METALS AND MATERIALS, Vol. 6, No. 1 (2000),pp. 17-24

Alloy Design of Intermetallic Dispersion Strengthened Aluminum Systems by Mechanical Alloying for High Temperature Applications Sung I1 Park and Hyuck Mo Lee Department of Materials Science and Engineering Korea Advanced Institute of Science and Technology 373-1 Kusong-dong, Yusong-ku, Taejon 305-701, Korea Good interfacial bonding between the matrix and the second phase is very important and may be guaranteed by the low lattice mismatch between them. In this study, trialuminide intermetallics such as tetragonal D022 and D023 with A13M (M=Ti, V, Zr) or cubic L12 with A16~Mng(Ti,Zr)25 were chosen as second phases to be dispersed on the basis of MEAM (Modified Embedded Atom Method) simulations. Specific compositions were determined based on thermodynamic calculations. Alloys were mechanically alloyed using AI powders and intermetallic powders that were separately fabricated. Hot pressed samples under vacuum were analyzed in terms of phase formation during heat treatment and were tested in compression at room temperature and 693K. Specimens showed enhanced strength with a yield strength of 800-920MPa at room temperature and 190-260MPa at 693K. Observation of no phase change during heat treatment indicated high thermal stability of the microstructure.

Key words:A1 composite, trialuminide, EAM, thermodynamic calculation, mechanical alloying

1. I N T R O D U C T I O N In order to enhance mechanical properties of A1 alloys at low or high temperatures, second phases that strengthen the A1 matrix need to satisfy several requirements such as high strength, fine scale, homogeneous distribution, thermal stability so as not to transform to other undesirable phases and good bonding at the interface [1]. Trialuminide phases, A13M (M is a transition element), with structures of D022, D023 or L12 have been regarded as candidates for dispersion. Aside from a high melting point and high strength, the structures and lattice parameters are similar to those of A1, which lower the lattice mismatch of the interface with the A1 matrix [1-4]; as a result, they form good bonding at the interface. Methods of fabrication to disperse these intermetallic compounds into the A1 matrix can be classified into two types: RSP (Rapid Solidification Process) [5-8] and MA (Mechanical Alloying) [9-12] with alloys made by the former method, it is hard to attain a large and optimum 20-25% volume fraction of the dispersed phase [1]. The alloy made by the latter maintains fine dispersion and apparently achieves an ideal distil-

bution. In conventional MA, elemental powders are mixed together to induce phase transformation toward the formation of the second phase, which cannot produce a nonequilibrium second phase such as trialuminide L12. Therefore, this study devises a new concept in order to widen the selection range of trialuminide intermetallic second phases including non-equilibrium phases, and to control volume fraction without considering phase equilibria. The production of intermetallic powders precedes MA and then intermetallic powders are mixed with A1 powders. In this work, elements and their compositions are selected to control the structure of intermetallic compounds through the EAM (Embedded Atom Method) [13,14] and thermodynamic calculations in the spirit of alloy design. A1 composites designed and fabricated in this way are investigated by the phase analysis method and by the compression test. 2. A L L O Y D E S I G N C O N C E P T The selection criteria of the second phase for strengthening in the A1 matrix lie in the resistance to deformation, thermal stability and interracial properties. The three fac-

This article is basedon a presentation made in the symposium"The 1~KIM-JIMJoint Symposium:High Strength Ratio Aluminum Alloys",held at Inha University,Inchon, Korea, October22. 1999 under the auspicesof The Korean Institute of Metals and Materials and The Japanese Institute of Metals.

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Sung II Park and Hyuck Mo Lee

tors depend on the structure and composition of the second phase: thermal stability may be easily estimated by thermodynamic calculations, resistance to deformation is indirectly measured by Young's modulus or bulk modulus, and interfacial properties can be estimated by the interfacial energy and by the lattice mismatch, although the real strength of atomic bonding at the interface is difficult to calculate and solid-solid interface energies are also hard to obtain. Therefore, theoretical calculations such as EAM are useful to surmise necessary physical properties and to design alloys. The compositional effect on phase equilibria is provided by thermodynamic calculations. A13M intermetallic compounds are selected as the second phase because of high thermal stability and the estimation of the resistance to deformation and interracial properties are numerically performed through atomic energy calculations by EAM modeling.

3. EAM AND THERMODYNAMIC CALCULATIONS In this study D0:: or D0:3 A13(Ti, V, Zr) and LI2 A166Mn9 (Ti, Zr)2~ based on the ordered fcc structure as shown in Fig. 1 were selected as the second phase. D022 is a stable configuration for A13Ti and A13V as is D02~ for A13Zr. The stable configuration of the mixed type, AI~(Ti, V, Zr), depends on the ratio of transition elements, which may be a single phase of D022 or D0:3 or two phases of D022+D023. Table 1 shows the lattice parameters of each phase. According to Fig. 1 and Table 1, the fact that the intermetallic compounds have a similar to those of A1 may impart the second phase enhanced interfacial properties.

Table 1. Lattice parameters of A1 and A13M phases Phase

Structure

AI AI3Ti AL V ALZr

Fcc D022 D022 D02~

Lattice Parameter (•) a=3.849 a=3.780 a=4.013

~=4.05 at 698K c=8.610[3] c=8.32213] c=17.321 [3]

To explore the energy for the separating (111) interface between intermetallic compounds and the A1 matrix, atomic simulations have been made by EAM. A large value of interface energy means strong bonding which restricts the creation of and resists the propagation of cracks through the interface. Good creep resistance is also expected. Calculated values were compared to the (111) interface energy of A1. EAM calculations exclude the V addition because the simulation of a system with more than three components has not proved itself to be reliable and the main concern of this study lies in Ti and Zr. The topic of V addition leaves room for further study. EAM is a new advanced energy model constructed to overcome the limit of the pair potential model [ 13,14]. The following Eq. 1 represents the energy of an indexed j atom, which consists of pair energy ~ j ( l ~ j ) and embedding energy

F~(0k ). Ek = Fk(P~) + 21 E ~k,j(Rk,J ) k(~.i)

(1)

The embedding energy of the host k site, the core of the EAM spirit, is a function of the total background electron density (Pk) that is the sum of electron density contributed by surrounding atoms at the j site. The EAM model is especially useful for describing energy and the structure of defects such as interface and grain boundary since simulations reveal the exact relationship between bonding energy and the coordination number. This study has adopted an advanced EAM model, MEAM (Modified EAM) [15] which overcomes the problem of EAM being applicable only to metallic bonding, as it considers angle dependent terms to mimic covalent bonding. A brief explanation follows. The embedding energy takes a logarithmic form in the MEAM model as in Eq. 2.

(2) k ~e j

2[J e ff

The partial electron density contributed by surrounding atoms takes an exponential term as in Eq. 3. Fig. 1. Atomic configuration of various structures: (a) LI2, (b) D0:~ and (c) D0~3.

9j(h)(R) = exp(-b*) b= ~tO(R/RO- l )

h=0-3

(3)

Alloy Design of Intermetallic Dispersion Strengthened Aluminum Systems by Mechanical Alloying Table 2. MEAM parameters for A1-Tiand AI-Zr systems A-B A1-Ti A1-Zr

OtA~ 5.0 5.0

A,,. 2.73 2.85

RA~ 0.5 0.5

p!~ 1.0 1.0

p~ 1.9 1.2

]9

Table 3. Calculated energy for separating (111) interface and bulk modulus Fcc

A1 A13Ti ALZr

D023

D022

S

B

897

0.48

S

B

S

L12 B

S

B

1126 0.76 1211 0.75 1058 0.75 1146 0.87 1168 0.85 1249 0.78

S:energyforseparatinginmfface(erg/cm2),B:bulkmodulus(eV/A3)

Fig. 2. Calculated energy and force for separating (111) interlace between AI and ALTi.

resistance to deformation and good interfacial properties. Thermodynamic calculations play a significant role as guide in the phase selection because transition compositions in A13(Ti, V, Zr) determine if a single phase field of D022 or D023 or a two-phase field of D022+D023 is stable while MEAM estimates interfacial properties and the measure of deformation. The phase equilibria of the A13 (Ti, V, Zr) system have been studied by using thermodynamic parameters assessed by the authors [17]. A brief explanation for numerical expressions of thermodynamic models follows. (4)

G P = G r e f - T A S id + A G ex

Firstly, the total partial background electron density is calculated using Eq. 3, which is made up of four components including angle dependent terms, and secondly the total background electron density P is obtained using each total partial background electron density and then the embedding energy is calculated using Eq. 2. The pair energy can be obtained by using Roses universal energy function [16]. Reported parameters for A1, Ti and Zr [15] were used and interaction parameters for AI3Ti and ALZr were determined in this work as shown in Table 2. Detailed expressions for various MEAM parameters and calculation methods can be found in the work of Baskes [15]. Fig. 2 shows the calculated energy for the separating (111) interface between A1 and ALTi. Calculations were repeated at each location as they move apart from each other starting from the interface step by step. The force curve shows the location of the total consumed energy for the separating interface and maximum critical force. The energy at the location in Fig. 2 where no force exists is equivalent to the interface energy. Strictly speaking, since the estimated interface energy is the difference between non-separated and separated interfaces, a high value means a high stability of interface and large resistance to separation. Calculations of other structures and ALZr also show similar behavior. The bulk modulus was calculated, which explains the degree of resistance to deformation. Table 3 implies that intermetallic compounds have a large

General thermodynamic models describing the Gibbs free energy of a certain phase, P, are expressed in Eq. 4. Proper models for AldTi, V, Zr) may be established considering two sublattices in which one sublattice consists of an A1 site and the other of an M (M=Ti, V, Zr) site. The following Eqs. 5-8 are applicable to the two sublattice model. For convenience, the (Ti, V, Zr) site is designated by Ti site. Gref Ti Ti Ti 9 = YTiGTi.'AI + YV GV:AI + YzrGzr:AI

(5)

GTi..AI=&GfAI3T i

(6)

GV..AI=A~I3V

GZr..AI=A~I,Zr

Y~: site fraction of j atom in M sublattice id

Ti

Ti

Ti.

Ti

Ti 9

Ti

AS =--0.25R(YTilnYTi+YV mYv + Yz,.mY'zr)

(7)

AGeX

Ti Ti Ti Ti g7 = YTiYZrGTLzr.AI + YV YZr~V, Zr.'AI 0 GTi, Zr:AI = GTi, Zr:AI

0 yTi . Ti, ~1 GV, Zr.'AI = GV, Zr.'AI + ( V - YZr)CiV, Zr:AI

(8)

G,,N in Eq. 5 represents the Gibbs free energy of formation when M atoms occupy the (Ti, V, Zr) site and N atoms occupy the (A1) site. Interaction parameters (AG~') are effective among Ti, V and Zr at the (Ti, V, Zr) site. Underlined terms were originally unknown and therefore were assessed by the authors [17,18]. Reported values were used for Gibbs free energy of formation for stable

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Sung I1 Park and Hyuck Mo Lee

Fig. 3. Simulated A13Ti-A13V-AI3Zrpseudo ternary phase diagram at 1273K by using the thermodynamic variables shown in Table 4.

Table 4. Assessed thermodynamic variables, from ref. [ 17,18] Element

A1

Ti

V

Zr

Reference state

Fcc

Hop

Bcc

Hcp

Phase

D022

D02~

Parameter Value [J/gram atom K]

AG-fALZr *

-39293+ 11.59"T

0 Gv, Zr:AI

10051-6.9347"T

1

GV, Zr:A l

-72

0 --(TrZr,ri:A1

-4188 2.3715*T

AG4Al Ti *

-38666+10.34"T

AGfl v *

-22678+3.80"T

o GV, Zr.'AI

- 1355+4.5388"T

i GV, Zr:A I

-341

0 Gzr, Ti:AI

-5464+ 1.856"T

*: counter phase

phases [19,20]. Fig. 3 shows the A13Ti-A13V-AI~Zr phase diagram calculated at 1273K in this way. Intermetallic compounds of D022 or D023 can be fabricated on the basis of the original composition. The L12 phase diagram was not calculated because any composition in the whole range of Ti and Zr did not result in a stable L12 phase field.

4. EXPERIMENTAL PROCEDURES Three types of intermetallic compounds

-- D023-A13(Tio.4

V0x~Zro56), D022-A13(Tio.lV0.sZro.,) and L12-A166Mn9(Tio4Zro6)2s

--were prepared by arc melting from pure elements, A1 (99.995%), Ti (99.98%), V (99.9%) and Zr (99.2%) and Mn (99.8%). Specimens cut from the buttons were encapsulated in silica tubes that were back-filled with argon gas. Encapsulated specimens were heat treated at 1273K for 24 hr and then quenched in water. The resulting monolithic intermetallic alloys were pulverized and then mixed with pure A1 powders. The ball milling was performed in an ambient environment with 3/16 inch stainless steel balls by an attritor mill. About 60 g of powder mixture was loaded with a 1/20 weight ratio of ball to powder. A 2% stearic acid was added as a process control agent. The amount of intermetallic powders was designed to maintain the volume fraction of the second phase at about 20% when mixed with A1 powders. Mechanically alloyed powders were compacted and hot pressed under vacuum at 698K and 100MPa for 10 min. Final products are of cylinder shapes with a diameter of l lmm and a height of 20 mm. They were heat-treated at 698K for 100, 200 and 400 hr to investigate the thermal stability of the second phase. SEM (Scanning Electron Microscopy) was employed to observe the microstructure. Phases were identified by XRD (X-ray Diffractometry) analyses. Specimens were cut into a tetragonal shape the size of 3x3• (ram 3) and were tested in compression with a strain rate of 2x104s -' at room temperature and at 693K, which is 75% of the absolute melting temperature of pure A1. 5. R E S U L T S Specimens heat treated at 1273K for 24hr and quenched in water were analyzed by XRD and identified to have structures of D022, D023 and L12. Fig. 4 shows such typical patterns. Among these, vacuum hot pressed A1-A166Mn9 (Tio~Zr06)25 composites were heat treated at 698K for 100, 200 and 400 hr. The back scattered SEM images are shown in Fig. 5. Overall, there is no appreciable change in microstructures with time, which implies that the thermal stability is rather high mainly due to the high melting point of the intermetallic L12 phase. White regions observed all around the specimen were identified to be a transition element rich zone by EDX (Energy Dispersive X-Ray). Dark regions were identified to be an AI rich zone. Figs. 6-8 show XRD patterns of AI- D023 A13(rio.4go.04 7x,,56), AI- L12 Al,~Mn~(Ti,,4Zro&,and A1- D022Al3(rlo.lgo.eZro.,), respectively, under several conditions. They reveal the presence of intermetallic compounds after vacuum hot pressing and it has been confirmed that a new phase has not formed. In the mechanically alloyed powders, it is probable that severe strain may prohibit the detection of

Alloy Design of Intermetallic Dispersion Strengthened Aluminum Systems by Mechanical Alloying

2]

Fig. 4. XRD patterns of intermetallic compound D0~_3Al3(T~/40VowZro56), L12 Al~,6Mm(Tio~Zro025and D022 AlffTi,,,V~Zro0 produced by heattreating at 1273K for 24hr, followed by water quenching.

Fig. 6. XRD patterns of the A1-D0_~Al~(Ti04,No.o4Zr,,50alloy depending upon processing condition and heat-treating time.

Fig. 5. Microstructures of the vacuum hot pressed and heat-treated (back-scattered image) A1-LI: Ak6Mng(Tio4Zr~,02, alloy at 698K: a) VHP, b) 100hr, c) 200hr and d) 400hr.

-ray peaks related with intermetallic regions. This is why the peaks of the intermetallic phases are weak compared to the peaks of intermetallic powders only. The separation of L12 peaks with heat treatment may be vaguely observed in Fig. 7. The enlarged area of separated peaks is clearly shown in Fig. 9. This has been brought about by the existence of two L 1, phases with different lat-

rice parameters, one being a Ti-rich phase and the other a Zr-rich phase. Two L12 phases originated from the effect of the compositional segregation of Ti and Zr in the LL intermetallic compounds when they were fabricated as bulk, as was reported earlier [21]. Fig. 10 shows the compression test of the vacuum hot pressed A1- D02~A13(Tio.4Vo.04Zro~)alloy and A1- L12 A166 Mng(Ti04Zr0.6)25 alloy at room temperature and 693K. Compressive yield strengths were measured to be 920 and 800MPa at room temperature, respectively, and 260 and 190MPa at 693K, respectively. Maximum compressive strains are 0.06 at room temperature and 0.27 at 693K in case of A1- D023A13(Ti0.4V0ooZr050alloy and 0.11 at room temperature and 0.31 at 693K in the case of the A1- LI~_A16~Mng(Ti0~Zr0025alloy. Table 5 is a comparison between this study and published results of mechanically alloyed A1 alloys. As shown in Table 5, the mechanical properties of intermetallic-dispersionstrengthened A1 alloys are as good as those in other research and could be improved through the proper control of the volume fraction of the second phase and the adoption of a proper fabrication process, for example, hot extrusion.

22

Sung II Park and Hyuck 340 Lee

Fig. 9. Detailed XRD patterns of (100) plane in the L12 AIo6Mn,

(Ti04Zl~&~intennetallic phase.

Fig. 7. XRD patterns of the A1-L12 Ale6Mng(Tio4Zro6)2salloy depending upon processing condition and heat-treating time.

Fig. 10. Compressivestress-strain curve of A1 alloys.

6. D I S C U S S I O N

Fig. 8. XRD patterns of the AI-D022A13(Ti,~,V0~Zr,,,)alloy depending upon processing condition and heat-treating time.

The origin of the high strength of intermetallic-dispersion-strengthened AI alloys at high temperatures can be explained as follows. The main source of the high strength comes from the high strength of intermetallic compounds, which means a large resistance to deformation, a strengthening effect by carbides formed by impurity and the homogeneous fine distribution of second phases. High thermal stability guarantees the maintenance of the fine distribution of the second phase without decomposition, coarsening and transformation, which guarantees no degradation of mechanical properties at high temperatures. The type of second phase was chosen based on the simulated phase diagram that was obtained through thermodynamic calculations. This is important because structures of second phases do not change with vacuum hot pressing and heat treatment by the peculiar fabrication method devised in this study. This enables alloy designers to control the properties of A1 alloys. As inferred by EAM

Alloy Design of Intermetallic Dispersion Strengthened Aluminum Systems by Mechanical Alloying

23

Table 5. Comparisonof A1 alloy (MA)

Dispersoid CompressiveYield Strength (MPa)

Volume Fraction of Dispersoid

Reference

Process

Matrix

Srinivasan et al. [9]

A1powder+Tipowder MA--->VHP

Al-base

AI~Ti

600MPa at 25~ 100MPa at 500~ (0.0002s.l)

40%

Lee and Moon [11]

AI powder+Tipowder MA--->VHP--~ Hot Extrusion

Al-base

A13Ti

500MPa at 25"C t67MPa at 400"C (0.0005s ~)

20%

This study

A1 powder+ AI3(Ti,V,Zr) powder ---~MA---~VHP

Al-base

AI3(Ti,V,Zr)

920MPa at 25"C 260MPa at 420~ (0.0002s ')

20%

This study

A1powder+ AI~,,,Mng(Ti, Zr) powder -->MA-->VHP

800MPa at 25~ 190MPa at 420~ (0.0002s ~)

20%

Al-base AI~6Mng(Ti,Zr)

calculations, the large resistance to deformation increases the strength of A1 alloys while interfacial properties are not easily verified. However, Fig. 10 casts a hint that high temperature strength implies that a small portion of deformation is being carried by intermetallic compounds and that high ductility shows good transfer of deformation crossing interfaces. This behavior is not possible without good bonding at the interface. Low ductility at room temperature may be due to the high density of defects and voids and not due to the bonding at the interface, which of course needs further study. In this work MEAM calculations with compositional variations have not been made because of the technical problems that relate to the reliability of energy and of extension to multi-component systems that have not yet been resolved. These problems are expected to be overcome soon and more sophisticated calculations with compositional variations will improve the alloy design work.

7. SUMMARY The interface energy and resistance to deformation of A13M phases as a dispersoid in the A1 matrix were estimated through MEAM calculation. Selection of transition element compositions in A13(Ti, V, Zr) compounds was determined on the basis of the simulated phase diagram through thermodynamic calculations. The core idea of intermetallic-dispersion-strengthened MA is the fabrication of intermetallic compound powders that were chosen by theoretical calculations and the mixing of them with A1 powders. Second phases were finely distributed and thermally stable in the vacuum hot pressed conditions. In compression tests intermetallic-dispersion-strengthened A1 alloys show high compressive yield strengths of 800-920MPa at

room temperature and 190-260MPa at 693K, and high compressive ductility of 0.27-0.31 at 693K. The origin of enhanced mechanical properties is high strength, high thermal stability, large volume fraction and fine distribution of intermetallic compounds.

ACKNOWLEDGMENT This study has been financially supported by the Center for Interface Science and Engineering of Materials (CISEM) at KAIST.

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