74-89 at.% Ta 28-44 at.% Nb 20-80 at.% Mo

look for a technique capable of generating cooling rates as high as possible. At present, IM can be considered as the technique capable of revealing the m a x i m u m GFA a m o n g the currently available glass producing techniques. The next issue is how to correlate the n o m i n a l G F R observed by IM and the intrinsic GFA of the system. Based on IM study, a parameter, namely MPAR, has been proposed by the author to predict the m a x i m u m possible G F R of a system. According to the definition, M P A R = 1 0 0 % - ( o r + 13), where c~ and/~ stand for the m a x i m u m solid solubilities at the two metal ends (Liu et al. 1987, Liu 1987). As IM can frequently form some supersaturated solid solutions, which would actually reduce the total width of the two-phase regions (Liu 1986), an empirical formula is therefore suggested for correlating the possible G F R and MPAR, i.e. G F R = 0.7MPAR, and it is reasonably compatible with the IM results in either positive or negative A H f systems, as shown by those examples listed in Tables 8.3 and 8.4. W h e n MPAR=0, corresponding to a system with a continuous solid solution phase diagram, G F R = 0 and no glass has so far been obtained in such systems, which are therefore considered as non-glass-forming ones.

213

Ion-Mixing

8.6. FORMATION OF METASTABLE CRYSTALLINE ALLOYS

8.6.1 Structural classification of the metastable crystalline phases Up-to-now, a number of metastable crystalline phases (MX) have been formed by ionmixing in both negative and positive A H i systems and they are classified into 5 structural categories (Liu et al. 1995), i.e., supersaturated solid solutions, the h.c.p.-I phases in h.c.p.- or f.c.c.-based alloys, the f.c.c.-I phases in h.c.p.-based alloys, the h.c.p.-II and f.c.c.-II phases in b.c.c.-based alloys. 8.6.1.1 Solid solutions

Ion-mixing of multilayered films can extend the solid solubility of the binary metal alloys resulting in the formation of a supersaturated solid solution always having the same simple crystal structure as the major alloying metal. Most of these alloy phases are identified as solid solutions. Their lattice constants agreed quite well with those predicted by Vegard's law, and they are therefore substitutional solid solutions (Liu and Jin 1987).

8.6.1.2 h.e.p.-I and f.e.e.-I phases A metastable crystalline phase with a similar h.c.p, structure has been formed by ionmixing in five binary (A-B) metal systems (Co-Au, Ti-Au, Co-Mo, Ni-Mo, and Ni-Nb) (see Table 8.5) where A refers to the first entry in the alloy designation. The MX phases were formed in A-rich multilayered films with an overall composition near A3B. These MX phases were named as h.c.p.-I, since they have all about the same spacing (dcpp) of the close packed planes (Table 8.5). Their formation was attributed to the valence electron effect. For a close-packed hexagonal structure, the maximum number of electron states per atom, n, in the Jones zone (including doubling for spin) can be calculated by the following equation (Jones 1960):

n=2-~

(a)2[ 1(a)2] c

1-~

c

'

(8.11)

where a and c are the lattice constants of the h.c.p, structure. According to this calculation, the n values of these five phases are almost identical, i.e., n -- 1.73 to 1.74, which is very close to the value of 7/4 = 1.75, corresponding to the well-defined Hume-Rothery 7/4 electron compound. These A3B alloy phases can therefore be considered as metastable electron compounds (Liu 1983). Table 8.6 shows that MX f.c.c.-I phases have formed in some h.c.p.-metal-ricia multilayered films by ion-mixing. As both h.c.p, and f.c.c, structures are built up by the stacking of close-packed atomic planes differing only slightly in the stacking order, one can easily transform into another. For example, sliding of atoms on the planes parallel to the (0002)hcp plane along (1 i00)hcp directions by a vector of 1/3(1 i00)hcp can transform the h.c.p, structure into f.c.c., and the reverse procedure transforms the f.c.c, structure to h.c.p (Liu and Zhang 1994a).

B. X. L i u

214

T a b l e 8.5 Structure data of the h.c.p.-I phase formed in five binary systems by ion-mixing. These phases were considered as metastable electron compounds. M X phase System Co-Au

Constituent metals

lattice constants

dcpp

lattice constants

dcpp

(nm)

(nm)

(nm)

(nm)

h.c.p,

d002 = 0.260

Co h.c.p,

d002 = 0.2023

a = 0.324, c = 0.521

a = 0.254, c = 0.407

c/a = 1.61

c/a = 1.59

Ref. Liu and Nicolet (1983)

Au f.c.c.

Ti-Au

h.c.p,

d002 = 0.263

a = 0.408

dl 11 = 0.2355

Ti h.c.p,

d002 = 0.234

a = 0.327, c = 0.526

a = 0.259, c = 0.469

c/a = 1.61

c/a = 1.59

Liu et al. (1982)

Au f.c.c.

Co-Mo

h.c.p,

d002 = 0.261

a = 0.408

d l l l = 0.2355

Co h.c.p,

d002 = 0.2023

a = 0.327, c = 0.522

a = 0.254, c = 0.407

c/a = 1.60

c/a = 1.59

Liu et al. (1983b)

M o b.c.c.

Ni-Mo

h.c.p,

d002 = 0.261

a = 0.3147

d110 = 0.2225

Ni f.c.c,

d l l l = 0.2034

a = 0.325, c = 0.521

a = 0.352

c/a = 1.60

Mo b.c.c. a = 0.3147

Ni-Nb

h.c.p, a = 0.328, c = 0.522

d002 = 0.261

Liu et al. (1983b)

d110 = 0.2225

Ni f.c.c. a = 0.352

d l l I = 0.2034

Kung et al. (1983)

Nb b.c.c.

c/a = 1.59

a = 0.3307

d110 = 0.233

Table 8.6 f.c.c.-I M X phases formed in Co-, Y-, Zr-, and Hf-rich multilayered films by room temperature 200 keV xenon ion mixing. System and composition

Mixing dose

Lattice parameter (nm)

Co94Mo6 ~ Co82Mo18

3 x 1015 X e + / c m 2

0.355 ~ 0.360

Co98Ta2 ~ Co80Ta20

3 • 1015 X e + / c m 2

0.355 ~ 0.365

Y73Mo27 ~ Y83MoI7

7 • 1015 X e + / c m 2

0.510

Y85TaI5 ~ Y75Ta25

1 x 1015 X e + / c m 2

0.518

Y25Zr75 ~ Y85ZrI5

3 • 1015 X e + / c m 2

0.467

Zr68Nb32 ~ Zr88Nb12

7 • 1014 X e + / c m 2

0.480

Hf75Nb25 ~ H f 9 0 N b l 0

5 • 1014 X e + / c m 2

0.455

Ion-Mixing

215

Table 8.7 h.c.p.-II and f.c.c.-II MX phases formed in Nb-, Ta-, Mo-rich multilayered films by room temperature 200 keV xenon ion-mixing. A-B

f.c.c.-b.c.c.

h.c.p.-b.c.c.

b.c.c.-b.c.c.

System and composition

h.c.p.-II

f.c.c.-II

Ni23Mo77 Ni20Nbs0 Ni35Ta65 Ag30Mo70 Ag33Nb67 Co20Mo80 Co23Nb77 Co30Ta70 Y22Mo78 Y25Ta75 Zr7Mo93 Zrl9Nbs1 Zr20Ta80 Hf17Nb83 Hf30TaT0 Fe24Mo76 Fe20Nbs0 Fe20Ta80

Yes

Yes

Yes No No No Yes Yes No Yes Yes Yes Yes No No No Yes

Yes Yes Yes Yes Yes Yes Yes Yes Yes Yes Yes Yes Yes Yes Yes

Yes

Yes

No

Yes

8.6.1.3 h.c.p.-H and f.c.c.-H phases based on b.c.c, metals In some Nb-, Ta-, and Mo-rich multilayered films, two MX phases, labeled h.c.p.-II and f.c.c.-II, were formed by ion-mixing, and their compositions were around AB3-4, in which B stands for the b.c.c metal. (see Table 8.7). Some of the systems have negative AHf, while combining with metals like Fe, Co and Ni and some have positive AHf when alloying with Ag, Y, Zr and Hr. Carefully performed experiments revealed that the h.c.p.II phase was observed as an intermediate phase in the process of forming the f.c.c.-II phase. Thus a two-step transition of b.c.c. --~ h.c.p. -+ f.c.c, was proposed and named as reverse martensitic transformation. First b.c.c, phase transforms to a h.c.p structure through a shearing mechanism. During shearing, the (111)bcc acted as the habit plane, while [ l l l]bcc (or [2110]hcp) acted as the shearing axis. The second step is to transform the h.c.p, phase to the f.c.c, phase. The shear can be achieved by sliding the atoms on planes parallel to the (0002)hcp along (1100)hcp directions by a vector of 1/3 (l l00)hcp. Although the relaxation period of ion-mixing is very limited, it is thought that it could meet the requirements of a shearing mechanism. The lattice parameters of h.c.p.-II and f.c.c.-II phases were deduced according to the two-step transition and these are in good agreement with the experimental results (Zhang 1995, Liu and Zhang 1992a).

B. X. Liu

216

8.6.2 Free energy calculation of the MX phases A metastable crystalline phase is of ordered structure, and so its entropy term can be neglected. The free energy of the MX phase is therefore A GMX = A H~I X -~- A Hi,ix -~- A H ~ x .

(8.12)

The expression of the chemical term is as follows A HI,cix = A HampX A I/" " A2/3 f A B ,

(8 13)

in which VA is the atomic volume of component A in the alloy, A Hamp is an amplitude concerning the magnitude of the electron redistribution interaction and is a constant for one specific metal system, faB is a function which accounts for the degree to which atoms of type A are surrounded by atoms of type B. In the calculation of faB, an empirical constant y is used to describe the short-range-order difference of solid solutions, amorphous phases, and ordered compounds, for which it is usually taken to be 0, 5, and 8, respectively. As an MX phase, except the solid solutions, is commonly considered as compoundlike, the value of y is therefore taken as 8 in the calculation. The elastic term of the MX phase can be derived from the equation A H ~ I X - - X A X B [ X A A H Me X ( B i n A )

where

-'F X B A

e , HMX(AinB)]

(8 . 14)

is the elastic contribution to the heat of solution of i in j, which can be calculated from the elastic constants of the constituent metals. The expression for the structural term is written as AOMx(iinj)e

A H ~ x = E(Z)

-

xAE(ZA)

--

xBE(ZB),

(8.15)

where E(Z), E(ZA) and E(ZB) are the structural stabilities of the alloy phase and the pure A and B metals, and Z, ZA and ZB are the mean number of valence electrons of the alloy phase and the number of valence electrons of pure metals A and B, respectively. For an MX phase at a composition in the vicinity of an intermetallic compound, the elastic term should be included. This term was from the difference in structures between the MX phase and that of the intermetallic compound. This is for the case that the MX phases were formed at A3B stoichiometry. Another situation is that the MX was formed at compositions near the eutectic point, where no equilibrium compound exists. In this case, the MX phase can be considered as ordered, so as the elastic term can be relaxed down to zero because of ordering. This situation corresponds to the formation of the MX phases with a composition of AB3. In short, with the above assumption, the formula for calculation for the A3B MX phase AGMx -- A G ~ x 4- A G ~ x + A G ~ x ,

(8.16)

and for the AB3 MX phase AGMx-

AG~x + AG~x.

(8.17)

Ion-Mixing

217

~" 12001-1 bcc(Mo) o.. I]

i) 91(P~

~ .

80(,[]

~

MoN

~- 6o0[/

i3

Y

4(t(t U

. . . . . . . . . 2(t 40 6(t

Mo

8()

Ni

(a)

~

"~ +8 E

bcc solution

fcc solution

..x +4

=

0

r

hcp-ll ~

~,, -4

~kJ "~

fcc-ll --

Mo

(b)

,A ....

9

20

9

J

40

9

,~

60

,

9

9

9

84)

Ni

Ni concentration (at.%)

Figure 8.7. Equilibrium phase diagram of the Ni-Mo system; (b) Calculated free energy diagram of the Ni-Mo system.

The free energy curves of the MX phases formed in the Ni-Mo system are calculated, using the above equations, as an example. Figure 8.7(a) presents the equilibrium phase diagram. Figure 8.7(b) shows the constructed free-energy diagram of the system, in which the free-energy curves of all the concerned phases are included. One sees that the diagram is a little complicated because some newly formed MX phases are also included. Before employing the diagram to interpret the alloy phase formation, an important question to be considered is if the calculated results are relevant, since the model used in the calculations is still semi-quantitative. The author has proposed an alternative way for checking the relevance of the calculation, i.e., the steady-state thermal annealing of the multilayered films at some important compositions. The idea was to see whether the phase evolution in

218

B. X. Liu

the films upon thermal annealing matches with that deduced from the calculation. Let us consider two examples in which the phase appearance in the samples is as follows:

300 ~ 450 ~ 650 ~ 800 ~ Niv5Mo25 30min > Am~ min > hcp-I20 min > Ni + hcp-II30 mi n > Ni + fcc-II

350~ 500~ 600~ Ni25Mo75~ > Amor.~ > hcp-II~ > fcc-II 60 min 30 min 30 min Apparently, the phase evolution sequences of these two reactions are in good agreement with that deduced from the calculated free energy diagram, thus confirming the applicability of the Miedema's model in this respect (Zhang and Liu 1994a). The author's group has also performed similar annealing experiments for several positive A H f systems and these results are also in support of the calculations (Chen and Liu 1996). The use of calculated free energy diagrams can explain the alloy phase formation upon IM, as discussed in the previous sections by some representative examples.

8.7. I N T E R F A C E - G E N E R A T E D

S O L I D - S T A T E V I T R I F I C A T I O N IN S Y S T E M S W I T H A

P O S I T I V E H E A T OF F O R M A T I O N

In 1983, Schwarz and Johnson reported for the first time that solid-state reaction (SSR) of multilayers consisting of two different crystalline metals (La and Au) can lead to the formation of an amorphous alloy. They argued that spontaneous vitrification was expected to take place in metal systems with a negative A H i and one metal diffusing anomalously fast in the other. However, the author's group has found very recently that in some binary metal systems with positive A H f , vitrification could be achieved by SSR, e.g., in the HfTa (Liu et al. 1996), Y-refractory metals (Chen and Liu 1997) and Cu-Ta (Pan et al. 1998) systems. Consider the Y-Nb system as an example, Figure 8.8 shows the X-ray diffraction patterns of the Yv2Nb28 multilayered films before and after thermal annealing. It should be emphasized that similar solid-state vitrification was also observed by author's group very recently in some other systems with a positive A H f (Chen and Liu 1997a). To test whether a large size difference is a decisive prerequisite in solid-state amorphization, the Au-Ta system with a negative A H f o f - 4 9 kJ/mol (Pan et al. 1995) was selected as the atomic radius ratio and the volume ratio between Au and Ta are 1.008 and 0.936, respectively. When the Auz3Ta77 multilayers were heated up continuously at a rate of 5 K/min to 200~ partial amorphization was observed. After maintaining at 200~ for 30 min, the annealing temperature was raised to 350~ and kept for 30 min. The structure of the multilayers turned entirely into an amorphous state. When the temperature was further raised to about 650~ and kept for 20 min, diffraction lines from metals Au and Ta emerged again, indicating crystallization of the Auz3Ta77 amorphous alloy. These results suggest that solid-state vitrification can also be achieved in a system with a small atomic size difference.

Ion-Mixing

219

J

.1500

(a) .

.

.

.

0i 0

.,,.i

I,=,

.e

(b)

~0

i

30

|

I

40 50 20 (degree)

60 -

-

Figure 8.8. X-ray diffraction patterns of the Y72Nb28 multilayered films with 121 layers and a total thickness of 350 nm before and after thermal annealing: (a) as-deposited; (b) amorphized by annealing at 300~ for 1.5 hours.

8.8. C O N C L U D I N G R E M A R K S

(1) In the past 15 years, study of IM of multilayered films scheme has provided extensive data, which promoted significantly the non-equilibrium processing of materials. IM has so far synthesized a great number of new amorphous as well as new MX alloys. It is anticipated that this technique is capable of synthesizing more new metallic glasses in those not yet studied systems known as non-glass-forming ones. (2) Miedema's theory was employed for free energy calculation and steady-state thermal annealing results showed that the calculated energetic sequence of the related phases was relevant, at least in its outline, suggesting the applicability of Miedema's theory in this respect. (3) Taking into account the interfacial free energy, a unified thermodynamic model was proposed for predicting the possibility of metallic glass formation by IM in binary metal systems with either negative or positive AHz, and for a reasonable explanation for the formation of MX phases. (4) The interface concept was also tested by solid-state reaction experiments in some positive A Hf systems and it turned out that metallic glasses were indeed obtained by thermal annealing in the properly designed multilayered films with enough fraction of interfacial atoms. These results offer a neat evidence demonstrating the decisive role of the interfacial free energy in amorphization taking place in the multilayered films.

220

B. X. Liu

(5) Kinetically, due to the very restricted conditions available in IM for a crystalline phase to nucleate and grow, only simple-structured phases can be formed and they feature a close crystallographic correlation with their matrix. (6) Based on the IM and SSR studies, a new glass-producing technique is developed, i.e., the multilayer technique, in which the multilayered films are firstly designed according to the interfacial free energy calculation and then either by thermal annealing or ion irradiation. (7) Further studies are still needed in this area, especially in the systems with a positive A Hf. For instance, detailed growth of the amorphous phase, local microstructures and the abnormal properties of the new metallic glasses formed in such systems are highly interesting topics to be investigated by IM as well as SSR technique.

ACKNOWLEDGMENTS

The author is grateful to his current and previous students for their continuous contributions to this research project, especially to Mr. Y.G. Chen for his substantial assistance in writing and finalizing this chapter. Special thanks go to the Editor of this book, Professor C. Suryanarayana, for his critical reading of the manuscript of this chapter and for many helpful suggestions. The continuous financial support from National Natural Science Foundation of China under several projects and by the Administration of Tsinghua University are greatly appreciated.

REFERENCES Alonso, J. A. and Simozar, S. (1983) Solid State Comm., 48, 765. Alonso, J. A., Gallego, L. J. and Simozar, S. (1990) Nuovo Cimento, D12, 587. Andersen, U. A., BCttiger, J. and Dyrbye, K. (1990) Nucl. Instru. Meth. Phys. Res., B51, 125. Bcttiger, J., Pampus, K. and Torp, B. (1987) Nucl. Instr. Methods Phys. Res., B19/20, 696. Bcttiger, J., Dyrbye, K., Pampus, K. and Poulsen, R. (1989) Phil Mag., A59, 569. Brenier, R., Thevenard, P., Capra, T., Perez, A., Treilleux, M., Romana, L., Dupuy, J. and Brunel, M. (1987) Nucl. Instr. Methods Phys. Res., B19/20, 691. Carter, G. and Grant, W. A. (1976) Ion Implantation of Semiconductors. (Edward Arnold Ltd., London). Chen, Y. G. and Liu, B. X. (1996) Appl. Phys. Lett., 68, 3096. Chen, Y. G. and Liu, B. X. (1997a) J. Alloys and Compounds, 261, 217. Chen, Y. G. and Liu, B. X. (1997b) Nucl. Instr. Meth. in Phys. Res., B127/128, 145. Chen, Y. G. and Liu, B. X. (1997c) J. Appl. Phys., 82, 3815. Chen, Y. G. and Liu, B. X. (1988) Mater Sci. Eng., B ( to be published). Chen, Y. G., Zhang, Q. and Liu, B. X. (1997) Nucl. Instr. Methods Phys. Res., B124, 523. Cheng, G. A. (1986) Thesis, Tsinghua University, Beijing, China. Cheng, Y. T., Johnson, W. L. and Nicolet, M.-A. (1984) in Proceedings of MRS Fall Meeting, Boston, MA. Chu, W. K., Mayer, J. W. and Nicolet, M.-A. (1978) Backscattering Spectrometry. (Academic Press, New York). Clemens, B. M. (1986) Phys. Rev., B33, 7615. Davies, H. A. (1983) in Amorphous Metallic Alloys, ed. Luborsky, EE. (Butterworths, London), p.13. de Boer, F. R., Boom, R., Miedema, A. R., Niessen, A. K., and Mattens, W. C. M. (1989) Cohesion in Metals: Transition Metal Alloys. (North-Holland, Amsterdam). Dearnaley, G., Freeman, J. H., Nelson, R. S., and Stephen, J. (1973) Ion Implantation. (North-Holland Publishing Company, Amsterdam). Ding, J. R., Che, D. Z., Zhang, H. B., Tao, K. and Liu, B. X. (1992) Appl. Phys. Lett., 60, 944. Gerkema, J. and Miedema, A. R. (1983) Surf. Sci., 124, 351. Giessen, B. C. and Whang, S. H. (1980) J. Physique, C8, 95.

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