The Role of cobalt and nickel intermetallic phases on the mechanical properties and microstructure evolution of Al-Zn-Mg-Cu alloys

Abstract

In this study, the effects of cobalt and nickel additions on microstructure evolution and mechanical properties of Al-Zn-Mg-Cu alloys produced by chill casting process were investigated. Al-Zn-Mg-Cu aluminum alloys containing Ni and Co additives were homogenized at different heat treatment conditions, aged at 120ºC for 24 h, and retrogressed at 180ºC for 30 min then re-aged at 120ºC for 24h. Comparison of the ultimate tensile strength of the Al-Zn-Mg-Cu base alloy underwent RRA temper process with that of the base alloy underwent a similar heat treatment process after the addition of 0.5 wt. % Ni and 0.5 wt. % Co showed that the gains of tensile strengths of the orders of 100 MPa and 135 MPa were attained respectively. These improvements are attributed to phase dispersion particles of the formation of Ni and Co-rich such as Al7Cu4Ni, Al75Ni10Fe15, Al3Ni2 and Al70Co15Ni15. The dispersed particles of nickel or/and cobalt have improved the ultimate tensile strength to reach 630 MPa and 665 MPa and the hardness HV values to 250 MPa and 265 MPa respectively. Besides dispersion, Ni and Co intermetallic dispersoids retarded the grain growth, led to grain refinement and precipitation hardening and resulted in further strengthening of the Al alloy. Microstructure characterization of the alloys was carried out using optical microscopy (OM), scanning electron microscopy (SEM), energy dispersive X-ray spectroscopy (EDS) and X-ray diffraction (XRD).

Al-Zn-Mg-Cu alloys; nickel additives; Dispersion; RRA treatments

REGULAR ARTICLES

The Role of cobalt and nickel intermetallic phases on the mechanical properties and microstructure evolution of Al-Zn-Mg-Cu alloys

Haider Tawfiq Naeem^*; Kahtan Sadeq Mohammed; Khairel Rafiez Ahmad

School of Materials Engineering, University Malaysia Perils, Muhibbah, 02600 Jejaw, Perlis, Malaysia

ABSTRACT

In this study, the effects of cobalt and nickel additions on microstructure evolution and mechanical properties of Al-Zn-Mg-Cu alloys produced by chill casting process were investigated. Al-Zn-Mg-Cu aluminum alloys containing Ni and Co additives were homogenized at different heat treatment conditions, aged at 120°C for 24 h, and retrogressed at 180°C for 30 min then re-aged at 120°C for 24h. Comparison of the ultimate tensile strength of the Al-Zn-Mg-Cu base alloy underwent RRA temper process with that of the base alloy underwent a similar heat treatment process after the addition of 0.5 wt. % Ni and 0.5 wt. % Co showed that the gains of tensile strengths of the orders of 100 MPa and 135 MPa were attained respectively. These improvements are attributed to phase dispersion particles of the formation of Ni and Co-rich such as Al₇Cu₄Ni, Al₇₅Ni₁₀Fe₁₅, Al₃Ni₂ and Al₇₀Co₁₅Ni₁₅. The dispersed particles of nickel or/and cobalt have improved the ultimate tensile strength to reach 630 MPa and 665 MPa and the hardness HV values to 250 MPa and 265 MPa respectively. Besides dispersion, Ni and Co intermetallic dispersoids retarded the grain growth, led to grain refinement and precipitation hardening and resulted in further strengthening of the Al alloy. Microstructure characterization of the alloys was carried out using optical microscopy (OM), scanning electron microscopy (SEM), energy dispersive X-ray spectroscopy (EDS) and X-ray diffraction (XRD).

Keywords: Al-Zn-Mg-Cu alloys, nickel additives, Dispersion, RRA treatments

1. INTRODUCTION

Due to rapid development in aerospace, military industry and ground transportation vehicles, persistent efforts have been exerted to improve the mechanical properties of Al-Zn-Mg-Cu alloys to satisfy the application challenges. The requirement was to go to high strength tough alloys. Working at high temperature properties implies good stress corrosion cracking resistance^1-5. In order to achieve this goal, microalloying with transition or rare earth elements such as Ni, Co, Sc, Zr, Yb, Cr and Er^1,6-9 is an effective method in attaining this objective. Addition of transition elements to high strength aluminum alloys enhances their mechanical properties. All the strengthening mechanisms of the high strength aluminum alloys depend primarily on precipitation hardening^10,11. It is found that trace elements in aluminum can form a supplementary strengthening afforded in the most cases by dispersions which improve the recrystallization resistance and introduce a controlled amount of the strain^5,6. The strong bonding of the second phase to the matrix is very important in impeding dislocation motion through the matrix and hence in strength enhancement^12,13. Recently, many researchers have been focusing on the rare earth metals. Wenbin et al.¹⁴ found that the addition of scandium (Sc) to the Al-Zn-Mg-Cu-Zr alloy led to ultimate tensile strength improvement. The Sc as Al₃(Sc, Zr) particles are effective in refining. Seyed et al.¹⁵ observed that the addition of Zr to the new super high strength (Al-12.24Zn-3.25Mg-2.4Cu) aluminum alloy decreases the average grain size, causes uniform dispersion of its constituent. Owing to the formation of the Al₃Zr particles, it limits the area fraction of recrystallization and eventually improves the UTS and the yield strength (YS).

However, the applications of the rare earth metals-containing alloys are extremely restricted due to their high cost. Therefore, more attention has been paid to the transition metals such as nickel and cobalt since they are cheap. Additionally, it was found that phases containing cobalt and nickel plus iron were more resistant to coarsening than phases containing Cr, Ti and Zr^7,12,16. Harder Al matrix and accelerated artificial aging rate increase the hardness peak. However, the influence of additions nickel on the mechanical properties of Al-Zn-Mg-Cu alloys was investigate by numerous the techniques such as; Ingot metallurgy and the rapid solidification, Lei et al.¹⁷ found that additions of nickel with the Al-Zn-Mg-Cu alloy (C912) which produced by ingot metallurgy, Ni promotes the development of the matrix precipitates besides enhanced good the strength and stress corrosion crack resistance. On the other hand, Liang et al.¹⁸ reported that addition of nickel to an Al-Zn-Mg-Cu aluminum alloy produced by spraying deposition resulted in Ni-rich particles of Al₃Ni_2, Al₇Cu₄Ni and MgNi₂ dispersoids which enormously attribute to the precipitation hardening of the alloy. To the best of knowledge of the authors, the effect of nickel and cobalt micro - alloying additions coupled with semi-direct chill casting process on the microstructure evolution of Al - Zn - Mg - Cu aluminum alloys has not been studied previously. The objective of this work is to investigate the influence of nickel and combination of nickel plus cobalt elements on mechanical properties and microstructure evolution of new ultra-high strength aluminum alloys under different heat treatments.

2. EXPERIMENTAL PROCEDURES

2.1. Research material

The present study was carried out on Al-Zn-Mg-Cu alloy ingots provided by ALCAN GLOBAL AEROSPACE. The slabs were 13 mm thick and 20 mm wide. Nickel and Cobalt of 99% purity both as additives were provided by Merck KGaA. The nominal compositions of the studied alloys are listed in Table 1.The terms "Base alloy", "Alloy 1" and "Alloy 2" refer to as received alloy, alloy with 0.5 wt . % Ni and alloy with 0.5 wt . % Ni plus 0.5 wt . % Co additions respectively. The chemical composition analysis was carried out using the arc-spark spectrometer.

The additions of cobalt and nickel were chosen after different preliminary amounts of additions within the experiment reaches its sufficient state, as in our previous studies¹⁹ where it was used 0.1 Ni % due to the solubility nickel is limited in aluminum matrix in a solid state and the required inter compounds precipitates formed as shown in the XRD plots in Figure 1 and Figure 2.

The three alloy types were re-melted in a graphite crucible at 830°C in electrical resistance furnace (with accuracy about ±5°C). The samples were produced by a semi - direct chilling (DC) casting process conducted in a cylindrical iron steel mold of Øin - out of 35-45 mm and height of 150 mm. The mold was preheated to 250°C prior to the casting process. After casting, the alloys were homogenizing treated according to step No.1 in Table 2, followed by quenching in cold water immediately after each step of the homogenizing treatment. After quenching, the extrusion process was performed on the cast bars. To do so, the alloy bars and the extrusion mold were preheated to 450°C for 30 min and to 400°C in electrical resistance furnaces. After extrusion, the samples were quenched in cold water. The extrusion rate was about 3.5 and waxes was used as a lubricant material during the extrusion process. The extruded bars were held separately at 450°C and 470°C for 1 h, and then kept at 480°C for 30 min for further homogenizing treatment. Finally they were quenched in cold water immediately after each step. In this study all extruded samples were tempered according to the procedure of T6 and RRA temper in step No.2 & 3 in Table 2. After each step of the aging at T6 temper and the retrogression and reaging (RRA), the specimens were quenched in cold water.

The homogenizing and heat treatments carried out in this study have been chosen according to the optimize treatments in the previous studies^20-22.

2.2. Microstructures characterizations

The microstructures were analyzed by the optical microscopy (OM) using Olympus PMG3 optical microscope. The specimens were extracted from position of ½ height of the ingot, ground and polished according to the standard methodology method. They were etched with Keller's reagent. The grain size analysis was carried out using the linear intercept method . In order to investigate the effects of additives on the microstructure scanning electron microscopy (SEM) coupled energy dispersive X-ray spectroscopy (EDS) and X-ray diffraction analysis (XRD) were used.

2.3. Mechanical testing

The HV microhardness measurements were carried out on the specimens using microhardness tester "Mitutoyo - model DX256 series". Indentation force was set to 30N and 10 second dwelled time. To ensure clean liness the surfaces of the extruded samples were polished prior to HV measurement. Each reading was an average of at least ten separate measurements taken randomly. The highest and the lowest values of the ten readings were disregarded. The tensile test was carried out at ambient temperature on round tensile specimens using an INSTRON testing machine at a ram speed of 10 mm/min, and a load of 500 kN. The tensile test specimens were prepared according to the standard methodology method.

3. Results

The optical micrographs in Figure 3a, b and c reveal the microstructure for the samples of the as cast; base alloy, alloy 1 and 2 respectively. They consist of primary dendrites of aluminum-rich solid solution and interdendritic networks of intermetallic compounds around the primary grains. Additionally, there are fine equiaxed grains encompassing columnar grains between them.

However, as shown in Figure 3a, the same direct-chill (DC) casting of the base alloy sample has a fine equiaxed grain structure throughout their entire cross sections of about 45µm as an average grain size. Previous studies on conventional casting conducted by Yongdon et al.²³ gave grain size of about 250mm for as-cast Al-Zn-Mg-Cu alloy. In another study, Ying et al.²⁴ reported a grain size of 121mm for the as-cast 7050 Aluminum alloy. The outcome of the semi direct-chill (D.C) casting process adopted in this study proved its effectiveness on grain refinement and eventually on the mechanical property improvement. The average grain size of the as-cast samples was about 38 and 31 mm for alloy 1 and alloy 2 respectively. The grain size refinement of the as quenched alloy 1 and alloy 2 is attributed to the additions of both Ni and/or the dual Ni with Co elements to the aluminum melt which acted as nucleation centers for the grains. These additions block the columnar growth which usually starts from the mold wall and forms favorable sites for the nucleation of the first forming aluminum phases during the solidification process^25,26.

The optical micrographs of the microstructures for the samples underwent hot extrusion combined with T6 temper (which is similar to RRA) are shown in Figure 4a, b and c respectively, they indicated that the average grain size for these microstructures was measured to be about 37, 31 and 26 mm for base alloy, alloy 1 and alloy 2 respectively. The reduction in the grain size of the three samples was a consequence of the extensive hot extrusion process. Furthermore, the interaction of the T6 temper and the added nickel and cobalt has enhanced the grain size reduction wherein additive compounds have formed nucleation sites and promote high solidification rate.

Figure 5 shows different the X-ray diffraction (XRD) patterns for the samples of the base alloy; as-cast, T6 and RRA. It can be observed that the as quenched alloy is mainly composed of α-(Al) solid solution and other secondary phases. These are mainly; T-AlMg₄Zn₁₁, S-Al₂CuMg, and η- MgZn₂ phases. Generally the accepted precipitates' formation sequence for the 7000 series alloys are as reported in^27,28. Metastable the h′phase stands as the main strengthening precipitate for these alloys. The main precipitations in the matrix after 120 °C for 24 h aging treatments are GP zones and h′phase. As shown in the XRD plots in Figure 5b, the sample T6 of the base alloy reveals the increasing existence of the η-phase of MgZn₂ and h′-phase of Mg₂Zn₁₁ which dissolved during homogenizing then precipitated during aging at T6. The XRD plot of the base alloy RRA processed sample is shown in Figure 5a. In this plot the MgZn₂ and Mg₂Zn₁₁ phases peaked higher. The main precipitation phases within the RRA processes have been detailed in our reports^29,30. Finally, the number of the nucleus that can promote re-precipitation of the GP zones and h′ phase in the RRA is increasing as shown in the XRD results in Figure 5a the major peaks of the h′ phase, through the present study. On the other hand, rapid quenching rate from above the solvus line forms metastable supersaturated solid solutions which in turn increase the GP zones precipitation and its compounds³¹.

The SEM image in Figure 6a and b depict the microstructure of the as-cast sample for the base alloy. The dark areas are the primary solid solution, and the bright areas are the non-equilibrium eutectic solid solution between grains which appears as a lattice pattern in the enlarged image in Figure 6c. There are dispersal gray particles indicated by the encircled region (f) representing the stoichiometry close to in the EDX analysis in Figure 6d. The labeled region in Figure 6a represents the matrix of the base alloy which contains the α(Al) eutectic phase. The EDS scan analysis in Figure 6b reveals the chemical composition for this labeled region wherein the elements suppose to be close to the T-(AlMg₄Zn₁₁) and S-(Al₂CuMg) phases.

The SEM image and EDS scan analysis in Figure 7a and b depict the microstructure of the base alloy T6sample. After various processes including: homogenization, extrusion, solution treatment and ageing at 120°C for 24h, the SEM image reveals the existence of precipitates (bright spots) indicated by the encircled region (g) in the matrix as clarified by the supplemented enlarged image in Figure 7c. The EDS scan analysis results for the labeled region indicate as shown in Figure 7b.

The SEM and the EDS scan analysis in Figure 8 depict the microstructure of the base alloy sample after undergoing an the RRA process. The SEM image depicts the existence of precipitates (bright spots) indicated by the encircled region (h) in the matrix as clarified by the SEM enlarged image in Figure 8c. The EDS analysis reveals the chemical composition of these bright regions. They are probably the stoichiometry for this labeled region close to the T-AlMg₄Zn₁₁, S-Al₂CuMg, η-phase MgZn₂ and h′- Mg₂Zn₁₁ phase compounds.

Figure 1 shows different the XRD patterns for alloy 1 samples i.e., the as quenched, the T6 and the RRA. The patterns for the as-cast alloy1; sample confirms that the primary eutectic is mainly consisted of α(Al) solid solution and intermetallic compounds of the T-Al₅Mg₁₁Zn₄, and Al₃Ni₂ phases. The XRD plots of alloy1 T6 sample (b) in Figure 1 reveal the existence of dispersoid compound phases of Al₅Mg₁₁Zn₄, Al₇₅Ni₁₀Fe₁₅, Al₃Ni₂, MgZn₂ and Mg₂Zn₁₁ . These dispersive phases peaked high due to the intensive dissolution of the alloying elements and the nickel additive in the alloy brought about by the extrusion process and the subsequent heat treatments as shown by the T6 temper in Figure 1b.

Li et al.³² proposed that Ni addition to Al-Zn-Mg-Cu aluminum alloys suppresses the formation of the MgZn₂. The outcome of this study came on the contrary and provides compelling evidence on the existence of the MgZn₂ phase as shown in Figure 1a and b by the RRA and the T6 respectively.

The RRA process has resulted in high diffraction dispersed precipitate phase peak of Mg₂Zn₁₁ phase as shown in the XRD plots in Figure 1a.

The SEM and EDS scan analysis in Figure 9 depict the microstructure of the as-cast sample of alloy1. The dark regions are the primary solid solution and the bright areas are the non-equilibrium intergranular eutectic solution as indicated in Figure 9a, its detail appeared in the supplemented high magnification micrograph in Figure 9c . Furthermore, the EDS scan analysis in Figure 9b reveals the stoichiometry of the bright encircled area.

The SEM image and the EDS scan analysis in Figure 10 depict the microstructure of alloy 1 T6 temper sample. The distinguished bright region illustrated by encircled region (k) in the SEM image is an elongated rod like shape phase in Figure 10c. This phase is explicated by the supplemented high magnification SEM image and the EDS scan analysis in Figure 10b. The EDS scan result of this labeled region assumes its affiliation to the stoichiometry of the Al₃Ni₂, MgZn₂ and Mg₂Zn₁₁ phases. This assumption is also supported by the XRD analysis in Figure 1. During the extrusion process the precipitates elongated aligned along the grain boundaries and reoriented along the extrusion direction. The SEM image in Figure 11a depicts the morphology of alloy1 RRA sample. The encircled region (i) clarified by the supplemented enlarged SEM image representing chained precipitates of flaky structure is shown in Figure 11c. The EDS scan analysis of this region in Figure 11b.

Figure 2 shows XRD plots for different samples of the alloy 2; the as-cast, the T6 and the RRA. The plot of the as quenched sample of alloy 2 in Figure 2c is mainly consists of Al₅Mg₁₁Zn₄, Al₇Cu₄Ni, MgZn₂ phases. The XRD plots of the alloy 2 T6 sample in Figure 2 breveals the appearance of new dispersed compounds of the Mg₂Zn₁₁, Al₃Ni₂, Al₇₅Ni₁₀Fe₁₅, Al₃Co, Al₇₀Co₁₅Ni, Al₅₀Mg₄₈Ni₇ phases. The XRD plot of alloy2 RRA sample in Figure 2a showed that the RRA process has resulted in high diffraction peaks of the dispersed precipitate the Mg₂Zn₁₁ and Al₃Ni₂ phases.

The SEM in Figure 12 adepicts the microstructure of the as-cast alloy2 sample. The dark regions are the primary solid solution, and the bright regions are the non-equilibrium intergranular eutectic solution as elucidated by the enlarged micrograph in Figure 12c. The EDS scan analysis of the bright encircled region (m) shows its chemical composition.

The SEM image in Figure 13 adepicts the microstructure of alloy2 T6 sample. In this image there are two encircled regions (m) and (q) . Region (q) represents the matrix. The EDS scan analysis results of this region in Figure 13b supports the likelihood of its attribution to the chemical composition. The region (m) explicated by the enlarged SEM image in Figure 13c exhibits a rod like shape phase. The EDS scan analysis in Figure 13d of this region. In the extrusion process the grains elongated, aligned and reoriented along the extrusion direction.

In Figure 14a and b defined by the regions (s) and (p), the rod like shape phase in a region (s) explicated by the supplemented enlarged SEM image in Figure 14c. The EDS analysis in Figure 15d of this region are related to the stoichiometry of phases. The dispersed particle in the encircled region (p) is explicated by the supplemented high magnification SEM image in Figure 14e. Figure 14b indicated the EDS analysis in for this labeled region . The EDS results of the samples under T6 temper and the RRA process are consistent with the results XRD in Figure 2.

Figure 15

Figure 16 reveals the variations in the Vickers hardness values of the three examined alloys after different heat treatment conditions. After performing the T6 and the RRA heat treatments on the three alloys samples extruded, maximum gains of about (110 and 120) MPa were attained in UTS for all alloys respectively. Principally the hardness scale depends on the extrusion process which led to reduced the porosity and grain size as well as the ageing precipitation whose impact is included in the distribution of precipitates or dispersion phases of additives in the matrix during these processes.

4. DISCUSSION

The results indicated that substantial improvements in the YS, UTS and Hv hardness of the base alloy were attained after following certain heat and mechanical treatments of T6 temper, the RRA and extrusion processes. The strengthening mechanism of the ageing at T6 (120°C for 24h) for the base (Al-Zn-Cu-Mg) alloy is attributed to the precipitation hardening i.e., the effects of the Guinier-Preston zones coherent with nanometer sized h′ metastable precipitates. These precipitates act as pinning points which impede the dislocation movement^31-33. The XRD and the EDS results have revealed the existence of the η and η′ phases. The YS and UTS of the base alloy sample after RRA process showed a distinguished improvement as opposed to that of the T6 temper specimen. This improvement is attributed to the partial dissolution of the pre-existing GP zones and h′ phase. The GP zones can act as nucleation sites for h′ particles and the remaining h′ phase carries on growing during the RRA process. Additionally, after the RRA treatment the solute atoms dissolve in the matrix, precipitate again and produce more small size GP zones and h′ phase. Hence it led to the augmentation of the phases η and h′ by the RRA process more than that which resulted from the T6 process.

The strengthening mechanisms achieved by the addition of the Ni and Co microalloying elements into the Al-Zn-Cu-Mg alloy can be divided into precipitation strengthening derived from the alloying elements in the base alloy already explained and the dispersoid and fine-grain strengthening derived from the Ni and Co - rich particles subjected to the extensive extrusion process. In the later strengthening mechanism the dislocations are hindered by the Ni or/and Co dispersed within the slipping planes. The hindering level depends on the even distribution of the particle phase, size, spacing and degree of coherency. The dispersoid phase particles are looped, by passed and/or sheared by the dislocation through the Orowan mechanism³⁴. Moreover, The fine grains of Ni or/and Co contained dispersoids restrict recrystallization and inhibit the grain growth as shown in Figure 3b and c. Furthermore, the extensive extrusion process has led to grain size reduction as shown in Figure 4b and c. Thus grain refinement according to Hall-Petch equation, means higher material strengths^5,9.

5. CONCLUSIONS

1. The RRA process for the base alloy improved the yield strength more than what the T6 temper did; it pushed the ultimate tensile strength to a higher level. UTS of 530 MPa and Hv hardness value of 223 were attained by the application of this process.

2. Regarding the mechanical properties, the YS, UTS and the Hv hardness for the base alloy sample of 0.5 wt. % Ni additions underwent RRA process increased by 95 MPa, 100 MPa and 27 MPa respectively.

3. With appropriate microalloying additions of 0.5 wt. % nickel plus 0.5 wt. % of cobalt, alloy2 after RRA process exhibited the highest strength and hardness values. YS of 615 MPa, UTS of 665 MPa and Hv hardness of 265MPa were attained. The rapid increase in strength and hardness was due to the Ni+Co dispersed of Al₇₀Co₁₅Ni₁₅ phase which have high hardness and strength.

4. Generally, the incremental increase in the strength of alloys 1 band 2 in the present study was due to different factors working together. These are precipitation hardening, Orowan strengthening and the Hall-Petch strengthening resulted from grain size refinement brought about by the heat and mechanical treatments.

ACKNOWLEDGEMENTS

This work is supported under the grant No. 9001-00338 of the Universiti Malaysia Perlis (UniMAP). The authors gratefully acknowledge the outstanding support provided by the technicians of the workshop in the Materials Engineering School, UniMAP.

Received: July 29, 2014

Revised: November 12, 2014

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Publication Dates

• Publication in this collection
  10 Feb 2015
• Date of issue
  Dec 2014

History

• Received
  29 July 2014
• Reviewed
  12 Nov 2014